INDUSTRIAL CORROSION AND
CORROSION CONTROL TECHNOLOGY
H.M. Shalaby
A. Al-Hashem
M. Lowther
J. Al-Besharah
(Editors)
Published By
Kuwait Institute for Scientific Research
INDUSTRIAL CORROSION AND CORROSION CONTROL
TECHNOLOGY
1996
Sponsored by the Kuwait Institute for Scientific Research (KISR), the Kuwait
Foundation for the Advancement of Science (KFAS), the Kuwait National
Petroleum Company (KNPC), the Kuwait Oil Company (KOC), Ministry of
Electricity and Water (MEW), Kuwait University (KU), Ministry of Oil (MO),
the Gulf Cooperation Council-General Secretariat (GCC), Kuwait Chemical
Society (KCS), Organization of Arab Petroleum Exporting Countries (OAPEC),
and Petrochemical Industries Company (PIC).
INDUSTRIAL CORROSION AND CORROSION
CONTROL TECHNOLOY
Proceedings of the 2nd Arabian Corrosion Conference
Kuwait, October 12-15, 1996
Editors
H.M. Shalaby, A. Al-Hashem, M. Lowther and J. Al-Besharah
PUBLISHED BY KUWAIT INSTITUTE FOR SCIENTIFIC RESEARCH
P.O. BOX 24885, 13109 SAFAT, KUWAIT
Published by
Kuwait Institute for Scientific Research
P.O. Box 24885, 13109 Safat, Kuwait
Publication Number: KISR 4890
Copyright ® 1996 by Kuwait Institute for Scientific Research
The papers were reviewed for their technical contents. Editing was restricted to matters of
format, general organization and retyping. The editors assume no responsibility for the
accuracy, completeness or usefulness of the information disclosed in this book. Unauthorized
use might infringe on privately owned patents of publication right. Please contact the
individual authors for permission to reprint or otherwise use information from their papers
This book was printed in Kuwait
The 2nd Arabian Corrosion Conference
FOREWORD
The 2nd Arabian Corrosion Conference was held in the state of Kuwait during the period 1215 October, 1996 under the auspices of H.H. Sheikh Saad Al-Abdullah Al-Salem Al-Sabah,
Kuwait’s Crown Prince and Prime Minister. The present conference was scheduled to be
held in Kuwait during 27-30 April, 1991, however, it was postponed due to the events that
encompassed Kuwait and the Gulf region in 1990-1991.
The 1st Arabian Corrosion Conference was held in Kuwait during 4-8 February, 1984. It was
attended by over 300 scientists and engineers, representing 26 countries. The conference
proceedings were published in two volumes by Pergamon Press under the title “Corrosion:
Industrial Problems, Treatment and Control Techniques”. The conference provided a forum
for the exchange of ideas between scientists and engineers from the region with their
counterparts from the industrialized countries.
The patronage of the present conference, the organizing bodies, and the emphasis on
industrial corrosion and corrosion prevention reflect the keen interest of the countries in the
region in actively combating corrosion problems. This also reflect the recognition of the
economic impact resulting from the corrosion of materials.
Kuwait and the other Arab countries rely heavily on the utilization of metallic materials in
their oil-based industries. Seawater derived from the Arabian Gulf is used in water
desalination and as an industrial cooling media. The salinity of the Arabian Gulf seawater is
very high when compared to other seawater bodies. The Arabian Gulf countries are located in
an arid environmental zone where the temperature during the summer months could reach
50oC and the humidity during the autumn season could become 80% in some of the Gulf
states. All these factors contribute to the enhancement of the rate of corrosion of metals
and/or cause unpredictable service failures.
The program of the present conference includes a field visit to one of Kuwait’s modern
refineries and a trip to one of Kuwait’s oil fields. The success of the conference is perhaps
difficult to assess. However, the quality of the papers in this volume provides some
indication.
The Editors
v
PREFACE
The technical program of the present conference includes five plenary lectures and fifty three
scientific presentations from about twenty two countries. A number of honorary speakers,
carefully selected from high ranking officials and policy makers, were also invited to address
the conference. The honorary speakers are expected to provide an overview of the magnitude
of corrosion related problems in the Middle East as well as the avenues of linkage between
corrosion science and industrial applications. The conference papers were carefully selected
to include a blend of fundamental and applied research, and industrial experience. Such a
blend was thought to be essential for providing the participants from both industry and
academia with a chance to become familiar with the challenges facing each group and the
preventive actions to meet them. The papers were refereed in terms of scientific and
technical content and format in accordance with internationally accepted standards.
The papers in the proceedings are grouped in the following sections for quick reference:
•
•
•
•
•
•
•
•
•
Plenary Lectures
Oil Field Corrosion
Corrosion in Refinery and Petrochemical Industries
Seawater Corrosion
Corrosion in the Building Industry
Fundamental Aspects
Corrosion Protection and Monitoring
Corrosion Management
Novel Techniques
The plenary papers are mostly reviews covering important topic related to the objectives of
the conference. The remaining papers cover various topics of major importance to corrosion
in general and particularly to the oil-based and desalination industries. A good number of
papers delt with corrosion protection and new techniques for corrosion monitoring.
The task of editing this volume was facilitated by the efforts of the International Advisory
Committee and the Scientific Committee for the conference who reviewed all the papers. The
editorial board gratefully acknowledge these efforts; the cooperation, time and effort of all
authors ; and the management of the Kuwait Institute for Scientific Research for allocating
the required resources to prepare the manuscript of this volume.
The Editors
vi
TABLE OF CONTENTS
Foreword............................................................................................................................................v
Preface...............................................................................................................................................vi
Organizing Committees......................................................................................................................xi
Acknowledgement.............................................................................................................................xii
PLENARY LECTURES
Corrosion Management
V. Ashworth........................................................................................................................................1
The Deterministic Prediction of Damage
D.D. Macdonald...............................................................................................................................17
Relevance of Laboratory Corrosion Tests in Corrosivity Assessment
and Materials Selection: Case Studies
R.D. Kane.........................................................................................................................................37
Corrosion of Condensers in Multi Stage Flash
Evaporation Distillers
A.M. Shams El Din...........................................................................................................................49
Correct Materials Selection for Desalination
-The Key To Plant Reliability
J.W. Oldfield....................................................................................................................................67
OIL FIELD CORROSION
Corrosivity Prediction for Co2/H2s Production Environments
S. Srinivasan and R.D. Kane.............................................................................................................89
Testing of Drilling Fluids Formulated From Tabuk Formation Clays
M.N.J. Al-Awad, A.S. Dahab and M.E. El-Dahshan........................................................................111
Preventing Sulfate Scale Deposition in Oil Production Facilities
C.J. Hinrichsen, M.J. McKinzie, S. He, J. Oddo, A.J. Gerbino, A.T. Kan, and M.B. Tomson...........127
Concerns Over the Selection of Biocides for Oil Fields and Power Plants:
A Laboratory Corrosion Assessment
J. Alhajji and M. Valliappan...........................................................................................................135
Evaluation of Microbially Influenced Corrosion Risks and Control Strategies
in Seawater and Produced Water Injection Systems, Kuwait
P.F. Sanders, M. Salman and K. Al-Muhanna.................................................................................149
Hydrogen Degradation of Steel - Diffusion and Deterioration
M. Farzam...................................................................................................................................................165
Control Strategies for Thermophilic Sulphate-Reducing Bacteria
P.F. Sanders, H.M. Lappin-Scott and C.J. Bass...............................................................................179
Corrosion Evaluation of Austenitic and Duplex Stainless Steels in Simulated
Hydrogen Sulphide Containing Petrochemical Environments
K. Saarinen and E. Hamalainen......................................................................................................191
vii
Damage of Pump Linkages and Tool Joints Caused by Crack Corrosion
A. Kinzel.........................................................................................................................................201
Analysis of Soils Possibility to Give Rise to Pipe Metal Stress Corrosion Cracking
V.G. Antonov and S.A. Loubenski....................................................................................................209
A Mysterious Downhole Corrosion Failure in an Oil Well
A. Husain and A. Hasan..................................................................................................................215
CORROSION IN REFINERY AND PETROCHEMICAL INDUSTRIES
Methodologies for Assessment of Crude Oil Corrosivity in Petroleum Refining
S. Tebbal and R.D. Kane.................................................................................................................225
New Nickel Alloys Solve Corrosion Problems of Various Industries
D.C. Agarwal and W.R. Herda....................................................................................................................233
Macro-Micro Segregation Bands (MMB) as a Main Factor Influencing
Steel Applicability for the Petroleum Industry
A. Mazur.........................................................................................................................................245
Fluid Catalytic Cracking Interstage and High-Pressure Cooler Corrosion
S.M. Halawani................................................................................................................................255
Assessment of Cracks in a High Pressure Multilayered Reactor
for its Fitness for Purpose
A.M. Askari, M.I. AL-Kandari and P.K. Mukhopadhyay.................................................................263
Polythionic Acid Stress Corrosion Cracking of Incoloy 800: Case Study
and Failure Analysis
M.S. Mostafa and S.A. Hajaj...........................................................................................................273
Corrosion of Tube Heaters in Refineries: Symptoms and Cures
A. Attou , A. Rais and H. Smamen...................................................................................................283
SEAWATER CORROSION
Super Duplex Grade UNS S32750 for Seawater Cooled Heat Exchangers
P.A. Olsson and M.B. Newman.......................................................................................................289
Evaluation of Aluminum Alloy 5083 Weldments to Stress Corrosion Cracking
in Seawater
A. Saatchi, M.A. Golozar and R. Mozafarinia.................................................................................301
Cavitation Corrosion Behavior of Some Cast Alloys in Seawater
A. Al-Hashem, P.G. Caceres and H.M. Shalaby..............................................................................311
Microbiologically Induced Corrosion of a Stainless Steel Pipe
H.H. Lee, M. Ali and K. Al-Omrani................................................................................................323
A Laboratory Study of Service Failure of Al-Brass Tubes in Arabian Gulf Seawater
H.M. Shalaby, W.T. Riad and V.K. Gouda......................................................................................329
CORROSION IN THE BUILDING INDUSTRY
Corrosion of Reinforced Concrete Structures and the Effects of the Service Environment
S. Al-Bahar and E.K. Attiogbe........................................................................................................341
Corrosion of Concrete in Seawater
viii
M. Pakshir and S. Esmaili...............................................................................................................353
Concrete Quality and its Effect on Corrosion of Steel Reinforcement
E.K. Attiogbe and S. Al-Bahar........................................................................................................361
The Effect of the Type of Copper on its Corrosion Behavior in Kuwait’s Soft Tap Water
H.M. Shalaby and F.M. Al-Kharafi.................................................................................................371
FUNDAMENTAL ASPECTS
Corrosion Behavior of Vanadium in Aqueous Solutions
W.A. Badawy, F.M. AI-Kharafi and M.H. Fath-Allah......................................................................383
The Effect of UV Irradiation on Passive Films Formed on Type 304 and 316 Stainless Steels
M.S. Al-Rifaie, C.B. Breslin, D.D. Macdonald and E. Sikora..........................................................395
Kinetics of High Temperature Corrosion of a Low Cr-Mo Steel in Aqueous NaCl Solution
W.A. Ghanem, F.M. Bayyoum and B.G. Ateya................................................................................407
Corrosion and Passivation Behaviour of Aluminium and Aluminium Alloys:
Mechanism of the Corrosion Process
F.M. AI-Kharafi, W.A. Badawy and A.S. El-Azab............................................................................417
The Susceptibility of Molybednum and Vanadium-Bearing Austenitic Stainless Steel
Weldments to Intergranular Corrosion
M.K. Karfoul..................................................................................................................................431
Effect of Crystallization on the Corrosion Behavior of Amorphous
FeCr9P6C3Si0.2 Alloy in 1 M H2SO4
F. Hajji, S. Kertit, J. Aride and M. Ferhat.......................................................................................441
CORROSION PROTECTION AND MONITORING
Experience With VOC-Compliant Waterborne and High Solids Coatings
in Corrosive Environments
P Kronborg Nielsen........................................................................................................................449
Anticorrosive Film-Forming Nonpolluting Products Achieved in Romania
R. Serban, N. Moga and E. Stockel.................................................................................................461
Cathodic Protection Under Disbonded Coatings of 56 Inch Gas Pipeline
Along the Kangan-Shiraz
M. Pakshir......................................................................................................................................471
Synergistic Effect Existing Between and Among a Phosphonate, Zn2+, and Molybdate
on the Inhibition of Corrosion of Mild Steel in a Neutral Aqueous Environment
S. Rajendran, B.V. Apparao and N. Palaniswamy...........................................................................483
Evaluation of Corrosion Inhibitors for Carbon Steel, Monel 400 and Stainless Steel 321
in a Monoethanolamine Environment Under Stagnant and Hydrodynamic Conditions
J. Carew, H. Al-Sumait, A. Abdullah and A. Al-Hashem..................................................................493
Laboratory Evaluation of the Effects of Ozone on Corrosion Rates and Pitting
of Engineering Alloys
S. Nasrazadani...............................................................................................................................501
A Critical Comparison of Corrosion Monitoring Techniques Used in Industrial Applications
M.S. Reading and A.F. Denzine......................................................................................................511
ix
Detection, Localization and Monitoring of Stress Corrosion Cracking, Hydrogen Embrittlement
and Corrosion Fatigue Cracks During Service Conditions Using Acoustic Emission
L. Giuliani......................................................................................................................................521
Electrochemical Monitoring of Aerobic Bacteria and Automation of Biocide Treatments
L. Giuliani......................................................................................................................................533
Corrosion Monitoring for Integrity of Pipeline
G.L. Rajani.....................................................................................................................................543
Power and Desalination Plants: Pumps, Corrosion and Maintenance
H. Hosni, N.J. Paul and A. Masri...................................................................................................555
CORROSION MANAGEMENT
Impact of Metallic Corrosion on the Kuwait Economy Before and After the Iraqi Invasion:
A Case Study
F. Al-Matrouk, A. Al-Hashem, F.M. AL-Kharafi and M. EL-Khafif.................................................567
Corrosion Problems in a Steam Condensate System and Treatment of Condensate for Recovery
G.L. Rajani.....................................................................................................................................581
Improved Cathodic Protection of Above Ground Storage Tank Bottoms: MAA Refinery Experience
A.K. Jain, L. Cheruvu and M.E. Al-Ramadhan................................................................................597
Impact on Ship Strength of Structural Degradation Due to Corrosion
M.A. Shama....................................................................................................................................615
NOVEL TECHNIQUES
Contact Electric Resistance (CER) Technique for Monitoring of Process Plants
and for Solving Practical Corrosion Problems
K. Saarinen and T. Saario..............................................................................................................627
Design of Radio Frequency Methods for Corrosion Processes Monitoring
Yu.N. Pchel’nikov, Z.T. Galiullin and A.S. Sovlukov.......................................................................637
A New, Rapid Corrosion Rate Measurement Technique for All Process Environments
A.F. Denzine and M.S. Reading......................................................................................................647
Assessing Corrosion of Thick Marine Paints by Surface Corrosion Potential Mapping (SCM)
and AC Impedance Spectroscopy (EIS)
A. Husain........................................................................................................................................657
Optics and Lasers in Corrosion Laboratory
K. Habib and F. Al-Sabti................................................................................................................669
Author Index...................................................................................................................................677
Subject Index..................................................................................................................................679
x
ORGANIZING COMMITTEE
Jasem Al-Besharah
Khaled Al-Muhailan
Abdulhameed Al Hashem
Hamdy M. Shalaby
Abbas Ali Khan
Hussain Shareb
Jamal Al-Hajji
Khaled Shehab
Khalifa Al-Feraij
Abdel Monem Bedair
Mohammad Ashkanani
Mohammad Al-Rasheed
Mohammed Al-Qalaf
Abdul Khaliq Mustafa
Khawla Al-Rifaee
Chairman
Rapporteur
Coordinator
Member
Member
Member
Member
Member
Member
Member
Member
Member
Member
Member
Member
KISR
KFAS
KISR
KISR
KFAS
OAPEC
KU
KNPC
MEW
PIC
KOC
GCC
KCS
KISR
MO
INTERNATIONAL ADVISORY COMMITTEE
Ahmed M. Shams El Din
John Oldfield
Russel D. Kane
Digby D. MacDonald
Member
Member
Member
Member
UAE
UK
USA
USA
Chairman
Rapporteur
Member
Member
Member
Member
Member
Member
Member
Member
KISR
KISR
KISR
KISR
KU
KOC
KOC
KOC
KNPC
KNPC
SCIENTIFIC COMMITTEE
Hamdy M. Shalaby
Abdulhameed Al Hashem
Khalid Habib
Adel Hussein
Waheed Badawi
Afkar Hussain
Emad Al Naser
Eman A. Razzak Al-Shayji
Lakshmipati Cheruvu
Fahed Al-Otaibi
xi
ACKNOWLEDGEMENT
The Organizing Committee was deeply honored by the patronage of H. H. The Crown Prince
and Prime Minister Sheikh Saad Al-Abdullah Al-Salem Al-Sabah, which reflects his keen
interest in science and technology.
The Committee was also grateful for the financial support of the Kuwait Institute for
Scientific Research, Kuwait Foundation for the Advancement of Science, Kuwait National
Petroleum Company, Kuwait Oil Company, Ministry of Electricity and Water, Kuwait
University, Ministry of Oil, the Gulf Cooperation Council, Kuwait Chemical Society,
Organization of Arab Petroleum Exporting Countries, and Petrochemical Industries
Company.
The Committee would also like to extend its deep appreciation for the effort and time put
forth by the distinguished honorary speakers, the members of the International Advisory
Committee, and the Scientific Committee.
We would like to thank our colleagues, the members of the working committees, at the
Kuwait Institute for Scientific Research and the chairmen and cochairmen of the sessions,
who provided unlimited assistance at times when it was really needed. Finally, we feel
deeply indebted to the authors of papers and participants for their valuable contribution to the
success of the conference
Jasem Al-Besharah
Chairman, Organizing Committee
xii
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
CORROSION MANAGEMENT
V. Ashworth
Global Corrosion
The White House, Victoria Road, Shifnal, England, TF11 8AF
ABSTRACT
The consequences of corrosion are often very costly. Little surprise, therefore, that a substantial
engineering effort is directed towards its prevention and control. By contrast, little consideration
seems to be directed towards making anti-corrosion effort cost-effective. This paper addresses the
problem of ensuring value-for-money corrosion engineering and the possible limitation of unnecessary
corrosion control activities.
Corrosion in itself is not important, but the consequences of corrosion failure may well be. So
the first step in corrosion management is a corrosion risk assessment to evaluate the risk associated
with failure in any item. This is not an evaluation of the risk of failure alone, but of the consequences
should that failure occur.
Given an assessment of risk, a strategy of corrosion management can be constructed. This might
involve lifetime corrosion control for items identified as producing a high risk. A less rigorous, but
monitored, level of protection might be adopted for medium risk items, whilst no action at all may be
considered necessary in the case of low risk items. Thus, resources are distributed according to the
risk.
Once a strategic approach has been defined, the tactics of corrosion management may be
determined. These will include not only the specific corrosion control activity or activities that will be
used in any given case, but also any monitoring and inspection requirements that are necessary. The
object is to maintain corrosion within acceptable limits at minimum cost in all parts of the facility and
throughout the facilitiy’s life.
Key Words: Risk, probability, monitoring, inspection, corrosion management
INTRODUCTION
The purpose of industry is to make a profit from the production of supplies and artefacts.
In an increasingly competitive world, there is continuing pressure on prices. If the selling
price is under pressure, profitability may only be maintained or increased by cutting costs.
Any factor that serves to increase costs represents a tax on profits.
Corrosion is one certain consequence of using engineering materials. Commonly,
corrosion will be modest, but not always. In the hydrocarbon production and processing
industries and the chemical industry, for example, the exception almost becomes the rule.
Since corrosion brings a cost, it impacts profits.
1
Plenary Lectures
THE COST OF CORROSION
The first formal attempt to assess the cost of corrosion to a nation was made in the UK in
1970 [1]. Since that time, similar studies have been published, in Australia [2], the US [3]
and elsewhere. One surprising outcome is that the cost of corrosion to an industrialized
nation is relatively constant at approximately 3.5% of the gross domestic product (GDP). To
put the matter in context, this is substantially higher than the cost of fires which in the UK is
put at ~0.5% GDP.
Sedriks [4] reported the experience of the Dupont Company in the period 1968-71. After
examining 685 plant failures, it was concluded that 55% were due to corrosion and 45% to
mechanical failure. This may be regarded as a remarkable outcome given that, by the
standards of the time, Dupont was corrosion aware and the greater proportion of the material
that failed was stainless steel.
The Dupont experience was mirrored by that of Britoil in the UK during the period
1978-88 [5]. As Table 1 shows, 33% of the failures that were analysed were attributed to
corrosion.
Table 1. Analysis of Oilfield Failures [5]
Type
Corrosion (all forms)
Fatigue
Mechanical damage/overload
Brittle fracture
Fabrication defects (not welding)
Welding defects
Other
Frequency (%)
33
18
14
9
9
7
10
Not infrequently, corrosion hits the headlines because some particularly dramatic failure
occurs resulting in the loss of life and property. At Flixborough Works in the UK, a chemical
explosion related to a corrosion failure resulted in 28 fatalities, 36 serious injuries, virtual
destruction of the plant and damage to some 2000 third party properties [6]. In Guadalajara,
in the early 1990's, stray current corrosion of a water pipe produced a failure that caused
erosion-corrosion of an adjacent gasoline line [7]. The leaking gasoline caught fire,
producing an explosion in which tens of local inhabitants were killed.
The accumulated corrosion failures at Dupont and Britoil were potentially costly and the
two accidents were certainly so. That cost is ultimately borne by the community, but in the
short term it falls on the industry concerned.
There is a growing awareness in industry of the costs of corrosion. This engenders a
desire for more effort and expenditure on corrosion prevention and control. Industry finds
willing allies in meeting this goal from the companies that sell anticorrosion materials and
systems. Their products and services are not free.
2
Ashworth
THE COST OF CORROSION PREVENTION AND CONTROL
Expenditure on corrosion prevention and control is no less a tax on profits than the cost
of corrosion itself. It is, therefore, entirely appropriate to ask if this expenditure is necessary.
An accountant can usually produce figures to illustrate the impact of a corrosion failure
on profit. Some tangible, and less tangible, inputs to the calculation are given in Table 2.
Table 2. Cost of Corrosion Failure
Safety hazards
Loss of capital plant or equipment
Fire/explosion
Loss of production capacity
Loss of product quality
Maintenance/repair/replacement
Loss of stored/entrained product
Pollution clean-up costs
Increased insurance premiums
Loss of consumer confidence
Alienation of workforce
Increased scrutiny by statutory bodies
Public image
Accountants
By contrast, the accountant is rarely moved to make an assessment of the cost of not
having a failure. If a plant and equipment operate without breaking down, everybody is
usually well satisfied. It is rare to question whether the cost of achieving that performance
has been excessive or even worthwhile. Table 3 lists some sources of possible over-spending
in endeavouring to avoid corrosion failures.
Table 3. Costs of Over-Protection
Unnecessarily expensive materials
Overdesign of metal sections
Excessive weight
Excessive inhibitor consumption
Excessive monitoring
Excessive inspection
Excessive data handling
Overdesigned CP systems
Over-operated CP systems
Excessive replacement stocks
Premature retirement of equipment
Over specified protective coatings
The items that relate to evident over-engineering in this list may be readily understood.
However, two areas, monitoring and inspection, are worth singling out because they are so
often regarded as a good thing, i.e., they have intrinsic merit. This is far from the case.
Industrial corrosion monitoring is commonly excessive both in terms of the extent of
monitoring and the sophistication of the equipment used. We need to remind ourselves that
corrosion monitoring has never controlled any corrosion. This author believes that corrosion
monitoring should be used almost exclusively in a process control function, the process, in
this case, being corrosion. Thus, a monitoring device should only be used in circumstances
where the output from it can validly be used to adjust some corrosion controlling function,
e.g., inhibitor injection or cathodic protection (CP) output. It follows that probes should not
be installed where such action cannot be taken, nor should they be installed where probes
3
Plenary Lectures
installed elsewhere fulfil essentially the same function. In practice, these rules are observed
more in their breach than in their application. Likewise, for reasons that are well understood,
corrosion probes provide precise information on what is happening on the probe and, often,
comparatively little about what is happening on the pipe or vessel wall. The value of the
probe is that it detects change and prompts review and, possibly, action. Very often simpler
means of detecting change than the use of corrosion-measuring devices will serve the same
function at less cost, e.g., reference electrodes, pH probes, dissolved oxygen meters and
moisture meters in the gas phase. They also have the merit, where relevant, of permitting
continuous readout which allows the identification of the precise 'upset' that produces
corrosion.
Inspection is similarly open to over-engineering. How often is inspection carried out
because we have the opportunity ? It is remarkable that upon shutdown, the internal
inspection of tanks and vessels will often be considered mandatory. Yet pipelines or
pipework that carry the same fluids are not inspected. It has been pointed out [8] that the cost
of inspection is high, often equivalent to 2-6% of the invested capital. That is a significant
tax on profits. What is often overlooked is that inspection can often be potentially dangerous
and may even produce conditions conducive to corrosion, e.g., when sulphuric acid tanks are
opened up for inspection.
If corrosion costs money and corrosion control costs money, how do we target the
optimum approach that strikes the correct balance between ignoring corrosion and seeking to
control it ? The answer lies in
• assessing the risk of corrosion failure on an item-by-item basis
• developing a lifetime corrosion management plan for each item to contain
corrosion at an acceptable level of risk
This latter implies the need for risk modification, and sometimes, a defined degree of
corrosion control activity. The goal is to maintain corrosion at an acceptable level. This begs
the question: What is acceptable ?
RISK
It is a very dangerous game to talk about risk, largely because it is an ill-defined subject,
and yet, everybody has a perception of what it is. What is clear is that risk is bad. We always
associate risk with the likelihood of an undesirable or catastrophic event occurring. Thus, we
talk of the risk of climbing, flying or crossing the road, but never of the risk of a traffic-free
journey to work. Moreover, not everybody sees a particular risk in the same way. For
example, we know that individuals are prepared to take a greater risk if they feel that they
have some control over the process, or if the risk is associated with some activity considered
to be beneficial [9]. That is why climbers do not appear to recognise the risk of climbing that
is so self-evident to the rest of us. They feel a measure of personal control and perceive a
personal benefit.
In short, risk is subjective. This is a worrying matter if, before we can proceed to a
corrosion management plan, we need to make a corrosion risk assessment.
It is necessary to remove a little fuzziness. The most appropriate definition of risk is
4
Ashworth
Risk = Probability x Consequence
(1)
Despite its formality, this is not a very precise equation as we shall see. It does,
however, indicate that risk is not simply a reflection of the probability of something bad
happening.
Probability
Probability has the appearance of precision because it is a mathematical quantity. It
derives from the stochastic nature of the frequency of the occurrence of events. Given
sufficient failure data, a classic probability may be calculated to reflect the likelihood of a
particular event occurring. In using the probability, it is important to be sure of its validity.
Consider above-ground pipelines. The probability of failure, taking the overground pipeline
population as a whole, is much less than the probability of failure of a small diameter (150250 mm) line. Considerable error can arise from using the former probability in the latter
case. Nevertheless, given valid and relevant failure data, a useful quantitative probability can
be assessed.
Very commonly the failure data from which probability is calculated do not exist. The
so-called Bayesian technique [10] can then be used to compute a probability. The technique
uses prior knowledge (e.g., failure rates in similar, but not identical, circumstances elsewhere
and the view of experts) and refines it steadily as specific information becomes available with
plant operation. In the limit, of course, when the specific information database builds up
sufficiently, the classic approach becomes more reliable.
Probability has one other unfortunate characteristic: uncertainty. The probability of an
occurrence may be low, but it can happen tomorrow.
Consequences
The consequences element in Eq. 1 relates to the perceived magnitude of the loss if the
failure occurs. This is a very subjective matter since different people rate the various
consequences of an individual event differently, and many even disagree about the
consequences that derive from that event. Nevertheless, it is possible for these individuals to
list the potential consequences and to rate each consequence on a scale of 1-10 to compute a
consequence for Eq. 1. The number that emerges is entirely subjective.
We see that the probability we use in Eq. 1 may hide a degree of uncertainty because of a
lack of failure data, and it may include educated guesswork. Similarly, the consequence is a
subjective valuation of the consequences of a failure. The outcome is a value for risk which,
although numerical, is not exact.
CORROSION RISK ASSESSMENT
The foregoing does not seem to suggest that any form of risk assessment is likely to be
productive. Yet the experience is that, in the case of corrosion, it can be helpful and
rewarding.
In carrying out a corrosion risk assessment it is axiomatic that corrosion does not matter,
but its consequences do. The assessment then aims to combine objective estimates of the
possibility of a corrosion failure with the operating company's view of the level of
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Plenary Lectures
undesirability of the consequences of it occurring. It will be seen below that the probability
component of the risk is rendered quantitative and that the consequence component remains
subjective, but reflects accurately the perceptions and ambitions of the people that own and
run the plant. Subjective it may be, arbitrary it is not.
Methodology
The methodology of conducting a corrosion risk assessment has been discussed in detail
elsewhere [11]. Only a brief outline is presented here.
If we take the example of a corrosion risk assessment in a refinery or chemical plant, it
may be as coarse or as refined as the operator wishes. First, the plant is divided into systems,
e.g., the gas sweetening unit. Second, each system is broken down into items. These usually
comprise individual components, e.g., a vessel, a heat exchanger, a pump or a specifically
identified length of pipework associated with the system. An item may be more widely
defined in a coarse corrosion risk assessment or more closely defined in a fine analysis. In
the latter case, for example, it may involve considering a vessel as a number of discrete items
according to the known variation in fluid composition with height. Equally, it may be
necessary to single out non-stressed relieved welds as separate items and, from an internal
corrosion point of view, each dead leg.
What follows is an outline approach to corrosion risk assessment which has proved to be
successful. Other methods are available that operate somewhat differently [12,13], but aim to
achieve the same objective.
Life Factor
The aim of the corrosion risk assessment is to assign a risk number to each item using a
risk equation similar to equation (1).
There are a variety of ways to deal with the probability element. The experience of the
author's company is that it is best dealt with by assigning a life factor (L) that relates to the
residual corrosion life of the item.
The residual corrosion life is the anticipated time required for corrosion at the predicted
rate, or rates, to lead to failure to perform the required mechanical duty. Given information
on the materials of construction, the exposure environment, and the relevant circumstances
(e.g., temperature, pressure, flow, heat transfer, and stress), the morphology of corrosion can
be predicted with confidence. Where uniform attack is expected, maximum penetration rates
can be calculated using conventional corrosion engineering practices, including public
domain algorithms [14,15,16], in-house database information and, if necessary, modelling. If
localized corrosion is expected (e.g., pitting attack or one of the cracking modes of failure),
probabilistic analyses of failure [17 ] are more useful.
For risk assessment purposes, the estimate of residual life in years is transposed to a
dimensionless L. For example, an anticipated time to failure shorter than the time to the next
shutdown would be assigned an L = 3. Anticipated lives beyond that point would attract L =
2, except where the residual life is put at >10 years in which case L = 1. Of course, this
breakdown is arbitrary and the individual cut-offs, and the relative scoring, can be selected to
match the requirements of the plant owners.
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Ashworth
Consequence Factors
The point has been made that the consequence of a corrosion failure are more important
than the failure itself. Thus, the consequences that bear on plant operators' minds include:
•
•
•
•
•
•
•
•
Safety,
Production,
Emergency repair,
Operability,
Environment,
Third party interests,
Customer perception, and
Public perception.
Adverse effects on any, or all, of these may often flow from an isolated failure.
The consequence factor (C) is a numerical assessment of the perceived consequences of
a corrosion failure. The number is arrived at using structured group discussions with plant
management, operations personnel, maintenance engineers, loss prevention officers etc. It
elicits a subjective assessment. However, because individuals work towards a consensus in a
group, and the methodology of subsequent analysis is rigorous, the rankings produced
accurately reflect, in a quantitative way, the operating aspirations of the company concerned.
Thus, the C numbers provide the relative importance attached to any consequence. Since the
risk numbers that finally emerge are not absolutes but reflect perceived risk in a relative
manner, the subjectivity of the consequence analysis is not only permissible, but desirable.
The risk assessment becomes plant specific. That is, identical plants operated by different
companies or in different locations will produce different risk assessment results.
There are two elements in establishing the value of C:
• The individual consequence, and
• The events that can lead to that consequence.
Some consequences will always be regarded by company personnel as more undesirable
than others; to that extent, the staff can develop a point loading to be applied to each. This
gives a consequence rating (F); a typical set to emerge in one case is given in Table 4.
Table 4. Typical Consequence Rating (f)
Consequence
risk of safety to personnel or public
loss of production
pollution
loss of produce quality
loss of consumer confidence
Rating (F)
10
9
3
1
1
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Plenary Lectures
It should be noted that the company in question is not concerned about contamination of
the product by corrosion or any consequent alienation of the customer. This is a typical
response from a primary producer; quite different numbers would have arisen in an
assessment made in a food or pharmaceutical factory.
The events that lead to a given consequence produce an event rating (P). It is clear that a
number of events which might occur in a process plant, may lead to the same consequence.
The company staff are able to identify and rank these events according to their perceived
undesirability, as shown in Table 5. In this case, the table relates to two plants owned by the
same company in which one uses the product of the other.
Table 5. Typical Event Rating (P)
Consequence
Safety
Outage
Pollution
Quality
Event
Crack in a toxic line or equipment
Pinhole in a melt line
Crack in a flammable line or equipment
Crack in other HP line or equipment
Pinhole in a flammable line or equipment
Pinhole in other HP line or equipment
Other cracks
Pinhole in toxic line or equipment
Other leaks
Falling objects
Plant no. 1 - no standby
Plant no. 2 (HP) - no standby
Plant no. 2 (LP) - no standby
Plant no. 1 - standby
Plant no. 1 - non-critical - no standby
Plant no. 2 - standby
Plant no. 2 - non-critical
Plant no. 1 - non-critical - standby
Marine
Atmospheric
Final product (colour only)
Intermediate product
Rating (P)
10
10
9
7
7
6
5
4
2
1
10
9
7
5
5
4
3
3
10
5
10
5
It is not uncommon when considering safety, for staff to take into account the inventory
of a system. Thus, they will commonly regard a crack in a system or item with a high
inventory and, therefore, a high potential for damage, as more significant than one where the
inventory is small. Different values of P may then arise according to the volume of an
unisolatable part of the system. Any item included in the unisolatable part attracts the P value
for that part.
The values of F and P are combined to yield C:
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Ashworth
C = ∑Cx =
x=n
∑x = 1
fn (Fx, Px)
(2)
Where the subscript x refers to each of the consequences, e.g., safety, pollution etc. in turn.
The Risk Equation
The risk equation must reflect the operating company's perception of risk associated
with various forms and rates of failure. The equation, which is derivative of Eq. 1, produces a
numerical assessment of risk (R) and takes the form:
R = fn (L, C)
(3)
The shape of the function linking the life and consequence factors is determined by a
formalized heuristic procedure. The function is modified through a series of computer
iterations with the effect on the value of risk being assessed after each iteration.
Allocation of Risk Classes
The numerical value of risk can be calculated for each item in the plant. The higher the
value of R for any item, the greater the risk and the more attention that must be focused on
the local corrosion situation.
In practice, the spread of numerical risk values amongst all the items within a plant
usually proves to be a discontinuous spectrum. That is, the risk numbers tend to fall into
clusters with distinctive breaks between. This is an inevitable consequence of data like those
recorded in Tables 4 and 5. It permits a convenient reduction of the numeric data into risk
classes (e.g., low, medium and high). Such a sub-division aids communication of the
outcome of the corrosion risk assessment either on a narrative basis or as colour coded P and
ID's. It also assists with establishing a corrosion management programme.
Risk Modification
The fact that, in assessing a new or existing plant, areas of high risk have been identified,
does not mean that the risk must be tolerated. The aim should be to moderate the risk and to
move to a more acceptable condition.
Some methods of corrosion risk assessment [12] do not proceed as far as a risk equation
or a risk number, but rather consider separately the perceived severity of the probability (L in
this case) and the consequence (C). This produces a risk matrix as shown in Table 6.
Table 6. Risk Matrix
Consequence
H
M
L
H
H
HM
M
Probability
M
HM
M
ML
L
M
ML
L
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Plenary Lectures
Risk categories can then be devised as shown in Table 7:
Intelligent risk modification aims to move towards Zone 3. This is not to aim at zero
failure but to achieve, by good management, a tolerable level of risk.
Table 7. Risk Categories
High consequence
High probability
1
High consequence
Low probability
4
Consequence
Low consequence
Low consequence
High probability
Low probability
2
3
Probability
Using the risk equation approach, the aim is to concentrate resources on areas of high
risk in order to reduce the risk number (i.e., modify the risk). Some care has to be taken here.
It will often be the case that a high risk number will place the specific item in Zone 1 of Table
7. Clearly, for these items, it is important to move towards Zone 3 by means of corrosion
control activities that reduce the risk number.
Somewhat lower risk numbers may fall into either Zone 2 or Zone 4. Indeed, the same
risk number may apply to either a low consequence/high probability situation or a high
consequence/low probability. The former simply represent failures that will be an irritation;
pinholing in a seawater cooling line. The latter are certainly more serious; pinholing in a dry
flammable gas line, for example.
In the case of Zone 2, the high probability means that sufficient data were available to
assess the probability fairly accurately. By contrast, in the case of Zone 4, the reverse is true,
and there may be considerable uncertainty.
If the probability (in our case, L) has been calculated using tried and tested tools, then
identical risk numbers that derive from high probability/low consequence and low
probability/high consequence events are equally reliable. Where the L calculation has used
limited data, an uncertain algorithm or stochastic techniques, that level of reliability is absent.
Thus, in moving from Zone 1 towards Zone 3, it is often better to achieve Zone 2 rather than
Zone 4. Equally, it may be more important to move from Zone 4 than to move from Zone 2.
In general, the consequences of failure are usually not amenable to modification; thus,
risk can only be modified by changing the probability. For that reason, and to introduce
security, risk modification needs to pay attention to moving to situations where the
probability is known or can be determined with some degree of confidence.
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CORROSION MANAGEMENT
Corrosion risk assessment is not an end in itself. It identifies areas where corrosion may
be safely ignored and where it must be attended to. It even provides the pointer to where
resources will be spent with greatest reward. Thus, it provides the evidence that permits the
construction of a cost-effective corrosion management programme. The objectives of a
programme relate to the whole life of a facility and are to
• Maintain corrosion within predetermined acceptable limits at minimum cost,
• Develop and facilitate rapid access to, records showing the corrosion status of
each item within the facility in order to form a basis for future corrosion
management decisions, and to provide assurance for managers, owners and
statutory bodies, and
• Ensure that corrosion upsets are quickly identified and appropriate remedial
action is implemented, if necessary, to minimize the consequences of any
failure.
There is no universal corrosion management programme. Targeting these objectives is a
unique exercise for every facility. However, the philosophy of corrosion management is
common to them all. An overall strategy for corrosion management must first be agreed
upon, and then the tactics become self-evident.
Strategic Considerations
The corrosion risk assessment will have produced a risk ranking for all items of a plant.
This will enable a strategy for corrosion management to be set down. Table 8 illustrates a
strategy that might be drawn up for an industrial facility.
Table 8. A Corrosion Management Strategy
Assessed Risk
High
Medium
Low
Alternative Corrosion Management Options
Corrosion prevention, or corrosion
control for life, or corrosion control to
meet planned maintenance or planned
replacement
Corrosion control for life, or planned
maintenance
No action, replace if required
It will be noted that corrosion prevention, or careful corrosion control, is dictated by a
high risk classification. By contrast, a low risk classification justifies no corrosion
controlling action. A medium risk requires some action. Thus, corrosion management
involves a spectrum of activity from no action to considerable action according to the risk.
However, taking no action, or taking action, is not corrosion management unless the decision
to follow the particular course has been based on an assessment of risk. Action where it is not
needed, like inaction where it is, represents a waste of resources and a tax on profits.
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Plenary Lectures
It will be recalled that corrosion risk assessment is carried out by dividing the plant into
systems and items. Ultimately, the output relates to individual items. It is possible for an
item within a system to have a high risk classification whilst other items in the same system
belong to a lower risk class. The decision must then be made whether to apply corrosion
control to the system in order to preserve the item, or to ensure corrosion prevention for the
item (say, by the use of more corrosion-resistant material) and avoid dealing with the system.
The application of the broad strategy does allow, and requires, some flexibility in the tactics
adopted.
Tactical Considerations
The complete elimination of the chance of corrosion failure, i.e., corrosion prevention, in
a high risk area is rarely possible in an existing plant. Invariably, it would require a
significant engineering change, for example, replacement of existing materials by corrosionresistant alloys or modification of the process (e.g., addition of gas dehydration). Even with
new plants, such proposals might raise major design and engineering problems, not to
mention cost.
It is much more likely that active corrosion control will be adopted with the objective of
extending the time to failure of an item beyond the planned life of the plant, or up to some
planned maintenance shutdown. The adoption of this tactic requires that:
• the performance targets for the corrosion control are defined, and
• procedures are put in place to ensure the targets are met.
The performance target may be set in terms of an allowable rate of metal penetration.
This approach will most commonly be adopted when uniform corrosion is anticipated.
Alternatively, limits may be set on some parameter that is an indication of fluid corrosivity,
e.g., electrode potential in anodic and cathodic protection systems, dissolved oxygen in
oilfield water injection or boiler feedwater, pH, temperature, or dewpoint. Irrespective of
which approach is adopted, it will be necessary to obtain on-line information to make
adjustments as required. Thus, corrosion monitoring is necessary, and it then forms an
essential segment of the corrosion management plan. Table 9 lists some monitoring
techniques and indicates how they may be used in corrosion management.
Table 9. Corrosion Monitoring in Corrosion Management
Corrosion Control Strategy
Examples of Adjustments
and Activities Based on
Data Monitoring
Inhibition of crude oil On-line
probes
(e.g., Adjust inhibitor dosage,
pipelines
coupons, electrical resistance change inhibitor type,
probes)
discontinue inhibition
De-oxygenation of boiler O2 probes
Adjust oxygen scavenger,
feed-water
check pump seals, etc.
Impressed current CP
Potential
Adjust system output
Anodic
protection
of Potential
Adjust system output
sulphuric acid plant
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Monitoring Technique
Ashworth
Dehydration of process gas
On-line probes,
moisture detection
Temporary inhibition,
overhaul dehydrator
The key to effective corrosion management is information since it is on the basis of that
information that on-going adjustments to corrosion control are made. Information is valid
data. Thus, to make effective corrosion management decisions on a day-to-day basis, the
monitoring data must be valid. This is not simply a requirement for the probes to be
operating correctly. It requires that they be placed in the most appropriate places, i.e., at
those points where the corrosion controlling activity might be expected to work, but where it
might equally be expected to be least effective, e.g., remote from the inhibitor injection point.
In many cases specially designed traps are introduced into a plant so that corrosion probes
may be inserted. These often produce their own microenvironment, atypical of the plant
itself, and with little hope of effective entry for an inhibitor. Data from a probe in such a
location are unlikely to be relevant to corrosion management elsewhere in the system. Invalid
data leads to ineffective corrosion management.
Keeping Track
In any facility the means of corrosion management will vary from place to place. In one
location a corrosion resistant alloy may be used; in another, CP allied to coating may be
employed, whilst elsewhere no action may be taken because the consequences of any failure
are regarded as unimportant. In short, no corrosion management action is taken that does not
contribute positively to meeting the objective of containing risk whilst maintaining the level
of action at the minimum necessary.
It is important to ensure that the targets are being met. Overshooting the target will
involve excessive corrosion control costs, whilst undershooting the target may lead to a
situation that cannot economically be recovered. Corrosion monitoring is not appropriate for
the purpose since it rarely provides evidence of the metal loss from a pipe or vessel wall.
That is, aggregating the output from probes over time does not give any indication of the loss
of a section. The value of corrosion probes is that we rapidly develop experience so that we
can be reasonably sure that when the probes read a given value, we are on target, and that a
change in reading requires consideration of an adjustment to the corrosion controlling
activity. Reference electrodes, pH probes, moisture meters etc. often fulfil the same function.
Thus, corrosion probes and the like, do not provide quantitative performance assessment.
That can only come from inspection and nondestructive testing (NDT). These activities are
part of corrosion management since they provide reassurance, identify wasteful corrosion
control activity, and permit reassessment of the corrosion management programme. The
same critical approach that was adopted in setting up the corrosion control strategy must be
applied to the inspection strategy. That is, the resources must be applied according to the risk.
If we have attempted to modify the risk by instituting some corrosion control activity, we
should be tracking the success, or excess, of the activity in our inspection programme. Thus,
inspection is not based on convenience, inspecting because an item is accessible (at
shutdowns, for example). It should be based on the premise that if the consequences of
failure are to be avoided and the cost of control is to be minimized, inspection is necessary.
There must, therefore, be a clear connection between the risk assessment output, the corrosion
management strategy, the tactics of corrosion management and the inspection programme.
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Plenary Lectures
The key point approach to NDT is particularly effective. Here a limited number of
points, in areas where validation is required, that are regarded as typical and extreme, are
identified for NDT inspection at regular intervals. This provides, on a temporal basis, a
readout of the progress of corrosion which will validate, or otherwise, the targeting achieved
by, say, inhibitor injection.
Similarly, during internal inspection, the risk assessment will have identified particular
areas of concern, e.g., tube/tube sheet assemblies, tube baffles, and non-stress relieved welds,
which must become the focus of activity. Again, this will confirm whether the corrosion
control is adequate or perhaps is insufficient or excessive.
The data that are produced from the inspection activities must be valid and limited in
volume so as not to deter analysis or hide anomalies. Thus, it is important to restrict key
points and inspections to critical positions and to limit the frequency of inspection and survey
work. The time to the next inspection or survey should be indicated by the outcome of the
current work. That is, a lifetime fixed interval programme will usually prove wasteful;
inspection and survey should be carried out on an as-needed basis. A valuable template
giving an approach to the re-classification of in-service inspection is to be published in 1996
[18] and has been reviewed in reference [12].
Review
From time to time a corrosion management programme should be reviewed at both the
strategic and tactical level. In human affairs, things change. The management of a facility
will always be alive to current market trends, competitors activities, interest rate movements
and so on. Inevitably, it may be necessary to revise the management objectives from time to
time. Since the corrosion management programme was constructed to meet the objectives of
an earlier plant management plan, it will be necessary to review the programme and possibly
to alter it. Likewise, the pace of technological change is rapid compared to the anticipated
lifetime of most facilities. Thus, newer, more effective, cheaper means of achieving the same
ends may emerge, and indeed, it may be possible to adopt them in place of existing tactics
within the corrosion management programme. Thus, the programme is not a fixed blueprint,
but a means to an end that must be reviewed and revised to meet the current management
objective.
One objection that is raised to corrosion management planning comes from the corrosion
engineers themselves. They draw attention to the fact that by fixing permissible rates of
metal loss, the lifetime of the facility is effectively determined. Further, that management
will often, at a later stage, decide to extend the required operating life. There is then a
mismatch. The argument seems to be that corrosion management planning should ignore the
present requirements and anticipate the future requirements of the management. This is an
extremely wasteful approach. Certainly there is a possibility that a mismatch will occur and
will need to be overcome. That will be achieved at some cost. That cost must be attributed to
the decision to go for life extension and is, therefore, a natural consequence of that extension.
It needs to be included in the cost benefit analysis of extension, not hidden in lifetime
overspending in anticipation that life extension might be required. It may not be.
ILLUSTRATIONS
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Ashworth
Two recent examples illustrate how corrosion risk assessment provided important results
for the clients. In the first instance, the assessment of a petrochemical complex in the Middle
East found the plant to be extremely well engineered from the corrosion standpoint. It was
constructed predominantly in carbon steel, with excursions into more exotic metallurgy only
where the process conditions demanded it. However, the assessment highlighted, somewhat
to the client's surprise, the cooling water system as a high risk area.
By using a closed system with secondary cooling by seawater and specifying high
quality primary water with corrosion inhibitor injection, the designers had judged that it was
possible to construct the majority of the cooling system in carbon steel. Certain that the
primary heat exchangers were, however, constructed in a stainless steel due to the
aggressivity of the process fluid, the corrosion risk assessment identified modes whereby the
quality of the cooling water could be adversely affected (e.g., by leakage of seawater at
secondary plate exchangers). Failure to maintain cooling water within specification would
very rapidly lead to stress corrosion cracking of one of ten critical stainless steel process
exchangers, failure of any one of which would halt production.
In view of this, Global Corrosion put forward recommendations for modest on-line
monitoring of cooling water quality. Tied to this was the setting up of a formal action plan to
be followed in the event a sudden deterioration in water quality should be detected. The
client accepted and implemented these recommendations but, unfortunately, not before one
failure of the type predicted occurred.
The second example derives from an installation, also in the Middle East. The corrosion
risk assessment concluded that the absence of CP on water storage tanks, together with the
prevailing soil conditions, would result in high tank bottom corrosion rates. Since an
adequate supply of water was essential to maintain production, the assessment concluded that
the prospective failure of the tanks constituted a high risk and it was strongly recommended
that CP be installed.
In the event the client was reluctant to accept and act upon the outcome of the report.
Just under a year later the raw water tank perforated due to soil-side corrosion. The resulting
loss of water caused a two week interruption in production, prompted a belated decision to
install CP and engendered in the client a heightened appreciation of the benefits of corrosion
risk assessment and the need for effective corrosion management.
CONCLUSIONS
Corrosion cannot be ignored for it will not go away. However, there is little merit in
controlling corrosion simply because it occurs, and none in ignoring it completely. The
consequences of corrosion must always be considered. If the consequence of corrosion can
be lived with, it is entirely proper to take no action to control it. If the consequences are
unacceptable, steps must be taken to manage it throughout the facility’s life at a level that is
acceptable. To manage is not simply to control.
Good corrosion management aims to maintain, at a minimum life cycle cost, the levels of
corrosion within predetermined acceptable limits. This requires that, where appropriate,
corrosion control measures be introduced and their effectiveness ensured by judicious, and
not excessive, corrosion monitoring and inspection. Good corrosion management serves to
support the general management plan for a facility. Since the latter changes as market
15
Plenary Lectures
conditions, for example, change, the corrosion management plan must be responsive to that
change. The perceptions of the consequences and risk of a given corrosion failure may
change as the management plan changes. Equally, some aspects of the corrosion
management strategy may become irrelevant. Changes in the corrosion management plan
must, inevitably, follow.
REFERENCES
1. T.P. Hoar (Chairman), Report of the Committee on Corrosion and Protection, HMSO,
London, 1971
2. B.W. Cherry and B.S. Skerry, Corrosion in Australia - the Report of the Australian
National Centre for Corrosion Prevention and Control Feasibility Study, Monash
University, 1983.
3. L.H. Bennett, National Bureau of Standards Special Publication 511.1, NBS,
Washington, 1978.
4. A.J. Sedriks, Corrosion of Stainless Steels, Wiley, 1979, p. 7.
5. Kermani, An Overview of Wet H2S Attack: Types, Causes and Problems, in Papers of
the Conference on Wet H2S Attack on Steels, Institute of Mechanical Engineers,
London, 1996.
6. F. Lees, Loss Prevention in the Process Industries, Vol 2, p863, Butterworth (1989)
7. J.M. Malo, V. Salinas and J. Uruchurtu, Materials Performance 33, 8, 1994, p. 63.
8. C. Edeleanu and J.G. Hines, Materials Performance 29, 12, 1990, p. 68.
9. P. Slovic, Science 236, 17 April 1987, p. 280.
10. M.E. Giuntini, Proceedings of Fourth Space Logistics Symposium, Florida, November
1992.
11. V. Ashworth and W.R. Jacob, Proc. Corrosion 32, Australasian Corrosion
Association, 1992.
12. B. Spalford, Carbon steel equipment in wet H2S service, Papers of Conference on Wet
H2S Attack on Steels, Institution of Mechanical Engineers, London, 1996.
13. Private communication, Shell-Expro, UK
14. C. de Waard, V. Lutz and D.E. Milliams, Corrosion 47, 1991, 976.
15. F.A. Posey and A.A. Palko, Corrosion 35, 38 (1979)
16. J.W. Oldfield, G.L. Swales and B. Todd, Proc. 2nd BSE/NACE Corrosion
Conference, Bahrain, 1981.
17. M. Akashi, Proc. Conference of Life Prediction of Corrodable Structures, NACE,
1991.
18. EEMUA publication 179, A Working Guide for Carbon Steel Equipment in Wet H2S
Service (to be published in 1996)
16
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
THE DETERMINISTIC PREDICTION OF DAMAGE
D.D. Macdonald
Center for Advanced Materials
The Pennsylvania State University
517 Deike Building, University Park, PA 16802, USA
ABSTRACT
As our industrial and infrastructural systems (refineries, power plants, pipelines, etc.) age,
considerable economic incentive develops to avoid unscheduled outages and to extend operation
beyond the design lifetime. The avoidance of unscheduled outages is of particular interest, because
the failure of even a minor component can result in the complete shutdown of a facility. For example,
the unscheduled shutdown of a 1000 Mwe nuclear power plant may cost the operator between US $1
million and US $3 million per day, depending upon the cost of replacement power and other factors.
However, if component failures could be accurately predicted, maintenance could be performed
during scheduled outages, the cost of which has already been built into the price of the product. With
regard to life extension, the successful extension of operation beyond the design life translates into
enhanced profits and the avoidance of costly licensing and environmental impact assessments
associated with the development and construction of a new facility. In this case, as well, the key to
successful operation is the ability to avoid downtime, and hence, to maintain production. Eventually,
the frequency and severity of unscheduled outages will render operation uneconomic, and at that
point, replacement of the facility is necessary. In order to develop effective inspection and
maintenance scheduling and life extension technologies, it is first necessary to predict the evolution of
damage into the future as a function of various system variables. The only effective prediction
technologies are those based on determinism, in which the system behavior is described in terms of
natural laws. In this paper, the deterministic prediction of damage, via damage function analysis
(DFA), which provides a robust technology for estimating the damage function at future times, is
described. The application of DFA to the prediction of pitting damage is illustrated by reference to
pitting damage in condensing heat exchangers.
Key Words: Corrosion damage, determinism, prediction, pitting corrosion.
INTRODUCTION
Corrosion is a major cause of component failure, and hence, in the occurrence of
unscheduled downtime, in complex industrial systems. In particular, the various forms of
localized corrosion, including pitting corrosion, crevice corrosion, stress corrosion cracking
(SCC), and corrosion fatigue (to name the common forms) are particularly deleterious
because they frequently occur without any outward sign of damage, and because they often
result in sudden and catastrophic failures. Thus, the development of effective corrosion
damage prediction technologies is essential for the successful avoidance of unscheduled
downtime and for the successful implementation of life extension strategies.
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Plenary Lectures
Corrosion damage is currently extrapolated to future times using damage tolerance
analysis (DTA). In this strategy, known damage is surveyed during each subsequent
inspection, and the damage is extrapolated to the next inspection period allowing for a
suitable safety margin. We have argued [1] that this strategy is inaccurate and inefficient, and
that in many instances it is too conservative. Instead, we argue that damage function analysis
(DFA) is a more effective method for predicting the progression of damage, particularly when
combined with periodic inspection. DFA is based upon the deterministic prediction of the
rates of nucleation and growth of damage, with particular emphasis on the compliance of the
embedded models with natural laws. Although corrosion is generally complicated
mechanistically, a high level of determinism has been achieved in various treatments of both
general and localized corrosion.
The application of DFA is illustrated by reference to the development of damage due to
pitting corrosion of stainless steels in condensing heat exchangers. Deterministic models
have been developed for both the nucleation and the growth of damage, and these models
have allowed us to calculate the damage function as a function of exposure time and system
conditions.
FUNDAMENTAL CONCEPTS
In this paper, I outline a deterministic method for predicting the damage function for
pitting corrosion in condensing heat exchangers [1,2]. This method is considered to be
potentially superior to empirical (including stochastic and probabilistic) techniques, because
it is mechanistically-based and hence provides analytical relationships between the damage
function (i.e., number of pits versus pit depth presented in the form of a histogram [1]) and
the damaging variables (e.g., chloride concentration and combustion parameters).
Accordingly, deterministic methods are expected to be more efficient at using databases,
because a lesser need exists to establish the damage function/damaging variable relationships
empirically.
Any deterministic model must account for the fact that localized corrosion involves
nucleation and growth phenomena which occur sequentially for a single site but that tend to
occur in parallel for an ensemble of pits. Furthermore, the model must account for the
experimental observation that the parameters that characterize the breakdown event are
distributed, due to the fact that the population of sites on any real surface is not homogeneous.
Outlined below is one model that satisfies these (and many other) conditions related to the
nucleation and growth of damage resulting from localized corrosion. While the model may
not be complete (or even correct), it is deterministic in that the distribution function and the
relationships between the model parameters and the damage function are analytic and follow
from the natural laws. In illustrating this technology, I have chosen to discuss the prediction
of the damage function for pitting corrosion, because this form of attack is almost ubiquitous
in condensing heat exchangers. Furthermore, pitting corrosion displays most of the features
of all forms of localized attack, including an induction time and the autocatalytic
development of the damage.
The algorithm developed in this study to estimate the damage functions for condensing
heat exchangers contains five modules as outlined in Fig. 1. Also indicated are the parameters
that propagate from one module to the next. The output of the algorithm can be specified in
three forms:
18
Macdonald
• For a specified probability of failure, the algorithm estimates the damage
function as a function of exposure time and computes the number of pits with
lengths exceeding the condenser wall thickness to predict the service life.
• For a specified probability of failure and design life, the algorithm calculates
the wall thickness to ensure acceptable performance.
• For a specified wall thickness and design life, the algorithm calculates the
failure probability.
MODEL INPUTS
Duty
Cycle
Chloride
Concentration
Condensate
Temperature
Flue Gas
Composition
Condensate Chemistry
Model
pH.[Cl-]*
Mixed Potential
Model
Ecorr.pH. [Cl-]
Pit Nucleation
Model
N(t).Ecorr.pH. [Cl-]
Pit Growth
Model
N(tobs) vs. n(u)
Damage Function
Model
Service Life
Figure 1.
Wall Thickness
Specifier
Failure
Probability
Structure of the algorithm for the prediction of damage function
(*parameters propagated from one model to the next)
Below I describe the various modules in this algorithm; however, due to the limited
space available, I outline only the principles of these modules.
19
Plenary Lectures
The Condensed Chemistry Module (CCM)
The composition of the flue gas will differ from burner to burner. With this in mind, we
developed a generalized condensate model for the condensate environment. This model
assumes the flue gas to be a mixture of CO, CO2, H2S, NO, NO2, SO2, SO3, and H2O. The
relative proportions of these components may vary widely from furnace to furnace, depending
on the nature of the ambient air, the air/gas ratio, and the impurities of the gas. The goal of
the condensed chemistry module (CCM) is to calculate the pH and the composition of the
condensate on the condenser surface. The pH is a key parameter in controlling the rates of pit
nucleation and pit growth. The concentrations of species in the liquid layer determine the
ionic conductivity of the solution, which has great impact on the pit growth rate. The module
employs an equilibrium model along with mass balance and charge balance constraints, and
computes ion activity coefficients using the extended Debye-Huckel theory. It is assumed
that the condensed liquid film is in equilibrium with the ambient environment, so that
equilibrium calculations are applicable. The details of this module are described in the
literature [3].
A typical gas-fired heat exchanger is schematically shown in Fig. 2a [4]. The
temperature ranges from approximately 308oK in the cold end to 353oK in the hot end,
depending on the design of the heat exchanger. Typical values of the pH and chloride
concentration in these different zones are given in Fig. 2b [4]. It is shown that the condensed
liquid phase is enriched in chloride to the extent of approximately 150 ppm in the hot end.
Acidification of the condensed thin liquid layer also occurs, in that pH values as low as 2.7
and 3.3 are found at the hot end and the cold end, respectively. In Fig. 2c, the computed pH
for a typical composition of the flue gas and the chloride content of the condensate are
presented. The calculation shows a variation in pH from 2.93 to 3.32 from the hot end down
to the cold end. Recognizing the wide range of operating conditions and designs of
condensing heat exchangers, it is concluded that good agreement is observed between the
experimental data and theoretical prediction.
Cooling Air
Exhaust
Heat Exchanger
Simulator
Flue Gas
Flow
Zone 1 Zone 2
Zone 3
Zone 4
T = 353-326oK
Zone 5
T = 308-326oK
Cooling Air
Figure 2a. Schematic diagram of a typical heat exchanger in a gas-fired furnace [4]
20
Macdonald
Figure 2b. Characteristics of flue-gas condensate from different zones [4]
Figure 2c. Calculated pH in the condensate as a function of temperature and
chloride concentration (PCO = 2.60 x 10-5 atm, PCO2 = 6.80 x 10-2
atm, PSO2 = 9.37 x 10-9 atm, PSO3 = 2.00 x 10-9 atm, PH2S = 1.20 x
10-4 atm, PNO = 1.74 x 10-8 atm, PNO2 = 2.67 x 10-8 atm, PO2 = 1.00
x 10-4 atm)
The Mixed Potential Module (MPM)
The mixed potential module (MPM), which is based on the Wagner-Traud hypothesis [5]
for free corrosion processes, was developed to calculate the corrosion potentials of alloys in
corrosive environments. The theory outlined here is essentially identical to that developed by
Macdonald et al. for calculating corrosion potentials for stainless steel components in the heat
transport circuits of boiling water reactors (BWRs) [6,7]. The theory is based on the physical
condition that charge must be conserved in the system.
21
Plenary Lectures
Because electrochemical reactions transfer charge across a metal/solution interface at a
rate measured by the partial currents, charge conservation demands that
Σ iR/O,j (E ) + icorr (E ) = 0
j=1
(1)
where iR/O,j is the partial current density due to the j-th redox couple in the system, and icorr
is the corrosion current density of the substrate. The currents are written as functions of the
potential E to emphasize the fact that the partial currents depend on the potential drop across
the metal/solution interface. Indeed, the solution to Eq. 1 provides the quantity that we seek
(i.e., the corrosion potential). Note that in deriving Eq. 1, the surface of the alloy is assumed
to be equally accessible to all reactions in the system.
Figure 3. Calculated and measured corrosion potentials as a function of
temperature and oxygen partial pressure for alloy A129-4C (the solution
composition for the experimental measurement: [HCl] = 200 ppm, [HF]
= 40 ppm, [H2SO4] = 20 ppm, PCO = 2.60 x 10-5 atm, PCO2 = 6.80 x 10-2
atm, PSO2 = 9.37 x 10-9 atm, PSO3 = 2.00 x 10-9 atm, PH2S = 1.20 x 10-4
atm, PNO = 1.74 x 10-8 atm, PNO2 = 2.67 x 10-8 atm, PO2 = 1.00 x 10-4 atm)
Experimental polarization curves and kinetic data for the reduction of oxygen on Alloy
Al29-4C were used as input to the MPM. The module then calculates the corrosion potential
of the alloy under the service condition, as shown in Fig. 3. The calculation indicates that if
the solution is saturated with oxygen, the corrosion potential is only weakly dependent upon
case of an operating condensing heat exchanger, the corrosion potential varies from -290 mV
to -460 mV as the temperature increases from 300o to 370oK.
The Pit Nucleation Module (PNM)
As a result of an intensive effort over the past decade to develop an understanding of the
breakdown of passive films, we have derived theoretical distribution functions for passivity
22
Macdonald
breakdown that are in good agreement with experimental data [8-11]. We derived these
distribution functions from our point defect model for the growth and breakdown of passive
films, by assuming that breakdown occurs when a critical concentration of cation vacancies
accumulates locally at the metal/passive film interface, such that decohesion occurs between
the barrier layer and the metal substrate. Subsequently, localized dissolution and/or
mechanical instability leads to rupture of the film [8]. We further assumed that the
breakdown sites are normally distributed with respect to the diffusivity of cation vacancies
within the film [9-11]. The full derivation of these distribution functions can be found in the
literature [8-11].
From the point defect model [8], the breakdown voltage and the induction time for a
single pit nucleation site on the surface are given by
Vc =
4.606RT
⎡ J
⎤ 2.303RT
log o m−χ / 2 −
log(a x )
⎣ J uˆ
⎦
χα F
αF
and
(2)
−1
⎡ ⎛ χα FΔV ⎞ ⎤
tind = ξ ′⎢ exp
− 1 +τ
⎣ ⎝ 2RT ⎠ ⎥⎦
(3)
respectively, where χ is the film stoichiometry ( MO χ / 2 ), α is the dependence of the potential
drop across the film/solution interface on applied voltage, ax is the activity of halide ion in the
solution, τ is a relaxation time (τ ~0), Δ V = V-Vc, V is the applied voltage (or the corrosion
potential under open circuit conditions),
(4)
o
J = aˆ D
and
N
ˆa = χ K⎡ v ⎤
⎣Ω⎦
1+
χ
2
[
exp − ΔGs /RT
o
]
(5)
The parameters u and Jm are defined in the original derivation [8].
Assuming that the breakdown sites are normally distributed in terms of the cation
vacancy diffusivity,
dN
=
dD
1
2πσ D
[
exp −(D − D) / 2σ D
2
2
]
(6)
where D is the mean diffusivity and σ D is the standard deviation, the point defect model
yields the following expressions for the distributions in the breakdown voltage and induction
time [9-11]:
2
2
γˆD
dN
=−
⋅ e−(D− D) / 2 σ D
(7)
2πσ D
dVc
and
23
Plenary Lectures
ˆ
dN ⎡ −ξu χ / 2 ⎤ −(D −D )2 / 2 σ D2
e − γV
=
⋅ χ /2
e
2
dt ind ⎢⎣ 2πσ D aˆ ⎥⎦
a x (t ind − τ )
where
γˆ = χFα/2RT
(8)
(9)
Eqs. 7 and 8 contain four important system parameters: D and σ D which describe the
transport properties of cation vacancies in the passive film, and α and β which appear in the
expression for the dependence of the potential drop across the film/solution interface on
applied voltage and pH
φf/s = α V + βpH + φf/so
(10)
All four parameters can be determined by independent experiments: D from electrochemical
impedance spectroscopy, σ D from passivity breakdown induction time measurements [8], and
α and β from film growth data, much of which exists in the literature. Accordingly, Eqs. 7
and 8 represent analytical distribution functions for the nucleation of pitting attack, provided
that the assumptions in the model hold. Our previous work [9-11] has demonstrated the
quantitative nature of these expressions for representing experimental distribution in Vc and
tind, for those cases where sufficient experimental data are available for analysis.
The Pit Growth Module (PGM)
The pit growth module (PGM) computes the growth rate of an individual pit. Figure 4
shows schematically a typical pit that develops on the surface of a metal in contact with a thin
liquid layer. The module developed in this study calculates the pit growth rate for a pit of
cylindrical geometry. The details of the theoretical approach can be found in the literature
[12,13], except that we developed a transmission line analog of the external environment
from which the distributions in the electrostatic potential and the current in the external
environment, as a function of distance away from the pit, can be estimated.
z
Gas
Thin electrolyte film
h
a0
r
L
Metal
24
Macdonald
Fe + n 4 H 2 O − Fe
( OH ) (n 2 − n
Cr + n 5 H 2 O − Cr ( OH
) (n
Ni + n 6 H 2 O − Ni ( OH
) (n 2 − n
H2O − H
+
+ OH
4
)−
4
3−n
5
)−
6
)
5
6
+ n4H
+
+ 2e
−
+ n5H
+
+ 3e
−
+ n6H
+
+ 2e
−
−
−
Figure 4. Schematic diagram of a pit on the surface of the condensing heat exchanger
The principle of the transmission line approach is shown in Fig. 5, which yields the
following equations for the distributions in the electrostatic potential (φs ) and the current in
the external environment
d 2φs 1 dφ s
ρ
−
φ =0
2 +
dr
r dr hZs s
(11)
d 2 I 1 dI
ρ
−
I= 0
2 −
dr
r dr hZs
(12)
where ρ is the resistivity ( Ω⋅ cm ) of the solution, Zs is the specific impedance ( Ω⋅ cm 2 ) of
the external surface, h is the electrolyte film thickness, and r is the radial distance from the
C
center of the cylindrical pit. Note that Zs is a function of distance [Zs(r) =- (φ s − φm )/iN , where
i CN is the net cathodic current density]. The value of Zs(r) was determined iteratively when
solving Eq. 11 by substituting for the net current density the following expression
e−(φ s −φ s )/b a − e(φ s −φ s )/b c
i =
+i φ
1 1 -(φ s −φ se )/b a 1 ⎛⎜⎝φ s −φ se⎞⎟⎠ /b c p ( s )
+ e
− e
il,r
i o il,f
e
e
C
N
(13)
where the first term on the right-hand side is the generalized Butler-Volmer equation for the
oxygen electrode reduction
2H 2 O ⇔ O 2 + 4H+ + 4e−
(14)
and the second is the polarization current of the substrate, both of which are functions of the
potential difference across the interface. The parameters φse , io, il,f, il,r, ba, and bc in Eq. 13 are
the negative of the equilibrium potential for Reaction 14, io is the exchange current density, il,f
and il,r are the limiting current densities for Reaction 14 in the forward and reverse directions,
respectively, and ba and bc are the corresponding Tafel constants. Note that the signs of the
exponents in the first term in Eq. 13 are opposite to those normally defined because we have
written the current in terms of the electrostatic potential in the solution with respect to the
metal. We used the finite difference method to solve Eqs. 11 and 12 for φs (r) and i c(r) ,
respectively.
25
Plenary Lectures
The distribution of the electrostatic potential within the pit is obtained by solving
Laplace's equation, assuming that the environment within the pit confine is electrically
neutral,
∇φ= 0
2
(15)
The solution to Laplace's equation (Eq. 15) yields the following expression, assuming
that the potential variation in the radial direction is negligible compared to that in the
longitudinal direction:
Z
Electrolyte film
Flue gas (O2)
r
r+dr
Crevice
Metal
Figure b
(a) Element of electrolyte film on the metal surface.
I
Rdr
φS
I-dI
Z(r)/dr
φm
(b) Element of transmission line for calculating current and potential
distributions radially from the crevice mouth.
Figure 5. Transmission line model for thin electrolyte film on the metal surface
φs (z) = (φ s0 − φ s−L ) +φ s0
z
L
(16)
where φs−L is the electrochemical potential at the pit tip, L is the pit depth, and z is a negative
quantity. We also apply the Butler-Volmer equation to the electrodissolution reaction
occurring at the pit bottom to yield the electrochemical potential at the pit tip as [14]
⎛ i 00 A ⎞
⎟
⎝ I0 ⎠
φs−L = φ s00 + ba1n ⎜
(17)
where φs00 is the (negative of the) standard electrochemical potential for the dissolution of the
metal, i 00 is the standard exchange current density, ba is the anodic Tafel constant for metal
dissolution, and At is the effective active surface at the pit tip. The model outlined above is a
variant of the Coupled Environment Fracture Model (CEFM) that we developed some time
ago [15] for describing crack growth in stainless steel piping in nuclear power reactor heat
26
Macdonald
transport circuits. Thus, following our previous work [15], Eqs. 11, 12, 16, and 17 are solved
for the unknowns φs (r) , i c(r) , and Io, such that charge conservation, expressed as
I0 + ∫ i C dS= 0
N
(18)
s
is obeyed, where I0 is the (positive) current exiting the pit mouth, and dS is an increment of
the external surface (dS = 2π rdr ). Because the cathodic current due to oxygen reduction
predominates on the external surface, the second term on the left side of Eq. 18 is negative.
Once I0 is known, then the pit growth rate is calculated using Faraday's law:
dL
M I0
=
dt 2ρ m Z FA
(19)
where ρm is the density of the metal (g/cm3), M is the composition-averaged atomic weight
of the alloy, and Z is the composition-weighted oxidation state of the metal dissolving at the
pit tip. Finally, the pit length is calculated as a function of time using the recursive formula:
L(t) = L(t − 1) +
dL
Δt
dt
(20)
where L(t-1) is the depth of the pit calculated from the previous time (t-1), and Δt is the
increment in time.
The Damage Function Module (DFM)
By combining Eqs. 8 and 20 for a fixed density of potential breakdown sites (No,
number/cm2), it is possible to estimate the pitting damage function. Thus, if one observes the
system at time tobs, then the number of pits that nucleate over the time increment Δt at tind is
ΔN , as determined from Eq. 8. However, these pits will have grown to a depth L(t), as given
by Eq. 20, at the time the system is examined. By moving the increment Δt from t = τ to
t = t obs , the damage function is then generated in the form of the number of pits versus the
depth of the pits. If this procedure is repeated for different observation times, a family of
damage functions is generated that extends to greater depths with increasing tobs.
By specifying the surface area of interest, it is possible to define the service life as the
time taken for one or more pit to grow to the critical length, which in this case corresponds to
the wall thickness of the condensing heat exchanger. The number of pits with lengths
exceeding the critical dimension is simply calculated as
L max
N L ≥L crit = S∑ N(L) ⋅ ΔL
(21)
L crit
where N(L) is the density of pits per unit surface area and per unit increment in pit length
(number/cm3) in the damage function, S is the surface area of interest (cm2), L is the pit
27
Plenary Lectures
length, ΔL is the increment of the pit length in the damage function, and Lcrit is the critical
dimension. The service life is simply the time at which N L ≥L crit = 1.
DISCUSSION
The procedure outlined above for estimating damage functions for localized corrosion is
currently being developed to explore the impact of corrosion on condensing heat exchangers
in domestic and industrial gas-fired furnaces. The practical problem lies in selecting the most
cost-effective alloys for the condensing stages of heat exchangers, because of the highly
competitive nature of the furnace manufacturing business. Consequently, little room exists
for over-designing furnaces by employing highly-alloyed, costly materials to fabricate the
condensing sections. Therefore, selection of materials with adequate pitting resistance, and of
acceptable cost, is of prime concern to furnace manufacturers and users alike. It is evident,
then, that the design and materials specifications for condensing heat exchangers would
greatly benefit from the development of a deterministic method for predicting localized
corrosion damage functions. This, in turn, could reduce the cost of the alloy by decreasing
the required database, through the availability of deterministic relationships between the
damage function and important environmental variables (including pH, [Cl-], and gas
composition).
In this study, I present predictions of the model in comparison with experimental damage
functions measured on Type 304L stainless steel by G. Stickford, B. Hindin, and A.K.
Agrawal of The Battelle Columbus Laboratories. For the experimental data, the damage
functions are measured on condensing heat exchanger tubes after a given number of cycles, at
the hot end (temperature ranging from 326o to 353oK) and at the cold end (temperature
ranging from 308o to 326oK). Each cycle consists of 240 second with the burner on and 480
sesond with burner off, which represent the dry and wet conditions, respectively. It was
shown in a previous study [4], that no significant difference exists in the damage functions
between the hot end and the cold end; the damage functions are, therefore, plotted without
distinction between the hot ends and cold ends. Based on this experimental finding, the
model is constrained to the case where the surface is covered at all times by a condensing
liquid phase (wet condition). However, I choose the appropriate temperature at the hot end to
calculate the damage functions in order to avoid underestimating the damage.
The experimental data reported by Battelle were measured at three levels of chloride
concentration (3, 26, and 225 ppm) on a number of different candidate alloys. I present in
this study only the damage functions for Type 304L stainless steel, as shown in Figs. 6, 7, and
8, as a function of chloride concentration. Not surprisingly, fewer pits were observed at the
lower chloride concentration (3 ppm, Fig. 6). At higher chloride levels (26 and 225 ppm), the
number of pits increased substantially (Figs. 7 and 8) and led to perforation of the wall in
shorter time, thus reducing the service life. However, the experimental data show some
inconsistencies, which are due to the fact that different tubes were used to determine the
damage functions in each case. The chemical composition, the metallurgical history, and the
surface state may vary from tube to tube.
Because, the kinetics of the cathodic oxygen reaction on Alloy Al29-4C are considered
to be essentially identical to those on Type 304L stainless steel, the parameters for Al29-4C
were chosen for calculating the corrosion potential used in estimating the damage functions
28
Macdonald
(Table 1). Calculated damage functions are presented in Figs. 9 through 12 for chloride
levels of 3, 10, 26 and 225 ppm, respectively. The calculations clearly indicate the
progressive nature of the nucleation and growth of pits on the alloy surface. It is predicted
that at lower chloride concentrations (3 ppm), fewer pits exist on the surface of the steel,
while at higher chloride concentrations (26 to 225 ppm), the number of pits increases
substantially, leading to the majority of the pits perforating the wall thickness in a short
period of time.
The predicted service life is presented as a function of chloride concentration in Fig. 13,
in comparison with the experimental data. The calculations indicate that the service life of a
condensing heat exchanger is highly sensitive to the chloride level in the condensate,
especially at the lower chloride concentration (3 ppm). The principal effect of increasing
chloride is to accelerate pit nucleation, so that, in the limits of very high chloride
concentration, (~>100 ppm) in the condensate, the failure time is dominated by pit growth.
Because the pit growth rate is dominated by the conductivity of the external environment (i.e.,
the condensate film), for any given pH and oxygen concentration, and because the
conductivity is dominated by non-chloride species, the failure time becomes constant at
sufficiently high chloride levels. This corresponds to the situation where the entire service
life is determined by the time required for the pits that nucleate on initial exposure of the
alloy to condensate and grow through the condenser wall. Noting that the service life for the
case shown in Fig. 13 is calculated to decrease from 1.55 x 108 s (4.92 years) for a chloride
concentration of 3 ppm to 2.34 x 107 s (0.74 year) for a chloride concentration of 225 ppm, it
is evident that the time required for an active pit to perforate the wall is about three-quarters
of a year, corresponding to an average pit growth rate of 0.7 mm/year. Clearly, then, the
increase in the service life on lowering the chloride concentration is due almost entirely to an
increase in the initiation time, and it would seem that substantial service lives for this alloy
can only be obtained if nucleation becomes the dominant phase in the development of
damage. Finally, due to the fact that different tubes are used in determining the damage
functions, the experimental data are rather scattered. In recognition of this observation,
relatively good agreement is claimed between the experimental data and the theoretical
prediction.
Table 1. Values for Parameters Used in Calculating Damage Functions
Parameter
χ
Ω
ΔG AO −1
φf/sO
O
Δ Gs
τ
ε
α
β
ξ
(Passive film stoichiometry)
(Mole volume of passive film)
(Gibbs energy of Cl- absorption)
(Constant)
(Gibbs energy of cation vacancy formation)
(Relaxation time)
(Electric field strength)
(Constant)
(Constant)
(Critical areal concentration of vacancies)
Value
3
30
-60
-0.375
20
0
1 x 106
0.25
-0.001
1 x 1016
Units
cm3/gm cation
kJ/mol
V
kJ/mol
s
V/cm
V
No/cm2
29
Plenary Lectures
Jm
D
σD
(Vacancy flux in metal phase)
(Standard deviation in cation
diffusivity)
(Standard deviation in cation
diffusivity)
0.12 x 107
vacancy 1.0 x 10-18
vacancy
Vacancies/cm2.s
cm2/s
0.5 D
cm2/s
The influence of the oxygen partial pressure on the development of damage functions has
been calculated, and is shown in Fig. 14. The calculations indicate that oxygen has a great
impact on the service life of heat exchangers. This is because oxygen, in the condensed
liquid phase on the external surface, consumes the positive current associated with the pit tip
dissolution process, thereby driving the growth of the pit. By decreasing the partial pressure
of oxygen from 10-4 atm to 10-8 atm, the service life of a heat exchanger having the
characteristics assumed in this work could be extended from 3.02 x 107 s (1 year) to 1.05 x
109 s (approximately 30 years).
Figure 6. Measured pitting damage functions for Type 304L stainless steel heat
exchanger tubes under condensing conditions: [Cl-] = 3 ppm,
temperature ranging from 308o to 353oK, pit counting interval = 2.54 x
10-3 cm. a) Tobs = 4.32 x 106 s; b) Tobs = 8.63 x 106 s; c) Tobs = 1.73 x
107 s; tube thickness = 5.34 x 10-2 cm, as indicated by the dashed line
Figure 7. Measured pitting damage functions for Type 304L stainless steel heat
exchanger tubes under condensing conditions: [Cl-] = 26 ppm,
temperature ranging from 308o to 353oK, pit counting interval = 2.54 x
10-3 cm. a) Tobs = 4.32 x 106 s; b) Tobs = 8.63 x 106 s; c) Tobs = 1.73 x
107 s; d) Tobs = 3.46 x 107 s; e) Tobs = 5.18 x 107 s; tube thickness = 5.34
x 102 cm, as indicated by the dashed line
Figure 8. Measured pitting damage functions for Type 304L stainless steel heat
exchanger tubes under condensing conditions: [Cl-] = 225 ppm,
30
Macdonald
temperature ranging from 308o to 353oK, pit counting interval = 2.54 x
10-3 cm. a) Tobs = 4.32 x 106 s; b) Tobs = 8.63 x 106 s; c) Tobs = 1.73 x
107 s; tube thickness = 5.34 x 10-2 cm, as indicated by the dashed line
Figure 9. Calculated pitting damage functions for Type 304L stainless steel heat
exchanger tubes under condensing conditions: [Cl-] = 3 ppm,
temperature = 350oK, PCO = 2.60 x 10-5 atm, PCO2 = 6.80 x 10-2 atm,
PSO2 = 9.37 x 10-9 atm, PSO3 = 2.00 x 10-9 atm, PH2S = 1.20 x 104 atm,
PNO = 1.74 x 10-8 atm, PNO2 = 2.67 x 10-8 atm, PO2 = 1.00 x 10-4 atm, pit
radius = 1.00 x 10-4 cm, initial pit depth = 1.00 x 10-3 cm, pit counting
interval = 2.54 x 10-3 cm, a) Tobs = 6.30 x 107 s; b) Tobs = 1.58 x 108 s;
tube thickness = 5.34 x 10-2 cm, as indicated by the dashed line
Figure 10. Calculated pitting damage functions for Type 304L stainless steel heat
exchanger tubes under condensing conditions: [Cl-] = 10 ppm,
temperature = 350oK, PCO = 2.60 x 10-5 atm, PCO2 = 6.80 x 10-2 atm,
PSO2 = 9.37 x 10-9 atm, PSO3 = 2.00 x 10-9 atm, PH2S = 1.20 x 10-4 atm,
PNO = 1.74 x 10-8 atm, PNO2 = 2.67 x 10-8 atm, PO2 = 1.00 x 10-4 atm,
pit radius = 1.00 x 10-4 cm, initial pit depth = 1.00 x 10-3 cm, pit
counting interval = 2.54 x 10-3 cm, a) Tobs = 3.15 x 107 s; b) Tobs =
4.10 x 107 s; c) Tobs = 4.41 x 107 s d) Tobs = 5.36 x 107 s e) Tobs = 6.30
x 107 s; tube thickness = 5.34 x 10-2 cm, as indicated by the dashed
line
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Plenary Lectures
Figure 11. Calculated pitting damage functions for Type 304L stainless steel heat
exchanger tubes under condensing conditions: [Cl-] = 26 ppm,
temperature = 350oK, PCO = 2.60 x 10-5 atm, PCO2 = 6.80 x 10-2 atm,
PSO2 = 9.37 x 10-9 atm, PSO3 = 2.00 x 10-9 atm, PH2S = 1.20 x 10-4 atm,
PNO = 1.74 x 10-8 atm, PNO2 = 2.67 x 10-8 atm, PO2 = 1.00 x 10-4 atm,
pit radius = 1.00 x 10-4 cm, initial pit depth = 1.00 x 10-3 cm, pit
counting interval = 2.54 x 10-3 cm, a) Tobs = 4.32 x 106 s; b) Tobs =
8.63 x 106 s; c) Tobs = 1.73 x 107 s d) Tobs = 2.52 x 107 s e) Tobs = 2.84
x 107 s; f) Tobs = 2.99 x 107 s; g) Tobs = 3.15 x 107 s; h) Tobs = 3.46 x
107 s; i) Tobs = 5.19 x 107 s; tube thickness = 5.34 x 10-2 cm, as
indicated by the dashed line
Figure 12. Calculated pitting damage functions for Type 304L stainless steel heat
exchanger tubes under condensing conditions: [Cl-] = 225 ppm,
temperature = 350oK, PCO = 2.60 x 10-5 atm, PCO2 = 6.80 x 10-2 atm,
32
Macdonald
PSO2 = 9.37 x 10-9 atm, PSO3 = 2.00 x 10-9 atm, PH2S = 1.20 x 10-4 atm,
PNO = 1.74 x 10-8 atm, PNO2 = 2.67 x 10-8 atm, PO2 = 1.00 x 10-4 atm,
pit radius = 1.00 x 10-4 cm, initial pit depth = 1.00 x 10-3 cm, pit
counting interval = 2.54 x 10-3 cm, a) Tobs = 1.58 x 107 s; b) Tobs =
2.21 x 107 s; c) Tobs = 2.52 x 107 s d) Tobs = 3.15 x 107 s; tube
thickness = 5.34 x 10-2 cm, as indicated by the dashed line
Figure 13. The measured and calculated service life for Type 304L stainless
steel heat exchanger tubes as a function of the chloride
concentration (parameters are identical to that for Figs. 6-8 for
experimental data and Figs. 9-12 for calculation)
33
Plenary Lectures
Figure 14.
The calculated service life for Type 304L stainless steel heat
exchanger tubes as a function of the partial pressure of oxygen
(temperature = 350oK, PCO = 2.60 x 10-5 atm, PCO2 = 6.80 x 10-2 atm,
PSO2 = 9.37 x 10-9 atm, PSO3 = 2.00 x 10-9 atm, PH2S = 1.20 x 10-4 atm,
PNO = 1.74 x 10-8 atm, PNO2 = 2.67 x 10-8 atm)
SUMMARY AND CONCLUSIONS
A deterministic model has been developed to predict the damage functions for
condensing heat exchangers in gas-fired furnaces. The model incorporates calculations for
the condensed chemistry environment, the electrochemical corrosion potential of the alloy,
and mechanistic treatments of the nucleation and growth of pits.
The model predicts that the chloride concentration in the condensed liquid layer has
great impact on the service life of the condensing heat exchanger, particularly at low chloride
concentrations. At high chloride concentrations, the service life of the condensing heat
exchanger is predicted to be relatively independent of the chloride concentration,
corresponding to the dominance of pit growth in determining the failure time. The service
life for the condensing heat exchanger with Type 304L stainless steel tubes is predicted to
decrease from 1.55 x 108 s (4.92 years) for a chloride concentration of 3 ppm to 2.34 x 107 s
(0.742 year) for a chloride concentration of 225 ppm.
The model predicts that the service life of the condensing heat exchanger also depends
strongly on the oxygen content in the flue gas; by decreasing the oxygen partial pressure from
10-4 atm to 10-8 atm, the service life of the condensed heat exchanger can be extended from
3.02 x 107 s (1 year) to 1.05 x 109 s (approximately 30 years).
Recognizing the scattered nature of the experimental data, I conclude that the algorithm
developed in this work provides estimates of the service life that are in good agreement with
the available experimental data, even though no a priori fit of the experimental data to the
model was made.
ACKNOWLEDGEMENTS
The author gratefully acknowledges the support of this work by the Gas Research
Institute (GRI) through Contract No. 5090-260-1969, and G. Stickford, B. Hindin, and A.K.
Agrawal at The Battelle Columbus Laboratories for supporting the experimental damage
functions used in this study.
REFERENCES
1. D.D. Macdonald and M. Urquidi-Macdonald, "The Corrosion Damage Functions:
Interface between Science and Engineering," 1992 Whitney Award Address, NACE,
Nashville, Tenessee, submitted to Corrosion, 1992.
2. R. Razgaitis, J.H. Payer, S.G. Talbet, B. Hindin, E.L. White, D.W. Locklin, R.A.
Cudnik, and G.H. Stickford, Condensing Heat Exchanger Systems for
Residential/Commercial Furnaces and Boilers, Phase II, Battelle Report to DOE/BNL,
BNL Report No. 51943, October, 1985.
34
Macdonald
3. D.D. Macdonald, M. Urquidi-Macdonald, S.D. Bhakta, N. Khalil, and H. Yashiro,
Development of Analytical Methods for Predicting Damage Functions for Pitting
Corrosion in Condensing Heat Exchangers, Final report to the Gas Research Institute,
GRI No 5090-260-1969, January, 1992.
4. G.H. Stickford, B. Hindin, S.G. Talbert, A.K. Agrawal, M.J. Murphy, R. Razgaitis,
J.H. Payer, R.A. Cudnik, and D.W. Locklin, Technology Development for CorrosionResistant Condensing Heat Exchanger, Final Report to the Gas Research Institute,
GRI-85/0282NTIS PB86-172038, October, 1985.
5. C. Wagner and W. Traud, Z. Electrochem. 44, 1938, p. 391.
6. D.D. Macdonald, Corrosion 48, 1992, p. 194.
7. D.D. Macdonald, Proc. 5th Int. Symp. Environ. Degrad. Mat. Nucl. Power Systs:
Water Reactors, Monterey, California, NACE, August, 1991.
8. L.F. Lin, C.Y. Chao, and D.D. Macdonald, J. Electrochem. Soc. 128, 1981, p. 1194.
9. D.D. Macdonald and M. Urquidi-Macdonald, Electrochim. Acta 31, 1986, p. 1079.
10. D.D. Macdonald and M. Urquidi-Macdonald, J. Electrochem. Soc. 134, 1987, p. 41.
11. D.D. Macdonald and M. Urquidi-Macdonald, J. Electrochem. Soc. 136, 1989, p. 961.
12. D.D. Macdonald, M. Urquidi-Macdonald, C. Liu, S. Bhakta, N. Khalil, and H.
Yashiro, Proc. Int. Gas Res. Conf., Orlando, Florida, November, 1992.
13. D.D. Macdonald, M. Urquidi-Macdonald, and C. Liu, Paper No. 173, CORROSION
93, New Orleans, Louisiana, March, 1993.
14. D.D. Macdonald and M. Urquidi-Macdonald, Corros. Sci. 32, 1991, p. 51.
35
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
RELEVANCE OF LABORATORY CORROSION TESTS IN CORROSIVITY
ASSESSMENT AND MATERIALS SELECTION: CASE STUDIES
R.D. Kane
CLI International, Inc.
14503 Bammel-N. Houston, Suite 300, Houston, TX USA
ABSTRACT
Laboratory tests are a convenient means for simulating service environments for the purpose of
evaluating both corrosiveness of the environment and material performance. Such tests can provide
substantial information upon which engineering decisions can be made. These decisions generally will
be made with greater confidence and a more efficient materials design. Measurements can be made
under more controlled and reproducible conditions than are possible in field or plant situations. The
consequences of process changes can be evaluated in advance of the actual situation. However, when
conducting tests in simulated service environments, it is most important that the experimental methods
be designed with the specific intent to provide meaningful, representative, correlative results. This
presentation will review the various aspects that must be included when designing, conducting, and
interpreting corrosion tests in simulated service environments. Several case studies will also be
presented that involve petroleum applications, including the selection of corrosion-resistant alloys in
sour petroleum production, inhibition of multiphase flow lines, assessment of crude oil corrosiveness,
and corrosion under insulation.
Key Words:
Laboratory testing, correlation, oil and gas, multiphase production,
corrosion under insulation, electrochemistry
INTRODUCTION
There are many types of laboratory corrosion tests utilized for various purposes. These
tests can range from simple glassware immersion tests involving freely corroding, nonstressed coupons to highly sophisticated exposure tests involving dynamic replenishment,
heat and/or mass transfer and high pressures maintained only through autoclave, flow loop or
pilot plant operations. Furthermore, they may also involve ancillary techniques which
include a variety of DC or AC electrochemical methods and mechanical loading
configurations. To extend the predictive capabilities of the tests results, modeling techniques
and statistical experimental design programs can be employed that allow for identification or
verification of mechanisms or for establishment of linkages between the results of the
laboratory test and the intended service application while minimizing the number of tests
required.
LABORATORY TESTING: A BASIC AND IMPORTANT TOOL
In many areas of technical activity, emphasis is now being placed on reducing the overall
time, funding and effort allocated for testing, research, development and engineering. In
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Plenary Lectures
some cases, corrosion testing appears as only an afterthought. At the same time, there are
pressures to provide systems with higher reliability and reduced operating costs. In simpler
words, these engineering concepts can be referred to in four basic terms: Better, Cheaper,
Faster, Safer [1]. One of the most important aspects that must be realized is that these are not
necessarily mutually inclusive concepts. Usually, substantial scientific and technological
developments are required before these terms can come together in a complex engineering
system.
Laboratory testing is one of the basic tools available to investigate complex interactions
of variables that exist in real service applications. Data can be developed which both have
applied engineering significance and provide insight into fundamental relationships in
engineered systems that can in turn be used as stepping stones to achieve quantum leaps in
efficiency, reliability and safety [2]. Furthermore, when compared to the costs generally
associated with corrosion (which can be around 10% of revenues [3]), and those saved
through application of corrosion testing, there can often be a cost reduction of between one
and two orders of magnitude and, in some cases, more.
NEED FOR SIMULATED-ENVIRONMENT TESTS
The need for simulation varies greatly depending on the purpose of the test. For example,
corrosion tests are usually conducted for three reasons [4]: screening-comparison of the
response of two or more materials or material conditions relative to a particular form of
corrosion (e.g., general corrosion, local pitting or crevice corrosion, and stress corrosion
cracking); qualification-verification that the material has a required conformance to
composition and that the metallurgical or fabrication processes have resulted in a
microstructure that will provide adequate corrosion performance; evaluation-assessment of
the influence of process changes (e.g., temperature, additives, inhibitors, and product purity).
In many cases, relatively simple, standardized tests found in NACE, ASTM, and ISO
documents can be conducted that are useful for these purposes. These are relatively simple
environments often involving combinations of acids and salts that can be handled in standard
glassware. [5-7]. However, a limitation common to many standardized tests is the difficulty
in obtaining a direct correlation between the performance in such tests and actual in-service
performance. To expand the applicability of standardized tests, substantial development work
is often required which establishes such correlations. This work usually includes one or more
of the following: review of service or failure records on a range of material conditions or
over a range of process variables, analysis of field or in-plant tests in actual service
environments, and use of laboratory exposure tests conducted under simulated service
conditions. In many situations, the laboratory provides the most convenient avenue since it
yields firsthand data that allows the engineers to make the fine distinctions in performance
required to select the most cost-effective corrosion mitigation or control methods, and thereby
gives the maximum cost benefit.
THE REASONABLE WORST-CASE SCENARIO
In the study of corrosion in simulated environments, the concept of a reasonable worstcase scenario has been a guiding light for those involved. Being reasonable includes two
aspects: reasonable levels of corrosive severity and reasonable involvement of expense and
time. It is often necessary to refine the set of exposure conditions by a process of
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Kane
prioritization so that only the most important variables related to the operable corrosion
mechanisms are taken into account and thereby minimize testing costs. Keeping these two
aspects in mind, one must first define and preserve the active in-service mechanism(s) of
corrosion. Once this has been completed, one must identify the aggravating and mitigating
factors that combine in the actual service condition to determine the severity of corrosion.
SIMULATED ENVIRONMENT TEST PROCEDURES
Several case studies are presented herein which highlight a few of the basic, yet
important, concepts which provide the link between the laboratory and field conditions.
Case Study No. 1: Simulation of Conditions of Deaeration or Aeration
One of the most important and universally applicable situations that must be evaluated
when conducting laboratory tests in simulated service environments is the need to produce a
reasonable representation of the level of deaeration or aeration found in the actual
environment. The main reason for the importance of this effect is that corrosivity, in many
service applications, can change dramatically with changes in oxygen content. Aeration
accelerates anodic corrosion processes with a concomitant increase in localized corrosion
activity (i.e., pitting, crevice attack and stress corrosion cracking). Examples of oxygen
effects can be seen in applications such as seawater injection, the use of heavy brine
completion fluids in oilfield operations and desalination. As shown in Fig. 1, the corrosion
rate of steel increases by an order of magnitude going from 10 ppb to just 100 ppb [8]. It only
takes very low levels of oxygen contamination (about 1% of normal atmospheric saturation
levels) to greatly accelerate corrosion.
Furthermore, due to the sensitivity of corrosion reactions even in low levels of aeration,
oxygen contamination can produce excursions to higher corrosion rates that have prolonged
effects [9]. The increase in localized anodic attack produced by aeration can be illustrated by
its interaction with other species such as chlorides and sulfides. An example, in terms of
susceptibility to stress corrosion cracking (SCC), is the interaction between dissolved oxygen
and chlorides in elevated temperature applications involving alkaline phosphate treated boiler
water [10]. As the availability of oxygen increases above the 0.1 ppm (100 ppb) level, the
tolerance for chloride is reduced resulting in a dramatic increase in susceptibility to SCC.
Figure 1. Corrosion rate vs temperature for various oxygen levels
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Plenary Lectures
The situation in aqueous sulfide-containing environments is even more complicated
since oxygen can result in the formation of elemental sulfur, acids, and in some cases,
polysulfide species. These can synergistically interact with the oxygen effects mentioned
previously to produce quite severe limitations on the corrosion and SCC performance of even
very highly alloyed materials. Such conditions can be found in applications which involve
pumping of sour oilfield brines, injection of wastewater and flue gas desulfurization. A
comparison of the minimum required pitting resistance equivalent (PRE) for conditions
involving a simulated sour oilfield service can be seen for identical situations (0.7 kPa H2S,
138kPa CO2, 2 meq/L HCO3, 30,000 ppm Cl-, 65oC) except one is aerated and one deaerated
[11]. For the deaerated condition, the minimum pitting resistance equivalent [PRE = Cr +
3.3Mo + 11N + 1.5(W + Cb)] is only 12. This would indicate successful use of materials
with > 12 Cr. However, this same environment under aerated conditions yields a minimum
PRE value of 30. Under evaporative conditions, this can increase still further.
Understanding the conditions of aeration in the service application is necessary to
reproduce similar conditions in the laboratory corrosion test. For example, most geochemical
systems naturally contain less than 10 ppb oxygen. By comparison, mechanical deaeration
techniques usually will not go below 100 ppm. Multiple vacuum, ultra-low oxygen inert gas
purge cycles and prolonged gas purges are usually required to get below 50 ppb oxygen. In
some cases, oxygen scavengers must be used to obtain complete deaeration. However, these
must be used carefully because they may, in some cases, add other chemical species into the
environment that can complicate electrochemical measurements.
Case Study No. 2: Simulation of Corrosion in Multiphase Environments
There are many factors that need to be considered when conducting corrosion
assessments in multiphase environments. These include important factors related to the
dynamic or flowing nature of the fluids which determine the mode of flow [12] and the
kinetic shear forces that are imparted by the flowing fluids on the pipe wall. There have been
several major studies involving very sophisticated simulations of three-phase flow. These
studies are particularly capital intensive and costly since major investments must be made in
the handling, pumping and disposal facilities required for such tests. However, there are no
real alternatives for investigating questions involving the direct effects of flow regime such as
measuring the shear forces developed by particular flow regimes and operating conditions
and the movement of inhibitors [13].
On a more practical basis, more simple yet reasonable approximations of multiphase
flow conditions can be obtained using pseudo-three phase systems such as the flow loop
shown in Fig. 2 [14]. These systems provide for the establishment of three phase conditions
(gas/oil/water) in a reservoir autoclave. Under these conditions, the primary corrodents are
dissolved gases (e.g., CO2, H2S and sometimes O2) and an aqueous brine or condensed water
phase. Facilities for replenishment of both the gas and liquid phases must be considered
depending on the exact nature of the environment. Simulation of the affects of a flowing
environment is usually based on modeling the shear stress produced in service on the metal
surface by the flowing liquid containing the dissolved gases using the equations given in
Table 1 [15]. The main assumption utilized in this approximation is that the major
contribution to the wall shear stress is usually made by the liquid phase. In most cases, this
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Kane
technique is valid since the contribution of liquid phase density and viscosity on the resultant
shear stress predominates over that of the gas phase.
Figure 2. MAPS™ - Multiphase Autoclave Pipeline Simulator
Of significance in most flowing multiphase systems is the handling of slug flow which is
the predominate flow regime for horizontal and near horizontal flow applications. The main
attribute of slug flow is the very high shear stresses and accompanying high turbulence in the
region of the flow just ahead of the moving slug [16]. This effect results in levels of shear
stress much greater than those produced by the bulk fluid. It has been proposed that this is the
location where excessive corrosion is generated as a result of the effect of locally high shear
stress and turbulence on both corrosion and inhibitor films. Investigations have recently
focused on techniques such as flow loops and jet impingement to reproduce accurate
simulations of such highly turbulent conditions for assessment of corrosion resistance and
inhibitor performance [17].
Another major effect that must be addressed in multiphase systems is the potential role
of the oil phase as a possible mitigation factor in terms of reducing the corrosion rate [18].
The properties of the hydrocarbon/liquid phase significantly influence the severity of the
environment with respect to weight loss corrosion (see Fig. 3). In a typical case, oil/water
mixtures remain relatively non-corrosive under flowing conditions of up to about 30% water
cut resulting mostly from the preferential wetting and persistence of the oil phase on the metal
surface. However, depending on the nature of the liquid hydrocarbon phase, some cases
become corrosive with very low water cuts (<5%) while other cases do not become corrosive
until more than 50% water is reached. The exact variables that relate to these mitigating
41
Plenary Lectures
effects have only been qualitatively investigated. Therefore, when simulating service
applications, the exact nature of the oil phase should be considered since it may play a major
role in the overall corrosivity of the system and in the efficacy of inhibitor treatments.
Furthermore, the presence of the nonconductive oil phase in multiphase tests can also produce
confusing results when electrochemical techniques are used as the sole basis for evaluation.
Care must be utilized in the selection of corrosion monitoring techniques and the comparison
of this data with the results of physical examination, mass loss and localized corrosion
measurements.
Table 1. Flow/Shear Stress Relationships
Description
Relationship
ρVD
Determine dimension-less
Re =
μ
parameters to describe fluid flow
characteristics (e.g., Reynold’s
number) to account for mass
transfer effects
f = z (Re, e/D)
Determine friction factor, f, to
account for pipe wall roughness
(from Moody diagrams)
fρV 2
Determine wall shear stress, t, as
τ=
2
a function of the friction factor
and other flow properties.
Vt − Vs
Determine flow regime (annular,
fr =
g . heff
stratified, bubble, slug, etc.) to
estimate correction factors (e.g.,
for slug flow, Jepson et al. use
the Froude number as a basis to
estimate turbulent intensity).
⎡r⎤
Laboratory simulation: Jet
τ = 0.0112ρV R
⎢r ⎥
impingement (Giralt and Trass)
⎣ ⎦
Laboratory simulation: Rotating
τ = 0.0791Re−0 . 3 ρr 2ω 2
cylinder electrode (D. C.
Silverman)
Summary: The corrosion rate in fully developed turbulent pipe flows
computed from field parameters can be simulated in the laboratory
produced. can be expressed in terms of wall shear stress. Wall shear stress
through experimental methods, and hence similar corrosion rates.
−2 . 0
2
−0 .182
e
0
Case Study No. 3: Simulation of Geometry of Exposure
One of the factors that can have a great impact on corrosion severity is the geometry of
the service application. Obvious cases are those involving crevices, seams, laps and welds
where the formation of an occluded cell can result in differences between local and bulk
solutions. In some cases, the whole service condition may bring together somewhat unique
combinations of solution and geometric variables which cannot be accurately simulated by
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Kane
simple immersion or atmospheric tests. One such case that illustrates this situation is
exhibited by corrosion under insulation (CUI).
Figure 3. Effect of hydrocarbon/liquid phase on weight loss corrosion
CUI can result from a build-up of water and contaminants in the annular space between
the metal surface and the thermal insulation. It is compounded by situations such as hot wall
effects and alternate wetting and drying. The problem typical of CUI is that corrosion rates
are typically greater than predicted based on aqueous corrosion data produced from either
open or closed system measurements [19]. In an open system, corrosion rates are generally
low due to the decreasing solubility of oxygen with increasing temperature. The CUI
situation more closely represents a closed system; however, prior studies attempting to
simulate CUI by these methods have generally been unsuccessful.
Recently, experiments were conducted with a special test cell designed to model CUI
(see Fig. 4) [20]. This novel approach included the use of an internally heated metal tube and
isolated ring specimens surrounded by insulating material. The annular space was filled with
a simulated atmospheric condensate. Corrosion was assessed using ring specimens that could
be monitored using linear polarization resistance (LPR) techniques per ASTM G59, mass loss
per ASTM G1 and localized corrosion rate per ASTM G46 [21-23]. Tests incorporated
isothermal conditions, thermal cycling and alternate wet dry conditions.
Figure 5 shows the comparison of isothermal and cyclic tests. The mass-loss corrosion
rates show values comparable to those associated with CUI in field and plant operations. Of
particular interest is the variation in corrosion rate with time for the cyclic tests. The trend
indicates that periods of maximum corrosivity involve the periods during re-wetting of the
metal surface following the dry cycle. The peaks in corrosion rate are 2 to 3 times the steady
state corrosion rates. Furthermore, for cyclic wet-dry conditions, the steady state corrosion
rate also increases with time. The benefit of protective surface treatments which results in
much lower rates of corrosion versus time can also be seen.
Case Study No. 4: Need for Environment Replenishment
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Plenary Lectures
Two major concerns when simulating service conditions in the laboratory are the
changes in severity that may be caused by depletion of reactive constituents in the corrosive
environments and build-up of corrosion products or by-products. Both can modify the
corrosivity of the environment to produce a variation in the severity of corrosion from that in
the service application. Therefore, periodic monitoring of the laboratory environment may be
required to determine the rate of consumption or build-up of various species. Additionally,
replenishment may be necessary to eliminate the undesirable effects that they can produce.
Figure 4. Corrosion under insulation (CUI) cell designed by CLI International, Inc.
Figure 5. Instantaneous and mass-loss corrosion rates for a corrosion under insulation (CUI)
system
An example illustrating these situations is industrial applications having high partial
pressures of carbon dioxide (100-600 psia) in combination with low to moderate hydrogen
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Kane
sulfide partial pressures (0.01 to 1.0 psia). These conditions illustrate the importance of both
liquid and gas phase replenishment in these high CO2/low H2S systems. In such systems, the
steel corrosion product (Fe+2) is soluble in the aqueous phase at low to intermediate
temperature (< 60 oC). Additionally, this reaction is accompanied by an increase in the HCO3 concentration, which has a buffering effect resulting in increased pH. These factors can
result in the premature formation of a protective FeCO3 scale. Furthermore, corrosion will
also tend to consume the initial low-level supply of hydrogen sulfide. These three effects, if
not controlled, will generally result in an artificially low corrosion rate for steel when
compared to the service application. An autoclave procedure is needed for replenishment of
the gaseous and/or liquid phases so that the test duration can be prolonged and accurate
corrosion assessment can be achieved [24]. Additionally, in most cases, replenishment
procedures should be combined with the careful use of a large solution volume to specimen
surface area ratio to achieve optimum results.
In cases where a low hydrogen sulfide partial pressure is being utilized, special care must
be taken to maintain the intended amount of this reactive constituent. The difficulty in this
process increases directly with the corrosion rate of the materials being tested, the total
specimen surface area and decreasing hydrogen sulfide partial pressure. The case shown in
Fig. 6 is for replenishment involving tests of corrosion resistant alloys [25]. It can be seen
that following the first gassing, the desired hydrogen sulfide partial pressure was achieved,
but it decreased to a very low level after a short exposure period. At least two more
replenishments were required to achieve acceptably constant levels of hydrogen sulfide in the
test environment.
Figure 6. H2S partial pressure vs time
Case Study No. 5: Acceleration of Corrosion Processes
One of the greatest needs in laboratory corrosion testing is the ability to attain accurate
simulation of the test conditions while simultaneously achieving acceleration. The test
conditions must produce a reasonable mechanistic simulation yet achieve a degree of
acceleration which allows the laboratory test to predict future in-service events in a
reasonably short period of exposure time. This is perhaps the most difficult combination of
requirements. Oftentimes, tests that are accelerated produce artifacts in the data related to the
influence of corrosion mechanisms which are not present in the actual service. Such tests
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Plenary Lectures
must be approached and conducted with caution. Examples of beneficial acceleration
techniques are electrochemical techniques such as controlled polarization, mechanical
techniques such as slow strain rate and fracture mechanics testing, and elevated temperature
short duration tests on polymeric materials using Arrhenius. modeling techniques.
An example of electrochemical acceleration in a simulated environment was used to
produce corrosion films on wear test specimens in a short period of exposure (< 10 days) that
were comparable to those produced on actual components over a long period of service (1.5
years). The first step was to achieve accurate simulation of the service environment which, in
this case, was high purity cooling water for a boiling water reactor (BWR). Since the
intended fluid flow rate was low (2.2-2.8 ft/sec), a rotating cage setup located inside of the
autoclave reservoir was utilized [26] (see Fig. 7). The cage employed a special contactor
system to allow for application of a controlled anodic current while monitoring
electrochemical current. To minimize corrosion of the internal fixtures and application of the
anodic current to only the specimens, the fixture was constructed from pre-oxidized Zr-alloy
parts.
Figure 7. Rotating cage setup
An extensive literature/experience survey was conducted to determine the rate of steel
corrosion with time in the BWR environment and the chemical structure of oxide that would
be expected [27]. Based on this information, it was estimated that the corrosion film would
be composed of a-Fe2O3 near the surface and Fe3O4 near the metal/oxide interface and, after
18 months of exposure, it would be about 24,000 Å thick. Using the simulated environment,
a test was developed that utilized a slight anodic current to accelerate corrosion. Following a
series of qualification tests, corrosion films were produced which were a the mixed iron oxide
composition very close to the requirements. After an exposure period of ten days, the films
were found to be between 15,3000 and 30,600 Å in thickness. These test specimens were
subsequently utilized to conduct frictional wear tests so that the actuating force required to
manipulate control valves could be estimated.
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ACKNOWLEDGMENTS
I wish to thank the staff of CLI International, Inc. for their hard work and dedication and
also the CLI clients which have provided financial support for many technical investigations.
Both have been essential, and have contributed greatly to the development and use of the
techniques highlighted in this paper. I also give my appreciation to Ms Delia Cuellar, who has
worked and collaborated with me for many years. Thanks are also given to those key people
who have helped to provide a guiding light into the practical applications of laboratory
testing: Dr. J Brison Greer, Mr. Walter K. Boyd, Professor Joe Payer, Dr. Peter Rhodes and
Mr. Bill Ashbaugh.
REFERENCES
1. K. Lewis, Corrosion and Failure Prevention Using Appropriate Materials Selection
during Design: Better, Cheaper, Faster, Safer, Presentation at the Golden Gate
Materials Technology Conference, San Francisco, California, February 1-3, 1995.
2. J.H. Payer, Increased Reliability and Useful Life through Better Understanding of
Corrosion Processes, Plenary Lecture, Seventh Middle Eastern Corrosion Conference,
Bahrain Society of Engineering/NACE International, Manama, Bahrain, February 2628, 1996, pp. 50-53.
3. B.C. Syrett, Cost Effective Corrosion Control in Electric Power Plants, Plenary
Lecture, Seventh Middle Eastern Corrosion Conference, Bahrain Society of
Engineering/NACE International, Manama, Bahrain, February 26-28, 1996, pp. 1-18.
4. J.W. Spence, et.al., Planning and Design of Tests, Corrosion Tests and Standards, R.
Baboian, ed., MNL 20, ASTM, West Conshohocken, Pennsylvania, 1995, pp. 33-39.
5. ASTM G31, Standard practice for laboratory immersion corrosion testing of metals,
ASTM Annual Book of Standards, Section 3.02, West Conshohocken, Pennsylvania.
6. ASTM A262, Standard practices for detecting susceptibility to intergranular attack in
austenitic stainless steels, ASTM Annual Book of Standards, Section 3.02, West
Conshohocken, Pennsylvania.
7. ASTM G48, Standard test methods for pitting and crevice corrosion resistance of
stainless steels and related alloys by the use of ferric chloride solution, ASTM Annual
Book of Standards, Section 3.02, West Conshohocken, Pennsylvania.
8. J.W. Oldfield and B. Todd, Corrosion considerations in selecting metals for flash
chambers, Desalination 31, 1979, pp. 365-383.
9. A. Ikeda et.al., Corrosion Behavior of Low and High Alloy Tubular Products in
Completion Fluids for High Temperature Deep Wells, Paper No. 46, NACE
Corrosion/92, March 1992, NACE, Houston, Texas.
10. M. Fontana, Corrosion Engineering, 3rd edition, McGraw-Hill, Inc., New York, 1986.
11. R.D. Kane and S. Srinivasan, Socratestm: Selection of Corrosion Resistant Alloys
through Environmental Specification, CLI International, Inc., Houston, Texas, 1996.
12. C.A. Palacios and J.R. Shadley, CO2 Corrosion of API N-80 Steel at 71oC (160oF),
Paper No. 476, NACE Corrosion/91, March 1991, NACE, Houston, Texas.
13. X. Zhou and W.P. Jepson, Corrosion in Three Phase Oil/Water/Gas Slug Flow in
Horizontal Pipes, Paper No. 26, NACE Corrosion/94, March 1994, NACE, Houston,
Texas.
47
Plenary Lectures
14. Course Materials, MAPStm - Multiphase Autoclave Pipeline Simulator, Short Course
on Corrosion Test Methodologies for Inhibitor Evaluation, CLI International, Inc.,
Houston, Texas, 1995.
15. S. Srinivasan, Internal report on flow modeling for laboratory simulation of field
effects, CLI International, Inc., Houston, Texas, 1995.
16. W.P. Jepson, The Effect of Multiphase Flow on the Performance of Corrosion
Inhibitors in Oil and Gas Pipelines, Seventh Middle Eastern Corrosion Conference,
Bahrain Society of Engineering/NACE International, Manama, Bahrain, February 2628, 1996, p. 80.
17. K.D. Efird et al., Experimental Correlation of Steel Corrosion in Pipe Flow with Jet
Impingement and RCE Laboratory Tests, Paper No. 81, NACE Corrosion/93, March
1993, NACE, Houston, Texas.
18. K.D. Efird, Petroleum Testing, Corrosion Tests and Standards, R. Baboian, ed., MNL
20, ASTM, West Conshohocken, Pennsylvania, 1995, p. 354.
19. W.G. Ashbaugh, Corrosion under thermal insulation, Metals Handbook Volume 13:
Corrosion, 9th edition, ASM International, Materials Park, Ohio, 1987, p. 1145.
20. Private Communication, W.G. Ashbaugh, D. Abayarathna, R. Kane (CLI
International, Inc., Houston, Texas) N. McGowen (Elisha Technologies - Moberly,
MO), March 1996.
21. ASTM G59, Standard practice for conducting potentiodynamic polarization resistance
measurements, ASTM Annual Book of Standards, Section 3.02, West Conshohocken,
Pennsylvania.
22. ASTM G1, Standard practice for preparing, cleaning and evaluating corrosion test
specimens, ASTM Annual Book of Standards, Section 3.02, West Conshohocken,
Pennsylvania.
23. ASTM G46, Standard practice for examination and evaluation of pitting corrosion,
ASTM Annual Book of Standards, Section 3.02, West Conshohocken, Pennsylvania.
24. Course Materials, MARStm - Multiphase Autoclave Replenishment System: Short
Course on Corrosion Test Methodologies for Inhibitor Evaluation, CLI International,
Inc., Houston, Texas, 1995.
25. Private Communication, M.S. Cayard to R.D. Kane (CLI International, Inc., Houston,
Texas, March 1996.
26. R.D. Kane, High Temperature and High Pressure Testing, Corrosion Tests and
Standards, R. Baboian ed., MNL 20, ASTM, West Conshohocken, Pennsylvania,
1995, p. 108.
27. T. Honda, et al., Corrosion of ferrous materials and deposition of trace metal ions in
high purity water at high temperature, Corrosion Engineering 36, 1987, pp. 257-266.
48
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
CORROSION OF CONDENSERS IN MULTI STAGE FLASH
EVAPORATION DISTILLERS
A.M. Shams El Din
Material Testing Laboratory, Water and Electricity Department,
Abu Dhabi, UAE
ABSTRACT
All Arab Gulf States suffer from shortages in freshwater resources. The deficiency is made up for
by desalinating seawater, mainly by the multi stage flash (MSF) evaporation technique. Depending on
its design, an MSF distiller has between 16 and 39 cells, arranged in one or two decks. The basic
element of a cell is the condenser which converts the vapor into the distillate. The condenser is a
structure incorporating many components, each susceptible to a variety of corrosion forms.
The lecturer will describe the fine features allowing the characterization of the types of attack
affecting the water boxes, tube-plates, tube supports and the tubes themselves. Counter measures
necessary to remedy the defects and/or eliminate their causes will be mentioned.
Vapor side corrosion of condenser tubes made of copper-base alloys will be high lighted. Attack
causes the thinning of the tube walls and the contamination of the distillate by copper above
permissible limits.
A number of operational processes affect the corrosion of MSF condensers. The roles of
seawater chlorination, ferrous sulphate dosing, scaling and ball cleaning, acid cleaning and distiller
outage will be treated in some detail.
Key Words:
Water boxes, tubes, tube supports and plates, vapour side corrosion,
chlorination, ferrous sulphate dosing, scaling, ball cleaning, acid
cleaning, shut-down
INTRODUCTION
All Arabian Gulf States suffer from a scarcity of freshwater resources. The available
underground water is either nonrenewable, or gradually builds up salinity as a result of
exhaustive pumping. On the other hand, rainfall over the area is scant and unpredictable, and
extends over a fairly long period of time. Potable water deficiency is remedied by
desalinating seawater. Two techniques, viz. distillation and reverse osmosis, are commonly
employed. Distillation, which represents the largest share, is generally carried out by multi
stage flash (MSF) evaporation.
An MSF distiller is a mammoth metallic structure operating on seawater at temperatures
between ambient and 100-115°C. It incorporates a variety of components, each designed to
perform a definite task in the overall process of distillation. The number of stages of an MSF
unit is a design parameter and may vary between 16 and 39. A schematic presentation of an
49
Plenary Lectures
18 stage (cell) MSF distiller is given in Fig. 1. Fresh, chlorinated, settled seawater is passed
through the condensers of cells 18, 17 and 16 (heat rejection section) to raise its temperature
to 42°C. Part of this water, to be used as makeup water, is deaerated and mixed with return
brine in the flash side of cell 18. Between 15 and 30% of this mixed water is rejected as
blowdown, and the rest is pumped into the condenser tubes of cell 15, which represents the
start of the heat-gain section. As water passes through cells 14, 13, 12, etc., its temperature
rises gradually. Following cell 1, the water passes through the brine heater to bring its
temperature up to the top value for distillation. Coming out of the brine heater, the water
goes through the flash chambers of the various cells in the reverse direction of flow in the
condensers, i.e., cells 1, 2, 3, and so forth. The water distills under vacuum, and its
temperature progressively decreases.
The basic element of a cell in an MSF evaporation unit is the condenser which converts
the vapor into distillate. The condenser is a multicomponent structure comprising the water
boxes, the tube plates, the tube supports and the tubes themselves. A typical condenser is
represented diagrammatically in Fig. 2. During operation, all parts of the condenser suffer
corrosion in a variety of ways, depending on the material they are manufactured of. It is
estimated that between 80 and 85% of distillers’ outages are the result of defects occurring in
the condenser system, in particular the tubes. The shape of failure can in many instances be
the same, but its causes may differ widely. The choice of the necessary countermeasures
depends, therefore, on the successful deciphering of the fine features of attack.
The present paper describes the different forms of corrosion which affect the four basic
components of condensers in MSF distillers. The material included is largely based on
observations made by the author in the Umm Al Nar Power and Desalination Plant (Abu
Dhabi, UAE) during the past 14 years. The presentation is of an orientative nature and is
intended to serve a wide spectrum of people. For the normal operator, with little or no
experience in corrosion, it is an introduction to the subject. For the operation engineer it
evokes awareness and interest in material performance. For the corrosion specialist, it lists
alternative countermeasures. No attempt is made to deal with the theory of metal corrosion.
For the first two categories of persons this is not necessary, while for the specialist the
information is assumed to be well known. The reader must, however, be well aware of the
terms associated with the common forms of corrosion, viz. general- (uniform-), galvanic(bimetallic-), crevice-, pitting-, intergranular-, dealloying- and erosion-corrosion, as well as
corrosion cracking.
Further, it is beyond the scope of the present paper to enlist all materials suitable for
constructing part or all of the condenser system of MSF plants. Broadly speaking, these fall
into four classes: mild- and carbon steels, stainless steels, copper-base alloys and titanium
alloys. Details of the composition of these materials, their physical properties and their
corrosion characteristics are well documented [1-15].
50
Shams El Din
Figure 1. Schematic presentation of an 18-stage MSF distiller
51
Plenary Lectures
Figure 2. Schematic presentation of an MSF condenser.
CORROSION OF WATER BOXES
A variety of materials is used to manufacture water boxes. In some old, small units cast
iron was used. Cast iron undergoes severe general attack in flowing seawater, which is
intensified by contact to copper alloys used as tubes and tube plates. For this reason,
unprotected iron water boxes were designed with above-normal thicknesses to ensure longer
service life. Corrosion of iron protects the copper-base components of the distiller and raises
the resistance of copper tubes against erosion-corrosion [16-18]. Excessive attack on the iron
water boxes leads to graphitization and the loss of mechanical strength. Graphitization also
promotes the galvanic corrosion of copper alloys.
Means of preventing attack on iron water boxes involve their isolation from seawater
with appropriate paints or rubber-lining. This approach is ideal so long as the cover remains
intact. Once it has failed, though, excessive corrosion takes place as a result of the area ratio.
Parts of the failed linings might adhere to tube plate and/or block tube entrances, causing
crevice and pitting attack.
Corrosion of iron water boxes by seawater can be harnessed through cathodic protection.
Sacrificial zinc anodes are commonly employed. These must be monitored to ensure their
presence and active functioning.
Nowadays, MSF water boxes are manufactured of copper-base alloys; e.g., aluminium
bronze and gunmetal. Solid or claded cupro-nickels are more customary. There is no
difference between the corrosion resistance of these materials in the massive and the cladded
states, so long as seawater does not creep behind the clading. When this happens, severe
galvanic corrosion of the base plate takes place.
Aluminium bronze, and the cupro-nickels are relatively immune to general attack by
seawater. They suffer, however, erosion-corrosion when the water flow exceeds a certain,
material-dependent velocity. Attack is more severe if the water carries sand, shell fragments,
entrapped air or vertices. This type of attack is recognized by the marks developed on the
metal surface which depict the water flow pattern. Erosion-corrosion is reduced, or even
eliminated, by improving the intake filtration systems and/or by reducing the flow velocity.
Eroded areas are cleaned and filled with fiberglass or similar coatings, resistant to
impingement.
The water boxes undergo corrosion when parts of the surface are shielded by
nonconducting materials (e.g., plastic sheets and pieces of failed coatings) brought in with the
water. In the heat-rejection cells, where temperature still allows them to flourish, barnacles
and shells escaping the band screen can attach themselves to the surface of the water box.
There they not only promote crevice attack, but they also secrete harmful sulphur- and
nitrogen-containing chemicals.
Failure of water boxes might also result from material defects. These include alloy
inhomogeniety, defect welds, phase changes resulting from preferential cooling, foreign
inclusions, stresses and cracking. There is no general remedy for all types of faults. Frequent
inspection of the metal surface for microcracks and pits is recommended, especially following
commissioning. Attacked areas may be ground and filled with a neutral filling. Excessive
52
Shams El Din
corrosion will require removal of the defective area and careful repair welding. Most cases of
water box failure can be eliminated through cathodic protection using sacrificial iron anodes.
The released ferrous ions also improve the corrosion resistance of the copper condenser tubes
and tube plates.
CORROSION OF TUBE PLATES
Tube plates are commonly manufactured from stainless steel or copper-base alloys (i.e.,
brasses, bronzes, and cupro-nickels). These materials are used as such or are clad over Csteel. The choice of plate material is made on the basis of mechanical characteristics, cost
considerations and compatibility to condenser tubes.
Depending on their composition, and the operational and environmental conditions, tube
plates undergo different forms of corrosion:
Erosion-Corrosion
This results from the impact of seawater at velocities higher than design the values, from
impingement attack by entrained gas bubbles or from abrasion with sand loaded water. The
metal surface assumes a rough touch and acquires a shiny silver or golden luster due to the
loss of the natural protective film. Attack need not cover the whole plate since turbulence and
vortices are not equally distributed on the surface. Countermeasures against erosion-corrosion
involve the reduction of flow velocity, the removal of suspended material from the circulating
water and the release of entrapped gases. Severely attacked plate areas are ground and filled
with a neutral coating. Ferrous sulphate dosing improves surface film resistance to erosioncorrosion.
Crevice Corrosion
Tube plates may undergo crevice attack when covered by nonconducting material (e.g.,
polythene sheet, paint flakes etc.). In heat-rejection cells, barnacles and shell animals
continue to strive and stick to the surface of tube plates. During their life the organisms produce
organic nitrogen and sulphur compounds which intensify the attack on copper alloys. The
released copper ions eventually kill the animal, and crevice attack continues thereafter by way
of an oxygen concentration cell mechanism.
As a result of the large thickness of the tube plate, it is improbable that crevice attack
will continue undetected till it cuts through the metal. The water box is likely to be opened,
for one reason or the other, at which time preventive and corrective measures are undertaken.
These involve screening water properly at the intake, raising the chlorination level to
discourage the entry of marine organisms, and cleaning and filling present crevices with
neutral filling followed by a coat of hard epoxy resin.
Crevice/Galvanic Corrosion
Another type of tube-plate failure develops on claded plates with imperfect tube holes.
Expansion of the tubes leaves behind narrow spaces which are filled with seawater.
Corrosion starts as crevice corrosion until the water reaches the base metal under-covered
where the attack is accelerated as galvanic corrosion. Left for a long time, attack might eat
through the thickness of the plate. During its early stages, when the crevice corrosion is only
operative, corrosion might escape the attention of the inexperienced or hasty eye. Careful
53
Plenary Lectures
examination will reveal, however, tiny recessions between the tubes and plate. A steel needle
tip will measure the depth of attack. During the galvanic phases however, corrosion is
recognized by the brown stains of iron oxides oozing from between the tubes and the plate.
Counter-measures against crevice/galvanic attack involve the rerolling (re-expansion) of the
tubes. In most cases, this proves to be adequate for eliminating the clearance between the
tubes and the plate. If this fails, more stringent measures are undertaken. These involve the
thorough cleaning of the inside of the dried crevice with the help of steel needles until it is
free of corrosion products. An air compressor producing a micro-jet might prove useful.
This is followed by the intrusion of a natural fill and a cover with a hard epoxy resin. When
the claded tube plate and the tubes are of the same material, welding after cleaning presents
another practical alternative.
Corrosion by Dealloying
Tube plates made out of uninhibited brasses (i.e., copper-zinc alloys) undergo rapid and
severe dezincification by seawater, especially at high temperatures. Dealloying occurs in
patches in what is commonly referred to as plug-type attack.
Naval brass (i.e., brass with 0.5-1.0% tin) undergoes differential attack in which the
defected areas are covered with easily detachable scales of the corrosion products. This type
of attack is designated as exfoliation or layer-type corrosion.
Uninhibited brasses should not be used as tube plates in condensers operating on
seawater. Their use in small units is not due to good performance but to over-dimensions
which prolong their life span.
Dezincification is remedied by grinding the dealloyed areas and filling the depressions
with an appropriate neutral material. The whole plate surface, and preferably, the tube ends
as well, are coated with an approved epoxy resin. Cathodic protection with sacrificial mild
steel anodes will ensure stiffing of corrosion should the coat fail.
CORROSION OF CONDENSER TUBES
Failure of condenser tubes constitutes the largest cause of distiller outages, so the choice
of tube material is accordingly crucial. Tubes have to have reasonably high thermal
conductivity, sufficient ductility to expanded into the tube plate, and their corrosion
performance should be well understood. Last but not least is the question of cost. Three
types of materials present themselves. These are copper-base alloys, stainless steels, and
titanium. Each possesses its own merits and limitations, and will be treated separately.
Corrosion of Copper-Base Condenser Tubes [19-29]
Copper-base alloys suitable for manufacturing condenser tubes encompass the brasses,
the bronzes and the cupro-nickels. Simple brasses are 70/30 copper-zinc alloys, occasionally
with 0.5-1.2% tin or lead [12]. These must be inhibited by minute amounts of arsenic,
antimony or phosphorus to prevent dezincification. Aluminium brasses are a family of
copper-zinc alloys containing between 1.8 and 2.5% aluminium, with minor quantities of
iron, nickel, manganese and lead. Tubes of these brasses exhibit excellent resistance to brines
at high temperature. They do, however, suffer erosion-corrosion at flow rates above about 2
m⋅s-1, and are not immune to polluted (i.e., H2S and NH3) seawater. Their relatively low price
is, however, an appreciated asset.
54
Shams El Din
Bronzes, on the other hand, are copper-tin alloys. As a rule, simple bronzes are superior
in property to simple brasses. The reader must be aware of the confusion associated with the
nomination of bronzes [12]. Materials like aluminium bronzes (i.e., C 63700 and C 65200)
and manganese bronzes (i.e., C 67500 and C 67800) are in fact high-tensile brasses,
designated as bronzes just to praise their properties. True bronzes contain up to 13% tin. Gun
metals contain both tin and zinc; types A and LG 3 are resistant to erosion-corrosion. The
resistance is further improved through additions of 1.0-2.0% nickel as in Admiralty gun metal
G1. An important material for condenser tubes running on seawater is the so called AP
bronze, which contains 8% tin, 1% aluminium and 0.1% silicon. In contrast to aluminium
brass, AP bronze withstands corrosion by H2S-polluted seawater. Its life expectancy is
estimated to be 2-4 times that of aluminium brass [30,31]. In nonpolluted seawater,
aluminium brass is preferred to aluminium bronze on the basis of cost [32].
Copper and nickel are freely soluble in one another and form an extended solid solution.
A large number of alloy compositions were prepared and their resistance to corrosion by
seawater was evaluated [33]. Two alloys, viz. the 90/10 and the 70/30 were specially
attractive, particularly after the discovery of the beneficial effect of small additions (up to
2%) of iron and manganese. As tube materials, the cupro-nickels withstand flow velocities
up to 6.0 m⋅s-1 [34] and are fairly resistant to fouling in seawater. They are, however, easily
affected by polluted water. Nowadays condensers with tubes made of 90/10 and 70/30 cupronickels represent some 85% of all copper-base condensers. Copper-nickel alloys containing
0.5% chromium were recently reported [35,36]. These are described as exhibiting high
tolerance to erosion and impingement attack. Similar effects are achieved through additions
of solutionized iron [37].
Copper-base condenser tubes fail much in the same way; the only difference being the
extent and rate of attack. Failure is commonly manifested as a pore leading to the admixing
of the brine with the distillate. Many reasons are responsible, however, for pore production.
A clear understanding of pore morphology is, therefore, essential in establishing the cause of
corrosion and in ensuring the choice of appropriate measures to prevent the further escalation
of the problem.
From the water side, condenser tubes may fail by:
Crevice Corrosion
Crevice corrosion in condenser tubes results wherever parts thereof get covered by nonconducting materials such as seashells, barnacles, sand, clay, and damaged sponge balls.
Damage from barnacles attaching themselves to tube walls is limited to the cells of the heat
rejection section where environmental conditions (i.e., temperature, oxygen content, and
nutrients) still allow their growth. Barnacles can anchor anywhere along the tube’s length
and induce corrosion by way of depleting oxygen and secreting harmful chemicals. Due to
gravity, under-deposit attack develops at the bottom of the tubes. Very seldom, do they cross
the five and seven o’clock limits. Another feature characterizing crevice attack is that the
pore lies in the centre of a clear imprint of the cover that initiated the attack. Elongated pores
follow, in most cases, the flow direction although oval cuts with their major axes
perpendicular to the flow have been recorded. A linear succession of pores along the tube’s
length indicates a thick sand or clay deposit. Contrary to some claims [38], crevice attack
55
Plenary Lectures
does not develop under alkaline scales. The main constituents of the scales, CaCO3 and
Mg(OH)2, act as anodic inhibitors.
Crevice corrosion of condenser tubes can be stopped or greatly reduced by controlling
the entry of undesirable material. This requires the proper functioning of the band screen and
the appropriate chlorination of the feedwater to discourage marine life inside the lowtemperature cells. Dead shells and barnacles do not adhere to tube walls. Prolongation of the
time of stagnation in the settling basins helps prevent the accumulation of sand and silt inside
the tubes. The use of Tapproge ball cleaning, brush cleaning and back-flushing of the tubes
during shutdown, and the eventual acid-washing of the distiller remove tenaciously held
deposits and bodies.
Stress-Corrosion Cracking (SCC)
The expansion of tubes in the tube plate results in the accumulation of stresses along the
diameter-change zone. Although these stresses are within the range of the design values, they
induce, on the long run, the cracking of the tubes. Cracking takes the form of a
circumferential cut, perpendicular to the direction of flow. The cut is located only a few
centimeters from the tube edge and can readily to be ascertained with an inspecting finger.
Metallographic examination reveals that cracking is mainly intergranular [28] but
transgranular attack is not uncommon. The frequency of failure at tube inlets exceeds by far
that at outlets, suggesting that turbulance assists cracking. Stress corrosion cracking (SCC) of
aluminium-brass tubes has been recorded in the first cells of the heat-gain section of the
distiller [39]. In the high temperature cells, tubes of 70/30 cupro-nickel sometimes exhibit
failure features very similar to those of aluminium-brass. Corrosion of 70/30 cupro-nickel by
SCC has not been reported before, and the matter requires further investigation [40].
SCC requires special agents for initiation. Three materials for copper-base alloys are
recognized, viz., ammonia, hydrogen sulphide and mercury. The first two result from
decomposition of marine organisms and/or sewage discharged into the sea. Pollution by
mercury comes mainly from industrial activities in the vicinity of the desalination plant.
Chlorination of seawater at the intake removes ammonia and hydrogen sulphide. Sewage
treatment on land before discharge into the sea removes or reduces the amounts of the two
pollutants. Application of special paints to tubes’ ends isolates the areas susceptible to attack.
In Abu Dhabi the problem of SCC of aluminium-brass tubes was controlled through cathodic
protection of the water boxes, tube plates and tube terminals by sacrificial iron anodes [39].
Already perforated tubes are recovered through tube inserts. These are either metallic, in
which case they must be galvanically compatible to the tubes and tube plate, or of man-made
materials. In this last case they should not be made of nylon or amide-based polymers. These
substances readily hydrolyze in seawater to yield ammonia which promotes cracking.
Erosion-Corrosion
Erosion-corrosion of copper-base condenser tubes, occurs whenever the naturallydeveloped- or artificially induced film on their surface is destroyed [41-43]. This happens
when the water flow rate exceeds the upper limits of tolerance of the material. Erosion is
manifested in the form of grooves, gullies, waves, rounded holes or horse-shoes, lined up in
the direction of the flow. Erosion is more prominent at tube inlets due to turbulence. It
produces a rough surface which ends where the flow turns laminar. Erosion-corrosion can
proceed till complete failure occurs, producing pores elongated along the direction of flow.
56
Shams El Din
The rate and extent of attack increases when the water carries suspended sand or entrained
gases.
Counter-measures against flow-induced attack involve reduction of the water flow rate,
extension of the time of stagnation in sedimentation basins, and controlled use of ball
cleaning. Excessive ball cleaning, especially when of the hard type, removes the inhibiting
film from inside the tubes. Filming by ferrous sulphate or through sacrificial iron anodes will
restore the film. Too much chlorination of seawater at the intake is also detrimental to the
passivating film. Eroded tube inlets are repaired with appropriate tube inserts.
Corrosion by Polluted Waters
Polluted seawater contain ammonia and/or hydrogen sulphide. Both materials are
products of the decay of dead animals and organisms. The two pollutants attack copper
condenser tubes.
Hydrogen sulphide reacts spontaneously with copper tubes to produce a black, porous
copper sulphide film. Being nonprotective, the film allows further attack until the metal is
eaten through [44-49]. Attack is more rapid if the polluted water is loaded with slime and silt.
Under the slime, attack by sulphide is considered a special type of deposit corrosion in which
metal deterioration is initiated by the presence of hydrogen sulphide rather than by a
deficiency in oxygen.
Attack by sulphide-polluted water is identified by the black coloration of the tube inside.
Treatment of the black film with dilute acid sets free the hydrogen sulphide, known by its
characteristic bad odor. On the micro-scale range, the presence of sulphide is ascertained
through its catalyzing the evolution of nitrogen gas bubbles from a drop of iodine-sodium
azide solution. The test is carried out under a magnifying lens or microscope.
Sulphide corrosion can be dealt with in a variety of ways. Suitable tube material can be
chosen if it is known from the start that only polluted water is available. Extension of the
intake facilities away from the source of polluted water is a second alternative. When present
in small amounts, hydrogen sulphide is oxidized by chlorination. This must, however, be
carefully controlled. The same also applies to rubber ball cleaning if sulphide-polluted sludge
is in abundance.
Seawater polluted with ammonia affects the corrosion of copper base condenser tubes in
two distinct ways: it induces SCC and the rate of crack propagation increases with pollutant
content; and ammonia and ammonium salts enhance the general attack by dissolving and
complexing with the copper ions. The tubes lose their protective film and acquire a shiny
appearance. Eventually the tubes fail when they become perforated at weak points. Failure
can occur anywhere along the tube length and the resulting pore takes any form. Ammonia
attack is suspected when daily analysis of the brine shows a constant, high level of copper.
There is no direct solution to the problem of corrosion by ammonia-polluted seawater. If
the source of pollution is permanent, relocation of the seawater intake might prove practical.
If pollution is widespread, a change to ammonia-resistant tubes, e.g., titanium or stainless
steels, might be considered. Chlorination of ammoniated seawater is effective in removing
low levels of pollution. The two agents react to produce chloramines which are weaker
disinfectants than chlorine itself. The formation of chloramines reduces, but does not
completely eliminate the problem of ammonia-corrosion since chloramines readily hydrolyze
57
Plenary Lectures
to the original harmful material. Ferrous sulphate dosing retards the aggressive action of
small amounts of ammonia by forming an inhibiting ferric oxide film.
Corrosion of Stainless Steel Condenser Tubes
There is a wide variety of stainless steels that can be used as condenser tubes in MSF
distillers [7,50-52]. In flowing seawater, the performance of stainless steel tubes is quite
satisfactory. They withstand general corrosion and are not affected by the flow velocities
operating in a distiller. Normal steels are susceptible to SCC in solutions of high chloride
content at high temperatures. As the tubes are not subject to any noticeable stresses, this type
of attack is unlikely to develop. In stagnant solutions, on the other hand, normal stainless
steels are liable to undergo pitting corrosion [53-56]. During a long outage, the distiller must
be carefully drained and flushed with ample quantities of potable water to prevent chloride
attack. Stainless steel tubes are also prone to crevice corrosion, which develops under
barnacles, stuck rubber balls and the like. The same procedures described in case of copperbase tubes apply in such cases.
Weld areas and the heat affected zones are weak points where intergranular and pitting
attack readily propagate on stainless steel in a desalination unit [57,58]. Spot and arc welding
is to be applied, and annealing should be carried out whenever feasible. Finally, stainless
steel tubes mounted on copper-base tube plates will lead to their rapid deterioration by
galvanic action [59-62]. Cathodic protection using sacrificial mild steel anodes overcomes
this problem.
Corrosion of Titanium Condenser Tubes
Titanium tubes exhibit excellent resistance to general-, pitting- and crevice corrosion.
They withstand erosion-corrosion and impingement attack at flow velocities as high as 20
m⋅s-1. They are also immune to SCC and are unaffected by polluted seawater. Titanium has,
however, a low thermal conductivity, and this is resolved by reducing the tube wall thickness.
To compensate for the loss in rigidity, a larger number of tube supports are installed.
Titanium tube plates are expensive, and the expansion of titanium tubes, therein, can lead to
crevice corrosion [65,66] at temperatures at and above 80°C. Organic sealants (e.g.,
dimethacrylate) used to fill the tube/ tube plate interface worsen rather than improve the
situation, as crevice corrosion starts at about 80°C [66]. Two alternatives can be applied to
prevent titanium/titanium crevicing. The first involves seal welding and this has the
advantage of preventing crevice attack completely. It has, however, the drawbacks of being
expensive and requiring skilled personnel. The second procedure is the coating of the
titanium tubes/tube plates with a palladium oxide (PdO)/titanium oxide (TiO2) mixture. The
coating is prepared in situ from a mix solution of palladium chloride (PdCl2) and titanium
chloride (TiCl3) followed by thermal oxidation in air [67]. This type of sealant prevents
crevice attack in 15% NaCl of pH 3 at 120°C [67].
The mounting of titanium tubes on copper-base tube plates prevents the crevice
corrosion of the tubes. However, because of the large difference in their free corrosion
potentials, titanium induces the galvanic corrosion of the tube plate [62,66,68]. To stop
attack, cathodic protection of the plate and tube ends by impressed currents [69,70] or
sacrificial anodes [71,72] has been suggested. Unless very carefully controlled, protection
can do more harm than good. Protection should not reduce the potential of titanium to lower
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Shams El Din
than -0.50 V versus the saturated calomel electrode (SCE) [66] or -0.65 V versus SCE [71].
Polarization to more negative potentials will cause the absorption of hydrogen in titanium and
the formation of titanium hydride. This causes a loss of ductility and cracking of the tubes.
Normal mild steel is not recommended as anode material as it shifts the potential of titanium
to values more negative than those recommended for safe protection. An iron/9% nickel
alloy is recommended to keep the potential at the desired value [66]. A recent study from this
Laboratory [73] revealed that hydrogen uptake by titanium can take more than one form,
depending on the type of anode used.
Titanium tubes are not recommended for use in the high temperature cells and in the
brine heater, where they are likely to pick up hydrogen from the hot brine. In their place,
tubes of 70/30 cupro-nickels can be installed.
CORROSION OF TUBE SUPPORTS
Tube supports are elements of MSF distillers not readily accessible for inspection and
repair. The functions of tube supports are to keep the tube bundle in position, to act as guide
for long tubes and to prevent the tubes from sagging and vibrating. Because tube supports are
exposed to deaerated steam, they were assumed to suffer no corrosion. Accordingly, they
have been made of the cheapest material, viz. mild steel. In fact, however, mild steel tube
supports are heavily corroded by reaction with steam as well as by galvanic contact with the
nobler material of the tubes. By the time the supports change into iron oxide lumps which
make it extremely difficult, if not completely impossible, to extract failed tubes or retube the
condenser. There is nothing that can be done to save C-steel supports in contact with
stainless steel, copper alloys or titanium tubes, from completely failing. Nowadays tube
supports are specified to be of the same material as the tubes themselves.
VAPOR-SIDE CORROSION
Until now, we have considered the forms of corrosion affecting components of MSF
condensers in contact with seawater. Cases have also been recorded, where corrosion occurs
on the vapor side. One such case has already been mentioned, namely, the corrosion of Csteel tube supports bearing condenser tubes.
Occasionally reference is made to the vapor side corrosion of copper-base condenser
tubes [74,75]. Such attack is recognized through shining tube surfaces, an outside reduction
in wall thickness and an above-normal (>0.1 ppm) increase in the Cu2+ content of the
distillate. This occurs mainly in the high temperature cells of the distiller. The general belief
[76] is that the vapor is largely contaminated with CO2 resulting from the high temperature
decomposition of the HCO -3 ion of the brine [95]. The formed carbonic acid, H2CO3, is
assumed to be sufficiently strong at high temperatures to attack the copper tubes. Apparently
this explanation is an over simplification of a more complex process. At the temperatures of
the first heat gain cells, the solubility of CO2 in the vapor is both limited and transient.
However, neither copper nor its alloys displace hydrogen from acid solutions. If H2CO3 is to
exert a corroding action it must do so by dissolving the cuprous and cupric oxides formed on
the metal surface. Practical experience confirms this conclusion. Condenser tubes
undergoing overhaul acquire black coloration as a result of air oxidation. When returning to
service, the distillate exhibits above-normal copper-content and has to be dumped back into
59
Plenary Lectures
the sea. Considerable time (up to 60 hours [75]) elapses before the copper concentration in
the distillate returns to permissable limits. For copper tubes to continue corroding, their
surface has to reoxidize. This occurs through the oxygen (air) dissolved in the flashing brine.
As this is limited, the rate of corrosion falls to low values. Higher corrosion rates point to
the presence of tiny leaks in cell gaskets. A vacuum leak test (or pressure leak test) on the
distiller will confirm this assumption.
Vapor-side corrosion of copper condenser tubes can be serious if seawater is polluted by
ammonia or hydrogen sulphide. What has been said regarding water-side corrosion applies to
vapor-side corrosion too.
OPERATION PROCESSES AFFECTING CORROSION IN MULTI STAGE FLASH
(MSF) DISTILLERS
In running an MSF plant, certain processes are carried out the aim of which is to
overcome an operational difficulty and/or to improve production efficiency. These processes
affect the corrosion of distiller components in variety of ways. In some instances, the role of
a certain parameter can be beneficial at one level and detrimental at another. The successful
operation of the distiller depends, therefore, on a proper understanding and careful execution
of the process. The following operations deserve special mention:
Chlorination of Seawater
Chlorination of seawater at the intake is carried out to discourage marine organisms from
entering the distiller and to prevent bifouling. There is no single chlorination procedure that
applies to all units. The adoption of a certain course of action evolves from trial and error and
depends on factors such as geographic location, bioactivity, temperature, and water purity.
For example, in Umm Al Nar (Abu Dhabi) chlorination is carried out continuously to a level
of 0.4 ppm. This ensures a residual chlorine content of about 0.25 ppm in the condenser
tubes. To prevent organisms from building up resistance, shock dosing at 1 ppm is achieved
for 1 hour. once a week [78]. Chlorine reacts with water to produce hydrochloric and
hypochlorous acids : Cl2 + H2O = HCl + HClO. Both acids are directly neutralized by
seawater alkalinity. The hypochlorous ion, ClO-, is a strong oxidizing agent which raises the
free corrosion potentials of alloys to higher values. This raises the susceptibility of normal
stainless steels to pittings- and crevice corrosion, particularly in regions of stagnation [79].
Similarly, chlorination of seawater interferes with the process of ferrous sulphate dosing.
Chlorination is stopped before, during and for some time after dosing to allow the iron-oxide
film the chance to build up properly. On the other hand, chlorine injection destroys pollution
by hydrogen sulphide, and neutralizes, to some extent, ammonia polluted seawater.
Ferrous Sulphate Dosing
Copper-base condenser tubes have little resistance to erosion-corrosion and to corrosion
by hydrogen sulphide-polluted seawater. Both types of attack are greatly inhibited by ferrous
sulphate dosing [80]. The ferrous ion readily hydrolyzes and oxidizes to colloidal FeOOH,
which adheres strongly to tube surface. The iron oxide film improves the naturally formed
copper oxide and offers better protection. Two to three hours before sulphate dosing,
chlorination of the seawater is stopped. The sulphate is injected at a rate of 3 ppm for one
hour per day for one month at tube inlets. Chlorination can be resumed after two hours.
During the second and third months, ferrous sulphate dosing is reduced to 2 and 1 ppm for
60
Shams El Din
one hour per day, successively. During the whole period of treatment, sponge ball treatment is
reduced to one or two times per week for one hour. The success of sulphate dosing is
recognized through the uniform coloration of the tube plates and tubes with a light- to deepbrown film, which does not peel off when scratched with a nail. A number of factors
interfere with film formation. These involve waters flowing at high rates or carrying large
amounts of sand, slime or deposits; excessive or continuous ball cleaning; chlorination
directly before, during or after sulphate dosing and/or the presence of entrapped, polluted
dead waters.
Scaling in Condenser Tubes
Scale formation in the condenser tubes is one of the major problems encountered during
the operation of MSF distillers. Scaling impairs heat transfer, causes tube blockage and can
induce corrosion. Scale formation is an inherent product of the composition of seawater,
which contains HCO -3 , Ca2+ and Mg2+ ions. The thermal decomposition of the HCO -3 leads
to the deposition of calcium carbonate and magnesium hydroxide. Two alternative
techniques are employed to prevent (retard) scale formation. The first involves the controlled
acidification of the makeup water to convert the HCO -3 ion into CO2. Either sulphuric or
hydrochloric acids is used; the first is preferred on the basis of cost considerations. An
addition of 100-200 ppm sulphuric acid is normally used, and the evolved CO2 is stripped off
either in a separate degassing tower or in the MSF vacuum deaerator. Unless very carefully
carried out, acidification can lead to serious corrosion of the condenser tubes and the
deaerator system. To overcome this problem under-acidification has been proposed [81].
Laboratory experiments have shown that 80% of the acid required to completely neutralize
the bicarbonate ions is sufficient to prevent scale formation for a long time [82]. Another
technique to prevent corrosion involves the neutralization of excess acid. Following
degassing, the brine pH is brought back to ~7.5 through the controlled addition of caustic
soda. This technique is elaborate and expensive [83].
The second approach for scale prevention is the use of antiscaling agents. These are
polymeric, surface-active substances which adsorb on active centers of CaCO3 and Mg(OH)2
crystallites. This inhibits them from forming a continuous layer. A small amount of the
inhibitor (2-5 ppm) is needed to retard scale formation, and the method is known as a
threshold treatment [84]. A large variety of antiscalants is available on the market, and the
nature of the scale depends on the compound used. Polyphosphates, for example, give rise to
dense, fluffy, grayish-brown deposits [85]. Maleic anhydride polymers, on the other hand,
yield thin, hard scales [86].
A hybrid technique for scale control has been suggested [87]. It involves the use of lessthan-stiochiometric quantities of a mineral acid together with small quantities of a threshold
antiscalant. Under-acidification eliminates the problem of acid corrosion of tubes, while the
removal of the largest part of the water’s alkalinity allows lower concentrations of the
antiscalant to be used. This represents a sizable reduction in cost.
Sponge Ball Cleaning
The retardation of scale formation in the condenser tubes through dosing of antiscaling
agents is usually coupled with the technique of sponge ball cycling (Tapproge ball cleaning).
The rubber balls are slightly larger than the interior diameter of the tubes and are forced
61
Plenary Lectures
through by the pressure of circulating water. As they travel through the tubes, the balls wipe
out the tube’s insides and prevent scale crystallites from building a solid layer. As the balls
pass randomly through the tubes, they are cycled for enough time to ensure the cleaning of
the maximum number of tubes. There is no standard procedure for ball cycling. This varies
between one or more cycles per day and continuous cycling. The duration of a cycle also
differs from one plant to another. On the market are sponge balls with different surface
hardness. The choice of the appropriate type depends on the seawater purity, type and
concentration of the additive used, nature of the deposit and distiller top temperature. A trial
to optimize sponge ball cleaning was recently published [88].
Sponge balls can affect tube corrosion in many ways. Soft balls and/or short treatment
times might be ineffective and allow sand, silt and scale deposition which lead to crevice
attack. The same result might be noted when the number of balls is insufficient. On the other
hand, balls that are too hard can induce erosion-corrosion by stripping the protective film
from the metal surface. Also, it is not uncommon for balls to get stuck inside the tubes,
promoting crevice attack. Finally, as mentioned above, ball cleaning damages the fresh iron
hydroxide film resulting from ferrous sulphate dosing.
Acid Cleaning
Antiscaling agents do not inhibit completely the formation of alkaline scales; they only
retard their growth. Even when their use is coupled with sponge ball cleaning, a scale film
continues to grow inside the condenser tubes. Due to their bad thermal conductivity, a
situation is eventually reached when the gained output ratio of the distiller drops below a
preset value. The operation of the distiller becomes impractical (noneconomical) and an acidwash of the distiller is necessary.
The washing process is carried out by circulating warm (~65°C) fresh seawater through
the water boxes and condenser tubes of the cells and the brine heater. Enough inhibited acid
is added to the water to bring its pH value to 1.8-2.0. For copper-base condensers either
hydrochloric or sulphuric acids can be used. The pH of the water is monitored, and its value
rises as result of reaction with the scale. Extra acid is added to bring the pH to its lower value
and the process is repeated until no further increase in pH is recorded for a long time (usually
two hours). The acid water with the accumulated sludge is discharged, and the distiller is
flushed clean with fresh seawater. The use of an improper corrosion inhibitor [89] or
insufficient quantities of a suitable one leads to general attack on the tube material.
The same washing procedure described above applies to condensers with titanium tubes
as well. However, because of the high tendency of titanium to absorb hydrogen, weak
organic acids, e.g., citric [90] or sulfamic [91] acids are used instead. A few organic
compounds marketed under trade names, e.g., Galvane® (ICI) and IBIT [91], are said to retard
acid attack on titanium.
Distiller Outage
As is clear by now, seawater is an extremely aggressive medium which attacks all of the
metallic components of an MSF unit. The fact that distillers operate for long times with little
damage is due to two main reasons:
62
Shams El Din
• the exploitation of certain material properties manifested during operation;
• the application of protective and preventive measures during operation.
To the first reason belongs:
• The ability of stainless steels to withstand the corrosive action of flowing seawater,
while being quickly and severely attacked by the same water under stagnation, and
• The readiness by which copper-base alloys develop protective surface films which
offer some resistance to the action of flowing water.
Preventive and protective measures that can taken during operation are mentioned in
detail in previous sections. These involve degassing to remove oxygen (depolarizer) and
carbon dioxide (acidity) from the brine, ferrous sulphate dosing to increase the protection of
copper tubes, chlorination to discourage marine organisms from gaining access to the inside
of the distiller, use of antiscalants and ball cleaning to minimize scale formation and the
application of cathodic protection to prevent tube failure. Accordingly, as long as these
measures are observed and the distiller is operating, the corrosion of its components is largely
under control.
This well balanced system of protection is lost once the distiller is shut down. Upon
opening the water boxes and the flash chambers, air (oxygen) fills the entire unit, and
accelerates ongoing corrosion. On the other hand, brine stagnation inside condenser tubes
causes the settling of sand and silt, and occasionally a few sponge balls. This initiates crevice
corrosion. Stagnation also promotes the pitting corrosion of stainless steel components.
Unless the shutdown is for a short time (i.e., a maximum of two days) precautions against
attack should be applied. These involve the draining of the stagnant brine from the distiller,
followed by a thorough flushing with potable or distilled water. The washings should likewise
be drained out. During long outages, all inlets and outlets should be left open to remove
humidity and speed up the drying of the distiller’s inside. The various components of the unit
should be examined and the necessary corrective measures taken.
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Plenary Lectures
9. B. Todd, Proceedings 25th Annual Conference of Metallurgists, Toronto, 1986.
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39. A.M. Shams El Din, Wsia J. 11, 2, 1984, p. 1.
40. A.M. Shams El Din, WED, Abu Dhabi, unpublished results.
41. W.E. Heaton, Br. Corros. J. 13, 1978, p. 57.
42. T. Sydberger and U. Lotz, J. Electrochem. Soc. 129, 1982, p. 276.
43. J.M. Popplewell and E.A. Thiele, Corrosion 80, Houston, 1980, Paper No. 30.
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45. J.N. Al-hajji and M.R. Reda, Corrosion science 34, 1993, p. 163.
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1263.
64
Shams El Din
47. K. Habib, Desalination 89, 1992, p. 41.
48. B.C. Syrett, Corrosion 80, Houston, 1980, Paper No. 33.
49. G.A. Gehring, R.L. Foster and B.C. Syrett, Corrosion 83, Anaheim, 1983, Paper No.
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50. Trent Tube, Stainless Steel Pipe and Tube Alloy Handbook, Product Promotion
Brochure, 1992.
51. K.R. Fröhner, Desalination 21, 1977, p. 147.
52. A.J. Sedriks, Corrosion 45, NACE, 1989, p. 510.
53. T. Hodgkiess and A. Asimakopoulos, Desalination 38, 1981, p. 247.
54. T. Hodgkiess, W.T. Hanbury and M.N. Hejazian, Desalination 44, 1983, p. 223.
55. H.P. Hack, Materials performance 22, 1983, p. 24.
56. T.S. Lee, R.M. Kain and J.W. Oldfield, Materials performance 23, 1984, p. 9.
57. T. Rogne and J.M. Drugli, Corrosion 86, NACE, 1986, Paper No. 230.
58. H. Al Zahrani, S. Somuah and N.M.A. Eid, 12th Inter. Symp. Desalination and Water
Reuse, Malta, 1991, vol. 4, p. 349.
59. B. Wallen and T. Andersson, ACOM 2, 1987, p. 1.
60. G.A. Gehring, C.K. Kuester and J.R. Maurer, Corrosion 80, NACE, 1980, Paper No.
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61. G.A. Gehring and J.R. Maurer, Corrosion 81, NACE, 1981.
62. G.A. Gehring and R.J. Kyle, Corrosion 82, NACE, 1982, Paper No. 60.
63. J.A.S. Green, B.W. Gamson and W.F. Westerbaan, Desalination 22, 1977, p. 359.
64. A.R. Morris, Desalination 31, 1979, p. 387.
65. S. Kido and T. Shinohara, Desalination 22, 1977, p. 369.
66. T. Fukuzuka, K. Shimogori, H. Satoh and F. Kamikubo, Desalination 31, 1979, p.
389.
67. K. Shimogori, H. Satoh, F. Kamikubo and T. Fukuzuka, Desalination 22, 1977, p.
403.
68. P.D. Simon, Corrosion 83, NACE, 1983, p. 60.
69. J.P. Fulford, R.W. Schutz and R.C. Lisenbey, Joint ASME/IEEE Power Generation
Conference, Miami Beach, Florida, 1987, Paper No. 78 - JPGC-Pwr-F.
70. J.I. Lee, P. Chung and C.H. Tsai, Corrosion 86, NACE, 1986, p. 259.
71. T. Moroishi and H. Miyuki, Titanium 80, vol. 4, 4th Intern. Conf. on Titanium, Kyoto,
1980, p. 2713.
72. K. Kohsaka, K. Kitaoka, Y. Masuyama, M. Oshiyama, M. Yamamoto and K. Kashida,
Seawater Desalination Group, Development Committee of Japan Titanium Group,
Product Promotion Brochure, May 1984, and April, 1986.
73. A.M. Shams El Din, T.M.H. Saber and A.M. Taj El Din, Paper presented before the
IDA Congress on Desalination and Water Sciences, Abu Dhabi, 1995.
74. R. Heaton and T.A. Douglas, Desalination 41, 1982, p. 71.
75. E.A. Al-Sum, Sh. A. Aziz, A. Al-Radif, M.S. Said and O. Heikal, Proc. IDA and
WRPC World Conf. Desal. and Water Reuse, Yokohama, 1993, vol. I, p. 501.
76. A.H. Khan, Desalination Processes and Multi-Stage Flash Distillation Practice,
Elsevier, 1986, p. 441.
77. A.M. Shams El Din nd R.A. Mohammed, Desalination 99, 1994, p. 73.
78. A.M. Shams El Din, B. Makkawi and Sh.A. Aziz, Desalination 97, 1994, p. 373.
79. B. Wallen, ACOM 4, 1989, 1990.
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Plenary Lectures
80. T.W. Bostwick, Corrosion 17, 1961, p. 12.
81. Office of Saline Water, Res. Develop. Report, 1968, Vol. 559.
82. J.W. McCutchan, UCLA, Dept. Eng. Rept., 1967, No. 67-1.
83. M.N. Elliot, Desalination 6, 1969, p. 87.
84. K.S. Spiegler and A.D.K Laird, Principles of Desalination, 2nd ed., Part B., New
York: Academic Press, 1980, p. 672.
85. A.M. Shams El Din, Desalination 61, 1987, p. 89.
86. A.M. Shams El Din, Desalination 69, 1988, p. 147.
87. F. Butt, F. Rahman, A. Al-Abdallah, H. Al-Zahrani, A. Maadhah and M. Amin,
Desalination 54, 1985, p. 307.
88. F. Al-Bakeri, F. and H. El Hares, Desalination 94, 1993, p. 133.
89. T.M.H. Saber, A.M. Tag El Din and A.M. Shams El Din, Br. Corros. J. 27, 1992, p.
139.
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91. Japan Titanium Society, Product Promotion Brochure, September 1991.
66
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
CORRECT MATERIALS SELECTION FOR DESALINATION:
THE KEY TO PLANT RELIABILITY
J.W. Oldfield
Cortest International
23 Shepherd Street
Sheffield, S3 7BA, UK
ABSTRACT
Large scale desalination really began with the multi-stage flash (MSF) plant built by Weir
Westgarth in Kuwait in 1960. This and similar plants elsewhere used mainly carbon steel and brass
alloys in their construction. Corrosion problems with these early plants led to costly shutdowns and
maintenance. As knowledge of the requirements of materials has grown there has been a steady
upgrading of materials and this, together with improved control of operation has resulted in much
better reliability and performance.
This paper presents a state of the art review of factors which are important in the selection and
use of materials in MSF and reverse osmosis (RO) desalination plants. Economics play an important
part in materials selection and are given consideration in general terms.
Key Words: Desalination, MSF, RO, corrosion, materials.
INTRODUCTION
Large scale desalination really began with the multi-stage flash (MSF) plant built by
Weir Westgarth in Kuwait in 1960. This and similar plants built elsewhere used mainly
carbon steel and brass alloys for their construction. Experience in these early plants,
particularly when acid-dosing was used as an antiscalant, was that corrosion problems were
often extremely severe leading to costly shutdowns and maintenance. As knowledge of the
requirements of materials to meet the conditions in MSF plants has grown, there has been a
steady upgrading of materials, and this, together with improved control of operation, has
resulted in much better reliability and performance.
This paper presents a state of the art review of the factors which are important for
selection and use of materials in MSF and reverse osmosis (RO) desalination processes. As
the risk of corrosion is always present in desalination, corrosion plays a major part in
determining material selection. Corrosion is a process involving the reaction of a material
with its environment, and this paper reviews materials for usiing incomponents in the
environments encountered in MSF and RO processes. In MSF plants these environments are
sea water, deaerated sea water and brine, distillate and incondensible gases; in RO plants they
are sea water, brackish waters and brines. Economics play an important part in materials
selection and are given consideration in general terms. For general reference an appendix is
included giving typical compositions of materials used in desalination plants.
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Plenary Lectures
BEHAVIOUR OF MATERIALS IN SEA WATER AND BRINE ENVIRONMENTS
General Considerations
Sea water is the most corrosive of the natural environments that materials have to
withstand, but it is much less corrosive than many environments encountered in industry,
such as mineral acids. This situation has important implications for desalination materials in
that there is a wide range of corrosion resistant alloys readily available, which have been
developed for the chemical and process industries, and which are resistant to sea water.
However, many of these materials are much more expensive than those used in industries
which have traditionally handled sea water, such as shipping and power plant. In these
industries, carbon steel, cast iron, copper base alloys and standard grades of stainless steel
have been the usual choice for sea water applications. The desalination industry followed
traditional practice and whilst in many cases this proved satisfactory, in others it did not.
Upgrading, therefore, largely involved economic decisions-as the better materials were
always available, the main problem was to decide what level of cost and performance was
acceptable.
Corrosion of Carbon Steel in Sea Water
Sea water is a complex environment consisting of a mixture of inorganic salts, dissolved
gases and organic compounds [1]. However, it also supports living matter in the form of both
macro-organism (e.g., fish, shellfish, seaweed etc) and micro-organisms. All of these can
have an influence on corrosion processes. Sea water is a well buffered solution, so the pH
remains fairly constant at about 8. This means that most corrosion processes are dependent on
the presence of oxygen. For carbon steels immersed in sea water, the rate of corrosion is
mainly dependent on the oxygen content and the temperature of the sea water, the
composition of the steel has relatively minor influence even when small amounts of alloying
elements are added. Table 1 gives data on several steels exposed in the Pacific Ocean near
Panama.
Table 1. Corrosion of Steel in Sea Water [2]
Steel
Nominal
Composition
Carbon
Copper-Bearing
0.2 5% C
0.22% C
0.31% Cu
2% Ni 0.6% Cu
Corrosion Rate
(mm/yr)
1 year
16 years
0.15
0.075
0.15
0.077
Low Alloy Cu0.15
0.076
Ni
Chromium Steel 0.08% C 3% Cr
0.05
0.110
Nickel Steel
5% Ni
0.16
0.080
Grey Cast Iron
0.24
0.147
Samples immersed below minimum low tide. Tidal flow 0.3
m/sec. Water temperature 15.5-32.2°C. Surfaces pickled before
exposure.
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Oldfield
In most engineering applications sea water flows over the metal surfaces so, it is
important to have corrosion data under these conditions. Figure 1 [3] shows how flow rapidly
increases corrosion of carbon steel in sea water. This is mainly due to the increase in the
mass flow of oxygen to the corroding surface. In applications where continuous flow is
required, a corrosion rate of about 1 mm/yr can be assumed for carbon steel and cast iron.
Figure 1. Effect of velocity on corrosion of carbon steel in sea water
Figure 2. Estimates of corrosion of carbon steel in deaerated sea water
Deaeration, by reducing the oxygen content, would be expected to reduce corrosion, and
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Plenary Lectures
in general, this is found to be the case (Fig. 2). However, in MSF plants, where deaeration
occurs as part of the process, two other factors have an important effect. These are
temperature and flow rate. Few test data have been measured under controlled oxygen, flow
and temperature conditions, but Oldfield and Todd [4] have modelled the process and
compared calculated and actual test data.
Comparison with measured data [5] show that even with good deaeration, high corrosion
rates can be experienced at high flow rates and high temperatures. At low corrosion rates, the
measured and calculated corrosion rates are in good agreement, but at higher rates the
calculated values are significantly above the measured values. This was attributed to the
buildup of rust scales on the steel. Although these are not protective, they do provide a
barrier to oxygen access to the surface and reduce the rate of attack.
Corrosion of Copper Base Alloys in Sea Water
Copper alloys have been used traditionally in marine engineering for heat exchanger
tubing and for cast and wrought components in pumps and valves. These alloys form good
protective films in sea water, and provided these films are undamaged, corrosion is slight.
However, the protective films are susceptible to damage by fast flowing sea water, and this is
an important factor in selecting these alloys for applications involving flow-for example, heat
exchanger tubing. Table 2 gives data on some commonly used heat exchanger alloys.
Table 2. Jet Impingement Tests on Copper Base Alloys
Alloy
Admiralty Brass
Aluminium Brass
90/10 Cupronickel
7 0/30 Cupronickel
66/30/2/2 CuNiFeMn
Depth of Attack (mm).
(jet velocity = 5.5 m/s)
0.60
0.13
0.06
0.03
0.025
(jet velocity: 12 m/sec)
In deaerated sea water copper alloys can still form protective films and show higher
resistance to corrosion than in natural water. This is because the potential of copper alloys in
sea water is much higher than the potential at which hydrogen can evolve; thus, in the
absence of oxygen, corrosion is negligible. However, as in the case of carbon steel, low
oxygen levels with high flow rates can cause impingement attack as shown by Anderson [6],
but the rate of attack was much lower than in aerated sea water. Further data on corrosion of
heat exchanger tubing alloys in deaerated sea water is given in Fig. 3 [7]. These data indicate
that at low temperatures and oxygen levels, corrosion of the three alloys tested is acceptably
low. However, at high temperatures with high oxygen levels, although the cupronickels
continue to show low corrosion rates, aluminium brass is attacked. Thus, in MSF plants,
where aluminium brass tubing is used, it is confined to the lower temperature recovery stages
with the cupronickels being used in the higher temperatures areas.
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Oldfield
Figure 3. Effect of oxygen content on corrosion of Cu base alloys in deaerated sea water
Corrosion of Stainless Steels in Sea Water
When chromium is added to steel, there is a marked increase in corrosion resistance, and
at about 12% chromium, the alloys form a protective passive film and are referred to as
stainless steels. Although the iron-chromium ferritic stainless steels are used commercially,
the greatest tonnage of stainless steels are the austenitic iron-nickel-chromium alloys. These
have better ductility and weldability than the ferritic alloys, and as they usually contain about
18% chromium, they have better corrosion resistance.
In sea water, general attack on these alloys is negligible; however, they are prone to
localised attack due to their high chloride content. This attack is particularly severe in
crevices such as occur with overlapping surfaces, under seals and gaskets, and in similar
areas. Resistance to this type of attack is improved by adding molybdenum to the alloys. The
most commonly used grade of stainless steel in marine environments contains about 17%
chromium, 12% nickel and 2.5% molybdenum and is commonly referred to by its US AISI
designation, Type 316.
Localised attack is stimulated by the difference in oxygen level within the pit or crevice
and that in the surrounding area. In deaerated sea water, where the oxygen level is low, the
risk of localised pitting and crevice corrosion of stainless steels is greatly reduced. Table 3
gives data on some stainless steels in aerated and deaerated sea water. These show that
pitting in deaerated sea water and high chloride-containing brines is much less severe than in
aerated sea water.
Another characteristic of stainless steels is their ability to remain passive even in very
fast flowing sea water. Table 4 gives data on some materials in fast flowing sea water.
Carbon steel and cast iron are rapidly attacked. The copper base alloys suffer erosioncorrosion, but the stainless steel and nickel-base alloys are virtually unattacked.
Table 3. Stainless Steels in Aerated and Deaerated Sea Water and Brine
Alloy
Environment
Type 316 North Atlantic Ocean
Type 316 Sea water (100°C, 25 ppb 02)
Type 304 Sea water (100°C, 25 ppb 02)
Type 316 130 g/l Cl- (pH 7, 8.25 ppb 02)
13-4 CrNi 130 g/l Cl- (pH 7, 8.25 ppb 02)
P = Pitting
G = General corrosion
Velocity
(m/s)
0
0
0
40
40
Max Depth of
Attack (mm)
2.400 P (486 days)
0.170 P (547 days)
0.600 P (547 days)
0.027 G (/year) [8]
0.180 G (/year) [8]
Table 4. Corrosion in Fast Flowing (35-42 m/s) Natural Sea Water
Alloy
Carbon Steel
Cast Iron
Corrosion Rate (mm/yr)
4.50
13.20
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Gunmetal (85/5/5/5 CuSnZnPb)
Ni Aluminium Bronze
Type 316SS
Ni Cu Alloy 400
1.32
0.97
<0.01
0.01
Where the alloys are to be welded, the low carbon or stabilised grades (L grades) should
be used as the heat of welding can cause carbide precipitation in the heat affected zone, in
conventional grades.
Although Type 316 stainless steel and its variants, such as Type 316L, represent the
greatest tonnage of stainless steel used in desalination, higher grades, with improved
resistance to pitting and crevice corrosion are available and are used particularly in RO
systems. These alloys have higher molybdenum contents than Type 316, and nitrogen is also
added as this improves resistance to localised attack and is inexpensive. In assessing
resistance to pitting and crevice corrosion, a simple formula relating the chromium,
molybdenum and nitrogen contents is often used:
Pitting resistance equivalent number (PREN) = Cr% + 3.3Mo% + 16N%
(1)
The PREN is useful for ranking the resistance of different stainless steels, and it is
generally accepted that if it exceeds 40 the alloy can be considered resistant to pitting and
crevice corrosion in sea water. Austenitic stainless steels with approximately 20% chromium,
20-25% nickel, 6% molybdenum and 0.2% N (PREN = 43) are now widely used in critical
applications in aerated sea water.
Another important group of stainless steels finding increasing application, are the duplex
alloys. These have a mixed ferrite/austenite structure which provides increased strength.
They normally have higher chromium and lower nickel contents than the austenitics, and
when alloyed with molybdenum and nitrogen, they have similar resistance to pitting and
crevice corrosion. Alloys with about 25% chromium, 7% nickel, 3.5% molybdenum and 0.2%
nitrogen (PREN = 40) have similar resistance to the 6% molybdenum austenitics.
Nickel-base alloys, based on nickel-copper (often referred to as MonelTm alloys) and
nickel chromium, have characteristics similar to the stainless steels in sea water. They are
prone to pitting under static and low flow conditions and become passive in fast flowing sea
water. The nickel copper alloys are used in pumps, particularly for shafts, where their high
corrosion-fatigue strength and good corrosion resistance are advantageous.
Titanium
Titanium has excellent resistance to sea water under static, low and high flow conditions
and does not suffer crevice attack. There are many titanium alloys, but the type used in
desalination is pure titanium containing low levels of iron, oxygen and nitrogen. It is
generally referred to as Grade 2 (ASTM B338 Grade 2). It has good resistance to deaerated
sea water and brine but can suffer crevice corrosion in hot deaerated brine [9]. As a result,
laboratory investigations were carried out to define safe limits for the use of titanium (Fig. 4)
[10].
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Oldfield
Figure 4. Crevice corrosion of titanium - safe use limits
Titanium is cathodic to most other metals, and care is needed in multi-metal systems. As
titanium is expensive, it is often used for only part of a system, e.g., for tubing with copper
alloy tube plates and water boxes. In such cases, the copper alloy may suffer severe galvanic
attack. Gehring [11] measured attack on aluminium bronze Type C61400 of 7.4 mm/year
when coupled to titanium in a simulated condenser tube/tube plate test. Simon [12] reported
severe corrosion on tube plates in the same alloy in a steam condenser with titanium tubes.
Galvanic corrosion can be controlled by fitting cathodic protection in the water boxes.
Iron anodes are normally used as these avoid any problem of hydriding of titanium which has
been experienced with impressed current systems [13]. Iron anodes can be used at up to 70°C
in deaerated sea water, but above this temperature hydriding can occur. Anodes of 9% nickel
steel can then be used above 70°C to avoid hydriding [14].
BEHAVIOUR OF MATERIALS IN VAPOUR SIDE ENVIRONMENTS
General Considerations
Sea water contains dissolved gases notably oxygen and carbon dioxide and these are
evolved during thermal processes both to reduce corrosion and to improve heat transfer. At
high temperatures the bicarbonate ions present in sea water decompose releasing large
quantities of carbon dioxide in the first few stages. Where acid dosing is used for antiscaling,
carbon dioxide is released externally in a decarbonator, and less carbon dioxide is generated
during the process.
Oxygen is sometimes removed in a separate deaerator, and levels of 10-20 ppb in the sea
water feed can be achieved reliably. In some cases, the final reject stage is used as a
deaerator, and although this reduces the oxygen level, it is difficult to achieve constantly the
low levels obtained with a separate deaerator.
In addition to the normal gases dissolved in sea water other gases such as ammonia and
hydrogen sulphide are sometimes present and may be evolved during the process. Also, most
sea water is treated with chlorine to control marine growth leaving a low residual amount in
the plant feed. In most cases, this does not cause any problems, but in acid dosed plants, it
can give rise to bromine emission, and in additive plants, where ammonia is present in the sea
water, bromamine and chloramine can be evolved.
Carbon Dioxide Corrosion
Carbon dioxide dissolves in water forming a weak acid, and when sufficient is present to
lower the pH to less than 5, rapid attack on steel can occur [15]. In the case of copper base
alloys, the acidity caused by carbon dioxide in condensates can remove protective films but as
the potential of these alloys is higher than that required for hydrogen release, oxygen must be
present to cause corrosion. However, there is usually sufficient oxygen available to cause
corrosion if the pH falls to a level which damages the film. In this context the problems occur
only where incondensable gases are allowed to accumulate in areas of the vapour space where
venting is poor. It can also occur in vent condensers where the gases removed from the plant
are concentrated prior to exhausting to the atmosphere. Stainless steels and titanium are
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Plenary Lectures
resistant to corrosion by carbon dioxide and are often used in venting systems for this reason.
The solubility of carbon dioxide decreases with increases in temperature, so that for a
given carbon dioxide concentration, the lowest pH levels occur at low temperatures. Figure 5
[16] illustrates this effect for various carbon dioxide partial pressures and temperatures.
Corrosion data on the alloy materials used in desalination are very sparse as vapour
space corrosion has not been studied extensively.
Vapour Side Corrosion due to Halogen Emission
Most desalination plants using sea water are treated with chlorine to control marine
growth in the intake systems. Ideally the chlorination is controlled so that only a low
residual, i.e., 0.1-0.2 ppm, is left at the evaporator, but use of shock dosing and higher than
required injection levels often produce higher levels. When added to sea water, chlorine
reacts with bromides, which are always present, and within a few seconds the chlorine has
been converted to bromine species. As these have a similar antifouling effect to chlorine, in
most cases, this reaction is of no significance. However, the compounds actually present at a
given sea water pH differ. This is illustrated for bromine in Fig. 6 [17] where it can be seen
that bromine gas in solution is present even at pH 7, whereas, as shown in [17], chlorine in
solution is not present above about pH 5. This has important consequences for acid-dosed
plants as the pH is lowered to 4-5 by this treatment and even after decarbonation is still
usually about pH 6. Under these conditions, when deaerated, the bromine gas in solution is
stripped out with other gases and enters the venting system where it can cause severe
corrosion.
Figure 5. Effect of partial pressure of carbon dioxide on pH
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Oldfield
Figure 6. Distribution of bromine in sea water
In additive-dosed plants, where there is no reduction in pH, stripping out of bromine gas
does not occur. However, there have been corrosion problems in the venting systems of these
plants due to halogen compounds evolving by a different mechanism [18]. In this case the
presence of ammonia in the sea water led to the formation of chloramines and bromamines.
As these compounds are volatile, they were stripped out into the venting system where they
decomposed forming acidic halogen compounds.
As chlorine and bromine compounds are highly corrosive, the best method of corrosion
control is to limit their presence by maintaining as low a residual chlorine level in the sea
water feed as is compatible with controlling fouling. This residual can, if necessary, be
removed by adding bisulphite to the feed.
CORROSION IN PRODUCT WATER
Untreated product water from distillation processes is slightly acidic due to dissolved
carbon dioxide. Also, it contains very little dissolved solids, which in natural waters often
provide a scaling and buffering effect, stabilising the pH. As long as it remains deaerated, its
corrosive effects on construction materials such as carbon steel are slight. However, within
the plant, it is usual to handle product water in stainless steel in order to maintain its purity.
Stainless steels are resistant to corrosion in fresh water, and normally no corrosion occurs;
however corrosion problems have been experienced as a result of both hydrotesting and
disinfectant procedures.
When aerated, product water is corrosive to the construction materials commonly used in
distribution systems. It will dissolve lime in concrete and asbestos cement piping, and attack
carbon steel and cast iron resulting in 'red water' when the dissolved iron precipitates. In
order to avoid these problems, the water is hardened by adding calcium carbonate and
bicarbonate.
Product water from RO processes may contain up to 500 ppm dissolved, solids and as
this process does not normally involve deaeration, this water can be corrosive. As RO
membranes remove larger ions such as calcium and magnesium more effectively than those of
sodium and chlorine, the product water may be relatively higher in these aggressive nonscaling ions. In the case of carbon steel and concrete piping, the scaling tendency of the
water should be checked, and if necessary, the water hardened as for distillate. For stainless
steels general corrosion is not a problem but some grades may suffer pitting and crevice
corrosion. Selection of the appropriate grade can be made by means of a computerised expert
system [19].
SELECTION OF MATERIALS FOR MULTI-STAGE FLASH (MSF)
DESALINATION PLANTS
General Considerations
The MSF process is a materials-intensive process, approximately 25 Kg of heat
exchanger tubing is required for each cubic meter of output per day. Evaporator bodies, tube
plates, tube support plates, waterboxes, piping and pumps all require considerable amounts of
materials. In order to optimise materials selection in terms of cost and performance, it is
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Plenary Lectures
necessary to use a variety of materials, restricting the more expensive alloys to those areas of
the plant where they are needed. As the cost of alloy materials, such as stainless steels, have
decreased in cost relative to carbon steel and maintenance costs, there has been a tendency for
the industry to make greater use of alloy materials to achieve higher reliability and lower
overall water costs. Thus, factors influencing materials selection are considered for the main
components of MSF units in the following sections.
Heat Exchanger Tubing
A typical MSF unit can be considered in three sections as regards technical requirements
for tubing. These are the reject section, brine heater and recovery section. Table 5 [20]
shows that the failure rates for tubing in these three sections of the plant vary significantly.
Table 5. Failure Rates (Including Tube Replacements)
for All Alloys in Distillation Plants Tubing
Component
Brine Heater
Heat Recovery
Head Reject
Failure Rate (%)
4.90
0.81
2.46
Heat Rejection Section. This section handles natural sea water, and copper-base alloys
have often been selected for tubing. The most common cause of failure in these alloys is
impingement attack, which can be caused by partial blockage due to debris passing the
screens, unsatisfactory flow conditions in the water boxes and unsatisfactory entry conditions
to the tube, i.e., any condition that can lead to local turbulence and high velocities (i.e., the
normal design velocity is about 2 m/sec.) Table 2 gives data on the resistance of copper base
alloys to this type of attack, and it is useful to note that the cost of these alloys is on the same
order as their resistance to impingement corrosion.
In unpolluted sea water with good intake screening and flow conditions, the copper base
alloys perform well. However, modifications to the sea water composition by pollutants or
by over-chlorination can give rise to corrosion problems. These are considered below.
Effect of Sulphides. Sulphides may be present in sea water from the decomposition of
organic matter, effluents from nearby sewage works or action of sulphate reducing bacteria
(SRB). Normally when sulphides form in the sea water, there is a reduction in oxygen
content, and in some cases the oxygen is completely removed. In this case, corrosion is not
severe. However, in most cases the sulphides exist in aerated sea water, and this mixture of
sulphides and oxygen, or preexposure to sulphides followed by exposure to aerated sea water,
is very damaging to most copper base alloys. The sulphides enter the protective films and
reduce their resistance so that they are easily damaged by flowing sea water at velocities
which they would normally withstand. Table 6 [21] illustrates this effect on four widely used
heat exchanger alloys. Sulphide levels as low as 20 ppb can be damaging to copper base
alloys.
Although copper base alloys such as cupronickels are not prone to severe pitting or
crevice corrosion in static natural sea water, they can suffer this type of attack if sulphides are
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Oldfield
present. Where the heat exchanger has operated in clean sea water and has formed good
protective films, these alloys can endure occasional exposure to sulphides. Ferrous sulphate
dosing can also reduce the effects of sulphides [22]. However where sulphides occur
regularly, the use of titanium can often be justified.
Table 6. Effects of Sulphides and Chlorine on Impingement Attack in Sea Water
Alloy
Depth of Attack (mm)
Sea Water
0.1 ppm
0.25 ppm
Sulphides
Chlorine
Aluminium Brass
0.26-0.28
0.56-0.90
0.07-0.20
90/10 Cupronickel
0.04-0.06
0.24-0.44
0.09-0.10
70 /30 Cupronickel
0.07-0.12
0.62-1.03
0.12-0.15
66/30 /2/2 CuNiFeMn
0.01-0.02
0.98
0.05
Velocity in test zone 7.5 m/sec.
Effect of Ammonia. When organic matter decomposes in sea water, ammonia forms as
well as sulphides. It may be present from other sources such as effluent from ammonia
plants, of which there are several in the Arabian Gulf, and sewage plants. The main effect of
ammonia is to cause severe pitting under crevice conditions with heat transfer [23]. Table 7
gives data on the effect of ammonia under crevice conditions. The pitting is often
accompanied by deposition of copper, sometimes outside the pit. Addition of iron markedly
reduces this type of attack, as shown by data in the table.
Table 7. Crevice Corrosion in Sea Water Containing Ammonia
Alloy
Depth of Attack (mm)
No Iron
Plus Iron
Ammonia 0 ppm 2 Ammonia 0 ppm 2
Aluminium brass
0.010
0.090
0.010
0.000
90/10 Cupronickel
0.015
0.100
0.015
0.000*
70/30 Cupronickel
0.015
0.075
0.000
0.015
66/30/ 2/2 CuNiFeMn
0.000
0.065
0.005
0.000*
*Incipient pits-too small to measure; average of two tests in each case.
Two month test in Campbell test rig. 0.042 ppm iron added continuously.
Effect of Chlorine. Chlorine, added to control marine growth, can influence corrosion of
copper base alloys. As can be seen in Table 8, low levels of residual chlorine are not
damaging, indeed in some cases, they reduce corrosion. However, at higher levels, chlorine
increases the effect of flow, and impingement attack becomes more likely [24, 25].
Effects of sand. Suspended matter, such as sand, can cause corrosion effects on copper
base alloys in sea water. If it forms deposits in tubes, then some deposit attack, which is a
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form of crevice corrosion, can occur beneath the deposit. Under flowing conditions, sand can
cause damage to the protective films on copper alloys giving rise to impingement attack. In
this context, two factors are important, namely, the sand content and the size of the sand
particles [26]. The cupronickels have better resistance to sand than aluminium brass. The
70/30 cupronickel with 2% iron and 2% manganese was developed specifically for use in
waters with high sand content, e.g., 4000 ppm, and is often specified for reject section tubing.
Titanium has excellent resistance to sand erosion and is used where very high sand content
occurs.
Table 8. Effect of Carbon Steel and Ni Resist Type II on Crevice
Corrosion of Type 316 Stainless Steel in Sea Water
Area Ratio
(SS:C. Steel or Ni Resist)
Crevice Attack
No of Sites Maximum Depth
Initiated
of Attack mm
1:0 (Control)
42
3.18
10:1 (Steel)
0
0.00
50:1 (Steel)
0
0.02
50:1 (Ni-Resist)
0
0.00
30-day test; multicrevice assembly with 120 crevice sites; flow velocity 0.5 m/sec
Heat recovery Section. The data from Table 5 indicate that this section of the plant
presents fewer corrosion problems than others. This is to be expected as in the recovery
section, the sea water and brine are usually deaerated so that there is little oxygen to support a
corrosion reaction. The main corrosion risk is from vapour side corrosion, and as this is more
likely in the first few stages, where carbon dioxide is evolved by the decomposition of
bicarbonates in the sea water, these are normally vented directly to the vacuum system. Also,
the tubing is made from materials with higher resistance to vapour side corrosion, that is to
say, 70/30 cupronickel rather than the 90/10 CuNi alloy or aluminium brass.
When vapour side corrosion occurs, it takes the form of fairly uniform thinning of the
tube wall which eventually perforates. The attack is often concentrated at the tube plates or
tube support junction with the tube and with acidic condensates. This is sometimes reported
as a galvanic effect, for example, when the tube support plates are stainless steel, but it can
occur when these are of the same material as the tubing or are even of carbon steel.
Brine Heater. The data in Table 5 indicate that corrosion conditions in the brine heater
are the most severe in the plant. However, if the causes of these failures are examined [20], it
can be seen that most failures are due to mechanical damage during descaling rather than to
corrosion. The need, therefore, for tubing in this section is for mechanical strength and for
this reason, 70/30 cupronickel is the usual choice, and it is normally used at at 1.2 mm
thickness.
Tube Plates
The main requirements of tube plates are mechanical strength and corrosion resistance.
They should be galvanically compatible with the tube material and preferably slightly anodic
to it so as to give some cathodic protection. As they are much thicker than the tube, some
corrosion can be tolerated provided it does not disturb the tube/tube plate joint which is
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Oldfield
usually made by roller expansion.
Water Boxes
The most popular type of MSF design is the cross-tube type, and this requires a large
number of water boxes with interconnecting piping. For large units, the most economic
material for water boxes is carbon steel, but this needs protection if severe corrosion is to be
avoided. Paint coatings have not proven to be reliable means of protection as they are easily
damaged and require increased maintenance and inspection. Thick (3-5 mm) rubber linings
can be used for the reject section but are not reliable for the hot recovery sections where they
can become detached under vacuum. The most common solution for large water boxes is the
use of 90/10 cupronickel clad steel with a 2-3 mm thickness of the alloy material. This has
been used in many plants and is an economic and reliable form of construction.
Evaporator Bodies
The most common material of construction for evaporator bodies is carbon steel. In early
plants, this was the only material used, but experience has shown that it could suffer severe
corrosion, even in deaerated sea water and brine and alloy materials were introduced.
The use of stainless steels such as Type 316 (low carbon or stabilised grades) began with
attempts to repair areas of severe attack and proved very successful. In the absence of oxygen,
the well-known tendency of these alloys to pit in sea water was inhibited, and provided care
was exercised at shutdowns, when oxygen had access to the plant, stainless steel proved
reliable.
An alternative to stainless steels for chamber lining is 90/10 cupronickel. This has been
used successfully for many years for evaporator shells for small plants [27] and in clad steel
plate form can be used to construct large evaporator shells.
Components in the brine spaces such as brine gates, weirs and nozzles, are normally
made from stainless steel even in unlined spaces. The galvanic effect is slight and does not
usually give rise to any problems. However, where only part lining with stainless steel is
used, there is a risk of galvanic attack where the large area of stainless steel meets the carbon
steel. This can be controlled by coating the stainless steel in the final lined stage.
Vapour Spaces
The mixture of water vapour, carbon dioxide and air entering the vapour spaces causes
condensation on the walls and roof of the vapour spaces. These may cause corrosion as they
are slightly acidic (due to the dissolved carbon dioxide), and if sufficient oxygen is present, it
may add to the corrosion rate. Corrosion of steel in these spaces has been studied [16], and it
was concluded that the main cause was oxygen in-leakage to the vacuum stages. Corrosion in
these spaces is normally found in the mid-plant areas, and the highest rates of attack are in
those stages immediately before the stage where cascaded gases are vented. Attack is also
greatest above the distillate transfer trough where gases from the higher temperature stages
are released as the distillate flashes on passing down the plant. Serious cases have been seen
in acid dosed plants with excellent deaeration. These findings, together with the absence of
carbonates in the corrosion product, which is mostly magnetic iron oxide Fe304 led to the
conclusion that carbon dioxide played only a small part in the attack and that air in-leakage
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Plenary Lectures
was the basic cause. Careful attention to air tightness in one affected plant led to a 60%
reduction in the corrosion rate.
As stainless steels have very high corrosion resistance to carbonic acid and low-chloride
distillate, even when oxygen is present, the use of this material for interstage walls and roof
lining can control this type of attack. Stainless steels are often used in these spaces for parts
such as tube support plates, distillate troughs, tube bundle stay bars and vent piping.
However attention to air in-leakage is advised where carbon steel is used.
Demisters
Wire mesh demisters are normally fitted to prevent carryover of droplets of brine.
Because the wire is very thin, usually 0.1-0.2 mm diameter, it is necessary to use materials
with high resistance to general corrosion. Originally the nickel-copper Alloy 400 was widely
used and gave good performance except where sulphides were evolved from the brine. This
alloy is still used but most plants now use Type 316 stainless steel which performs well even
when sulphides are present.
Venting System
Piping and Ejectors. Vacuum on the evaporator is normally produced by steam ejectors
drawing gases from different sections of the plant. The first three stages are normally vented
directly to the system, and gases evolved in lower temperature stages are cascaded to two or
three points before being drawn off into intermediate ejectors and condensers. Where the last
stage is used for deaeration or if a separate deaerator is fitted, the gases are drawn off into the
deaerator condenser. Stainless steels Types 304 or 316 are normally used for the vent piping,
headers within the piping system and baffles within the tube bundles. At lower temperatures
some use has been made of glass-reinforced plastic piping (GRP), but stainless steels are
needed for the higher temperatures, and most plants use them throughout.
The ejectors are also made from stainless steel with nozzles in grades such as Alloy 20.
Provided the chloride content of the gases remains low, the stainless steels work well, the
only risk being from external stress corrosion cracking at temperatures above about 60°C. To
minimise this risk, painting of the stainless steel is advised, particularly under insulation
where chlorides can accumulate.
Condensers. Two types of condenser are used in venting systems, namely, shell and
tube type, and barometric. As they have different materials requirements, they are considered
separately.
The shells of shell and tube type condensers are normally made in Type 316L stainless
steel to resist the incondensible gases. In some cases, copper alloys such as aluminium
bronze are used, but this is not common. For tubing, copper base alloys have a limited life
due to vapour side attack in the high concentration of gases drawn from the plant which are
mixed with ammonia from the ejector steam when hydrazine is used in the boilers. As the
attack is usually fairly uniform, several years of life are possible, and some plants use copper
base alloys with regular renewal. However in most cases it is more economic to use a
resistant material and thin wall welded titanium tubing 0.7 mm thick is the usual choice.
High alloy stainless steel such as UNS S31254 has been used successfully for tubing in
ejector condensers in the Middle East. These were roller expanded into an existing naval
80
Oldfield
brass tube plate and were in perfect condition after one year of service [28]. The advantage
of using this alloy is the ease of descaling compared to titanium.
For shell and tube units, tube plates and water boxes require similar considerations as for
the reject section stages. In general copper base alloys are used with anodes in the water
boxes to control galvanic effects when titanium or stainless steel tubes are fitted.
Barometric condensers are essentially vessels into which sea water is sprayed. The large
deaerator condenser can be made from GRP as temperatures are low; however, at higher
temperatures, metals are necessary. Standard stainless steels such as Type 316L are prone to
pitting in the sea water spray. Anodes are sometimes fitted to prevent this, but in a spray
situation, they are not reliable. As the higher temperature units are relatively small, it is
economic to make them from sea water resistant alloys such as the 6% molybdenum stainless
steels.
Sea Water Pumps
Wet pit vertical lift pumps are normally chosen for large MSF units. Pump parts must
withstand fast flowing sea water, and the normal choice for those parts exposed to the most
severe conditions, such as impellers, is stainless steel. The reason for this is evident from the
data in Table 4 where stainless steel Type 316 is seen to passivate in very high velocity sea
water. The nickel copper Alloy 400 behaves similarly. Although stainless steels are
excellent for flow conditions it is also necessary to consider static conditions, and as can be
seen from Table 6, they are then subject to severe pitting. Thus, an all stainless steel pump
would have problems at shutdown. Such problems have been reported [29].
One method of overcoming this is to make the static parts of the pumps of a material
which gives cathodic protection to stainless steels at shutdowns. The material normally
chosen is Ni-Resist, a high alloy cast iron. Table 8 [30] gives data on Ni-Resist and stainless
steel galvanic couples showing that this combination can control pitting and crevice corrosion
under static conditions.
The usual grade of Ni-Resist used is Type D-2W. This has good sea water corrosion
resistance but is subject to stress corrosion cracking in warm sea water and brine. Failures
have occurred in service on castings which have not been stress relieved. However, when
stress relief heat treated for a few hours at 650-675°C and furnace cooled the alloy has proven
resistant to cracking. It is essential, therefore, to apply this treatment before service and after
any weld repairs.
In recent years, there has been a trend towards the use of duplex stainless steels in sea
water pumps. These are stronger than the standard austenitic alloys such as Type 316 and the
grades used have higher resistance to pitting and crevice corrosion. In this case, to protect
against pitting at shutdowns, it is advisable to operate the pumps for a short period every few
days.
Brine Recycle and Blowdown Pumps
These are in many ways similar to the sea water pumps described above, but are barreltype or canned pumps. The impeller, shaft and other internals are often made from Type 316
stainless steel or its cast equivalent. The casings are often made of Ni Resist Type D-2W, and
these again need to be stress relieved as cracking, particularly of diffusers, has been
81
Plenary Lectures
experienced in service. The barrels or cans are usually fabricated from clad steels. The
cladding is normally Type 316L stainless steel, but 90/10 cupronickel is also used. In some
cases, thick, 3-5 mm, vulcanised rubber is used.
Distillate Pumps
The distillate leaving the plant is slightly acidic but not strongly corrosive until it
becomes aerated. However, to maintain its purity, it is usual to use stainless steels for the
pumps. Type 316 stainless steel is often used. Although this is not really necessary to handle
low chloride distillate, there is always a risk during hydrotesting that it will be exposed to
high chlorides. If this risk can be eliminated, then stainless steels such as Type 304 or 420
could be used.
SELECTION OF MATERIALS FOR REVERSE OSMOSIS (RO) SYSTEMS
General
RO is a membrane process, and it is important to ensure freedom from corrosion
products and similar material in the feed which might block the membranes. The use of
stainless steels for piping and other components in the plant is, thus, advantageous as these
alloys suffer virtually no general corrosion in waters. The problem is to ensure freedom from
pitting, as most waters contain chlorides, are aerated and depending on their concentration
and the other salts in the water, can cause pitting and crevice corrosion. In the case of sea
water, selection is fairly straightforward and is often based on the PREN as described earlier.
In the case of brackish waters, which can vary greatly as regards the concentration of salts
present and their relative amounts, selection of the optimum grade of stainless steel is more
difficult. Some of the factors involved in the selection of alloys for important components in
RO systems are considered in the following sections.
High Pressure Piping: Sea Water Systems
The high pressure piping consists of large-bore pipes which convey high pressure feed
from the pumps to the membrane banks. The feed is then distributed to rows of membrane
cells by headers with connections to individual cells. A similar arrangement handles the
effluent brine and permeate.
Compression fittings, sometimes incorporating rubber O-rings, are used to connect
various sections of the piping, and these provide sites very favourable to crevice corrosion. It
is necessary, therefore, to use stainless steels with high resistance to crevice corrosion in sea
water. Many plants have used a 6% molybdenum austenitic stainless steel. Nordstrom and
Olsson [31] list 13 plants using this type of alloy.
High Pressure Sea Water Pumps
These pumps operate at about 70 bar pressure, and like the piping, should not add
corrosion products to the feed. The problem is to avoid pitting and crevice corrosion and, as
Ni-Resist is not suitable for these pumps (as its strength is too low), this means that alloys
such as 6% molybdenum austenitics or duplex stainless steels with high PRENs should be
used. However, the main risk of pitting occurs at shutdown, and by ensuring that the pumps
do not stand idle for more than a few days, lower grades can be used. Draining and flushing
with low chloride permeate is needed for longer periods of shutdown.
82
Oldfield
Recovery Turbines: Sea Water Systems
The effluent brine from these plants is at high pressure and can be used to recover
energy. The machines used for this purpose are impulse turbines, i.e., Pelton Wheels, or
reaction turbines which are centrifugal pumps running in reverse. Materials selection for the
reaction-turbine type is based on the same criteria as for the high pressure feed pumps. For
Pelton Wheels, duplex stainless steels are favoured. In this case resistance to cavitation
corrosion is important as sheets of cavitation can form at the discharge nozzles, and if these
impinge on the buckets of the wheel, they can cause severe damage.
High Pressure Piping: Brackish Water Systems
Selecting the optimum grade of stainless steel is more difficult in this case as the
compositions of these waters can vary greatly. This problem has been addressed by Oldfield
and Todd [32] using a corrosion engineering guide (CEG) based on a mathematical model
which has been verified by exposure testing. In the mathematical model of crevice corrosion,
the main corrosion problem in these systems has been converted into a user friendly computer
program, which, once given the necessary input data makes predictions of performance.
The CEG also enables the user to see how minor variations in their conditions might
influence the prediction, thus it enables them to see if the material selected is near its limit of
usage or whether there is a margin of safety to allow for variations in conditions. Although
the CEG takes into account many more factors than alternative methods of prediction, there
are factors which it does not consider, e.g., the presence of sulphides, unusual anions. When
such conditions are encountered, additional advice should be sought. However, in most
circumstances the CEG will select an appropriate grade of stainless steel.
High Pressure Pumps: Brackish Water Systems
The selection of suitable grades of stainless steel is based on the same criteria as
selection for piping in the previous section.
SUMMARY AND CONCLUSIONS
When selection of materials is based on their known behaviour in relevant corrosion
environments such as natural sea water, and when this is coupled with service experience,
reliable performance can be obtained in desalination processes. This treatise has documented
relevant data and experience, and has indicated how materials can be selected for the main
components of MSF and RO process plants.
Selection of materials is a continually changing process as new materials are developed,
experience is gained and relative costs of materials and labour change. However, the need to
make the optimal selection of materials will always remain, and if based on sound technical
data and experience, can be achieved.
REFERENCES
1. J. Lyman and R. Babel, Chemical aspects of physical oceanography, J. Chemical
Education 35, 3, 1938, pp. 113-115.
2. C.R. Southwell, J.D. Bultmann and A.L. Alexander, Corrosion of metals in tropical
83
Plenary Lectures
environments: Final report of 16-year exposures, Materials Perform., July 1976, pp. 9-26.
3. H.R. Copson, Effects of velocity on corrosion, Corrosion 16, 1960, pp. 86-92.
4. J.W. Oldfield and B. Todd, Corrosion considerations in selecting metals for flash
chambers, Desalination, 1979, 3, pp. 365-383.
5. F.W. Fink E.L. White and W.K. Boyd, U.S. Department of Interior R&D Progress Report
No. 255, December 1966.
6. D.B. Anderson, Copper-nickel and other alloys for desalination, Inco Publication No.
4319, pp. 39-47.
7. Desalination Materials Manual, Dow Chem for the U.S. Office of Water Research and
Technology, May 1975.
8. G. Pini and J. Weber, Materials for pumping sea water and media with chloride content,
Sultzer Technical Review, 1975, p. 158.
9. S. Kido and T. Shinohara, Corrosion under heat flux encountered in desalination plant,
Desalination, 1977, 22, pp. 369-378.
10. K. Shimogori, H. Sato, F. Kamikubo and T. Fukuzuka, Corrosion resistance of titanium in
MSF desalination plant, Desalination, 1977, 22, pp. 403-413.
11. G.A. Gehring and R.J. Kyle, Galvanic corrosion in steam surface condensers tubed with
either stainless steel or titanium, NACE Corrosion '82, Paper No. 60, 1982.
12. P.D. Simon, Tube sheet corrosion and mitigation techniques in a sea water cooled
titanium-aluminium bronze condenser, NACE Corrosion '83. Paper No. 77, 1983.
13. J.P. Fulford and R.W. Shutz, Characteristics of titanium condenser tube hydriding at two
Florida Power and Light Company plants, ASME/IEEE Power Generating Conference,
Miami Beach, Oct. 1987.
14. T. Fukuzuka, K. Shimogori, H. Sato and F. Kamikubo, Corrosion problems and their
prevention in desalination plant with titanium tube, Proc. of Intl. Congress in Desalination
and Water Re-use, Nice, October 1979, pp. 389-397.
15. P.F. George, J.A. Manning and C.D. Schrieber, Desalination Materials Manual, May
1975.
16. J.W. Oldfield and B. Todd, Vapour side corrosion in MSF plants, Desalination, 1978, 66,
pp. 171-184.
17. J.W. Oldfield and B. Todd, Corrosion problems caused by bromine formation in MSF
desalination plants, Desalination 1981, 38, pp. 233-245.
18. W.S.W Lee, J.W. Oldfield and B. Todd, Corrosion problems caused by bromine
formation in additive dosed MSF plants. Desalination 1983, 44. pp. 209-221.
19. Crevice Corrosion Engineering Guide, Nickel Development Institute, Publication and
Computer Disk D-0003.
20. E.H. Newton, J.D. Birkett and J.M. Ketteringham, Survey of materials in large desalting
plants around the world, A.D. Little & Co., March 1972.
21. R. Francis, The effect of sulphide and chlorine on the corrosion of copper alloy heat
exchanger tubing. BNF Paper R421/5, September 1984.
22. H.P. Hack and T.S. Lee, David Taylor Naval Ship and Research Development Centre
Report DTN SRDC/SME - 81/91, January 1982.
23. R. Francis, The effect of ammonia and chlorine on the corrosion of copper alloy heat
exchanger tubes. BNF paper R241/4, June 1984.
24. S. Sato, Sumitomo Light Metal Industries Review, July 1962.
25. R. Francis, The effect of chlorine additions to cooling water on corrosion of copper alloy
84
Oldfield
condenser tubes, Materials Performance, August 1982, p. 44.
26. S. Sato, Recent aspects of corrosion protection in condenser tubes. Boshoku Gijutsu
(Corrosion Engineering), 24, 6, 1975. pp. 313-331.
27. D.A. Rayney, Choose CuNi tubes, Ni steel for ascension desalination, Nickel 6, 3, 1991.
28. J. Olsson and B. Wallen, Experience with a 6% Mo stainless steel in saline water,
BSE/NACE Conference, Bahrain 1984.
29. S. Zaharani, B. Todd and J.W. Oldfield, Bimetallic corrosion in MSF desalination plants.
Galvanic Corrosion, ASTM STP978 HP Hack Ed, ASTM, 1988, pp. 323-335.
30. T.S. Lee and A.H. Tuthill, Guidelines for the use of carbon steel to mitigate crevice
corrosion of stainless steel in sea water, NACE Corrosion '82. Paper No. 63, 1982.
31. A.J. Sedriks and K.L. Money, Corrosion fatigue properties of nickel-containing materials
in sea water, International Nickel Publication A1258, 1978.
32. J. Nordstrom and J. Olsson, Which stainless steel to use for SWRO plants, Symposium
Internacional De Desalacion y Reuso Del Aqua, Canagua 1992.
33. J.W. Oldfield and B. Todd, Economic selection of stainless steels for desalination and
industrial water applications using a computerised corrosion engineering guide (CEG),
6th Middle East Corrosion Conf. Vol. 2, January 1994, p. 771.
85
Plenary Lectures
APPENDIX
Nominal Composition of Alloys
Copper Base Alloys
Alloy
UNS No
Cu
Ni Fe Mn Zn Al Sn Pb Other
(%) (%) (%) (%) (%) (%) (%)
(%)
90\10 Cupronickel
70\30 Cupronickel
66\30\2\2
CuNiFeMn
Aluminium Brass
C70600
C71500
C71640
Rem
Rem
Rem
10
30
30
1.5
1
2
1
1
2
-
-
-
-
C68700
Rem
-
-
-
22
2
-
-
Admiralty Brass
C44300
Rem
-
-
-
29
-
1
-
Naval Brass
Aluminium Bronze
NiAl Bronze (W)
Admiralty Gunmetal
Leaded Gunmetal
NiAl Bronze (C)
Rem = Remainder
W = Wrought
C = Cast
C46400
C61400
C63000
C90300
C83600
C95800
Rem
Rem
Rem
Rem
Rem
Rem
5
5
2
4
4
1
1.5
1
37
2
5
-
7
10
9
1
10
5
-
5
-
0.04A
s
0.04A
s
Titanium
UNS No Ti (%)
Grade 2
Fe
(%)
R50400
Rem
0.3
Maximum % in all cases
O( %) N( %) H (%) C (%)
0.25
0.03
0.015
0.1
Stainless Steels
a)
Austenitic
Alloy
304
304L
316
316L
317
317L
317LMN
86
UNS No
S30400
S30403
S31600
S31603
S31700
S31703
S31726
C (%)
0.08
0.03
0.08
0.03
0.08
0.03
0.03
Cr (%) Ni (%) Mo (%)
18
10
18
10
17
12
2.5
17
12
2.5
19
13
3.5
19
13
3.5
18
13
4.5
Other
0.18N
Oldfield
904L
b)
0.03
UNS No
S31254
N08926
N08367
S31254
C%
0.03
0.03
0.03
0.02
20
25
4.5
1.5Cu
Superaustenitics
Alloy
254SM O
1925 hMO
AL-6XN
HR25 4
c)
N08904
Cr (%) Ni (%) Mo (%) N (%) Other
20
18
6
0.2 0.7Cu
21
25
6
0.19
21
25
6.5
0.2
20
18
6
0.2 0.7Cu
Duplex and Superduplex Alloys
Alloy
2205
3RE6O
Ferralium
Zeron 100
SAF 2507
DP3W
UNS No
S31803
S31500
S32550
S32760
S32750
S32974
C (%) Cr (%) Ni (%) Mo (%) N (%)
Other
0.02
22
5.5
3
0.17
0.03
18.5
5
2.5
0.1
0.04
25
6
3
0.18
2Cu
0.03
25
7
3.5
0.25 0.8W + 0.8Cu
0.03
25
7
4
0.3
0.02
25
7
3.2
0.3
2W + 0.5Cu
Nickel Base Alloys
Alloy
Alloy 400
Alloy 500
A lloy 825
Alloy 625
UNS No C (% max) Cr (%) Ni (%) Mo (%) Cu (%)
Other
N04400
63 min
31
2Fe
N05500
63 min
31
2.7Al + 0.6Ti
N08825
0.05
21
42
3
2.5
30Fe
N06625
0.1
21
Rem
9
3.5Nb
Rem = Remainder
Ni Resist Alloy Cast Irons
Alloy
UNS No C (% max) Cr (%) Ni (%) Other
Type 2
F41002
3.0
2
20
Type D-2
F43000
3.0
2
20
Type D-2W
3.0
2
20
0.15Nb
Graphite Form
Flake
Spheroidal
Spheroidal
87
Industrial Corrosion and Corrosion control Technology
sbalaby, H.M et aI. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
CORROSIVITY PREDICTION FOR
CO2/H2S PRODUCTION ENVIRONMENTS
CLl International, Inc. 14503
Bammel N. Houston # 300 Houston,
TX 77014, USA
One of the most fundamental issues in current day corrosion research is
assessment of corrosion rates in steels and determination of corrosivity of
typical operating environments in oil and gas production. Such assessment
requires an understanding of the role of primary environmental and
metallurgical variables and underlying mechanisms of corrosion. This paper
presents a novel hierarchical approach to assess system corrosivity and predict
corrosion rates in carbon steels in production environments containing CO2
and/or H2S. Critical environmental parameters that influence system
corrosivity are identified and the effects of these parameters on corrosion are
examined. Modeling for synergistic assessment of system corrosivity as a
function of relevant operating parameters is presented and is accompanied by
a description of a computer program to capture the model described.
Key Words: CO2 H2S, Corrosivity, Prediction, Corrosion rates,
Computer Model
INTRODUCTION
CO2/H2S corrosion in oil and gas production environments represents one
of the most important areas of corrosion research. It is so because of the
criticality of the need to assess corrosive severity as a means to ensure safe
utilization of steels, which have wide application in just about every sphere of
oil and gas production and refining. Even though CO2/H2S corrosion and
concomitant mechanisms have been areas of significant work over the last
thirty years, there still exists a need to accurately predict corrosivity of
CO2/H2S environments from a standpoint of defining limits of use for carbon
steels. Even though numerous predictive models have been developed and are
being developed [1,2]; most of the available predictive models tend to be
either very conservative [3] in their interpretation of results or focus on a
narrow range of parametric effects, thereby limiting the scope of the model's
application in realistic assessment of corrosivity and corrosion rates. Often
times, data required by the models are not easily accessible or available to the
operators who need to employ the model, thereby limiting the applicability of
the models to situations of reduced practical importance [4,5]. In this context,
the issue of corrosivity assessment for carbon steels can be restated in terms of
the following critical requirements:
- Development of a predictive model that utilizes commonly available
operational parameters,
- Utilization of existing laboratory/field data and theoretical models to
obtain realistic assessments of corrosivity and corrosion rates, and
- Development of a computational approach that integrates both
numerical (read "laboratory trends") and heuristic (i.e., field data and
experience) information and knowledge about corrosivity prediction.
In this paper, a methodology to determine system corrosivity and to predict
corrosion rates in steels is described, consistent with the objectives stated
above. The method adopted here attempts to capture both the effect of critical
parameters on corrosion rates as well as that of parameter interactions. The
model described in this paper has been encoded into a Widows™-based
computer program, Predict™, that would allow the end user to predict
corrosion rates.
The primary variables in corrosivity prediction in the model are the acid
gases CO2 and H2S that contribute to the typically acidic pH found in
production environments. The model uses the widely accepted de WaardMilliams [3] relationship for CO2 corrosion for an initial determination of
CO2-based corrosion rates. However, the effective CO2 partial pressure in the
system is not based on the operating partial pressure but on one obtained from
the system pH. This rate is further refined to account for the presence of H2S,
corrosion products, temperature effects etc. A technical description of
different corrosivity modeling parameters and their effects is given in ensuing
sections of this paper. The underlying idea here has been to develop a
prediction model that accurately represents the state-of-the-art in theoretical
analyses as well as parametric correlations based on laboratory and field data.
The model has also been compared with actual field conditions in an effort to
compare system predictions [6] with field observations.
In developing any corrosivity model, it is important to recognize the role of
superposition of different parameters. Such recognition requires a clear
understanding of independent parameter effects and how corrosion rate
progresses when subjected to the effects of two or more variables. While the
current prediction model is primarily concerned with environmental
constituents and their corrosive effects, it is also important to recognize the
significant role of metallurgy in fashioning appropriate corrosion behavior.
The influence of compositional and alloying elements has been chronicled but
has hitherto not been rigorously studied in assessing resistance to system
corrosivity. A brief discussion of metallurgical factors in corrosivity
determination is provided elsewhere in this paper.
CO2/H2S-BASED CORROSION: TECHNICAL BACKGROUND AND
LITERATURE REVIEW
CO2-based corrosion has been one of the most active areas of research,
with several predictive models for carbon steel corrosion assessment. These
efforts range from a predictive model that begins with CO2 corrosion [2,3] to
models that focus on specific aspects of the corrosion phenomena (such as
flow-induced corrosion or erosion corrosion) [4,5] and models that
empirically relate corrosion rates to gas production and water production rates
[7]. Crolet and Bonis [8] use the. physical chemistry of the corrosive medium
as the key notion and take
into account ionic strength, pH and specific ionic species as relevant factors.
Other relevant efforts include those by Ikeda et al. [9] that look at the
influence ofH2S and O2 on CO2-based corrosion as well as those by Adams
et al. [10]. Many of these efforts suffer from significant drawbacks in that
- They focus on a narrow range of parametric effects, for example, there is
relatively little published information on the effects of H2S in production
systems and how sulfide scaling can affect the CO2 corrosion process,
- Some models focus on just one component of corrosivity, such as erosional
effects, wall shear stress effects or flow effects, and have opted to ignore
effects of chemical species (i.e., factors such as pH, H2S, CO2 etc.), and
- Other models totally rely on laboratory data for predictive modeling, with
the consequence that the simplifying assumptions made in developing
laboratory models often lead to results that can be far removed from what
is observed in the field.
The current model attempts to integrate laboratory data and field
experience within the framework of the relevant controlling parameters that
are most prominent in oil and gas production. It is important to realize that
while arcane theoretical models are interesting from an academic standpoint,
the controlling parameters in a model must also represent data easily available
to oil and gas production personnel. The current model attempts to integrate
principles hitherto delineated in developing the predictive model.
While there have been several studies focusing on the exact mechanism of
metal dissolution in CO2-containing waters, the efforts of de Waard and
Milliams, and others [2,3,9] present a commonly accepted representation
wherein anodic dissolution of iron is a pH dependent mechanism as given by
Bockris [2], and the cathodic process is driven by the direct reduction of
undissociated carbonic acid. These reactions can be represented as [3]
The buildup of the bicarbonate ion can lead to an increase in the pH of the
solution until. Conditions promoting precipitation of iron carbonate are
reached, leading to the reaction given below
Iron carbonate solubility, which decreases with increasing temperature, and
the consequent precipitation of iron carbonate are significant factors in
assessing corrosivity. The charge transfer controlled reaction involving
carbonic acid and carbon steel (or Fe) can be represented in terms of the
concentration or partial pressure of dissolved CO2 in the medium to
arrive at a corrosion rate equation that incorporates the order of the reaction
and an exponential function that approximates for Henry's reaction constant's
temperature dependence. This corrosion rate equation is given as [2]
The corrosion rate obtained by Eq. 4 has typically been seen as the
maximum possible corrosion rate without accounting for iron carbonate
scaling. A nomogram representing Eq. 4 is given in Fig. 1 [2], which also
includes a scale factor to account for the formation of protective carbonate
films that lead to a reduced corrosion rate at higher temperatures.
The above correlation describes CO2-based corrosion. There have been
other significant efforts to demonstrate the effects of other environmental
variables such as pH, H2S, chlorides, bicarbonates, water/gas/oil ratios,
velocity etc. Effects of H2S on corrosion rates in the laboratory have been
studied and presented by Videm and Kvarekval [11], and Ikeda et al. [9].
Ikeda et al.' s work indicates that the preferential formation of an iron sulfide
film can decelerate the corrosion rate, especially at temperatures above 20°C
and extending up to 60°C. Above I50°C, the corrosion reaction falls back to
the standard CO2-based corrosion with an FeC03 film that is more stable than
the FeO2 film. Videm and Kvarekval's work supports the theory that even
small amounts of H2S can provide instantaneous protection at temperatures in
the range of 70-80°C.
•
Lotz et al. [12] have chronicled the role of the hydrocarbon condensate
in providing corrosion mitigation in specific production systems. The role of
the type of oil or gas condensate is important from the standpoint of accurate
assessment as reported by Choi et al. and Efird [13,14]. Other studies
evaluating effects of critical parameters such as pH and velocity on CO2
corrosion include those by Dugstad and Lunde [15] as well as Lotz [16].
Other predictive models also include those by Gunatlun [17] and Bonis and
Crolet [18] wherein a combination of the parameters discussed herein along
with electrochemical considerations have been utilized to arrive at a
determination of the corrosion rate.
The primary objective of the corrosivity prediction model described in
this paper is to address the need of developing a predictive method that
would synthesize different parametric relationships based on information
from literature, laboratory research/data and practical experience/expertise.
It has often been observed that laboratory data and the ensuing models
represent poor and often inadequate simulation of field conditions [19]. It is
also necessary to understand that field data is typically sparse and can be
negated by other production data. The need to integrate field data/experience
and laboratory models stems from the fact that the laboratory data can
provide significant pointers and trends that can be used in conjunction with
field data and experience. The idea is to develop a methodology that can
integrate analytical and heuristic models. To this end, this predictive system
mirrors other successful development efforts undertaken by the authors in
the areas of evaluation of CRAs and cracking in steels [20,21]. The central
theme is to develop a computer program that can bring together different
types of modeling knowledge to provide a realistic solution to the significant
question of predicting corrosion rates in typical production environments.
CORROSIVITY PREDICTION MODEL DESCRIPTION
A flowchart delineating the hierarchical reasoning structure of the
predictive model is given in Fig. 2. The first step in corrosivity
determination is computation of the system pH, since it is the hydrogen ion
concentration that drives the anodic dissolution. Further, the role of pH in
promoting or mitigating CO2-based corrosion has been extensively
chronicled [19,22]. For production environments, where it is the dissolved
CO2 or H2S that contributes significantly to a suppressed pH, the pH can be
determined as a function of acid gas partial pressures, bicarbonates and
temperature, as shown in Fig. 3 [23]. From a practical standpoint, the
contribution of H2S or HC03, or temperature to pH determination is another
way of representing effective levels of CO2 that would have produced a
given level of pH.
This type of pH determination has been found to be quite accurately
applicable in other modeling efforts involving verification of the relationship
given by Bonis and Crolet [23]. While it has been documented that the CO2
corrosion mechanism is dissimilar to that of strong acids like HCI (as CO2
corrosion is now understood to progress through direct reduction of H2C03
to HC03 - rather than reduction of II' ions), and that carbonic acid corrosion
is much more corrosive than that obtained from a strong acid such as HCI at
the same pH [19], there is also significant agreement that lower pH levels
obtained from higher acid gas presence leads to higher corrosion rates.
Conversely, higher levels of pH obtained through buffering in simulated
production formation water solutions have been shown to produce
significantly lower corrosion rates even at higher levels of CO2 and/or H2S
[24]. Data about the effects of pH from another study are shown in Fig. 4
[15]. Hence, it is more meaningful to determine the
Where Cl and C2 are constants, pH2S and pCO2 are partial pressures in bar
and the HCO3 concentration is represented in meq/l (61 mg/l).
The system pH is given by the larger number between PHI and pH2.
Correspondingly, if the temperature is higher than 100°C, there is a slight
reduction in the hydrogen ion concentration as shown in Fig. 3, but the
change in pH can be accounted for by a change in the value of the constants
in Eq. 5 and 6 above. Once the system pH is determined, the effective CO2
partial pressure can be determined from Eq. 5 as,
where pCO2. is the effective partial pressure of CO2 in a production system
that can produce the prevalent level of hydrogen ion concentration.
The effective CO2 partial pressure from Eq. 7 can be used in Eq. 4 to
determine an initial corrosion rate for C02-based corrosion. The corrosion
rate so obtained is modified to account for the formation of a FeCO3 film
(Fe304 at higher temperatures), the stability of which varies as a function of
the operating temperature. The scale correction factor shown in Fig. 1 is
used to determine the initial corrosion rate from the nomogram in Fig. 1 [2].
It is generally estimated that this corrosion rate presents a maximum
corrosion rate even though it has been reported that the rates computed by
the nomogram are reached or exceeded in systems with high flow rates. It is
important to recognize that this corrosion rate has to be modified to account
for the effect of other critical variables in the system. Further, this rate does
not indicate modality (general or localized), but rather, represents the
maximum rate of attack.
As mentioned earlier, it is necessary to superposition the effects of
other critical system parameters. The flowchart in Fig. 2 provides a list of
the sequential effects that are important from a standpoint of corrosivity
determination. In addition to the system pH, these include:
H2S partial pressure, maximum operating temperature, dissolved chlorides,
gas-to-oil ratio, water-to-gas ratio/water cut, oil type and its persistence,
elemental sulfur/aeration, fluid velocity, type of flow, and inhibition type
and efficiency.
H2S
Oilfield production environments, in recent years, have been
characterized by an increasing presence ofH2S and related corrosion
considerations. Even though H2S is probably the most significant concern in
current-day corrosion and cracking evaluation, the role ofH2S in corrosion
in steels has received much less attention when compared to the widely
studied CO2 corrosion. [26]. However, H2S-related corrosion and cracking
has remained one of the biggest concerns for operators involved in
production because of the significance of H2Srelated damage [27].
In the current modeling effort, in addition to the contribution in pH
reduction, H2S has a threefold role:
- At very low levels of H2S « 0.01 psia), CO2 is the dominant corrosive
species, and at temperatures above 60°C, corrosion and any passivity is a
function of F eC03 formation-related phenomenon, and the presence of
H2S has no realistic significance.
- In CO2 dominated systems [26,28], the presence of even small amounts
of H2S (ratio of pC02/pH2S > 200), can lead to the formation of an iron
sulfide scale called mackinawite at temperatures below 120°C. However,
this· particular form of scaling, which is produced on the metal surface
directly as a function of a reaction between Fe++ and S· ., is influenced
by pH and temperature [27]. This surface reaction can lead to the
formation of a thin surface film that can mitigate corrosion. The authors
are currently pursuing laboratory studies to characterize the stability and
formation of mackinawite in sour systems.
- In H2S dominated systems (ratio of pCO~pH2S < 200), there is a
preferential formation of a metastable sulfide film in preference to the
FeC03 scale; hence, there is protection available due to the presence of
the sulfide film in the range of temperature of 60-240°C. Here, initially it
is the mackinawite form of H2S that is formed as a surface adsorption
phenomenon. At higher concentrations and temperatures, mackinawite
becomes the more stable pyrhotite. However, at temperatures below 60°C
or above 240°C, the presence of H2S exacerbates corrosion in steels
since the presence of H2S prevents the formation of a stable FeC03 scale
[9,29]. Further, it has been observed that FeS film itself becomes unstable
and porous and does not provide protection. Also, the scale factor
applicable for CO2 corrosion with no H2S (shown in Fig. 1) becomes
inapplicable. Even though there is agreement amongst different workers
that there is a beneficial effect of adding small amounts of H2S at about
60°C, Ikeda et al. [9] and Videm and Kvarekval [11] present divergent
results at higher concentrations and higher temperatures.
The effect ofH2S adopted in the predictive model reflects work published
by Murata et al. [29] for CO2dominated systems. Figure 5 [29] shows the
combined effects of temperature and gas composition on the corrosion rate
of carbon steels. Figure 6 [9] shows the effect of varying degrees ofH2S
contamination on C02 corrosion. It is to be noted that the role ofH2S in C02
corrosion is a complex issue governed by the film stability of FeS and
FeC03 at varying temperatures and is an area of further active research by
the authors.
Temperature
Temperature has a significant impact on corrosivity in CO~2S systems.
Corrosion rate as a function of different levels of CO2 and temperature is
given in Fig. 7 [2]. It has to be noted that once the corrosion products are
formed, there is a significant mitigation in corrosivity. It is also apparent that
the carbonate film is more stable at higher temperatures and affords greater
protection at higher temperatures. Figure 7 also shows that at temperatures
beyond 120°C, the corrosion rate is almost independent of the CO2 partial
pressure of the system. The carbonate film may, however, be weakened by
high chloride concentrations, or it can be broken by high velocity. In H2Sdominated systems, because no carbonate scale may be formed and the FeS
film becomes porous and unstable at temperatures beyond 120°C, significant
localized corrosion may be observed.
Chlorides
Produced water from hydrocarbon formations typically contains
varying amounts of chloride salts dissolved in solution. The chloride
concentration in this water can vary considerably, from zero or a few parts
per million in condensed water to saturation in formation waters having high
total dissolved salts/solids (TDS). In naturally deaerated production
environments, the corrosion rate increases with increasing chloride ion
content over the range 10,000 ppm to 100,000 ppm [30]. The magnitude of
this effect increases with increasing temperature over 60°C (1 50°F). This
combined effect results from the fact that chloride ions in solution can be
incorporated into and penetrate surface corrosion films, which can lead to
destabilization of the corrosion film and increased corrosion. This
phenomenon of penetration of surface corrosion films increases in
occurrence with increases both in chloride ion concentration and in
temperature.
Bicarbonates
Bicarbonates in the operating environment have a significant impact on
corrosion rates. On one hand, high levels of bicarbonates can provide higher
pH numbers leading to corrosion mitigation even when the partial pressures
of CO2 and HzS are fairly high. Bicarbonates, which can be present in
substantial quantities in formation waters (up to 20 meq/l) [31], have a
natural inhibitive effect when present. Condensed water in production
streams typically contains no bicarbonates.
Velocity
Next to the corrosive species that instigate corrosion, velocity is
probably the most significant parameter in determining the corrosivity of
production systems. Fluid flow velocities affect both the composition and
extent of corrosion product films. Typically, high velocities (> 4 m/s for
uninhibited systems) in the production stream lead to the mechanical
removal of corrosion films, and the ensuing exposure of the fresh metal
surface to the corrosive medium leads to significantly higher corrosion rates.
Corrosion rate as a function of flow velocity and temperature is shown in Fig.
8 [15].
In multiphase (i.e., gas, water, and/or liquid hydrocarbon) production,
the flow rate influences the corrosion rate of steel in two ways. First, it
determines the flow behavior and flow regime. In general terms, this is
manifested as static conditions (i.e., little or no flow) at low velocities,
stratified flow at intermediate conditions, and turbulent flow at higher flow
rates. One measure which can be used to define the flow conditions is the
superficial gas velocity. In liquid (i.e., oil / water) systems, this is replaced
with the liquid velocity.
Velocities less than 1 m/s are considered static. Under these conditions
corrosion rates can be higher than those observed under moderate flow
conditions. This occurs because under static conditions, there is no natural
turbulence to assist the mixing and dispersion of protective liquid
hydrocarbons or inhibitor species in the aqueous phase. Additionally,
corrosion products and other deposits can settle out of the liquid phase to
promote crevice attack and underdeposit corrosion.
At velocities between 1 and 3 m/sec, stratified conditions generally still
exist. However, the increased flow promotes a sweeping away of some
deposits, and increasing agitation and mixing. At 5 m/sec flow velocity,
corrosion rates in uninhibited applications start to increase rapidly with
increasing velocity [31]. The data shown in Fig. 9 [31] demonstrate the
effects of velocity on corrosion rate for both inhibited and non-inhibited
systems. For inhibited applications, corrosion rates of steel increase only
slightly at velocities between 3 and 10 m/sec, as a result of mixing of the
hydrocarbon and aqueous phases. At velocities above about 10 m/sec,
corrosion rates in inhibited systems start to increase due to the removal of
protective surface films by the high-velocity flow.
Flow-related effects on corrosivity have been linked to the wall shear stress
developed and is an area of intense research in the community [32]. Flowinduced corrosion is a direct consequence of mass and momentum transfer
effects in a dynamic flow system where the interplay of inertial and viscous
forces is responsible for accelerating or decelerating metal loss at the
fluid/metal interface. While flow-induced corrosion is a significant
component of predictive modeling discussed herein, the topic of flow-related
effects is being actively researched by the authors and forms the focus of
another publication. Another relevant aspect of flow- or velocity-induced
corrosion is erosion-corrosion [33] and refers to the mechanical removal of
corrosion product films through momentum effects or through impingement
and abrasion. Guidelines for velocity limits with respect to erosional
considerations are given in API-14E in terms of the density of the fluid
medium [34].
Water/Gas/Oil Ratios
The predictive model classifies systems as oil dominated or gas dominated
on the basis of the gas/oil ratio (GOR) of the production environment. If the
environment has a GOR < 890
m3/m3 (5000 scf/bbl in English units) [35], the tendency for corrosion and
environmental cracking is often substantially reduced. This is caused by the
possibly inhibiting effect of the oil film on the metal surface, which
effectively reduces the corrosivity of the environment. However, the
inhibiting effect is dependent on the oil phase being persistent and acting as
a barrier between the metal and the corrosive environment. The persistence
of the oil phase is a strong factor in providing protection, even in systems
with high water cuts. In oil systems with a persistent oil phase and up to
45% water cut, corrosion is fully suppressed, irrespective of the type of
hydrocarbon [12]. Relative wet ability of the oil phase versus the water
phase has a significant effect on corrosion [36]. Metal surfaces that are oil
wet show significantly lower corrosion rates [37].
The predictive model described in this paper provides for a significant
reduction in the corrosion rate (up to a factor of 4) based on the type of oil
phase, i.e., persistent, mildly persistent and not persistent. However, the
degree of protection can be quantified only as a function of water cut and
velocity. The persistence determination is a more complex task and requires
knowledge of the kerogen type and hydrocarbon density. It is important to
understand the type of crude oil in terms of the organic compounds that
make up the crude to determine wettability effects. Figure 10 shows data that
relate the acid number of the crude to oil wettability, and Fig. 11 shows
corrosion rate as a function of produced water content for different crude
oil/produced water mixtures [36]. While the effect of persistence of the oil
medium is significant on corrosion rates, it is even more difficult to quantify
precise compositional elements of oil medium that contribute to wettability
and persistent oil film formation. Such quantification is possible by rigorous
laboratory testing of different actual, uncontaminated (read "deaerated")
production water samples, so as to determine the extent of protection.
In oil systems, the water cut acts in synergy with the oil phase to
determine the level of protection from the hydrocarbon phase. However, at
very low water cuts « 5%), the corrosive severity of the environment is
lessened due to the absence of adequate aqueous medium to promote the
corrosion reaction.
In gas-dominated systems, there are two measures to evaluate the
availability of the aqueous medium. If the operating temperature is higher
than the dew point of the environment, no condensation is going to be
possible, leading to highly reduced corrosion rates. Corrosion under
condensing conditions (i.e., operating temperature less than the dew point) is
a function of the rate of condensation and transport of corrosion products
from the metal surface [38]. If the total water in a condensing system as
measured by the water-to-gas ratio is < 11.3 m3/Mm3 (2 bbl water/MSCF
gas), corrosivity is substantially reduced.
Aeration/Sulfur
The presence of oxygen significantly alters the corrosivity of the
environment in production systems. Old field and Todd [39] chronicled how
the presence of oxygen could significantly increase corrosion rates due to
acceleration of anodic oxidation. While the corrosion rate increases with
oxygen, the rate of oxygen reduction as a catholic reaction is further
exacerbated by
- Increased operating temperature,
- Increased fluid flow leading to increased mass flow of oxygen to the
metal surface, and
- Increased oxygen concentration
Data showing increases in corrosion rate as a function of oxygen
concentration for differing temperatures are shown in Fig. 12 [39]. The
presence of elemental sulfur is similar to that of free oxygen since elemental
sulfur also acts as a strong oxidizing agent.
Inhibition/Inhibition Effectiveness
Appropriate inhibition is a critical criterion for effective use of carbon steels
in corrosive production systems. Inhibition has been typically found to be
viable in flows with velocity in the range of 0.3-10 mfs. Requirements for
the type of inhibitor and the method of delivery depend on the type of
system (Le., production tubing or horizontal flow lines) to be inhibited.
Inhibition efficiency (IE) describes the efficacy of an inhibitor treatment in
mitigating weight loss corrosion and is an important factor in assessing
corroslV1ty. It is based on either laboratory or field data where inhibited and
uninhibited corrosion rates are compared using the following equation:
Values of IE near 1.0 represent conditions with maximum efficacy of the
inhibitor treatment. Conditions which affect IE include:
-
Inhibitor concentration,
Severity of the corrosive environment,
Service temperature,
Solubility of the inhibitor in the aqueous phase,
Phase behavior of the inhibitor and carrier fluid in the service
environment, and Persistence of the inhibitor on the metal surface.
The predictive model evaluates inhibition efficacy on the basis of
velocity, hydrocarbons to-water ratio and dissolved chloride levels. The
method of delivery (e.g., batch, continuous, pigging, etc.) is also an
important factor in determining the appropriateness of inhibition for a given
set of operating conditions.
The corrosion rate predicted in the current model can be represented in
terms of three broad rules that guide the computer model's decision making:
- Effect of fundamental system variables such as CO2, H2S, pH,
temperature, and velocity on corrosion rate;
- Effect of parameter interactions on corrosivity, such as, the influence
of temperature on the carbonate or sulfide film stability, or flow
effects on corrosion products and the ensuing loss of protective films
as a function of velocity, temperature, acid gases and pH; and
- Effects of system modifiers such as oil film persistence (or lack of it)
or the crude type, water cut, dew point, aeration and inhibition.
Corrosion rate, thus predicted, incorporates the synergy of the effects of
all the critical system variables and provides a more realistic estimation of
corrosivity than would be available with conservative theoretical models that
focus on a limited number of parameters. The significance of the reasoning
in the predictive model stems from the fact that the decisions made
synthesize different types of corrosion knowledge:
- Theoretical models that provide the effects of different parameters,
- Data from laboratory tests that provide insight on parametric
correlations and trends of parametric effects, and
- Experience-based heuristics that facilitate proper interpretation of data
from both the laboratory and field.
The predictive model in this paper has been implemented as a Windowsbased computer program with an interface as shown in Fig. 13. Based on
data specified for different parameters, the system will instantaneously
display the following results:
- System pH,
- Predicted corrosion rate called corrosion index (in mpy or mmpy),
- A textual recommendation in the results box indicating whether the
predicted corrosion rate is within the specified allowance for the
particular system, and
- A corrosion index bar that graphically represents the corrosion rate.
The user can specify data for any of the parameters and watch the effect
of that parameter on the corrosion rate in the system instantaneously. The
system starts with a set of default values and calculates a corrosion rate
based on any changes to the displayed values. A typical consultation will
involve the following five steps:
- Specification of pH-Related Data: At the outset, the system determines
a corrosion rate only if the operating environment is acidic or has
aeration. If the specified environment has no acid gases or there is
sufficient buffering to produce a pH higher than 7.0, the system will
predict zero or very low corrosion rates, except under conditions of
aeration. So, the first step in consulting the system involves the
specification of the acid gas (H2S and CO2) partial pressures as well
as the bicarbonate content of the environment.
- Temperature/Gas-Water Ratios: Temperature has a significant impact
on corrosion rates as described in the previous section. Corrosion rates
typically increase with increasing temperature. If the gas-to-oil ratio
indicates gas-dominated conditions (as opposed to an
oil-dominated system), the system uses the water-to-gas ratio and the dew
point as the means to determine the availability of an aqueous medium to
measure corrosion. So, depending on the value entered for the gas-to-oil
ratio, the system will let the user specify the relevant water-related
parameters. If the gas-to-oil ratio is < 5000 scribble (which denotes an oil
well), the system uses the water cut and oil persistency to determine the
wetness effect.
- Chlorides/Sulfur: Chloride and sulfur typically make corrosion worse if
the process has been initiated by the presence of acid gases. Their role,
while not as critical as that ofH2S or CO2, is significant because these
parameters can significantly increase corrosion rates in mildly corrosive
systems.
- Velocity/type of Flow: Flow parameters are very critical in both
determining and controlling corrosion effects. Erosion corrosion as well
as the protection (or the lack of it) from corrosion films is very much a
function of fluid velocity.
- Inhibition/Corrosion Allowance: Inhibition choices in the system allow
the user to select applicable methods of inhibition for vertical or
horizontal flow and determine the extent of corrosion mitigation. In some
cases, the system might provide no protection due to inhibition because
of high velocities or chloride concentrations. The system's rules assess
the appropriateness of the method of inhibition delivery for a given set of
conditions.
FUTURE WORK
Predicting the corrosivity of production environments is a complex and
challenging task from several standpoints. While the system described in this
paper captures the effects and interaction of several critical parameters,
significant opportunity exists for enhancing the system's analytical and
modeling capabilities. Further work in refining the predictive model
described herein is governed by the following factors:
Many of the fundamental mechanisms driving corrosion are well understood;
however, there is still a large body of ongoing research grappling with
providing accurate phenomenological models. Formation of sulfide films
and their stability as a function of temperature and pH are areas that require
quantification and better modeling.
A large number of parameters influence the corrosion process, and a
complex set of parameter interactions exacerbate or mitigate corrosion. New
data generated in the laboratory or field is critical to refining the existing
model.
The current work does not include mass transfer and momentum transfer
effects on surface corrosion. However, further work is aimed at including
flow models that capture the effects of inertial and viscous forces in single
phase and multi-phase flows. Wall shear stress developed as a function of
prevalent flow regimes has a direct influence on corrosion rate and is the
focus of considerable research in both industry and academia [32,40].
The metallurgy of steels used -m CO:z/H2S production environments is
critical to determining environmental corrosivity.Metallurgical factors
include microstructure, material processing (e.g., annealed, quenched and
tempered, normalized etc.) and other morphology-related factors like
hardness of welds and residual stresses. The addition of residual and
alloying elements (e.g., Cr, Cu, Ni, etc.) has been shown to have a
significant impact on corrosion performance [41,42]. While there is some
data available for understanding the effects of metallurgy on corrosivity in
CO2 environments, very little information is available on metallurgical
effects versus corrosivity in H2S environments, and this represents an area
of active research. The authors have currently initiated a research program
on testing to quantify metallurgical effects as they relate to corrosion in
typical production environments [43].
CONCLUSIONS
Predicting the corroslV1ty of CO:z/H2S production environments is a
complex and challenging task requiring a clear understanding of the role of
several critical parameters from theoretical and practical standpoints. While
theoretical models are valuable from a perspective of mechanistic
comprehension, it is necessary to integrate different kinds of data knowledge
and experience-based expertise to provide a realistic basis for corrosion
prediction. A hierarchical predictive model has been developed to integrate
the effects and interactions of several critical parameters enrooted to
determining system corrosivity. The model has been implemented as a
Widows-based computer program and incorporates a framework that
facilitates further refinement.
The authors would like to recognize and thank the contributions of
researchers and pioneers in corrosion modeling, whose efforts, available in
the public domain, provided the firmament on which the model presented in
this paper has been built. The authors also acknowledge the contributions of
numerous workers within the authors' organization whose untiring efforts
have contributed significantly to the development of the corrosivity model.
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Application and Interpretation, R. Baboian (Eds.), ASTM, 1995, pp. 350358.
- John S. Smart ill, Wet ability: A major factor in oil and gas system
corrosion, Corrosion/93, Paper No. 70, New Orleans, Louisiana, 1993.
- S. Olsen, Corrosion under dewing conditions, Corrosion/91, Paper No.
472, Cincinnati DOhio, USA, 1991.
- J. Old-field and B. Todd, Corrosion considerations in selecting metals for
flash chambers, Desalination 31, 1979, pp. 365-383.
- E. Dayalan et al., Influences of flow parameters on CO2 corrosion
behavior of carbon steels, Corrosion/93, Paper No. 72, New Orleans,
Louisiana, USA, 1993.
- Dugstad et al., Influence of alloying elements upon the CO2 corrosion
rate of low alloy carbon steels, Corrosion/91, Paper No. 473, Cincinnati,
Ohio, USA, 1991.
- M. Kimura et al., Effects of alloying elements on corrosion resistance of
high strength line pipe steel in wet CO2 environment, Corrosion/94,
Paper No. 18, 1994
- R.D. Kane et al., Prediction and assessment of corrosivity for use of
steels in multi-phase CO2/H2S environments, Multi-Client Proposal, CLI
International, Inc., Oct. 1995.
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
TESTING OF DRILLING FLUIDS FORMULATED FROM
TABUK FORMATION CLAYS
M.N.J. Al-Awad1, A.S. Dahab2 and M.E. El-Dahshan3
1
Petroleum Engineering Department, College of Engineering, King Saud University
P.O. Box 800, Riyadh 11421, Saudi Arabia
2
3
Petroleum Engineering Department, College of Engineering,
Cairo University, Giza, Egypt
Chemical Engineering Department, College of Engineering, King Saud University,
P.O. Box 800, Riyadh 11421, Saudi Arabia
ABSTRACT
There are huge drilling operations in most areas of the Kingdom of Saudi Arabia nowadays
searching for potential hydrocarbon reservoirs. The increasing expense of importing drilling fluid
components for oil and gas wells, drilling, especially clays, necessitates the testing of some local clays
which are less expensive and are abundantly available. Therefore, representative Tabuk formation clay
samples from the northern and central provinces of Saudi Arabia have been studied. The clays were
mineralogically investigated using x-ray diffraction and scanning electron microscopy. The
physicochemical properties of clay suspensions with and without common drilling fluid additives were
measured. These properties include rheology, filtration loss, density and pH. The thermal stability of
clay minerals and suspensions was studied. Since sample 2 of the Tabuk clay produced better results,
its thermal stability was further studied by autoclaving its suspensions at temperatures between 25 and
121oC after the addition of high-viscosity carboxy methyl cellulose (HV-CMC) and XC-polymer. A
concentration of 10% by weight of sample 2 of the Tabuk formation clay plus 1% NaOH and 0.5 XCpolymer produced the best rheological, filtration and thermal stability.
Corrosion in drilling operations is strongly related to the composition of drilling fluids. A
concentration of 10% by weight of sample 2 of the Tabuk formation clay plus 0.5% NaOH and 0.5
XC-polymer developed minimum corrosion rates as revealed by corrosivity tests (static and dynamic)
and surface morphology investigations
Key Words: Corrosivity, drilling fluids, thermal stability, rheology, Tabuk formation, clays
INTRODUCTION
Clays in Saudi Arabia are represented by numerous commercial stocks [1-2]. The
mineralogical, chemical and mechanical analysis of these clays led to their use in different
industrial applications. The use of these national clays in the drilling of oil, gas or water
wells in the Kingdom as well in other Arabian Gulf countries will save millions of dollars that
otherwise would be spent on purchasing these materials. In the Khurays and Qasim regions,
many kinds of clays are found and can be studied [3]. Other unexplored areas for commercial
utilization of clay and shale are abundant in the northern region of the Kingdom, especially in
111
Oil Field Corrosion
the Tabuk formation. The clay deposits of AsSarat in the southern region of the Kingdom
represent another possible source of clay minerals [4]. The geological studies of the Ummer
Radhuma and Dammam formations have indicated the presence of numerous types of clay
minerals, especially palygorskite, in these formations [5,6,7]. In Saudi Arabia, the utilization
of clays, especially in the drilling of oil and water wells, has been very limited. The intense
exploration for oil and water in the Kingdom as well as other Gulf countries, necessitates the
use of local raw materials in this important industry [8].
EXPERIMENTAL PROCEDURE
This study specifically aims to achieve the following objectives:
• Geological sampling of good clay occurrence in the northern province of Saudi
Arabia,
• (Characterization of clay and non-clay minerals by mineralogical and mechanical
analysis of clay samples, and
• Testing of the application of Tabuk formation clays in drilling fluids by:
a) Measuring the rheological behavior of clay suspensions, and correlating the
mineralogical analysis and rheological properties,
b) Studying the stability of formulated drilling fluids at high temperatures and
in chemically complex environments,
c) Activating those clays shown to be of lower filtration, yield or rheological
properties, and
d) Study of the corrosivity of the formulated drilling fluids.
Figure 1 represents a flow diagram of tests and analysis conducted in this study.
METHODS OF CHARACTERIZATION
Oriented clay films were prepared on glass slides after having been dried at 40-50oC in
an oven. The slides then were analyzed with a Philips x-ray diffractometer with a graphite
monochromator and a Cu-target x-ray tube operated at 35 kV and 15 mA. The precise
determination of the shape of the clay particles was made by the electron microscope. The
samples were mounted on stubs to be coated with gold for 3 minutes at 15 on an EMSCOPE
SC-500 sputter coater, and then examined with a JEOL-35 FC scanning electron microscope.
All rheological properties, both at room and high temperatures, were investigated using the
standard devices offered by courtesy of Baroid Petroleum Services. The corrosivity of the
formulated drilling fluids was investigated using various laboratory test equipment including
a dynamic flow loop (Fig. 2), a static test reservoir, a pH meter and a low power microscope
with a camera.
112
Al-Awad et al.
Figure 1. Schematic diagram of tests and analysis
Figure 2. Schematic diagram of the closed dynamic loop
RESULTS AND DISCUSSION
Characterization
The importance of the granulometric analysis goes back to the fact that the clays which
contain more than very small amounts of non-clay minerals, especially in the size of sand and
silt, will not be suitable for drilling fluid applications. Sand can import undesirable properties
to drilling fluids and have an erosive action on pumps and drill strings. The granulometric
analysis of Tabuk formation clays showed that more than 95% of the samples had fineness
fractions of < 37 mm as shown in Table 1. The good rheology of the suspensions prepared
from Tabuk formation clays is caused in part by the fineness of the Tabuk formation clays.
The results obtained from x-ray diffractograms (Fig. 3) showed that the samples consisted
exclusively of illite and kaolinite, interpreted by the presence of the principal peaks
characterizing both clays. The non-clay minerals characterized by the diffractograms were
mainly quartz. For further investigation, the samples were examined by the scanning electron
microscope to study the texture of the clay minerals present. The electron micrographs of the
Tabuk formation clays indicated the presence of disordered illite between kaolinite particles,
as shown in Fig. 4. Dehydration curves of the Tabuk formation clays are shown in Fig. 5.
These clays lost their absorbed water at 100oC, whereas the interlayer water was lost between
100 and 400oC. The hydroxyle water was lost in the range of 400-600oC. It was very
difficult to distinguish the clay minerals from such curves because mixtures of different
minerals occur in the samples.
Table 1. Granulometric Analysis of Tabuk Formation Clays
113
Oil Field Corrosion
Sieve Size
(mm)
500
300
150
100
63
40
Fines
114
Weight (%)
Sample 1
0
0.2
0.3
0.3
1.6
2.6
95
Weight (%)
Sample 2
0
0
0
0
0.3
0.5
99.2
Al-Awad et al.
Figure 3. X-ray diffractogram of Tabuk formation clays, raw samples 1 and 2, at 25oC
(a) Sample 1
115
Oil Field Corrosion
(b) Sample 2
Figure 4. Scanning electron micrographs of Tabuk formation clays
Figure 5. Dehydration curves of Tabuk formation clays
Rheology
The rheological properties of drilling fluids play a very important role in determining the
overall success of drilling operations. In order to obtain properly functioning drilling fluids,
the rheological properties should be carefully investigated. Thus, the ability to use material in
a drilling fluid depends on the material’s flow, filtration and electrochemical characteristics.
The testing of other properties can be achieved in many ways, the most important of which is
the measurement of rheological behavior. In the case of mineral suspensions,this type of
measurement is influenced among other things by the temperature, weight percent of solids,
quantity and sign of superficial charges, clay mineralogy, and type of mixing water [9,10]. It
is practically impossible to study individually the effect of each of the previous parameters on
the rheological behavior of clay suspensions. However, the establishment of rheograms and
viscosity measurements can indicate modifications in suspension’s behavior. Figures 6 and 7
give typical examples of shear stress-shear rate relationships obtained when formulating
Tabuk formation clays in freshwater. These figures indicate the presence of relations which
accurately describe the flow characteristics of a drilling fluid over the shear rate ranges
normally encountered in the wellbore. The suspensions possessed a pesudo-plastic behavior
116
Al-Awad et al.
even with concentrations of clay solids below 10%. Table 2 gives the typical relationships of
plastic and the apparent viscosities, yield point and density of Tabuk formation clays
suspensions in freshwater. Such rheological behavior is typical for clay-water systems
formed using illite and kaolinite clays. The filtration loss was too high for it to be used as a
drilling fluid without using fluid loss control additives. In order to obtain a drilling fluid
exhibiting minimum corrosion rates, NaOH was added, and concentrations of 0.5 to 1.0% by
weight, were found adequate to develop the desired properties. 0.5 percent XC-Polymer by
weight was added to improve the fluid rheology and filtration loss. Figure 8 shows the
reheograms of the drilling fluids prepared from 10% by weight Tabuk formation clay
(sample 2) plus 1.0% NaOH with the addition of various additives; the filtration behavior of
the best additives are shown in Fig. 9. Figure 10 shows a comparison between the drilling
fluid formulated from Tabuk formation clay suspension (10% Tabuk formation clay sample 2,
0.5% XC-polymer, and 1% NaOH) with common drilling clays.
117
Oil Field Corrosion
Figure 6. Shear stress-shear rate relationships for Tabuk formation clay (sample 1)
suspension using freshwater
Figure 7.
118
Shear stress-shear rate relationships for Tabuk formation clay (sample 2)
suspension using freshwater
Al-Awad et al.
Figure 8. Shear stress-shear rate relationships for 15% Tabuk clay (sample 2) suspension
with various additives
Figure 9. API and HT-HP filteration loss of 10% by weight of Tabuk formation clay (sample
2) suspension in freshwater
Figure 10. Natural and conditioned (10% Tabuk clay + 0.5% XC-polymer + 1% NaOH) clay
119
Oil Field Corrosion
suspensions formulated from Tabuk formation clay (sample 2) compared to
common clays
Table 2. Rheological Properties of Tabuk Formation Clays
Clay
weight
(%)
10
15
20
Density
Yield Point
Apparent Viscosity
Plastic Viscosity
(ppg)
Sample Sample
1
2
8.9
8.9
9.3
9.1
9.7
9.4
(lb/100 ft2)
Sample Sample
1
2
4.5
3.0
5.0
4.0
7.5
6.3
(cp)
Sample Sample
1
2
2.94
2.94
3.74
3.74
5.27
5.27
(cp)
Sample Sample
1
2
1.1
1.1
2.14
2.14
2.15
2.15
Corrosivity
Corrosion is noted for its destructive effect on materials, and its consequences on the
economy. Corrosion has been a problem in the petroleum industry throughout its history.
Drilling fluid components play major roles in corrosion processes. To understand the
influence of drilling fluids on the rate of corrosion, certain factors should be studied, such as
their physical and chemical properties. Corrosion within a drilling well was simulated by
immersing test specimens in the drilling fluid at a fixed exposure time and temperature in
both static and dynamic conditions. The loss in weight of the test specimen served to
measure the rate of corrosion. Drilling fluids formulated from Tabuk formation clays (sample
2) were investigated for their corrosivity on three different types of casing alloys, namely:
mild steel, J-55 casing grade and K-55 casing grade. 10% by weight suspension of Tabuk
formation clay (sample 2) was chosen to study the corrosivity of Tabuk formation clays with
and without the addition of NaOH. Coupons made from the chosen alloys were made and
appropriately placed inside either the static cell or the dynamic loop, depending on the type of
test to be performed. Then the formulated drilling fluid was added to the test apparatus. Both
tests were run for two week periods at 58oC. After termination of the test, the samples were
cleaned with inert fluid, and their weights were recorded and their surfaces were carefully
examined using a low-power microscope. The following formula was used to calculate the
corrosion rate [11].
⎛ W∗L ⎞
⎟
mpy = 22273 ∗ ⎜
⎝ D∗A∗T ⎠
where
mpy
W
D
A
T
120
= penetration rate in mm-inch/year,
= weight loss in grams,
= density in gm/cm3,
= area in contact with the test fluid in in2, and
= exposure time in hours.
(1)
Al-Awad et al.
where mpy is the penetration (corrosion) rate in mm/year. As shown in Table 3, mild steel
had corrosion rates under dynamic conditions that were two-fold greater than in static
conditions. This can be attributed to the effect of shear forces on the surface of the metal
caused by the
121
Oil Field Corrosion
Figure 11. Surface morphology of steel coupons tested under static conditions for 2 weeks
using a drilling fluid formulated from 10% Tabuk formation clay (sample 2) +
0.5% XC-polymer with and without the addition of NaOH
122
Al-Awad et al.
Figure 12. Surface morphology of steel coupons tested under dynamic conditions for 2 weeks
using a drilling fluid formulated from 10% Tabuk formation clay (sample 2) +
0.5% XC-polymer with and without the addition of NaOH
123
Oil Field Corrosion
flow of the test fluid. Mild steel was found to have the lowest resistance to corrosion by
activated Tabuk formation clays, while K-55 had the highest resistance to corrosion under the
same test conditions. In the static tests, all three samples corroded at the same rate. The
corrosion rates were significantly reduced when NaOH was added to the test fluids. This was
due to the ability of NaOH to change the test fluid’s environment from being acidic (pH = 5)
to being alkaline (pH = 11), as shown in Table 3. The surface morphology of the tested steels
before and after the addition of NaOH for both static and dynamic tests is shown in Figs. 11
and 12. It is clear from the figures that the metal surfaces were more severely corroded
before the addition of NaOH.
Table 3. Corrosion Tests Results of Tabuk Formation Clays
Test
Coupon
Static Test
(mpy)
Dynamic Test
(mpy)
A
B
C
D
Mild Steel
0.26
0.10
1.7
0.255
K-55
0.25
0.10
1.0
0.212
J-55
0.23
0.10
0.8
0.189
A: Static test (10% sample 2 without additives, pH = 5)
B: Static test (10% sample 2 + 0.5% XC-polymer + 0.5% NaOH, pH = 11)
C: Dynamic test (10% sample 2 without additives, pH = 5)
D: Dynamic (10% Sample 2 + 0.5% XC-polymer + 0.5% NaOH, pH = 11)
CONCLUSIONS
1. The tested clays made from Tabuk formation clays contain more than 95% fines by
weight (< 40 mm). They are exclusively composed of the clay minerals illite and
kaolinite together with small amounts of interstratified minerals and quartz.
2. The rheological properties of the Tabuk formation clay suspensions indicate that they lie
in the range of typical native clays. After the addition of common drilling fluid additives,
the suspensions had more or less the same flow behavior as commercial clays. Therefore,
it is recommended that this clay be used in formulating low-weight drilling fluids.
3. The formulated drilling fluids made from Tabuk formation clay (sample 2) are thermally
stable up to the tested temperatures of 25-121oC.
4. Corrosion tests on the drilling fluids formulated from Tabuk formation clays (sample 2)
indicated that the fluids possessed minimal corrosivity.
5. The application of Tabuk formation clays in oil well drilling is recommended especially
after performing economical feasibility studies.
REFERENCES
1. A.S. Dahab, Evaluation of some Saudi shales for use in drilling fluids, First Conference
on Indigenous Raw Materials and Their Industrial Utilization in the Gulf Region, Kuwait,
November 1-4, 1986.
2. M.M. Aba-Husayn and A.H. Sayegh, Mineralogy of Al-Hasa desert soils, Saudi Arabia,
124
Al-Awad et al.
Clay and Clay Minerals, 1977, pp. 138-147.
3. A.S. Mashhady, M. Reda, M.J. Wilson and R.C. MacKenzie, Clay and silt mineralogy of
Saudi soils from Qasim, Saudi Arabia, International Soil Soc. 31, 1980, pp. 101-115.
4. W.C. Overstreet, D.B. Stoeser, E.F. Overstreet and G.H. Goundarzi, Tertiary Laterite of
the As Sarat Mountain, Asir Province, Saudi Arabia., Bulletin No. 21, DGMR, Geddah,
1977, p. 30.
5. N. Guven and L.L. Carney, The hydrothermal transformation of Sepiolite to stevensite
and the effect of added chlorides and hydroxides, Clay and Clay Minerals 27, 1979, pp.
253-260.
6. S.S. Sayari and J.G. Zotl, Quaternary Period in Saudi Arabia, Vo.1, Springervertag, New
York, 1978.
7. S.Y. Lee, J.B. Dixon and M.M. Aba Husayen, Mineralogy of Saudi Arabia Soil, Eastern
Region, Soil. Society of America Journal 47, 1983, pp. 321-326.
8. M.N.J. Al-Awad, Rheology, thermal Stability and Corrosivity of Saudi Clays from the
Central province, M.Sc. thesis, Department of Petroleum Engineering, College of
Engineering, King Saud University, Riyadh, Saudi Arabia, 1990.
9. A.S. Dahab, Y. Champetier and J.F. Delon, Quelques Argiles Egyptiennes dans Le
Domaine des Boues de Forage Petrolieurs, Min. Ind. J., France, April 1985, pp. 183-187.
10. M.G. Fontana and D.G. Norbert, Corrosion Engineering, Updated Textbook Edition,
McGraw Hill, 1978.
125
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
PREVENTING SULFATE SCALE DEPOSITION
IN OIL PRODUCTION FACILITIES
C.J. Hinrichsen1, M.J. McKinzie2
S. He3, J. Oddo4, A.J. Gerbino5, A.T. Kan3 and M.B. Tomson3
1
Texaco E&P Technology Department
P.O. Box 770070, Houston, TX 77215-0070, USA
2
Texaco E & P
P.O. Box 8992, Corpus Christi, TX 78468, USA
3
Rice University, Department of Environmental Science and Engineering
P.O. Box 1892, Houston, TX 77251, USA
4
Water Research Institute, Inc.
P.O. Box 980636, Houston, TX 77098, USA
5
OLI Systems, Inc.
American Enterprise Park
108 American Rd., Morris Plains, NJ 07950, USA
ABSTRACT
Scale deposits are a concern for oil producers not only throughout the United States and United
Kingdom (North Sea) but also within parts of the Middle East, primarily in Egypt and the Arabian
Gulf region. In addition to plugging perforations and reducing production volumes, certain types of
sulfate scale deposits are becoming an environmental issue.
Considering the varied problems often associated with the chemical and mechanical removal and
disposal of sulfate scale deposits, the best, most cost-effective method for dealing with sulfate scale is
simply to try to find a way to prevent the scale deposits from occurring in the first place. In order to
achieve this goal, an effective sulfate scale inhibition program must be developed.
As part of its efforts in this area, Texaco has actively participated in a research consortium which is
developing new technology specifically aimed at preventing the deposition of barium sulfate scales in oil
field production facilities. Barium sulfate is a particularly troublesome form of scale due, in part, to its
extreme insolubility. Removing barium sulfate with chemical scale dissolvers is often a very difficult
process. This paper will present the results of a recent case history in which Texaco is currently applying
scale control technology with notable success. Specifically, the process of screening and selecting a scale
inhibitor suited to the temperature, mineralogy, and brine chemistry of the producing formation will be
described. The design and implementation of the scale squeeze treatment will then be discussed. This
technology can also be applied in preventing other sulfate-type scales and for inhibiting calcium carbonate
scale as well.
Key Words: Scale inhibition, barium sulfate, scale inhibitors, squeeze treatment, scaling
tendency, scale prediction, laboratory testing, NORM
127
Oil Field Corrosion
INTRODUCTION
The A.E. Guerra No. 43 well, located in South Texas, produces 8.5 MMSCFPD, 50
BOPD, and 120 BWPD. The reservoir is relatively deep (approximately 14,000 ft). The
bottom hole temperature is typically about 340°F, and the reservoir pressure is approximately
6997 psi. The reservoir formation is composed of calcite (20 - 40%), quartz (25 - 30%),
feldspar (20%), and clay minerals (10 - 35%). Among the clay minerals, illite is typically
80% and chlorite is around 20%. The porosity is about 18%. The permeability varies from
0.01 to 0.05 mDarcy, and it is a fracture-stimulated reservoir. The content of carbon dioxide
in the gas phase is about 0.5%.
Two months after production commenced, scale buildup was discovered when a tubing
caliper survey was run. Shortly thereafter, gas production dropped from 10 MMSCFPD to 5
MMSCFPD. Two types of water-formed scale deposits, barium sulfate and calcium
carbonate, were encountered in the production system (mainly in the downhole tubing). Hard
deposits of barium sulfate scale pose a severe operational problem to production operations
since they cannot be easily removed once deposited and have occasionally been associated
with naturally occurring radioactive materials (NORM) [1]. In this case, fortunately, the
barium sulfate scale was not contaminated with NORM.
Before the squeeze treatment, discussed in this paper, barium sulfate and calcium
carbonate scales had to be periodically removed by chipping or broaching with wire-line
tools. This operation is expensive, time-consuming, and risky considering the possibility of
losing the tools in the well. The annual cost for the frequent mechanical removal of the scale
deposits was estimated to have been US $87,000.
Under the sponsorship of the Gas Research Institute (GRI), the Brine Chemistry
Consortium at Rice University and Water Research Institute conducted an integrated study of
the scaling condition in the subject well and designed a treatment plan to control barium
sulfate scaling. This paper summarizes the laboratory design and field implementation of a
scale inhibitor squeeze treatment. First, a brine analysis and scaling tendency are described.
Next, laboratory tests of scale inhibitors are presented. Finally, the field squeeze treatment is
detailed and the results are shown.
BRINE CHEMISTRY AND SCALING TENDENCY
Brine samples were taken at the wellhead and analyzed in the laboratory. The chemical
composition of the brine and other information concerning the well are listed in Table 1. The
chemical analysis was performed at room temperature and pressure. The concentration of
each cation was analyzed using inductively coupled plasma spectrometry.The chloride
concentration and bicarbonate alkalinity were determined by titration. The concentration of
sulfate was determined by a turbidimetric method and found to be very small (< 5 mg/1).
The scaling tendency of the brine, expressed as a saturation index (SI), was calculated
under various production conditions using the EQPITZER program [2]. EQPITZER is a
computer program for calculating the SI of brines with respect to common water-formed scale
deposits such as calcite, gypsum, anhydrite, celestite, and barite. The SI for a sparingly
soluble salt, M+A-, is defined as the logarithm of the ratio of the ionic activity product divided
by the thermodynamic solubility product:
SI = log[(aM+ aA-)/Ksp]
(1)
128
Hinrichsen et al.
An SI > 0 indicates a supersaturated solution and, hence, a probability for scale deposition.
No scale is anticipated, however, when the calculated SI < 0 (i.e., the brine is
undersaturated).
Table 1. Chemical Composition of the A.E. Guerra Brine
Species
+
Na
Mg2+
Ca 2+
Sr2+
Ba2+
Fe (total)
ClSO42Alkalinity (HCO3-)
Ionic Strength (M)
pH (meas.) at surface
CO2(g) in the gas phase
Concentration
(mg/1)
19871.8
54.0
6500.0
700.0
550.0
12.0
43000.0
<5.0
281.0
(mmol/l)
864.4
2.2
162.2
8.0
4.0
0.2
1228.6
<0.052
4.6
1.42
7.10
0.5%
The calculated saturation indices are presented in Table 2. It can be concluded that a
serious calcium carbonate scaling potential will develop (the SI is greater than 1.0) as the
reservoir pressure drops. This is verified by the occurrence of calcium carbonate scale in the
tubing near the perforations where there is a significant pressure reduction.
Table 2. Calculated Saturation Index (SI) for the A.E. Guerra Brine
with Respect to Common Water-Formed Scale Deposits
Parameters
Temperature (°F)
Pressure (psi)
pH (calc)
SI (calcite)
SI (barite)
SI (celestite)
SI (anhydrite)
Downhole
236
6997
5.34
-0.27
-0.23
-1.60
-2.05
Wellhead
236
397
6.71
1.21
0.01
-1.37
-1.75
Surface
236
14.7
7.94
2.24
0.03
-1.35
-1.72
77
14.7
7.91
1.94
1.17
-1.70
-2.70
In the case of barium sulfate, however, a decrease in pressure to atmospheric conditions
is expected to only shift the SI up to 0.03 (near equilibrium). In fact, barium sulfate scale was
noted within downhole sections of the tubing string despite the fact that the calculated SI
suggested near equilibrium (SI = 0.01) conditions would prevail either downhole or at the
wellhead. The calculated SI for barium sulfate predicts that a potential for barite scale will
129
Oil Field Corrosion
only occur when the temperature is reduced to 77°F (SI = 1.17). The discrepancy between
the calculated SI and field observations could be due to the fact that the original brine
analysis may have incorrectly reflected the true downhole concentrations of barium and
sulfate.
INHIBITOR EVALUATION AND SQUEEZE SIMULATION
In view of the declining gas production rate and the potential for further scale deposition,
it was decided to treat the subject gas well by a scale squeeze treatment. Given the US $4
MM cost to drill and complete the well and the additional US $1.5 MM cost for the fracture
operation, a carefully planned series of laboratory tests was conducted to ensure not only that
the scale squeeze treatment would be effective but also that the chemical treatment would not
contribute to any formation damage.
Several commercially available scale inhibitors were tested in the laboratory for their
efficiency at preventing barium sulfate scale deposition. Both static and dynamic tests were
conducted.
The static testing of inhibitor efficiency was based on the measurement of the nucleation
induction period for barium sulfate in the presence of inhibitors [3], specifically, the relative
prolongation of the nucleation induction period in the presence of various scale inhibitors,
each at the same concentration. The static test results were then used to rank the scale
inhibitors according to their efficiency in delaying the nucleation process.
The dynamic testing of scale-inhibitor effectiveness was performed in the laboratory
using a high-temperature and high-pressure flow-through apparatus which is designed to
simulate a production system [4]. The dynamic test results are useful in ranking the scale
inhibitors based upon the minimum effective dose for each inhibitor (The minimum effective
dose is the minimum concentration required to prevent any scale deposition).
Several common commercial scale inhibitors were tested. These inhibitors include 1hydroxyethylidene-1,1-diphosphonic acid (HEDP), nitrilotrimethylene phosphonic acid
(NTMP), hexamethylene diamine tetramethylene phosphonic acid (HDTMP), diethylene
triamine pentamethylene phosphonic acid (DTPMP), bis-hexamethylene triamine
tetramethylene phosphonic acid (BHTMP), polyacrylates (PAA, molecular weight from 1000
to 7000), phosphinopolycarboxylates (PPPC, molecular weight from 1900 to 3800), and
sulfonated polyacrylic acid (SPA, molecular weight of 3500). Based on the static and
dynamic test data, BHTMP was found to be effective and superior to other scale inhibitors,
and was recommended for squeeze-treatment use.
Once the scale inhibitor testing was complete, a simulation of the squeeze process was
conducted in the laboratory in order to obtain information concerning the retention and
release of the BHTMP inhibitor onto and off of the A.E. Guerra core material. First, the
BHTMP inhibitor solution was pumped into a column packed with synthetic materials having
a mineralogical composition similar to the formation rock. The column was then shut in for
two days. Next, the column was turned around and a synthetic brine was pumped through the
packed column from the opposite direction. These two steps are designed to simulate the
inhibitor injection and subsequent return flow. The concentration of BHTMP in the return
flow was continuously monitored for over 60 pore volumes. The flow rate was 10 ml/min.
The inhibitor return data is presented in Fig. 1.
130
Hinrichsen et al.
a
100000
0.5%
5%
10000
BHTMP (Acid Active, mg/l)
1000
100
10
1
0
20
40
60
80
100
Pore Volumes
The return concentration of the inhibitor
b
30
20
Inhibitor Return (%)
10
0.5 %
5%
0
0
20
40
60
80
100
Pore Volumes
The inhibitor return percentage
Figure 1. The squeeze simulation result of BHTMP in synthetic core materials.
The open circles represents data from squeezing 0.5% inhibitor acid,
while the solid diamonds represent data from squeezing 5.0% inhibitor
acid
131
Oil Field Corrosion
Two sets of simulations were performed using two different injection concentrations (5%
and 0.5% BHTMP). The inhibitor concentration remained over 1 mg/l for over 100 pore
volumes (Fig. 1a) and the percentage of inhibitor returned was less than 30% (Fig. 1b).
INHIBITOR SQUEEZE DESIGN
Based on the information derived from both the laboratory evaluation of commercially
available scale inhibitors and from the column simulation of an inhibitor squeeze under
simulated field conditions, a scale inhibitor squeeze treatment was designed for the A.E.
Guerra No. 43 gas well. The squeeze process included five phases, which are listed in Table
3.
Table 3. Design of a Scale Inhibitor Squeeze Treatment
Phase
*
Process
Volume
Additive
(bbl)
Composition
Concentration
1
(Preflush)*
25
HCl
0.5%
2
Pill
270
BHTMP (110 gals)
3537 mg/l
3
Overflush
250
Filtered field brine
4
Shut-in (48 hours)
5
Production
Since an acid treatment had been used two days prior to the actual squeeze treatment,
in order to remove any existing scale, no acid preflush was used for this squeeze.
Normally, however, an acid preflush would be recommended.
The preflush solution (i.e., 0.5% hydrochloric acid with a corrosion inhibitor) was used
to clean the production tubing by removing calcium carbonate scale deposits. The squeeze
pill consists of the scale inhibitor (i.e., bis-hexamethylene triamine tetramethylene
phosphonic acid, BHTMP in acid form) dissolved in filtered produced water. An overflush
was used to push the inhibitor pill farther into the reservoir formation. A shut-in period of
two days was necessary to allow adsorption of the inhibitor onto the formation rock through a
reaction of the inhibitor acid with the formation material.
INHIBITOR SQUEEZE LIFETIME AND ECONOMIC IMPACT
The inhibitor squeeze was performed on December 21, 1993. After the two-day shut-in,
the well was returned to production. Long-term monitoring of the inhibitor return has been
performed since the squeeze.
The concentration of the inhibitor returned as a function of the cumulative volume of
brine produced since the squeeze is presented in Fig. 2. The inhibitor concentration remained
above 1 mg/l for the initial 20,000 bbl of brine produced (approximately 167 days) and
maintained around 0.5 mg/l for over 50,000 bbl of brine produced (about 417 days) (Fig. 2a).
After 18 months, the amount of inhibitor returned was less than 20% of the total amount of
the inhibitor squeezed (Fig. 2b).
132
Hinrichsen et al.
a
1000
100
10
1
BHTMP (Acid Active, mg/l)
.1
0
10000
20000
30000
40000
50000
Cumulative Volume of Brine (bbl)
The concentration of BHTMP in produced brines.
b
20
15
10
5
Inhibitor Return (%)
0
0
10000
20000
30000
40000
50000
Cumulative Volume of Brine (bbl)
The percentage of BHTMP in produced brines.
Figure 2. The return of inhibitor BHTMP as a function of brine flow back in the
A.E. Guerra well after the squeeze
133
Oil Field Corrosion
After the squeeze, the Guerra well produced gas and oil with few problems due to scale
formation for about 18 months. During this period, the well was periodically tested for
downhole scale. Recently, the gas production fell off due to a blockage caused by scale
formation. Scale samples were recovered and found to be mostly barium sulfate. A new
squeeze treatment is planned. Since a light deposit of calcium carbonate was found in the end
of the production tubing, the next squeeze treatment will include a mixed inhibitor treatment
for preventing both calcium carbonate and barium sulfate scale deposition.
The cost savings for the chemical squeeze are estimated to be over US $80,000/year for
this well alone. Furthermore, 24 days of added production per year has been realized by not
having to shut in the well for scale removal. This has produced an additional US
$150,000/year in cash flow.
CONCLUSIONS
The problem of barium sulfate scaling in downhole tubing in the subject gas well was
eliminated by a scale inhibitor squeeze using BHTMP. The squeeze has lasted for 18 months.
Cost savings are estimated to be more than US $80,000/year. The following conclusions can
be drawn from this case study:
1. BHTMP is an efficient inhibitor of barium sulfate scaling, especially in high calcium
brines.
2. Inhibitor squeeze treatment is an effective and economic method for controlling scale
deposition in downhole conditions in an oil and gas production system.
3. In the inhibitor squeeze treatment, blends of inhibitors may be needed instead of a
single inhibitor to control the formation of mixed mineral scales, such as barium
sulfate and calcium carbonate. Such combinations are presently being tested.
ACKNOWLEDGMENTS
This work was supported, in part, by the Gas Research Institute and by a consortium of
companies including Texaco, Inc.; Conoco, Inc.; Champion Technologies, Inc.; FMC
Corporation; Product Additives Division; and Zapata, Inc. The authors would like to thank
Texaco, Inc. for permission to publish this paper. The authors would also like to express their
deep appreciation to Saudi Arabian Texaco for supporting this presentation at the Second
Arabian Corrosion Conference.
REFERENCES
1. J.E. Oddo and M.B. Tomson, Algorithms can predict; inhibitors can control NORM
scale. Oil and Gas Journal, 1994, pp. 33-37.
2. S.L. He and J.W. Morse, Prediction of halite, gypsum and anhydrite solubility in
natural brines under subsurface conditions, Computers and Geosciences 19, l, 1993,
pp. 1-22.
3. S.L. He, J.E. Oddo and M.B. Tomson, The inhibition of gypsum and barite nucleation
in NaCl brines at temperatures from 25 to 90°C, Appl. Geochemistry 9, 1994, pp. 561567.
134
Hinrichsen et al.
4. J.E. Oddo and M.B. Tomson, Elevated temperature-pressure flow simulator, U.S.
Patent No. 5,370,799.
135
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
CONCERNS OVER THE SELECTION OF BIOCIDES FOR OIL FIELDS AND
POWER PLANTS: A LABORATORY CORROSION ASSESSMENT
J. Alhajji and M. Valliappan
Kuwait University
P.O. Box 5969, Safat 13060, Kuwait
ABSTRACT
Control of microbially induced corrosion (MIC) involves the application of biocides that
eliminate the deleterious microorganisms, resulting in a reduced corrosion rate of the material of
interest. However, the electrochemical properties of such chemicals in promoting or inhibiting
corrosion processes are not well understood. The present investigation is focused on the evaluation of
biocides commonly used in the oil fields and power sectors, from an electrochemical point of view.
This study examines the use of oxidizing and nonoxidizing biocides in sulfide contaminated saline
solutions. Gluteraldehyde (1,5-pentanedial) was chosen as a nonoxidizing biocide and ammonium
chloride, which is the principal chemical used in the in-situ generation of chloramine, as the oxidizing
biocide. A chemical environment of sulfide in seawater in a deaerated state was set up to simulate the
conditions occurring during biological processes. The biocides investigated were found to alter the
corrosion tendencies of mild steel and pitting tendency of 316L stainless steel in this environment.
Key Words: Mild steel, stainless steel, glutaraldehyde (1,5-pentanedial), ammonium
chloride, sulfide, biocides, corrosion, microbially induced corrosion,
MIC
INTRODUCTION
Extensive available literature [1-4] relating to microbially induced corrosion (MIC) has
indicated that MIC certainly remains as an incessant problem for the corrosion world.
Moreover, MIC is regarded as an intolerable corrosion problem, and multi-disciplinary effort
have been expended to undermine its detrimental effects and to understand the phenomenon
and its mechanisms [5-9]. All the engineering materials in general use, notably carbon steel
and stainless steels in oil fields and the power sectors are susceptible to some form of
microbial corrosion which usually arises from the activities of a wide range of
microorganisms and their usually oxidizing metabolic products.
The biological influences can be divided into three general categories [10]: production of
differential aeration or chemical concentration cells, production of organic and inorganic
acids as metabolic by products, and production of sulfides under oxygen-free (i.e., anaerobic)
conditions. Several studies have focused on MIC of carbon steels [3] and steel alloys
containing 2-3% Mo [6-8]. Control of MIC usually involves the utilization of chemical
biocides. These chemicals are sometimes misunderstood and misused in processes involving
aqueous environments due to a lack of understanding of the nature of the problem. Very little
effort [11] has been made to investigate the effects of biocides that are used in oil fields and
135
Oil Field Corrosion
power generation systems on the electrochemical behavior of metallic materials. Most of the
research [12-13] has been concerned with the biocidal properties of the chemicals that are
used to combat biologically induced corrosion problems and not the electrochemical
properties of those chemicals.
Biocides are categorized as either oxidizing or nonoxidizing toxicants. The selection and
application of suitable biocide treatment depends on the broad spectrum activity, pH,
economics, compatibility with the other chemicals used for treatment, and most importantly,
suitability with the materials of construction from a corrosion point of view. Glutaraldehyde
(1,5-pentanedial), a nonoxidizing biocide, is extensively used in oil fields. Chlorination is
commonly employed in the treatment of freshwater systems. Chlorine is the most widely
used industrial oxidizing biocide. Ammonium chloride, an oxidizing chemical, is commonly
used in the in-situ generation of chloramine as an oxidizing biocide.
The present investigation is focused on investigating the electrochemical nature of
glutaraldehyde, as a nonoxidizing chemical, on the corrosion effects of mild steel and 316L
stainless steel in deaerated synthetic seawater media containing sulfide to simulate the effects
of biogenic sulfides generated by sulfate reducing bacteria (SRB) the most documented
deleterious organism in MIC, under anaerobic conditions. Also, ammonium chloride, an
oxidizing chemical, is investigated under similar conditions simulating biogenic sulfides
generated by SRB.
EXPERIMENTAL PROCEDURE
Carbon steel (UNS G10200) and 316L stainless steel (UNS 316003) of exposed areas of
1 cm2 were used for this study. The working electrodes were polished with emery papers to a
600 grit finish. The polished specimens were successively rinsed with analar grade acetone
and double distilled water, and then air dried. Experiments for corrosion measurements were
conducted utilizing standard seawater. This was prepared with distilled water and standard
seawater salt. Standard seawater salt (Marinemix + Bio-Elements from Wiegandt GMBH &
Co., F.R. Germany) was used to reduce the variability of effects resulting from conducting
measurements using natural seawater. Experiments were also conducted in sulfide polluted
seawater. The sulfide was introduced using research-grade sodium sulfide (Na2S). The level
of sulfide in the seawater was checked by the iodimetric method of analysis. The synthetic
seawater solutions were deaerated using purified nitrogen gas. Sulfide was introduced to
these solutions as sodium sulfide in the concentrations of 1 and 10 ppm. Glutaraldehyde (1,5pentanedial) was selected as the nonoxidizing-type of biocide in concentrations of 10, 50 and
100 ppm and, ammonium chloride was selected as the oxidizing agent in the concentration
range of 1, 5 and 10 ppm. Electrochemical measurements employing linear polarization,
Tafel extrapolation, potentiodynamic polarization and AC impedance techniques were carried
out in the three-electrode cell assembly (EG&G). The reference electrode used was a
saturated calomel electrode (SCE), located in a glass tube fitted with vycor frit and the
electrochemical circuit was completed with an encapsulating cylindrical platinum counter
electrode.
Polarization resistance (Rp) values were obtained by using linear polarization applying
Ecorr + 0 mV across the working electrode’s surface. Tafel polarization studies were
performed by imposing Ecorr + 250 mV with respect to the open circuit potential (OCP). The
136
Alhajji and Valliappan
anodic polarization experiments were run at a scan rate of 0.5 mV/sec using a potentiostat
5
(EG&G model 273A). AC impedance data were generated as a function of frequency (10
Hz-10 mHz) by imposing a sinusoidal voltage signal of 5 mV across the interface. The
complex impedance data were acquired using a frequency response analyzer (Schlumberger
model SI1255) controlled by a computer. All the electrochemical measurements were made
o
under stagnant conditions at room temperature (i.e., 19 +
– 1 C).
RESULTS AND DISCUSSION
The Tafel polarization curves for mild steel in seawater and a sulfide-polluted system
under deaerated conditions are presented in Fig. 1a. Increasing the sulfide concentration
resulted in a shift in potential in the active direction. Also, increasing the sulfide level
resulted in a significant change in the values of ba and bc. At the same time, this resulted in
an increase in the corrosion rate for a higher concentration of sulfide, i.e., 10 ppm. Clearly
the presence of sulfide has a significant influence on the rate of the cathodic and anodic
reactions. From a comparison of the values of ba and bc at the two sulfide concentrations of
interest, it can be seen that increasing the sulfide concentration resulted in bc tending towards
increasing values which indicates a diffusion-limiting reaction occurring at the surface. This
is clearly a result of exposure to higher concentrations of sulfide ions which resulted in the
development of an oxidized layer acting as a diffusion barrier. Thus, it is possible that the
corrosion reaction was under mixed control. The limiting diffusion current was observed on
the cathodic branch of the polarization curve, as can be seen in Fig. 1a. This is consistent
with the conclusions of Iofa [14] that sulfide participates directly in the cathodic reactions and
is simply a catalyst which speeds up the discharge of hydrogen ions. Various mechanisms
have been proposed for the intensification of the corrosion rate due to the presence of
hydrogen sulfide [14-18]. An example of one such mechanism is that proposed by Panasenko
[16], i.e.,
Fe + HS
-
→
Fe(HS-)ads
Fe(HS-)ads
→
Fe(HS) + 2e
→
Fe
Fe(HS)
+
+
2+
+ HS-
(1)
(2)
(3)
It is clear that sulfide acts as a catalyst promoting ferrous ion generation. This will lead
to an increase in the concentrations of ferrous ions near the interface. The adsorption of HSions will produce a negative charge on the surface which will accelerate the hydrogen
discharge reaction through another simultaneous cathodic reaction [15]:
-
2-
HS + e
→
Hads + S
(4)
Hads + Hads
→
H2
(5)
137
Oil Field Corrosion
-500
-500
0 ppm Glutaraldehyde
10 ppm Glutaraldehyde
50 ppm Glutaraldehyde
100 ppm Glutaraldehyde
Seawater
-600
1 ppm Sulfide
SCE
10 ppm Sulfide
-700
Potential, mV
Potential, mV
SCE
-600
-800
-900
-1000
-700
-800
-900
-1000
-1100
-1100
-9
-8
-7
-6
-5
-4
-3
-2
-10
-9
-8
2
-7
-6
-5
-4
-3
-2
2
log (i[A/cm ])
log (i[A/cm ])
(a)
(b)
-500
0 ppm NH 4Cl
1 ppm NH 4Cl
5 ppm NH 4Cl
10 ppm NH 4Cl
-500
0 ppm Glutaraldehyde
10 ppm Glutaraldehyde
50 ppm Glutaraldehyde
100 ppm Glutaraldehyde
SCE
-600
Potential, mV
Potential, mV
SCE
-600
-700
-800
-900
-700
-800
-900
-1000
-1000
-1100
-1100
-10
-9
-8
-7
-6
-5
-4
-3
-10
-2
-9
-8
-7
-6
-5
-4
-3
2
2
log (i[A/cm ])
log (i[A/cm ])
(c)
(d)
-500
0 ppm NH 4Cl
1ppm NH
Cl
4
5 ppm NH
Cl
4
10 ppm NH
Cl
Potential, mV
SCE
-600
4
-700
-800
-900
-1000
-1100
-10
-9
-8
-7
-6
-5
-4
-3
-2
2
log (i[A/cm ])
(e)
Figure 1. Polarization diagrams for mild steel in seawater: (a) without biocides; (b)
with glutaraldehyde and 1 ppm sulfide; (c) with glutaraldehyde and 10
ppm sulfide; (d) with ammonium chloride and 1 ppm sulfide; and (e)
with ammonium chloride and 10 ppm sulfide
138
-2
Alhajji and Valliappan
2-
At the pH of the present solution (9.7 – 9.75) S is not stable and thus the reaction
2-
S
→
+ H2O
-
HS- + OH
(6)
takes place, by which hydroxide ions are produced, increasing the pH of the solution.
However due to the high concentration of ferrous ions near the interface, another partial
reaction will compete with the reaction in Eq. 4, that is the precipitation of the solid sulfide
according to the following reaction:
Fe
2+
-
+ HS
→
Various Iron Sulfides (Insoluble)
(7)
Thus in this case, the corrosion current measurements for steel have indicated significant
increases in the presence of low concentrations of sulfide which is due to the formation of a
corrosion product which acts as a physical barrier to the corrosion reaction. However, the
anodic polarization curves from Fig. 1a show an increase in the anodic dissolution for a
sulfide-added system.
To investigate the effect of oxidizing biocide on the mild steel-sulfide-polluted seawater
system, glutaraldehyde and ammonium chloride were introduced. The Tafel polarization
curves for mild steel in deaerated sulfide-polluted seawater with various concentrations of
either glutaraldehyde or ammonium chloride are presented in Figs. 1b-e. The parameters
obtained from the linear polarization and Tafel extrapolation methods are given in Tables 1
and 2. It can be seen from these tables that the addition of sulfide increases the corrosion rate
of mild steel, especially in higher concentrations of sulfide as reflected both in the
measurement of Rp and corrosion current, icorr.
Table 1. The Electrochemical Parameters of Mild Steel and Stainless Steel
in Unpolluted Seawater
Alloy
Mild Steel
Stainless Steel
Parameter
OCP (mV vs SCE)
Rp, LP (KΩ cm2)
icorr, (μA/cm2)
Rp,Imp (KΩ cm2)
Epit, (mV vs SCE)
Ipit (μA/cm2)
Rp, LP (KΩcm2)
Measured Value
–776
12.38
0.508
6.87
64
2.69
236.3
As the conditions in the present investigation involve is a simulated chemical
environment undergoing biological processes, the behavior of mild steel in the sulfidepolluted, oxygen-free system is similar to its behavior in an anaerobic sulphate reducing
bacteria (SRB) environment. Sulfide ions present in the oxygen-free chloride system are
aggressive to mild steel, and the resultant product is iron sulfide which has no inhibiting or
intact properties on the metal surface to prevent further dissolution. However, the addition of
139
Oil Field Corrosion
glutaraldehyde and ammonium chloride alterred the corrosion tendencies, as depicted in
Tables 1 and 2. In the lower concentration of sulfide, the addition of glutaraldehyde in the
concentrations of 10 and 100 ppm had no profound effect on the dissolution, but in the
intermediate concentration of 50 ppm, it increased the corrosion rate of mild steel. It is
interesting to note that the accelerating mode of corrosion in the presence of a low sulfide
concentration, i.e., 1 ppm and a glutaraldehyde concentration of 50 ppm, is completely
reversed in the high sulfide concentration, i.e., 10 ppm, where the minimum corrosion current
was measured, when compared with the other concentrations of glutaraldehyde.
Table 2. The Effects of the Investigated Biocides on the Electrochemical
Parameters of Mild Steel in Sulfide Polluted Seawater
Sulfide Concentration
1 ppm
Biocides
Investigated
None
10 ppm
Conc.
ppm
OCP, mV
vs SCE
-784
Rp (LP)
KΩ cm2
10.670
icorr,
μA/cm2
0.49
Rp (Imp)
KΩ cm2
6.333
OCP, mV Rp (LP),
icorr,
vs SCE KΩ cm2 μA/cm2
-794
3.012
1.29
Rp (Imp)
KΩ cm2
2.2610
Glutaraldehye
10
50
100
1
-810
-780
-799
-809
10.46
3.447
10.03
8.577
0.535
2.12
0.617
0.579
9.931
1.734
6.685
5.370
-816
-738
-816
-815
7.880
16.62
4.310
13.31
0.93
0.4
1.09
0.524
5.133
--4.383
---
Ammonium
Chloride
5
10
-806
-818
2.047
2.815
4.44
0.70
1.223
7.236
-782
-732
4.14
18.8
1.40
0.372
-----
The addition of ammonium chloride to the sulfide polluted system increased the
corrosion current and reduced the Rp values in the lower concentration of 1 ppm sulfide.
However, this scenario was not observed in the higher concentration of 10 ppm sulfide level
where decreases in the corrosion current values were noted for an increase in the ammonium
chloride concentration, except at 5 ppm. The Nyquist plots and Bode magnitude diagrams
obtained form the AC impedance measurements are presented in Figs. 2 and 3, respectively.
Because of the distinct advantage of the Bode magnitude plots over the Nyquist plots in the
low frequency range, the Rp values obtained from the Bode plots were taken into account for
the analysis and are presented in Tables 1 and 2. These measurements also indicated the
similar trends observed in the DC polarization measurements.
Some investigators [11] have shown that the analyses of results from the AC impedance
method are more appropriate than the DC polarization for the biocide added system where
potentiodynamic measurements yielded higher corrosion rates because of the change in the
surface of the system exposed to the biological environment during the polarization. It is a
general belief that the addition of biocide to the MIC system results in the elimination or the
control of the deleterious microorganisms, thereby reducing the corrosion rate of the material
of interest. The results from the previous investigations [19,28] on the glutaraldehyde-added
system were of a conflicting nature, because of the effect attributed to the release of
metabolic products during glutaraldehyde treatment which might change local pH and/or
inorganic species adjacent to the metal/solution interface, thereby interfering with the
corrosion of steel.
140
Alhajji and Valliappan
6
0 ppm Glutaraldehyde
10 ppm Glutaraldehyde
50 ppm Glutaraldehyde
100 ppm Glutaraldehyde
5
2
4
(KOhm.cm )
0 ppm Sulfide
1 ppm Sulfide
10 ppm Sulfide
3
3
imaginary
2
4
1
2
Z
Z
imaginary
2
(KOhm.cm )
5
1
0
0
1
2
3
Z
4
5
6
0
7
2
real
0
(KOhm.cm )
2
4
6
Z
8
10
12
2
real
(KOhm.cm )
(a)
(b)
6
2
(KOhm.cm )
2
8
imaginary
6
0 ppm NH 4Cl
1 ppm NH Cl
4
5 ppm NH 4Cl
10 ppm NH Cl
5
4
4
3
2
Z
4
Z
imaginary
10
(KOhm.cm )
0 ppm Glutaraldehyde
10 ppm Glutaraldehyde
50 ppm Glutaraldehyde
100 ppm Glutaraldehyde
12
1
2
0
0
0
2
4
Z
6
8
10
0
12
2
2
real
4
6
Z
(KOhm.cm )
(c)
real
8
10
12
2
14
16
(KOhm.cm )
(d)
0 ppm NH Cl
4
1 ppm NH Cl
4
5 ppm NH Cl
4
10 ppm NH 4Cl
12
10
8
6
4
Z
imaginary
2
(KOhm.cm )
14
2
0
0
2
4
Z
6
8
10
12
2
real
(KOhm.cm )
(e)
Figure 2. Nyquist plots for mild steel in seawater: (a) without biocides; (b) with
glutaraldehyde and 1 ppm sulfide; (c) with glutaraldehyde and 10
ppm sulfide; (d) with ammonium chloride and 1 ppm sulfide; and
(e) with ammonium chloride and 10 ppm sulfide
141
Oil Field Corrosion
4
0 ppm Sulfide
0 ppm Glutaraldehyde
10 ppm Glutaraldehyde
3.5
50 ppm Glutaraldehyde
100 ppm Glutaraldehyde
2
2
3
Log |z|, Ohms-Cm
Log |z|, Ohms-Cm
4
1 ppm Sulfide
10 ppm Sulfide
3.5
2.5
2
1.5
3
2.5
2
1.5
1
1
-4
-2
0
2
4
6
-4
-2
0
Log [frequency], Hz
2
4
6
Log [frequency], Hz
(a)
(b)
4.5
0 ppm Glutaraldehyde
10 ppm Glutaraldehyde
50 ppm Glutaraldehyde
100 ppm Glutaraldehyde
0 ppm NH 4Cl
4
1 ppm NH Cl
4
5 ppm NH 4Cl
2
3.5
4.5
Log |z|, Ohms-Cm
Log |z|, Ohms-Cm
2
4
3
2.5
2
1.5
3.5
10 ppm NH 4Cl
3
2.5
2
1.5
1
1
-4
-2
0
2
4
6
-4
-2
0
Log [frequency], Hz
2
4
Log [frequency], Hz
(c)
(d)
4.5
0 ppm NH Cl
4
4
1 ppm NH 4Cl
Log |z|, Ohms-Cm
2
5 ppm NH 4Cl
3.5
10 ppm NH Cl
4
3
2.5
2
1.5
1
-4
-2
0
2
4
6
Log [frequency], Hz
(e)
Figure 3. Bode plots for mild steel in seawater: (a) without biocides; (b) with
glutaraldehyde and 1 ppm sulfide; (c) with glutaraldehyde and 10 ppm
142
6
Alhajji and Valliappan
sulfide; (d) with ammonium chloride and 1 ppm sulfide; and (e) with
ammonium chloride and 10 ppm sulfide
It is believed that glutaraldehyde enters into interaction with sulfide and forms
thioglutaraldehyde that polymerizes on the metal surface. It is possible that organic
compounds are capable of entering into chemical interaction with the sulfide present, forming
on the metal surface insoluble compounds (i.e., Eq. 6) that constitute a peculiar kind of phase
barrier. It is possible that localized alteration of pH at the electrode surface (Eq. 4,) makes
this mechanism possible.
However, the addition of ammonium chloride can certainly decreases the resistance of
mild steel in any sulfide system because of its oxidizing nature. In the 1 ppm sulfide system,
the co-presence of ammonium chloride and oxidized forms of sulfide promotes the corrosion
process. Increasing the concentration of sulfide in a system containing NH4Cl, resulted a the
decrease in the corrosion current. This can be attributed to the formation of species such as
sulfate as the oxidized product of sulfide in the presence of NH4Cl. Also the increased
dissolution rate observed in the intermediate concentration of 5 ppm of ammonium chloride
compared with other concentrations call attention to the existence of a critical concentration
of oxidizing-types of chemicals, like passivators added to a neutral medium.
Stainless Steel
The anodic polarization curves for stainless steel in deaerated sulfide added to a biocide
system are presented in Figs. 4a-e. These curves yielded important parameters which are
presented in Tables 1 and 3, and are relevant to the assessment of 316L stainless steel in a
biocide added system. The pitting potential values shown in Tables 1 and 3 clearly indicate
the influence of sulfide on the passivity of stainless steel in a deaerated, biocide-free system.
It is well known that sulfide anions present in polluted seawater lead to the formation of an
oxide layer of poor protective characteristics, which facilitates the initiation of corrosion
attack [21]. Moreover, the presence of sulfide in a stainless steel-seawater system is
generally detrimental to the pitting resistance of all stainless steel grades. The biologically
generated sulfide can also modify the local chemistry of the marine environment and prevent
the repair of the passive film of the stainless steel [22].
Table 3. The Effects of the Investigated Biocides on the Electrochemical Parameters of 316L
Stainless Steel in Sulfide Polluted Seawater
Sulfide Concentration
1 ppm
10 ppm
Rp, Ω cm
Ip, μA/cm
275.7
10
50
100
1
5
10
Biocides
Investigated
None
Conc.
(ppm)
Glutaraldehye
Ammonium
Chloride
Rp, Ω cm
2.55
Epit mV
vs SCE
251
278.6
180.6
219.7
2.44
2.70
2.55
144.4
392.7
164.4
3.2
2.17
3.45
2
2
2
148.6
Ip,
μA/cm2
3.04
Epit mV
vs SCEE
142
159
95
198
206.7
238.9
169.3
3.4
2.52
3.08
223
125
193
104
172
158
226.9
172
152.1
3.02
5.06
3.13
264
196
169
143
Oil Field Corrosion
600
0 ppm Glutaraldehyde
10 ppm Glutaraldehyde
50 ppm Glutaraldehyde
100 ppm Glutaraldehyde
400
Seawater
400
SCE
1 ppm Sulfide
10 ppm Sulfide
200
Potential, mV
SCE
300
Potential, mV
100
0
-100
-200
200
0
-200
-300
-400
-400
-10
-9
-8
-7
-6
-5
-4
-10
-3
-9
-8
-7
-6
0 ppm Glutaraldehyde
10 ppm Glutaraldehyde
50 ppm Glutaraldehyde
100 ppm Glutaraldehyde
0 ppm NH Cl
4
1 ppm NH Cl
400
4
5 ppm NH 4Cl
10 ppm NH Cl
SCE
200
0
-200
4
200
0
-200
-400
-400
-10
-9
-8
-7
-6
-5
-4
-3
-10
-9
-8
2
-7
-6
-5
-4
2
log (i[A/cm ])
log (i[A/cm ])
(c)
(d)
0 ppm NH Cl
4
1 ppm NH Cl
400
Potential, mV
SCE
4
5 ppm NH Cl
4
10 ppm NH Cl
4
200
0
-200
-400
-10
-9
-8
-7
-6
-5
-4
-3
-2
2
log (i[A/cm ])
(e)
Figure 4. Polarization diagrams for 316L stainless steel in seawater: (a) without
biocides; (b) with glutaraldehyde and 1 ppm sulfide; (c) with
144
-3
(b)
Potential, mV
SCE
Potential, mV
-4
log (i[A/cm ])
(a)
400
-5
2
2
log (i[A/cm ])
-3
Alhajji and Valliappan
glutaraldehyde and 10 ppm sulfide; (d) with ammonium chloride and 1
ppm sulfide; and (e) with ammonium chloride and 10 ppm sulfide
145
Oil Field Corrosion
250
200
0 ppm glutaraldehyde
10 ppm glutaraldehyde
50 ppm glutaraldehyde
100 ppm glutaraldehyde
(KOhm.cm )
1 ppm Sulfide
10 ppm sulfide
200
2
150
150
imaginary
100
Z
50
100
50
Z
imaginary
2
(KOhm.cm )
0 ppm sulfide
0
0
0
50
Z
100
150
0
20
40
2
real
(KOhm.cm )
60
Z
80
100
120
2
(KOhm.cm
real
(a)
140
160
)
(b)
0 ppm Glutaraldehyde
10 ppm Glutaraldehyde
50 ppm Glutaraldehyde
100 ppm Glutaraldehyde
250
0 ppm NH 4 Cl
1 ppm NH 4 Cl
5 ppm NH Cl
4
10 ppm NH Cl
200
2
(KOhm.cm )
150
4
150
100
imaginary
100
50
Z
Z
imaginary
2
(KOhm.cm )
200
0
0
50
Z
100
50
0
150
0
50
2
real
(KOhm.cm )
Z
(c)
real
100
(KOhm.cm
150
2
)
(d)
200
150
100
50
Z
imaginary
2
(KOhm.cm )
0 ppm NH 4 Cl
0
0
50
100
Z
real
150
(KOhm.cm
200
2
250
)
(e)
Figure 5. Nyquist plot for 316L stainless steel in seawater: (a) without biocides;
(b) with glutaraldehyde and 1 ppm sulfide; (c) with glutaraldehyde and
10 ppm sulfide; (d) with ammonium chloride and 1 ppm sulfide; and
(e) with ammonium chloride and 10 ppm sulfide in seawater
146
Alhajji and Valliappan
6
6
0 ppm Sulfide
0 ppm glutaraldehyde
10 ppm glutaraldehyde
50 ppm glutaraldehyde
100 ppm glutaraldehyde
1 ppm Sulfide
5
2
5
Log |z|, Ohms-Cm
Log |z|, Ohms-Cm
2
10 ppm Sulfide
4
3
2
4
3
2
1
1
-4
-2
0
2
4
6
-4
-2
Log [frequency], Hz
0
2
4
6
Log [frequency], Hz
(a)
(b)
6
0 ppm NH Cl
4
1 ppm NH 4 Cl
5 ppm NH Cl
4
10 ppm NH 4Cl
2
5
Log |z|, Ohms-Cm
2
5
Log |z|, Ohms-Cm
6
0 ppm glutaraldehyde
10 ppm glutaraldehyde
50 ppm glutaraldehyde
100 ppm glutaraldehyde
4
3
2
1
4
3
2
1
-4
-2
0
2
4
6
-4
-2
0
Log [frequency], Hz
2
4
6
Log [frequency], Hz
(c)
(d)
6
0 ppm NH 4Cl
1 ppm NH 4Cl
5 ppm NH 4Cl
10 ppm NH 4Cl
Log |z|, Ohms-Cm
2
5
4
3
2
1
-4
-2
0
2
4
6
Log [frequency], Hz
(e)
Figure 6. Bode plot for 316L stainless steel in seawater: (a) without biocides; (b)
with glutaraldehyde and 1 ppm sulfide; (c) with glutaraldehyde and
10 ppm sulfide; (d) with ammonium chloride and 1 ppm sulfide; and
(e) with ammonium chloride and 10 ppm sulfide
147
Oil Field Corrosion
In the present investigation, the effect of sulfide on the passivity breakdown, thereby
increasing the pitting tendency, for the increase in sulfide concentration is clearly revealed.
The addition of glutaraldehyde, a nonoxidizing biocide, also resulted in a lower pitting
potential than the blank system. The least was observed for the intermediate concentration of
50 ppm. The analysis based on the anodic current density corresponds to the passive region
indicating that the susceptibility mode is more pronounced than in the biocide-free system.
The addition of ammonium chloride as an oxidizing chemical changed the entire scenario and
involved an increased susceptibility of stainless steels to pitting. The pitting potential values
from Tables 1 and 3 indicate this effect. This is not the case all the time or for all
concentrations investigated. Increases in the concentration of sulfide to 10 ppm with the
addition of NH4Cl resulted in improved pitting resistance based on the pitting potential values
compared to the blank one. The anodic current density values corresponding to the passive
region showed a high value for the concentration of 5 ppm of ammonium chloride. The same
concentration for mild steel in a sulfide-polluted system increased the dissolution rate. Under
inorganic, near neutral conditions, sulfide and other sulfur species are known to decrease both
the pitting potential and the repassivation potential of stainless steels [23]. Particularly,
chloride ions are aggressive to stainless steels. The pitting potential is reduced as the
concentration of chloride increases [24]. However, the effects of the aggressive anions are
reduced by the presence of inhibiting anions. Notably sulfate [24], hydroxide [25] and
acetate [25] inhibit the pitting corrosion of stainless steel to varying extents, and this was the
case observed in the present investigation for higher concentrations of sulfide where the
presence of an oxidizing agent led to the formation of sulfate as the resultant oxidizing
product. Figures 5a-e and 6a-e-show the complex plane impedance diagrams (Nyquist plots)
and the Bode plots for 316L stainless steel in seawater with and without a sulfide polluted
system (i.e., addition of glutaraldehyde and ammonium chloride. These plots do not allow a
reasonable estimate for Rp because the valid Rp values can only be obtained using the AC
impedance technique if the low frequency locus in the complex plane impedance diagram is
essentially complete. The behavior, however, is of the Warburg type for diffusion controlled
reactions.
CONCLUSIONS
1. The oxidizing and nonoxidizing biocides generally alter the corrosion tendency of
mild steel and the pitting tendency of 316L stainless steel in sulfide polluted
environments.
2. The increase in corrosion rate of mild steel was observed for the glutaraldehyde
concentration of 50 ppm added to the lower concentration of sulfide i.e., 1 ppm. This
trend completely reversed for a higher concentration of sulfide i.e., 10 ppm.
3. The increase in the concentration of sulfide, i.e., 10 ppm, for the oxidizing type of
chemical added to the synthetic seawater resulted in a decrease of the corrosion rate of
mild steel as well as a lower pitting susceptibility for stainless steel.
ACKNOWLEDGMENT
This work was supported financially by the Kuwait Foundation for the Advancement of
Science (KFAS) and Kuwait University. This support is gratefully acknowledged.
148
Alhajji and Valliappan
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3.
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R.G.J. Edyvean, Proceedings 6th International Congress on Marine Corrosion and
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R.E. Tantnall, Materials Performance 19, 8, 1980, p. 88.
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USA, 1986, p. 285.
149
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
EVALUATION OF MICROBIALLY INFLUENCED CORROSION RISKS AND
CONTROL STRATEGIES IN SEAWATER AND PRODUCED WATER INJECTION
SYSTEMS, KUWAIT
P.F. Sanders1, M. Salman2 and K. Al-Muhanna2
1
Oil Plus Limited, Hambridge Road, Newbury, Berkshire, RG14 5TR, England
2
Kuwait Institute of Scientific Research, P.O. Box 24885, 13109 Safat, Kuwait
ABSTRACT
Injection of seawater and produced (effluent) water for secondary oil recovery can encourage the
growth of bacteria, particularly in biofilms on pipewall surfaces. One particular problem in water
injection systems is the uncontrolled growth of sulphate-reducing bacteria (SRB) which leads to
increased corrosion of the process plant. Injection water sources in Kuwait range from brackish water
(total dissolved solids, i.e., TDS = 4,000 mg/l) through seawater (TDS = 30,000 mg/l) to highly saline
brines (TDS = 200,000 mg/l). In addition, some of these water sources are highly sulphide sour and
may require treatment to prevent scaling, corrosion or iron sulphide precipitation.
Studies were undertaken to evaluate the relative risks of microbially influenced corrosion in a
range of Kuwait’s water sources, using a combination of field sampling and laboratory biofouling
trials. Bacteria isolated from the field surveys were used to evaluate remedial treatments such as
chlorine, chemical biocides and mechanical removal. Recirculating biofouling loops were set up with
the appropriate site water, and inoculated with the bacteria from the system so that an active biofilm
was set up on small steel studs. These biofouled studs were treated with proprietary antibacterial
products under various dose regimes in order to select the most appropriate control regime for
particular water chemistries and process options.
Key Words: Microbially influenced corrosion, seawater, water injection systems, secondary oil
recovery, remedial treatments
INTRODUCTION
Pitting corrosion is characteristic of sulphate reducing bacteria (SRB) attack of steels,
with the pits being open and filled with soft black corrosion products in the form of iron
sulphides [1,2]. When the corrosion products are removed, the metal underneath is bright but
rapidly rusts on exposure to air. With buried pipelines in anaerobic environments, such as
waterlogged clay soils, the corroded area as a whole, will be covered with a film of iron
sulphides. In aerobic environments, SRB corrosion invariably occurs beneath deposits of
inorganic or organic detritus, microbial slimes, or tubercules caused by the action of the ironoxidising bacteria [3,4].
Unique features of corrosion caused by SRB are that it occurs at neutral pH and in
anaerobic environments. Oxygen is not involved, and the corrosion products include iron
149
Oil Field Corrosion
sulphides. The exact mechanisms of the corrosion reactions are still under debate, and a
range of possible mechanisms has been proposed [1,2].
Normal electrochemical corrosion would not be expected to occur under anaerobic
conditions, because the cathode becomes polarized by the build-up of a layer of atomic
hydrogen. SRB are potentially able to stimulate the normal electrochemical corrosion
mechanism by a number of possible means, the most widely accepted being
•
•
•
•
The enzymic removal of polarizing cathodic hydrogen,
The formation of iron sulphides which are themselves cathodic to steel,
The formation of elemental sulphur upon re-oxidation of the sulphides,
The formation of aggressive iron phosphides by SRB.
It is now generally accepted that, although all of these mechanisms may operate in
nature, the production of iron sulphides and cathodic depolarisation are the two most
important corrosion mechanisms.
Control of microbes can be effected by chlorination, ozonization, by means of ultraviolet
(UV) light and/or organic biocides. A combination of these methods is commonly used to
minimise microbially influenced corrosion and reservoir souring in oilfield water injection
and reservoir systems. For seawater injection dosing systems, chlorine injection at the
seawater lift pumps is the most common primary method of control [5]. However, chlorine is
generally lost at the deaeration stage by reaction with oxygen scavenger, and biological
control is therefore maintained in the remainder of the system by the use of organic chemical
biocides. The most common strategy is a regular (frequently weekly) batch treatment for 3-6
hours. The dose rate and contact time are both critical in order to optimise the killing effect
of the biocide;: in most cases laboratory trials are used to determine the best treatment.
Most water injection systems, therefore, employ chemical biocide treatments [6,7] with
varying degrees of success [8,9]. Great care must be taken in selecting and applying biocides
to ensure that they are compatible with the system, can be handled safety, and have no
environmental impact. Additionally, both batch and low-level continuous treatments of
organic biocides will ultimately lead to the build-up of resistant and tolerant strains of the
bacteria. Biocide effectiveness will, therefore, decline with time, and monitoring must take
place to identify when a change in biocide type is required.
In most cases, a proprietary organic biocide is used as a batch dose to replace the
continuous microbiological control exerted by chlorine. There are many biocides offered by
chemical supply companies, which may carry out field tests to determine which of their range
is best suited to a particular problem. An unbiased view can, however, only be given by the
operator or an independent laboratory carrying out a testing program for products from a
range of supply companies. Despite the apparent large range of biocides, there are, in fact, a
relatively limited number of molecules that are regularly used. The most widely used
biocides are listed below. Many products are mixtures of one or more biocides with
surfactant that aids penetration of deposits and provides cleansing action.
• Chlorine and chlorine-releasing compounds,
• Aldehydes, e.g., glutaraldehyde, formaldehyde, that combine with genetic
material to prevent cell replication and protein formation,
150
Sanders et al.
• Quaternary ammonium compounds and quaternary phosphonium compounds
which are good penetrating agents and attack the outer membrane of cells
causing rupture and leakage of their contents, and
• Amines, e.g. guanide compounds, that attack genetic material and prevent cell
growth and replication.
Miscellaneous compounds such as acrolein (substituted alkene) and isothiazalones have
been successfully tested against SRB. They are beginning to be used in water injection
systems for specific applications.
Biocide test methods for water injection systems fall into two main categories:
planktonic and sessile methods. The following outlines illustrate the basics of the methods.
In many cases, a combination of the two methods must be used in order to arrive at the most
effective and appropriate microbiological control regime for a specific system. Traditional
methodology has concentrated on determining the numbers of SRB and other bacteria
suspended in fluids, typically in water from injection, oil production and oil storage facilities
[10]. Recent work has shown clearly that these planktonic bacteria are a very small fraction
(typically less than 1%) of the total population [11]. Novel, rapid methods can sometimes be
used to assess the numbers or activity of bacteria [12], but they must be used with care to
ensure that results are reliable [9].
Planktonic Biocide Tests
These typically follow the methods of CCEJV [13] and have the basic weakness that
they do not measure the ability of the treatment to penetrate biofilms. Planktonic bacteria are
more susceptible to biocides than are sessile ones, and a lower effective concentration is
indicated from these tests than from sessile tests.
Bacteriostatic Test
This is a simple screening test to rank a large number of potential treatments. Biocide is
added to SRB growth media at a range of concentrations, followed by an inoculum of SRB
from the system, freshly grown in SRB media. The bottles are incubated and checked for
SRB growth by blackening. A minimum inhibitory concentration (MIC) can be determined
for each biocide. This test does not measure bacterial kill but merely the biocide’s ability to
prevent growth.
Time-Kill Test)
The time-kill (TK) test provides a more realistic test: system water samples are used, and
killing ability is determined. It is more time-consuming, however, and is typically used to
select the most suitable from a small number of biocides. Biocide is added in the desired
concentration to water samples from the system, followed by an inoculum of SRB from the
system, freshly grown in SRB media. Sub-samples are taken at intervals and tested by MPN
for surviving SRB.
After incubation, time-kill curves are plotted and a suitable
concentration/contact time combination can then be selected.
Sessile Biocide Tests
151
Oil Field Corrosion
Simple recirculating biofilm generators are set up in the laboratory to mimic system
conditions, allowing biofouling to occur with bacteria freshly isolated from the system. After
a period of biofouling, the biostuds are removed and placed in stirred test biocide solutions in
anaerobic seawater. After the desired contact time, these are removed and the surviving
bacteria are enumerated, thus enabling TK curves to be determined for each biocide.
Special Tests
Appropriate test methods should be devised for other systems; for example, biocides for
reservoir application should be tested at high temperature and pressure and with the presence
of sand. Sand-packed columns are also useful for selecting biocides for hydrocarbon
reservoir production system applications.
Field Test Methods
Both planktonic and sessile biocide trials can be carried out in the field on samples of
water, solids or biostuds, derived from the system. A side stream is ideal for this purpose as
biostuds can be used. Whilst laboratory derived results may relate to the system under study,
it is generally better to carry out such tests in the field, under system conditions, so that the
results can be directly related to the system [9,14,15].
In order to confirm the performance of a selected biocide regime, close monitoring of the
system should be carried out before and after any change is made, by analysing the biostuds
before and after biocide dosing. Once effectiveness is confirmed, monitoring can revert to a
normal schedule until biocide-resistant and tolerant strains build up in the system.
METHODS
Three different systems in Kuwait were studied for their biocide treatment requirements.
These were
• A seawater injection system,
• A complex produced water/effluent, and
• A highly sour aquifer water injection project.
The biocide test methods were selected to be as relevant as possible for each individual
system, ranging from MIC tests, through time-kill tests on biofouled studs to complex
dynamic tests in a recirculating loop system.
Bacteria were isolated from each site and grown in nutrient-enriched media. The media
were based on standard formulations, but the brine composition was, in each case, matched to
that of the system under investigation.
Minimum Inhibitory Concentration (MIC) Test
In order to obtain very basic information on the biocide, MIC tests were carried out [13].
A stock culture of mixed general aerobic bacteria (GAB), general anaerobic bacteria (GAnB)
and SRB was used. Each vial was inoculated with the active stock culture (0.1 ml). The vials
were then immediately injected with small volumes (ca 0.1 ml) of the biocide, diluted in
anaerobic, sterile, effluent water. In this way, a wide range of biocide concentrations was set
up to obtain general data on the various biocides.
152
Sanders et al.
Planktonic Time-Kill (TK) Tests
As with the MIC tests, these methods followed the CCEJV methods [13], using a
bacterial suspension in 125 ml Wheaton bottles. Water was filtered (Whatman No. 1),
dispensed in 100 ml volumes into 125 ml Wheaton bottles and autoclaved (121°C/15
minutes) before cooling under nitrogen and capping. Stock biocide solutions were made up
(10,000 ppm) and relevant volumes accurately dispensed into the Wheaton bottles using a
variable automatic pipettor. Each Wheaton bottle was then immediately inoculated with 1 ml
of the stock culture and stirred on a multipoint magnetic stirrer at 850 rpm. Samples were
withdrawn from the Wheaton bottles at 0, 1, 3 and 6 hour contact times, and MPN analyses
were conducted.
Chlorine was dosed (as a solution of sodium hypochlorite) into aerobic, unfiltered water.
Because the free chlorine concentrations were low and the chlorine demand was high,
hypochlorite was redosed midway through the contact period in an attempt to retain a residual
chlorine concentration.
Sessile Time-Kill (TK) Tests
Biofilm Generation
Biofilms were generated using the Biocide Evaluation Test Rig (BETR) system (Figs. 1
and 2). This had six sessile bacterial monitoring tubes (SBMTs), each with 24 carbon steel
studs. The BETR was run on a daily feed and bleed system.
Figure 1. Schematic diagram of 6 SBMT mounted in the recirculating BETR
153
Oil Field Corrosion
Figure 2. Detailed schematic diagram of 1 SBMT
Water was treated with a minimal concentration of nutrients to encourage bacterial
growth. It was then autoclaved to remove oxygen and cooled under nitrogen to maintain
anaerobic conditions. The reservoir was filled with 20 litres of appropriate water and dosed
with nutrients. The pH of the water was adjusted using dilute HCl and/or NaOH. Anaerobic
conditions were maintained throughout by purging with nitrogen during all manipulations.
After pH adjustment, the water was added to the main reservoir of the BETR under a
nitrogen blanket. Upon initial start-up, 20 litres of the relevant water plus nutrients was
inoculated with 100 ml of SRB, 100 ml of GAB and 100 ml of GAnB stock cultures for
enrichment. Each SBMT was run in parallel with the others. Six SBMTs were set up at the
same time and treated identically. Flow was maintained continuously at 8 litres per minute,
equating to 1.0 metre per second, this being a typical velocity for water injection systems.
Throughout the period of the test, the water flowing through the BETR was typically of
the following characteristics: dissolved oxygen was <10 ppb, temperature was 30-34°C, free
chlorine was 0, and residual bisulphite was 0.
The SBMT were run without any additional sulphide in an attempt to encourage bacteria
to colonise the rig. The effect of sulphide, if required, was accounted for during the biocide
dosing phase. Some tests studs were tested with sulphide at 75 mg/l, whilst parallel tests were
undertaken without any additional sulphide.
Biocide Dosing
Biocides were supplied in labelled bottles by the chemical vendors. Samples were
withdrawn from the bottles and used in the laboratory tests. Table 1 gives the details of the
products supplied for testing.
For stirred biocide tests, products were made up to the desired concentration in site water
in 125 ml Wheaton bottles; dispensing was carried out using adjustable automatic micro
pipettors and suitable dilution. For the sulphide-free tests, the studs were added directly to
the biocide solution; for the tests containing sulphide, a stock sulphide solution was made up
and the appropriate volume added to the Wheaton bottle to give the desired concentration,
prior to adding the stud to the biocide. Studs were removed from the SBMTs and suspended
in the biocide/sulphide solution before assessing surviving bacterial numbers.
For dynamic tests, biocides were injected directly into the SBMTs at the desired
concentration, and the studs were removed before and after dosing. Subsequently, biocide
doses were repeated at increasing concentrations to identify the most effective regime.
Bacterial Evaluation
The exposed surface of the stud was scraped carefully and all deposits were transferred
to 10 cm3 of sterile seawater diluent. The diluent tube was then treated for 30 seconds in an
ultrasonic bath followed by 60 seconds in a vortex mixer and a further 30 seconds in the
ultrasonic bath. This treatment disrupts biofilm/deposits but does not kill bacteria. The
treated diluent tube was used to inoculate triplicate series (to 10-7) of the following media,
154
Sanders et al.
specially formulated for the bacteria grown from each site: SRB - Postgates B; GAB glucose-based aerobic broth; Gan - glucose based anaerobic broth.
One cubic centimetre of the diluent was injected into each broth series and 1 in 10
dilutions were made. The procedure followed industry standard practice for MPN (most
probable number) tests [13]. GAB and GAnB results were taken after 5 days of incubation at
30°C. SRB results were taken after 28 days of incubation at 30°C.
CASE HISTORY 1: PRODUCED EFFLUENT WATER REINJECTION SYSTEM,
WEST KUWAIT
A complex system with various water sources for potential reinjection into the oil
reservoir was tested. A 50:50 mixture of produced water and aquifer water was selected to be
A full range of tests was carried out on 9 biocide products (see Table 1). MIC,
planktonic time-kill, sessile time-kill, chlorination and mechanical removal were all tested,
with and without 75 mg/l of sulphide.
Due to the high salinity and high H2S content of the water, bacterial fouling was low in
both the system and the laboratory BETR. In fact, it proved difficult to isolate bacteria from
the system.
The simple MIC tests yielded a generalized ranking of the products, but little information
could be gained about the realistic dose rates required. More information was obtained from
the planktonic time-kill tests, particularly with respect to speed of kill. The most realistic
testing (against sessile bacteria) in the sulphide-free scenario gave a very similar ranking to
the planktonic tests. Very importantly, however, the resulting recommended dose rates are
much higher against sessile bacteria because of the high bactericide demand caused by the
presence of inactivating organic materials in the biofilm.
Table 1. Results of Testing of Nine Biocide Products on a Complex Produced Effluent Water
Biocid
e
Formulation
Ranking
(MIC)
1A
Mixed Aldehydes
+ QAC + QPC
Mixed Aldehydes
+ QAC
Glutaraldehyde +
QPC
Aldehydes
+
QAC
Fatty Amines
Aldehydes
+
QAC + Surfactant
Glutaraldehyde
Mixed Aldehydes
+
QAC
+
Surfactant
Formaldehyde +
4
1B
1C
1D
1E
1F
1G
1H
1I
2=
Planktonic Time-Kill
(Sweet)
Ranking
Dose
(ppm/hrs)
6
100/6
5
100/6
8
2=
100/6
1
5=
2=
100/6
5=
7
2=
1
100/6
25/6
Sessile Time-Kill
(Sweet)
Ranki
Dose
ng
(ppm/hrs)
5
>>500/6
6
5
2=
5=
Planktonic Time-Kill
(Sour)
Ran
Dose
king
(ppm/hrs)
3
100/3
>>500/6
100/6
4
1
100/1
25/3
1
25/6
6
100/3
Sessile
(
Rankin
g
2
4
3
500/6
2
500/6
4
1
>500/6
500/3
3
5
1
6
155
Oil Field Corrosion
QPC
representative of the system. This mixture had a TDS of 130,000 mg/l. One stream had a
TDS of 259,000 mg/l, whilst another was of brackish water composition. One of the waters
contained a sulphide level of 75 mg/l, and therefore, sour conditions were also tested.
A similar picture emerged from the sour testing. As with the sulphide-free tests, the
sessile ranking was similar to the planktonic ranking for the sulphide-containing tests. The
only difference in the rankings was that Biocide 1A (effective in planktonic tests) was not
very effective in the sessile tests, probably due to a lack of penetration into the biofilm.
Again, a much higher demand for biocide is indicated when the sessile and planktonic tests
are compared. The comparison of the planktonic and sessile rankings (Table 2) indicates the
comparability between the two types of tests.
Chlorination at conventional levels will not, in fact, be effective in controlling planktonic
bacteria, even under the most favourable conditions. GAB numbers were initially 1.4 x
105/ml, and no measurable reduction was detected at 0.2 mg/l over 2 hours. Even at 2 mg/l
for 2 hours, GAB were only reduced to 4.5 x 104/ml (i.e., less than one order of magnitude).
GAnB showed slightly better kill rates, but 2 mg/l free chlorine only reduced numbers by one
order of magnitude (from 4.5 x 104/ml to 3.0 x 103/ml). SRB behaved similarly (1.5 x 104/ml
was reduced to 1.4 x 103/ml by 2 mg/l chlorine over 2 hours). Chlorination was, therefore,
not considered as a feasible treatment regime for this system.
Physical removal of corrosion deposits, biofouling, scale, settled solids, etc., in pipelines
and vessels is a common cleanup method. A wide range of cleaning tools and devices are
available to mechanically clean pipelines. These include foam spheres, nylon brush pigs,
wire brush pigs, flexible plastic pigs and (the most aggressive) milling pigs.
Normal oilfield operations combine chemical treatment programs (including biocides)
with regular pigging operations of pipelines. Different types of physical removal of biofilms
were tested, using the biofouled studs from the SBMTs. Initial bacterial populations
consisted of 2.5 x 104 GAB, 2.0 x 104 GAnB and 4.5 x 104 SRB per stud. The biofilm was
very thin and corrosion products were not in evidence. This meant that even a small physical
perturbation was sufficient to dramatically reduce the numbers of sessile bacteria.
Table 2. A Comparison of Planktonic and Sessile Rankings
Bactericide
Planktonic Ranking
Sessile Ranking
1H
1F
1D
1G
1A
1B
1
2=
2=
4
6
5
1
2
3
4
5
5
(a) Sweet (sulphide stripped)
Recommended Dose
(Sessile)
500 ppm/3 hours
500 ppm/6 hours
500 ppm/6 hours
> 500 ppm/6 hours
>> 500 ppm/6 hours
>> 500 ppm/6 hours
Sessile Ranking
Recommended Dose
(b) Sour (sulphide at 75 mg/l)
Bactericide
Planktonic Ranking
156
Sanders et al.
1G
1A
1E
1C
1F
1I
2
3
4
5
1
6
1
2
3
4
5
6
(Sessile)
500 ppm/3 hours
500 ppm/6 hours
500 ppm/6 hours
> 500 ppm/6 hours
>> 500 ppm/6 hours
>> 500 ppm/6 hours
A single pass of a stiff nylon brush was the most gentle of the treatments, probably
relating to use of a foam pig in a real system. Even this minimal treatment reduced the
biofilm population to 4.5 x 102 GAB, 1.5 x 103 GAnB and 1.5 x 102 SRB per stud, a reduction
of 1 or 2 orders of magnitude.
A more aggressive treatment with a wire brush produced even better results, with
residual populations of 1.5 x 102, 9.5 x 101 and 9.5 x 101 GAB, GAnB and SRB respectively.
This treatment would reflect the use of a wire brush pig in a pipeline, and the results
demonstrate the dramatic cleanup effects of such a treatment. Reductions were in the region
of 3 orders of magnitude for all 3 microbial groups tested, and very low residual microbial
populations were left on the surface of the cleaned steel.
Finally, vigorous use of a sterile stiff plastic scraper resulted in the most effective
removal of biofilm. Only 7 GAB, 7 GAnB and 75 SRB per stud remained after this
treatment. These numbers can be considered to be insignificant and the conclusion must be
that this treatment effectively removed all the biofilm, leaving only a few bacteria in
inaccessible micro-environments. Such treatment probably equates to the use of a combined
brush/scraper pig in the field.
Although the trial could not be carried out on long-term corroded surfaces (which would
be more difficult to clean by mechanical means), the results demonstrate that mechanical
removal, if practised regularly and frequently, would be an effective biofouling/microbially
induced corrosion preventative measure. This would be particularly true if pigging was
combined with chemical biocide treatment. Under normal circumstances, a biocide treatment
would be used to kill those bacteria easily accessible on surfaces and to loosen biofouling
deposits. The mechanical scraper (pig) would then be used a few hours later to complete the
removal process, perhaps followed immediately by a second biocide slug.
CASE HISTORY 2: AQUIFER/EFFLUENT WATER INJECTION SYSTEM, SOUTH
KUWAIT
This was another mixed water injection system, this time with sulphide in all of the
streams. A more limited test program was deemed adequate given the low biofouling
measured in the field and the laboratory.
Three separate waters were tested:
• Aquifer A - low sulphate, 186,000 TDS, 140 mg/l sulphide
• Effluent B - moderate sulphate, 76,000 TDS, 250 mg/l sulphide
• Mixed waters - moderate sulphate, 108,000 TDS, 220 mg/l sulphide
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Oil Field Corrosion
Eight biocides were submitted for testing; the products were pre-screened in the vendors
laboratories for tolerance to sulphide and salinity (see Table 3).
The testing in this case was undertaken at the suppliers dosages, on fouled studs from
SBMTs. Bacteria from the system were used to inoculate the SBMTs, one SBMT being set
up for each water chemistry.
Aquifer A
The aquifer water was very highly saline and contained very low bacterial numbers
during the site analysis. Bacteria were only detected when 100 ml of water was filtered; the
low numbers are due to
• The high temperature of the water,
• The high salinity of the water, and
• The high sulphide concentration of the water.
The bacteria were likely to be present in a dormant, inactive, state and bacterial growth
was evidently minimal under the aquifer conditions.
This low growth was confirmed by the minimal growth seen in the SBMTs. Even after
48 days of biofouling under ideal conditions (35°C, nutrients added, steady flow), bacterial
colonisation of the studs was miniscule. Under normal conditions, such rigs (with seawater
for example) would have a population of 106 GAB and 105 SRB after 14 days, and 107 GAB
and 106 SRB after 28 days. It is likely that the aquifer water is sulphate-limited which would
be an additional factor predisposing the water to low bacterial growth, particularly that of
SRB.
Table 3 Composition of Biocides Submitted for Testing
(a) Case History 1
Biocide
Formulation
1A
Mixed aldehydes + QAC + QPC
1B
Mixed aldehydes + QAC
1C
Glutaraldehyde + QPC
1D
Aldehydes + QAC
1E
Fatty amines
1F
Aldehydes + QAC + surfactant
1G
Glutaraldehyde
1H
Mixed aldehydes + QAC + surfactant
1I
Formaldehyde + QPC
QAC = Quaternary Ammonium Compound
QPC = Quaternary Phosphonium Compound
(b) Case History 2
Biocide
2A
2B
158
Formulation
Biguanidine + QAC
Aldehydes
Suppliers
Recommended
Dose/contact time
500 ppm/3 hours
500 ppm/3 hours
Sanders et al.
2C
Fatty amine arylquaternary +
fatty amine salt
Aldehydes + QAC
Aldehydes
QAC
Aldehydes + QAC
Glutaraldehyde + QPC
2D
2E
2F
2G
2H
(c) Case History 3
Biocide
3A.1
3A.2
3B.1
3B.2
3C.1
3C.2
3D.1
3E.1
3E.2
Formulation
Aldehydes
Aldehydes + amines
Aldehydes + amines
Glutaraldehyde + QPC
Fatty amines
Aldehydes + QAC
Formaldehyde + QAC +
QPC
Aldehydes
Aldehydes + QAC
400 ppm/4 hours
400 ppm/4 hours
500 ppm/3 hours
400 ppm/3 hours
300 ppm/4 hours
1000 ppm/4 hours
Suppliers Recommended
Dose/contact time
700 ppm 3 hours/week
700 ppm 3 hours/week
1000 ppm 4 hours/2 weeks
1000 ppm 4 hours/2 weeks
250 ppm 4 hours/week
250 ppm 4 hours/week
600 ppm 4 hours/week
200 ppm 4 hours/week
200 ppm 4 hours/week
Given the low biofouling population, it is not surprising that most biocides worked well,
reducing numbers to zero. For the sweet system, the ranking for the aquifer water (Aquifer
A) was
2H = 2G = 2E > 2C = 2B,
and for the sour system, (140 ppm sulphide), the ranking for the aquifer water was
2G = 2A = 2F > 2D > 2H
These rankings are based upon very small differences in performance.
Effluent B
The effluent SBMTs blackened due to the production of sulphide by SRB, but SRB
growth and biofouling was slow. Bacteria were present in the system in moderate numbers.
However, the salinity of the water was high (76,000 TDS), which would have reduced the
growth rate of any bacteria.
GAB were easier to kill then were SRB in this system. This is unusual (GAB are usually
more resistant to biocides in biofilms than are SRB) and suggests that an active biofilm was
not present on the studs. This was probably due to the combined effects of high salinity and
residual corrosion inhibitor from the site water.
Given the very low initial biofouling population, the rankings of the biocides were as
follows. For the sweet (no sulphide) situation in the effluent water (effluent B):
2H > 2G = 2E > 2C = 2B,
159
Oil Field Corrosion
and for the sour water (250 ppm sulphide) situation in the effluent water (Effluent B):
2H > 2G > 2F > 2A > 2D
As for aquifer A, the bacterial numbers were very low, and the differences between the
biocides were minor.
Mixed Water
This water appeared to be the worst case from a biofouling point of view. This was
likely to be due to the addition of sulphate to the aquifer water from the effluent water, the
dilution of the aquifer's high salinity, and the dilution of possibly biostatic corrosion
inhibitors from the effluent water.
Control numbers of 300-650 GAB and 115-450 SRB per stud were recorded on this
SBMT, and most biocides reduced numbers substantially. As with other SBMTs, GAB were
reduced more than SRB, suggesting that a coherent biofilm did not form on the studs. Most
biocides performed better in the sweet tests than in the sour tests, suggesting that many
biocides initially acted as sulphide scavengers, thus reducing the effective biocide
concentration.
For the non-sour tests (no added sulphide), with mixed water C (mixed effluent and
aquifer) the biocide ranking was
2H > 2C > 2G > 2E > 2B
and for the sour tests (sulphide added at 220 ppm), with water C (mixed 2H/effluent/aquifer),
the biocide ranking was
2H > 2G > 2D > 2F > 2A
CASE HISTORY 3: SEAWATER INJECTION SYSTEM, NORTH KUWAIT
This is a conventional seawater injection system and, in this case, dynamic trials were
undertaken by injecting alternate biocides directly into the pre-fouled, recirculating SBMTs.
One SBMT was used for each combination biocide treatment. Formulations of the biocides
are shown in Table 3.
Biocides were dosed to the other 5 SBMTs over a 4 week period, steadily increasing the
dose from 50% to 100% of the manufacturers recommended dose concentration, retaining the
recommended frequency and dosage time. In all cases, a pronounced saw-tooth effect was
evident, with bacterial numbers recovering rapidly once the biocide was removed from the
system. This saw-tooth effect is regularly seen in such trials and in field monitoring of water
injection systems. Biocides tend to kill only a proportion of the bacteria present in biofilms,
and the survivors rapidly grow to re-form the active biofilm. This confirms the need to dose
biocide frequently in order to keep sessile bacterial populations at a low level.
Biocide Regime A
Both biocides were each dosed up to 700 ppm of the product for 3 hours once per week
(i.e., two biocide treatments each week). Biocide 3A.1 was an aldehyde mixture that did not
foam, whilst 3A.2 was a mixture of aldehydes and amines that did form foam when agitated.
160
Sanders et al.
At half the recommended dose (350 ppm), biocide 3A.1 did not perform well, but
biocide 3A.2 reduced bacterial numbers to ca 103 per stud. At three-quarters of the
recommended dose (525 ppm), biocide 3A.1 was more effective, particularly against SRB,
but biocide 3A.2 was only marginally more effective. At 90% of the recommended dose (630
ppm), biocide 3A.1 reduced numbers to 104 per stud while biocide 3A.2 reduced numbers to
102-103 per stud. At 700 ppm, biocide 3A.1 reduced bacterial numbers to 103 per stud, and
biocide 3A.2 reduced them to 101 to 102 per stud.
Bacterial numbers appear to recover rapidly (within a few days) after all the biocide
doses. After the 75% dose of biocide 3A.2, however, numbers did not recover to their
original level: 104-105 cells of each bacterial type were present compared to 106 per stud in the
early stages of the trial. The data indicate that 3A.2 (containing an amine) was particularly
effective at 630 ppm and above, and that 3A.1 above 525 ppm was able to support the effects
of 3A.2. Continued treatment at 630-700 ppm with this biocide regime would be likely to
maintain (and probably substantially reduce) sessile bacterial populations under these
conditions.
Biocide Regime B
These biocides were dosed only once per week with the recommended concentration
being 1000 ppm; 3B.1 was an aldehyde and amine blend dosed for 4 hours, 3B.1 was a mixed
glutaraldehyde and THPS blend, again dosed for 4 hours. Biocide 3B.1 was strongly foaming
whilst biocide 3B.2 did not foam when agitated. This once-per-week treatment gave good
kill of all 3 bacterial types at 750 ppm (106/stud was reduced to 102-103 per stud), but within a
few days the initial bacterial population levels were reestablished. There appears to be no
significant difference in the data from the 1000 ppm dosing regime as compared to the 750
ppm dosing regime for GAB and GAnB (numbers were reduced to 101-103 per stud). Data
from the Regime B trial indicated that one biocide dose per week is insufficient, and that 750
ppm for 4 hours is as effective as 1000 ppm for 4 hours.
Biocide Regime C
The biocides were dosed twice per week. Biocide 3C.1 was a fatty amine (250 ppm, 4
hours) and biocide 3C.2 was an aldehyde/QAC mixture (250 ppm, 4 hours). Both biocides
formed foam when agitated, so they would not be suitable for dosing upstream of a deaerator
tower.
Biocide 3C.1 was very effective against GAB, GAnB and SRB;: even at 125 ppm,
bacterial numbers were reduced from 106 to 102-103. Above 187 ppm, 3C.1 reduced numbers
even further, in some cases to below 10 cells per stud. Biocide 3C.2 was, however, less
effective. Even at 250 ppm, numbers were reduced by only 1 or 2 orders of magnitude.
Despite the relatively poor performance of biocide 3C.2, however, there was a general
downwards trend in bacterial numbers over the course of the trial and continuation with such
a regime would probably lead to effective microbiological control. The most effective
biocide in this case appeared to be 3C.1. Biocide 3C.2 could be improved by increasing the
dosing concentration or changing to an alternative type.
Biocide Regime D
This vendor chose to submit only one chemical for testing, at a recommended dosing rate
of 600 ppm for 4 hours once per week. The biocide was a complex blend of formaldehyde,
161
Oil Field Corrosion
QAC and THPS. It formed foam when agitated, and thus, might only be dosed downstream
of a deaerator tower.
At 300 ppm, bacterial numbers were reduced to 103 per stud, but at 450 ppm they were
reduced to only 104. At 540 - 600 ppm, bacterial numbers were reduced to between 102 and
103 per stud, but only transiently. There was no apparent downwards trend in the data, and
effective microbiological control did not seem to have been imposed by this regime, probably
due to the infrequency of dosing.
Biocide Regime E
Regime E biocides were dosed twice per week. Biocide 3E.1 was a mixture of
aldehydes, dosed at 200 ppm for 4 hours, while biocide 3E.2 was an aldehyde/QAC blend,
again dosed at the recommended 200 ppm for 4 hours once per week. Biocide 3E.1 did not
foam and thus could be dosed upstream of deaerator towers, whilst biocide 3E.2 did form
foam and so would only be suitable for dosing downstream of such a tower.
This biocide regime was particularly ineffective (apart from some transient reduction in
SRB numbers early on). No consistent reduction in bacterial numbers was observed over the
course of the treatment. Even if the concentration of the products were increased, effective
control could not be guaranteed. The ranking derived from these trials is given in Table 4.
Table 4. Ranking of the Different Biocide Regimes
Most Effective
Least Effective
Regime C
Regime A
Regime B
Regime D
Regime E
Biocide 3C-1 effective at 50% dose; biocide 3C.2 less
effective even at 100% dose
Effective at 90% dose
Effective at 50% dose but needs twice per week
Effective at 90% dose but needs twice per week
Ineffective at 100% dose
CONCLUSIONS
It is clear from this trial that one biocide treatment per week is insufficient to achieve
good kill of biofilm bacteria. The data clearly show that two doses per week of an effective
chemical can give good control of a well established biofilm. Such dosing would also
minimize the buildup of biofilm in the first place if implemented from the startup of the
system.
In order to assess the suitability of any biocides for the system, other factors must be
taken into account:
1. System demand. Sufficient residual biocide must be present at the end of the
water distribution system to exert control there. A higher dose may be
required at the main treatment plant to ensure that the minimum effective
concentration is maintained at the system extremities, if there is a significant
162
Sanders et al.
biocide demand. System biocide demand must be determined as soon as the
plant is in operation.
2. Environmental impact. Certain of the biocide components may have an
adverse environment impact, and this must be assessed.
3. Safety considerations. All biocides are harmful but some may contain, for
example, carcinogens, and these may be banned from use in Kuwait.
4. Physical properties. Foaming tendency, for example, is critical in deciding
where in the system particular biocides can be dosed. Other factors such as
precipitation in the formation should also be considered.
5. Consistency of supply. Selected products must be continuously available so
that there are no periods when biocide cannot be supplied.
6. Chemical reliability. All products supplied in bulk should conform to the
samples tested without any changes in composition. Random testing of supply
for fingerprinting should be implemented to ensure consistency.
7. Local service base. Biocide vendors should ideally have a permanent base in
Kuwait to deal with technical questions and to provide a local testing and
advisory service.
REFERENCES
1.
2.
3.
4.
5.
6.
7.
8.
9.
A.K. Tiller, Aspects of microbial corrosion, In Corrosion Process, R.N. Parkins Ed.,
Applied Science, 1982, pp. 115-160.
W.A. Hamilton, Sulphate-reducing bacteria and anaerobic corrosion, Annual
Reviews of Microbiology 39, 1985, pp. 195-217.
G. Kobrin (Ed), A Practical Manual on Microbiologically Induced Corrosion, NACE,
1993.
G.H. Booth, Microbial Corrosion, M & B Monograph CE/1, Mills & Boon, 1971.
P.F. Sanders and D.L. Robinson, Corrosion control using continuous residual
chlorine, In Microbial Corrosion (Proceedings of the 2nd European Federation of
Corrosion Workshop on Microbially Induced Corrosion), European Federation of
Corrosion, Publication No. 8, Institute of Materials. C.A.C. Sequeira and A.K. Tiller
Eds., 1992, pp. 198-209.
P.F. Sanders, Monitoring and control of sessile microbes: Cost effective ways to
reduce microbial corrosion, In Microbial Corrosion 1, Elsevier Applied Science,
C.A.C. Sequeira and A.K. Tiller Eds., 1988, pp. 191-223.
W.J. Georgie, P.I. Nice and S. Maxwell, Selection, optimisation and monitoring of
biocide efficiency in the Statfjord water injection systems, In UK Corrosion ’91,
1991.
I. Ruseska et al., Biocide testing against corrosion causing oilfield bacteria helps
control plugging, Oil and Gas Journal 80, 1982, pp. 253-264.
P.F. Sanders and L. Latifi, On-site evaluation of organic biocides for cost-effective
control of sessile bacteria, In Proceedings of the Second International Conference on
Chemistry in Industry, American Chemical Society, Vol. I, Paper O-30, 1994, pp
242-257.
163
Oil Field Corrosion
10. Review of current practices for monitoring bacterial growth in oilfield systems.
CCEJV 1987. Document number 001/87. Corrosion Control Engineering Joint
Venture, CCEJV, UK.
11. J.W. Costerton and G.G. Geesey, Microbial contamination of surfaces, In Surface
Contamination (1),. K.L. Mittal Ed., Plenum Pub., 1979, pp. 211-222.
12. P.F. Sanders, Rapid methods for detecting microbial corrosion, In Proceedings of UK
Corrosion ’92, Volume 3, Session D. Institute of Corrosion, 1992.
13. Review of current practices for monitoring bacterial growth in oilfield systems, 1987,
Document No. 001/87, Corrosion Control Engineering Joint Venture CCEJV/NACE.
14. P.F. Sanders, Control of microbiologically induced corrosion using field and
laboratory methods. International Biodeterioration 24, 1988, pp. 239-246.
15. P.F. Sanders and J.F.D. Stott, Assessment, monitoring and control of microbiological
corrosion hazards in offshore oil production systems. NACE Corrosion ‘87. Paper
No. 367, 1987.
164
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
HYDROGEN DEGRADATION OF STEEL - DIFFUSION AND
DETERIORATION
M. Farzam
Faculty of Petroleum Eng.
University of Petroleum Industry, Ahwaz, Iran
ABSTRACT
Following aqueous surface electrochemical reactions and hydrogen reduction, atomic hydrogen will
diffuse through steel, to reside at the microstructural defects with possible catastrophic failures to follow.
Carbon steel, low alloy steel or high strength low alloy steel under no-load, static or cyclic loading will be
affected by the presence of hydrogen. Such an effect is mainly influnced by temperature, alloying,
residual and applied stress, and partial pressure of hydrogen.
Permeation tests were conducted by Devanathan's electrochemical method and a newly invented
electro-vacuum method. Experiments illustrated that as the microstructure changed the diffusion constant,
D, changed. Ingress of hydrogen increased with the reduction of voltage and pH. Thickness would not
have a realistic effect, but cold work reduced D. The difference in D measured by the two methods of
permeation was due to the variation in material, apparatus design and operation.
Hydrogen degradation during stress corrosion and corrosion fatigue tests is one of the two processes
of crack propogation. Immersion, dynamic polarization, stress corrosion (dynamic and static), corrosion
fatigue tests, cathodic protection and fractographic studies were conducted in seawater and sour (hydrogen
sulfide generated by sulfate reducing bacteria) environments. In all, it was concluded that with regards to
Nernest's equation, as the partial pressure of hydrogen increased, the hydrogen concentration gradient
(Fick's first law) increased and D increased. Furthermore, the speed of hydrogen diffusion increased. In
the sour environment, the overtaking mechanism of failure was found to be hydrogen embrittlement rather
than anodic dissolution.
Key Words:
Hydrogen degradation, hydrogen diffusion, diffusion constant, corrosion
fatigue, stress corrosion cracking
INTRODUCTION
The sources of hydrogen production in the industry are many:
water dissociation
+
H 2 O → H + OH
(1)
Fe-H2O reaction
2Fe+ 3 H 2 O
(2)
effect of pH and bacteria
+
H +e → H
(3)
165
Oil Field Corrosion
After atomic hydrogen production at the cathodic side of a corrosion cell, hydrogen will diffuse
into steel residing at the lattice interstitial sites, grain boundaries, vacancies, impurities and
alloying elements. The bonding energy of such interactions is different (see Fig. 1 and Table 1).
Trap sites are either low energy reversible (e.g., interstitial solutes 3-15 KJ/mol [1]) or high
energy irreversible (e.g., TiC particles 96 KJ/mol [1])
It is believed that steels with high energy traps will release hydrogen at about 2500C and at
higher temperatures. Generally, hydrogen will not saturate and therefore is less likely to damage
the austenitic stainless steels, but carbon steel, and ferritic or martensitic steels will be degraded,
and finally cracked.
Table 1. Trap Energy [1]
Trap
Interstitial Solutes (C, N)
Si Atom
Ti Atom
Vacancy
Y Vacancy
Dislocation Elastic Stress Field
Dislocation Core (Screw)
Dislocation Core (Mixed)
1/2 H2 (Vapour /Liquid)
Grain Boundary
Free Surface
AlN Interface
Fe3C Interface
TiC Interface
EB (KJ/mol)
NT (m-3)
~3-15
>20
26
46
126
20
20 - 30
59
29
~59
70 - 95
65
84
96
1025
1027
1027
<1023
1023
19
10 - 1026
1019 - 1026
1019 - 1026
--19
10 - 1023
1021
24
10 - 1025
1024 - 1025
1024 - 1025
B
Figure 1. Lattice trapping [1]
166
Farzam
Hydrogen degradation will appear in one of the following forms:
(1)
(2)
(3)
Hydrogen embrittlement,
Hydrogen blistering,
Hydrogen induced cracking.
As little as H<1 ppm is enough to damage steel. Figure 2 shows a 20"-oil line pipe
(Cheshmeh Khosh) damaged by hydrogen blistering.
Figure 2. Hydrogen blistering (Cheshmeh Khosh)
Taking into account the tensile strength of steel, the H2 gas pressure in the blister must have
been about 104 atm.
There are number of theories describing the mechanisms of H degradation:
(1)
(2)
(3)
(4)
(5)
(6)
Pressure theory,
Surface energy,
Enhanced plastic flow,
Transport model,
Hydride formation,
Decohesion theory.
Since the diffusion of hydrogen into austenitic stainless steel will not increase the internal
gas pressure, the pressure theory cannot be generally applicable. As far as the surface energy
theory is concerned, as the oxygen absorption at the surface will not reduce the surface energy
(with higher adsorption energy than H), this theory may not always be applicable. The third
theory, which discusses the enhancement of plastic deformation in the presence of H, has rarely
been approved and reported. The fourth theory describes hydrogen transportation by dislocations,
but in the absence of dislocations, hydrogen will also diffuse and be transported. There is an
interaction of hydrogen with alloys such as Nb, e.g., Nb-H which is a brittle compound (the fifth
theory). Such an effect cannot also be a prime factor of degradation.
167
Oil Field Corrosion
It is generally believed that the last theory; decohesion theory, is the major responsible
factor weakening the Fe-Fe bonds, and that a hydrogen failure is a result of the participation of all
six theories.
The temperature at which hydrogen will inflict its severest damage is at about 25oC [2].
Figure 3 shows that the crack advance at this temperature is maximum. The figure also shows
that the higher the steel’s strength, the faster the crack propagation rate. The effect of hydrogen
when coupled with static and cyclic stresses will be reflected as stress corrosion, corrosion
fatigue or a combination of the two, i.e., stress corrosion fatigue [4].
The hydrogen diffusion characteristics previously reported by Devanathan and Stachurski
[5] (Fig.4) showed that it takes an incubation period tb before hydrogen reaches the exit surface.
The amount of hydrogen reaching the exit side increases and reaches a steady-state with time.
Time tb and the area under the curve are representative of the metal’s characteristics.
Figure 3. Crack advance vs temperature in AISI 4130 [2]
Figure 4. Hydrogen diffusion curve [5]
168
Farzam
Hydrogen diffusion obeys Fick's first law:
J = -D
dc
dx
(4)
where J is the flux in mol/cm2S, D is the diffusion constant in cm2/S, C is the concentration in
mol/cm3 and X is the distance hydrogen travels in cm.
D = D0 exp
Q
RT
(5)
where D is a constant, Q is the activation energy in J/mol, R is the gas constant in J/mol K and T
is the temperature in K.
It has been shown experimentally that:
2
D=
L
6t lag
(6)
where L is the thickness in cm and tlag = 0.63 J steady-state. From Fig. 4, tlag is measured at the
intersection of the vertical line with the tangent of the slope at 0.63PS.S.. Some researchers [6]
have used the 0.83PS.S.. Devanathan and colleagues [5] showed that as the voltage decreased, the
amount of hydrogen diffused increased and D was not a function of thickness. Detailed research
has been conducted on the effect of microstructure on hydrogen diffusion [7].
The concluding remarks on such an effect are that the structural inhomogeneities; line
defects, impurities (TiC) or martensitic laths, will trap hydrogen and degrade the microstructure.
The lower the hydrogen-defect bonding energy and the less such microstructural
inhomogeneities, the less the amount of trapped hydrogen and the less the likelihood of hydrogen
degradation. Table 1 shows the variation of such energies with defects. Other researchers have
reported their findings on the increase of the number of dislocations and therefore the amount of
trapped hydrogen and trapping energies [8,9]. Having a homogeneous structure reduces the risk
of hydrogen damage. Variation of D with microstructure and alloying is definite, and it may be
stated that the lower the D, the lower the amount of diffused hydrogen [10]. Therefore, if the
amount of trapped hydrogen is lowered, the possibility of hydrogen damage is decreased. It may
be suggested that by controlling the microstructure, hydrogen degradation is controlled.
Considerable research has been previously conducted on the fractography of fractured surfaces.
In a stress corrosion cracking study in H2S, Huang and Shaw [11] showed that the crack
morphology was quasi-brittle and the observed white markings were considered as a sign of the
presence of hydrogen.
EXPERIMENTAL PROCEDURE
Diffusion measurements have been either conducted by electrochemical charging [5] or a
dry gas technique [8]. In the electrochemical method, a steel sample is situated between an
anodic cell and a cathodic charging cell. When a potential difference between the two sides is
169
Oil Field Corrosion
established due to the presence of hydrogen on the cathodic side, H diffusion takes place and
after awhile is detected at the anodic side. Hydrogen diffusion increases with time reaching a
maximum of a steady-state amount. As the negative charging potential is increased, the PS.S. is
increased (Nernest equ.).
This article represents more than ten years of reseach, beginning at Sheffield University
(U.K.) then at Heriot-Watt University (U.K.), and then in Iran. The electrochemical technique
was used at Sheffield, while a newly invented method which is a cross technique between the
electrochemical and vacuum (gaseous) techniques was employed in Scotland. The results were
then correlated with stress corrosion, corrosion fatigue, and case studies. Figures 5 and 6 show
the apparatus used in the present investigation.
Figure 5. Devanathan cell construction used for elecrochemical testing
170
Farzam
Figure 6. Apparatus used for electro-vacuum testing
The electro-vacuum testing used a cathodic electrochemical cell to charge the hydrogen into
steel, and a vacuum (10-7 torr) side leading to a mass-spectrometer to detect the amount of
hydrogen diffused through, using a computer. Table 2 gives details of the alloying elements of
the steels tested by the electrochemical method. The API 5LX-65 was used to transport liquid
gas. Hydrogen-induced cracking was observed in this steel during service, and cracking occured
at positions where the microstructure due to the segregation of Mn (= 2%), was martensitic. The
steel’s structure was similar to BS4360 (Table 2).
Table 2. Analysis of Steels Used for Hydrogen Permeation Test (Devanathan)
Material
API X-65 (Low S)
API X-65 (High S)
BS 4360
BS 4360 (Quenched)
C
Mn
Si
0.079
0.089
0.13
0.13
1.33
1.38
1.57
2.23
0.29
0.31
0.32
0.31
S
0.004
0.031
0.007
0.01
P
0.006
0.005
0.008
0.008
Cr
Ni
0.11
-
0.15
-
The steel samples were Pd-plated and tested in caustic soda (0.1N NaOH) at potentials
lower than -1300 mV. During these tests, the effects of thickness, voltage and plastic
deformation (after 72% R.A) on variation in D were recorded. Table 3 shows some of the results
obtained.
Table 3. Results of Devanathan Test (0.1 N NaOH, 2 mm thick)
Material
API X-65 (Low S)
API X-65 (High S)
BS 4360
BS 4360 (Quenched)
tlag (min)
D (cm2/S X 10-6 )
68.75
77.5
69
253
1.61
1.43
1.61
0.44
Table 4 gives details of the alloys tested by the electro-vacuum method.
Table 4. Steel Analysis Used for Hydrogen Permeation Test (Electro-Vacuum)
Material
API X-52
AISI 4340
C
Mn
Si
S
P
Cr
Ni
0.14
0.38
1.33
0.73
0.33
0.2
0.003
0.016
0.008
0.13
0.05
1.08
0.03
1.38
171
Oil Field Corrosion
API 5LX - 52 steel may also be used to carry liquid gas or other petroleum products, and
AISI 4340 is used as a well casing material. Table 5 shows the results obtained in 0.1N NaOH
for these latter two steels.
Table 5. Results of Electro-Vacuum Test (0.1N NaOH, 2 mm Thick)
Material
tlag (min)
D (cm2/S X 10-7)
242
541
4.58
2.05
API X-52
AISI 4340
For the reason of an actual simulation, the steel samples were not plated in the latter tests.
During these tests, the effects of environmental change, thickness and voltage were examined
(see Table 6).
Two cases of fractured surfaces were compared with one another to find footprints of
hydrogen.
1. A case of stress corrosion fatigue of high strength steel cables that failed in H2S
environment,
2. Another case of stress corrosion fatigue of N.I.O.C drill pipes.
It is worth mentioning that other stress corrosion and corrosion fatigue tests were conducted
that are not all stated here. The stress corrosion tests were conducted either statically (80%Syield)
or dynamically (slow strain rate = 10-5 S-1). The corrosion fatigue tests were conducted at 0.167
Hz frequency and constant stress range. These latter tests were conducted on AISI l055, 1075,
and 50B60 in putrid seawater (H2S = 200 ppm) [14]. As the emphasis here is on the kinetics of
hydrogen diffusion rather than on fracture morphology, only a few fractographs are presented in
this paper. Some of the linear polarization, potential and pH measurements conducted on the
cables show the rate and possible change in the types of reactions. Immersion tests were also
conducted in the sour environment.
Table 6. Results of Electro-Vacuum Test (0.5mm)
Material
AISI 4340
AISI 4340
AISI 4340
API X-52
API X-52
API X-52
API X-52
Inh. : Inhibitor
H2S : Seawater + H2S
172
Environment
tlag (min)
D (cm2/S X 10-7)
0.1 N NaOH
Seawater
Seawater + Inh.
0.1 N NaOH
Seawater
H2S
H2S + Inh.
82
62
56
90
62
28
42.5
0.84
1.12
1.24
0.77
1.12
2.48
1.63
Farzam
RESULTS AND DISCUSSION
Diffusion tests conducted by the Devanthan's electrochemical method (Table 3) showed that
with the increase in sulphur in X-65, D, the diffusion coefficient decreased from 1.61 X 10-6 to
1.43 X 10-6 cm2/S. This was due to the increase in MnS surface area. MnS would act as a trap
impeding hydrogen transport. The measured D for BS 4360 was similar in value to that for X-65
and was equal to 1.61 X 10-6 cm2/S.
Alloy BS 4360 with higher Mn when quenched showed an increased tlag and a decreased D
of 0.44 X 10-6 cm2/S (Table 3). This result was attributed to the presence of fine laths of
martensite which is saturated in carbon and is an internaly stressed structure (shear-induced
transformation). Figure 7 compares the behaviors of X-65 and the quenched 4360. The figure
shows 4360 to have the longer tb.
Recharging the test sample (second transient) after the first transient produced a shorter tb
and increased D. This was ascribed to the saturation of the traps during the first transient.
Lowering the steady-state flux was a sign of reduction in dc/dx according to Fick's first law.
However, this may change (i.e., increase the steady-state) for a different alloying element. When
BS 4360 was cold worked (72% reduction in area, 0.5 mm thickness), tb increased from 6 to 112
min, and D decreased from 4.62 x 10-6 to 5.34 x 10-7 cm2/S. This was due to increased
dislocation density (106 to 1012 cm/cm3) acting as strong traps. Vicker's hardness showed an
increase from 183 to 272 HV.
When the charging voltage was decreased from -1300 to -1700 mV; D increased from 4.3 x
10 to 6.9 x 10-6 cm2/S. This may be explained according to the Nernest's equation (increase in
hydrogen partial pressure):
6
Figure 7. Hydrogen diffusion measurements conducted by Devanathan’s method (2mm thick) (i)
X-65 (low S); (ii) X-65 (high S); (iii) BS 4360; and (iv) BS 4360 (Quenched)
-E = - E 0 +
RT
ln P H 2 - constant PH
ZF
(7)
As the thickness was decreased from 2 to 0.5 mm, D increased from 1.6 x 10-6 to 4.3 x 10-6
cm2/S. This change was only due to the kinetics of the electrode’s surface. The electro-vacuum
173
Oil Field Corrosion
tests conducted on API X-52 and AISI 4340 (Fig. 8) were made without Pd coating of the test
samples.
Published research has shown that the volume of hydrogen diffusing through would increase
with the increase in surface scratches as a result of the increase in surface area [12]. Thus, during
the present series of tests, all the samples were polished with 600 grade SiC paper and used
uncoated. Other investigators reported unsuccesful test results if the samples were not Pd-coated
[13]. In the present work, the measured D for X-52 was twice that of 4340. A comparison of the
alloying elements of the two metals indicates that AISI 4340 has a higher amount of C, Cr, Ni
and S. It is generally believed that elements on the right-hand side of Fe in the periodic table (C,
Ni, etc.) repel hydrogen while elements on the left-hand side of Fe (Cr, Nb, etc.) attract hydrogen.
A comparison of D values for the electro-vacuum tests with those of the electrochemical
experiments reveals that D was 10 times larger (faster) for the electrochemical tests. This is most
likely due to the surface impedance of the scales (absence of Pd) in addition to variations
resulting from the electro-vacuum test method. Table 6 shows that D changed with the changes
in the test environment. Testing AISI 4340 in 0.1N NaOH provided D = 0.84 x 10-7 cm2/s, which
is similar to that for API X-52. But when X-52 was tested in H2S, X-52, D increased from 0.77
to 2.48 x 10-7 cm3/S. Linear polarization tests conducted in H2S on AISI 1074 showed that FeS
resulted in cathodic depolarization (see Fig. 9). Thus, with the increase in cathodic reaction, the
amount of hydrogen reduction and diffusion would increase. This matches the results obtained
during diffusion (Table 6).
Figure 8. Hydrogen diffusion measurements by the electro-vacuum testing technique
174
Farzam
Figure 9. Linear polarization (cathodic side) in: (1) seawater + H2S = 300 ppm; and (b)
3.5% NaCl solution
The free corrosion potential of a steel in the presence of H2S will change with time. Figure
10 shows that with time (after 219 h) the potential of AISI 1055 increased from -720 to -560 mV.
This may be due to the change in FeS stoichiometry, perhaps to FeS2. When the corroded surface
was cleaned, it was noted that pits had formed under the black FeS layer (see Fig. 11).
Figure 10. Changes in Ecorr in different H2S concentrations for AISI 1055
175
Oil Field Corrosion
Figure 11. Pitting corrosion of AISI 1055 in H2S
Corrosion fatigue and stress corrosion tests were conducted on wires and wire-ropes (AISI
1055, 1074, and 50B60) in an H2S environment. In the fatigue tests, the number of cycles to
failure decreased as the maximum stress increased. This was considered as a sign of stress
corrosion fatigue failure. Cracks were noted to originate from the corrosion-pits formed under
the perforated FeS layer (Fig. 12). Thus, one can say that the crack initiation mechanism is
anodic dissolution.
Figure 12. Initiation of microcracks from surface pits (anodic dissolution)
176
Farzam
White markings as a result of hydrogen embrittlement were observed on the fractured
surface (Fig. 13). Thus, hydrogen embrittlement could be the mechanism of crack propagation.
Figure 13. Signs of hydrogen embrittlement (white markings)
Drill pipe failure (N.I.O.C) showed that in the alkaline environment of the drill mud (pH =
10), fracture was due to a combination of tensile and torsional cyclic stresses. The fractured
surface showed a thumbnail fatigue fracture with surface scratches and microcracks acting as the
initiating sites. Thus, the initiation and propagation mechanisms seem to be via anodic
dissolution. However, the presence of cyclic loading may imply stress corrosion fatigue failure.
CONCLUSIONS
1.
2.
3.
4.
5.
6.
With an increase in sulphur content, tLag increased.
Quenched BS 4360 had a D four times slower than pearlitic structure.
D measured by the electro-vacuum method was ten times that by Devanathan.
In the presence of H2S, D increased.
In the presence of inhibitors, D did not decrease.
Reduction in voltage increased D.
REFERENCES
1. R. Gibala and A.J. Kumnick, Hydrogen trapping in iron and steels, in: Hydrogen
Embrittlement and Stress Corrosion Cracking, Ed., R. Gibala and R.F. Hehemann, ASM,
1984, p. 61.
2. H.H. Johnson, Overview on hydrogen degradation phenomena, in: Hydrogen
Embrittlement and Stress Corrosion Cracking, Ed., R. Gibala and R.F. Hehemann, ASM,
1984, p. 3
3. F.P. Ford, Stress Corrosion Cracking, in Corrosion Processes, Ed., R.N. Parkins, Applied
Science Publishers, 1982, p. 271.
4. J. Congleton and I.H. Craig, Corrosion Fatigue, Applied Science Publishers, 1982, p. 209.
177
Oil Field Corrosion
5. M.A.V. Devanathan and Z. Stachurski, Mechanism of hydrogen evolution on iron in acid
solutions by determination of permeation rates, Journal of the Electrochemical Society,
1964, p. 619.
6. J. McBreen, L. Naris and W. Beck, A method for determination of the permeation rate of
hydrogen through metal membranes", Journal of the Electrochemical Society, Nov. 1966,
p. 1220.
7. I.M. Bernstein and A.W. Thompson, The role of microstructure in hydrogen
embrittlement, in: Hydrogen Embrittlement and Stress Corrosion Cracking, Ed., R.
Gibala and R.F. Hehemann, ASM, 1984, p. 135.
8. H.H. Johnson and R. Way Lin, Hydrogen and deutrium in iron, in: Hydrogen Effects in
Metals, Ed., I.M. Bernstein and A.W. Thompson, 1984, ASM, p. 3.
9. K.K. Kim and S. Pyon, Hydrogen permeation through 3.3% Ni-1.6% Cr steel during
plastic deformation, Scripta Metallurgica 22, 1988, p. 1719.
10. C.H. Tseng, W.Y. Wei and J.K. Wu, Electrochemical methods for studying hydrogen
diffusivity, permeability in AISI 420 and AISI 430 stainless steels, Materials Science and
Technology 5, Dec.1989, p. 1236.
11. H. Huang and W.J.D. Shaw, Cold work effects on sulphide stress cracking of pipeline
steel exposed to sour environment, Corrosion Science 34, 1, 1993. p. 61.
12. R.L. Reuben, Ph. D. Thesis, Hydrogen Permeation, Open University, U.K. 1980.
13. I.M. Bernstein and A.W. Thompson, Proc. Mechanisms of Environmental Embrittlement
in Materials Metal Society, 1978, p. 403.
14. M. Farzam, R. Brook and R.G.J. Edyvean, The fracture of steel wire, strand and rope in
marine environments, 31st Corrosison Science Symposium, 11-14 Sept. 1990, Newcastle,
England.
178
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
CONTROL STRATEGIES FOR THERMOPHILIC SULPHATE-REDUCING
BACTERIA
P.F. Sanders1, H.M. Lappin-Scott2 and C.J. Bass2
1
2
Oil Plus Limited, Hambridge Road, Newbury, Berkshire, RG14 5TR, England
University of Exeter, Hatherly Building, Department of Biological Sciences, Prince of Wales
Road, Exeter, EX4 4PS, England
ABSTRACT
Thermophilic (high-temperature requiring) sulphate-reducing bacteria (tSRB) are responsible for
both corrosion of oil production facilities and souring of produced oil. During oil recovery, a range of
chemicals is injected into oil reservoirs for various operational reasons. Such chemicals may be
designed to inhibit microbial growth, or to specifically inhibit general corrosion. Other chemicals,
such as scale inhibitors, oxygen scavengers, H2S scavengers and surfactants, may be also injected. To
date, there is little published data on the effect of any of these chemicals on the micro-flora that is
known to exist in the oil formation.
The effect on sulphide production using a tSRB isolated from an oil bearing chalk formation was
monitored in the presence of rock surfaces and individual chemicals. The presence of a surface
increased sulphide production, whilst oilfield chemicals produced a variety of responses, from
complete inhibition to significant enhancement of sulphide production. Differences were observed
within the various classes of treatment chemicals. Some molecules appeared to break down more
easily than others under reservoir conditions and thus served to stimulate sulphide production. The
data obtained will help to manage oil reservoirs in a more cost-effective and environmentally friendly
manner, minimizing both sulphide production and microbial corrosion by determining appropriate
chemical selection and dosing regimes.
Key Words:
Thermophilic sulphate-reducing bacteria, corrosion of oil production
facilities, oilfield chemicals, management of oil reservoirs
INTRODUCTION
The production of oil worldwide is adversely affected by the growth of bacteria,
particularly the sulphate-reducing bacteria (SRB). These may be mesophilic (mSRB),
growing optimally at 30°C, or thermophilic (tSRB), growing optimally at 60°C. Although
SRB are encountered in many anaerobic environments, their principal habitat is marine
situations, either in sediments or suspended in seawater, where sulphate concentrations are
high. SRB activity enhances corrosion of iron and steel in storage tanks, pipelines and
equipment as well as creating problems of toxicity and spoilage of the gas and oil. In the
reservoir itself, permeability reductions may be caused by growth of bacteria in the pore
spaces of the rock leading to reduced recovery efficiency [1-4]. Most attention has been paid
to mSRB growing in seawater treatment facilities, where biofilms containing mSRB are
responsible for enhancing the corrosion of steel [5]. Recently tSRB have been isolated from
179
Oil Field Corrosion
oilfield reservoirs and production systems [6-8], with significant implications for maintenance
and integrity of hot oil production pipework and associated plants [9]. tSRB have also been
detected in injection water and open seawater [7,10], together with other thermophilic
heterotrophic bacteria, indicating the presence of widespread consortia of thermophilic
bacteria in these environments. Temperatures may reach 120°C in some oil-bearing
formations [11], but are typically 80-100°C. tSRB have been shown to grow or survive in
this range of temperature [4,7].
Actively growing tSRB readily attach to surfaces, assisted by the production of large
amounts of exopolysaccharide [12] and survive adverse conditions by the formation of
dormant cells which have been shown to retain viability for long periods of time [8]. These
dormant cells may pass through the reservoir rock since they are small and produce less
exopolysaccharide. They thus contaminate both the deep reservoir and the production
system. In some situations, even starved cells may attach to surfaces and form part of a
developing biofilm [13]. Biofilms containing tSRB and other thermophilic bacteria can thus
develop on rock surfaces as well as on metal surfaces. When such a biofilm develops on
exposed steel, corrosion can be initiated by the production of sulphide or organic acids [14];
when it forms within rocks, permeability reduction can occur.
Many microorganisms are able to survive wide and rapid variations in environmental
conditions. Renewed interest in the ability of microorganisms to survive low nutrient status
has prompted research into bacterial starvation responses [15]. Oil/water mixes produced
from reservoirs frequently contain tSRB as they enter topside facilities. Nutrient availability
in such situations will vary with the origin of the produced water; therefore, any tSRB in
these locations will be responding directly to nutrient stress as well as temperature shifts.
Some cells will adopt a starvation/survival mode which may take the form of reduction in cell
size with reduced levels of metabolic activity [7]. Additionally, cells adhere to available
surfaces [10] and subsequently form biofilms in topside production systems. The ability of
tSRB to switch between different growth states, depending upon the environmental
conditions, allowing the bacteria to survive extreme or adverse conditions has been described
[11]. This, in turn, means that reservoir sulphide souring and corrosion in high temperature
systems will be encouraged if a change to more favourable conditions occurs.
Biofilms (consisting of vegetative, active, tSRB and other thermophilic anaerobic
bacteria) developing in hot oil production systems will give rise to increased sulphide
production, which may eventually lead to significant souring of the produced fluids [17]. In
addition, the activity of thermophilic bacteria on hot metal surfaces will lead to an increase in
microbially influenced corrosion (MIC) due to the local production of corrosive metabolites
such as sulphide and organic acids [9,18,19].
The studies described in this paper are part of an ongoing investigation into the survival
and activity of tSRB in oilfield reservoirs. A wide range of tSRB was enriched and isolated
from a series of surveys on and around North Sea installations. Samples included raw
seawater, injection water, well-head samples, production separators and overboard discharge
lines at three major oil field production/injection facilities. These isolates were used to
conduct a series of laboratory trials related to the control of reservoir souring and associated
corrosion implications. Particular emphasis was placed upon the relative importance of
starved, dormant (non-sulphide-producing) growth states, the active (sulphide producing)
180
Sanders et al.
vegetative state, and the availability of surfaces and chemical treatments on sulphide
generation rates.
METHODS
Thermophilic Sulphate-Reducing Bacteria Cultures
Two oilfields were selected for the isolation of high-temperature, hydrogen sulphideproducing bacterial enrichments. Both fields had fractured chalk reservoir rock, and were
experiencing seawater breakthrough and increasing hydrogen sulphide concentrations in the
produced fluids. Two consortia (EX251 and EX258) from hot oil producing systems were
studied for their tolerances to chemicals and reactions to surfaces.
Growth Media
Three anaerobic media selective for SRB were prepared: MCP, EX2 and SWM.
MCP is a modification of Postgate's E. It contained Na lactate (70% w/w solution 12.5 ml 11
), Na acetate (50 g l-1), Na propionate (1.0 g l-1), and Na butyrate (0.4 g l-1) with NaCl (25 gl1
). EX2 contained KH2PO4 (0.4 g l-1), NH4Cl (1.0 g l-1), CaCl2 (0.1 g l-1) MgSO4.7H2O (2.0 g
l-1) FeSO4.7H2O (0.5 g l-1), yeast extract (1.0 g l-1), Na2SO4.2H2O (1.0 g l-1), NaHCO3 (2.4 g l1
), Na pyruvate (6.0 g l-1) and NaCl (25.0 g l-1). SWM was used when a clear medium was
required. Na pyruvate (6.0 g l-1) and yeast extract (oxoid, 1.0 g l-1) was added to one liter of
artificial seawater solution (ASW). After these additions, the prepared medium was sterilized
by filtration (with Whatman's presterilized cellulose acetate filter, 0.45 mm pore size) and
dispensed in presterilized bottles in an anaerobic cabinet.
Spectrophotometric Sulphide Assay
The micro method described here is a modification of those described by Truper and
Schlegel [20] and relies on the liberation of sulphide by acidification and subsequent
development of methylene blue from n,n-dimethyl-p-phenylene diamine sulphate.
Influence of Surface Area on Sulphide Production
This experiment was designed to determine whether the area of available surface in a
culture vessel exerted an influence on the amount and rate of sulphide production of a culture
of tSRB cells. Culture EX251 was selected for this experiment. Several sets of four injection
vials were set up containing 0 (control), 0.01, 0.1 and 1 g of washed Bunter sand with 9 ml of
SWM. Each vial was inoculated with one milliliter of a three day old culture of EX251 and
incubated at 60°C. At predetermined sample times, three vials were removed from each
treatment set and sampled for sulphide production and metabolic activity.
Sulphide Production from EX251 in Long Term Contact with Oil Formation Rock
Chalk core chips were distributed in 15 g quantities in a series of presterilized, loosely
stoppered injection bottles and placed in an oven at 110°C for 36 hours. The bottles were
placed in an anaerobic cabinet overnight to deoxygenate the rock pore spaces. EX251 was
inoculated as follows: sterile, full strength SWM, 1/10 strength SWM (diluted with ASW)
and ASW. 1/10 strength SWM was added to the uninoculated control. Each bottle was crimp
sealed to maintain the anaerobic status of the contents and then incubated at 60°C. Samples
were withdrawn at intervals over several months for sulphide assay.
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Oil Field Corrosion
Treatment Chemicals and Sulphide Production
To determine whether the tSRB utilize a chemical additive, FWM was used. The low
nutrient level medium, FWM/8, was prepared by making up 125 ml of FWM to one litre with
ASW. Enhanced sulphide production, as compared with an untreated control, indicates the
utilization of the chemical by tSRB. Vials containing a two milliliter standard scoop of silica
sand (BDH, 100 mesh) were filled with either high or low nutrient medium (FWM or
FWM/8) in the anaerobic cabinet. Vials were stoppered, capped and autoclaved at 115°C for
10 minutes prior to use. Also, 90 ml of fresh C2 medium was inoculated with five milliliter
of a one-week old culture of EX258 grown in C2 medium and incubated at 60°C for three
days.
Chemicals were suitably diluted in either high or low nutrient medium to enable a one
milliliter addition to the appropriate individual test vials at the start of the test. On the first
day of the test, one milliliter of a three day old EX258 culture was injected into each vial.
This was followed by a one milliliter addition of each test chemical to the appropriate set of
vials. Each set was sampled for sulphide as soon as the chemical inoculations were complete.
The remaining vials were placed in the 60°C incubator for the remainder of the test. At
predetermined time intervals three vials were randomly selected from each chemical
treatment set for sulphide analysis as described above.
RESULTS AND DISCUSSION
Sulphide Production by Thermophilic Sulphate-Reducing Bacteria with Varying
Amounts of Available Surface
Figure 1 shows the rate and extent of sulphide production from one consortium
inoculated into vials containing defined amounts of crushed Bunter sandstone as a surface for
attachment and growth. During the first 24 hours, sulphide concentrations in all four
treatments (0, 0.01, 0.1 and 1 g sand) increased to between 30 and 40 ppm, with the least
sulphide assayed from the vials containing the most sand (i.e., one gram). However, after a
further 24 hours of incubation, the situation was reversed, with those vials containing the
most additional surface area yielding the most sulphide (35 ppm). The rest of the range of
experimental vials showed a range of sulphide concentrations in accordance with the
available surface area. This range was further emphasized at 72 and 96 hours. Sulphide
concentration appeared to be directly linked to the available surface area incorporated in the
experimental vials. Up to 30 hours of incubation, there was no significant difference between
the means. At 54 hours, the mean of data from the vials incorporating one gram of sand was
found to be significantly different from the other treatments and the control. By 102 hours of
incubation, there was a more significant difference, with the means of each of the sand
treatments being significantly different from the control. Thus, increasing the amount of
available surface significantly increased the amount of sulphide generated.
Sulphide Production by Thermophilic Sulphate-Reducing Bacteria in Long Term
Contact with Oil Formation Rock
The previously white chips of oil-bearing chalk which were inoculated with ten milliliter
of SWM grown culture of EX251 (containing little precipitated material) blackened overnight
in an anaerobic cabinet before the experimental nutrient additions were made. The colour
change was retained throughout the nine month course of this experiment and was not
182
Sanders et al.
observed in the control bottles containing only chalk chips and SWM/10 medium. The
sulphide production is plotted against time in Fig. 2 for each of the nutrient treatments given
to the EX251 cells in contact with the chalk. In the bottles with SWM at full strength, there
was a rapid increase in sulphide concentrations in the first two days, similar to the previous
experiment. This was followed by further increases over the ensuring four weeks until the
maximum assayed sulphide was 100 ppm. The concentration of sulphide then steadily
diminished to 55 ppm over a period of six months. Extra nutrient was added at the start of
week 34, and sulphide measurements continued. Due to dilution caused by the added
nutrient, the sulphide concentrations were reduced immediately after the fresh nutrient was
added; however, within two days, the SWM cultures showed renewed sulphide production
which peaked at 60 ppm within ten days of the extra nutrient being supplied.
Figure 1. Sulphide produced from EX251, grown in the presence of 0.01, 0.1
or 1 g sand compared with a control containing no additional surface. n = 3;
all error bars displayed (+/- SE of mean). At 54 h, p < 0.05 for 1 g sand
treatment and by 102 hours all sand treatments were significantly different
from the control (p < 0.01)
When SWM/10 was added to the bottles, there was a discernible increase in sulphide
output in the first three to four weeks; the maximum sulphide concentration achieved was 20
ppm and thereafter remained at about 15 ppm until the nutrient addition at week 38. After an
initial drop in assayed sulphide to 10 ppm, the sulphide concentration again rose in the
following ten days to 15 ppm, and as in the first part of the experiment, gradually declined.
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Oil Field Corrosion
In the experiments where cultures were offered ASW only, low levels of sulphide (c. 10
ppm) were observed when in contact with the chalk chips until the 36th week. The final two
readings, taken at 37 and 38 weeks, showed a reduction in assayed sulphide to 6 and 3 ppm
respectively. No sulphide was detected in the control series without added bacterial culture.
Figure 2. Long term survival and sulphide generation from cells of EX251
on oil-bearing chalk chips. Sulphide produced consistently available
nutrient at start of experiment. After further addition of nutrient at 34
weeks, renewed sulphide generation occurred. Results plotted are the
average of duplicate readings from two separate experimental vessels in
each case
Both cultures EX251 and EX258 secreted extracellular polysaccharide which aids
attachment to surfaces [12,21]. The presence of exopolysaccharide in corrosive biofilms on
metal surfaces containing SRB from Brazilian oilfield waters has been demonstrated by
ruthenium red staining techniques followed by electron microscopy [22]. Similar procedures
used in the authors’ laboratory have demonstrated the presence exopolysaccharide secretion
by tSRB.
The implication of the results is that thermophilic reservoir SRB reduce sulphate more
efficiently when attached to a surface such as sandstone, chalk or metals. Work by
Laanbrook and Geerlings [23] demonstrated an increase in sulphide production by SRB
cultures when clay particles were incorporated in the growth media. Other studies, on nonsubterranean mSRB taken from paddy fields and lake sediments, have suggested a
physiological advantage for bacteria attached to particulate material in sulphate deficient
situations [24]. It is also possible that attached cells were disadvantaged due to a reduction of
cell surface area available for solute transport at the physical points of contact in non-sulphide
limited conditions. This conclusion does not agree with other researchers [14, 25] who have
consistently reported the advantageous properties of commensal biofilm habitats. However, it
is clear from the work described here that there are factors affecting the metabolism of
sulphide-generating organisms in biofilms which have yet to be fully examined.
184
Sanders et al.
Effects of Treatment Chemicals on Thermophilic Sulphate-Reducing Bacteria
A wide range of treatment chemicals is employed in the oil production process. These
chemicals are designed to prevent scale formation (scale inhibitors), prevent corrosion
(corrosion inhibitors), and remove trace amounts of oxygen from the water (oxygen
scavengers), as well as the biocides which are designed to kill bacteria in the injection water
and the near wellbore area. The use of organic and inorganic chemicals could have a
significant impact upon the growth of bacteria in the system. A test method was devised
which assesses the impact of the chemicals on tSRB in different growth states. A wide range
of treatment chemicals was tested, as shown in Table 1. Future testing will make use of sandpod and cross-flow rock core flood equipment. Typical sulphide production curves from some
of the chemicals are shown in Figs. 3-18.
Sulphide production by EX258 was affected in a variety of ways by treatment with the
chemicals tested. In some instances there was a clear indication of a detrimental effect (i.e.,
enhanced sulphide production in comparison with the control). In others, the effect was to
reduce the amount of sulphide produced (i.e., an effective control of sulphide). In some cases
there was a switch in the effect noted, most frequently from an initially beneficial response to
a detrimental response. This may be due to degradation of the applied chemical or adaptation
of the tSRB to its presence.
Figure 3. Sulphide production by
EX258, treated with THPS at 100
ppm, under high nutrient conditions
Figure 4. Sulphide production by
EX258, treated with THPS at 25
ppm, under high nutrient conditions
Figure 5. Sulphide production by
EX258, treated with THPS at 100
ppm, under low nutrient conditions
Figure 6. Sulphide production by
EX258, treated with THPS at 25
ppm, under low nutrient conditions
185
Oil Field Corrosion
Table 1. Generic Types of Oilfield Treatment Chemicals Tested, Together with
Concentrations Typically in-Use and Tested
Injection Water Type
Chemical Function
Scale Inhibitor
Corrosion Inhibitor
Filter Aid
Oxygen Scavenger
Sea or Aquifer Water
Produced Water
Chemical Type and Dose Rate (ppm)
Polyacrylate
5
Polyacrylate
Phosphinocarboxylic acid
20
Polycarboxylate
Polyvinyl sulphonate
20
Phosphonate
Copolymers
20
Ethylene diamine tetra MPA
Organic phosphonate
5
Diethylene-triamine-penta
MPA
Amines
Imidazoline (oil sol)
Substitued alcohols
Amine salts (water sol)
Quaternary amines
5
Carboxylic acids
Polyamine
1
Ferric sulphate
Ferric chloride
Ammonium bisulphite
Sodium bisulphite
Polyacrylamide
Polyglycol
Phenol formaldehyde
alkoxylate
Polyol
Polyamine
Oxyalk phenolic + PAG
Oxyalk phenolic
Anthroquinone
QAC
Isothiazolone
THPS
Glutaraldehyde
Antifoam
20
20
20
10
2-10
10
30
20
2-4
Continuous
100-400
100-400
Scale Inhibitor Squeeze
Drag Reducer
Dispersant
20
7
14
Deoiler
Demulsifier
Biocide
20
20
Penta methylene phosphate
PCA
Hexaphosphonate
Sulphonated polymer
Polymer in mineral oil
70
Fatty acids & esters
(intermittent)
Silicones in HC
20
20
Under high nutrient conditions, EX258 consistently produced a maximum of 100-120
ppm sulphide after three days. Under low nutrient conditions, this sulphide production was
limited to 40-50 ppm. THPS-biocide (a quaternary phosphonium compound) fully inhibited
sulphide production at a 100 ppm dose rate under low nutrient conditions (Fig. 5); when high
nutrient conditions prevailed, there was an initial inhibition of sulphide production for four
days (Fig. 3). However, the biocide evidently acted as a biostat since sulphide generation
occurred at day 7. At 25 ppm, this effect was more pronounced, with sulphide production
occurring in both high and low nutrient conditions (Figs. 4 and 6).
186
Sanders et al.
Figure 7. Sulphide production by
EX258, treated with glutaraldehyde
at 100 ppm, under high nutrient
conditions
Figure 8. Sulphide production by EX258,
treated with glutaraldehyde at 25 ppm,
under high nutrient conditions
Figure 9. Sulphide production by
EX258, treated with glutaraldehyde
at 100 ppm, under low nutrient
conditions conditions
Figure 10. Sulphide production by
EX258, treated with glutaraldehyde
at 25 ppm, under low nutrient
conditions
Figure 11. Sulphide production by EX258, Figure 12. Sulphide production by
treated with ammonium bisulphite
EX258, treated with ammonium
(oxygen scavenger) at 7 ppm continuous, bisulphite (oxygen scavenger) at 7 ppm
under high nutrient conditions
continuous, under low nutrient
conditions
187
Oil Field Corrosion
Figure 13. Sulphide production by
EX258, treated with proprietary
surfactant, at 30 ppm slug, under
high nutrient conditions
Figure 14. Sulphide production by
EX258, treated with proprietary
surfactant, at 30 ppm slug, under
low nutrient conditions
Figure 15. Sulphide production by
EX258, treated with phosphonate
corrosion inhibitor, at 20 ppm
continuous, under high nutrient
conditions
Figure 16. Sulphide production by
EX258, treated with phosphonate
corrosion inhibitor, at 20 ppm
continuous, under low nutrient
conditions
Figure 17. Sulphide production by
EX258, treated with QAC corrosion
inhibitor, at 20 ppm continous,
under high nutrient conditions
Figure 18. Sulphide production by
EX258, treated with QAC corrosion
inhibitor, at 25 ppm continuous, under
low nutrient conditions
Glutaraldehyde showed a similar pattern of sulphide generation, but in this case, control
appeared to be less effective, and sulphide concentrations were higher than in the control after
two to three day period (Figs. 7-10). These results emphasize that under-dosing biocides can
lead to stimulation rather than reduction in the souring problem.
No effect on sulphide generation rates was seen when oxygen scavenger was used (Figs.
11-12). Although widely used in seawater injection systems for oxygen control, such
chemicals are thought to be unlikely to stimulate additional sulphide from SRB activity in the
formation.
Some surfactants (used in production systems) were effective as biocides against tSRB
(Figs. 13-14), and the use of produced water containing such compounds would exhibit a
good degree of microbiological control near the wellbore.
Corrosion inhibitors (added in seawater and produced water injection systems) showed a
wide range of effects on sulphide production by tSRB (Figs. 15-18). This is due to the variety
188
Sanders et al.
of active agents used in this class of chemical: phosphonate types appear to enhance sulphide
production, having no biocidal properties, whilst QAC formulations can act as very good
biocides.
The combined effects of chemical dosing can now be assessed with respect to the
potential for enhancing or depressing sulphide production by SRB.
ACKNOWLEDGEMENTS
We wish to thank Agip SpA, Chevron U.K. Ltd, Maersk Olie og Gas A.S., Nalco/Exxon
Energy Chemicals Ltd., the Saudi Arabian Oil Company and the UK Department of Trade
and Industry for their valuable support and help during this study.
REFERENCES
1.
B.N. Herbert and P.D. Gilbert, Isolation and growth of sulphate-reducing bacteria, in:
Microbiological Methods for Environmental Biotechnology, 1984, pp. 235-257.
2. W.P. Iverson and G.J. Olson, Problems related to sulfate-reducing bacteria in the
petroleum industry, in: Atlas R.M. (ed) Petroleum Microbiology, MacMillan
Publishing Company, New York, 1984, pp. 619-641.
3. B.N. Herbert, Reservoir souring. in: Hill Shennan, Watkinson (eds) Microbial
Problems in the Offshore Oil Industry, J Wiley and Sons Ltd, London, 1987, pp. 6372.
4. C.J. Bass, P.F. Sanders and H.M. Lappin-Scott, Starvation and survival of
thermophilic sulphate-reducing bacteria from North Sea oil reservoirs. in: Proc 6th
Int Symp Microbial Ecology, International Committee on Microbial Ecology
(ICOME), 1992, Paper C4-3-3.
5. W. Lee, Z. Lewandowski, P.H. Nielsen and W.A. Hamilton, Role of sulfate-reducing
bacteria in corrosion of mild steel: A review, Biofouling 8, 1995, pp. 165-194.
6. J.T. Rosnes, T. Torsvik and T. Lien, Spore forming thermophilic sulfate-reducing
bacteria isolated from North Sea oilfield waters, Appl Environ Microbiol 57, 1991,
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7. C.J. Bass, H.M. Lappin-Scott and P.F. Sanders, Bacteria that sour reservoirs: New
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Offshore Technol 1, 1993, pp. 31-37.
8. K.O. Stetter, R. Huber, E. Blochl, M. Kurr, R.E. Eden, M. Fielder, H. Cash and I
Vance, Hyperthermophilic archaea are thriving in deep North Sea and Alaskan oil
reservoirs, Nature 365, 1994, pp. 743-745.
9. P.F. Sanders, H.M. Lappin-Scott and C.J. Bass, Corrosion implications of
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10. S. Walsh, H.M. Lappin-Scott, H. Stockdale and B.N Herbert, An assessment of the
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11. M. Magot, C. Hurtevent and J.L. Crolet, Reservoir souring and well souring. in:
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12. C.M.L. Cutinho, F.C. Magalhaes and T.C. Araujo-Jorge, Morphology of the surface
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Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
CORROSION EVALUATION OF AUSTENITIC AND DUPLEX STAINLESS
STEELS IN SIMULATED HYDROGEN SULPHIDE CONTAINING
PETROCHEMICAL ENVIRONMENTS
K. Saarinen1 and E. Hamalainen2
1
2
VTT MANUFACTURING TECHNOLOGY
P.O. Box 1704, FIN-02044 VTT, Finland
Helsinki University of Technology, Laboratory of Engineering Materials
Puumiehenkuja 3, FIN-02150 Espoo, Finland
ABSTRACT
The combined effect of hydrogen sulphide, carbon dioxide and chlorides on the sulphide stress
cracking (SSC) behavior of several austenitic and duplex stainless steels was studied in high temperature
environments by using slow strain rate (SSR) tests and immersion tests of stressed specimens.
The initiation of SSC is considered to require the breakdown of the passive film whereby the
localized corrosion resistance can be related to the SSC resistance and to the pitting resistance equivalent
number (PREN) of an alloy. The evaluation of the usefulness of different PREN values showed that the
SSC resistance of a certain alloy can be evaluated by using the PREN values including the nickel content
of an alloy. According to the results, as the PREN values increase, the SSC resistance increases. In the
environments studied, the critical temperature for cracking of austenitic alloys having a PREN value under
45 was below 150°C, and as the PREN value was over 70 the critical temperature for cracking was higher
than 150°C. The nickel-based alloy having a PREN value of 140 was not susceptible to SSC even at a
temperature of 230°C with high partial pressures for H2S and CO2. When the SSR test was used, brittle
behavior was observed in the duplex alloy having a PREN value of 56, even with small amounts of H2S in
the chloride environment.
Key Words: Alloy steels, stainless steel, corrosion, stress corrosion, cracking (fracturing),
petrochemical environments, oil fields, hydrogen sulphide
INTRODUCTION
Corrosion of steels in sour gas environments is a complex process involving many reactions
between the environment and construction materials producing different kinds of corrosion
products [1]. For passivating alloys, the reduction of sulphur is the most feasible cathodic
reaction that accelerates localized corrosion in cathodically reducible corrosion systems [2]. The
anodic dissolution has been observed to increase and the passivation of corrosion-resistant alloys
has been observed to retard significantly as the hydrogen sulphide content of the environment
increases [3]. Besides hydrogen sulphide, other chemicals appearing in the solution have to be
considered [4].
In sour wells, the oil and natural gas are often contaminated by hydrogen sulphide, the
content of which may vary from a few parts per million up to 30% [5]. In the presence of
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Oil Field Corrosion
hydrogen sulphide, the atomic hydrogen that is produced on the surface of the steel by the
corrosion reactions tends to enter the steel and cause embrittlement. This phenomenon is called
sulphide stress corrosion cracking (SSC). SSC susceptibility is enhanced by chlorides, and the
results obtained by Fliethmann et al. indicate that the cracking mechanism is a mixed mechanism
of hydrogen- and chloride-induced SSC [6]. Besides cracking, considerable general corrosion
can be observed due to presence of carbon dioxide that acidifies the mixture of brine and oil [5].
In order to optimize the SSC behavior of the austenitic and duplex stainless steels, the chemical
composition and the ferrite/austenite ratio have to be considered [7]. Also the microstructure,
hydrogen diffusion coefficients [8] mechanical properties, heat treatment conditions, grain size
and form of nonmetallic inclusions play an important role in the SSC resistance of construction
materials [9,10]. Thus, for reasonable assessment of the applicability of a duplex stainless steel
in a specific sour environment, it is important to consider practical experience and to use suitable
SSC test methods, both regarding the loading procedure and the environment [11,12].
Localized corrosion resistance is related to the chemical composition, and thereby, to the
pitting resistance number (PREN) of an alloy. The initial process of SSC can be the breakdown
of passive film [7,13] whereby the increase in the chromium and molybdenum content of the
alloys increases the passivity of the surface layer. However, it has been pointed out that the
ferrite phase of duplex stainless steels is richer in chromium and molybdenum than the austenitic
phase, and hence, the PREN of the austenitic phase seems to play a more important role in pitting
resistance, and hence, in SSC resistance of duplex stainless steels than the PREN of the ferrite
phase [7]. The SSC susceptibility of duplex stainless steels is also greatly affected by the
ferrite/austenite ratio, distribution and grain size. The threshold stress for SSC in simulated sour
environments has been observed to be highest at the α-content of 40 to 45%, and an alloy with a
α-content exceeding 80% is reported to have a high susceptibility to pitting and intergranular
corrosion [14].
To evaluate the SSC behavior of selected austenitic and duplex stainless steels in high
temperature H2S solutions, slow strain rate (SSR) tests and immersion tests of stressed specimens
were used.
EXPERIMENTAL PROCEDURE
Test Materials
The austenitic stainless steels investigated in this study were delivered by Outokumpu
Polarit Oy, and the duplex stainless steels and the two experimental powder metallurgically
produced alloys were delivered by Rauma Materials Technology Oy. The reference materials
selected were Nicrofer 6020, Nicrofer 5716 and Nicrofer 5923; they were and manufactured by
VDM Nickel-Technologies AG and delivered by Cronimo Co. The chemical compositions of
the test materials and their PRENs are presented in Table 1. PREN1, PREN2 and PREN3 reflect
the effect of alloying elements, and PREN number 4 correlates with the ferrite content in a
duplex alloy so that the higher the PREN4 value, the lower the ferrite content.
Test Equipment and Test Environments
The high temperature sulphide stress cracking studies were carried out by using the SSR
method with round tensile test specimens (Fig. 1) and U-bend specimens. To insure the stability
of the test environment during testing, the Materials Research Shuttle Laboratory facility was
used. The facility consists of a high temperature and high pressure autoclave with a room
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Saarinen and Hamalainen
temperature storage tank and a recirculation loop. The facility is also equipped with an
environment monitoring system to control the water chemistry of the autoclave. The facility is
assembled into a freight container (Fig. 2) that was transferred outside the laboratory facilities in
order to avoid any indoor leakages of hydrogen sulphide. The corrosion potential measurements
were carried out using an Ni/NiSO4 reference electrode, and the results are presented in the scale
of a standard hydrogen electrode.
Table 1. Nominal Composition of the Materials Used in the Study as Provided by the
Manufacturers
Material
POLARIT 725
POLARIT 757
POLARIT 774
POLARIT 778
DUPLOK 21
DUPLOK 22
DUPLOK 25
DUPLOK 26
DUPLOK 27
DUPLOK 27 PM
Inconel 625 PM
Nicrofer 6020
Nicrofer 57161
Nicrofer 5923
C
0.027
0.049
0.02
0.024
0.012
0.013
0.015
0.023
0.01
0.04
0.02
0.01
0.01
0.01
Si
0.41
0.49
0.55
0.4
0.6
0.53
0.43
0.2
0.28
0.38
0.51
0.13
0.05
0.12
Mn
1.45
1.51
1.46
0.46
0.58
0.52
0.62
0.72
0.52
1.16
0.32
0.17
0.37
0.15
S
0,001
0.007
0.001
0.001
0.008
0.01
0.008
0.01
0.008
0.005
0
-
P
0.034
0.029
0.015
0.019
0.032
0.033
0.03
0.03
0.032
0.022
0
0.005
0.005
0.005
Cr
18
17.1
20
19.8
20.7
23
24.5
26.2
27.1
25.9
21.6
21.8
19
22.5
Ni
8.6
10.6
24.7
21.5
8.3
6.1
6.4
7.6
7.9
6.6
Bal.
Bal.
Bal.
Bal.
Mo
0.019
2.51
4.48
6.15
2.8
3
3
2.95
3
3.3
8.2
8.5
15
14
Cu
0.18
0.16
1.67
0.79
1.25
0.19
1.45
0.68
2.5
2.1
0.06
0.06
0.08
0.01
Material
Al
Co Nb
Fe
N
PREN1 PREN2 PREN3 PREN4
POLARIT 725
0.001 0.17 0
Bal. 0.0565 20
35
30
POLARIT 757
0
0.21 0
Bal. 0.046 27
42
45
POLARIT 774
0.022 0.22 0
Bal. 0.0765 37
76
73
POLARIT 778
0.024 0.37 0
Bal. 0.21
47
72
79
DUPLOK 21
0.03 0.05 0.06 Bal. 0.05
32
42
49
-8
DUPLOK 22
0.02 0.1 0.02 Bal. 0.15
38
40
50
-10
DUPLOK 25
0.03 0.07 0.04 Bal. 0.14
39
42
52
-11
DUPLOK 26
0.02 0.07 0.02 Bal. 0.22
43
46
55
-9
DUPLOK 27
0.01 0.05 0.03 Bal. 0.24
45
47
57
-10
DUPLOK 27 PM 0.02 0.05 0.01 Bal. 0.25
45
44
56
-9
Inconel 625 PM
0.02 0.06 3.7 4
0
49
172
140
Nicrofer 6020
0.16 0.05 3.4 3
0
50
173
142
1
Nicrofer 5716
0.09 0.1 0.04 6.5
0
69
172
171
Nicrofer 5923
0.2
0
0.03 0.5
0
69
170
167
1
The tungsten (W) content of the alloy Nicrofer 5716 is 4%
PREN1 = Cr + 3.3 Mo + 32 N. PREN2 = Cr + 2 Ni + 1.5 Mo [15]. PREN3 = Ni + 1.2 Cr + 5.5
Mo [16]. PREN4 = Ni + 0.5 Mn + 30*(C + N) - 1.1*(Cr + Mo + 1.5 Si) + 8.2 [14]
193
Oil Field Corrosion
The test environments were water-based deaerated solutions containing different amounts of
NaCl, H2S and CO2. The materials tested in each experiment and the test type are presented in
Table 2. Also given in Table 2, are the contents of the different species in each experiment and
the pH value of the test solutions. The exposing time of the U-bend specimens in test 1 was 100
hours and in test 2 it was 400 hours. After testing, the specimens were examined using an optical
microscope; the failures were reported. For failure cases, a fractographical inspection was
performed.
0.4
Rmin6.4
B
M8
Rmin6.4
0.03 A-B
M8
A
φ 3.81±0.05
φ 0.05 A-B
12.5
25.4±0.05
12.5
70
Figure 1. Dimensions of round tensile test specimens machined for
the SSR tests. Before testing the specimens were polished
electrolytically
Figure 2. Schematic representation of the Materials Research Shuttle
Laboratory used in the SSR tests. During testing, the test
194
Saarinen and Hamalainen
environment was kept constant with a water recirculation
system
Table 2. Test Matrix of the Studies
Test
1
NaCl Temperature
[wt-%]
[°C]
5
200
H2S
[bar]
8
CO2
[bar]
-
pH
3
2
20
150
5
-
3
3
4
5
6
25
20
20
20
230
93
93
93
70
20
0.5
0.1
50
35
40
85
2.5
2.7
2.6
4.2
Materials
Tested
POLARIT 725
POLARIT 757
DUPLOK 22
All in Table 1 except
Inconel 625 PM
Inconel 625 PM
DUPLOK 27 PM
DUPLOK 27 PM
DUPLOK 27 PM
Test
Type
U-bend
U-bend
SSRT
SSRT
SSRT
SSRT
RESULTS
In test 1, failures were observed in all of the alloys tested. In the austenitic alloys, the
fracture was brittle and intergranular. In the duplex alloy, the fracture partly followed the phase
boundaries between the ferrite and austenite faces, but mainly the fracture was transgranular and
brittle. The corrosion potential of all the alloys during the experiment was near -100 mV (SHE).
In test 2, failures were observed in the alloys POLARIT 725 and POLARIT 757. The
failure mode in these austenitic alloys was mainly intergranular, as observed in test 1. No
failures in the duplex alloys were observed despite clear general and pitting corrosion. Pitting
corrosion was observed in the duplex alloys having a PREN1 value smaller than 40. In the
austenitic alloys, no similar pitting corrosion was observed. The corrosion potential of all the
alloys during experiment was near -100 mV (SHE).
In test 3, no evidence of brittle behavior was found for Inconel 625 PM. The fracture was
ductile, and the elongation to fracture was 40% which corresponds to the elongation observed in
the SSR test conducted in a nitrogen gas reference environment.
In test 4, the fracture of the DUPLOK 27 PM specimen was brittle and transgranular. The
elongation to fracture was only 8%, and the reduction in area was 18%. In the reference test in
nitrogen gas, the elongation to fracture was 32% and the reduction in area 60%.
In test 5, the hydrogen sulphide partial pressure was much lower than in test 4. The fracture
was still brittle and transgranular. Also secondary cracking was observed (Fig. 3). The
elongation to fracture was 26%, and the reduction in area was 25%.
In test 6, the hydrogen sulphide partial pressure was still lower than in test 5. The fracture
was mainly brittle (Fig. 4), but ductile areas were also observed in the fracture surface (Fig. 5).
The initiation and the outermost circle of the fracture surface were brittle, whereas the center of
the fracture surface was ductile.
195
Oil Field Corrosion
Figure 3. SEM micrograph of the DUPLOK 27 PM specimen after test 5
secondary cracks are visible (35x)
Figure 4. SEM micrograph of the brittle fracture surface
196
Saarinen and Hamalainen
of the DUPLOK 27 PM specimen after test 6 (2000x)
Figure 5. SEM micrograph of the ductile fracture surface
of the Duplok 27 PM specimen after test 6 (2000x)
DISCUSSION
The corrosion resistance of passivating alloys is mainly determined by the stability of the
protective layer on the alloy surface. According to the high temperature pH-potential diagrams
for Fe-S-H2O, Cr-S-H2O and Ni-S-H2O systems [17], at pH 3, the surface layers in deaerated
systems consist mainly of nickel and iron sulphides. As the content of the alloying elements of
the alloys increases, the volume of the sulphides in the surface layer increases. In the high
temperature water containing sulphur, the breakdown of molybdenum sulphide greatly influences
the corrosion behavior of highly alloyed alloys containing molybdenum [17]. Despite the fact
that the SSC susceptibility of a certain alloy is the sum of several factors like chemical
composition, microstructure, mechanical properties and heat treatment conditions, in this
study the main considerations were the chemical compositions of the alloys studied.
It is known that the corrosion cracking susceptibility of austenitic alloys increases as the
temperature increases. The more alloying elements the material has, the higher the critical
cracking temperature is. In this study, a temperature of 150°C in test 2 was high enough to cause
intergranular cracking in the austenitic alloys having a PREN1 value under about 30, a PREN2
value under about 40 and a PREN3 value under about 45. Therefore, the critical temperature for
cracking of austenitic alloys having the PREN values mentioned is below 150°C. On the other
hand, no failures were observed in the austenitic alloys having a PREN1 value over 37, a PREN2
value over 72 and a PREN3 value over 73. Therefore, the critical temperature for cracking of
austenitic alloys having the PREN values mentioned is over 150°C. Since only iron and nickel
197
Oil Field Corrosion
sulphides are stable in the environments studied, the failure cases of the austenitic alloys were
apparently related to the various nickel contents of the different alloys. The nickel contents of the
unfailed specimens were clearly higher than those of the failed specimens. Therefore, for
materials selection purposes, the PREN equations including the nickel content of alloys are
recommended for use in H2S-containing environments.
It is also known that the cracking resistance of duplex alloys is enhanced compared to
austenitic alloys having the same level of alloying elements and mechanical properties.
Normally the initiation and the growth of cracks in duplex stainless steels take place in the ferrite
phase of an alloy. In this case, however, the cracking was observed to grow not only in the ferrite
phase, but also in the austenite phase. This was probably due to the high contents of hydrogen
sulphide, and therefore high contents of hydrogen, in the test environments. In test 1, the duplex
alloy failed although the PREN1 value of the duplex alloys was much higher than that of the
failed austenitic alloys. The PREN2 and PREN3 values including the nickel content of the alloys
were, however, at the same level in all of the failed specimens. Therefore, the correlation
between the content of alloying elements and the SSC resistance was better with the PREN
values including the nickel content of an alloy. In test 2, the difference between the PREN2 and
PREN3 values was noted. The PREN2 value showed no difference between the failed austenitic
alloys and the unfailed duplex alloy. The PREN3 values of the failed austenitic alloys were
lower than the PREN3 value of the unfailed duplex alloy. Therefore, the PREN3 value is
recommended for evaluating the SSC resistance of austenitic and duplex alloys. The usefulness
of the PREN4 value correlating with the ferrite content of a duplex alloy could not be evaluated
in this study due to the small number of experiments with different duplex alloys.
The SSR tests showed that the SSR method is aggressive since brittle behavior was
observed in the DUPLOK 27 PM alloy having a PREN3 value of 56 even with small amounts of
H2S in the environment. The fracture surface of DUPLOK 27 PM was mainly transgranular and
mainly in the ferrite phase. The nickel-based alloy Inconel 625 PM was on the contrary resistant
towards SSC even at very high temperature with high contents of hydrogen sulphide, carbon
dioxide and chlorides in the environment.
CONCLUSIONS
In this work the evaluation of the corrosion and SSC behavior of several austenitic and
duplex stainless steels in hydrogen sulphide containing environments was made using the SSR
test method and U-bend specimens. According to the results:
1. When evaluating the SSC resistance of a specific alloy, the PREN values including the
nickel content should be used,
2. The PREN3 value, which did not over emphasize the effect of nickel, gave the best
correlation between the content of alloying elements and SSC,
3. The critical temperature for cracking of austenitic alloys having a PREN3 value under 45
was below 150°C in aggressive H2S-containing environments,
4. The critical temperature for cracking of austenitic alloys having a PREN3 value over 70
was higher than 150°C in aggressive H2S-containing environments,
198
Saarinen and Hamalainen
5. A small decrease in H2S partial pressure had a greater effect on the SSC resistance of
DUPLOK 22 than bigger changes in the temperature and in the NaCl content of the
environment,
6. In the austenitic alloys, the fracture mode was mainly intergranular,
7. In the duplex alloys, the fracture occurred partly on the phase boundaries of the ferrite
and austenitic faces, but the fracture was mainly transgranular,
8. The nickel-based alloy Inconel 625 PM was not susceptible to SSC, even in extremely
aggressive H2S-containing environments,
9. SSR is an aggressive test method, and brittle behavior was observed in the DUPLOK 27
PM alloy having a PREN3 value of 56 even with small amounts of H2S in the
environment, and
10. As the hydrogen sulphide partial pressure was lowered to 0.1 bar, partly ductile fracture
was observed in the DUPLOK 27 PM alloy.
ACKNOWLEDGEMENTS
This study was financed by the Technology Development Centre of Finland, Outokumpu
Polarit Oy, Rauma Materials Technology Oy and VTT Manufacturing Technology. The authors
acknowledge these organizations for their financial support and for providing the test materials.
Also Cronimo Co. is acknowledged for providing the reference test materials.
REFERENCES
1. G.I. Ogundele and W.E. White, Some observations on the corrosion of carbon steel in
sour gas environments, Corrosion 42, 7, 1986, pp. 398-408.
2. A. Miyasaka, et al., Environmental aspects of SCC of high alloys in sour environments,
Corrosion 45, 9, 1989, pp. 771-780.
3. A. Miyasaka, et al., Prediction of critical environments for active-passive transition of
corrosion resistant alloys in sour environments, ISIJ International 31, 2, 1991, pp. 194200.
4. Z.A. Foroulis, Role of solution pH on wet H2S cracking in hydrocarbon production,
Corrosion Prevention & Control, August 1993, pp. 84-89.
5. R. Garber, et al., Sulfide stress cracking resistant steels for heavy section wellhead
components, Journal of Materials for Energy Systems, 1985, pp. 91-103.
6. J. Fliethmann, et al., Autoklaven-Untersuchungen der Spannungsribkorrosion von Fe-CrNi-Legierungen in NaCl/CO2/H2S-Medien, Werkstoffe und Korrosion 43, 1992, pp. 467474.
7. M.F. Brunella, et al., Stress corrosion cracking in sour environments of martensitic and
duplex stainless steels, Proceedings of International Conference on Stainless Steels,
Chiba, June 10-13, 1991. Tokyo 1991, ISIJ, Vol. 1. pp. 264-271.
8. Th. Bollinghaus, H. Hoffmeister and C. Middel, Scatterbands for hydrogen diffusion
coefficients in steels having a ferritic or martensitic microstructure and steels having
an austenitic microstructure at room temperature, Welding in the World 37, 1, 1996,
pp. 16-23.
199
Oil Field Corrosion
9. S.V. Artamoshkin, Effect of the microstructure and nonmentallic inclusions on the
susceptibility of low-alloy steels to sulfide stress corrosion cracking. Translated from
Fiziko-Khimicheskaya Mekhanika Materilov 27, 6, 1991, pp. 60-66.
10. A.A. Omar, et al., Factors affecting the sulfide stress cracking resistance of steel
weldments, CORROSION/81, April 6-10, 1981, Toronto. Canada. Paper No. 186.
11. H. Eriksson and S. Bernhardsson, The applicability of duplex stainless steels in sour
environments, Corrosion 47, 8, 1991, pp. 719-727.
12. M.C. Place, et al., Qualification of corrosion resistant alloys for sour service,
CORROSION/91, March 11-15.1991, Cincinnati, Ohio, USA, Paper No. 1.
13. A. Miyasaka and H. Ogawa, Corrosion performance and applications limits of corrosionresistant alloys in oilfield service, Corrosion 51, 3, 1995, pp. 239-247.
14. T. Kudo, et al., Stress corrosion cracking resistance of 22 % Cr duplex stainless steel in
simulated sour environments, Corrosion 45, 10, 1989, pp. 831-838.
15. K. Denpo, et al., A selection method for high alloy materials for sour service,
Proceedings of Engineering Solutions for Corrosion in Oil and Gas Applications, Milan,
Italy, November 14-17, 1989. pp. 2-1 - 2-14.
16. S. Azuma and T. Kudo, Crevice corrosion of corrosion-resistant alloys in simulated sour
gas environments, Corrosion 47, 6, 1991, pp. 458-463.
17. C.M. Chen, et al., Computer-calculated potential pH diagrams to 300°C, Electric Power
Research Institute, California, USA, Palo Alto 1983, EPRI NP-3137, Vols. 1-2.
200
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
DAMAGE OF PUMP LINKAGES AND TOOL JOINTS
CAUSED BY CRACK CORROSION
A. Kinzel
Amtliche Materialprüfanstalt für Werkstoffe des Maschinenwesens und Kunststoffe
Appelstr. 11 A, D-30167 Hannover, Germany
ABSTRACT
The Amtliche Materialprüfanstalt (AMPA) is an official materials testing institute and a
subordinate board of the Ministry for Economy, Technology and Traffic of Lower Saxony, which is a
federal state of the Federal Republic of Germany. The AMPA works in a lot of fields of materials
testing, quality testing, quality assurance and damage investigation. The special fields of the AMPA
include acceptance testing, supervision and inspection in the field of plant and pipeline construction
for natural gas and oil.
Because of this, the AMPA is often used by operators of pipelines or plants to give their expert
opinions in cases of damage. Two damage investigations will be given in this article. They are
damage to several pump linkages and damage to a tool-joint.
Key Words: Pump linkage, tool joints, damage investigation, acceptance tests, inspection
INTRODUCTION
Pump linkages and tool joints are used for pumping oil and gas from a depth of about
1,500 m in northern Germany. The pump linkages have lengths of about 6-10 m and
diameters of 7/8 in up to a depth of 400 m and 3/4 in beyond that depth. The pump linkages
reported here are directly connected. Because of tapping of different caverns from one
starting point, the drilling hole is often curved in to a horizontal drilling situation. The pump
linkages can follow these curves, because of their oscillating movement within the hole, the
linkages have proctection devices against mechanical damage made of plastic which are
placed at several defined areas along every linkage. The puming frequency is low-cycle and
depends on the length and weight of the equipment. The pumped medium is a mixture of
about 2% crude oil and 98% salty water.
These pump linkages are connected directly, whereas other pumping equipment uses tool
joints to connect their linkages. The tool joint has a length of about 0.5 m and a diameter of
about 0.2 m. The pumping rate is about 2,000 l/minute.
PUMP LINKAGE
The damaged pump linkages were made of 20 NiMo 8. These linkages can normally be
used for an operating time of about 18 months. For the case of damage in question, the pump
linkages broke after an operating time of 3-6 months. Operation was between 4,000 and
13,000 hours for at least 1.1-3.5 million oscillating units. All of them were broken at depths
201
Oil Field Corrosion
of 500-1,000 m, so the broken linkages were 3/4 in diameter. These pump linkages were
broken directly beneath the connecting areas of the linkages.
Besides this, the mode of operation of pumping was changed. The frequency of
pumping was halved, which led to increased vibration in the pump linkages. The mechanical
stress on pump linkages is oscillating stress because of its own weight, and therefore, mass
inertia, as well as bending stress especially in the transition sections to more bend-stiff areas.
Further on, corrosion is caused by aggressive components within the pumped oil (salty water)
especially if the corrosion protective layer is damaged.
Visual inspection showed that there was no visible coarse surface damage. All breaks
were in transition sections to more bend-stiff areas of the pump linkage (Fig. 1) and looked
very similar to fatigue fracture or crack corrosion (Fig. 2). The corrosion protective layer was
not regular in terms of thickness of layers and showed small blisters on the surface (Fig. 3).
Within a metallographic inspection, a faultless purity and texture was noticed. The
pump linkages had no faults in the material and its texture. Near the surface, a slight
reduction of carbon content was visible (Fig. 4) as well as a rough surface in the base material
with detached corrosion protective layer (Fig. 5).
Testing of the mechanical and technological properties of the puming linkages disclosed
no faults. The required tensile strength and extension could be reached by all test samples.
Also the materials composition met the demands.
Figure 1. Breaking of a pump linkage (K 2149/6)
202
Kinzel
Figure 2. Breaking area (K2149/11)
Figure 3. Blisters on the protective surface (K2165/20)
Figure 4. Reduction of carbon content near the surface (M 23370)
203
Oil Field Corrosion
Figure 5. Rough surface of base material with a detached corrosion protective layer
(M 233415)
TOOL JOINT
The object of this investigation was a broken tool joint (7 3/8 in OD x 4 9/16 in ID, 5 1/2
in IF connector with 5 ½ in drill pipe, grade G 105, Fig. 6). The built-in rope length was
3,600 m, the pumping rate was 2,000 l/min at a pressure of 180 bar, the coupled load was 160
tons, and the torque at top-drive about 41,500 N/m. The breaki was at a depth of 2,173 m and
occurred while clearing the hole. The linkage had been stored in open air from 1986 to
joining in 1995.
Figure 6. Broken tool joint (K 2182/27A)
The break was vertical to the tool joint axis and looked like a fatigue fracture (Fig. 7).
The starting point of the break was near the transition area to the pivot (Fig. 8), and the
fatigue fracture area was only a small part of the whole break area. Near the starting point, a
corrosion area with cracks was detected. A metallographic inspection revealed faultless
purity and texture. The materials analysis met the demand as did the mechanical properties of
the material.
204
Kinzel
Figure 7. Break area of the tool joint (K 2182/25A)
Figure 8. Starting point of the break (K 2185/1)
DAMAGE ANALYSIS
Pump Linkage
Within the transition sections to more bend-stiff areas, there is a combination of normal
oscillating stress and bending stress. Further on, and combined with an aggressive pumped
medium and a damaged corrosion protective layer, tension-induced corrosion had started.
After initiating a first crack, this had lead to crack corrosion. A second crack about 24 mm
below the first one was detected at one of the pump’s linkages, which is another indication of
crack corrosion (Fig. 9). Increased cracking is mostly determined by oscillating frequency,
tension amplitude and corrosion effects. The low-cycle operation of the pump linkages may
have had a negative effect. The area near the crack’s starting point had a low crackincreasing velocity and a lot of corrosion effects, whereas the break area showed a higher
crack-increasing velocity combined with reduced corrosion effects.
The damage was caused by an insufficient corrosion protective layer covering the rough
material’s surface with an irregular thickness. The small blisters may have been caused by
fouling or corrosion before being coated by the protective layer or by a lack of resistance to
the aggressive medium. Therefore, it can be concluded that the damage-causing event was a
coating process which did not meet the demand.
Tool Joint
Because the tool joint was stored in the open air for a long time, corrosion was initiated
within the damaged area. Besides this, the damaged area is a transition area to a zone with
increased tension. The combination of both, corrosion and the tension during operation lead
to the first cracks. Further operation and the tension peak lead to crack growth, crack
205
Oil Field Corrosion
corrosion and fatigue fracture. Because of the low percentage of fatigue fracture area, it can
be concluded that the operation tension must have been high to break the tool joint.
Figure 9. Corrosion crack about 24 mm below the break (M 23414)
CONCLUSIONS
Corrosion caused by inadequate protective layers or by incorrect storing leads to
initiation of cracks during operation. The oscillating of the parts increases the cracking. This
is very dangerous for operation because the failure could not be detected in most cases by
visual inspection before use. The failure seemed very surprising for the operation team, and a
cost-intensive repair must be done, while the pumping system is not in operation. Therefore,
attention has to be paid to effective quality management and quality inspection before using
such parts to avoid these failures.
206
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
ANALYSIS OF SOILS POSSIBILITY TO GIVE RISE
TO PIPE METAL STRESS CORROSION CRACKING
V.G. Antonov and S.A. Loubenski
All-Russian Scientific Research Institute of Natural Gases and Gas Technology
(VNIIGAS), p. Rasvilka, Leninsky rajon, Moscow region, 142717, Russia
ABSTRACT
The target of the present study is to analyse the possibility of soils, sampled along the routes of
gas mains laid without cathodic protection, to give rise to pipe metal stress corrosion cracking (SCC).
The tests included soils taken both on gas pipelines operated for 30 years without failure and on
pipelines that had failed due to SCC. The pipe steel specimens were tested by the constant rate slow
deformation method in soil suspensions taken on gas pipelines that had failed due to SCC. A change
of the main parameters which characterise steel resistance to cracking was observed. This change
included a decrease in the plastic properties of the metal, namely, a decrease in the relative elongation
and reduction of area, and the formation of surface corrosion cracks near the place of failure. The
results of the experiments show that the soils taken on gas pipelines which had failed due to SCC may
promote the formation of surface corrosion cracks.
Key Words: Stress corrosion cracking, gas pipeline, soil, constant rate slow
deformation method.
INTRODUCTION
Today, SCC is one of the serious problems in the gas industry. Last year, SCC was the
cause of some failures in gas pipelines in Western Siberia and Central and Northern Russia.
The attempts to explain SCC on gas mains only by hydrogen embrittlement or carbonate
cracking are not correct since many years of experience in the operation of offshore cathodic
protected gas pipelines and laboratory tests of the resistance of pipe steels to hydrogen
embrittlement show that the content of hydrogen absorbed by the metal under cathodic
polarization is less than the critical content which leads to embrittlement.
The Laboratory of Corrosion-Resistance Materials and Corrosion (VNIIGAS) has
conducted investigations which showed that SCC of pipe steels at potentials from 1.0-1.2 V
(SHE) in carbonate media with hydrogen numbers between 7.0 and 11.0 had not occurred.
THE TARGET OF THE INVESTIGATION
The target of the present study was to analyse the potential of soils sampled along the
routes of gas mains to give rise to pipe metal stress corrosion at corrosion potentials, i.e.,
without a cathodic protection system.
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Oil Field Corrosion
EXPERIMENTAL PROCEDURE
The investigations were carried out on soils taken on gas mains which had not been
subjected to failure due to SCC (soils A and B), on soils taken on failed gas pipelines having
been under operation from 6 to 18 years that failed due to SCC (C (Western Siberia), D
(Northern Russia), and E (Central Russia)), on aqueous solutions of carbonates and chlorides,
on a medium specially designed in VNIIGAS for conducting corrosion tests, and a selection
of pipe steels. The pipe steel samples were made according to NACE Standard TM 0177
(Testing of Metals for Resistance to Sulphide Stress Cracking at Ambient Temperatures).
The soil samples were dried and crushed in a mortar. A weighted sample was placed in a
flask followed by the addition of distilled water in the ratio of 1000 g water to 1000 g soil.
The suspension was placed in a cell and was given some time to settle.
The corrosion tests were conducted on round specimens 6.0 mm in diameter, using a
special test unit (MP- 5-8B type) with a travel rate of active entertainment of 1.8 x 10-6 m/sec.
The specimens were measured and marked before the test in order to determine their
relation elongation and reduction of area after the test. The specimens were degreased with
organic solvent using a common method and placed in a 250-cm3 cell. Then the cell was
filled with the test solution. The specimens were held for 72 hours followed by setting the
required deformation rate. All tests were carried out at natural aeration and ambient
temperature.
The tests on SCC resistance were conducted on not fewer than 4 specimens in each
medium. The tests ended when the specimen failed. Failed specimens were washed, dried
and measured. The test results were omitted when failure took place in the unit’s holder, or
due to metallurgical defects or along the marks.
The SCC resistance was estimated by the relative change in the properties of the metal
tested in air and in corrosive media, and by the formation of secondary corrosion cracks on
the specimens surface.
The elongation and reduction of area were calculated on the basis of GOST 1497-84
(Metals: Methods of Tensile Tests):
δ = (l - l0) x 100 / l0
(1)
where : δ = the elongation (%)
l0 = the initial design length of a specimen (mm)
l = the length of a specimen after its failure (mm)
ψ = (S0 - S) x 100 / S0
(2)
where: ψ = the reduction (%)
S0 = the initial cross-section area of a specimen (mm2)
S = the minimal cross-section area of a specimen after its failure (mm2)
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Antonov and Loubenski
RESULTS AND DISCUSSION
The test results are given in Tables 1, 2, 3 and 4. The electrochemical measurements
made on pipe steels in aqueous suspensions of soils showed the following:
• The pH values of the aqueous extracts of soils sampled on failed gas pipelines were
between 5.6 and 6.0 (C, D, and E), and the pH of soil suspensions (A and B) are
between 7.6 and 10.6. Failures of pipe steels due to SCC were not observed in soils
A and B.
• The corrosion potentials of pipe steels in the soils sampled on gas pipelines which
failed due to SCC are between -0.56 and -0.58 V (SHE) (C, D, and E), the corrosion
potential in soil with a pH of 7.6 (A) was -0.48 V (SHE), and the corrosion potential
in soil with a pH of 10.6 (B) was -0.20 V (SHE).
• The general corrosion rate for pipe steels in all types of soil did not exceed 0.01
mm/year.
Aqueous suspensions of soils C, D, and E contained hydrogen sulphide, acetic acid,
formic acid and compounds of lead, selenium and arsenic.
The SCC resistance tests showed that in such systems as aqueous solutions of chlorides
and carbonates and the soil suspensions A and B, the main parameters characterising steel
resistance under this test method are close to the parameters of the steel set tested in air.
Table 1. Resistance of Steel X-65 Specimens to SCC Determined by the Constant Rate
Slow Deformation Method in Media of Different Compositions
Corrosion
Potential, V
(SHE)
δ
ψ
(%)
(%)
-
-
20.4
65.0
No
Solution Developed by
VNIIGAS
5.1
- 0.43
10.8
25.8
Yes
3% NaCl
6.8
-0.44
23.6
59.9
No
NaHCO3 - NaCO3
9.6
-0.06
21.7
66.9
No
Soil of Pipeline A
7.6
-0.44
18.8
57.6
No
Soil of Pipeline B
10.4
-0.20
23.0
60.0
No
Soil of Failed Pipeline C
5.8
-0.52
15.3
42.2
Yes
Soil of Failed Pipeline D
5.8
-0.56
17.2
45.6
Yes
Medium
Air (Control)
pH
Presence of
Corrosion Cracks
211
Oil Field Corrosion
Soil of Failed Pipeline E
5.6
-0.56
21.3
56.1
Yes
When pipe steel specimens were tested in the soil suspensions sampled on the pipelines
which failed due to SCC, there was a change in the main parameters characterising steel
resistance to cracking, namely, a decrease in the metal ductility (decrease in elongation and
reduction of area) and the formation of surface corrosion cracks near the point of failure.
Table 2. Resistance of Steel X-70 Specimens to SCC Determined by the Constant Rate Slow
Deformation Method in Media of Different Compositions
Corrosion
Potential, V
(SHE)
δ
ψ,
(%)
(%)
-
-
25.2
46.1
No
Solution Developed by
VNIIGAS
5.1
- 0.43
7.4
11.0
Yes
3% NaCl
6.8
-0.44
21.0
43.1
No
Soil of Failed Pipeline C
5.8
-0.52
14.3
21.2
Yes
Medium
Air (Control)
pH
Presence of
Corrosion Cracks
Table 3. Resistance of Steel X-56 Specimens to SCC Determined by the Constant Rate Slow
Deformation Method in Media of Different Compositions
Corrosion
Potential, V
(SHE)
δ,
ψ,
(%)
(%)
-
-
28.1
59.4
No
Solution Developed by
VNIIGAS
5.1
- 0.43
21.0
38.1
Yes
3% NaCl
6.8
-0.44
21.2
50.1
No
Soil of Failed Pipeline C
5.8
-0.52
14.2
48.3
Yes
Soil of Failed Pipeline D
5.8
-0.56
16.1
38.1
Yes
Medium
Air (Control)
212
pH
Presence of
Corrosion Cracks
Antonov and Loubenski
Table 4. Resistance of Iron (99.9%) Specimens to SCC Determined by the Constant Rate
Slow Deformation Method in Media of Different Compositions
Corrosion
Potential, V
(SHE)
δ,
ψ,
(%)
(%)
-
-
24.2
77.5
No
Solution Developed by
VNIIGAS
5.1
- 0.43
19.3
53.5
Yes
3% NaCl
6.8
-0.44
23.9
72.8
No
Soil of Failed Pipeline C
5.8
-0.52
23.2
75.0
Yes
Medium
Air (Control)
pH
Presence of
Corrosion Cracks
CONCLUSIONS
The laboratory test results show that the soils sampled on gas pipelines which failed due
to SCC favour the formation of surface corrosion cracks, i.e., the main cause of SCC was the
physical-chemical interaction between the metal and the components of the soils (hydrogen
sulphide, acetic acid, formic acid and compounds of lead, selenium and arsenic).
213
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
A MYSTERIOUS DOWNHOLE CORROSION FAILURE IN AN OIL WELL
A. Husain and A. Hasan
Materials Application Department
Kuwait Institute for Scientific Research
P.O. Box 24885-Safat 13109, Kuwait
ABSTRACT
Recently, the failed downhole tubing from one of the oil wells in a Kuwaiti oil field was
examined at the Materials Application Department of the Kuwait Institute for Scientific Research.
This paper presents the results of a failure investigation carried out on the J55 steel tubing materials
used in many downhole oil wells in that particular oil field area. The results showed that the failure
of the tube was due to localized attack in the form of pitting and crevice corrosion that initiated on the
internal surface of the tubes. In contrast, no downhole tubing failures were reported for any of the
other oil wells located in the same field and exposed to the same aquifer environment and operation
conditions. The internal side of the failed tube plate was exposed to the oil/water stream with a 73%
water cut and high salinity content in addition to CO2 and H2S corrosive gases. The preliminary
results of the investigation indicate that the likely causes of this type of tubing corrosion failure may
be related to corrosive environmental effects combined with metallurgical or mechanical properties of
the steel tubing material. Moreover, the combined effects of both chloride and CO2 and H2S species
under certain condition are widely known to cause accelerated, premature corrosion failure in oil well
tubing. In this study, a material failure case in the oil field will be discussed. Several attempts were
made to determine the exact cause of the failure found in this oil well. However, no convincing
evidence as to the specific cause of the failure was found, due to the additional detrimental effects of
many other factors.
Key Words: Oil field corrosion, oil well tubings, downhole equipment corrosion, CO2
corrosion, H2S corrosion.
INTRODUCTION
Recently, two steel joints in a section of downhole tubing failed due to severe
perforations. Visual inspection revealed that the downhole tubing experienced severe
corrosion damage and complete perforations of the tube’s walls. To elucidate the causes of
the tubing failure, the company requested the Kuwait Institute for Scientific Research (KISR)
to carry out specific tensile, hardness, and scanning electron microscopy (SEM)
examinations. The problem was first thought to be an isolated example of a manufacturing
defect; however, the present study indicates that the cause of failure merits further
investigation.
Information gathered on the oil well and failed parts indicates that only one well, No. W41, experienced accelerated corrosion failure. It was indicated that the downhole tubing was
exposed to the oil/water stream for 0 days prior to failure. Two perforated tube joints were
215
Oil Field Corrosion
supplied by the company for an investigation into the causes of the perforations. The first
tube sample, designated as Joint No. 33, had one perforated hole (Fig. 1a). The second tube
sample, designated as Joint No. 60, had four perforated holes (Fig. 1b). An additional
nonperforated tube was also examined to compare its quality with that of the failed tubes.
Each tube measured 35 cm in length, 60.75 mm in outside diameter and 4.49 mm in wall
thickness.(Appendix A contains full details of the operations condition and production data).
This report presents the results of the investigation as well as a discussion of the likely
causes of failure and a possible solution to the problem.
LABORATORY EXAMINATIONS
Chemical Analysis and Metallographic Examinations of the Material
The material of the tube was chemically analyzed to confirm its composition and to
ensure that it complied with the requirements of the API J55 specifications. The analysis was
carried out on a cleaned section cut from the tubes.
A metallographic examination was carried out to identify the microstructures of the asreceived, perforated tubing samples. The examination was conducted on metallographically
prepared specimens cut from one of the failed tube joints. The metallographic preparation
included grinding the specimens with silicon carbide papers of the grade sequence: 320, 600,
1000 and 1200 grit. This was followed by polishing with 3 and 6 μm diamond abrasive
pastes, and rinsing with distilled water and acetone. Finally, the specimens were etched in
2% nital etchant solution to reveal the microstructure and the second phase precipitates or
inclusions.
Microscopic Examinations of Perforations and Energy Dispersive Spectroscopic (EDS)
Analysis
Optical microscopy and SEM were used to observe the perforations and the dark
hydrocarbon deposits, or scale, found along the internal surfaces of the tubes. The failed
tubes were initially sectioned into two halves in a longitudinal direction away from the
defective areas before the microscopic examinations. The examinations were conducted
before and after the removal of the deposits as well as on the etched surfaces. The purpose
of etching was to identify the microstructure at the initiation sites of the perforations.
Qualitative EDS analysis was carried out to determine the composition of the deposits
both inside and outside the perforations. X-ray diffraction (XRD) analysis was carried out to
determine the composition of the deposits scrubbed from the internal surfaces of the
perforated tubes.
Mechanical Properties of the Tubing Materials
Mechanical tensile and hardness tests were conducted to investigate the quality of the
failed as-received tubing material sample. These tests were carried out to confirm the
tubing’s mechanical properties as well as to ensure that they complied with the requirements
of the API J55 specifications for downhole tubing material.
216
Husain and Hasan
(a)
(b)
(c)
Figure 1. Perforation resulting from localized corrosion attack of internal tubing surface:
217
Oil Field Corrosion
(a) Joint No. 33; (b) Joint No. 60; and (c) condition of interior of tubing surface
RESULTS
Chemical Composition and Microstructure of the Material
Table 1 presents the results of the chemical analysis of the material of the perforated
tubes and the as-received un-used tubes. The alloy of both tube samples was carbon steel
with minor elemental constituents of Cu, Si and Cr. The carbon and sulfur contents of the
alloy were not properly determined due to analytical limitations.
Table 1. Chemical Composition of the As-Received and Perforated Tubes
Element
C
Si
S
P
Cu
Mn
Cr
Fe
AsReceived
Tube
(wt%
0.45-0.6
0.3
0.015
0.03
0.32
1.45
0.06
Balance
Perforated
Tube
(wt%
Error
0.45-0.6
0.33
0.05
0.03
0.33
1.56
0.06
Balance
±0.05
±0.05
±0.005
±0.005
±0.005
±0.005
±0.01
±0.01
Structurally, as indicated by the equilibrium diagram of iron and iron carbide, the alloy
for both tubes was composed of ferrite and pearlite. The structure of the iron-carbon alloy is
shown in the photomicrographs of Fig. 2a and b. It can be seen that the microstructure is
typical of carbon steel material in terms of the presence of equal proportions of ferrite and
pearlite phases.
Microscopic Examinations of Perforations and EDS Analysis
of Deposits in the Failed Tubes
Visual and optical microscopic examinations of the failed tubes revealed several
perforations that appeared to originate from inside the tubes. Examinations of the internal
surfaces, after splitting the tubes into halves, showed numerous crevices and pit-like shapes in
a somewhat straight line along the longitudinal axis. The internal surfaces also exhibited
many rounded, mutually intersecting pits lying partially beneath a mixture of dark scale of
hydrocarbon deposits with iron oxide and dry oil residue. The scales were observed inside
and around penetrated areas. Cleaning the metal’s surface revealed the presence of crevice
corrosion in addition to the observed pits. The crevice attack appeared the same beneath all
kinds of deposits, with an approximate depth of 1.21 mm. Some of these crevices or pit-like
shapes penetrated the full thickness of the tubing (Fig. 1c). Visual examinations of the asreceived tubing indicated the presence of mill scale that covered the internal surface of the
tube.
218
Husain and Hasan
EDS analysis of the deposits inside the pit-like shapes showed that they were enriched
mostly with sulfur, chloride, calcium and iron and, to a much lesser extent, with manganese
and silicon (Fig. 2c). Sulfides were present in the deposits and corrosion products.
Aggressive sulfur-and chloride-containing species were also concentrated beneath the iron
oxide layer (Fig. 2d).
(a) as-received tube
(c) EDS analysis inside of pits
(b) perforated tube
(d) EDS analysis beneath deposit layer
Figure 2: (a and b) SEM micrographs of etched specimens both showing ferrite
and pearlite phases distributed in equal proportion; and (c and d) EDS
analysis of compounds in the perforated side
219
Oil Field Corrosion
XRD analysis of the scrubbed scale obtained after cleaning the internal surfaces of the
perforated tubes indicated the presence of magnetite (Fe3O4), iron oxide hydroxide
[FeO(OH)], iron carbonate (FeCO3) and iron. This suggests that the layers of deposit present
on the internal surfaces of the tube are possibly from mill scale (as Fe3O4) and also from
corrosion by-products of carbonic acid combined with iron in the form of FeCO3.
Mechanical Properties of the Tubing Material
Table 2 presents the mechanical properties of the tubing material in terms of tensile
strength, yield strength, and hardness. The results of the mechanical testing showed that the
tubing materials fall within the same range of API J55 specifications.
Table 2. Mechanical Properties of the Tubing Material
Tubing
Material
As-received
Tube
Perforated
Tube
Tensile
Stress*
(MPa)
675
675
Tensile
**
Stress
(MPa)
695
Yield
**
Stress
(MPa)
480
+
+
Rockwell
Hardness
(HRB)
94
94
*Approximate tensile stress based on hardness measurements.
+ Standard tensile test samples were not prepared due to severe perforations of the tube wall
(data obtained with tensile test machine).
DISCUSSION
Both the optical microscopy and the SEM examinations indicate that the failure of the
tube was caused by a typical case of localized corrosion caused by differential acid
concentration cell. The fact that localized attack occurred on the internal surfaces of the
tubing at accelerated rates during 10 days of exposure indicates that the cause of this
corrosion phenomenon is related to the aggressive nature of the environment. Under normal
conditions in deaerated oil field environments of CO2 and H2S, one would not expect
corrosion perforation to occur at such a fast rate.
Generally, the corrosion rates of steel are not as high in sour water systems (i.e., those
usually observed in oil field systems with a high water cut) as they are in CO2 systems due to
the somewhat protective nature of sulfide scale (FeS mackinawite) relative to iron carbonate
(FeCO3) unless oxygen contamination has occurred [1].
Typically, deaerated systems with less than 0.02 bar CO2 partial pressure are not
considered excessively corrosive to steel, but they can exhibit corrosion rates of up to 0.20
mm/year (10 mpy). As the partial pressure increases, the corrosion rate increases. At 0.5 bar
CO2 partial pressure, the corrosion rate of steel (i.e., 1 mm/year (40 mpy)) is high enough to
require inhibition if the bicarbonate is low as in condensed water systems. At 2.0 bar CO2
partial pressure, the system would be at a pH in the range of 3.5-4.5, and can be considered
220
Husain and Hasan
severe from the standpoint of weight loss corrosion (i.e., a corrosion rate of > 2.5 mm/year
(100 mpy)).
The combination of hydrogen sulfide and carbon dioxide is more aggressive than
hydrogen sulfide alone and is frequently found in oil field environments. Once again, the
presence of even minute quantities of oxygen can be disastrous. In all cases, increased
velocity would be expected to increase the corrosion rate of steel [2].
The following ferrous corrosion products would form with H2S and CO2 in the presence
of oxygen and low solid water. For corrosion, they are the only products of concern:
Fe + H2S → FeS + H2
(Sour corrosion)
(1)
Fe + H2O + CO2 → FeCO3 + H2
(Sweet corrosion)
(2)
4Fe + 3O2 → 2Fe2O3
(Oxygen corrosion)
(3)
The iron sulfide produced by reaction (Eq. 1) generally adheres to the steel surfaces as a
black powder or scale. The scale tends to cause local acceleration of corrosion because the
iron sulfide is cathodic to the steel; this usually results in deep pitting during O2 reduction
reactions along this FeS layer. Oxygen will provide a high electrochemical potential because
it is a strong, rapid oxidizing agent. This means that it will easily combine with electrons at
the cathode, and allow the corrosion to proceed at a rate limited primarily by the rate at which
oxygen can diffuse to the cathode. The primary concern in this case would be when the H2S
partial pressure is ≥ 0.05 psi.
As a first approximation, calculations of the partial pressure of both carbon dioxide and
hydrogen sulfide are useful for the present case study in predicting the degree of corrosivity
of the oil well (i.e., No. W-41):
Partial pressure of CO2 = 1500 x 0.07 = 105 psi
Partial pressure of H2S = 1500 x 0.001 = 1.65 psi
Using the CO2 partial pressure as a yardstick to predict corrosion, it was found that the
CO2 partial pressure was above 30 psi, which usually indicates the onset of corrosion,
however, in this study, the volume fraction of CO2 over that of H2S was < 200. This
indicates that the presence of small concentrations of H2S play a prominent role in the
corrosion mechanism involved in the reactions (Eqs. 1-3). If the given volume fraction is >
200, then carbon dioxide will be the controlling factor for the reactions; this does not apply to
the present case and its oil field environment.
According to a computer modeling program [3] used in this study to estimate the
corrosion rate on the pitted tubing, and based on the analysis of the production data in
addition to the operation conditions given in Appendix A, the following two cases were
predicted:
221
Oil Field Corrosion
1. The production of oil and formation water with H2S and CO2 under deaerated
conditions (i.e., oxygen is excluded) would result in a corrosion rate resulting
in tubing perforation after more than two years of exposure, and
2. Changing the production conditions to include oxygen in addition to H2S and
CO2 would result in a worst-case corrosion rate leading to tubing failure in
about eight days (10 days in service to actual failure).
Based on the examination and this preliminary computer analysis from CLI
International, Inc., it was assumed that case 2 is likely to model the situation involved in the
tubing failure. However, according to on-site field inspection of the oil well and observation
of the huge amount of gas pressure released from the wellhead, when the pressure valve was
opened, the theoretical assumption based on the combined effects of oxygen and corrosive
gases could not be accurate because oxygen had no chance to diffuse into the well tubing
under such a high-pressure release of gas at the outlet. Moreover, there was no chance for
oxygen to be drawn into the well through leaky valves or the expansion of joints in this
solidly designed oil well tubing part. Therefore, the only solution was to look for other
causes of damage that would have enhanced the corrosivity of the gases (i.e., CO2 and H2S)
and could have contributed to the acceleration of the rate of corrosion attack.
On observing the case history of the oil well (see Appendixes A and B) and by
comparing the performance of this oil well with respect to the nearest oil wells located within
one kilometer in the same oil field, the following suspected factors can be assumed: The
failure of the tubing was initiated due to differential acid concentration cell (case 1) that had
been established beneath crevices that were hydrocarbon in nature, and being strengthened
with a very high propagation rate of corrosion attack caused by one or more of the following
suspected factors:
• Failure in the electrical submersible pump materials (ESP) as indicated in
•
•
•
•
•
Appendix B,
Electrical power problems (i.e., a rectifier power supply problem),
Galvanic coupling effects between the materials of the ESP and the internal
surface of the tube,
Failure including doglegs, splits, and leakage of the well casing materials at a
certain depth adjacent to the well tubing promoted after the burning of the well
during the Iraqi invasion, and
Faulty electrical ground connection (ground electrical contact was made to the
well tubings), and
Stray current effect from either the ESP or other exterior effects.
CONCLUSIONS
1. Metallographic examination, SEM and mechanical testing of the tubing did not show
any major differences between the as-received and the failed tubing in terms of
mechanical properties or the presence of second phase particles and inclusions. The
minor additions of Cu, Cr and Si did not contribute much to the poor corrosion
properties of the steel tubing; however, the presence of higher contents of
nonhomogenous MnS in both tubing specimens may also indicate the possibility of
222
Husain and Hasan
MnS inclusion caused by variable degrees of extrusion during pipe manufacturing.
Manganese sulfide inclusions are known to be detrimental to corrosion resistance.
2. Oxygen was not the most likely cause of the perforation.
3. A solution to combat the tubing corrosion problem is to use a more expensive
approach by carefully evaluating and selecting corrosion-resistant alloys or to inhibit
corrosion with chemicals to minimize corrosive attack on the steel tubing.
4. Internal coating of tubing using powder coating systems (e.g., Tuboscope Vetco
International) or fiberglass reinforcements have proven to be successful in many oil
industry and downhole applications.
REFERENCES
1. F.W. Smith, Structure and Properties of Engineering Alloys, New York, MacGrawHill, 1993.
2. NACE. Corrosion control in petroleum production. National Association of
Corrosion Engineers, Publication No. TPC 5, Houston, Texas, 1979.
3. Private communication with Dr. R.D. Kane during his visit to Kuwait, 1995.
APPENDIX A
Case History for Well No. W-41
Industry: A Kuwaiti oil production
Specimen location: Downhole tube from Well No. W-41
Specimen orientation: Vertical
Years/days in service: 10 days
Salinity: 80,000 to 90,000 ppm
Average water cut: 73%-Temperature: 116oF (47oC). Total pressure: 1250-1500 downhole
CO2 Concentration: 7 vol% H2S Concentration: 0.11 vol%
Failed samples given: Two failed tubing samples with tube specifications J55-2 3/8 in. O.D.
Sample designation: Sample No. 3 from Joint No. 33. Downhole tubing depth: ± 1008 ft
Failure condition: One perforated hole
Sample designation: Sample No. 4 from Joint No. 60. Downhole tubing depth: 1832 ft
Condition: 4 perforated holes
Engineering component connection: Electrical submersible pump (ESP) at 2,258 ft
Oil well fluid level: static 527 ft. Dynamic 558 ft
223
Oil Field Corrosion
APPENDIX B
.History of the Premature Failure of Downhole Equipment and Tubings in Well No. W-41.
Visual Inspection Findings of Pulled Equipment
Failure Failure Run Life
No.
Date (months)
1
Jan.
2.5
90
2
May
3.5
90
3
Oct.
5
92
___
burnt
leaking
___
___
shorted
leaking
damaged
corroded
burnt
___
___
___
___
pump
Joint.
#29
burnt
with
2 holes
burnt
Ext. corr.
erosion
corrosion.
___
good
___
good
#30
#59
External
corrosion
Erosion
corrosion
___
Ext.
Corr.
left
in
hole
5
Jul.
93
Jan.
94
6
___
6
good
7
Nov.
94
10
good
(red alloy)
8
Nov.
94
0.5
good
(red alloy)
9
Jan.
95
1.5
10
Feb.
95
10 days
11
Mar.
95
40 days
good
(carbon
steel)
good
(carbon
steel)
good
(carbon
steel)
12
Jul.
95
42 days
13
Sept.
95
Failed tubing Jts
(Times Failed)
Depth of tubing
Section Failed
Frequently
29
(1)
one
hole
one
hole
Iraqi
Invasion
DDI
while
RIH
Separator
good
2
___
Protector
3
6
___
Motor
Jan.
90
good
(carbon
steel)
good
(carbon
steel)
30
(1)
good
hole &
good
Ext. Corr.
(Red Alloy)
(carbon
steel)
reused
shorted
good
+big hole
stuck
(Red Alloy)
(carbon
(carbon
steel)
steel)
good
good
good
(carbon
(carbon
(Red Alloy)
steel)
steel)
good
good
good
(carbon
(Carbon
(red alloy)
steel)
steel)
wire
bolts<------corroded---connection
--------->red alloy
loose
(carbon
steel)
good
good
1/16 inch hole
(carbon
(red
(red alloy)
steel)
alloy)
good
red
red alloy
(carbon
alloy
steel)
33
37
39
(1)
(1)
(1)
900 ft- 1000 ft
Cable
Remarks
Pump
4
PSI
Failed Tubing
Joint
Joint No. of
No.
Holes
___
___
not
R.C.
work damage
___
___
shorted
(red alloy)
DDI
while
RIH
pump
Joint
___
F.C.
burned
&
parted
good
___
good
___
good
#59 3 holes DDI while
&#63
&1
RIH
hole
(KSRC)
#33
1 hole
KISR
&#60 &4 hole
___
___
___
one
hole
hi Volt
hi Amp
#37,
#39,
#50
___
#50,
#58,
#59
58
(1)
59
(3)
5 holes
3 holes
7 holes
red
alloy
50
(2)
60
(1)
N-80
grade
tubing
N-80
grade
tubing
63 pup
(1) joint
(2)
1800 ft- 1925 ft
Production data for Well No. W-41:
Operating Conditions:
-Average production rate = 2,750 BFPD
At subsurface (downhole)
-Average water cut = 73%
Temp.:116°C
-Salinity = 80,000- 90,000 ppm
Pressure:1250-1500 psi
-Setting depth of:
CO2 Conc.:7 Vol.%
Electrical Submersible Pump (ESP) = 2,258 ft.
H2S Conc.:0.11 Vol%
Dynamic fluid level = 558 ft. Static fluid level = 527 ft. -GOR = 635
224
DDI
while
RIH
At Surface
85°C
110 psi
15 Vol.%
1.1 Vol.%
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
METHODOLOGIES FOR ASSESSMENT OF
CRUDE OIL CORROSIVITY IN PETROLEUM REFINING
S. Tebbal and R.D. Kane
CLI International, Inc.
14503 Bammel N. Houston, Suite 300
Houston, Texas 77014-1149, USA
ABSTRACT
Lower quality “opportunity” crudes are now processed in most refineries, and the source of the
crudes may vary daily. These feedstocks, if not properly handled, can result in reduction in service
life of equipment as well as costly failure and downtime. Analytical tools are needed to predict their
high-temperature corrosivity toward distillation units. Threshold levels of total sulfur and total acid
number (TAN) have been used for many years as rules of thumb for predicting crude corrosivity.
However, it is now realized that they are not accurate in their predictive ability. Crudes with similar
compositions and comparable with respect to process considerations have been found to be entirely
different in their impact on corrosion. Naphthenic acid content, sulfur content, velocity, temperature,
and materials of construction are the main factors affecting the corrosion process. Despite progress
made in elucidating the role of the different parameters on the crude corrosivity process, the main
problem is in calculating their combined effect, especially when the corroding stream is such a
complex mixture. The TAN is usually related directly to naphthenic acid content. However,
discrepancies between analytical methods and interference of numerous components of the crude itself
lead to unreliable reported content of naphthenic acid. The sulfur compounds, with respect to
corrosivity, appear to be related more to their decomposition at elevated temperature to form hydrogen
sulfide than to their total content in the crude. This paper reviews the present situation regarding
crude corrosivity in distillation units, with the aim of indicating the extent of available information,
and areas where further research is necessary.
Key Words: Naphthenic acid corrosion, crude oil corrosion, high temperature corrosion
INTRODUCTION
The quality of crude oils processed around the world is worsening. The higher densities
of crudes and higher contents of sulfur, acids and other impurities found in crude oils have
increased the likelihood of corrosion failures in the processing plants. Moreover, it is now
realized that many of the old rules of thumb, namely threshold levels of sulfur and total acid
number (TAN) or neutralization number, do not appear to be accurate in their predictive
ability. Analytical tools are needed to predict their high-temperature corrosivity toward
distillation units. Sulfur at a level of 0.2% and above is known to be corrosive to carbon and
low alloy steels at temperatures from 230°C (450°F) to 455°C (850°F). When sulfur is the
only contaminant, McConomy curves [1], with the help of correction factors, are used to
predict the relative corrosivity of crude oils and their various fractions as well as the effect of
225
Corrosion in Refinery and Petrochemical Industries
operational changes on corrosion rates already experienced in the field. However, when
naphthenic acids are present, crude corrosivity prediction becomes more complex. Crudes
with similar compositions and comparable with respect to process considerations have been
found to be entirely different in their impact on corrosion.
MANIFESTATION OF NAPHTHENIC ACID CORROSION
Naphthenic acids are organic acids present in many crude oils from around the world,
especially those from California, Venezuela, the North Sea, Western Africa, India, China, and
Russia. They have the following generic chemical formula: R(CH2)nCOOH, where R is a
cyclopentane ring and n is typically greater than 12. Naphthenic acid corrosion is
differentiated from sulfidic corrosion by the nature of the corrosion (i.e., pitting and
impingement) and by its severe attack at high velocities in crude distillation units. Crude
feedstock heaters, furnaces, transfer lines, feed and reflux sections of columns, atmospheric
and vacuum columns, heat exchangers, and condensers are among the type of equipment
subject to this type of corrosion [2]. Damage is in the form of unexpected high corrosion
rates on alloys that would normally be expected to resist sulfidic corrosion. Isolated, deep
pits in partially filmed areas and/or impingement attack in essentially film-free areas is typical
of naphthenic acid corrosion [3,4]. In many cases, even very highly alloyed materials (i.e., 12
Cr, AISI 316 and 317, and in some severe cases even 6% Mo stainless alloys) have been
found to exhibit sensitivity to corrosion under these conditions.
PARAMETERS AFFECTING NAPHTHENIC ACID CORROSION
TAN (or neutralization number), sulfur content, velocity, degree of vaporization,
temperatures, and alloy composition (i.e., Cr and Mo) were found to be the main factors
affecting the corrosion process.
Total Acid Number
The TAN was related in early studies to the naphthenic acid corrosion rate, and its
threshold was believed to be around 0.5 mg KOH/g [3]. Analysis of crudes’ TAN
distribution as a function of the true boiling point showed where the acids concentrated in the
refinery and resulted in better correlation of TAN with corrosion experienced in the field [5].
However, correlating the TAN of specific cut to their corrosivity was still far from being
reached. Even naphthenic acid weight percent did not correlate to experienced corrosion [6].
A standard laboratory test was developed and showed promising crude corrosivity prediction.
An index called the naphthenic acid corrosion index (NACI) was calculated from the
exposure of a carbon steel coupon at 500°F (260oC) for 48 hours in the fraction to be tested
[7]. The index was the ratio of the corrosion rate of the coupon in mills per year (mpy) to the
weight of its corrosion film in milligrams per square centimeter. It was postulated that a
calculated ratio under 10 indicated sulfidic corrosion while a higher index indicated
naphthenic acid attack. However, this index was developed from static test results, and
correlations to high velocity might not be that simple and need to be verified. Moreover,
carbon steel was the material used in these tests while most of the refineries process
naphthenic crudes with a minimum alloy of 5 Cr.
Sulfur Content
226
Tebbal and Kane
Most crude oil feedstocks vary greatly in both the amount of sulfur and the type of
sulfide species present. It is believed that the sulfur content does not reflect the true effect of
sulfur. A more important factor may be the capability of these sulfur compounds to form H2S
during heating in the refining process [8]. At low temperatures, certain sulfur compounds in
the crude may reduce the severity of naphthenic acid corrosion [9]. In this case, the sulfide
film may offer some degree of protection from the acidic corrosion. At higher temperature
conditions, the presence of naphthenic acids was found to increase the severity of sulfidic
corrosion. Presumably, the presence of these organic acids disrupted the sulfide film, thereby
promoting sulfidic corrosion on alloys that would normally be expected to resist this form of
attack (i.e., 12 Cr and higher alloys). Figures 1 and 2 show the effect of chromium and
molybdenum content as well as TAN and sulfur on the corrosion rate of crude fractions tested
in the laboratory at a velocity of about 3 m/s (10 ft/s) and temperature of 370°C (700°F) for
three days. In the vacuum heater feeder line (VHFL) cut the 1.5 Cr and 5 Cr showed the
same corrosion rate, but in the long resid (LR) fraction the 5 Cr and 9 Cr behaved similarly.
In general, the corrosion rate decreased steadily with increase in Cr and Mo. However, both
TAN and sulfur content do not correlate in any way to corrosion rate. In the VHFL, the
fraction with medium TAN and sulfur was the most corrosive to low alloy steels. In the LR,
the fraction with the lowest TAN and sulfur was the most corrosive. In this case, blending a
high TAN crude to lower the acid content did not provide the expected results and aggravated
corrosion. These cases are common and refinery experience shows that prediction of
corrosion is more complex than it is believed to be.
Temperature
Naphthenic acid corrosion occurs primarily in high velocity areas of crude distillation
units in the 220 to 400°C (430 to 750°F) temperature range. No corrosion damage is found at
temperatures above 400°C (750°F) probably, because of the formation of coke at the metal
surface. The corrosion rate of all alloys of importance to the distillation units increases with
increases in temperature.
Velocity
The flow regime and the degree of vaporization have a significant effect on both sulfidic
corrosion and naphthenic acid corrosion. The higher the acid content generally, the greater
the sensitivity to velocity. In fact, in some cases, it appears possible to obtain very high
corrosion rates even at very low levels of naphthenic acid content (i.e., TAN ≈ 0.3) and low
sulfur content when combined with high-temperature and high velocity.
Materials of Construction
The normal materials of construction used in crude distillation units are carbon steel , 5
Cr, 9 Cr, 410SS, and 316SS [10]. If only sulfur is present and the temperature is above
288°C (550°F), 5 Cr or 12 Cr cladding is recommended for crudes over 1% sulfur when no
operating experience is available [2]. If hydrogen sulfide is evolved, 9 Cr minimum is
preferred. In contrast to high-temperature sulfidic corrosion, low-alloy steels containing up to
12% Cr provide no benefits over carbon steel in naphthenic acid service [1]. With 316 SS
(with 2.5% Mo minimum) or better, with 317 SS with a higher Mo content (3.5% minimum),
cladding of the vacuum column is recommended when TAN is above 0.5 mg KOH/g and in
an atmospheric column when the TAN is above 2.0 mg KOH/g [2].
227
Corrosion in Refinery and Petrochemical Industries
120
Corrosion Rate (mpy)
VHFL I (TAN = 0.35, S = 4.17%)
VHFL II (TAN = 1.64, S = 1.06%)
80
VHFL III (TAN = 0.54, S = 2.09%)
40
0
0
5
10
15
20
25
Cr + Mo (%)
Figure 1. Effect of Cr and Mo content on the general corrosion rate of alloys
in three vacuum heater feeder line oil cuts
50
Corrosion Rate (mpy)
LR I (TAN = 0.35, S = 0.6%, H2S = 2.0%)
40
LR II (TAN = 2.35, S = 0.9%, H2S = 2.0%)
30
20
10
0
0
5
10
15
20
25
Cr + Mo (%)
Figure 2. Effect of Cr and Mo content on the general corrosion rate of alloys
in two long resid oil cuts
MITIGATION METHODS
Mitigation of naphthenic acid corrosion includes blending, inhibition, and materials
upgrading [11]. Blending is the most preferred method and is accomplished by diluting a
high TAN crude with a low TAN one, thus reducing the acid content to a level which
corresponds to an acceptable corrosion attack. Injection of corrosion inhibitors may provide
228
Tebbal and Kane
adequate and economic protection if it is closely monitored and used for specific fractions
that are known to be particularly severe, or if it fluctuates with feedstock quality. When
possible, upgrading the construction materials to a higher chrome and/or molybdenum alloy is
the best solution for long term reliability.
CRUDE CHEMISTRY AND CORROSIVITY CORRELATION
Most of the laboratory studies and refinery experience so far have shown that crude
corrosivity prediction is very complex and that further studies are needed to correlate the
chemistry of crudes to refinery corrosion. Unfortunately, crude chemistry (sulfur and TAN),
process variables (temperature and velocity), and failure analysis (sulfidic and/or naphthenic
acid corrosion) are far from being assessed uniformly throughout the industry. To be able to
achieve correlations among refineries and between laboratories and plants, the measurement
method for each parameter needs to be defined precisely.
Naphthenic Acid Distribution
Naphthenic acid content is generally expressed in terms of TAN. ASTM D974 is a
colorimetric method., with reproducibility of 15% and interference from inorganic acids,
esters, phenolic compounds, sulfur compounds, lactones, resins, salts, and additives such as
inhibitors and detergents. ASTM D664 is a potentiometric method with reproducibility of 20
to 44% depending on the end point (i.e., buffer or inflection), type of oil (i.e., used or fresh),
and titration mode (i.e., automatic or manual), and the same interfering impurities as ASTM
D974. Both ASTM methods do not differentiate between naphthenic acids, phenols, carbon
dioxide, hydrogen sulfide, mercaptans, and other acidic compounds present in the oil. In
addition, the two methods were compared [5], and D664 yielded numbers that were 30 to
80% higher than D974. Thus, prediction of crude corrosivity based on TAN alone could be
misleading. Additionally, for assessment of plant corrosion effects, the naphthenic acid
content needs to be determined for each cut in order to predict exactly where the acids will
concentrate during the distillation of the crude. The isolation and analysis of naphthenic
acids from crude oil may be performed adequately with methods such as UOP 565
(potentiometric) and UOP 587 (colorimetric), chromatographic separations, or other available
analytical techniques [12,13]. The relative abundance of naphthenic acid and its average
molecular weight (i.e., boiling point) may be determined. In addition, the assays of crude
must be current [2]. Once steam flooding or other recovery method is begun in an oil field,
the specific gravity and the organic and sulfur content of the crude can change. Fire flooding,
when used in some fields, tends also to increase the naphthenic acid content.
Hydrogen Sulfide Evolution with Temperature
Sulfur is the most abundant element in petroleum other than carbon and hydrogen. It
may be present as elemental sulfur, hydrogen sulfide, mercaptans, sulfides, or polysulfides.
The total sulfur content is generally analyzed with the ASTM D4294 method using x-ray
fluorescence. Halides and heavy metals interfere with this method. The capability of these
sulfur compounds to form H2S during heating in the refining process, rather than their total
content, is believed to correlate to corrosion in the plants [8]. However, a standard procedure
for determining hydrogen sulfide evolution with temperature is not currently available.
229
Corrosion in Refinery and Petrochemical Industries
Figure 3. Schematic diagram of a rotating autoclave
Figure 4. Schematic diagram of a jet impingement apparatus
Wall Shear Stress
Fluid flow velocity has long been used as the parameter for comparing flow among
refineries and between laboratory and field. However, this concept was found to lack
predictive capabilities and was replaced by data related to fluid flow parameters such as wall
shear stress and Reynolds number [14]. Wall shear stress, rather than velocity, is the
parameter directly proportional to corrosion through the removal of normally protective films.
The wall shear stress in the field is proportional to (1) the density and viscosity of the liquid
and vapor in the pipe at temperature, (2) the degree of vaporization in the pipe, and (3) the
pipe’s diameter. Wall shear stress in the laboratory depends on the geometry and dimensions
230
Tebbal and Kane
of the laboratory apparatus. Figure 3 shows a schematic diagram of a rotating autoclave with
a condenser for the return of light components, and Fig. 4 is a schematic diagram of a jetimpingement laboratory setup. Both of these apparatus are used in the laboratory to simulate
corrosion in the field at high velocities. Table 1 compares the shear stress level between the
field and the laboratory. The results show that the wall shear stress changes drastically with
the degree of vaporization, and that identical fluid flow in the laboratory and the field do not
correspond to the same level of shear stress.
Table 1. Wall Shear Stress (in Pascals) Calculated at Different Velocities
(VGO cut of 0.7 Specific Gravity and 0.55 Centistokes Viscosity
at 700oF)
Velocity
3 m/s
(10 ft/s)
Rotoclave
Jet Impingement
18
4
DV = 0%
46
DV = 30%
32
DV = 70%
14
DV = Degree of vaporization
6 m/s
16 m/s
(20 ft/s)
(50 ft/s)
Laboratory Setup
58
N/A
13
80
Plant
158
691
111
514
47
286
33 m/s
(100 ft/s)
66 m/s
(200 ft/s)
N/A
298
N/A
1049
2830
2080
862
10291
5148
2210
Table 2. Effect of Velocity on General and Pitting Corrosion Rates
Alloy
General Corrosion Rate (mpy)
Pitting Corrosion Rate (mpy)
at 10 ft/s
at 200 ft/s
at 10 ft/s
at 200 ft/s
5 Cr
21.8
25.5
0.0
201.1
9 Cr
20.3
24.2
0.0
191.8
317 SS
3.29
6.09
0.0
28.7
Laboratory Testing
Laboratory studies are directed at the simulation of field conditions under controlled and
reproducible conditions. CLI International, Inc. is currently involved in a major multiclient
effort to provide more systematic understanding and a methodology for handling crude
corrosivity and naphthenic acid corrosion issues. The interpretation of laboratory results and
their correlation to the plant need to be analyzed carefully. The temperature of fluid and
specimens in laboratory studies most likely are equal. However, in furnaces and heat
exchangers of crude distillation units, temperature differences between the stream and the
metal skin may be as high as 85 to 100°C (150 to 200°F) [1]. Rates of corrosion found in
laboratory testing may correspond to the maximum corrosion rates found in the field. This
usually results from the short test duration in the laboratory. The corrosion rate is usually
high initially and then decreases with time because of the formation of protective films. The
laboratory corrosion rates may also be much lower than those experienced in the field if the
composition of the test solution changes with time as a result of degradation of its corrosive
components. The type and rate of corrosion may be easily calculated in the laboratory. Table
2 shows the effect of velocity on the general and localized corrosion rates of three alloys after
231
Corrosion in Refinery and Petrochemical Industries
three days of exposure. Differences in flow velocity on the general corrosion of the three
alloys was minimal. However, a significant increase in the pitting corrosion rate with
increases in flow from 10 ft/s to 200 ft/s was found especially for the 5 Cr and 9 Cr alloys.
Corrosion rates in the field are evaluated by on-line monitoring tools which indicate only
general corrosion rates unless the equipment is inspected for pitting and/or impingement.
CONCLUSIONS
It has clearly been proven through extensive laboratory and plant studies that predicting
crude corrosivity by using the general rules of total acid and sulfur content is not reliable
especially with the wide range of crude oil feedstocks being processed today. Naphthenic
acid content and distribution in side cuts, hydrogen sulfide evolution with temperature, wall
shear stress, temperature at the metal surface, and materials of construction are the main
factors affecting the crude corrosivity process. The exact mechanisms which are operating
are not precisely known at this point, and much research and testing is necessary to build a
more comprehensive understanding. Based on the complexity of the situation and the current
level of understanding, each case must be dealt with on an individual basis until a more
comprehensive methodology for assessment can be developed. It is possible, however, to
provide a practical assessment of the plant corrosion process by establishing a more
comprehensive database from both laboratory and field experience where the various
parameters affecting naphthenic acid corrosion can be more extensively and unambiguously
defined and quantified. This information will serve as a firm basis for materials selection
decisions, feedstock blending requirements and plant operating conditions.
REFERENCES
1. L. Garverick, Ed., Corrosion in the Petrochemical Industry, ASM International, 1994.
2. R.A. White and E.F. Ehmke, Materials Selection for Refineries and Associated Facilities,
NACE, Houston, Texas, 1991.
3. W.A. Derungs, Corrosion 12, 12, 1956, p. 617t.
4. J. Gutzeit, Materials Performance 16, 10, 1977, p. 24.
5. R.L. Piehl, Materials Performance, January 1988, and Paper No. 196, Corrosion/87,
NACE.
6. E. Babian-Kibala, et al., Naphthenic acid corrosion in a refinery setting, NACE
Conference, Corrosion/93, Paper No. 631, 1993.
7. H.L. Craig, Naphthenic Acid corrosion in the refinery, Paper No. 333, Corrosion/95,
NACE.
8. R.L. Piehl., Corrosion, June 1960, p. 305t.
9. Heller, Materials Protection, September 1963.
10. F. Blanco and B. Hopkinson, Experience with naphthenic acid corrosion in refinery
distillation process units, NACE Conference, Corrosion/83, Paper No. 99, 1983.
11. G.L. Scattergood and R.C. Strong, Naphthenic acid corrosion: An update of control
methods, NACE Conference, Corrosion/87, Paper No. 197, 1987.
12. Tseng-Pu Fau, Energy and Fuels 5, 3, 1991, p. 371.
13. I. Dzidic, et al., Analytical Chemistry 60, 13, July 1, 1988, p. 1318.
14. K.D. Effird, et al., Experimental correlation of steel corrosion in pipe flow jet
impingement and rotating cylinder laboratory tests, Corrosion/93, Paper No. 81, NACE.
232
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
NEW NICKEL ALLOYS SOLVE CORROSION
PROBLEMS OF VARIOUS INDUSTRIES
D.C. Agarwal1 and W.R. Herda2
1
VDM Technologies
11210 Steeplecrest, # 120, Houston, Texas 77065, USA
2
Krupp-VDM GmbH
P.O. Box 1820, 58778 Werdohl, Germany
ABSTRACT
The materials of construction for the modern chemical process and petrochemical industries not only
have to resist uniform corrosion caused by various corrodents, but must have sufficient localized corrosion
and stress corrosion cracking resistance as well. These industries have to cope with both the technical and
commercial challenges of rigid environmental regulations, the need to increase production efficiency by
utilizing higher temperatures and pressures, using more corrosive catalysts, and at the same time possess
the necessary versatility to handle varied feed stock and upset conditions. Over the past 30 years
improvements in alloy metallurgy and a fundamental understanding of the role of various alloying
elements has led to newly developed Ni-Cr-Mo and Ni-Mo alloys, which not only extend the range of
usefulness of existing alloys by overcoming their limitations, but are also cost-effective and open new
avenues of applications. This paper presents the various metallurgical, thermal stability and corrosion
resistance characteristics of some newer alloys along with actual case histories, where these alloys have
solved specific problems.
Key Words:
Nickel alloys, chemical process industries, petrochemical applications,
localized corrosion, uniform corrosion, acid corrosion, Alloy 59, Alloy 31,
Alloy 33, Alloy B-2, Alloy B-4
INTRODUCTION
At one time one, of the major factors in any material selection was initial cost with little
thought given to maintenance and the cost associated with lost production due to unscheduled
equipment downtime. In today's economic environment, increased maintenance costs and
downtime have placed a greater emphasis on the need for reliable, safe and versatile performance
of process equipment.
Prior to the 1950's, the alloy choices to combat corrosion were very limited. The latter half
of the 20th century saw a phenomenal growth in the development of new nickel-based alloys due
to improved melting and thermo-mechanical processing innovations, a better fundamental
understanding of the role of various alloying elements and their effect on both corrosion and
physical metallurgy behavior. Even though the standard austenitic stainless steels (Alloy 304L
and 316L) have been and continue to be the workhorse of many industries, their vulnerability to
localized corrosion and chloride stress corrosion cracking (SCC) has been a major problem in
233
Corrosion in Refinery and Petrochemical Industries
many chemical processes. The knowledge that chromium and molybdenum improved the
localized corrosion resistance and increased nickel enhanced the chloride SCC resistance led to
different alloys with varying nickel, chromium, molybdenum and iron contents. The knowledge
that in certain low nickel containing alloys nitrogen could be added to impart certain unique
mechanical, metallurgical and corrosion characteristics was used to come up with a completely
new 6% Mo alloy class which was very cost-effective and in certain cases approached or equaled
the corrosion resistance of the more expensive high nickel containing alloys. Two alloys of this
6% Mo stainless steel class, with varying chromium and nickel contents, were developed to
bridge the performance gap between the standard austenitic stainless steel and the very high
performance nickel-based alloy of the Ni-Cr-Mo family such as Alloy C-276, Alloy C-22 and the
most recent and advanced development, Alloy 59. Table 1 gives some of the alloy groupings of
the materials of construction used today, whereas Table 2 lists the effects of the various alloying
elements in the Ni-Cr-Mo-containing alloys. As one moves upward from Type 304 stainless
steel metallurgy (Group I of Table 1) to higher alloys groupings such as Groups II, IV and VI,
corrosion resistance improves, as evidenced by their chemical composition and the higher
calculated pitting resistance equivalent (PRE) shown in Table 3. This higher PRE accounts for
the improved corrosion resistance of the various alloys in a variety of environments. For
interested readers, a list of references is provided for further in depth reading to gain a better
understanding of these newer developments [1-16].
Table 1. Some Austenitic Corrosion Resistant Alloys for Combating Aqueous Corrosion
Alloy Group
I
II
III
IV
V
VI
Generic Description
Iron-based 18-8 austenitic SS alloys
High performance austenitic SS alloys
Ni-based general purpose alloys
6% Mo superaustenitic SS alloys
Ni-based special alloys
Nickel based high performance alloys
VII
Chromium based high performance
wrought super austenitic SS
* Newer developments in the last 6 years.
Typical Alloys
304, 316, 317
904L, 20, 28, 825
200, 400, 600, 800
1925hMo, 31*, 254SMo,
B-2, B-3*, B-4*
G-3, G-30, 625, C-276,
C-4, C-22, 59*, 686*
33*
C-FAMILY Ni-Cr-Mo ALLOYS
Alloy C, the oldest alloy of this family (now obsolete), was superseded by Alloy C-276 in
the early 1960's, due to improvements in melting technology. Between 1983 and 1994, three new
alloys of this family were introduced to the market place: Alloy C-22 in the mid 1980's, Nicrofer
5923hMo-Alloy 59-in the late 1980's, and Alloy 686 in the early 1990's. Alloy 59 has the highest
pitting resistance equivalent and the lowest iron content (Table 3), which provides for improved
corrosion resistance over other alloys in a variety of standard laboratory environments, as shown
in Table 4. Eliminating tungsten and reducing the iron content to very low levels, formulated an
alloy with superior thermal stability characteristics, as shown in Table 5. This data clearly shows
the detrimental effects of tungsten on the thermal stability of the various alloys, all of which
234
Agarwal and Herda
contain tungsten, except Alloy 59. Not only is the uniform corrosion behavior and the thermal
stability improved, its localized corrosion resistance as measured by the standard ASTM G48 test
method in 10% FeCl3 is also enhanced. Table 6 clearly shows the beneficial effects of higher
PRE due to the highest Cr plus Mo content in Alloy 59. Further details on Ni-Cr-Mo alloys are
provided elsewhere [1-3,10,11,14,15,16].
Applications of Alloy 59
Due to its highest PRE and these unique corrosion-resistant properties, Alloy 59 has found a
number of successful applications, where other nickel-based alloys have been either inadequate
or marginal in performance. Some of these applications are described below.
Pollution Control
Combustion gases from burning fossil fuels or from waste incineration of municipal or
hazardous waste contain acidic pollutants such as SO2, SO3, HCl and Nox, which must be
scrubbed before the gases are released into the atmosphere. The equipment most frequently
employed to achieve this is a wet scrubber. Operating conditions in critical sections of wet
scrubbing systems can be extremely severe: for example, in condensates chloride levels may
approach 100,000 ppm at pH values below 1 and temperatures of 80°C. Laboratory and field
testing has shown Alloy 59 to be one of the very few metallic materials able to withstand such
aggressive corrosive conditions. In one lignite fired power station in Germany, 40 month field
test rack data showed Alloy 59 to be the only alloy free of any localized attack.
Table 2. Alloying Elements and Their Major Effects
Alloying
Element
Ni
Cr
Mo
W
N
Cu
Ti, Cb, Ta
Fe
Main Features and Benefits
Provides metallurgical compatibility to various alloying elements.
Improves thermal stability and fabricability. Enhances corrosion
in mildly reducing and alkali media, and improves chloride SCC.
Provides resistance to oxidizing corrosive media. Enhances
localized corrosion resistance.
Provides resistance to reducing (nonoxidizing) corrosive media.
Enhances localized corrosion resistance and chloride SCC.
Provides solid solution strengthening.
Behaves similar to Mo, but is less effective. Is detrimental to
thermal stability. Provides solid solution strengthening.
Austenitic stabilizer-economical substitute for nickel. Enhances
localized corrosion resistance, thermal stability and mechanical
properties.
Improves resistance to seawater. Enhances resistance to H2SO4
and HF containing acid environments.
Carbon stabilizers. Improves HAZ corrosion resistance.
Provides matrix for metallurgical compatibility to various
alloying elements. Enhances resistance to oxidizing corrosive
235
Corrosion in Refinery and Petrochemical Industries
media. Reduces cost by replacing nickel, and enhances scrap
utilization.
Table 3. Typical Chemical Composition of Some Ni-Cr-Mo Type Alloys
UNS #
Alloy
Ni
Cr
Mo
Fe
Others
S30400
304
8
18
72
18
S31603
316L
12
17
2.3 66
24
N08904
904L
25
21
4.8 48
Cu
N08020
20
38
20
2.4 34
Cu, Cb
N08825
825
40
22
3.2 31
Cu
N08028
28
31
27
3.5 36
Cu, Cb
N08926
1925hMo
25
21
6.5 46
Cu, N-0.2
N08031
31**
31
27
6.5 32
Cu, N-0.2
N06985
G-3
48
23
7
20
Cu, Cb
R20033
33**
31
33
1.6 32
Cu, N-0.4
N06625
625
62
23
9
3
Cb
N10276
C-276
57
16
16
5
W
N06022
C-22
57
22
13
3
W
N06686
686**
56
21
16
2
W
N06059
59**
59
23
16
1
* PRE = Pitting Resistance Equivalent = % Cr + 3.3 (% Mo) + 30 N
** Recent Alloy Developments
PRE*
37
29
32
38
48
54
45
50
52
69
65
74
76
Table 4. Comparison of Some Ni-Cr-Mo Alloys in Various Boiling Corrosive Environments
Alloy
Media
C-276
ASTM 28A
168
ASTM 28B
55
Green Death
26
10% HNO3
19
65% HNO3
750
10% H2SO4
23
50% H2SO4
240
1.5% HCl
27
10% HCl
239
10% H2SO + 1% HCl
87
10% H2SO4 + 1% HCl (90°C)
41
* To convert to mm/y, multiply by 0.0254.
236
Uniform Corrosion Rate (mpy)*
Alloy
Alloy
C-22
686
36
60
7
12
4
8
2
52
231
18
<5
308
34
392
354
92
-
Alloy
59
24
4
5
2
40
8
176
15
179
70
3
Agarwal and Herda
Table 5. Thermal Stability Per ASTM G28B After Sensitization Treatment at 1600°F
**
Corrosion Rate (mpy)*
Alloy 22**
Alloy 686**
339
17
313
85
1000
Not tested
Sensitization (hr)
Alloy C-276
Alloy 59***
1
>1000
4
3
>1000
4
5
>1000
17
* To convert to mm/y, multiply by 0.0254
** Alloys C-276, 22 and 686 - Heavy pitting attack with grains falling due to deep intergranular
attack.
*** Alloy 59 - No pitting attack.
Table 6. Critical Pitting and Crevice Corrosion Temperature Per ASTM G-48
Alloy
Critical Pitting Corrosion
Critical Crevice Corrosion
Temperature (oC)
Temperature (oC)
316
15
<0
904L
45
25
u20
15
<10
825
30
<5
G-3
70
40
1925hMo
70
40
625
77.5
57.5
33
85
40
31
>85*
65
22
>85*
58
C-276
>85*
>85*
686
>85*
>85*
59
>85*
>85*
* At temperatures exceeding 85°C, 10% FeCl3 chemically breaks down.
PRE
24
32
29
32
45
38
52
50
54
65
69
74
76
The municipal incinerator of Essen-Karnap in Germany originally had a scrubber with a
rubber lining, installed in 1987. After some 20,000 hours of service, the rubber lining failed,
when liquid permeated it and attacked the underlying carbon steel substrate. A decision to install
a metal lining was taken end of 1991. Following extensive laboratory and field tests in other
plants belonging to the same owners, Alloy 59 was selected for this lining. In 1992, 55 tons of
alloy 59 were used for this project. Examination after two years of operation revealed no
detectable loss of thickness or any localized attack of Alloy 59. A further advantage is that the
quantity of deposits retained on the lining of the absorber was reduced by a factor of one
237
Corrosion in Refinery and Petrochemical Industries
thousand, which significantly reduced periodic cleaning costs. Many hundreds of tons of Alloy
59 have been ordered in recent years for flue gas desulphurization systems of both power stations
and incinerators throughout the world [14].
Synthesis of Acrylates and Methacrylates
One process for the synthesis of acrylic or methacrylic esters involves reacting the
corresponding acids with fatty alcohols, in the presence of para-toluene sulphonic acid as a
catalyst. The reaction temperature is 130°C, and the reaction is carried out under oxidizing
conditions. Heating is by an internal steam coil. Following rapid failure of the material
previously used for the steam heating coil, a series of plant tests was made with alloys including
904L, 28, G-3, 625, C-276, 31, and 59. The only alloy, which showed no pitting or crevice
corrosion, and a corrosion rate of less than 0.01 mm/yr, was Alloy 59. A steam heating coil made
of Alloy 59 was installed in 1993 and has operated without any problems ever since.
Aluminum Refining
When aluminum scrap is remelted, the molten metal is protected from oxidation by a layer
of sodium and potassium chlorides. During the refining process this salt layer becomes
contaminated with ammonium chloride. These chloride salts then have to be purified and
recovered. This is done by dissolving them in water, and then recrystallizing the solution. In one
European plant, the solution thus obtained contains 20-25% NaCl, 6-8% KCl and 5-8% NH4Cl.
The pH is in the range 4.5 to 6. The evaporator operates at a temperature of 107°C. The initial
plant was built in rubber-lined steel, and failed rapidly due to cracking of the rubber lining and
subsequent corrosion of the underlying carbon steel. A plant test in 1994 with Alloy 59 showed
that after some 3800 hours of operating time, no corrosion could be detected. The
recystallization plant has since been rebuilt in Alloy 59.
Metals Processing
In a copper plant, the SO2-rich gas from the flash furnace is scrubbed with a solution of 5%
H2SO4 at a temperature of 45°-60°C. The acid produced has a concentration of typically around
50-55% H2SO4 and a temperature of about 75°C. The chloride and fluoride contents of this acid
are both high, at about 7000 ppm. Tests were carried out using both Alloy 59 and Alloy 31.
Corrosion rates for both alloys were below 0.013 mm/yr with no localized corrosion. Following
these tests, Alloy 31 was purchased for the scrubber internals handling the produced acid, and
Alloy 59 for the induced draft fans. These have been in successful operation for the last two
years with no detectable corrosion. Since then another order has been placed.
Citric Acid Production
Citric acid is produced in one plant by reacting calcium citrate with 95%-98.5% sulphuric
acid at 95°C-97°C. A pilot installation of Alloy 254SMo failed rapidly. A three month test with
Alloy 59 gave a corrosion rate of 0.05 mm/yr. The first of four reactors made of Alloy 59 was
installed in 1990, and continues to operate well with no problems.
Effluent Treatment
Effluent from an acetic acid derivatives plant is cooled in Alloy C-276 plate heat
exchangers. These require frequent replacement. Initial tests suggested that Alloy 59 might be a
better alternative, so more extensive testing was carried out. The test conditions, and the
238
Agarwal and Herda
corrosion rates observed during testing led to the selection of Alloy 59 to replace Alloy C-276 for
the new effluent treatment plant.
Fine Chemicals Production
At one major chemical company, the production of fluorinated organic chemicals requires a
halogen exchange reaction in which one fluorine atom is substituted for chlorine in the molecule.
This reaction is carried out at about 100°C in the presence of ammonium fluoride and a
proprietary catalyst. Because of the severely corrosive conditions, extensive tests were made with
Alloy 59 and with other alloys of the Ni-Cr-Mo and Ni-Mo family. The lowest corrosion rate
was exhibited by Alloy 59. A 2600 US-gallon reactor (9.8m3) (Fig. 1) was built to ASME code
requirements and has been giving excellent performance over the past 24 months. It is expected
that a life cycle cost analysis will show Alloy 59 to be at least 50% cheaper than the next best
candidate alloy belonging to the Ni-Mo family, Alloy B-2. Due to the excellent performance of
Alloy 59, another bigger ASME vessel has now been ordered by the same chemical company
(4,000 gallons capacity).
Figure 1. ASME code vessel made of Alloy 59 producing chlorinated and fluorinated chemicals
Alloy 625 was giving only three years life in a column in a fine chemicals plant.
The operating conditions were a temperature of 140°C and a medium consisting of 83.1% water,
14.3% sodium bisulphate, 0.34% sodium sulphate, 0.02% acetone, 0.46% isopropanol, 0.06%
copper sulphate, 0.04% DCNB, and 1.5% various organics. Tests were carried out both at the
inlet to the column, and at the foot of the column. Based on the results of these tests, an inquiry
was issued for a new column to be built in Alloy 59.
239
Corrosion in Refinery and Petrochemical Industries
There are many other applications of Alloy 59 which continue to find increasing usage and
specification in the various industries throughout the world. This alloy is covered under
appropriate ASTM, AWS and ASME specifications.
6% Mo ALLOYS
These alloys, such as Alloy 1925hMo, were derived from alloy 904L metallurgy by
increasing the molybdenum content from 4.5% to 6.5% and fortification with 0.2% nitrogen.
This addition of nitrogen provided added benefits of improved localized corrosion resistance,
thermal stability and mechanical properties. These alloys are readily weldable with over-alloyed
filler metals, such as Alloy 625, Alloy C-276 or Alloy 59 to compensate for segregation of
molybdenum occurring in the interdentritic regions of weldments. A higher chromium-nickel
version of these alloys known as Alloy 31 further improves the corrosion resistance
characteristics in a variety of media. Its corrosion resistance in sulfuric acid in medium
concentration range is superior to even that of Alloy C-276 and Alloy 20 (Table 7). However,
one must be careful, when specifying this alloy for higher concentrations and temperatures. At
80% concentration and temperatures above 80°C, Alloy 31 exhibits active behavior. The 6Mo
alloys have found extensive usage in pulp and paper, phosphoric acid, copper smelters, sulfuric
acid production, pollution control, rayon production, specialty chemicals production, marine and
offshore applications, heat exchangers using seawater and brackish water as coolant, pickling
baths and many other applications. These alloys are covered under appropriate ASTM and
ASME specifications. More details on these 6Mo alloys are presented elsewhere [4-8]. Some
applications of 6Mo alloys are presented in Tables 8 and 9.
Table 7. Corrosion Resistance in Sulfuric Acid
60°C
H2SO4
Alloy Alloy Alloy
Alloy
(%)
20
C-276 31
20
<5
<1
<0.1
40
<5
<2
<0.1
60
>5
<2
<0.1
80
5
<1
0.2
To convert to mm/y multiply by 0.0254
Corrosion Rate (mpy)
80°C
Alloy Alloy Alloy
100°C
Alloy Alloy
20
10
10
11
18
20
>25
>25
>50
>50
C-276
4
3
4
15
31
<0.1
<0.2
0.4
0.8
C-276
>1
10
11
240
31
0.3
0.6
1
249
Table 8. Some Typical Industrial Applications for Alloy 1925hMo
Offshore/Marine
Seawater lines
Product coolers
Reverse osmosis
Desalination plants
Chemical Process Industry - Organics
240
FGD
Pulp/Paper
Scrubbers
Bleach washers
Fans
Pulp lines
Ducts
Recovery boiler scrubbers
Dampers
Chemical Process Industry -Inorganics
Agarwal and Herda
Ethyl acetate production
Thermoplastic rubber (catalyst strippers)
TDI and MDI production
Organic intermediates (chloride catalysts)
Fine chemicals (pharmaceuticals, agrochemicals)
H2SO4 distribution systems
Metasilicate production
Sodium perchlorate crystallizers
Hydrofluoric acid producton scrubbers
Ammonium chloride evaporators
Catalyst strippers
Wet exhaust fans
Phosphoric acid digestion systems
Table 9. Some Industrial Applications of Alloy 31
Chlorine dioxide bleach washers
Acid pickling industries
acid plants
Waste water reclamation from uranium ore leaching process
Phosphoric acid production
Heat exchangers in chloride media/seawater/brackish water
Dampers
Sulfuric
Fine chemicals
Mist eliminators
Many others
ALLOY 33
Alloy 33 the most recent innovation [12,13], is a chromium-based, fully austenitic wrought
super stainless steel (33 Cr, 32 Fe, 31 Ni, 1.6 Mo, 0.6 Cu, 0.4 N). This alloy has excellent
resistance to both acidic and alkaline corrosive media, mixed HNO3/HF acids, localized
corrosion and stress corrosion cracking. Due to its high nitrogen content, this alloy has excellent
mechanical properties. Its high PRE (Table 3) makes it a very cost-effective alloy in comparison
to the Ni-Cr-Mo alloys such as G-3, G-30 and 625. Its localized corrosion resistance is equal to
or better than some of the Ni-Cr-Mo alloys (Table 6).
Table 10 shows some of the corrosion resistance data of Alloy 33 in various corrosive
environments. In comparison to other high chromium alloys such as Alloys G-30, 690 and 28,
Alloy 33 shows excellent corrosion resistance behavior. This alloy on a cost/performance basis
has the potential of being an excellent alternative to many alloys currently in use such as the 825,
904L, 20, 28, 6Mo alloys, G-3, G-30 and in some cases even Alloy 625. Some of the testing
done with Alloy 33 in sulfuric acid environments, nitration involving nitric acid, phosphoric acid,
acid pickling, nuclear waste reprocessing, and the pulp and paper industry, has shown it to give
excellent results. The alloy has performed exceptionally in concentrated sulfuric acid at high
temperatures as shown by the data in Table 10. Many companies are seriously considering using
this alloy in their various processes.
Table 10. Corrosion Resistance of Alloy 33 and Others in Various Media
Media
Sulfuric Acid
98% H2SO4
Temperature
100°C
150°C
200°C
Corrosion Rate (MPY)
Alloy
Alloy
Alloy
33
A611
28-4-2
1.6
0.8
1.2
3.2
32
21
1.6
24
2.8
241
Corrosion in Refinery and Petrochemical Industries
Alloy
Alloy
Alloy
Alloy
Phosphoric Acid
Temperature
33
G-30
28
690
85% H3PO4
100°C
8
12
8
50
154°C
43
53
56
Alloy
Alloy
Alloy
Alloy
Mixed Acid
Temperature
33
G-30
28
690
12% HNO3 + 0.9% HF 90°C
10
11
230
24
12% HNO3 + 3.5% HF 90°C
48
48
>500
252
32% HNO3 + 0.4% HF 90°C
11
20
38
58
56% HNO3 + 0.4% HF 90°C
66
96
135
187
Table 11. Chemical Composition of Ni-Mo Alloys
Alloy
(UNS No.)
B (N10001)
B-2 (N10665)
B-3 (N10675)
B-4 (N10629)
B-4(C)
(A)
Maximum
Decade
Introduced
Ni
1920s
Bal.
1970s
Bal.
1990s
65(B)
1990s
Bal.
1990s
Bal.
(B)
Minimum
Composition (wt. %)
Mo
Fe
Cr
26-30
4-6
1(A)
(A)
26-30
2
1(A)
27-32
1-3
1-3
27-30
2-5
0.5-1.5
28
3
1.3
(C)
Typical composition
C
0.05(A)
0.01(A)
0.01(A)
0.01(A)
0.006
Table 12. Laboratory Test Results and Potential Hazards
Ni-Mo Grade
Standard Alloy B-2
UNS N10665
Fe (2% Maximum)
Cr (1% Maximum)
Test Results
Tensile El (%)
ASTM G 30
at 700°C after
U-Bend 1 h, 700°C
1 h (700°C)
10% H2SO4 for 100 h
5
Failure
Potential Hazards
Cracking
SCC
during
during
fabrication
service
High risk
High risk
Controlled-Chemistry
Alloy B-2
UNS N10665
Fe (1.6-2.0%)
Cr (0.5-1.0%)
42
Reduced risk
None
Alloy B-4
UNS N10629
Fe (2-5%)
Cr (0.5-1.5%)
48
Significantly
reduced risk
None Significantly
reduced risk
B-FAMILY Ni-Mo ALLOYS
242
Reduced risk
Agarwal and Herda
Alloy B, the original alloy in the Ni-Mo family, developed in the 1920's, suffered from HAZ
corrosion in nonoxidizing acids (i.e., acetic, formic and hydrochloric) due to its higher carbon
content. In the decade of the 1960's, improved AOD melting technology led to development of
Alloy B-2. This alloy solved the HAZ problem, but suffered from poor fabricability. Recent
developments of controlled chemistry Alloy B-2 (Table 12) and, Nimofer(R) 6629 - Alloy B-4 UNS N10629 - (Table 11 and Table 12) solved both these problems by eliminating/reducing
formation of detrimental intermetallic phases, with further improvement in corrosion resistance
behavior. Table 11 gives the basic developments of the alloys in the Ni-Mo family, and Table 12
shows the improvements in fabricability and corrosion resistance. Greater details on fundamental
behavior and understanding of Ni-Mo alloy systems are presented elsewhere [1,9,10]. Alloy B-2
has been successfully used in production of acetic acid, pharmaceuticals, alkylation of ethyl
benzene, styrene, cumene, organic sulfonation reactions, melamine, herbicides, and many other
products. Alloy B-4, the improved version of alloy B-2, is being tested and considered for
various applications in many other industries.
CONCLUSIONS
The understanding of the role of alloying elements in nickel-based alloys has led to the
recent innovations in both high and low nickel-containing Ni-Cr-Mo alloys. Higher
molybdenum and chromium contents, together with nitrogen fortification, have opened up an
entirely new class of 6Mo alloys with unique properties, and has provided some specific benefits
to industry. A better understanding of the physical metallurgy of Ni-Cr-Mo and Ni-Mo alloy
systems has further contributed to the development of the newer improved alloys, which has
helped to expand the range of usefulness of existing alloys for the modern chemical process and
other industries requiring higher performance levels in the present day aggressive environments.
Alloy 59 many successful applications and its ever increasing specifications in various industries
clearly proves the above facts. A new chromium-based fully wrought super stainless steel (Alloy
33) and the two alloys of the 6Mo family (Alloy 1925hMo and Alloy 31) show excellent promise
of solving many corrosion problems of today's chemical process and other industries in a costeffective manner. The newest alloy of the Ni-Mo family, Alloy B-4 and the controlled chemistry
Alloy B-2 have not only solved the fabricability problems associated with the old Alloy B-2, but
have also improved upon its corrosion resistance behavior.
REFERENCES
1. W.Z. Friend, Corrosion of Nickel and Nickel Base Alloys, John Wiley & Sons, Inc., New
York, 1980, pp. 248-367.
2. R. Kirchheiner, M. Kohler and U. Heubner, A new highly corrosion resistant material for
the chemical process industry, flue gas desulfurization and related applications,
Corrosion/90, Paper No. 90, NACE International, Houston, Texas, USA, 1990.
3. D.C. Agarwal, et al., Cost effective solution to CPI corrosion problems with a new Ni-CrMo alloy, Corrosion/91, Paper No. 179, NACE International, Houston, Texas, USA,
1991.
4. D.C. Agarwal, M.R. Jasner and M.B. Rockel, 6% Mo austenitic stainless steel selection
for offshore applications, Offshore Technology Conference 1991, Paper No. 6598,
Houston, Texas, USA, 1991.
243
Corrosion in Refinery and Petrochemical Industries
5. D.C. Agarwal, et al., The 6% Mo superaustenitics: The cost effective alternative to nickel
alloys, Proceedings first Pan American Corrosion and Protection Congress, Mar del Plata,
Argentina, November 25-30, 1992, Vol. 1, pp. 103-114.
6. M. Rockel and M. Renner, Pitting, crevice and stress corrosion resistance of high
chromium and molybdenum alloy stainless steels, Werkstoffe und Korrosion 35, 1984, p.
537.
7. CroniferR 1925hMo, An advanced high alloy austenitic stainless steel for offshore
hydrocarbon and seawater applications, VDM Report No. 10/2, July 1992, Krupp-VDM,
Werdohl, Germany.
8. U. Heubner, R. Kirchheiner and M. Rockel, Alloy 31-A new high alloyed Ni-Cr-Mo steel
for the refinery industry and related applications, Corrosion/91, Paper No. 321, NACE
International, Houston, Texas, USA, 1991.
9. D.C. Agarwal, U. Heubner, M. Köhler and W. Herda, UNS N10629: A new Ni-28% Mo
alloy, Materials Performance 33, 10, 1994, pp. 64-68.
10. D.C. Agarwal, U. Heubner and W.R. Herda, Fundamental considerations during
fabrication and construction of nickel alloy components for CPI, Conference
Proceedings, First International Symposium on Process Industry Piping, December 1417, 1993, Orlando, Florida, USA.
11. Corrosion characteristics and applications of newer high and low nickel containing NiCr-Mo alloys, Conference Proceedings, 12th International Corrosion Congress,
September 19-24, 1993, Houston, Texas, USA.
12. M. Kohler, U. Heubner, K.W. Eichenhofer and M. Renner, Alloy 33, A new corrosion
resistant austenitic material for the refinery industry and related applications,
Corrosion/95, Paper No. 338, NACE International, Houston, Texas, USA, 1995.
13. M. Kohler, et al., Progress with Alloy 33, A new corrosion resistant chromium-based
austenitic material, Corrosion/96, Paper No. 428, NACE International, Houston, Texas,
USA, 1996.
14. D.C. Agarwal, Alloy selection methodology and experiences of the FGD industry in
solving complex corrosion problems: The last 25 years, Corrosion/96, Paper No. 447,
NACE International, Houston, Texas, USA, 1996.
15. D.C. Agarwal and W.R. Herda, Alloying effects and innovation in nickel base alloys for
combating Aqueous corrosion, VDM Report No. 23, February 1996, Krupp-VDM,
Werdohl, Germany.
16. F.E. White, G.K. Grossmann, H. Decking and D.C. Agarwal, Experience with the Use of
Alloy 59 in Industrial Applications, Corrosion/96, Paper No. 433, NACE International,
Houston, Texas, USA, 1996.
244
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
MACRO-MICRO SEGREGATION BANDS (MMB) AS A MAIN FACTOR
INFLUENCING STEEL APPLICABILITY FOR THE PETROLEUM INDUSTRY
A. Mazur
The Academy of Mining and Metallurgy
Al. Mickiewicza 30, Krakow, Poland
ABSTRACT
From a corrosion engineering point of view, in spite of common opinion, correctly welded steel
joints are not the weakest points of reliable constructions. The results of long-term tests have shown
the leading role of MMB structure in corrosion crack initiation and propagation. Such a band
structure often exists in weldable steel sheets selected for pressurized systems in the petroleum
industry. A practical aspect of the results reported is connected with steel quality recommendations
for reliable construction. The lack of an MMB structure or an MMB width limited to below 0.01 mm
[1] should guarantee the steel’s resistance to hydrogen-induced cracking.
Key Words: Segregation band, corrosion cracking, line-pipe steels, branching effects,
hydrogen sulfide
INTRODUCTION
The pressurized installations of the petrochemical industry belong to those most reliable
constructions that minimize the danger of catastrophic failure, sometimes of large-scale
losses. The probability of material failures is (in a simple way of thinking) proportional to the
total length of the pressurized system, e.g., piping-net, pressure vessels, etc. Of course, it also
depends on inspection system’s efficiency. Therefore, many petrochemical plants have been
compactly built to minimize the length of piping and energy losses.
However, the history of failure analysis shows that an ideal inspection system does not
exist in spite of serious improvement in testing equipment and maintenance. Thus, the
probability of catastrophic failures increases with plant operation time, as is shown in Fig. 1.
In the past, petrochemical plants were constructed based on the burst-before-leak (BBL)
concept. This concept is no longer acceptable because any of the pressurized installations can
produce a chain-like reaction damaging sometimes costly and vital parts of a compactly built
plant. Many such-constructed plants are still in operation. The accumulation of information
on the environment-material - stress relationships shows that the BBL concept is no more in
use.
Since about 1970, a new concept of leak-before-burst (LBB) has been widely applied in
the designing and engineering calculations for reliable construction. This concept takes into
consideration the role of material flaws in the brittle fracture phenomenon and the influence
of corrosive environments on catastrophic failures. Thus, the material’s susceptibility to
245
Corrosion in Refinery and Petrochemical Industries
corrosion fracture as a time-dependent parameter in terms of a critical stress intensity factor
for stress and environment action (KISCC) values should be taken into account.
Figure 1. Probability of catastrophic failures versus plant operation time
The results presented in this paper consist of three main parts: 1) failure analysis of a
gas-plant’s serious accident due to corrosion crack initiation and propagation, 2) persistent
macro-micro segregation bands (MMB) role in the delayed corrosion fracture of pressurized
installations, and 3) the concept of weldable steel’s susceptibility to corrosion crack initiation
in parent metal and welding joints.
FAILURE ANALYSIS
A line-pipe 30″ (762 mm) in diameter transporting hydrocarbon condensate at about 430
psi (3 MPa) pressure exploded after about 17 years of gas-plant operation. The lengthwise
split shown in Fig. 2a produced a gigantic torch that heated up the other installations nearby
until a few blasts happened. The resulting damage destroyed nearly the whole main parts of
this plant (Fig. 2b).
Figure 2a. Lengthwise split of 30″ line-pipe
surrounding plant
246
Figure
2b.
Damage
to
Mazur
installations
During the gas-plant’s operation, the hydrocarbon condensate from time to time also
contained some (about 500 ppm) H2S which was dependent on the actual natural gas sources.
That detrimental compound reacted with the steel, producing internal cracks and
delaminations in the pipe-wall (Fig. 3). The corrosion cracks penetrated into the inner surface
of the pipe-wall, as shown in Fig. 4.
Figure 3. Internal delaminations in a 30″
line-pipe wall.
Figure 4. Corrosion crack traces on the inner
surface of the line-pipe wall
The active hydrogen atoms produced by the reaction of:
Fe + H2S → FeS + 2H
(1)
diffused through the iron crystal lattice decreasing the steel’s ductility.
A bend test specimen taken from the burst 30″ diameter pipe exhibited almost no
ductility (Fig. 5). The fracture surface showed fully brittle features (Fig. 6).
Figure 5. Lack of ductility of bend-test specimen
Figure 6. Fracture surface of test
247
Corrosion in Refinery and Petrochemical Industries
specimen after bend-test.
The very low energy values of the Charpy impact test confirmed the bend and tensile test
results. The fracture surface after the impact test also showed a brittle character and a number
of internal fissures (cracks along laminations) which are clearly visible in Fig. 7. Such
fissures, spreading along the steel’s rolling direction, were the narrow cracks produced by
hydrogen-induced high molecular pressure [2].
The evidences of the detrimental reaction of H2S with the steel’s surface were obtained
by using the sulfur prints method. Such a print after optical magnification (shown in Fig. 8)
illustrates the heavy FeS deposit (black spotted layer) on the inner pipe-wall surface. The
outer surface is free of such a layer (Fig. 9). The black spots are randomly distributed in the
steel-wall’s cross section. They are the traces of non-metallic inclusions of (Fe, Mn)S-type
left after the metallurgical processes typical for weldable steel production. The internal
fissure completely filled by FeS deposit is shown in Fig. 10.
Figure 7. Fracture surface of Charpy-V specimen
with internal fissures
Figure 8. Sulfur print tracing heavy FeS
deposit as black spots on the
the inner pipe-wall surface
Figure 9. Outer surface of the pipe free of
FeS deposit
Figure 10. Internal corrosion crack
(fissure) completely filled by
FeS deposit
248
Mazur
All mechanical and metallurgical tests clearly showed that hydrogen atoms diffused into
the steel-wall dramatically decreased the steel’s ductility and produced a great number of
cracks (fissures). During the long period of gas-plant operation, these cracks penetrated
through the pipe-wall cross section until the last steel fiber was broken. At that moment, the
pipe catastrophically burst producing the long opening crack visible in Fig. 2a.
A very detailed material testing of the other pipes operating under the same conditions
during the same period of time (as the burst pipe) also showed the hydrogen’s action, but it
was limited to a relatively thin layer on the inner surface of the pipe. FeS deposit was found
there, but the steel’s ductility was still high, according to the standards.
The question is, why one pipe was so greatly affected by hydrogen sulfide (and stresses)
but another one operating in the same environment was not susceptible to corrosion cracking.
The most probable answer was suggested already by Mazur [1], but the next experiments
gave more information about the influence of structural constituents (especially segregation
bands) on the susceptibility to hydrogen crack initiation and propagation in weldable pipeline steels.
SEGREGATION BANDS IN STEELS
During laboratory investigations, two types of segregation bands were found. The first
one was connected with persistent chemical segregations inherited by hot-worked steel after
crystallization. It is well known that the most detrimental elements in steel are sulfur and
phosphorus. Sulfur is present as non-metallic inclusions (sulfides) elongated along the rolling
direction. They act as the traps for diffusing hydrogen atoms through the crystal lattice.
Phosphorus is present in steel as dissolved atoms in solid solution and can be traced
metallographically by the use of special etchants. The Oberhoffer reagent [3] does not react
with steel if the phosphorus content is high (white band) but strongly etches the surface of
low-phosphorus steel. Such a wide macro-micro segregation band (MMB) in pipe-steel with
localized corrosion cracking is shown in Fig. 11.
Figure 11. Wide MMB with localized
corrosion cracking
Figure 12. Schematic concentration of
phosphorus in narrow and
wide
segregation
bands
(MMB)
249
Corrosion in Refinery and Petrochemical Industries
As the MMB becomes wider, the phosphorus concentration should be relatively higher.
Assuming sinusoidal change of the phosphorus concentration in steel, the schematic situation
in two steels of similar average phosphorus content, but different MMB width, is shown in
Fig. 12. The phosphorus concentration in the wide MMB is higher than in the narrow one.
Therefore, the susceptibility of wide MMB to corrosion crack initiation and propagation is
higher because of the embrittling influence of phosphorus. Usually in wide MMB, nonmetallic inclusions are present and act as internal stress-risers, accelerating corrosion crack
initiation and propagation.
The second type of structural band is often observed in ferrite-pearlite hot-worked
weldable steels. An example of such a band structure is shown in Fig. 13 after etching in a
typical solution of 4% HNO3 in alcohol (nital); the dark constituent is pearlite, and the white
is ferrite.
Figure 13. Structural bands in hot
worked pipe steel.
nonmetallic
Figure 14. Wide MMB with short
corrosion
cracks
and
inclusions (Oberhoffer etchant)
That type of band structure is common for carbon segregation during austenite to ferrite
and pearlite phase transformation. Usually the band structure is more developed as the steel
contains cosegregants like manganese, chromium, and silicon. These elements effectively
change the carbon atoms activity during their diffusion in the solid solution.
250
Mazur
Figure 15. Structural constituents present
in the specimen shown in Fig. 14
across
after nital etching
Figure 16. Linear microprobe analysis
of phosphorus content
a wideMMB
Both types of segregation bands are often present in weldable steels. But from a safety
point of view, the persistent MMB-type should be treated as the weakest area for corrosion
crack initiation and propagation. When the MMB, are wide, the susceptibility of the steel to
corrosion fracture is high. Such a wide MMB with a number of non-metallic inclusions is
shown in Fig. 14. Corrosion cracks are shown by arrows. A different area of the same
specimen, but nital-etched, is shown in Fig. 15. There is visible band structure, but in marked
areas there are also elongated nonmetallic inclusions and microcracks which most probably
belong to the wide MMB. These features are not clearly visible in the micrograph because of
the type of etchant (nital) used. A microprobe linear analysis made across the wide MMB
confirmed an increase in the phosphorus content at these sites (Fig. 16).
CORROSION CRACKS IN WELDED JOINTS
The influence of MMB directionality on corrosion crack propagation was tested on the
arc-welded specimens. For the laboratory tests, an A-52 grade weldable steel was selected.
The calculated carbon equivalent, CE , was 0.38; thus, the ordinary arc-welding procedure
was used. From the welded K-joints (Fig. 17), the crack-opening-displacement (COD)-type
specimens were taken. The MMB structure, with an average width of 57 microns, was
distributed along the rolling direction in the parent steel sheets.
The constant load cantilever bend test [4] for stress corrosion susceptibility measurement
of the welded joints was used. NACE water solution (5% NaCl, 0.5% CH3COOH) saturated
by H2S up to pH 3.5-3.8 acted on a short fatigue precrack made at the mechanically cut Vnotch’s front. The COD specimens were taken in such a manner (Fig. 18) that precracks were
likely to initiate corrosion cracks in different zones of the welded joint, i.e., fusion zone (FZ),
heat affected zone (HAZ) and parent metal (PM).
At the beginning, the lowest stresses acted on the precracked notch. As the test time was
increased, the real cross section decreased due to corrosion crack propagation. Thus, the
stresses that acted at the crack front steadily increased. Therefore, at the beginning of testing
and because of low stresses, corrosion crack propagation was controlled by a dissolution
mechanism. At some crack distance, when the crack length was long enough, its propagation
became controlled by a mechanical factor (stress).
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Corrosion in Refinery and Petrochemical Industries
Figure 17. Arc-welded steel sheets.
Figure 18. Mechanical notch placed in
parent metal (PM) heat affected zone
(HAZ) and fusion zone (FZ)
The above analysis explains why the corrosion crack at the beginning of testing (low
value of the initial stress-intensity factor KIi) propagated perpendicularly to the surface in the
form of a wide fissure. Over distance, KIi increased significantly to such a value that the
corrosion crack split into two branches propagating more or less parallel to the specimen’s
surface. This change in the direction of the corrosion crack coincided with the rolling and
MMB directions. At this stage, the stress-structure factor (mechanical) began controlling the
further process of corrosion crack development.
Such untypical branching effect for weldable carbon steels was first promoted by the
presence of MMB in the tested specimens.
Figure 19. Branching process starting at point A in the parent metal (unetched)
Figure 19 shows the branching effect in parent metal. From point A, the primary
corrosion crack split into two branches parallel to the rolling direction. The most interesting
observations were found for the specimens with a V-notch placed in a heat-affected zone
(HAZ). The micrograph shown in Fig. 20 was taken for a specimen not fully broken during
corrosion crack propagation
After using Anczyc etchant [5], revealing MMB as black bands, it is quite clear that up
to area A the corrosion crack propagated as a wide fissure perpendicular to the specimen’s
surface; similar to the primary crack shown in Fig. 19. From point A (Fig. 20), the primary
corrosion crack split into two cracks parallel to the rolling direction. One very short branch
was arrested by the fusion zone where no band structure exists, but the longer crack
propagated towards the parent metal along the MMB. As the time of the corrosion test was
increased, these long cracks, parallel to the rolling direction cracks and propagated in the
wide bands, joined the deeper placed MMB by short step-wise cracks visible in area B (Fig.
20).
252
Mazur
At some distance of the corrosion cracks in the cross section, the critical stress-intensity
factor (KIC) was reached and the specimen had to break by fast fracture mechanism. After
branching, the corrosion crack rate became constant due to stress reduction at the multiplied
crack fronts in spite of the increasing values of the stress-intensity factor (KIi). The
da
and KIi is shown in Fig. 21.
relationship between the corrosion crack rate
dt
Figure 20. Primary corrosion crack perpendicular to the specimen’s surface; the crack
initiated in the HAZ and split into two longitudinal branches with the longer
branch propagating towards the PM and the shorter one towards the FZ (Anczyc
etchant)
Figure 21. Corrosion-crack rate versus stress-intensity factor (KIi).
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Corrosion in Refinery and Petrochemical Industries
In all tested specimens, a dominant role of the MMB structure in branching was found.
The corrosion cracks always propagated towards the parent metal where MMB existed, but
the branches propagated towards the weld-bead were arrested due to the lack of MMB
structure there. Therefore, the corrosion cracking of the weldable steels depends on more or
less developed MMB structure. With greater MMB width, is greater the steel is more
susceptible to corrosion crack initiation and propagation.
The results of KISCC measurements for specimens with precracks (notches) placed in the
three different zones of welded joints have shown that parent metal (steel sheet) exhibited the
highest susceptibility to corrosion cracking in terms of a KISCC to KIC ratio that was 0.72 only.
For HAZ and FZ this ratio was close to 1.0, i.e., both zones are nearly immune to applied
environment and stresses. Most of the published results did not analyze the role of structural
bands in branching effects. Published micrographs of branching cracks [6] most probably
have propagated along the rolling directions of austenitic steels and high strength aluminium
alloys.
ACKNOWLEDGMENT
This research was partially sponsored by The University Status Research Fund of The
Academy of Mining and Metallurgy, Krakow, Poland.
REFERENCES
1.
2.
3.
4.
5.
6.
254
A. Mazur, Materials Performance 34, 1995, pp. 52-54.
A. Ikada, Y. Morita, F. Terasaki, and M. Takeyama, On the hydrogen induced
cracking of line pipe steel under wet hydrogen sulphide enviroment, Proc. 2nd
International Congress, Pergamon Press, Paris, 6-10 June 1977.
L. Habraken, J.L. de Brouver, De Ferri Metallographia, CNRM, Vol. 1, Bruxelles,
Press Académiques Européennes, 1977, p. 41.
B.F. Brown, The application of fracture mechanics to stress-corrosion cracking,
Metallurgical Reviews 13, 1968, p. 21.
K. Przybylowicz, Metaloznawstwo, Part I, Struktura metali i stopow, metody badania,
Krakow, AGH Publ., 1989, p. 179.
M.O. Speidel, Branching of Stress Corrosion Cracks in Aluminium Alloys, in: The
Theory of Stress Corrosion Cracking in Alloys, Ed., J.C. Scully, Brussels, NATO
Publ., 1971, pp. 345-354.
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
FLUID CATALYTIC CRACKING INTERSTAGE AND
HIGH-PRESSURE COOLER CORROSION
S.M. Halawani
Saudi Aramco - Jeddah Refinery Department
Engineering Division - Operations Engineering Unit
Admin. Bldg. Room # 218, P.O. Box 5250, Jeddah 21441, Saudi Arabia
ABSTRACT
The fluid catalytic cracking unit (FCCU) is one of the most profitable process units in the
petroleum refining industry. The FCCU converts less valuable heavy hydrocarbons into lighter, more
valuable components by cracking the larger molecules in the presence of a catalyst. During the
conversion process, many highly corrosive chemical species are evolved from the catalytic reactions.
One of the principal areas of the FCCU that comes into contact with such corrosive chemicals is
the unsaturated gas concentration section. This section receives corrosive hydrocarbon gases from the
FCCU Fractionation section, which come into contact with various pieces of equipment. This paper
describes the economic penalty of frequent shutdowns or reduced throughput that have been caused by
this type of corrosion at the Jeddah Refinery.
The specific example addressed in this paper is corrosion of the wet gas compressor interstage
and high-pressure coolers. This costly problem was resolved by a simple and easy solution over
several other more costly alternatives. The solution that was selected involved proper distribution and
adjustment of water wash to the shell side of the interstage and high-pressure coolers. This water also
serve to prevent the disposition of ammonium chloride salts on the external surfaces of the cooler
tubes.
Key Words: Fluid catalytic cracking unit, molecules, catalyst, conversion process,
unsaturated gas concentration section, shutdowns, ammonium chloride
INTRODUCTION
Equipment deterioration is one of the major expenses in the petroleum refining industry.
This problem can have several different causes, but is generally mechanical (erosion) or
chemical (corrosion) in nature. This paper describes one of the corrosion problems that
occurred in the fluid catalytic cracking unit (FCCU) at the Jeddah Refinery.
This corrosion problem caused throughput reductions, safety hazards, equipment
deterioration and high maintenance costs. The resolution of the problem took several years to
identify the causes, and to develop and implement corrective action. In the end, the solution
was very simple and easy to implement.
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Corrosion in Refinery and Petrochemical Industries
DISCUSSION
The Jeddah Refinery is located in the city of Jeddah, Kingdom of Saudi Arabia. The
majority owner and operator of the refinery is Saudi Aramco, which is the world’s largest
company in the petroleum industry. The Jeddah Refinery processes light and heavy Arabian
crude oil, and has a capacity of 90,000 bbls/day. The major process units consist of
Atmospheric and Vacuum Distillation, Hydrodesulfurization, Catalytic Reforming, and
Catalytic Cracking units. The unit under discussion in this paper is the Fluid Catalytic
Cracking Unit or FCCU.
Feed
The FCCU processes a variety of different feedstocks, basically from Vacuum
Distillation units at the Refinery, and from the Lube Oil Refinery located adjacent to the
refinery. The feedstocks include heavy vacuum gas oils, slack waxes, and 100 SUS distillate.
Types of Corrosion in FCCUs
Several types of corrosion can take place in FCCUs. The types can be grouped into two
major categories [1]:
Hydrogen charging; and
Carbonate cracking
Hydrogen charging includes sulfide cracking, blistering, hydrogen induced cracking, and
stress-oriented hydrogen-induced cracking.
Carbonate cracking is also known as
intergranular stress corrosion cracking.
Table 1. Materials of Construction of Coolers
Cooler Part
Shell & Heads
Tube Sheet Stationary &
Floating
Tubes ( Straight Type)
Material
A-515 Gr.60 (CS for high temp service)
A-181 Gr. I (Rolled Steel)
ASTM B-111, No. 687 (Cu-Ni Alloy)
Interstage and High Pressure Cooler System
Serious corrosion problems were experienced in the exchangers that cool the FCCU wet
gas compressor’s first and second stage discharge streams. This system can be described as
follows (Fig. 1):
• The first stage compressor discharge is cooled by two parallel banks of exchangers,
each consisting of two exchangers operating in series, labeled G-E1 A/B/C/D.
• The second stage compressor discharge is also cooled by two parallel banks of
exchangers, each consisting of two exchangers operating in series, labeled G-E2
A/B/C/D.
• Seawater passes through the tube side at lower pressure than the hydrocarbon gases
that pass through the shell side.
• All coolers are TEMA type AES, which consist of four tube passes per cooler.
256
Halawani
Figure 1. Schematic of water wash line of fractionation system FCCU
257
Corrosion in Refinery and Petrochemical Industries
• GE-1 A/B/C/D tube number = 164 (each)
• GE-2 A/B/C/D tube number = 452 (each)
• The coolers are constructed from the materials shown in Table 1.
Water Wash Role
One of the factors that plays a major role in the system is the wash water to the
hydrocarbon inlets of G-E1 A, G-E1 C, G-E2 A, and G-E2 C (Fig. 1).
The vapor stream from the main column overhead receiver at the FCCU (compressor
suction) contains various contaminants which may cause corrosion, plugging, or fouling.
These contaminants include ammonia, sulfides, cyanides, chlorides and phenols. Since most
of these contaminants are ionic or polar species and are readily soluble in water, a wash-water
stream is used to concentrate them in the aqueous phase.
The contaminated water stream is subsequently removed from the system. A wash-water
rate of about 7 vol. % of the fresh feed is recommended for washing these exchangers.
The water should be clean, preferably steam condensate, to prevent adding more
problems, such as salts or dissolved oxygen, to the system. The water is injected after the
compressor first stage (i.e., to GE-1 A/B/C/D ) and is pumped out of the interstage suction
drum to the second stage outlet (i.e., to GE-2 A/B/C/D ). The water is finally collected at the
high-pressure receiver boot and transferred to the main column overhead coolers for further
washing.
There is always water present in the main column and the gas concentration section from
stripping steam. If the wash water is not used to flush out the corrodants, the water present
can become highly corrosive from absorption of these corrodants. Sulfides levels in excess of
20,000 ppm have been reported in the overhead receiver water.
Hydrogen blistering and general corrosion attack may become quite severe, especially if
feed sulfur is greater than 1% wt., or nitrogen is greater than 1000 ppm. While the main
column overhead receiver water may be basic (i.e., pH greater than 7.0), most of the ammonia
that is responsible for this high pH drops out in the main column receiver.
The water in the gas concentration section may become acidic from hydrogen sulfide and
cyanides. If there is any oxygen present, elemental sulfur may be formed from oxidation of
the sulfides. This elemental sulfur will cause problems in meeting gasoline product
specifications. Wash water will solve many of these problems by diluting the corrosives, and
keeping the water’s pH in the range of 8.0-9.0, where sulfide oxidation is greatly reduced.
TROUBLESHOOTING
The corrosion problem that took place in the subject coolers caused the following
drawbacks:
- Hydrocarbons leakage into the seawater, which caused safety and environmental
hazards;
- Decreased FCCU throughput to repair these leaks;
- Erosion and corrosion of tubes ends;
- Broken bolts on the floating heads at the coolers;
258
Halawani
- Shortened service life for tubes (i.e., 4 months period);
- Bulging of cladding.
The specific corrosion for each group of coolers was as follows:
Corrosion of GE-1 A/B/C/D
• Bulging of cladding at the floating head;
• Leakage of the brass tube caused by corrosion.
Corrosion of GE-2 A/B/C/D
•
•
•
•
Bulging of cladding at the floating head;
Erosion in the tube ends in the floating tube sheet with marine life deposit;
Erosion in top pass of floating tube sheet.
Fracture of bolts in floating head.
Several potential causes of the corrosion problems were discussed with a reputable
consultant. The areas considered included:
•
•
•
•
•
•
•
Foreign materials in seawater causing corrosion;
Contamination of seawater by hydrogen sulfide (H2S);
H2S leakage from shell side;
H2S stress corrosion cracking of bolts;
H2S blistering of cladding;
Wrong design of floating head;
Bacteria causing iron sulfide (FeS) deposits in the seawater supply.
RECOMMENDATIONS
The following recommendations were considered for this corrosion problem:
• Seawater velocity in the tube side to be maintained between 0.9-1.8
•
•
•
•
•
•
•
•
meters/Second;
Installation of sacrificial iron anodes on the channel and floating heads;
Injection of a filming inhibitor, i.e., iron sulfate (FeS04), into the wash water;
Removal of foreign materials from the seawater by filtering;
Installation of plastic inserts at tube inlets to prevent erosion;
Eddy current examination during the inspection of the unit,
Review of the design of the floating head;
Checking of the hardness of the bolts which should be less than 235 BHN
(Brinnel Hardness Number);
Changing of tube material to a corrosion-resistant type. (i.e., titanium).
Many recommendations were made, but one very important process related point was not
taken into consideration. This point concerned the water wash to the compressor coolers to
wash off ammonium chloride salts and other corrosion species.
259
Corrosion in Refinery and Petrochemical Industries
Figure 2. Schematic of proposed water wash line of fractionation system FCCU
after introducing modifications
260
Halawani
Originally, the total volume of wash water was routed to the first stage intercooler. In
checking the existing system, it was determined that one-half of the total wash water quantity
was sent to another user (M/C overhead cooling system). This caused the amount of wash
water sent to the interstage coolers to be half of the recommended quantity (20 gpm). This
deviation was corrected by injecting the recommended wash water quantity to the first stage
intercooler and by installing restriction orifices on each of the parallel cooler trains (Fig. 2).
CONCLUSIONS
After careful consideration, the following steps were taken to correct the corrosion
problems in the FCCU wet gas compressor interstage and high-pressure coolers:
1. The wash water rate was adjusted to 20 gpm as provided for in the original design.
2. Restriction orifices were installed on each branch to ensure equal distribution of wash
water to each branch.
After completing these actions, only very rare leaks have been observed in the
compressor interstage and high-pressure coolers. In summary, when troubleshooting
problems, all related items must be taken into consideration, even if they appear to be very
minor at the time. Obviously, it was much less expensive to adjust the wash water rates and
distribution than to install titanium bundles, which was one of the recommendations
considered.
REFERENCES
1. R.C. Strong, V.K. Majestic and S.M. Wilhelm, Basic steps lead to successful FCC
corrosion control, NALCO Reprint SL-47 (reprinted from Oil and Gas Journal), Sept.
30 and Oct. 7, 1991.
261
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
ASSESSMENT OF CRACKS IN A HIGH PRESSURE
MULTILAYERED REACTOR FOR ITS
FITNESS FOR PURPOSE
A.M. Askari, M.I. Al-Kandari and P.K. Mukhopadhyay
Industrial Safety Division, Shuaiba Refinery, Kuwait National Petroleum Company
ABSTRACT
The high-pressure Kero hydrodesulfurizer reactor in the Shuaiba Refinery of the Kuwait National
Petroleum Company was commissioned in 1968. By 1973, routine inspection revealed small cracks
up to 3/16 in. (4.8 mm) deep which gradually increased in size necessitating weld repair in 1981.
However, subsequent inspection in 1985 revealed reappearce of the cracks, and by 1990, the cracks
were about 1 in. (25.4 mm) deep and 13 in. (330.2 mm) long. This reactor was built in the 1960s
without adequate control of the chemistry (J-factor) and impact properties to reduce temper
embrittlement. In view of the this and possible damage due to hydrogen embrittlement after 20 years
of service, further in-situ repair by welding was ruled out. Instead, a fitness-for-purpose study was
made utilizing recent developments in fracture mechanics techniques and available software to ensure
safe and reliable operation by pressure down rating the reactor.
Key Words: Kero hydrodesulfurizer reactor, J-factor, temper embrittlement, hydrogen
embrittlement, fitness-for-purpose, fracture mechanics
INTRODUCTION
Chromium-molybdenum steels used for the construction of hydro-processing reactors are
known to undergo material degradation in service. The most important material degradations
are temper embrittlement and hydrogen embrittlement. The first generation of reactors, built
in the 1960s when the current controls on chemistry and impact properties were not exercised,
are prone to temper embrittlement. Therefore, when cracks are found in a reactor during
routine inspection, an assessment to predict the continued safe and reliable operation of the
reactor requires serious consideration. In the past, in the absence of adequate tools to
quantify the impact of such defects, plant engineers faced difficult decisions in such
circumstances. The decision, in some cases, may be very conservative requiring long-term
plant shutdown for costly repair/replacement. The availability of computer softwares and
standards providing guidance on methods for assessing the acceptability of flaws, e.g., British
Standard BSI PD6493 and other such standards, have provided engineers with much needed
tools upon which a more careful recommendation can be based for continued safe operation
of such equipment. This paper describes the assessment and current status of the cracks
found in a Kero hydrodesulfurizer reactor at the Shuaiba Refinery, Kuwait National
Petroleum Company, and highlights the importance of such assessment techniques.
CONSTRUCTION AND HISTORY OF THE REACTOR
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Corrosion in Refinery and Petrochemical Industries
The Kero hydro-unifiner reactor is a multilayered reactor consisting of 14 layers each ½
in. (12.7 mm) thick made of a high strength material (JSW 2HS) and a core plate of SA 387
Gr D i.e., 2 1/4 Cr 1 Mo steel. At either end of the multilayered shell is a forged ring to
which hemispherical heads (SA 387 Gr D) are welded. The reactor has 347 stainless steel
type cladding on the inside surfaces The construction details are shown in Fig. 1. The plant
was commissioned in 1968, and the design pressure and temperature of the reactor are 2800
o
psig. (19.3 Mpa) at 750°F (399 C).
The cracks were first detected in this reactor by ultrasonic flaw detection (UFD) during
routine inspection in 1972. The cracks were up to 3/16 in. (4.8 mm) deep on the inner side of
the welds between the bottom forged ring and the heads. The configuration of the welds and
approximate locations of the cracks can be seen in Fig. 2.
These cracks were monitored on a run-to-run basis till 1978 when UFD indicated flaws
up to 1.5 in. (38.1 mm) deep at some locations. These cracks were successfully repaired and
after repair no significant defect was detected by UFD. Subsequent inspection in 1986
revealed the reappearance of the cracks. During a 1990 inspection, cracks were detected up
to 1 in. (25.4) deep at places and 13 in. (330.2 mm) long in the circumference in the weld
between the forged ring and the bottom head. Alarmed by the dimensions of the cracks and
realizing the little chance for further weld repair, a fitness-for-purpose study was conducted to
determine the reactor’s continued safe and reliable operation with the existing cracks.
ASSESSMENT FOR BRITTLE FRACTURE
To assess failure due to the existing cracks by brittle fracture due to embrittlement of the
material, the stress-intensity approach was used. This approach is based on the assumption
that crack propagation will occur when the stress intensity (K) at the crack tip reaches a
critical value Kc. Under plane-strain condition, this critical stress intensity for tensile loading
is termed as KIC. A general expression for K [1] can be written as
K=Mσ
πa
(1)
where M = a constant specific to a given flaw size and geometry,
σ = stress in the system, remote from the crack, and
a = dimension of the crack.
As the cracks at the weld joint between the bottom forged ring and head were the worst,
the assessment of the cracks was done for these cracks, although cracks were present in other
joints.
The structural integrity assessment was carried out by the use of the licensed software,
Crackwise, from the Welding Institute (TWI), UK [2]. Several of the inputs required were
arrived at as follows:
• Flaw dimensions, as found by UFD, were grouped as per the guidelines provided
in BSI Standard PD 6493 [3].
264
Askari et al.
Figure 1. Sketch of the kero hydrodesulfurizer reactor
265
Corrosion in Refinery and Petrochemical Industries
Figure 2. Weld configuration and approximate cracks location
266
Askari et al.
Figure 3. FAD for the existing crack under 2800 psig for A387-D (level 1)
• Geofac software licensed by the Fracture Search Inc. [4] was utilized for
calculation of the M factor for the existing crack.
• Fracture toughness data was obtained from the co-relation given in PD 6493 [3]
between fracture toughness and charpy V-notch values, where the charpy V-notch
values were measured by removing boat samples from a similar reactor.
The analysis was carried out for both high-pressure operation at 2800 psi (19.3 Mpa) and
medium-pressure operation at 1300 psi (8.96 Mpa) at level 1, which is a somewhat
conservative assessment. The analysis was done for both materials, i.e., 2HS and SA 387 Gr
D as the cracks were present at the weld between these two materials. The fracture analysis
diagrams (FADs) for the 387Gr D material are presented in Figs. 3 and 4. As can be seen
from the FAD, the cracks are considered to be safe from failure by brittle fracture at 1300
psig (8.96 Mpa) operation, but not safe when the pressure is raised to 2800 psig (19.3 Mpa).
267
Corrosion in Refinery and Petrochemical Industries
Figure 4. FAD for the existing crack under 1300 psig for A387-D (level 1)
ASSESSMENT FOR HYDROGEN EMBRITTLEMENT
The second assessment was done for the combined effect of hydrogen embrittlement and
temper embrittlement, as several studies have shown that the threshold stress intensity for
cracking in hydrogen (i.e., KIH) can be appreciably reduced in steel that has been subjected
to prior temper embrittlement. Wadate et al. [5] have provided a correlation between KIH
and embrittlement properties measured by fracture appearance transition temperature (FATT)
as follows:
KIH = .0014 FATT 2 - 0.421 FATT + 57.0
where KIH is expressed in Mpa
(2)
m and FATT in degrees celsius.
In the absence of actual data on FATT, the correlation between the J-factor i.e., (Si +
Mn) x (P + Sn) x 10-4 and FATT has also been provided [6].
Table 1. Data Analysis for the Existing Crack Under 2800 psig for A387-D (level 1)
268
Askari et al.
Table 2. Data Analysis for the Existing Crack Under 1300 psig for A387-D (level 1)
269
Corrosion in Refinery and Petrochemical Industries
In the present study, a detailed chemical analysis was done for boat samples removed
from a similar reactor which has a similar manufacturing and operating history. This resulted
in a J-factor of 350, and a corresponding KIH of 30 MPa m0.5 after 100,000 hours of
operation. The stress intensity factor (KIc) for 1300 psi (8.96 Mpa) operation was calculated
to be 35.6 Mpa m0.5 for the worst crack present.
Since KIc > KIH the cracks are of the propagating type. However, the rate of crack
-24
11.7
growth, i.e., da/dt = 2.4 x 10 x K
to be considered very insignificant.
-6
[7], works out to be 3.39 x 10 mm/year which was
DISCUSSION
From the above analysis, it was concluded that the reactor was not safe for operation at
its originally designed pressure of 2800 psi (19.3 Mpa) (Fig. 3 and Table 1). However, it was
considered safe to operate at 1300 psi (8.96 Mpa) (Fig. 4 and Table 2) at which the plant was
found to be still capable of producing products to standard specifications. At 1300 psi (8.96
Mpa) operating condition, from a hydrogen embrittlement point of view, the chances of crack
growth is a possibility, but as the rate of crack growth is found to be quite low, it was
considered safe to operate the reactor by down grating the pressure to 1300 psig (8.96 Mpa).
No significant crack growth could be detected after operating at the lower pressure of 1300
psig (8.96 Mpa) between last two UFD inspections made in 1992 and 1995. It is, however,
fully recognized that in the event of further crack growth, in-situ welding repair will not be
possible due to the metallurgical degradation of the reactor’s materials. Also, due to the more
stringent specifications required for kerosene in respect to smoke point (i.e., 28 mm
minimum), the reactor needs to be operated at high-pressure at around 2300 psi. Our analysis
shows that with the existing crack, operation of the reactor will not be safe at 2300 psi.
Accordingly, it has been decided to procure a new reactor for operation at 2300 psi. Till such
time, the reactor will continue to operate at 1300 psi with run-to-run monitoring by UFD.
CONCLUSIONS
The above discussion highlights the advantages offered by fracture mechanics techniques
for finding acceptable solutions at the plant level. Such solutions can be achieved by
analyzing the fitness-for-purpose of the equipment to continue operation without
compromising the safety and integrity of such important equipment.
ACKNOWLEDGMENT
The authors wish to thank the management of KNPC for permission to publish this
paper. Encouragement and support provided by Refinery Manager Mr. A.L. Al-Houti is
sincerely acknowledged.
REFERENCES
1. Viswanathan, Damage Mechanisms and Life Assessment of High Temperature
Components, ASME International, Ohio, 1989, pp. 21-57.
2. Crackwise Software, The Welding Institute (TWI), UK.
270
Askari et al.
3. PD 6493, Guidance on Methods for Assessing the Acceptability of Flaws in Fusion
Welded Structures, BSI, 1991.
4. Geofac Software, The Fracture Search Inc., Ohio.
5. Wadate, J. Wantanable and Y. Tanka, Prediction of the remaining life of high temperature
pressure reactors made of Cr-Mo steels, Trans. ASME, J. Pressure Vessel Tech. 107, Aug.
1985, pp. 230-238.
6. Wadate, T. Nomura and J. Watanabe, Hydrogen effect on remaining life of hydroprocessor reactors, Corrosion 44, 2, 1998, p. 106.
7. Viswanathan, Damage Mechanisms and Life Assessment of High Temperature
Components, ASME International, Ohio, Ohio, 1989, pp. 329-382.
271
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
POLYTHIONIC ACID STRESS CORROSION CRACKING
OF INCOLOY 800: CASE STUDY AND FAILURE ANALYSIS
M.S. Mostafa and S.A. Hajaj
Industrial Safety Division, Mina Abdulla Refinery
Kuwait National Petroleum Company, Kuwait
ABSTRACT
A high-pressure warm-separator (HPWS) off gas air cooler made of Incoloy 800 failed at one of
its inlet nozzles. Inspection and nondestructive testing (NDT) showed cracks in the shape of sun rays
surrounding the weld of the chemical cleaning nozzle connected to the inlet nozzle. Chemical
analysis, metallographic examinations, and microhardness measurements indicated that the cracks
were due to polythionic acid stress corrosion cracking. Cracking was promoted by the presence of
carbide precipitation and residual stresses resulting from the manufacturing and welding practices.
Key Words: Polythionic acid, stress corrosion cracking, Incoloy 800, case study, failure analysis
INTRODUCTION
An atmospheric residue desulphurization unit (ARDS) was commissioned in the early
1988 as a part of the modernization project of Mina Abdulla Refinery. The unit is designed
to process 65,900 BPSD of atmospheric residue. At the heart of the unit is a pair of parallel
fixed-bed reactor trains containing desulphurization catalyst. The chemical basis for the
process is the catalytic reaction of hydrogen with sulfur-containing hydrocarbons to produce
sulfur-free hydrocarbon and hydrogen sulfide. The incoming residue feed is preheated by
exchange with the reactor effluent and a fired heater. The feed is mixed with hot-recycled
gas, and is sent to the reactor where it is desulphurized. Finally, the reactor effluent goes to a
series of separators to release the dissolved light ends (Figs. 1 and 2).
The high-pressure warm-separator (HPWS) off gas air cooler consists of a series of eight
banks air cooler, operating at 180°C under a gas pressure of 21 bar, and handling a mixture of
H2, NH3, H2S and hydrocarbon. Each bank has two inlet and two outlet nozzles, 4 in.
diameter. There are two chemical cleaning nozzles, 1.5 in. diameter, per bank, one of each in
the inlet and outlet nozzles. The air cooler headers, nozzles and tubes are all made of Incoloy
800 in accordance with ASTM B 407-UNS-N08800 (see Tables 1 and 2). Wash water is
injected into the air cooler effluent to prevent the deposition of solid ammonium salts (i.e.,
NH4Cl and NH4OH). The ammonium salts are readily removed by the wash water (see Table
3).
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Corrosion in Refinery and Petrochemical Industries
Figure 1. A schematic flow diagram of the atmospheric residue desulphurization unit
(ARDS)
274
Mostafa and Hajaj
Figure 2. A schematic diagram of high pressure warm separator off gas air cooler (HPWS)
275
Corrosion in Refinery and Petrochemical Industries
Table 1. Specified Mechanical Properties of Incoloy 800
(ASTM B407, UNS No. 8800)
Tensile Strength (ksi)
65
Yield Strength (ksi)
25
Elongation (%)
30
Table 2. Specified Chemical Composition of Incoloy 800
(ASTM B407, UNS No. 8800)
Element
Nickel
Chromium
Iron (min.)
Manganese
Copper
Silicon
Sulfur
Aluminium
Titanium
Amount (%)
30-35
19-23
39.5
1.5
0.75
1.0
0.015
0.15-0.60
0.15-0.60
Table 3. Specifications of Wash Water
Component
Total dissolved solids (max. ppm)
Dissolved oxygen (max. Ppm)
Sodium (max. Ppm)
Potassium (max. Ppm)
Total hardness (max. Ppm)
Organic material (max. Ppm)
pH
Conductivity ((μmattscm-1)
Amount
3.0
0.05
0.10
0.10
0.10
1.0
7.0-8.0
5.0
HISTORICAL DATA AND BACKGROUND OF THE UNIT
Two successive incidents of hazardous on-stream leakage developed from the HPWS off
gas air cooler, one during the commissioning period and the other 11 days after the
introduction of fuel oil feed. The tubes and headers are made of duplex stainless steel
conforming to SA-669 and UNS S31500, respectively.
Inspection revealed extensive cracking of many of the header’s partition plate/tube sheet
welds and also of the header’s corner welds. In all cases, cracking included tube end welds
that failed and caused leaks to occur. Metallurgical examination revealed almost 100% ferrite
coarse-grained heat affected zone (HAZ) near the welds due to improper design, fabrication
276
Mostafa and Hajaj
and welding technique. Hydrogen embrittlement of the HAZ developed under normal
operating conditions leading to weld cracking and subsequent leakage.
Repair of the cracked headers of the eight banks is impractical, costly and not guaranteed
to lead to solving the problem. Furthermore, the use of the remaining uncracked banks is not
advisable given the potential for similar problems. They are also not recommended for
reconditioning for the same reasons. It was recommended to use a completely new air cooler
with improved design. The new air cooler was fabricated from Incoloy 800, in a similar
manner to those used in other neighboring refineries.
CRACKING IN INCOLOY 800 OFF GAS AIR COOLER NOZZLE
The new air cooler made of Incoloy 800 conformed to ASTM B 407-UNS 08800. It was
delivered in 1992 and was installed to replace the then existing carbon steel air coolers. In
early January 1995, one four inch diameter inlet nozzle of one of the air cooler banks leaked
during start-up of the unit after a maintenance shutdown. The leak was located around the
weld of the 1.5 in. diameter chemical cleaning nozzle connected to the 4 in. diameter inlet
nozzle. All 32 inlet and outlet nozzles, 16 of which had chemical cleaning nozzles, were
checked with fluorescent dye, and 8 more 4 in nozzles were found to have cracks. Nine
nozzles (inlet and outlet) were totally replaced with 321 stainless steel. The new nozzles
were not provided with chemical cleaning nozzles.
Metallography and microhardness tests were carried out to determine the reason for
failure. Some sections were given to an outside consultant and to the manufacturer.
Initial Observations and Nondestructive Testing (NDT)
The leaky 4 in. diameter nozzle was removed from the header and the internal surface
was checked with fluorescent dye. Cracks were observed on the internal surface of the 1.5 in.
diameter nozzle weld. The cracks were seen to initiate from the weld toe and were
perpendicular to the weld (Fig. 3).
Cracks were also observed in the 4 in. diameter nozzle at the wall deposit to the weld.
Small indications of cracks were also noted at the nozzle-to-flange weld.
The leaky 4 in. diameter nozzle was first cut longitudinally into two halves. The half
containing the 1.5 in. nozzle was further cut and one quadrant was given to the representative
of the designer, while another identical quadrant was used for the in-house investigation. The
other half was preserved for further analysis by an outside party.
Metallography
The in-house sample containing cracks was first cut longitudinally. It was then cut into
two transverse cross sections. The first sample revealed intergranular corrosion of the base
metal (i.e., in the 4 in. diameter nozzle). Severe carbide precipitation was noticed at the grain
boundaries (Fig. 4). The examined section was away from the weld. Severe disintegration
and also loss of grains was also noticed (see Fig. 5), rendering the structure very weak. The
ASTM grain size was found to be 4.5-5. Also, the microhardness of the base metal was 169
BHN, which is considered normal.
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Corrosion in Refinery and Petrochemical Industries
Figure 3. A photograph showing internal view of nozzle section following the
application of dye penetrant. Cracks are seen to initiate from the weld
toe and are perpendicular to the weld
Figure 4. A stereo micrograph showing intergranular corrosion and grain
boundary precipitation (400x). Etchant: HNO3-HCl-Acetic
278
Mostafa and Hajaj
Figure 5. Severe intergranular corrosion with disintegration and loss of grains in
the base metal (50x). Etchant: HNO3 - HCl - Acetic.
Figure 6. Intergranular cracking in the longitudinal section, showing crack initiation
from weld toe and loss of grains
279
Corrosion in Refinery and Petrochemical Industries
The second sample revealed intergranular cracking in the longitudinal cross section. The
cracks were in the base metal (i.e., in the 4 in. diameter nozzle). The cracks were noticed to
initiate from the weld toe and to travel to the other edge through the base metal (Fig. 6).
Carbide precipitation, disintegration and loss of grains were observed once again . The
ASTM grain size of the second sample was 4.5-5 and the mircrohardness of the base metal
was 165 BHN, which is considered normal. The second sample also revealed intergranular
corrosion and cracking on the transverse cross section. The cracks appeared to propagate
along the grain boundaries. Carbide precipitation, disintegration and loss of grains were also
noticed within the examined cross section.
Fractography
The fractured surfaces were examined using scanning electron microscopy (SEM).
Crystalline facets were noticed (Fig. 7), especially in the region of crack initiation.
Figure 7. SEM fractograph showing separated grain facets and secondary cracking
Chemical Analysis
The chemical composition analysis confirmed that the material was identical in
composition to Incoloy 800. Energy dispersive spectroscopic analysis revealed the presence
of sulfur.
Reasons for Failure: Analysis and Ven Diagram for Stress Corrosion Cracking
Intergranular corrosion (IGC) of the base metal (i.e., 4 in. diameter nozzle) indicated that
carbides had precipitated on the grain boundaries during the manufacturing process, because
280
Mostafa and Hajaj
the operating temperature of 340°F is too low for such precipitation to occur during service.
Precipitation of carbides also occurred in the HAZ during welding, as was observed during
the examination.
All 1.5 in. diameter chemical injection nozzles were replaced at the time of installation
due to the leakage of the original nozzles during hydrotesting. This resulted in the HAZ
being affected twice. However, IGC was noticed on the base metal far away from the weld
which confirms the existence of microstructural deficiencies during manufacturing.
Precipitation of carbides makes the grain boundaries prone to corrosion and cracking.
The presence of a corrosive medium such as polythionic acid would promote corrosion of the
grain boundaries under such conditions. The presence of IGC suggested the formation of
such acid during shutdown, which is very likely in the absence of neutralization before
shutdown. Although Incoloy 800 is considered superior to austenitic stainless steels, it is still
susceptible to polythionic acid corrosion attack and the system should be neutralized before
opening, as recommended in NACE RPO1-70-93. Understandably, the presence tensile
stresses is required to initiate and propagate cracks along the corroded grain boundaries. In
the present case, stresses can be present in the form of residual stresses as a result of the
manufacturing practices and welding. Furthermore, the 1.5 in. diameter nozzle welds
appeared under bending stress.
CONCLUSIONS
The material (Incoloy 800) was received in a sensitized condition which is considered to
be inferior and unacceptable. Welding of the chemical cleaning nozzles considerably affected
the HAZ. Formation of polythionic acid during shutdown promoted the IGC process.
Indeed, energy dispersive analysis confirmed that sulfide was present. The pattern of
cracking was consistent with the radial stress around the weld. Based on the investigation, a
recommendation was made to recheck all the 4 in. inlet and outlet nozzles for cracking with
fluorescent dye penetrant as well as the 321 stainless steel. Considering the hazards involved,
it was preferable to replace all of the original Incoloy nozzles, even if they were not cracked.
Because there was a strong possibility that all nozzles from the same manufacturing batch
would contain harmful precipitations, in-situ metallography (replica) was highly
recommended for all the nozzles and headers to determine the presence of carbide
precipitation. Thus, a firm specialized in the field of in-situ metallography was hired to test
all of the inlet and outlet nozzles, headers, drains, vents and tubes.
281
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
CORROSION OF TUBE HEATERS IN REFINERIES:
SYMPTOMS AND CURES
A. Attou, A. Rais and H. Smamen
Entreprise Nationale de Raffinage de Petrole
Raffinerie d’Arzew BP 37 Arzew, Algerie.
ABSTRACT
Tube heaters (THs) in oil refineries play a strategic role in heating hydrocarbons. However, they
may suffer from corrosion damage if service conditions differ from those they are designed to operate
under and/or if not enough attention is given to the composition/heat treatment of the tube materials
and the composition of the combustible products.
We experienced this problem with a cylindrical vertical heater made from convection tubes made
of material ASTM A106, and radiation tubes made of ASTM 335 Gr P11. After seven years in
service, the heater’s tube surfaces showed severe crevice corrosion damage. A number of analyses
and tests were carried out to study the material’s characteristics and to compare them with those of a
sound heater. Tests included metallurgical analysis, microhardness, grain compactness, carbon
concentration and chemical analyses of the combustible products.
This paper describes the results of each testing technique and identifies the origin of the
corrosion damage as grain compactness/heat treatment and the presence of corrosive agents in the
combustible products and in the air mixed with vapours coming from a nearby cooling tower.
Key Words: Tube heaters, hydrocarbons, corrosion, grain compactness, combustibles.
INTRODUCTION
Tube heaters (THs) are widely used in the process of refining crude oils. Different
products are extracted at different temperatures: Naphtha up to 200°C, Kerosene between
200°C and 300°C, and Diesel between 20°C and 380°C. Other extracted products are usually
referred to as fuel. To reach these temperatures, a number of heating methods are used.
•
•
•
•
Steam heating,
Electrical heating,
Heat exchangers, and
By direct firing (THs).
THs are chosen when other methods fail to give the desired elevated temperature or for
economical reasons. They were first invented in the USA at the beginning of the 20th century
and are classified either by their heat transfer method, i.e., radiation or radiation/convection,
or by their external appearance, i.e., box heaters, cell heaters, cylindrical heaters, etc. The
tube materials used in these heaters are usually: A 335 Gr P11 (1¼ Cr ½ Mo), ASTM A 200
Gr P 22 (2 ¼ Cr ½ Mo), A335 Gr P5 (5% Cr, ½ Mo), or A106 Gr B. The combustibles used
283
Corrosion in Refinery and Petrochemical Industries
are fuel gas, fuel oil or natural gas. In the literature, corrosion damage of THs is well
documented [1,2]. Most authors agree on the following causes:
•
•
•
•
•
Type of tube material,
Hydrogen [H2] damage,
H2-induced cracking,
Hydrogen sulphide [H2 S] damage, and
Oxidation/decarburization.
This study is part of a plant life-expectancy study and concerns two heaters after seven
years of service.
CORROSION INVESTIGATION OF TWO FURNACES: TH1 AND TH2
Operation Conditions
The operating conditions are summarized in Table 1.
Table 1. Operating Conditions of TH1 and TH2
Characteristic
Design Temperature (oC)
Operation Temperature (oC)
Design Pressure (kg/cm2 G)
Operating Pressure
Fluid
Material
Dimensions (mm)
Operating Time
Skin Temperature (°C)
Combustible
H2 Partial Pressure
2
(kg/cm G)
Steady State
TH1
TH2
452
400
225/355
280
27.0
39
21.16-23.7
29
HC + H2
Oil + H2
+ light HC
A335 Gr P11 A335 Gr P11
141.3 x 6.6
88.9 x 7.62
1972 - to date 1983 - to date
315/504
507
90% H2+10%
CxHy
0.8 max
0.8 max
Start-Up State
TH1
TH2
200
200
H2
H2
Natural gas
+ Hydrogen
Although the two furnaces appear to be identical, the corrosion damage in TH2 is very
pronounced compared to that in TH1.
Visual Examination
284
Attou et al.
After 60,000 hours of service, the outer surface of the tubes of TH1 showed a layer of
black scale 0.4 mm thick (Fig. 1). The corrosion rate was estimated to be 0.06 mm per year.
This is the usual oxidation/decarburization mechanism.
Black Scale
Oxide Layer
Base Metal
Figure 1. A photograph showing oxidation of the outer surface of tubes (100x)
The TH2 tubes showed a similar pattern. However, in TH2, certain regions of the outer
surface showed traces of yellow deposit on the oxide layer, just like a drip of a runny paint on
a wall. The yellow deposit was found to contain 0.2 1% sulphur. Under the oxide layer, a
significant number of voids and blisters were detected (Fig. 2). These voids tended to
concentrate near the weld, the heat affected zone and bends.
Base metal
2. H2-filled voids
3. Sulphur deposits
1. Oxide layer
After 60,000 hours service
Figure 2. Schematic diagram showing the degradation mechanism of the TH2 tubes
Metallographic Examination
The examination results for TH1 and TH2 (Figs. 3 and 4, respectively) are given in
Table 2.
285
Corrosion in Refinery and Petrochemical Industries
Figure 3. A photograph showing the
outer surface structure of TH1 (50x)
Figure 4. A photograph showing the
TH2 structure (50x)
Table 2. Characteristics of the Outer Surface Structure of TH1 and TH2
Characteristics
Skin
Grain Structure
Base
TH 1
Black scale 4 mm thick
Decarburization up to 0.25 mm
TH 2
Idem
-
Ferrite-Pearlite fully annealed
Ferrite-Pearlite normalised
0.041
0.088
512
128
Average Grain Diameter
(mm)
Grain Density (grains/mm2)
Hardness tests
TH1 : 155 HV5
mid - thickness
130 -160 HV5 skin
TH2 : 140 HV5
skin
Chemical Analysis of the Combustible Products
The chemical analysis of the combustible products is given in Table 3.
Table 3. Chemical Composition of Furnace Fuels
COMPOSITION
TH1
(Reforming Gas)
<150
<10
86.95
6.28
3.88
2.27
0.33
0.05
-
H2S
(ppm)
HCl
(ppm)
H2
(% vol)
C1
(% vol)
C2
(%vol)
C3
(% vol)
C4
(% vol)
C5
(% vol)
N2
He
* The TH2 fuel is a mixture of 95% vol. natural gas + 5%
vol. hydrogen fuel-gas
286
TH2
(Natural Gas)*
>150
83.54
7.656
1.951
0.544
0.199
5.621
0.19
Attou et al.
DISCUSSION
The Nelson curves, Fig. 5, as compiled in API 941[2] indicate the safe area of operation
for carbon steel and low alloy steel, as related to hydrogen partial pressure and temperature.
Referring to these curves and the operating conditions in Table 1, it can be stated that both
THs operate safely as regard to hydrogen damage.
However, the presence of H2S in the combustible product of TH2, acts as a catalyst in
maintaining H2 in the atomic state, which diffuses easily in the relatively coarse grain
structure [3]. The accumulation of H2 in voids, many of which form at grain boundaries,
results in the development of high stresses that ultimately blister the metal (Fig. 6).
1.Oxide
2.Blistered surface
3.Moisture + sulphates = H2SO4
Figure 5. Operating limits for steel in hydrogen
Figure 6. Blistered outer surface of
service
TH1 tubes
During shutdown, the moisture and condensing vapor coming from a nearby cooling
tower react with the sulphur deposited on tubes’outer surfaces to form H2SO4 [4]. The
difference in conditions between the newly blistered surface and the scaled surface gives rise
to a galvanic cell [5], and hence to electrochemical corrosion. The frequency of start-stop
cycle for TH2 is much higher than that for TH1, by a ratio of 5 to 1 per year.
CONCLUSIONS
In refineries and chemical plants, furnace availability has become a major business
problem. This paper describes corrosion damage on two tube furnaces when service
conditions differ from those of the design and/or when not enough attention is given to the:
1. Chemical composition of structure of tube material,
2. Composition of the combustibles, and
3. Site position of the furnace.
The two furnaces discussed are made from the same material, A335 Gr P11, and
apparently operate under similar conditions. After 60,000 hours of service, both furnaces
showed corrosion damage with oxidation up to 0.4 mm thick and decarburization up to 0.25
287
Corrosion in Refinery and Petrochemical Industries
mm deep. However, one of the furnaces revealed additional corrosion damage in the form of
surface blistering. A number of tests were carried out and indicated the following causes:
1. Presence of H2 and H2S in the combustible product,
2. Low grain density, i.e., coarse grain structure of tube materials,
3. Position of the heater near a cooling tower which favors vapors condensation
during shutdown periods, and
4. Number of stop-start cycles.
REFERENCES
1. J. Baas and R. Warner, How much life is left in your olefin unit ?, Hydrocarbon
Processing, December 1992, pp. 81-87.
2. API Publication 941, Steel for Hydrogen Service at Elevated Temperatures and Pressures
in Petroleum Refineries and Petrochemical Plants, 4th edition, April 1990.
3. A.C. Million, L'Hydrogene dans les aciers et dans les joints soudés, DUNOD, Paris 1971.
4. Fired heaters and stacks, Guide for Inspection of Refinery Equipment, 2nd edition, 1976,
Chapter 1X, API, 1976.
5. J.A. Dean, Langes Handbook of Chemistry, 13th edition, McGraw-Hill Company, 1972,
pp. 6-33.
288
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
SUPER DUPLEX GRADE UNS S32750 FOR SEAWATER
COOLED HEAT EXCHANGERS
P.A. Olsson and M.B. Newman
R&D Centre, AB Sandvik Steel, S-811 81 Sandviken, Sweden
ABSTRACT
The super duplex stainless steel grade UNS S32750 is presented as a candidate material for
seawater-cooled heat exchangers. Due to its duplex structure, with roughly equal amounts of ferrite
and austenite, the material possesses high strength and superior resistance to erosion-corrosion. In
addition, the high alloying content of the material gives rise to excellent resistance to pitting and
crevice corrosion as well as to stress corrosion cracking (SCC).
The benefits of a super duplex stainless grade for this type of application are reviewed, mainly as
a comparison with other materials such as different copper-based alloys, low alloyed stainless steels
and titanium. The better resistance to erosion-corrosion is emphasized as well as a close description of
why stainless steels, despite the fact that their thermal conductivity is much lower than that of copperbased alloys, still can compete as heat exchanger material in systems operating with seawater as a
cooling medium.
Key Words: Seawater, heat exchangers, super duplex stainless steel, erosion, heat
transfer, corrosion resistance
INTRODUCTION
Selecting a proper tube material for heat exchangers working with seawater as a cooling
medium is a difficult task. Traditionally different copper-based alloys such as brasses,
bronzes and copper-nickel alloys as well as carbon steels have been used most extensively.
For severe steam-side environments high alloyed nickel-based alloys or titanium have been
chosen. In applications where copper-based tubulars have failed, the stainless steel grade
UNS S32750 (SAF 2507) is an alternative to the more expensive titanium and nickel based
alloys for retubing.
UNS S32750 is a high alloyed super duplex stainless steel, with excellent material
properties. Together with high strength, hardness and resistance to erosion-corrosion, UNS
S32750 possesses excellent resistance to localized corrosion such as pitting and crevice
corrosion as well as stress corrosion cracking (SCC), in a seawater environment. The
chemical composition and mechanical properties of UNS S32750 are given in Tables 1 and 2
below.
Table 1. Chemical Composition of UNS S32750 (Weight-Percent)
Seawater Corrosion
C (max.) Si (max.) Mn (max.) P (max.) S (max.)
0.030
0.8
1.2
0.035
0.02
Cr
25
Ni
7
Mo
4
N
0.3
Table 2. Mechanical Properties of UNS S32750
Rp0.2 (min)
550Mpa
Rp1.0 (min)
640Mpa
Rm
800-1000 MPa
A (min)
25%
Hardness (Vickers)
Average 290
SEAWATER
It is often stated that the composition of seawater is more or less the same world-wide.
The variation is, however significant. The total dissolved solids, mainly sodium chloride,
vary from approximately 8000 ppm in the Baltic Sea to 60,000 ppm in bay areas of the
Arabian Gulf. The composition on which artificial seawater is based is about 35,000 ppm
total salt which is a typical value for most seawater. The pH value in the water is
approximately 8. Seawater temperature varies between only a few degrees Celsius at great
depths and in the polar regions, and 30-35°C in areas close to the equator and in the Arabian
Gulf. When studying the corrosiveness of seawater, the main environmental factors are
chloride content, pH, temperature, residual chlorine and oxygen contents, fouling, and
galvanic and biological action as well as flow rate. In addition, the presence of impurities,
mainly sulphides, must be considered for copper-based materials used in heat exchangers
with seawater as cooling medium. The situation regarding polluted seawater will be dealt
with below.
The nominal saturation for dissolved oxygen in seawater is about 6-8 ppm at 25-30°C.
At higher temperatures the oxygen content decreases, and at lower temperatures it increases.
The dissolved oxygen may also decrease to essentially zero due to bacterial action or
biological or chemical oxygen demand. Copper-based alloys as well as stainless steels
perform best in water having enough oxygen to keep fish alive (3-4 ppm) [1]. However,
copper-based alloys do not stand up well in polluted seawater, in which dissolved oxygen has
been consumed in the decay process and sulphides are present. However, tubes of UNS
S32750 or titanium can serve successfully in such water.
Polluted Seawater
The deleterious effect of sulphides on the corrosion resistance of copper-nickel alloys in
seawater has been reported by many investigators [2]. The most important parameters,
excluding the dissolved sulphide content, affecting this behavior include the hydrodynamic
conditions, oxidation products of dissolved sulphide, pH value, degree of turbulence,
exposure time and effects of chlorine and other pollutants such as ammonia. The oxidation of
dissolved sulphide by molecular oxygen not only results in a decrease of the oxygen
concentration available for corrosion but also may cause considerable change in the pH of the
environment [3]. The corrosion rates of both copper-based alloys and stainless steel grades
are dependent on pH, making pH control of the water a major concern for operating units.
Olsson and Newman
Dissolved sulphide ions greatly reduce the corrosion resistance of Cu-Ni alloys in
aqueous systems. This effect, in flowing deoxygenated seawater, was shown by MacDonald
et al. [3]. Only a low concentration of dissolved sulphides (<0.85 ppm) is necessary to cause
a dramatic decrease of the polarization resistance. In the case of a 90Cu-10Ni alloy, this
sulphide concentration was sufficient to lower the polarization resistance by a factor of 8,
whereas for a 70Cu-30Ni alloy it was lowered by a factor of 70. Higher sulphide levels have
little further impact.
The upper sulphide concentration limit for copper-based alloys is controversial. Lee et
al. [4] found that 0.1 ppm and even lower levels of sulphide are detrimental to copper-nickel
alloys, and Carew and Islam [5] showed that small additions of sulphides dramatically affect
the resistance to erosion-corrosion of copper-nickel alloys. In contrast to this, Woods et al.
[6] concluded that 90Cu-10Ni alloys exposed in partially aerated and deaerated seawater
containing up to 8 ppm sulphides did not suffer from corrosion attacks but instead built up a
protective layer of Cu2S.
Brasses have better resistance to sulphide-induced corrosion than Cu-Ni alloys have, but
not as good as stainless steel in general and UNS S32750 in particular. UNS S32750 may
also suffer from sulphide-induced corrosion attacks under certain circumstances but not at the
levels expected in even highly polluted seawater. Tests performed in NACE solution (5%
NaCl and 0.5% CH3COOH and saturated with H2S) at room temperature, an environment
much more aggressive than polluted seawater, have shown that UNS S32750 is very resistant
to sulphide-induced cracking. No cracking occurred on the UNS S32750 samples irrespective
of the applied stress. Depending on the sulphide and oxygen levels in the water, at pH levels
far below 4, sulphides may initiate corrosion of UNS S32750. This situation is, however,
most unlikely to occur in seawater-cooled heat exchanger systems, and no such case has so
far been reported.
Microbially Induced Corrosion (MIC)
Microbiological biofilms develop on all surfaces in contact with aqueous environments
[7]. Water temperature, pH, content of organic and inorganic ions affect the microbiological
settlement. MIC is used to designate corrosion caused by the presence and activities within
biofilms. The reactions are usually localized and can include sulphide, acid or ammonia
production as well as metal deposition, oxidation or reduction.
Copper-based alloys are vulnerable to biocorrosion. Differential aeration, under deposit
corrosion and cathodic depolarization are some of the reported corrosion mechanisms for
MIC on copper-based alloys [7]. CO2, H2S and NH3 accelerate localized attack. It has also
been shown [8] that Alloy 400 (UNS N04400) is highly susceptible to MIC in Arabian Gulf
seawater. Stainless steels, mainly the low alloyed 304 and 316 austenitic grades but also the
more highly alloyed grade 904L, suffer from MIC. Favorable conditions, in general, for it to
occur are stagnant conditions, organic nutrients in the water, sediment, absence or neglect of
chlorination practices as well as the presence of chlorides and sulphates.
Most troublesome for both copper-based alloys and stainless steel materials are sulphatereducing bacteria (SRB). The SRB are notorious for interfering with the formation of and
actually removing the protective surface oxide film. Severe localized pitting can occur on
iron-bearing 90Cu-10Ni if exposed even briefly to anaerobic seawater [9]. Some SRB use
hydrogen, depolarizing the cathodic surface and accelerating attack on a crevice site. SRB-
Seawater Corrosion
induced corrosion is often characterized by an encrusted deposit over a deep pit with a black
powdery sulphide corrosion product underneath.
At temperatures below approximately 35°C, a biofilm present on a stainless steel surface
will raise the open circuit potential to +300-400 mV (SCE). At this temperature, the potential
increases the risk for localized attack on standard austenitic grades but not for a high alloyed
material such as UNS S32750. At temperatures above 40°C, the micro-organisms in the
biofilm are no longer active resulting in a potential drop to essentially 0 mV (SCE) for the
stainless steel material. Since localized attacks such as pitting are dependent mainly
(excluding solution content) on potential and temperature the potential drop at higher
temperatures is advantageous. Thus at low temperatures, the open circuit potential is elevated
but not so much that it will initiate localized attack on UNS S32750, and at high temperatures
the potential is kept low so corrosion attack will for that reason not initiate. Tests carried out
at the Sandvik Steel R&D Centre have shown that UNS S32750 does not suffer from
corrosion attacks at 95°C in natural seawater. There have not been any reports of MIC
attacks on UNS S32750 in operation.
Chlorination
In order to prevent the buildup of biofilm and to remove already attached biological
species, it is common practice to chlorinate the seawater by the addition of a hypochlorite
solution to the system. The usual objective is to keep residual chlorine levels below 0.5 ppm
at the inlet tube sheet, but this level sometimes is exceeded and the residue is normally
higher at the point of injection [1]. There is a marked difference in behavior between copperbased alloys and high alloyed stainless steel. Zanoni et al. [10] investigated the iron and
manganese alloyed copper-nickel alloy C71640 and its behavior in flowing natural seawater
with different chlorination levels. It was shown that 0.5 ppm free chlorine is the upper limit
for the material to maintain good resistance to localized corrosion.
With intermittent chlorination, which is used for economic and environmental reasons,
higher doses of chlorine are used but during limited time periods. A common time period for
chlorination is 2 x 15 minutes each day. The amount of chlorine added to the system varies,
but levels up to several parts per million are not uncommon for heat exchanger cooling
systems. This can make low alloyed stainless steels and copper-based alloys unsuitable for
these systems. Grade UNS S32750 can, however, maintain its resistance to corrosion even at
these high levels of chlorination.
When added continuously, the chlorination of seawater increases the corrosion potential
for the material to approximately +600 mV (SCE). At this high potential level, localized
corrosion attacks will initiate on UNS S32750 surfaces at 75-80°C in stagnant conditions and
at approximately 10°C above this at flow rates typical for seawater cooled heat exchanger
units. When adding chlorine intermittently, the open circuit potential will be lower than with
continuous chlorination but will vary with sharp potential peaks at the time of chlorine
addition.
EROSION-CORROSION
Olsson and Newman
Erosion-corrosion can be defined as the acceleration of attack caused by a rapidly
flowing corrodent sometimes containing solid particles capable of causing erosion or wear.
Soft metallic materials such as copper, brass, aluminium and lead are prone to this form of
attack, but most metals are susceptible to erosion-corrosion to some extent in particular flow
situations. This form of attack is the main problem that besets copper alloys in seawater
cooling systems. As the flow rate increases, brasses and bronzes become more prone to
impingement attack. Aluminium brass and copper-nickel alloys offer greater resistance to
higher flow rates, but both have maximum limits which must not be exceeded or the surface
film will be destroyed. The resistance of some copper-based alloys can be improved with
small quantities of iron present in the alloy or the water. The iron apparently produces a
tougher film. This has led to the use of iron sacrificial pieces in the water boxes of heat
exchangers using copper-based tubes and tube plates [11].
Normal flow velocities within heat exchangers are in the range between 1.5 and 3.0 m/s,
but in extreme cases, the velocity of the seawater can rise to 4.5 m/s [1]. Most of the copperbased alloys cannot resist the erosion-corrosion of seawater at flow velocities above 3 m/s, a
fact that excludes most brasses and bronzes, which perform best at water velocities below
2.2-2.5 m/s, from proper seawater cooling service for heat exchangers. Copper-nickel alloys
have better erosion-corrosion resistance than brasses and bronzes but are limited in their
applications.
In polluted seawater, the situation is worse. Carew and Islam [5] performed erosioncorrosion tests in seawater with varying sulphide additions (0, 1, 3, 5 ppm). In the test rig,
the seawater had velocities in the range between 0.2 and 8 m/s and except for 90-10 and 7030 Cu-Ni alloys also a duplex grade (UNS S32550) and a few nickel-based alloys were
tested. It was concluded that the 90-10 Cu-Ni alloy suffered from severe erosion damage
with no sulphides added. With a 1 ppm sulphide concentration, erosion damage was also
seen on the 70-30 Cu-Ni alloy. Neither on the duplex grade nor on the nickel-based alloys
could erosion damage be seen. At the higher sulphide levels, the erosion-corrosion seemed to
decrease.
Seawater used in cooling systems for heat exchangers and condensers often contains
small amounts of sand particles. This severely increases the risk of erosion damage to
copper-based alloys. Most pronounced is the effect on units equipped with brass tubing
material, and attacks are frequently found at the inlet ends of condenser tubes and heat
exchangers. Stainless steels, in general, and duplex grades, in particular, have much better
erosion-corrosion resistance than all copper-based alloys. There is no risk of erosioncorrosion damage of UNS S32750 at the velocities occurring in seawater-cooled heat
exchangers. Erosion-corrosion testing performed at SINTEF [12] showed that UNS S32750
had low erosion-corrosion rates in seawater containing 0.25% of silica sand particles at a flow
rate as high as 18.3 m/s.
HEAT TRANSFER
The thermal conductivity of stainless steel is much lower than that of copper-based
alloys. If this were the only criterion for selecting a proper alloy for the heat exchanger
application, stainless steel would never have been selected. In Table 3, the thermal
conductivities of some different materials are given [13].
Seawater Corrosion
Table 3. Thermal Conductivities (W/mK) of Some Materials
(the values are all approximate)
Carbon and
Low
Alloyed Steel
40-55
Stainless Steel
Including
UNS S32750
15
Brasses
120
Bronzes
70
Cu-Ni
Alloys
Glass
30-50
0.9
As can be seen from Table 3, UNS S32750 has a thermal conductivity 2-10 times lower
than all copper-based alloys. However, the heat transfer properties of tubular products cannot
be based exclusively on the thermal conductivity of the material itself. Other factors such as
the steam film, corrosion products, deposits and cooling water film must be considered.
Actually, the thermal conductivity of the tubular material in a heat exchanger has so little
influence on the overall heat transfer performance of the system that this term can be
neglected when calculating heat transfer rates. Table 4 illustrates the various factors affecting
resistance to heat transfer in actual service. It is obvious that surface films affect overall
performance to a far greater degree than the metal wall, which accounts for no more than 2%
of the total resistance to heat flow.
Table 4. Influence of Various Factors on the Heat Resistance of a Water-Cooled
Heat Exchanger [14]
Steam Side
Water Film
18%
Steam Side
Fouling
8%
Tube Wall
2%
Water Side
Fouling
33%
Water Side
Film
39%
Despite poorer thermal conductivity properties, UNS S32750 offers as good heat flow
properties as copper-based alloys do for seawater-cooled heat exchanger and condenser units.
GALVANIC CORROSION
Galvanic processes occur between different metals when they are coupled to each other
via a conductive media such as seawater. Current will flow with oxidation taking place at the
anode, and reduction (normally oxygen reduction) occurring at the cathode. If the two metals'
free corrosion potentials, Ecorr, differ by hundreds of millivolts from one another, the metal
with the more noble Ecorr value will become predominantly cathodic and the metal with the
more active Ecorr value will be the predominantly anodic one and will corrode. If the
difference is of the order of tens of millivolts, galvanic corrosion is less likely. In Table 5, the
corrosion potentials in natural seawater for some metals and alloys are given. The values shall
all be seen as approximate values, since the corrosion potentials differ with both flow rate and
temperature of the seawater.
Olsson and Newman
Table 5. Galvanic Series in Seawater (the values are all approximate)
Metal/
Alloy
Voltage
(mV/SCE)
Zn
Steel
Brasses
-1050
-600
-300
Cu-Ni
Alloys
-200
Ti
0
UNS
S32750
+50
Pt
+250
According to the text above and Table 5, galvanic corrosion is likely to occur copperbased tubulars are replaced with stainless steel or titanium tubing if the tube sheets are not
replaced at the same time. Investigations made on severely corroded, copper-based, alloy
tube sheets in connection with stainless steels and titanium have shown that the whole inside
of the noble alloy tubes becomes effective as a cathode in copper-based alloy coupled to
stainless or titanium [15].
Stainless steel or titanium tubing increases copper-based alloy tube sheet attack to a
point where impressed current cathodic protection normally is required to control tube sheet
corrosion. If titanium is to be used as replacement tubes, the cathodic protection unit must be
controlled to avoid hydriding of the tubes. The oxide film that normally covers the surface of
titanium has been shown to be an effective barrier to hydrogen penetration. In environments
where this film is unstable or where the oxide film has been damaged by abrasion, hydrogen
absorption can occur [16]. At potentials below approximately -700 mV (SCE), production of
hydrogen will occur and the risk for hydriding of titanium increases. If zinc is used as a
cathodic protection material, the potential will be significantly lower than this in seawater, 1000 mV (SCE), and hydrogen absorption is most likely to take place in the titanium tubulars.
The R&D Centre at AB Sandvik Steel has performed tests on UNS S32750 in connection
with zinc in solutions comparable to seawater at a potential of -1000 mV (SCE). The tests
were conducted at a temperature of 80°C, and the results show that hydrogen embrittlement
will not harm the material in normal service situations. If pure iron is used instead of zinc as
a cathodic protection material the risk for hydrogen embrittlement will totally disappear for
UNS S32750. In contrast, it has been observed [16] that hydriding of titanium occurs most
frequently in solutions containing H2S where galvanic couples with iron exist. Thus, when
choosing UNS S32750 instead of titanium when replacing copper-based alloy tubulars in heat
exchangers, the problem with hydrogen absorption and embrittlement is reduced to a
minimum if the tube sheets are correctly cathodically protected.
STRESS CORROSION CRACKING
Low alloyed stainless steel grades, such as types 304 and 316, are prone to chlorideinduced SCC at temperatures above 50°C which is one of the reasons why these materials
cannot be used in many of the seawater-cooled heat exchanger systems operating today.
Duplex grades do, however, possess very good resistance to stress corrosion due to their dual
microstructure and alloying contents. In seawater cooling systems for condensers and heat
exchangers, UNS S32750 can be considered as being immune to SCC, and no such case has
been reported. Neither has titanium suffered from SCC in seawater-cooled heat exchangers.
Copper-based alloys do not normally suffer from SCC in these applications, but brasses have
a clear tendency towards SCC in the presence of ammonia or ammonium salts [17].
Seawater Corrosion
LOCALIZED CORROSION
Stainless steels suffer from localized attacks, i.e., pitting and crevice corrosion, in
chloride-containing solutions such as seawater. Crevice sites are unlikely to occur if the
chlorination procedures are carried out correctly and a continuous flow is maintained inside
the tubes. Much of the sediment entrained in cooling water deposits out in the bottom of
condenser and heat exchanger tubing at velocities below about 1-1.5 m/s and during
shutdowns. Any unit designed or operated at lower velocities is a prime candidate for under
sediment crevice attack, MIC or both [15]. Therefore, it is of great importance to rinse the
inside of the tubes during shutdowns with freshwater, independent of material used for the
tubulars.
As for crevice attacks, the primary cause of pitting is the presence of chloride ions.
Often pitting attacks are initiated at defects at the surface like slag inclusions or chromiumrich precipitates around which a chromium depleted zone is present. Welds are very likely
the initiation sites for localized attacks because of their sometimes inhomogenous structure
and the presence of intermetallic precipitates in the melted or heat-affected zones adjacent to
the welds. Therefore, seamless tubes are preferred to welded ones.
Pitting is temperature-dependent, but will not occur in UNS S32750 in seawater at
normal potentials. However, chlorination of the seawater increases the open circuit potential
to approximately +600mV (SCE). At this high potential, pitting occurs very rapidly in low
alloyed stainless steels. For UNS S32750, the risk of pitting in seawater systems is very low,
provided the temperature is kept below 75-90°C, depending on the flow rate of the cooling
water.
Tests performed at the R&D Centre at AB Sandvik Steel have shown that the critical
pitting temperature (CPT), i.e., the lowest temperature at which pitting initiates in a given
environment, increases with flow rate. The tests were performed in synthetic seawater with
the material being put to a potential of +600 mV (SCE). For stagnant conditions at this
potential, the CPT is 75-80°C. For flow rates above 1.6 m/s, the lower flow rate limit for
most heat exchangers, it was shown in the tests performed at Sandvik that pitting would not
initiate at temperatures below 85°C; CPTs at 90°C and above were measured for flow rates at
2.2 m/s. According to Quik and Gedeuke [18], super duplex grades of the UNS S32750 type
can be used in unchlorinated seawater cooling systems up to 100°C provided a constant flow
can be assured.
CASE HISTORY 1: SEAWATER CONTAINING H2S
In a condenser equipped with copper-nickel material (90Cu-10Ni tubes and 70Cu-30Ni
tube sheet), the tubes started leaking after less than four weeks in service. The seawater
contained H2S in an amount such that the environment became very corrosive to copper-based
alloys. As replacement tubing, the main candidate materials were UNS S32750 and titanium.
Due to the risk of hydriding of the titanium tubes because of hydrogen generation when
cathodically protected (to suppress the galvanic action between the titanium and the Cu-Ni
tube sheet) and vibration effects due to the lower rigidity of the thin-walled titanium tubes,
thereby causing risk of fatigue failure, UNS S32750 was chosen as the replacement material,
and the tubes were installed in October 1995.
Olsson and Newman
CASE HISTORY 2: SEAWATER CONTAINING SAND
Seawater containing sand can be a very severe environment for copper-based alloys,
particularly for brasses. An example of this is a vacuum overhead condenser operated by a
refinery in Singapore which failed in such an environment. At flow rates as low as 1.5 m/s
the admiralty brass system started to leak because of erosion-corrosion effects. The
maximum seawater temperature when the failure occurred was 60°C.
Flow rates between 1.0 and 1.5 m/s are seen by most operators as the minimum rate to
prevent sediment build-up in seawater cooling systems. These two facts, that flow rate
should be kept above 1.0-1.5 m/s and that admiralty brass failed at a fluid velocity of 1.5 m/s,
imply that admiralty brass must be seen as an unsuitable material for cooling systems
working with seawater containing sand particles. In 1993, UNS S32750 tubes were installed,
and they have performed perfectly since.
CASE HISTORY 3: SAF 2507 LIMITING PARAMETERS
In a heating medium dump cooler operating on an oil platform UNS S32750 tubes
suffered from localized attack and started leaking after having been exposed to seawater at
too high a temperature. Initially the cooler was equipped with UNS S31254 tubes which
failed after 18 months in service. UNS S32750 tubes were used as replacement material. The
service conditions inside the tubes, when working properly, i.e., with a continuous flow, were
aerated, chlorinated (0.8 ppm) seawater at a temperature of T=27-33°C.
However, stagnant conditions or very low flow rates occurred every now and then in the
cooler which made the seawater increase its temperature because of a warmer (100-240°C)
mineral oil on the shell side. Temperatures above 80°C at certain "hot spots" adjacent to the
baffles, combined with chlorides from the seawater, made pitting of the tubes possible. After
30 months of service, the UNS S32750 material started leaking.
DISCUSSION
By emphasizing different aspects of seawater, and its influence on the heat exchanger
material, it has been shown that within certain limits UNS S32750 is a better choice than
mainly copper-based alloys, and also, due to the duplex grade's lower price, a better choice
than titanium and nickel-based alloys for seawater-cooled heat exchanger applications. UNS
S32750 possesses extremely good erosion-corrosion properties compared to copper-based
alloys, which have been shown to suffer from this form of attack even at very low flow rates
of the fluid. Not only the fluid velocity of the water, but also the sand, H2S, CO2 and NH3
contents play an active role in the erosion and corrosion behavior of the copper-based alloys.
To prevent biofilm formation of the metal surface exposed to seawater, chlorination of
the water is widely used. With constant chlorination the levels of residual chlorine in the
system most often are in the range between 0.1 and 0.4 ppm. The upper limit for copperbased alloys is approximately 0.5 ppm chlorine, whereas UNS S32750 can tolerate as much
as several parts per million [19]. The trend today is toward the use of intermittent
chlorination, which will lead to high chlorine levels (several parts per million) for shorter
periods each day. With the use of UNS S32750, there should be no corrosion problem due to
these higher levels.
Seawater Corrosion
If exclusively taking thermal conductivity levels of different materials into account, one
can conclude that UNS S32750 has approximately 2-10 times worse properties than copperbased alloys. As shown in this paper, however, the thermal conductivity of the tube material
contributes only 2% of the overall heat transfer of the system and can thus be neglected.
Seawater can, depending on its content and temperature, be corrosive to copper-based
alloys as well as to high alloyed stainless steels such as UNS S32750. In natural seawater
with no chlorine added, temperatures as high as 95°C have been shown to not be high enough
to initiate corrosion attacks alone. When normal continuous chlorination procedures are used,
however, the open circuit potential rises to approximately +600 mV (SCE) for UNS S32750.
At this potential, the CPT for UNS S32750 is 75-80°C for stagnant conditions and 10°C
higher if a continuous flow above 1.5 m/s is maintained.
CONCLUSIONS
1. Copper-based alloys suffer from erosion-corrosion damage at flow rates occurring in
seawater-cooled heat exchangers. UNS S32750 does not suffer from erosion at fluid
velocities far above those likely to be maximum for heat exchanger applications.
2. Normal chlorination procedures, continuous or intermittent, will not initiate corrosion
of UNS S32750 as long as temperature and flow rate are kept within specified ranges.
3. Despite the lower thermal conductivity of UNS S32750 compared to copper-based
alloys, the overall heat transfer capacities of the different material systems are equal.
4. UNS S32750 does not suffer from pitting corrosion attacks at 95°C in unchlorinated
seawater or at 75-80°C in stagnant chlorinated solutions. With flow rates above 1.5
m/s, typical for heat exchanger systems, tests have shown the critical temperature
limit to be above 85°C. These test results are good indications of the temperature
intervals within which UNS S32750 can be used.
5. With much better erosion and corrosion properties than copper-based alloys and with
properties at least as satisfactory as those of titanium and nickel-based alloys, UNS
S32750 must be considered to be one of the very best alternatives for seawater-cooled
heat exchangers in need of tube replacement.
REFERENCES
1.
2.
3.
4.
5.
6.
7.
8.
Tuthill, Chemical Engineering 97, 1, 1990, pp. 120-124.
Reda and Alhajji, British Corrosion Journal 30, 1, 1995, pp. 56-62.
MacDonald, Syrett and Wing, Corrosion 35, 8, 1979, pp. 367-378.
Lee, Hack and Tipton, Proceedings 5th International Congress on Marine Corrosion
and Fouling, May 1980, Barcelona, Spain, Comité International Permanent pour la
Recherche sur la Preservation des Materiaux en Mileau Marin.
Carew and Islam, Materials Performance 34, 4, 1995, pp. 54-57.
Woods, Amos and McNeil, Corrosion/92, April 1992, Nashville, Tennesse, USA,
Paper No.180.
Wagner and Little, Materials Performance 32, 9, 1993, pp. 65-68.
Gouda, Banat, Riad and Mansour, Corrosion 49, 1, 1993, pp. 63-73.
Olsson and Newman
9. MTI, Performance of tubular alloy heat exchangers in seawater service in the
chemical process industries, Publication No.26, St. Louis, Missouri, USA, Materials
Technology Institute of the Chemical Process Industries, Inc. & NiDi, August 1987,
pp. 8-9.
10. Zanoni, Gusmano, Montesperelli and Traversa, Corrosion 48, 5, 1992, pp. 404-410.
11. Tretheway and Chamberlain, Corrosion for Students of Science and Engineering, 2nd
ed., Hong Kong, Longman Group Limited, 1990, pp. 285-288.
12. Rogne and Solem, Erosion-corrosion testing of stainless steels, SINTEF Corrosion
Center, Publication No. STF34 F93017, February 1993.
13. Hedner (ed.), Formelsamling i Hallfasthetslara, The Department of Mechanics of
Materials, The Royal Institute of Technology, Stockholm, Sweden, 1990, pp. 205-213.
14. Davison and Miska, Chemical Engineering 86, 2, 1979, pp. 129-133.
15. NiDi, Guidelines for selection of nickel stainless steels for marine environments,
natural waters and brines, Publication No.11003, NiDi reference book series, 1987,
pp. 3-6
16. Covington, Corrosion 35, 8, 1979, pp. 378-382.
17. Rückert, Werkstoffe und Korrosion 47, No.2, 1996, pp. 71-77.
18. Quik and Gedeuke, Stainless Steel Europe 6, No.10, 1994, pp. 46-51.
19. Gartland and Drugli, Crevice corrosion of high alloyed stainless steels in chlorinated
seawater, Part I: Practical aspects, Corrosion /91, March 11-15, 1991, Cincinnati,
Ohio, USA, Paper No.510.
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
EVALUATION OF ALUMINUM ALLOY 5083 WELDMENTS
TO STRESS CORROSION CRACKING IN SEAWATER
A. Saatchi1, M.A. Golozar1 and R. Mozafarinia2
1
Materials Engineering Department, Isfahan University of Technology,
Isfahan, Iran
2
Defense University, Shahin-Shahr, Iran
ABSTRACT
The susceptibility of aluminum-alloy 5083 and its weldments to stress corrosion cracking (SCC)
was investigated using constant strain and slow strain rate (SSR) tests in actual seawater, in the Bandar
Bushehr region, and in 3.5% NaCl solution, respectively. Also, electrochemical polarization and
immersion tests were performed to study the corrosion rate of the aluminum alloy and its weldments.
In addition, long-term tests, on actual sections of a marine vessel for a period of one year, were used to
investigate the corrosion behavior of alloy 5083 and its weldments in Bushehr seawater. Scanning
electron microscopy (SEM) was employed to characterize the nature of corrosion and SCC in the
specimens tested under various conditions.
The results indicated higher corrosion rates in 3.5% NaCl solution than in seawater. Also, the
corrosion rates of welded specimens were higher than those of unwelded specimens. In addition, the
results revealed no susceptibility to SCC of the alloy and its weldments in Bushehr seawater.
Key Words: Aluminum alloy 5083, stress corrosion cracking, weldments, Persian Gulf
INTRODUCTION
Aluminum alloy 5xxx series are non-heat-treatable alloys containing Mg and are widely
used for the construction of marine structures and vessels. On the other hand, during the
welding of these alloys, Mg2Al3 precipitates along the grain boundaries. Mg2Al3 is anodic
with respect to the base metal, and thus may render the alloy susceptible to stress corrosion
cracking (SCC). SCC often causes unexpected costly failures. This is due to a lack of visual
evidence of corrosion as compared with uniform and localized corrosion. SCC is a cracking
phenomenon under the conjoined action of tensile stress and corrosion, which happens for
some metals in specific environments. In other words, there are three main factors acting
together in SCC: tensile stress, which can be working or designed stress, or residual and
thermal stresses due to welding and other operations; the corrosion environment which is
affected by major and minor chemical composition, temperature, velocity or turbulence, and
even biological activities of the environment; and the susceptibility of the metal to SCC,
which can be affected by the chemical composition, microstructure, and thermal and
mechanical operations.
301
Seawater Corrosion
Seawater is a complex biochemical broth. Its nominal chemical composition is 2.9%
NaCl plus 0.4% MgSO4, but in reality it contains traces of just about everything one can
imagine [l]. Its high chlorinity makes it corrosive to most metals and alloys. Chloride also
causes localized corrosion and SCC in many metals. The action of biological organisms also
create a complex situation for predicting the performance of alloys in seawater. So whereas
there are cases in which the results of SCC of steel in nominal seawater are identical to the
results in 3.5 NaCl solution [2,3], there are other cases where different behavior has been
reported [4].
The purpose of this study was to evaluate the SCC susceptibility of aluminum alloy 5083
and its weldments in Persian Gulf water. The test station was located in Bandar Bushehr,
toward the north of the Persian Gulf. Table 1 shows the chemical composition of the Bushehr
seawater. The salinity of the water is high, which is manifested by its 2.34% chloride
content. In the laboratory, slow strain rate (SSR) tests were performed in natural seawater
and in 3.5% NaCl solution. In the field, constant strain rate tests were performed in Bandar
Bushehr seawater for 12 months at various depths below the water’s surface. Also, two
specimens simulating the most critical sections of a marine vessel with regard to welding,
manufacturing and internal operational stresses were immersed at a depth of 5 m below the
water’s surface for 1 year. In addition, the corrosion rates of the base metal and the weld
metal were measured using weight loss and electrochemical polarization tests. Optical and
scanning electron microscopy (SEM) were also employed to investigate the microstructure of
the weld metal, heat affected zone (HAZ) and fracture surfaces of the specimens.
Table 1. Chemical Composition of Bandar Bushehr Seawater
2+
2+
Dissolved
solids (%)
Total Hardness
(ppm)
Ca
(ppm)
Mg
(ppm)
pH
6052
1335
7387
4.09
7.8
Dissolved
O2 (mg/l)
5.98
Salinity
(gm/l)
Cl
(%)
42.48
2.34
EXPERIMENTAL PROCEDURE
Specimens Preparation
The specimens for the various tests were cut from a 5 mm plate of the Al alloy 5083-O.
The chemical composition and mechanical properties of the alloy are shown in Tables 2 and
3, respectively. For weldments, the one-sided V-notch was used. Tungsten metal arc with
inert gas (TIG) and filler material 5183 were used for welding.
Table 2. Chemical Composition of Aluminium Alloy 5083-O (in wt.%)
Si
0.4
Fe
0.4
Cu
0.1
Mn
0.4-1.0
Mg
4.0-4.9
Cr
0.05-0.025
Zn
0.25
Table 3. Mechanical Properties of Aluminum Alloy 5083-O
302
Ti
0.15
Al
balance
Saatchi et al.
Temper
O
Tensile Stress
(MPa)
290
Yield Stress
(MPa)
145
Elongation
(%)
22
Immersion Tests
Immersion tests were used to measure the corrosion rate of the base metal, with and
without welding, using the ASTM Gl standard. The specimens dimensions were 20 x 50 x 5
mm.
Polarization Tests
The cathodic and anodic polarization of the alloy and weld metals in actual seawater and
3.5% NaCl solution were determined by a potentiodynamic technique based on ASTM G587. The corrosion rates of the specimens were also calculated using the linear polarization
technique. The electrochemical measurements were performed using a standard potentiostat,
Wenking model ST72, with a Ag/AgCl reference electrode.
Figure 1. The details of specimens used for constant strain stress corrosion
tests
(a) unwelded; (b) cross weld; and (c) longitudinal weld
303
Seawater Corrosion
Constant Strain SCC Tests
Figure 1 shows the details and types of specimens which were used for these
experiments. The test plates on these fixtures were either plain, i.e., without welding (Test
Piece a) or welded (Test Pieces b and c). The welding direction was either parallel to the
rolling direction of the plate or perpendicular to it. Constant strain SCC tests were performed
according to ASTM G39. The specimens were exposed to Persian Gulf seawater off the
Bandar Bushehr coast at depths of 5, 7, and 9 m for 12 months. The applied stress on all the
specimens was the yield stress of the alloy. In order to avoid galvanic and crevice corrosion,
the H-beam, nuts and bolts at both ends of the specimens were made from the same alloy.
Also, two specimens (Part Nos. l and 2) simulating the most critical condition of a marine
vessel (shown in Fig. 2) were immersed at a depth of 5 m below water’s surface in Bandar
Bushehr for one year. Part No. 1 had severe forming stresses along with welding residual
stresses. Part No. 2 was a hollow vessel with extensive welds containing air at a pressure of
3.5 atm.
Figure 2. Part Nos. l and 2 simulating the most critical condition of a marine vessel
Slow Strain Rate (SSR) Tests
Figure 3 shows the details of the specimens which were used in SSR tests. For these
tests an SSR testing machine, based on the recommendations of Parkins [5], was designed and
-6 -1
constructed. The strain rate used was 3 x l0 s . This was based on the previous results
obtained [6,7]. Stress-strain curves for specimens without welding and welded in the gauge
304
Saatchi et al.
length area were obtained. These curves were plotted in various conditions including: air,
seawater and 3.5% NaCl solution. The fracture surfaces of the SSR specimens were studied
using SEM (JEOL model S6400).
Figure 3. Details of the specimens used in the SSR tests: (a) unwelded; and (b) welded
Intergranular Corrosion Tests
In order to determine the extent of precipitation of Mg2A13 and its effect on
intergranular corrosion, intergranular corrosion tests were performed on specimens with the
dimensions 50 x 6 x 5 mm in concentrated nitric acid based on ASTM G67-80. The weight
loss was measured.
RESULTS AND DISCUSSION
Figures 4 and 5 show the potentiodynamic polarization curves of the base metal and
weld metals in Bandar Bushehr seawater and in a 3.5% NaCl solution. All the data in these
curves are summarized in Table 4. The corrosion rates obtained for the base metal and weld
metal using the linear polarization technique and immersion tests are also summarized in
Table 4. In laboratory tests, specimens show a lower corrosion rate in seawater than in 3.5%
NaCl solution. This could be due to the lower salinity of the seawater compared to the 3.5%
NaCl solution. In all the test methods, welded specimens showed higher corrosion rates than
unwelded ones. This is due to the precipitation of anodic Mg2Al3 in grain boundaries in the
HAZ during welding. Table 3 shows that the corrosion rates obtained by electrochemical
methods were an order of magnitude higher than those obtained in either immersion tests in
the laboratory or in actual field conditions. This is obviously due to the fact that the duration
of the immersion tests in actual field conditions was long enough for hard marine fouling to
develop on the surface of the metal and thus protect it from further corrosion. Table 4 also
305
Seawater Corrosion
shows that the corrosion rates in actual field conditions were higher than the corrosion rates
in the laboratory immersion tests. This clearly indicates that the corrosion conditions in the
sea were more intense than in the laboratory situations.
Figure 4. Potentiodynamic polarization
curves for aluminium alloy 5083 in
seawater and 3.5% NaCl solution
Figure 5. Potntiodynamic polarization
curves for welded alumimium alloy
5083 in seawater and 3.5% NaCl
solution
Table 4. Polarization and Immersion Test Results
Specimen
type
Environment
Polarization Test Results
Corrosion
Corrosion
Corrosion
Potential
Rate
Current
Density
(mA/cm2)
0.054
(mV)
-859
-2
0.105
-765
1.2 x 10
-2
0.066
-848
1.1 x 10
-2
0.095
-800
9.2 x 10
(mpy)
Unwelded
Seawater
2.4 x 10
Unwelded
3.5% NaCl
Solution
4.6 x 10
Welded
seawater
2.9 x 10
Welded
3.5% NaCl
Solution
4.4 x 10
Immersion
Tests in
Laboratory
Corrosion Rate
(mpy)
-2
8.4 x 10
-4
-3
-3
-4
Immersion Tests in
Bandar Bushehr
corrosion Rate
(mpy)
1 x 10
-3
--
1.54 x 10
-3
--
Figure 6 shows typical constant strain rate specimens and Part Nos. l and 2 that were
immersed at various depths below the water’s surface in Bandar Bushehr region for 12
months before cleaning. Severe hard fouling can be seen. After cleaning the specimens, they
were inspected for cracks. No cracking was observed in either of the specimens. These
results, summarized in Table 5, indicate that neither the base metal, nor the welded specimens
show any cracking under stress after 12 months of exposure to Bandar Bushehr seawater.
306
Saatchi et al.
Figure 6. Constant strain test specimen and Part Nos. 1 and 2 before cleaning exposed in
Bandar Bushehr seawater for 12 months
Table 5. Results of Constant Strain Tests and Simulated Service Condition Tests in Bandar
Bushehr Seawater for 12 months
307
Seawater Corrosion
Specimen Type
Unwelded
Transverse weld
Longitudinal weld
Unwelded
Transverse weld
Longitudinal weld
Part No. 1
Part No. 2
Depth (m)
Splash zone
Splash zone
5
5
5
5
5
5
Results
No cracks
No cracks
No cracks
No cracks
No cracks
No cracks
No cracks
No cracks
Figures 7 and 8 show the typical stress strain curves for the aluminum alloy 5083 with
and without welding obtained by the SSR testing machine. The specimens were tested in air,
as well as in seawater and 3.5% NaCl solution. The results obtained from these tests are
summarized in Table 6. There is no indication of SCC susceptibility from these test results.
Figure 7. Load versus strain for aluminum
alloy 5083 in air, seawater, and 3.5%
NaCl solution
Figure 8. Load versus strain for welded
aluminum alloy 5083 in air, sea
water, and 3.5% solution
Figure 9 shows a typical SEM fractograph of the fractured surface of the SSR test
specimens. All the fractured surfaces showed dimple fracture and no indication of
susceptibility to SCC in the environments tested.
Table 7 summarizes the intergranular test results in nitric acid, according to ASTM G6780. The data shows that the average weight loss of the unwelded specimens was 2.46
mg/cm2, and that of the welded specimens was 3.54 mg/cm2. These rates are below the range
of integranular susceptibility according to the ASTM standard.
Table 6. SSR Test Results on Aluminum Alloy 5083
308
Saatchi et al.
Specimen
Environment
Unwelded
Welded
Unwelded
Welded
Unwelded
Welded
Air
Air
Seawater
Seawater
3.5% NaCl solution
3.5% NaCl solution
Failure Time
(h)
24.40
17
24.30
21.16
26
20.43
Max Load
(N)
1900
1100
1600
1802
1540
1610
Failure Load
(N)
1490
1050
1380
1690
1320
1600
Failure Stress
2
(N/mm )
99.3
70
92
113
88
106.6
Figure 9. Typical SEM fractograph of the fractured surface of an SSR test specimen in
seawater: (a) unwelded; and (b) welded
Table 7. Intergranular Test Results on Welded and Unwelded Aluminum Alloy 5083
Specimen
Area
(mm )
1333
Initial
Weight
(gr)
4.0915
Final
Weight
(gr)
4.0515
Weight
Loss
(gr)
0.04
Weight
Loss
2
(gr/cm )
3
1227
3.8998
3.8765
0.0233
1.91
2.46
1552
1025
6.2633
2.8906
6.2028
2.8581
0.0605
0.0325
3.9
3.17
3.54
2
Unwelde
d
Unwelde
d
Welded
Welded
Average
Rating
2
(mg/cm )
acceptable
acceptable
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Seawater Corrosion
CONCLUSIONS
1. Constant strain specimens and also simulated marine vessel parts made from aluminum
alloy 5083 with and without welding, immersed in Persian Gulf seawater at various
depths for 12 months did not show any cracking tendency.
2. SSR test results and SEM fractographs of the fractured surfaces did not show any
indication of SCC susceptibility of this alloy with or without welding.
3. Welding aluminum alloy 5083 with TIG, increased the corrosion rate, due to grain
boundary precipitation, but the intergranular corrosion rate remained in the acceptable
range.
REFERENCES
1. G.A. Gehring, Materials Performance 9, 1987, p. 9.
2. G. Sandoz, Stress Corrosion Cracking in High Strength Steels and in Titanium and
Aluminum alloys, B.F. Brown, ed., Naval Research Laboratory, 1972.
3. G. Sandoz, Metallurgical Transactions 2, 1971, p. 1055.
4. H.P.Lockie and A.W. Loginow, Corrosion 14, 1968, p. 291.
5. R.N. Parkins, ASTM STP 665, 1987, pp. 2-25.
6. M. Sobhani, M.Sc. Theses, Materials Engineering Department, Isfahan University of
Technology, 1990.
7. R. Mozafarinia, M.Sc. Theses, Materials Engineering Department, Isfahan University of
Technology, Iran, 1994.
310
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
CAVITATION CORROSION BEHAVIOR OF SOME
CAST ALLOYS IN SEAWATER
A. Al-Hashem, P.G. Caceres and H.M. Shalaby
Materials Application Department
Kuwait Institute for Scientific Research,
P.O. Box 24885, 13109 Safat, Kuwait
ABSTRACT
This paper summarizes the results of extensive laboratory research carried out on the cavitation
corrosion behavior of three cast alloys in seawater. The alloys studied were nodular cast iron (NCI),
nickel-aluminum bronze (NAB) and duplex stainless steel (DSS). The work included measurements
of free corrosion potential, potentiodynamic polarization and mass loss in presence and absence of
cavitation. The cavitation tests were made using an ultrasonically induced cavitation facility at a
frequency of 20 kHz. Cavitation caused an active shift in the free corrosion potential and increased
the anodic polarization current for the three cast alloys. In quiescent seawater, the rates of mass loss
were very low for all of the alloys studied. Cavitation increased the rates of NCI, NAB and DSS by
930, 186 and 20 times, respectively. However, the application of cathodic protection decreased the
rates of NCI, NAB and DSS by about 50%, 47% and 19%, respectively. Morphological examinations
revealed that thin scales were formed on NCI in quiescent seawater, while NAB suffered from
selective corrosion of the copper-rich α phase. On the other hand, DSS was almost free of corrosion
in quiescent seawater. Cavitation made the surfaces of the alloys very rough, exhibiting large size
cavities and ductile tearing. The cavitation damage started at the graphite nodules and ferrite matrix in
NCI, the boundaries of α columnar grains in NAB, and the ferrite/austenite boundaries and the
austenite islands in DSS. Cross-sectional examinations revealed microcracks in the bulk of the alloys.
The formation of microcracks was attributed to cavitation stresses and preferential phase corrosion.
Key Words:
Nodular cast iron, nickel-aluminum bronze, duplex stainless steel,
cavitation corrosion, microstructure, corrosion morphology, seawater
corrosion
INTRODUCTION
During the last few years, extensive research was made in our laboratory in the
investigation of the flow-dependent corrosion behavior of cast alloys used in the
manufacturing of seawater valves and pumps. This research was due to the repeated failures
of such components in the seawater systems of the petroleum refineries in Kuwait. Failure
analysis [1-3] revealed that the damage was due to a combination of several factors that
included cavitation corrosion, erosion-corrosion and preferential phase corrosion. This paper
summarizes the results of laboratory cavitation studies made on nodular cast iron (NCI),
nickel-aluminum bronze (NAB) and duplex stainless steel (DSS). The work was aimed at (1)
determining the relative susceptibility of the three cast alloys to cavitation damage; (2)
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Seawater Corrosion
determining the effect of the application of cathodic protection in minimizing the degree of
damage; and (3) characterizing the role played by the different phases of the alloys during
cavitation.
Cavitation may be defined as the growth and collapse of vapor bubbles due to localized
pressure changes in a liquid. The collapse process takes place extremely rapidly, producing
strong shock waves that damage the material. Because a few million bubbles may collapse in
a second, damage quickly occurs. Published work [4,5] has shown that both mechanical and
electrochemical factors are involved in cavitation. Many investigators [4-7] have indicated
that cavitation can be reduced by applying cathodic protection to the cavitating body.
EXPERIMENTAL PROCEDURE
A study was made on specimens machined from the base of a failed NCI valve, the upper
casing of a failed NAB pump, and commercially cast DSS. Table 1 shows the chemical
composition of the alloys studied.
Table 1. Chemical Composition of Cast Alloys Used in the Cavitation Studies
Element (%)
Cast Alloy
NCI
NAB
DSS
C
3.54
0.04
Si
2.26
0.57
Mn
0.40
1.20
0.54
P
0.05
-
Cr
24.40
Ni
4.90
5.60
Mo
2.50
Cu
80.00
-
Al
9.00
-
N2
0.12
Fe
Bal
4.90
Bal
NCI: Nodular cast iron
NAB: Nickel-aluminum bronze
DSS: Duplex stainless steel
The microstructure of NCI after polishing and etching in 4% nital solution consists of
spheriodal graphite in a ferrite matrix. The microstructure of NAB after etching in a 10%
solution of ferric nitrate consists of columnar grains of α phase, which is a face-centered
cubic (fcc) copper-rich solid solution; a small volume fraction of lamellar eutectoidal phases
of β phase or retained β, which is martensitic; and intermetallic κ phases of basically four
morphologies (designated as κI, κII, κIII, and κIV). The κI, κII, and κIV phases are all ironrich precipitates with a body-centered cubic (bcc) structure based on aluminum-ironintermetallic (Fe3Al), while the κIII phase is a nickel-rich precipitate based on the structure
of aluminum-nickel intermetallic (NiAl) [8]. The microstructure of DSS after etching in 15%
ethanolic solution consists of a ferrite matrix in which austenite grains were present as
islands.
All the specimens used in the present work were in the form of discs, measuring 1.6 cm
in diameter and 0.27 cm in thickness. However, in the case of the cavitation testing, the
specimens had threads, as they were machined in accordance with the requirements of ASTM
G32-92 [9]. Before testing, all the specimens were mechanically ground with silicon carbide
paper, cleaned, degreased, and then weighed. For the morphological studies, some specimens
were etched before testing.
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Al-Hashem et al.
The study was made in untreated seawater collected from the Arabian Gulf. This
seawater is very saline (46,000 TDS), and it has a pH of about 8.0. The composition of this
seawater has been previously reported [10].
The cavitation tests were carried out at a frequency of 20 kHz and an amplitude of 25 μm
using an ultrasonically induced cavitation facility. The facility consists of a generator, a
converter, and a disrupter horn. The test specimens were mounted on a special holder which
was placed at a distance of 0.125 cm from the apparatus horn. The specimen’s area facing the
horn, and thus exposed to cavitation, was 1 cm2. The seawater was contained in an open
600-ml Pyrex beaker surrounded by a copper cooling coil in a water bath. The seawater
electrolyte was maintained at 25 + 1oC by controlling the flow rate and the temperature of the
circulating chilled water. The cavitation tests were made under free corrosion and cathodic
protection (CP) conditions. In the later tests, the specimens were kept at a single potential
value more negative than the free corrosion potential by 50 mV for NCI and NAB and 100
mV for DSS. The specimens were weighed after different time intervals in order to determine
the rate of mass loss as a function of cavitation time. In order to compare the rate of mass
loss in the presence of cavitation with that obtained in its absence, the rate of mass loss was
also determined for specimens immersed in quiescent seawater under free corrosion
conditions.
In order to evaluate the role of cavitation on the electrochemical behavior of the
materials, open-circuit potential and potentiodynamic polarization measurements were made
under quiescent and cavitation conditions. The polarization tests were carried out using
standard equipment at a scan rate of 0.5 mVs-1 starting from a potential value more negative
than the open-circuit potential. A saturated calomel reference electrode (SCE) and a graphite
counter electrode were used in these experiments.
To follow up on the development of cavitation damage and to identify the susceptibility
of the constituent phases to cavitation, detailed morphological examinations were made on all
tested specimens using scanning electron microscopy (SEM).
RESULTS AND DISCUSSION
Effect of Cavitation on Rates of Mass Loss
The rates of mass loss under quiescent conditions were very small for all alloys. In the
presence of cavitation, the mass loss behavior was quite similar for the three cast alloys. In
all cases, the rate of mass loss initially increased till it reached a maximum. Then, it sharply
decreased with the increase of cavitation time, followed by a slow decrease till a steady-state
was reached. This behavior remained unchanged in the presence of CP. However, the rates
of mass loss of specimens cavitated under free corrosion conditions were substantially higher
than those obtained in presence of CP. Figure 1 is given as an example to show the rates of
mass loss obtained for NAB and DSS, while Table 2 gives a summary of the rates of mass
loss at the steady-state for the three cast alloys after various test conditions.
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Seawater Corrosion
(a)
(b)
Figure 1. Rates of mass loss of: (a) NAB; and (b) DSS specimens, exposed
to seawater under various conditions.
Table 2. Rates of Mass Loss (in mg.hr-1.cm-2) for Cast Alloys Exposed
to Seawater Under Various Test Conditions
Test Condition
Cavitation under free corrosion
Cavitation under CP
Quiescent
NCI: Nodular cast iron
NAB: Nickel-aluminum bronze
DSS: Duplex stainless steel
NCI
3.12
1.56
3.36 x 10-3
NAB
1.50
0.80
8 x 10-3
DSS
0.64
0.52
2.5 x 10-4
It is clear from Table 2 that the rates of mass loss of the cavitated specimens under CP
were lower than those under free corrosion conditions by about 50% for NCI, 47% for NAB
and 19% for DSS. This difference should be equivalent to the electrochemical corrosion rates
of the alloys. A comparison of the mass loss rates obtained under quiescent conditions with
those obtained under cavitation conditions indicates that the kinetics of electrode process are
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Al-Hashem et al.
enhanced by cavitation. In other words, the greatest portion of damage during cavitation is
due to mechanical breakdown of the material surface and subsequent particle separation and
ejection from the metal. However, the results clearly show that DSS is quite resistant to
electrochemical corrosion in seawater. Thus, although cavitation has exposed fresh metal
surfaces to the aggressive environment, the enhancement of electrochemical corrosion in DSS
was not as significant as in NCI or NAB.
In the present study, CP was applied to the cavitated specimens in order to separate the
electrochemical corrosion component from the mechanical one. The present results indicate
that the application of CP during cavitation reduces the rates of mass loss for all the
investigated alloys. Since the rates of mass loss in quiescent seawater were very low, this
decrease cannot be solely attributed to the elimination of electrochemical dissolution. Thus,
the reduction in the rate of mass loss during cavitation under CP could be due to a
combination of two factors, namely: bubble collapse cushioning due to the development of
cathodic gas and diminished electrochemical dissolution [6].
Effect of Cavitation on the Electrochemical Behavior of Cast Alloys
The potentials of NCI and NAB specimens exposed to seawater under cavitation and
quiescent conditions were measured as a function of time. Table 3 shows the initial potential
values obtained upon immersion and the final potentials at steady-state.
Table 3. Corrosion Potentials (in mV vs. SCE) of Cast Alloys Exposed to
Seawater under Quiescent and Cavitation Conditions
Alloy
Quiescent
Initial Values At Steady-state
NCI
-560
-740
NAB
-260
-245
NCI: Nodular cast iron
NAB: Nickel-aluminum bronze
Cavitation
Initial Values
At Steady-state
-560
-650
-325
-300
It is clear from Table 3 that cavitation causes an active shift in the corrosion potential of
NCI and a noble shift in the case of NAB. However, it is worth noting that the application of
cavitation causes an immediate active shift in the initial corrosion potential of NAB. In order
to validate this effect, the measurements were repeated on NAB and DSS under a pulsating
cavitation condition. In these experiments, the materials were alternatively exposed to
periods of cavitation and quiescence for one hour each. Figure 2 is given as an example to
show that the application of cavitation immediately shifted the potential to a more active
value. When cavitation was stopped, the potential rapidly shifted back to that of the quiescent
condition. The rapid activation of the corrosion potential can be explained in terms of the
destruction of the passive film and the creation of fresh metal surfaces. On the other hand,
the rapid reversal of the free corrosion potential when cavitation was stopped suggests that
the repassivation kinetics is quite fast. A similar potential shift was previously reported for
copper-manganese-aluminum alloys cavitated in seawater [11].
315
Seawater Corrosion
Figure 2. Effect of cavitation on the free corrosion potential of DSS in seawater
The potentiodynamic polarization curves of NCI and NAB were similar in terms of a
gradual increase in the anodic current with the increase in potential in the absence of an
active-passive transition for the cavitated and non-cavitated specimens (see Fig. 3a). The
presence of cavitation, however, increased both the cathodic and anodic currents. Moreover,
it caused a small active shift in the corrosion potential. In the case of DSS, the alloy
underwent a direct transition from active dissolution to passivity without exhibiting an
active/passive transition (Fig. 3b). Cavitation was seen to have a small effect on the
polarization behavior of DSS.
The gradual increase in the anodic current of NCI and NAB with the increase in potential
in the absence of cavitation can be ascribed to a mass transfer phenomenon controlled by two
processes, namely: formation of soluble corrosion products during anodic polarization and
diffusion processes through adherent corrosion products or film. In the case of NAB, film
formation is mostly controlled by Al ion diffusion in the Alox phase [12]. In the presence of
cavitation, the corrosion product layer is destroyed by the collapse of bubbles. Under these
conditions, the anodic current increases as a result of activated electrochemical dissolution
and the cathodic current increases as a result of increased charge transfer for oxygen
reduction.
It was rather surprising to find that cavitation had a small effect on the polarization
behavior of DSS in seawater. The direct transition from active dissolution to passivity
indicates that the alloy spontaneously passivated in seawater. Within the anodic region, there
were two competing processes of passive film destruction and repassivation. However, the
presence of cavitation possibly led to a slow-down in the repassivation process. Thus, the
polarization behavior was not significantly affected, except for the small increase in current.
(a)
316
(b)
Al-Hashem et al.
Figure 3. Potentiodynamic polarization curves for: (a) NAB; and (b) and DSS,
immersed in seawater under quiescent and cavitation conditions
Effect of Cavitation on Surface Damage
In stagnant seawater, a loosely adherent, corrosion product layer was formed on NCI.
After cavitation for 2-10 minutes, erosion markings were visually observed. When SEM
examinations were made, the NCI surface was found to have suffered from localized surface
damage, as early as after 1 minute of cavitation. The damage was in the form of
fragmentation and removal of graphite nodules. This led to the formation of circular craters
of cavities; while the surrounding matrix remained smooth and unattacked (Fig. 4a). After 5
min., microcavities (erosion pits) developed in the ferrite matrix along with other cavities
caused by the loss of graphite. After 10 to 30 minutes, the attacked areas became larger and
engulfed the microcavities. Longer cavitation testing (5.5 hours) showed ductile removal of
material from the matrix (Fig. 4b). These characteristics did not change in the presence of
CP. However, the time needed to reach a specific stage of damage was longer under CP.
Examinations of cross-sections revealed cracks, 3-21μm long, that propagated into the matrix.
These cracks originated at the bottom of deep craters.
In general, cavitation corrosion increases in corrosive liquids such as seawater due to the
combined action of electrochemical corrosion and the fluid mechanical component. Arabian
Gulf seawater seems to be more aggressive than other water bodies, as it contains higher salt
and different pollutants [10]. In the present work, erosion pits appeared randomly in the
ferrite matrix of NCI due to ductile tearing caused by the sudden collapse of liquid bubbles.
This indicates that cavitation damage in cast iron is not only initiated at graphite nodules as
mentioned by Okada et al. [13], but also in the ferrite matrix.
(a)
(b)
317
Seawater Corrosion
Figure 4. SEM micrographs of NCI after cavitation testing in seawater. (a) after 1
minute, fragmented graphite nodules and formation of cavities; and (b)
after 5.5 hours, ductile removal of material from ferrite matrix
The surface of NAB specimens exposed to stagnant seawater appeared visually darker
than the unexposed specimens. SEM examinations indicated that the α phase was
preferentially attacked at the α/κIV interfaces and the precipitate-free zone did not suffer
from any attack (Fig. 5a). Dissolution of the α phase was also noted around the dendritic
(rosette-like) κII precipitates and at the globular and lamellar κIII precipitates. The fact that
the κ precipitates were not attacked, indicates that the κ precipitates are cathodic with respect
to the copper-rich α phase. Thus, corrosion of NAB alloy in quiescent seawater is galvanic in
nature.
SEM examinations revealed that the surface of the NAB specimen became slightly rough
after 3 hours of cavitation. After 13 hours of cavitation, the surface roughness increased and
several microcavities appeared in the attacked area (Fig. 5b). As the cavitation time increased
to 25 and 40 hours, the surface became very rough and contained large size cavities.
Cleaning the specimen that had been cavitated for 40 hours in acetone revealed ductile tearing
and grain boundary attack. In the presence of CP, the number of microcavities increased, but
grain boundary attack was absent. Examination of a cross-section of a specimen cavitated for
58 hours revealed microcracks, 5-10 μm in length, emanating from the bottoms of cavities
(Fig. 5c). The microcracks appeared to initiate and propagate in the α phase, favoring sites
adjacent to κ phases.
(a)
318
(b)
Al-Hashem et al.
(c)
Figure 5. SEM micrographs of NAB after: (a) exposure to quiescent seawater for
48 hours, showing preferential attack of the α phase; (b) cavitation
testing for 13 hr, showing increased surface roughness; and (c)
cavitation testing for 58 hr, showing cavities and microcracks in crosssection
The presence of cavities and ductile tearing in cavitated NAB specimens is readily
explainable in terms of the known devastating effects of cavitation. On the other hand, the
presence of grain boundary attack indicates that electrochemical dissolution due to structural
heterogeneity indeed contributes to the surface damage. The presence of microcracks in the
bulk of the material is similar to that reported for the NAB vertical-type pumps and sleeve
valves that prematurely failed in Arabian Gulf seawater [1,2]. Although selective phase
corrosion may have played a role in causing the microcracks, the contribution of cavitation
stresses should not be discounted.
The DSS specimens exposed to quiescent seawater under free corrosion conditions were
found to be free of any corrosion attack. When SEM examinations were made on the DSS
cavitated in seawater under free corrosion conditions, damage was seen after 30 minutes of
cavitation. At this stage, a few small, shallow voids were seen inside the austenite islands
and at the phase boundaries. Slip lines were also observed in the ferrite matrix with
microvoids developing at emerging slip steps. Furthermore, some austenite islands partially
extruded and experienced metal loss (Fig. 6a).
(a)
(b)
Figure 6. SEM micrographs of DSS after cavitation testing in seawater. (a) after 30
minutes, void formation at emerging slip steps and partially extruded
austenite; and (b) after 90 minutes, total removal of an extruded
austenite and extension of damage to ferrite
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Seawater Corrosion
After 70 minutes of cavitation, the formation of voids along the slip steps and the
material removal from the extruded austenite became more extensive. Slip lines were also
seen within the ferrite matrix and cleavage-like facets within the damaged austenite. After 90
minutes of cavitation, the extruded austenite was totally removed and the loss in material
extended to the ferrite matrix (Fig. 6b). As the cavitation time was increased to 4 hours, the
whole metal surface suffered from extensive damage and the formation of cavities. At this
stage, ductile tearing, cleavage-like facets, river patterns and crystallographic steps were seen
on the damaged surface. When the examinations were repeated on a specimen cavitated for
40 hours under CP, the surface roughness appeared to be somewhat less than under free
corrosion conditions. Laboratory-etched cross-sections of specimens cavitated for 40 hours
revealed microcracks in the bulk of the alloy. The cracks appeared to initiate in the ferrite
matrix. The propagation of the cracks appeared to be impeded by the austenite islands and to
branch along parallel slip systems. In some cases, however, cracks were seen extending into
the austenite grains.
The high ductility of the austenite could explain the extrusion of surface austenite. The
presence of cleavage facets within the damaged austenite suggests the occurrence of a brittle
mode of failure in addition to the ductile mode. This was rather surprising. The brittle mode
of failure in the austenite may be due to a strain-induced martensitic transformation [14].
The localized nature of cavitation seems to have led to an increase in work hardening of
the ferrite matrix, activation of slip systems and initiation of surface cracks. Thus, the earliest
deformation caused by the slip is expected to occur on the {110} plane and in the [111]
direction. The localization of plastic strain would also facilitate material removal by ductile
shearing of surface asperities. The repetitive stress mode of cavitation causes the cracks to
propagate laterally, giving rise to the observed cleavage-like facets. The cracks in the bulk of
the material express the fact that the plastic shock waves produced by cavitation can travel
deep into the material. The role of the austenite islands in impeding crack propagation has
been previously reported in stress corrosion cracking [6] and hydrogen embrittlement studies
[15]. It is feasible that passage of the propagating crack through an austenite obstacle is
facilitated by the formation of ε-martensite.
CONCLUSIONS
1. Cavitation caused an active shift in the free corrosion potential and increased the
anodic polarization current of NCI, NAB and DSS.
2. The rates of mass loss of the three cast alloys were very low in quiescent seawater.
When compared with quiescent conditions, cavitation increased the rates of mass loss
at the steady state by 930, 186 and 20 times for NCI, NAB and DSS, respectively.
The application of CP decreased the rates of mass loss by about 50%, 47% and 19%,
respectively. This reduction was attributed to bubble collapse cushioning due to
cathodic gas and diminished electrochemical dissolution.
3. In quiescent seawater, thin scales were formed on NCI, while NAB suffered from
selective corrosion of the copper-rich α phase at the interfaces with the intermetallic κ
precipitates. DSS was almost free of corrosion.
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Al-Hashem et al.
4. Cavitation made the surfaces of the alloys very rough, with large size cavities and
ductile tearing. The damage started at the graphite nodules and ferrite matrix in NCI.
In NAB, the damage started at the boundaries of α columnar grains. In the case of
DSS, slip lines formed in the ferrite matrix, austenite islands extruded, and voids
developed at emerging slip steps. Cleavage facets were evident within the damaged
austenite. The extrusion of the austenite was attributed to its high ductility, while the
presence of cleavage facets was explained in terms of strain-induced martensite.
Cross-sectional examinations revealed microcracks in the bulk of the cast alloys. The
formation of microcracks was attributed to cavitation stresses and preferential phase
corrosion.
ACKNOWLEDGMENT
The authors of this paper would like to acknowledge the support given by the Kuwait
Foundation for the Advancement of Science, the Kuwait National Petroleum Company, the
Petrochemical Industries Company, and the Abu Dhabi National Oil Company (Project No.
86-08-02).
REFERENCES
1.
2.
3.
4.
H.M. Shalaby and J.K. Cheriyan, KISR Report No. 2799, September 1988.
M. Islam, R. Abdul Wahab and S. Al-Kharraz, KISR Report No. 2913, January 1989.
V.K. Gouda, H.M. Shalaby and W.T. Riad, Materials Performance 28, 8, 1989, p. 53.
H.M. Shalaby, A. Al-Hashem, H. Al-Mazeedi and A. Abdullah, British Corrosion
Journal 30, 1, 1995, p. 63.
5. J.G. Auret, O.F.R.A. Damm, G.J. Wright and F.P.A. Robinson, Corrosion 49, 11,
1993, p. 910.
6. C.M. Preece, Cavitation erosion, in Treatise on Material Science and Technology,
Vol. 16, C. C.M. Preece, ed. (New York, NY: Academic Press, 1979), p. 249.
7. A. Al-Hashem, P.G. Caceres, W.T. Riad and H.M. Shalaby, Corrosion 51, 5, 1995, p.
331.
8. F. Hasan, J. Iqbal and N Ridley, Materials Science Technology 1, 1985, p. 312.
9. ASTM G32-92, Standard Test Method for Cavitation Erosion Using Vibratory
Apparatus (Philadelphia, Pennsylvania: ASTM, 1992).
10. H.M. Shalaby, S. Attari, W.T. Riad and V.K. Gouda, Corrosion 48, 3, 1992, p. 206.
11. K.R. Trethewey, T.J. Haley and C.C. Clark, British Corrosion Journal 23, 1, 1988, p.
55.
12. A. Schussler and H.E. Exner, Corrosion Science 34, 11, 1993, p. 1803.
13. T. Okada, Y. Iwai and A. Yamamoto, Wear 84, 1983, p. 297.
14. J.A. Venables, Philosophical Magazine 7 1962, p. 35.
15. W. Zheng and D. Hardie, Corrosion 47, 10, 1991, p. 792.
321
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
MICROBIOLOGICALLY INDUCED CORROSION
OF A STAINLESS STEEL PIPE
H.H. Lee, M. Ali and K. Al-Omrani
Sabic Industrial Complex for Research & Development
PO Box 42503, Riyadh 11551, Saudi Arabia
ABSTRACT
Microbiologically induced corrosion (MIC) is recognized as a major cause of metallic
component failure in many natural waters and soils. Most recently, a MIC failure occurred in the
stainless steel piping system in a chemical plant in Saudi Arabia. This failure case was discussed and
compared with the documented case histories in the published literature involving MIC of stainless
steels.
Key Words: Microbiologically induced corrosion (MIC), stainless steel, localized corrosion
INTRODUCTION
In late July of 1995, leakage was detected in the newly installed stainless steel (Grade
321 stainless) pipe in a chemical plant in Saudi Arabia. The leakage was observed in the
stainless steel core pipe which was enveloped by an outer jacket pipe for heat transfer
medium oil. Evidence of pinholes was found in the lowest sloping, horizontal section of the
core pipe, always around the 7 o'clock to 5 o'clock positions and about 0-8 mm from the weld.
The remainder of the pipes at a higher elevation have been found to be acceptable with no
leakage.
The original hydrotest using untreated fresh water was conducted for the newly
constructed pipe line in January of 1995, followed by air blowing to remove excess water.
This line was turned to the operating personnel of the chemical plant in March of 1995.
System hydrostatic leak testing was conducted by the plant personnel in April of 1995.
Pinhole leakage was observed in late July of 1995. Between April of 1995 and the leakage
was detection, the pipeline was never in use. In August of 1995, an investigation was initiated
on the cause of this failure. The investigation and analysis results are reported in this paper.
VISUAL AND MICRO EXAMINATION
One small pipe sample about 4 cm in length was cut from the failed stainless steel
pipeline for examination and evaluation. The sample included the weldment and pinholes.
This sample was cut along the longitudinal direction to expose the inner pipe surface for
examination. The measured outside diameter (OD) of the stainless steel pipe sample in asreceived condition was 88 mm with 3 mm in wall thickness. There was no observed
corrosion on the outer surface, as shown in the photograph in Fig. 1a. However, there were
reddish-brown corrosion products mixed with some light black oxides at and near the welded
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Seawater Corrosion
zone of the inner surface, especially at the 5 to 7 o'clock position. One pit is clearly visible at
approximately the 6 o'clock position adjacent to the weldment, as shown in the photograph in
Fig. 1b. The area including the pit was cut for further examination under scanning electron
microscope (SEM). Figure 2a is a SEM photograph of the pit clearly indicating that the pit
was initiated from the inner surface at area near the weldment. Figure 2b is a SEM
photograph of the front view of
(a)
(b)
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Lee et al.
Figure 1. Photographs of: (a) outer surface; (b) inner surface of the failed pipe sample,
indicating a pit growing nd propagating sideways from the inner surface adjacent to
the weldment
the pit indicating that the pit was turning sideways under the inner pipe surface. Figure 3 is a
SEM photograph of the side view of the pit looking from the cross-sectional direction of the
pipe showing a large under-surface cavity. This pit under examination did not penetrate the
whole pipe wall thickness. It clearly indicates that the pit was generated due to corrosion and
not from material defect.
ANALYSIS OF THE FAILURE MECHANISM
Based on the laboratory examination result, there are two possible environments which
could produce the type of pitting corrosion of stainless steel pipe as observed in the present
case. These include environment containing appreciable concentrations of chloride (Cl-), and
environment containing bacteria which could microbiologically induce corrosion of stainless
steels. These possibilities are further discussed in reference to the present failure case:
Pitting Corrosion of Stainless Steel in Environment Containing Chloride
It is known that stainless steel will develop pitting corrosion in environments containing
chloride ions because of the breakdown of surface passivity [1]. Generally, this requires an
appreciable concentration of chloride ions. For examples, stainless steels exposed to seawater
will develop surface pits in a matter of months to several years depending on the steel’s
composition. The tendency is greater in the martensitic and ferritic stainless steels than in the
austenitic steels. Austenitic stainless steels containing molybdenum (such as types 316 and
317) are more resistant to pitting corrosion in seawater. Pitting of these alloys will develop in
seawater within a period of 1-2.5 years [1]. The present failed stainless steel pipe was never
exposed to an aqueous environment containing appreciable amount of chloride ions prior to
the detection of the leakage. The type of pitting corrosion observed at this chemical plant is
not caused by chloride ions from the environment.
Localized Biological Corrosion of Stainless Steels
Biological corrosion is defined as the deterioration of metals as a result of the metabolic
activities of microorganisms [2]. Many case histories involving MIC of stainless steels are
well documented in the published literature [2]. Most of them involved sulfate-reducing
bacteria which accelerate the localized corrosion of stainless steels especially at welded areas.
One case history, as cited by Kobrin [3], is considered most relevant. It occurred in the early
1980s in a new chemical plant in Texas, USA.
Type 304 stainless was used for resistance to corrosion by wet carbon dioxide.
Included were underground sump tanks and drainage piping. Untreated well water
was used for hydrostatic testing after installation. No special effort was made to
remove the water or dry the systems after testing was complete. Initial indication of
a problem occurred several months after the hydrotest, following removal of a tank
to a fabrication shop for minor modifications. Leaks developed during a routine
shop hydrostatic test. Internal visual and liquid dye penetrant inspection showed pits
and cracks under rusting colored nodular deposits primarily along welds. Pits had
small 'mouth' at the surface, opening to large, bottle-shaped cavities below. As plant
personnel opened the stainless vessels and the pipes for inspection, they found many
325
Seawater Corrosion
pits, primarily at welds, and noted rust-colored deposits in mixture with brownishwhite slime deposits. Limited analyses showed the presence of sulfate reducing and
iron bacteria in the various deposits [4].
(a)
(b)
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Lee et al.
Figure 2. SEM photographs of: (a) the pit adjacent to the weldment (7X); and (b) the pit as
seen from the inner pipe’s surface (25X)
Figure 3. SEM photograph of the pit hole from the cross-sectional direction indicating the pit
growing sideways under the surface (25X)
The above reported case history is very similar to what has been observed at the
chemical plant in Saudi Arabia for the present pipeline pitting failure. It appears that some
water still remained in the lowest sloping horizontal section of the core pipe at the chemical
plant after the initial hydrotest conducted in January of 1995. Obviously the untreated water
used for the initial hydrotesting contained some sulfate reducing bacteria which caused the
subsequent biologically induced corrosion.
As showed in the SEM photograph in Fig. 3, some rod-shaped material is visible inside
the pit. It is possible that the rod-shaped material could be a colony of the corroding bacterial
consortium. According to the literature, there are many bacteria found in the corroded metals
in the form of filaments. One example is the filament-form thermophillic bacterium found in
the failed brazed fillet reported by Walch and Mitchell [5]. This filament form of organism
has variable lengths from 20 to over 200 μm.
Although Sabic R&D does not have the facilities to positively identify the type of
organism observed, it is fairly clear that the present pipeline failure observed at the chemical
plant was caused by biologically induced corrosion. Furthermore, the shape of the pit, i.e.,
with a small mouth on the surface and a large bottle-shaped cavity below, is a positive
indication of bacterial corrosion.
CONCLUSIONS
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Seawater Corrosion
Based on the examination results and the evidence obtained, the cause of the pitting
corrosion in the stainless steel piping system observed was caused by biologically induced
corrosion. After the piping system was repaired or replaced as required, it was recommended
that hydrotesting be conducted using demineralized water to eliminate the possibility of
microbiologically induced corrosion. The piping system was thoroughly drained and dried
immediately after testing. There was no reoccurrence of MIC since.
REFERENCES
1. H.H. Uhlig and R.W. Revie, Corrosion and Corrosion Control, John Wiley & Sons, 1985.
2. Metals Handbook, Ninth Edition, Vol. 13, Corrosion, p. 2, ASM International, 1987.
3. G. Kobrin, Reflections on microbiologically induced corrosion of stainless steels, in:
Biologically Induced Corrosion, Conference Proceedings, National Association of
Corrosion Engineers, 1986, pp. 33-46.
4. ibid., p. 37.
5. M. Walch and R. Mitchell, Microbial influence on hydrogen uptake by metals, in:
Biologically Induced Corrosion, Conference Proceedings, National Association of
Corrosion Engineers, 1986, pp. 201-214.
328
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
A LABORATORY STUDY OF SERVICE FAILURE OF AL-BRASS
TUBES IN ARABIAN GULF SEAWATER
H.M. Shalaby1, W.T. Riad1 and V.K. Gouda2
1
2
Kuwait Institute for Scientific Research
P.O. Box 24885, 13109 Safat, Kuwait
National Research Center, Dokki, Giza, Egypt
ABSTRACT
Failure analysis and a laboratory study were made on Al-brass tubes. The failure investigation
was conducted on tubes that had failed during service in Arabian Gulf seawater. The investigation
revealed that the tubes failed due to perforations after suffering from various forms of corrosion.
Crevice attack and erosion-corrosion of the horseshoe grooving type were the most serious corrosion
forms. Other corrosion forms were also observed such as inlet edge cutting, inlet edge beveling,
pitting, dezincification and intergranular attack. Lodgments, local high velocities and turbulence were
considered the most important factors causing the perforations.
In order to understand the role played by the presence of seawater pollutants, a systematic
laboratory study was conducted to evaluate the effects of manganese (5 ppm) and chlorine (4 ppm) on
the corrosion behavior of Al-brass in stagnant and flowing (0.1 and 2.2 m/s) seawater. Under stagnant
conditions, both pollutants had little effect on the free corrosion potential. In flowing seawater,
manganese caused a noble shift in the corrosion potential at 0.1 m/s and an active shift at 2.2 m/s.
Chlorine addition caused an electropositive shift at both flow velocities. During polarization, the
presence of either pollutant caused cathodic depolarization and elimination of the active/passive
transition. Microscopic examination revealed that the film formed in the presence of manganese was
thicker and more porous than that formed in clean seawater. Scattered nodules of basic copper
chlorides formed on the tube’s surface in the presence of chlorine. As in the field failed tubes, crevice
corrosion, intergranular attack and dealloying were found beneath the basic chloride nodules.
Key Words: Aluminum brass, corrosion failure, manganese, chlorine, pollutants
INTRODUCTION
In Kuwait, refineries and desalination plants utilize Arabian Gulf seawater in heat
exchangers, evaporators and condensers fitted with Al-brass tubes. Although the alloy is
recommended for the heat recovery stages in multistage flash distillation (MSF) plants [1], a
high incidence of tube failures have been reported [2-6].
The diagnosis of failure of Al-brass tubes in surface condensers, conducted by Gouda et
al. [4], indicated that the tubes failed due to severe erosion-corrosion in the form of
horseshoe-shaped grooves. On the other hand, Khatak and Gnanamoorthy [5] attributed leaks
developed at random locations in Al-brass condenser tubes after one year of operation to
crevice corrosion under seawater deposits. Furthermore, Chandrasekhariah and Mukherjee [6]
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Seawater Corrosion
attributed the failure of Al-brass condenser tubes after six months of service to crevice
corrosion resulting from lodgments of organic deposits.
The aim of this paper is to establish a link between a failure investigation of Al-brass
tubes and a laboratory study on the effect of seawater pollutants. It was hopped that
establishing such a link would throw more light on the mechanism of failure.
EXPERIMENTAL PROCEDURE
Failure analysis was undertaken on ten failed Al-brass tubes made according to ASTM
C68700. The tubes were removed from the heat recovery stage of a desalination plant that
used treated seawater. The design flow velocities were 1.5-2.14 m/s, and the temperatures
were 42°C at the inlet stage and 82.6°C at the outlet stage. Visual, macroscopic,
metallographic and scanning electron microscopic (SEM) examinations were conducted on
the tubes before and after the removal of corrosion products. X-ray diffraction (XRD) and
energy dispersive spectroscopy (EDS) were used for the analysis of the corrosion products
and deposits. Also, chemical analyses were made to verify the material’s conformity.
The laboratory corrosion tests were conducted under stagnant and flow conditions (0.1
and 2.2 m/s) using Arabian Gulf seawater at room temperature. The tests were made on fresh
(unused) Al-brass using clean seawater and seawater containing 5 ppm manganese or 4 ppm
chlorine. Chlorine was added in the form of sodium hypochlorite solution, and the
concentration of free chlorine was monitored using an ion selective analyzer and was
replenished daily. Manganese was added in the form of manganese chloride. The work
under stagnant conditions involved measurements of the open-circuit corrosion potential and
potentiodynamic polarization tests. The tests were made in duplicate using specimens in the
form of discs measuring 2 cm in diameter and 1 mm in thickness. Before the test, specimens
were mechanically ground using 800 grit silicon carbide paper, polished with 3 μm diamond
paste, and ultrasonically cleaned with detergent and acetone. Teflon holders was designed to
expose a specimen area of 1.6 cm2. The potentiodynamic polarization tests were carried out
at a scan rate of 30 mV/min starting from a cathodic potential value of -700 mV relative to the
open-circuit potential and scanning was stopped at a potential of + 700 mV. A potentioscan
with IR compensation controls, saturated calomel reference electrodes (SCE) and cylindrical
graphite counter electrodes were used in these experiments.
A circulating test rig in the form of a flow loop made of PVC pipes was used to carry out
the work under flow conditions. The rig was designed in such a manner that the Al-brass
tubes (10 cm in length, 1 mm in thickness and 1.6 cm in outer diameter) were exposed to
laminar flow at the same flow velocity. The tubes were internally ground with 600 grit
silicon carbide paper and soldered on the outside to copper wires to record the corrosion
potential. In this paper, average values of the open-circuit potential are reported.
At the end of the stagnant and flowing laboratory tests, the specimens were examined for
corrosion morphology and the composition of corrosion products. Again, SEM, EDS and
XRD techniques were used. The corrosion morphology was examined before and after
removing the corrosion products. The corrosion products were removed by immersion in 5%
citric acid for 24 hours, followed by ultrasonic cleaning in acetone.
RESULTS
Failure Investigation
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Shalaby et al.
Visual observation showed that the internal surfaces of the failed tubes were covered by
pale green and orange, nonhomogeneous deposits mainly located at the mouth of the tubes
and occasionally in the middle of the tubes. The orange product was generally more adherent
than the green deposits. XRD and EDS analysis of the corrosion products indicated the
presence of a nonhomogeneous mixture of cuprous oxide and copper hydroxychloride
(paratachamite), in addition to sea salts such as calcium carbonate and magnesium hydroxide.
Inspection of the tubes after removal of the scales and corrosion products revealed
various localized forms of attack at the tube’s mouth, in addition to metal thinning
downstream. In most of the tubes, the inlet edges suffered from inlet-beveling attack and
inlet edge cutting (Fig. 1), while at a distance of 3 cm from the inlet edges, localized craters
of various depths, shapes, sizes and angles (45°-90° to the water flow direction) were located.
Perforations were found in some craters. Also, in some tubes containing horseshoe-shaped
grooves, perforations were apparent on the metal surface underneath the orange colored
product. Another form of localized corrosion (minute pits) was identified in many of the
tubes under scattered, localized black deposits. Inspection was also conducted on unused Albrass tubes which showed the initiation of minute pits covered by a black corrosion product.
Metallographic cross-sections were prepared for areas that suffered from localized
corrosion and craters. Examinations of the cross sections revealed the occurrence of
dezincified and nonuniform porous surfaces as well as shallow pits in the α-brass structure.
Examinations of polished longitudinal sections at the craters revealed the presence of
intergranular attack at the crater’s borders. Transverse sections at the same location did not
reveal the penetration of intergranular cracking.
Laboratory Corrosion Tests in Stagnant and Flowing Seawater
As can be seen from Figs. 2 and 3, the potential-time behavior of Al-brass exposed to
pollutant-free seawater under stagnant and 0.1 m/s flow conditions was about the same. The
initial average potentials showed slight ennoblement during the first few days of testing,
followed by a small active shift until a steady-state was reached. When the flow velocity was
increased to 2.2 m/s, the corrosion potential in clean seawater gradually shifted to more noble
values (Fig. 4).
Visual examination of the Al-brass tested in stagnant clean seawater revealed a surface
that was mostly yellow-brown in color and contained a few small green precipitates which
increased in number and became darker in color on tubes exposed to seawater flowing at 0.1
m/s. On these latter tubes, beige and white precipitates were also observed. SEM
examination showed that the green precipitates were granular in nature, well compacted and
had formed on top of a thin, smooth, and adherent oxide film. EDS analysis indicated that the
green precipitates were rich in chloride and sulphur. After film removal with 5% citric acid,
the surface was free of any significant corrosion damage.
The tubes tested in clean seawater flowing at 2.2 m/s appeared to be similar to those
tested at 0.1 m/s. However, SEM examination showed that the surface film on the tubes
tested at 2.2 m/s had been stripped off from several areas, exposing the underlying bare metal.
Higher SEM magnification revealed that the bare surfaces had suffered from metal
dissolution that increased in intensity along the grain boundaries. A few scattered, small
shallow cavities were seen on the intergranularly attacked surface.
331
Seawater Corrosion
The potentiodynamic polarization of Al-brass in stagnant clean seawater is given in Fig.
5. A small active/passive transition peak is apparent ,and the active peak is followed by a
very small passive region and a small transpassive peak. The current slowly increased with
further increase in potential. Microscopic examination of the specimen showed general metal
dissolution associated with severe intergranular attack.
(a)
(b)
(c)
(d)
Figure 1. Photographs of forms of corrosion attack observed on the failed tube inlets: (a)
holes, craters and depression (1.1x); (b) horseshoe grooving and hole (1.8x); (c)
edge cutting and bevelling (8x); and (d) crater with elongated hole (10x)
The addition of 5 ppm manganese to stagnant seawater did not change the potential-time
behavior of Al-brass with respect to that in clean seawater (Fig. 2). On the other hand, the
presence of manganese increased the corrosion potential by 100 mV during potentiodynamic
332
Shalaby et al.
polarization in stagnant seawater (Fig. 5). It also shifted the cathodic branch of the
polarization curve towards higher current densities, eliminated the active/passive peak and
made the anodic peak to take place at slightly less positive potentials, suggesting selective
dissolution from the alloy. Optical microscopic examination of the sample used in the
polarization test revealed once again that general metal dissolution and intergranular attack
had occurred.
Figure 2. Potential-time behavior of Al-brass in stagnant seawater: (a) seawater; (b) seawater
containing 5 ppm manganese; and (c) seawater containing 4 ppm chlorine
Figure 3. Potential-time behavior of Al-Brass in seawater flowing at 0.1 m/s: (a) seawater;
(b) seawater containing 5 ppm manganese; and (c) seawater containing 4 ppm
chlorine
333
Seawater Corrosion
Figure 4. Potential-time behavior of Al-Brass in seawater flowing at 2.2 m/s: (a) seawater;
(b) seawater containing 5 ppm manganese; and (c) seawater containing 4 ppm
chlorine
Figure 5. Potentiodynamic polarization behavior of Al-brass in stagnant seawater: (a) clean
seawater; (b) seawater containing 5 ppm manganese; and (c) seawater containing 4
ppm chlorine
In seawater flowing at 0.1 m/s, the presence of manganese shifted the corrosion potential
to a peak value of about -110 mV during the first two days of testing. Then, the potential
gradually shifted towards a more active value, reaching a more or less steady state of about 210 mV after eight days of testing (Fig. 3). At 2.2 m/s, a different potential-time behaviour
was attained in the presence of manganese. In this latter case, the corrosion potential
remained much more active, being on the same order as that obtained within the steady-state
region under stagnant conditions.
The surface film formed in manganese-containing seawater flowing at 0.1 m/s was
morphologically similar to that formed in clean seawater under stagnant and flow conditions.
However, in the presence of manganese, the film was thicker and covered most of the
extrusion marks. At 2.2 m/s, the film formed in the presence of manganese became much
thicker and more porous when compared with the film formed in clean seawater at the same
flow velocity. Furthermore, the film stripped off the surface in many areas. Within the
stripped-off areas, the surface was intergranularly attacked and contained a few small,
shallow cavities (Fig. 6a).
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Shalaby et al.
The effect of the addition of 4 ppm chlorine on the potential-time behavior of Al-brass is
shown in Figs. 2-4. In stagnant seawater, the potential underwent an initial ennoblement
during the first few days of testing which was followed by a shift towards active values till a
steady state was attained after about ten days of testing (Fig. 2). At the steady state, the
corrosion potential in the presence of chlorine was close to that in clean stagnant seawater,
indicating a negligible effect of chlorine. Statistical hypothesis testing also confirmed that
there was no significant differences at the l and 5% levels of significance.
In seawater flowing at 0.1 m/s, the addition of chlorine shifted the corrosion potential to
a noble value of about -100 mV after two days of exposure, and the potential remained close
to that value for the remaining period of testing (Fig. 3). At a 2.2 m/s flow velocity, the
potential-time behavior in the presence of chlorine was to a good extent similar to that
obtained in clean seawater at the same flow velocity (Fig. 4). Although potential fluctuations
occurred in both cases, the potential values were close to each other and remained relatively
stable after about ten days of testing.
SEM examination of the Al-brass tested under open-circuit conditions in chlorinecontaining stagnant seawater revealed features similar to those noted in clean seawater. On
the other hand, the surface of tubes tested in chlorine-containing seawater flowing at 0.1 m/s
was slightly greenish in color and contained a few green nodules. EDS of the nodules
revealed that they were enriched with chloride and sulphur. The surface appeared (under
SEM) to be covered by an adherent film, however, a few areas seemed to be covered by a
smooth, wavy scale layer. After removing the surface film, the bare surface was found to
have suffered from minor general corrosion associated with increased surface roughness in
the form of micro pits.
When the seawater flow velocity was increased to 2.2 m/s, the film formed in chlorinecontaining seawater became a thick, scale layer containing scattered mounds or nodules (Fig.
6b). In some areas, the scale looked like a layer of dried, broken mud, and in others it was
multilayered. The scale was not stripped off as in the case of clean seawater and seawatercontaining manganese flowing at the same velocity (i.e., 2.2 m/s), which suggests that the
scale was adherent to the surface. XRD analysis of the corrosion products scrubbed off the
surfaces of tubes after testing showed major peaks of copper hydroxychloride in addition to
cuprous chloride. After removing the scale, SEM examination showed that the bare surface
contained a significant number of relatively large isolated and interconnected cavities in pitlike shapes (Fig. 6c). Inside and around the cavities, the bare surface was porous (spongy) in
nature. However, away from the cavities, the surface appeared to be quite similar to that
observed after testing in clean seawater or in the presence of manganese.
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Seawater Corrosion
(a)
(b)
(c)
Figure 6. SEM micrographs of Al-brass tested in seawater at 2.2 m/s, showing: (a) the
surface after testing in the presence of manganese; (b) nodules on top of the scale
formed in the presence of chlorine; and c) cavities observed after removing the
scale
The addition of chlorine to stagnant seawater shifted the cathodic branch of the
polarization diagram towards higher current densities (Fig. 5). It also caused the formation of
a small reduction peak. The addition of chlorine appears to diminish the passive region and
make the anodic peak take place at a slightly less positive potential than in clean seawater.
After the termination of the polarization test, microscopic examination revealed the same
featues as those observed in clean seawater.
DISCUSSION
Service Failure of Al-brass
It has been reported that the good performance of copper alloy tubing in desalination
plants is due to the formation of a protective film composed of cuprous oxide [1]. Other
deposits on top of the film, such as cupric oxychloride, cupric hydroxide, copper carbonates
and iron oxides, may also contribute to the corrosion protection. In fact, the corrosion
products on the surface of the field-failed tubes were a nonhomogeneous mixture of cuprous
oxide, copper hydroxychloride and paratachamite, in addition to sea salts such as calcium
carbonate and magnesium hydroxide. The nature of these scales appeared to be inadequate to
protect Al-brass in Arabian Gulf seawater.
Visual inspection of the failed tubes showed that inlet edge cutting and bevelling had
occurred in most of the tubes, while typical horseshoe grooving was evident in some tubes.
All these forms of attack were indicative of flow velocity related problems at the tube inlets
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Shalaby et al.
only, although the flow velocity inside the tubes was within the safe range specified for Albrass (2.4 m/second) [1].
The design of the water-box inlet and baffles may create local turbulent flow patterns
that exceed the maximum allowable flow velocity to avoid erosion-corrosion. Sato and
Nagata [7] reported that the water velocity passing a partial obstruction in the bore of a
condenser tube can reach 8 m/second even though the overall velocity is in the range of 2
m/second. The existence of lodgments such as shells, biofouling deposits or corrosion scales
may also create areas of localized high velocity and turbulence leading to localized corrosion
and under-deposit attack. In the present failure case, crevice corrosion was observed under
lodgments, and the surfaces of the creviced areas were porous in nature, suggesting the
occurrence of dealloying. Similar results were noted by Efird and Verink [8], and Schleithoff
and Schmitz [9] where the type of attack in the crevices was dealloying. Schleithoff and
Schmitz [9] indicated that the presence of magnesium in the alloy was responsible for the lose
of arsenic which was added to prevent dezincification. On the other hand, it has been
reported [10] that in the crevice attack of copper alloys exposed to flowing seawater, cuprous
ions are produced within the crevice and react with chloride ions to form ionic copper
chloride complexes, leading to local acidification and enhancement of corrosion.
Examinations of metallographically polished sections from areas that had suffered from
crevice attack revealed the presence of intergranular corrosion and cracking. Todd [11]
reported failure of Al-brass tubes exposed to polluted seawater in a desalination plant. The
failure of the tubes was due to severe intergranular corrosion. In the present failure case,
minute pits were also observed under scattered and localized black sulphide deposits. Similar
attack was reported for Al-brass tubes in another application utilizing Arabian Gulf seawater
[4]. It has been reported that a copper sulphide deposit formed in polluted seawater is more
cathodic than cuprous oxide formed in clean seawater. Therefore, galvanic corrosion
between the base metal at breaks in the large areas of cathodic sulphide film can result in
rapid failure by pitting attack [12,13].
Effect of Seawater Pollutants
It is well established that Al-brass tubes perform satisfactorily if the seawater is clean,
i.e., not contaminated with any pollutants. Previous work [14] has shown that the major
constituent of the film formed in clean seawater is of the hydrotalcite family of basic mixed
hydroxy-carbonates [Mg6Al2(OH)16CO3.4H2O]. In the presence of an intermediate chlorine
level, the composition of the normal surface film changes. Francis [14] suggested that this
change is achieved by substituting increasing amounts of copper for magnesium and chloride
ion for the hydroxyl ion, resulting in a film with low resistance to mechanical damage.
However, there are confliting results on the effect of chlorine. Studies under stagnant or low
flow velocities have often resulted in reduced corrosion rates in the presence of chlorine
[15,16], while other work at higher velocities [17,7], has shown dramatically increased
corrosion.
Sato et al. [18] reported catastrophic damage to Al-brass tubes in chlorinated seawater,
caused by malignant impingement attack. They explained this phenomenon on the basis that
manganese ions present in seawater react with the residual chlorine from chlorination, and
thus MnO2 is formed as colloidal particles depositing on the corrosion product film and
forming an active cathode in areas of turbulence around lodgments. Consequently, galvanic
337
Seawater Corrosion
action is enhanced, leading to intensified erosion-corrosion damage. Gouda et al. [19] also
considered that the presence of significant amounts of manganese on the surface of failed Albrass condenser tubes used in a refinery in Kuwait as one of the factors contributing to
accelerated erosion-corrosion of the tubes. In all these failure investigations, the exact role of
manganese was not studied or clearly defined.
The above discussion suggests that the present failure case cannot be attributed to the
wrong choice of material. However, when using Al-brass, several precautions have to be
taken to ensure the long life of the tubes. Pollutant-free water should be used for hydrotesting
tubing followed by thorough drying prior to boxing for shipping. The circumstances of the
present failure case and the confleting results with regard to the effect of pollutants made it
necessary to carry out a systematic laboratory study on the effects of manganese and chlorine
on the corrosion of Al-brass. It was hoped that such a study might throw more light on the
exact role of seawater pollutants in the failure of the Al-brass tubes. The effect of sulfide
pollution was not studied, as it is well covered in the literature [12,13].
In the present laboratory work, the addition of 5 ppm manganese caused an
electropositive shift in the free corrosion potential reached during testing in seawater flowing
at 0.1 m/s. It also caused a similar shift under potentiodynamic polarization in stagnant
seawater. These results can be explained on the basis of the mixed potential theory and
morphological examination. According to Pourbaix [20], a slightly oxidizing agent such as
that of oxygen, can oxidize manganous solutions with the formation of solid oxides: brownblack Mn3O4, black Mn2O3 or various varieties of anhydrous or hydrated MnO2 which are
brown or black. These oxides are more noble than Al-brass, and their presence on the surface
of the metal is similar to that of a galvanic couple or the addition of an oxidizer to the system.
Under such conditions, the corrosion potential of the system is shifted to more noble values,
the corrosion rate is increased, and the hydrogen evolution rate is decreased.
The fact that the free corrosion potential was of the same order in stagnant clean
seawater and in seawater containing manganese could possibly be due to the lack of sufficient
diffusion of manganese (or manganese oxides) to the Al-brass surface. On the other hand, the
activation of the corrosion potential noted at a 2.2 m/s flow velocity could be due to the
stripping of the protective film, exposing fresh metal to the corrosive environment. The
increase in thickness of the surface film as a result of the addition of manganese would
increase the corrosion rate of the metal. The increase in film thickness might help the
stripping action which might result from the shear stress between the surface and the layer of
seawater closest to it. The increase in film thickness and its associated increase in corrosion
rate could be a factor in the development of the observed intergranular attack. Thus, the
present results suggest that the presence of manganese in seawater affects the growth and
stability of the protective layer.
The present laboratory work revealed the strong oxidizing effect of chlorine. This was
apparent from the electropositive shift in the corrosion potentials (Figs. 3 and 4) and the
cathodic depolarization relative to that in clean seawater (Fig. 5). Furthermore, the
morphological examination showed that the surface film formed in Al-brass in the presence
of chlorine was different than that formed in clean seawater or in the presence of manganese.
However, another scale layer which was different than the primary film was observed in some
areas (Fig. 6b). XRD showed that the corrosion products scrubbed off the surface of Al-brass
338
Shalaby et al.
were basically copper oxychlorides and copper chlorides. It appears that the film formed on
Al-brass changes in the presence of free chlorine. The presence of a totally different scale
layer of basic copper compounds in the form of multiple layers or nodules in localized areas
could significantly change the corrosion rate. The basic chloride nodules could act as sites
for further precipitation of other compounds, such as calcium carbonate, leading to
nonhomogeneous mixtures of several compounds, as was observed in the present failure case.
The presence of these nodules increases the local acidity underneath the scale, leading to
crevice corrosion, intergranular attack and dealloying, as was observed in both the laboratory
work and the field-failed tubes. Furthermore, the presence of these nodules could induce
localized increases in velocity and turbulence, resulting in erosion-corrosion at high flow
velocities.
CONCLUSIONS
1. The failure investigation revealed that the Al-brass tubes failed due to various forms of
corrosion. Crevice attack and erosion-corrosion of the horseshoe grooving type were the
most serious forms. Inlet edge cutting, inlet edge beveling, pitting, dezincification and
intergranular attack were also observed. The failure of the tubes was ascribed to
lodgments, increased local high velocities and turbulence.
2. The laboratory study indicated that the addition of 5 ppm manganese or 4 ppm chlorine
had little effect on the free corrosion potential of Al-brass in stagnant seawater. In
flowing seawater, manganese shifted the corrosion potential to more noble values at 0.1
m/s and to more active values at 2.2 m/s. Chlorine addition caused electropositive shift at
both flow velocities.
3. During potentiodynamic polarization, the presence of either pollutant caused cathodic
depolarization, electropositive shift of the corrosion potential and elimination of the
active/passive transition.
4. SEM examination of the laboratory tested specimens revealed that the film formed in the
presence of manganese was relatively thick and porous. The addition of chlorine caused
the localized precipitation of a scattered scale layer in the form of nodules of basic copper
chlorides. Similar to the field-failed tubes, crevice corrosion, intergranular attack and
dealloying were found beneath the nodules.
REFERENCES
1. H. Tuthill, Materials Performance 26, 9, 1987, pp. 12-22.
2. A.D. Little, Survey of Service Behaviour of Large Evaporative Desalting Plants, US
Department of Commerce, Office of Water Research and Technology, 1981.
3. V.K. Gouda and W.T. Riad, Kuwait Institute for Scientific Research, Report No. 2767,
1988.
4. V.K. Gouda, J.A. Carew and J.K. Cheriyan, Kuwait Institute for Scientific Research,
Report No. 2605, 1988.
5. S. Khatak and J.B. Gnanamoorthy, Failure of an Aluminium Brass Condenser Tube:
Handbook of Case Histories in Failure Analysis, Vol. 2, ASM International, Materials
Park, Ohio, USA, 1993, pp. 192-193.
339
Seawater Corrosion
6. M.N. Chandrasekhariah and T.K. Mukherjee, Transactions of the Indian Institute of
Metals 31, 1978, p. 459.
7. S. Sato, and K. Nagata, Sumitomo Light Metal Technical Report No. 19, Japan, 1983, p.
296.
8. K.D. Efird, and E.D. Verink, Jr., Corrosion 33, 9, 1977, pp. 328-331.
9. K. Schliethoff and F. Schmitz, Practical Metallography 20, 2, 1983, pp. 88-91.
10. M. Schumacher, ed., Sea Water Corrosion Handbook,.Park Ridge, New Jersey, Noyes
Data Corporation, 1979.
11. B. Todd, Desalination 3, 106, 1967.
12. H.P. Hack and J.P. Gudes, Inhibition of sulfide-induced corrosion with a stimulated iron
anode, Corrosion/79, NACE, Atlanta, Georgia, USA, March 12-16, 1977, p. 10.
13. D.D. Macdonald, B.C. Syrett and S.S. Wing, Corrosion 35, 1979, p. 367.
14. R. Francis, Corrosion Science 26, 3, 1986, p. 205.
15. D.B. Anderson and B.R. Richards, Journal of Engineering for Power, July 1966.
16. I. Campbell and N.K. Searle, Conference on Fouling and Corrosion of Metals in Sea
Water, SMBA Dunstaffnage, April 1982.
17. R. Francis, Materials Performance 44, 2, 1982.
18. S. Sato, K. Nagata and S. Yamauchi, Evaluation of various preventive measures against
corrosion of copper alloy condenser tubes by seawater, Presented at Corrosion/81, NACE,
Toronto, Canada, April 6-10, 1981.
19. V.K. Gouda, S. Abo-Namous, W.T. Riad and A.M. Abdullah, Kuwait Institute for
Scientific Research, Report No. 3266, 1989.
20. M. Pourbaix, Atlas of Electrochemical Equilibria in Aqueous Solutions, 2nd ed., National
Association of Corrosion Engineers, 1974.
340
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
CORROSION OF REINFORCED CONCRETE STRUCTURES AND
THE EFFECTS OF THE SERVICE ENVIRONMENT
S. Al-Bahar and E.K. Attiogbe
Civil Engineering &Building Department
Kuwait Institute for Scientific Research
P.O. Box 24885, 13109 Safat, Kuwait
ABSTRACT
Corrosion of steel reinforcement is a major cause of concrete deterioration in Kuwait and other
Arabian Gulf countries. The selection of effective repair schemes for corroding structures requires
that the corrosion damage be assessed and quantified.
The symptoms of corrosion-induced deterioration in two reinforced concrete structures in Kuwait
are documented and discussed. The investigative techniques employed identified the nature of the
service environment and its role in promoting corrosion of steel reinforcement in the concrete
structures. The failure to account for the service environment of the structures through
implementation of corrosion preventive measures contributed to the deterioration of the structures.
Both ingress of chloride ions and carbonation of the concrete were the causes of the corrosion-induced
deterioration.
Key Words:
Carbonation, chlorides, concrete structures, reinforced concrete,
reinforcement corrosion, service environment
INTRODUCTION
The service environment plays an important role in the durability and serviceability of
reinforced concrete structures. The Arabian Gulf region provides one of the most aggressive
environmental conditions for concrete structures. This environment is characterized by high
temperature and humidity cycles, and severe ground and ambient salinity with high levels of
chlorides and sulfates in the soil and groundwater. Such environmental conditions have
promoted an extensive degree of deterioration in concrete structures within 10-15 years of
construction [1]. Structures exposed to the marine environment, groundwater conditions and
industrial pollution have suffered the most.
In Kuwait, as in other Arabian Gulf countries, the symptoms of deterioration in
reinforced concrete structures have shown that corrosion of steel reinforcement is the major
form of deterioration [2]. Studies conducted on the conditions of marine and non-marine
structures in Kuwait have shown corrosion to be caused primarily by ingress of chloride ions
aided by carbonation and sulfate attack, while carbonation of concrete is a major cause of
corrosion in some non-marine structures. Cracking of concrete caused by sulfate attack can
facilitate the ingress of chloride ions. A visual inspection of selected structures showed
evidence of corrosion to be extensive (Fig. 1) and widely spread [3]. The observations made
341
Corrosion in the Building Industry
during the visual inspection of these structures point to the vital role the service environment
played in the extent of corrosion-induced deterioration of the structural elements. The
discussions in this paper include information obtained from relevant data on the corrosioninduced deterioration of two reinforced concrete structures in Kuwait.
Figure 1. Spalled concrete and corroded steel reinforcement
MECHANISMS OF CORROSION IN REINFORCED CONCRETE
Chloride-Induced Corrosion
The alkaline environment in concrete (pH of approximately 13) protects reinforcing steel
from corroding through the formation of a passivating iron-oxide layer on the steel’s surface.
However, when concrete is exposed to chloride solutions, the steel’s passivating layer is
destroyed, and in the presence of an adequate supply of oxygen, the steel corrodes. The
destruction of the passivating layer depends on the molar ratio of chloride ions to hydroxide
ions (Cl-/OH-). When Cl-/OH- molar ratios are higher than 0.6, steel seems to be no longer
protected against corrosion, probably because the iron-oxide layer becomes either permeable
or unstable under these conditions [4].
It is generally accepted that chloride ions react with ferrous ions to form a soluble
complex which upon reaction with hydroxide ions leads to the formation of rust, Fe(OH)2.
The chloride ions are then released back into solution for further reaction with ferrous ions.
A typical reaction between ferrous and chloride ions is as follows
2+
Fe + 4Cl
2−
2−
−
FeCl4 + 2OH
(1)
FeCl4
−
Fe( OH )2 + 4Cl
−
(2)
In Kuwait, the groundwater in the coastal zone generally contains high levels of
chlorides and sulfates. Typical values for five residential areas within the coastal zone are
342
Al-Bahar and Attiogbe
presented in Table 1. Both the chloride and sulfate values for each area are either within or
above the ranges for onset of concrete deterioration in accordance with the Uniform Building
Code [5]. Whereas concrete distress caused by expansive sulfate attack is known to be
suppressed in the presence of chloride ions, the chlorides promote corrosion of steel
reinforcement [6,7]. As such, the groundwater in the coastal zone of Kuwait would be
expected to induce corrosion when in contact with unprotected reinforced concrete elements.
Table 1. Chloride and Sulfate Levels in Groundwater in Different Coastal Areas in Kuwait
Coastal Area
Sulaibikhat
Shuwaikh
Shaab
Salwa
Fahaheel
Ranges for onset of
concrete deterioration
Chloride (ppm)
10,423
642
41,126
752
3,092
500 - 1,500
Sulfate as SO4 (ppm)
6,004
2,134
1,685
1,677
6,647
150 - 1,500
Carbonation-Induced Corrosion
A second cause of corrosion in reinforced concrete is carbonation of the concrete cover
on the steel reinforcement. Atmospheric carbon dioxide (CO2) combines with moisture in the
concrete to form carbonic acid. This acid then reacts with the cement hydration products,
particularly calcium hydroxide (Ca(OH)2), to form calcium carbonate (CaCO3) and lower the
pH of the concrete.
CO2 + H2O + Ca(OH)2 → CaCO3 + 2H2O
(3)
When the pH is lowered below 11.5 at the level of the steel reinforcement, the
passivating layer is destroyed causing corrosion to occur [4]. The rate of carbonation of the
concrete is influenced by the concentration of CO2 in the atmosphere, and by the diffusion
rate of CO2 which is dependent on the permeability and moisture content of the concrete. If
the moisture content of the concrete is very low, carbonation does not occur because of the
lack of water. If the moisture content is very high (i.e., saturated concrete pores), there is
again very little carbonation because of the very low rate of CO2 diffusion in water.
Carbonation is greatest where the pores of the concrete are partially filled with water.
Carbonation-induced corrosion in the Arabian Gulf region is usually due to insufficient
concrete cover on the steel reinforcement [6]. The low cover thicknesses seem to be the
result of a lack of an adequate specified minimum cover and poor control during construction.
In the United Arab Emirates (U.A.E.) and Bahrain [6,8], carbonation-induced corrosion has
been documented in interior concrete elements which received little or no curing. Studies in
Kuwait have shown that carbonation due to poor construction practices was the cause of
corrosion of non-marine concrete structures [2].
343
Corrosion in the Building Industry
CORROSION PROTECTION MEASURES FOR NEW STRUCTURES
The measures implemented during the construction of reinforced concrete structures to
provide protection from corrosion-induced damage caused by the effects of the service
environment usually involve the following approaches: sealing the concrete surface to prevent
ingress of chlorides, modifying the concrete to reduce its permeability, and protecting the
reinforcing steel to reduce the effects of chlorides when they do reach the steel. Multiple
levels of protection may be applied consisting of two or more of these approaches. In the
Arabian Gulf region, the protection systems frequently used are silica-fume concrete, epoxycoated reinforcing bars, and calcium nitrite chemical admixture [9]. The full effectiveness of
these protection systems can only be achieved if good concrete construction practices are
adopted. In Kuwait, corrosion protection systems are not used as normal concrete
construction practice, enhancing the vulnerability of reinforced concrete structures to
corrosion-induced deterioration.
Good concrete construction practices, such as those specified by the American Concrete
Institute (ACI) [10] and the Construction Industry Research and Information Association
(CIRIA) [11], include limiting chlorides in the concrete ingredients, providing adequate
concrete cover on reinforcing steel, and allowing adequate curing to enhance hydration and
reduce the permeability of the concrete. The implementation of these practices is a first step
in the construction of corrosion-resistant concrete structures and in the effective use of
corrosion protection systems.
INVESTIGATION AND REPAIR OF CORROSION DAMAGE
The assessment of the effects of a given service environment on the extent of corrosion
damage requires a thorough investigation to be undertaken on the structure. This
investigation is a necessary first step for a successful repair of the damage. A typical
investigation includes the following: visual inspection of the structure to assess the condition
of the structural members and to select locations for obtaining concrete samples,
nondestructive testing to identify locations of extensive degradation in concrete properties,
and laboratory testing of material samples removed from the structure. These investigative
procedures were recently applied to two reinforced concrete structures in Kuwait, one a
commercial building and the other an oil refinery structure.
Visual Inspection
The visual inspection identified the nature of the service environment and its role in
promoting corrosion in the structure. The visual inspection of the 25-year old commercial
building revealed that the humidity level in the inspected area was high and was enhanced by
a lack of ventilation. Heavy deposits of salt were formed on the structural members due to
leakage from a brine tank in an upper floor laundry room. These are ideal conditions for
corrosion to occur, as evidenced by the severely corroded reinforcement in a portion of the
structure shown in Fig. 2. In the oil refinery structure, chloride-laden steam from a hot
seawater collector escaped freely into the environment around the structure. Figure 3 shows
that extensive cracking and brown rust stains were observed on the concrete surface.
344
Al-Bahar and Attiogbe
Figure 2. Deteriorated concrete with heavily corroded steel reinforcement
Figure 3. Extensive cracking and rust stains on concrete surface
Nondestructive Testing
Rebound hammer and ultrasonic pulse velocity tests were performed to evaluate the
quality of the concrete in different parts of the structures. The ultrasonic pulse velocity or
rebound hammer data were compared for specific parts of the structures to determine the
relative levels of degradation in concrete quality. A lower value for the ultrasonic pulse
velocity or the rebound hammer number indicates a lower concrete quality. The data showed
that, for the commercial building, the level of degradation in concrete properties was
extensive in structural elements adjacent to expansion joints when compared to elements
further from the expansion joints, as is illustrated in Fig. 4 by the ultrasonic pulse velocity
data for different structural members. This extensive degradation was due to exposure of the
concrete in the vicinity of the expansion joints to moisture from leakage through the joints.
345
Corrosion in the Building Industry
For the oil refinery structure, lower values for ultrasonic pulse velocity and rebound number
were obtained for structural elements with extensive cracking.
5000
Adjacent to Expansion Joint
Further from Expansion Joint
4000
3000
2000
1000
0000
Slab
Beam
Column
Structural Member
Figure 4. Ultrasonic pulse velocity data for concrete in different structural members
Concrete Cover and Carbonation Depth
A covermeter was used to determine the thickness of concrete cover provided on the
reinforcement to protect against corrosion. Core samples were removed from the structure to
determine the extent of carbonation of the concrete cover utilizing phenolphthalein solution.
Data obtained for the commercial building showed that carbonation depth exceeded the cover
thickness in some structural members, as illustrated in Fig. 5 for beams. In these members,
carbonation would be expected to contribute to corrosion of the reinforcement.
Concrete Quality
Cores were removed from selected structural members to assess water absorption and
voids content characteristics of the concrete in order to determine the ability of the concrete
to resist ingress of moisture and salts. The water absorption and voids content data indicated
that, overall, the concrete in the structures was highly porous. The water absorption values
for the concrete samples ranged from 6 to 8% and the voids content values ranged from 15 to
18%. These values are quite high and indicate that the concrete in the structures has a low
resistance to the ingress of moisture and chlorides.
Chloride Content of Concrete
Concrete powder samples within three depth ranges from the surface of the concrete (010 mm, 10-25 mm, and 25-50 mm) were obtained from selected structural members. These
samples were used to determine the acid-soluble chloride contents of the concrete to assess
the potential for reinforcement corrosion and concrete deterioration. The concentrations of
chloride in the structures were found to exceed the threshold values for initiation of corrosion.
346
Al-Bahar and Attiogbe
Figure 6 shows a typical chloride concentration profile in the structural elements of the
commercial building. For the concrete in the commercial building, the threshold chloride
value is estimated to be 0.025% by mass of concrete based on the ACI Committee 222recommended acid-soluble chloride limit of 0.20% by mass of cement [12].
5
Concrete Cover
Carbonation Depth
4
3
2
1
0
Beam 1
Beam 2
Beam 3
Structural Member
Figure 5. Concrete cover compared with carbonation depths in beams
0.10
0.08
0.06
0.04
Estimated limit for corrosion initiation = 0.025% mass of concrete
0.02
0.00
0.5
1.75
3.75
Chloride Penetration Depth (cm)
Figure 6. A typical chloride concentration profile in beams
347
Corrosion in the Building Industry
Chloride Diffusion: The chloride diffusion model based on Fick’s second law [13-15]
was used to determine the characteristics of chloride penetration into the concrete. The model
is expressed as
C = C o erfc
where C
Co
D
t
erfc
=
=
=
=
=
x
(4)
2 Dt
chloride concentration at a distance x from the concrete surface
surface chloride concentration
effective diffusion coefficient
time
a mathematical function (Gaussian error function complement)
The chloride data was used with Eq. 4 to calculate values for D and Co by performing a
nonlinear regression analysis. These values were used with Eq. (4) in conjunction with the
threshold chloride value and the thickness of concrete cover to estimate the time to initiation
of corrosion in the structural elements of the commercial building to be in the range of 2-6
years. This indicates that corrosion started early in the life of the 25-year old building and,
therefore, measures could have been taken earlier to control its progress.
Table 2. Steel Section Loss Due to Corrosion
Structural
Member
Slab
Beam
Column
Original
Diameter of
Bar, d (cm)
1.0
1.6
1.4
1.6
1.0
2.0
2.0
1.4
1.0
1.0
1.8
1.8
1.0
1.0
1.0
Mass of Corroded
Bar/Unit Length,
ms (g/cm)
5.73
15.76
10.08
15.48
5.92
22.64
22.39
10.56
3.17
2.35
13.80
13.05
5.92
5.88
5.87
Average Reduction in
Bar Diameter, Δd (μm)
366
22
1222
164
208
850
956
920
2834
3830
3048
3460
208
240
248
CrossSectional Area
Loss (%)
7.2
0.3
16.7
2.0
4.1
8.3
9.3
12.7
48.6
61.9
31.0
34.7
4.1
4.7
4.9
Loss of Steel Section: Some cores removed from the structural members contained
pieces of steel bars. These steel bars were cleaned using Clarke's solution [16] and weighed
to determine their mass loss per unit length. The relation between mass, volume and density
was used to calculate the average reduction in bar diameter based on the original diameter of
348
Al-Bahar and Attiogbe
the steel bars and a density of 7.86 g/cm3 for the steel. The values of average reduction in bar
diameter were used to calculate the cross-sectional area loss of the steel reinforcement at
different locations in the structural elements, as presented in Table 2 for the commercial
building. The estimates for steel section loss were as high as 62% in parts of the commercial
building and 88% in parts of the oil refinery structure. At these levels of steel section loss,
extensive cracking of the concrete would be expected, as was observed in the structures.
Eq. 4 was used with the estimated values of D and Co to calculate the maximum chloride
concentrations at the level of the steel reinforcement. The values for the commercial building
are summarized in Table 3, in addition to values of average thickness of concrete cover,
average depth of carbonation and maximum section loss of steel reinforcement. These results
show that both chlorides and carbonation were responsible for corrosion of the steel
reinforcement. Chloride-induced corrosion is indicated where the maximum chloride
concentration at the level of the steel reinforcement exceeds the threshold value of 0.025% by
mass of concrete and the carbonation depth is less than the thickness of the concrete cover.
Carbonation-induced corrosion is indicated where the carbonation depth exceeds the cover
thickness and the chloride concentration is less than the threshold value.
Table 3. Extent and Cause of Corrosion of Reinforcing Steel
Structural
Member
Slab
Beam
Adjacent to
Expansion
Joint
Max. Chloride
Concentration at
0.001
Steel Level (%
Mass of Concrete)
Average Concrete
4.0
Cover (cm)
Average
4.5
Carbonation
Depth (cm)
Max. Steel
16.7
Section Loss (%)
Type of
Carbonation
Corrosion*
-induced
corrosion
Column
Further from
Expansion
Joint
1.540
0.003
0.058
3.0
3.0
4.0
0.0
4.0
4.2
61.9
12.7
4.9
Chlorideinduced
corrosion
Carbonationinduced
corrosion
Chloride-induced
corrosion with
carbonation
contributing
* Based on chloride threshold value of 0.025% by mass of concrete.
Guidelines for Repair of Corrosion Damage
349
Corrosion in the Building Industry
The information gained from the corrosion damage investigation provides the basis for a
successful repair of the structure. In general, the arrest of chloride-induced corrosion requires
the removal of chloride contaminated concrete around the steel reinforcement. This would
require the mechanical removal of the contaminated concrete or the application of an
electrochemical method for chloride removal. Carbonation-induced corrosion can be arrested
by installing surface barriers. Surface applied barriers which have low vapor transmission
allow the uncarbonated concrete to re-alkalize or increase the pH of the carbonated concrete,
pushing the carbonation front back toward the surface. The carbonation process in cracked
concrete can be arrested by using elastomeric membranes for crack bridging or by using crack
sealants. Electrochemical technology can also be used to transport alkalies into carbonated
concrete [17].
The following general guidelines are appropriate when repairing corrosion-induced
damage in reinforced concrete structures.
• Cracked concrete should be repaired after an adequate preparation of the concrete
•
•
•
•
surface to ensure a good bond between the repair and the base concrete,
Whenever possible, cracked concrete cover should be removed to at least 2 cm
beyond corroded steel bars,
During repair of cracked concrete cover, all exposed, severely corroded
reinforcement (i.e., > 25% section loss) should be repaired by providing new
reinforcement,
After concrete repair, protective coating may be applied on the surfaces of the
structure to protect against the ingress of moisture, chlorides and gases, and
Periodic monitoring of the structure may be undertaken, following completion of
the repair work, to assess the extent of any further corrosion activity.
CONCLUSIONS
1. The following conclusions are drawn based on the discussions in this paper:
2. The investigative techniques employed provided information on the nature of the
service environment and its role in promoting corrosion of steel reinforcement in the
concrete structures.
3. Lack of adequate measures to mitigate the effects of the service environment
contributes significantly to corrosion-induced deterioration of reinforced concrete
structures in Kuwait.
4. Both chloride-induced corrosion and carbonation-induced corrosion are found to
occur in the structures.
REFERENCES
1. Rasheeduzzafar; F.H. Dakhil and A.S. Al-Gahtani, Deterioration of concrete structures in
the environment of the Middle East, ACI Journal 81, 1, 1984, pp. 13-20.
2. H.M. Shalaby, Case studies of corrosion and deterioration of reinforced concrete
structures in the State of Kuwait, First International Conference on Deterioration and
Repair of Reinforced Concrete in the Arabian Gulf, Bahrain, 1985, pp. 309-320.
350
Al-Bahar and Attiogbe
3. S. Al-Bahar and E. Attiogbe, Corrosion-induced deterioration of reinforced concrete
structures in Kuwait, Concrete Under Severe Conditions, Environment and Loading.,
Proceedings CONSEC ‘95, International Conference, K. Sakai, N. Banthia and O.E.
Gjorv Eds, Sapporo, Japan, Vol.1, August 1995, pp. 564-573.
4. P.K. Mehta, Concrete: Structure, Properties, and Materials, Prentice-Hall, Inc.,
Englewood Cliffs, New Jersey, 1986.
5. Uniform Building Code, International Conference of Building Officials (ICBO),
California, 1991.
6. Z.G. Matta, Deterioration of concrete structures in the Arabian Gulf, Concrete
International 15, 7, 1993, pp 33-36.
7. B.C. Gerwick, International experience in the performance of marine concrete,
Concrete International 12, 5, 1990, pp 47-53.
8. C. Le Sage de Fontenay, A study of the effect of concrete admixtures, concrete
composition and exposure conditions on carbonation in Bahrain, First International
Conference on Deterioration and Repair of Reinforced Concrete in the Arabian Gulf,
Bahrain, 1985, pp. 467-483.
9. Z.G. Matta, Protecting steel in concrete in the Persian Gulf, Materials Performance,
1994, pp. 52-55.
10. ACI 318, Building Code Requirements for Reinforced Concrete, American Concrete
Institute, Detroit, 1989.
11. CIRIA, Guide to Concrete Construction in the Gulf Region, CIRIA Special
Publication 31, 1984.
12. ACI Committee 222, Corrosion of Metals in Concrete, American Concrete Institute,
Detroit, 1989.
13. J. Crank, Mathematics of Diffusion, 2nd Edition, Oxford University Press, 1975.
14. N.S. Berke, D.W. Pfeifer and T.G. Weil, Protection against chloride-induced
corrosion, Concrete International 10, 12, 1988, pp. 45-55.
15. O.E. Gjorv, K. Tan and M. Zhang, Diffusivity of chlorides from seawater into highstrength lightweight concrete, ACI Materials Journal 91, 5, 1994, pp. 447-452.
16. ASTM G 1, Standard practice for preparing, cleaning, and evaluating corrosion test
specimens, Annual Book of ASTM Standards, Vol. 03.02, Philadelphia, Pennsylvania,
USA, 1990.
17. P.H. Emmons, Concrete repair and maintenance illustrated, R.S. Means Company,
Inc., Massachusetts, USA, 1993.
351
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
CORROSION OF CONCRETE IN SEAWATER
M. Pakshir and S. Esmaili
Department of Materials Science and Engineering
Shiraz University, Shiraz, Iran
ABSTRACT
Attack and disintegration of concrete structures under the influence of aggressive fluids, wet
gases, etc., are understood as corrosion of concrete. The corrosion of concrete is a chemical, colloidchemical process, yet often a physicochemical one, while that of concrete reinforcement is mainly
electrochemical.
Many sea harbor jetties, docking facilities and concrete-supported oil platforms are subjected to
the marine environment in the Persian Gulf, and therefore become corroded. In the present paper,
using an aerated synthetic seawater representative of Persian Gulf environment, we have tried to
investigate the time required for corrosive agents such as chloride and oxygen to penetrate through the
concrete. The influence of the concrete's cover depth and corrosion rate of the rebar has also been
studied.
Key Words: Concrete, reinforcement, concrete cover, seawater, corrosion
INTRODUCTION
In recent years, the performance of concrete in marine environments has assumed
considerable importance because of the offshore activity of gas and oil exploration in various
countries. Thus, concrete performance in marine environments has been investigated by
several researchers [1,2,31.
The interaction between the concrete’s service environment and the concrete itself can
lead to deterioration of reinforced concrete structures, and in some cases, render the structures
unsuitable for their designed purposes. The interaction is often with the chemical species in
the environment [4].
For concrete in a marine environment, there appears to be a direct correlation between
low permeability (i.e., high strength) and good durability. Therefore, concrete sea structures,
such as harbor structures and offshore platforms, are built using high quality concrete [5]. In
general, concrete is a fine, porous material. Pore sizes vary from a few angstroms to several
millimeters. This pore system is more or less filled with a solution which contains varying
quantities of salts.
The problems associated with the application of reinforced concrete in marine
environments are well known. These problems have led to extensive research on the
corrosion of metal in concrete structures. Concrete reinforcement is usually protected from
corrosion by the highly alkaline environment (pH about 11.5) of the concrete surrounding it.
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Corrosion in the Building Industry
As a result, the steel surface frequently develops a protective oxide layer which is difficult to
dissolve, and hence limits its disintegration [1,6,7]. Corrosion of the steel normally happens
either by the carbonation of the concrete around it which reduces the alkalinity, or through the
presence of even small quantities of chlorides in the concrete cover.
The depth of penetration of chloride is a function of time and the permeability of the
concrete [8]. Once the steel has become depassivated, the rate of corrosion will depend,
among other factors, upon the availability of moisture and oxygen near the steel surface
[9,10]. In the absence of either moisture or oxygen, corrosion will not occur.
The corrosion of steel causes the metal to be converted in different stages into various
ferric oxides and ferric hydroxide [11]. This change leads to an increase in volume. Damage
resulting from corrosion can also be seen in the form of cracks parallel to the embedded steel,
and finally cracking and spalling of the concrete which accelerates the rate of corrosion.
The corrosion of reinforcing steel in concrete results from an electrochemical process
[12] accompanied by anodic and cathodic reactions. Iron is transferred in the solution as
hydrated ions with two electrons left in the reinforcing steel. These two electrons are
transferred to the cathodic area to assimilate in the cathodic reaction.
EXPERIMENTAL PROCEDURE
The experiments performed were designed with two purposes: the first was to study the
time that a corrosive agent takes to penetrate through the concrete cover to the reinforcing
steel, and the second was to investigate the influence of the depth of cover for a fixed watercement (w/c) ratio.
Tests were performed on mortar specimens partially submerged in an aerated seawater
(Table 1). The mortar specimens were made with ordinary Portland cement and had w/c
ratios of 0.4, 0.5, 0.6, and 0.7, with a cement-sand ratio of 1:3.
Table 1 Chemical Composition of Artificial Seawater
NaCl
(g/l)
32
MgCl2.6H2
O (g/l)
6
MgSO4.7H2O
(g/l)
5
CaSO4.2H2O
(g/l)
1.5
KHCO3
(g/l)
0.2
The aggregate used was crushed stone, and it was used in a dry condition when the
concrete was mixed. The sand was ordinary local river sand. The specimens made were 20 x
55 x 80 mm with two round carbon steel bars 7 mm in diameter embedded symmetrically.
The steel bar was embedded 50 mm into the mortar, and to avoid the formation of a
differential aeration cell, the ends of the embedded bars were insulated with adhesive tape.
The Portland cement was cured at 20°C for 28 days. Since the procedure required a constant
2
current density (mA/mm ) over the reinforcement surface, the specimens were exposed to
seawater in the tank for several days to obtain a uniformity of moisture in the concrete before
the impressed current was applied. Direct current (DC) was impressed on the specimen from
a DC rectifier, and a potentiometer was used to control the current applied to each specimen.
354
Pakshir and Esmaili
Current direction was so arranged that the reinforcing steel bars served as anodes and a
stainless steel rod was employed as a counter electrode to serve as a cathode (Fig. 1). The
potential of each specimen was measured daily using a saturated calomel electrode with a
voltmeter twice a day, and any drift was corrected by adjusing the potentiometer. Potential
readings were taken with the power on. The readings represented the polarization potential of
the reinforcement plus the voltage drop across any corrosion product film at the
reinforcement /cement interface.
In order to study the effect of different cover depths on the corrosion rate of reinforcing
steel, 200 x 300 x 125 mm slab specimens with a 25 mm pond cast on top were made from
Portland cement at a w/c ratio of 0.5. Each specimen contained two carbon steel electrode
rods of 3 mm diameters embedded at four depths of cover, i.e., 10 , 20 , 30 and 40 mm.
The electrochemical corrosion measurements of the embedded steel reinforcement were
carried out by means of polarization resistance measurements and rest potentials. The
corrosion rates were calculated from the polarization resistance value using the Stern-Geary
equation, i.e.:
Icorr = Iappl / 2.3 ΔE (βcβa / (βc + βa))
(1)
Icorr was calculated on the assumption that both the anodic and cathodic Tafel constants, βa
and βc, had values of 120 mV/decade.
Figure 1. Schematic illustration of polarization measurements
355
Corrosion in the Building Industry
RESULTS AND DISCUSSION
The corrosion of concrete reinforcement in aerated seawater is expressed in terms of
potential and current in the embedded rebar as a function of exposure time. The temperature
and oxygen content of the seawater were checked regularly. The temperature varied from
20°C to 25°C, depending on the time of year, while the oxygen content was kept about 7 ppm
at all times.
The electrochemical potential was measured via a high-impedance
millivoltmeter against a calomel electrode.
Figure 2 shows the typical half-cell potential versus time behavior of embedded steel
reinforcement in aerated seawater. It shows that the time required for the reinforcing steel to
be in an active potential is about 3.5 weeks for a w/c ratio of 0.7 and increases to about 13
weeks for a w/c ratio of 0.4. In other words, the time for the corrosive agent to reach the
embedded steel is the time it takes the steel reinforcement to change from a passive potential
to an active potential, which occurs when sufficient chloride ions and oxygen are present to
cause corrosion of the steel bar. Therefore, one can assume that a high w/c ratio makes the
concrete permeable enough for the corrosive agent to penetrate to the reinforcing steel and
then depassivate it. Thus, as the w/c ratio increased, the potential of the reinforcing steel bar
became less noble.
A general decrease in the observed corrosion current was found to occur with decreases
in the w/c ratio (Fig. 3). In order to accelerate the corrosion process, a constant electrical
potential of 3.5 V was impressed into the embedded steel bar.
Figure 2. Corrosion potential
vs. exposure time
Figure 3. Corrosion current density
(Icorr) vs. exposure time
Concrete cracking due to corrosion of the reinforcement was observed after 14 weeks for
a w/c ratio of 0.4, after 10.5 weeks for a w/c ratio of 0.5, after 4.5 weeks for a w/c ratio of 0.6
and after 4 weeks for a w/c ratio 0.7. Therefore, the sharp rise in current in Fig. 3 is related to
the time the appearance of the cracks was observed. However, the failure of the concrete
356
Pakshir and Esmaili
specimen due to the corrosion of the steel reinforcement is related to the initial current of the
steel bar, and as can be seen, the initial current of the steel bar depends upon the w/c ratio of
the covered concrete. Therefore, for the reinforced steel embedded in the concrete with a
lower w/c ratio, the initial current was lower and it took a longer time for the corrosive agent
to reach the steel reinforcement and cause corrosion failure. The results of measurement of
the distribution of the corrosion potential and corrosion current density of the reinforcing
steel after 70 days of exposure for different cover thicknesses are shown in Figs. 4 and 5.
Figure 4. Corrosion current
density vs. exposure time
Figure 5. Corrosion potential vs.
exposure time
In the case of a cover thickness of 10 mm, the degree of corrosion was large, and the
specimen had a crack of about 0.2 mm due to the corrosion of the reinforcing steel. The
corrosion of the steel rebar at the location of the crack was particularly remarkable. On the
other hand, with an increased cover thickness of 40 mm, very little corrosion of the
reinforcing steel bar was observed.
Visual examination of the reinforcement bar after 70 days revealed a white deposit on
the surface. The amount of the deposit varied with the depth of the cover thickness. On
microscopic examination, the deposit appeared to be crystalline in structure and was very
dense in nature (Fig. 6).
Chemical analysis of the deposit showed that it consisted of a mixture of calcium
carbonate and magnesium hydroxide. Since seawater contains a lot of magnesium sulfate,
which is the most harmful salt as far as cement attack is concerned, it reacts with calcium
hydroxide and by substitution of the magnesium for calcium, secondary gypsum in flat prisms
and brucite Mg(OH)2 in superimposed platelets are formed (Fig. 7).
357
Corrosion in the Building Industry
(a) Crystallization of the brucite layer
aragonite layer
(b) Crystallization of the
(c) Aragonite deposit on the brucite layer
Figure 6. Morphology of calcareous deposits after seventy days of exposure to seawater
Figure 7. The dissolution of calcium hydroxide and the formation of secondary gypsum
CONCLUSIONS
Since seawater contains a lot of chloride in various forms as well as sulfate salts, the
initial period of attack is the time taken for the chloride and sulfate ions to penetrate from the
surface through the concrete cover and reach the steel reinforcement bar. During this period,
the steel is in the passive state until the threshold value of these aggressive ions is reached for
corrosion to initiate in the presence of oxygen.
A possible form of corrosion attack on the reinforcement in the initial stage was pitting
corrosion. There is good reason to assume that the number of pits was fairly high, as was true
in the case of a 10 mm cover thickness. Also, since the amount of available cathodic current
was constant, regardless of the number of pits, the average anodic current was small.
Therefore, the most important consequence of pitting is that oxygen is consumed in the
process, i.e., as oxygen is consumed, the steel potential drifts to a negative value and
corrosion takes place. Hence, the availability of oxygen is one of the main controlling factors
358
Pakshir and Esmaili
of the corrosion of steel rebar in concrete. The amount of corrosion discovered when the
reinforcement was removed was insignificant as the depth of cement cover increased, small
spots of corrosion products were in a few instances visible at the surface.
A high water-cement ratio makes concrete permeable enough for the aggressive ions to
penetrate to the reinforcing steel easily and then depassivate the reinforcement. Thus, as the
water-cement ratio of the concrete increases, the potential of the reinforcing steel becomes
less noble and the electric resistance of wet concrete drops due to yje low permeability which
accelerates the corrosion of reinforcing steel.
REFERENCES
1. P.K. Mehta and P.J. Monteiro, Concrete: Structure, Properties and Materials, 2nd ed.,
Prentice Hall, Englewood Cliffis, 1993.
2. P. Rodriguez, E. Ramirez and A. Gonzalez, Magazine Concrete Research 46, 167, June
1994, pp. 81-91.
3. O.A. Kayyali and M.N. Haque, Magazine Concrete Research 47, 172, Sept. 1995, pp.
235-242.
4. D. Bawega, H. Roper and V. Sirivivatnanon, Cement and Concrete Research Journal 23,
1993, pp. 1418 -1430.
5. H.H. Haynes, American Concrete Institute, Detroit, Michigan, USA, SP-65, 1982, pp. 2138.
6. C. Androde, Corrosion of Steel in Concrete: Monitoring Techniques, Report of the
Technical Committee 60-csc Rilem, Champman and Hall , New York, USA, 1988.
7. P. Lambert, C.L. Page and P.R.W. Vassie, Materials and Structures 24, 1991, pp. 351358.
8. C. Liam, S.K. Roy and D.O. Northwood, Magazine Concrete Research 42, 160, Sept.
1992, pp.205-215.
9. A. Gonzalez, A. Molina, E. Otero and W. Lopez, Magazine Conerete Research 42, 150,
March 1990, pp. 23-27.
10. C.L. Page and P. Lambert, Journal of Materials Science 22, 1987, pp. 942-946.
11. C.A. Lawrence, British Ceramic Proceedings, Cement and Concrete Association,
Wexham, UK, No.35, 1984, pp.277-293.
12. K. Okado and T. Miyagawa, American Concrete Institute, Detroit, Michigan, USA, SP65, 1982, pp.237-254.
359
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
CONCRETE QUALITY AND ITS EFFECT ON CORROSION
OF STEEL REINFORCEMENT
E.K. Attiogbe and S. Al-Bahar
Civil Engineering & Building Department
Kuwait Institute for Scientific Research
P.O. Box 24885, 13109 Safat, Kuwait
ABSTRACT
To minimize the incidence of premature deterioration of concrete, issues that need to be carefully
attended to include the establishment of material specifications that match the severity of the service
exposure conditions, and the implementation of construction practices that enable the desired level of
durability to be obtained. The paper discusses the characteristics of concrete ingredients that relate to
corrosion-induced deterioration of reinforced concrete structures. The role of supplementary
cementing materials in enhancing the resistance of concrete to reinforcement corrosion and the effects
of high ambient temperatures on the corrosion-related behavior of the concrete microstructure are
discussed. Selecting appropriate concrete materials that compensate for the aggressiveness of the
exposure condition is a critical factor in ensuring the durability of reinforced concrete.
Key Words: Concrete durability, concrete materials, concrete quality, reinforced
concrete, reinforcement corrosion
INTRODUCTION
The performance of concrete under harsh service conditions, such as those prevalent in
the Arabian Gulf region, is controlled by the interaction between the concrete and the
environment. Corrosion of steel reinforcement has been established as the major cause of
premature deterioration of concrete structures in the Arabian Gulf states [1-5]. Reinforced
concrete structures constructed with a design life of 50 years or more have become
structurally unsound within 10 to 20 years. The factors that promote this premature
deterioration of concrete structures include the use of unsuitable concrete mix ingredients,
poorly designed concrete mixes, and improper concrete placement and curing.
Chlorides are the primary cause of corrosion in reinforced concrete structures in Kuwait
and other Arabian Gulf states, with carbonation as a secondary cause [1-3]. A number of
studies have evaluated the role of different concrete ingredients in enhancing the corrosion
resistance of reinforced concrete, as well as the corrosion performance characteristics of
concrete materials used in the Arabian Gulf region [6-10]. In this paper, the characteristics of
concrete ingredients with respect to corrosion of reinforcement are discussed. In addition, the
paper presents the basis for producing durable concrete that is resistant to reinforcement
corrosion in the hot and arid environment of the Arabian Peninsula.
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Corrosion in the Building Industry
CONCRETE MATERIALS
Aggregates
Aggregates used for concrete mixes in Kuwait and other Arabian Gulf states are
typically contaminated with dust and salts. To enhance the durability of concrete made with
these aggregates, it is prudent to wash the aggregates to remove or minimize the chloride and
sulfate salts. However, due to the scarcity of water in the region, the aggregates may not be
washed prior to use in the concrete mix [1,3,9], enabling the corrosive action of the chlorides
to commence at a very early stage when the concrete is still very weak and porous.
In Kuwait, the sand is usually washed, whereas the coarse aggregate is sieved to remove
dust prior to use in concrete mix. Table 1 shows the chloride and sulfate levels in washed
sand and sieved coarse aggregate used for a residential building project. While the chloride
levels are lower than the limits for concrete deterioration as specified by the CIRIA Guide
[11], the level of sulfate in the sieved coarse aggregate is greater than the CIRIA Guide’s
limit. Where cracking of concrete takes place due to high levels of sulfate contamination,
chloride ingress from the salt-laden atmosphere in the coastal zone can occur leading to
reinforcement corrosion. Therefore, adequate steps must be taken to minimize both chlorides
and sulfates in concrete aggregates.
Table 1. Chloride and Sulfate Concentrations in Aggregates
Material
Sand (Washed)
Coarse Aggregate (Sieved)
Limits for Concrete
Deterioration
Chloride (% mass of agg.)
<0.001
<0.001
0.06 (sand)
0.03 (coarse)
Sulfate (% mass of agg.)
0.092
0.596
0.4
Mix Water
Potable or desalinated water is scarce in the Arabian Gulf region. As such, brackish
service water (nondesalinated tap water) is sometimes used in mixing concrete [1,9]. This
practice introduces quantities of chloride and sulfate into the concrete at the time of mixing.
In a study undertaken by Rasheeduzzafar et al. [7], the chloride and sulfate concentrations in
brackish water used to evaluate the influence of construction practices on concrete durability
were 1294 ppm and 375 ppm, respectively, compared with corresponding concentrations of
32 ppm and 32.5 ppm in potable water. The salt concentrations in the brackish water are
within the ranges specified by the Uniform Building Code [12] for onset of concrete
deterioration. These ranges are 500-1500 ppm for chlorides and 150-1500 ppm for sulfates.
The chloride concentration of 1294 ppm exceeds the maximum concentration of 500 ppm
recommended by the CIRIA Guide for water used in mixing or curing reinforced concrete.
Use of mix water of drinking quality would reduce the risk of corrosion in reinforced concrete
structures.
Portland Cement
362
Attiogbe and Al-Bahar
The major compounds in portland cement are tricalcium silicate (C3S), dicalcium silicate
(C2S), tricalcium aluminate (C3A) and tetracalcium aluminoferrite (C4AF). The C3A phase
has the ability to bind chlorides, resulting in the formation of insoluble calcium
chloroaluminate compound (3CaO.Al2O3.CaCl2.10H2O) or Friedel's salt. This reduces the
amount of free chloride ions in the concrete pores and, hence, lowers the corrosion risk.
Figure 1 shows that the concentration of free or unbound chlorides decreases with an increase
in the C3A content of cement [8]. The unbound chlorides were found to be 86, 58, 51 and
33% of the total chlorides in concrete made with cements having C3A contents of 2.04, 9.10,
11.02 and 14.00%, respectively.
Studies have shown that the time to initiation of corrosion increases with increasing C3A
content of cement [6,8]. However, the chloride-binding capacity of C3A decreases with an
increase in the chloride concentration, becoming almost ineffective when concrete is exposed
to high chloride concentrations [8]. In addition, the long-term effectiveness of chloride
binding by C3A in reducing corrosion risk is not yet clear, as the chlorides may be released
later in the life of the concrete structure due to carbonation or sulfate attack of the concrete
[13,14]. Thus, it appears that cements high in C3A should not be depended upon solely for
the purpose of reducing the risk of corrosion in reinforced concrete structures.
100
80
60
40
20
0
0
2
4
6
8
10
12
14
16
C3A Content of Cement, %
Figure 1. Effect of the C3A content of cement on chloride binding
The composition of modern portland cements, particularly that of cements used in the
Arabian Gulf region, usually is characterized by markedly higher C3S-to-C2S ratios than
those in older cements [8]. Therefore, the modern cements have a higher rate of hydration
and strength gain than the older cements. When concretes made with the modern cements are
specified in terms of 28-day strength only, they usually satisfy the specification at higher
water-cement ratios than concretes made with the older cements. In general, this would lead
to a more permeable concrete with a low resistance to reinforcement corrosion due to the
ingress of chloride solutions. Specifying concrete merely on strength considerations while
ignoring factors relevant to its durability would result in durability problems of varying
severity. Reducing the risk of reinforcement corrosion requires laying down specific
363
Corrosion in the Building Industry
provisions for using concrete permeability characteristics, along with strength, to match
concrete quality to environmental exposure conditions.
CONCRETE MIX CHARACTERISTICS
Water-Cement Ratio
Lowering the water-cement ratio (w/c) of concrete reduces its permeability to ingress of
chloride solutions. Lower w/c concretes have a higher electrical resistivity which impedes
the flow of corrosion current. The rate of oxygen diffusion is also substantially reduced the
lower the w/c, contributing to significant enhancements in polarization resistance and,
thereby, to reductions in the magnitude of corrosion current.
Low w/c concretes are, however, limited in their effectiveness in delaying the onset of
chloride-induced corrosion. A reduction in w/c from 0.50 to 0.35 improved the time to
initiation of corrosion by only 20% [10]. This marginal increase in the time to corrosion
initiation is attributed to the fact that the reduction in the permeability of the concrete due to a
lowering of the w/c is less effective in the case of diffusivity of chloride solutions than in the
case of diffusivity of plain water [15]. That is, chloride ions in solutions, particularly in high
concentrations, are able to diffuse through concrete made with plain portland cement at much
faster rates than with diffusion of water. Enhanced resistance to chloride ion ingress is
obtained when supplementary cementing materials, such as silica fume, are used in the
concrete mix in addition to lowering the water-cementitious materials ratio.
Supplementary Cementing Materials
The permeability of concrete is significantly reduced when supplementary cementing
materials such as condensed silica fume, fly ash or blast furnace slag are used with portland
cement. The decreased permeability substantially increases resistance to chloride penetration
and reduces the rate of carbonation. The supplementary cementing materials produce
concretes with low chloride diffusivity and high electrical resistivity. A 9% replacement of
cement by silica fume reduced the chloride diffusivity by a factor of about 5 [16].
Figure 2 shows the increase in electrical resistivity of concrete with increasing content of
silica fume [17]. The effect is more pronounced the higher the cement content. For a 400
kg/m3 cement content, the increase in resistivity was 550 and 1600% for silica fume additions
of 10 and 20%, respectively. The times to initiation of corrosion for silica fume concretes
compared with plain cement concretes are presented in Fig. 3 [6], showing that silica fume
significantly delays the onset of corrosion. Figure 3 indicates that there may not be any
significant advantage in increasing the cement replacement by silica fume in Type I portland
cement concrete (C3A content of 9 to 14%) to more than 10%. This is of considerable
practical significance in the Arabian Gulf region, where the cost of silica fume is reported to
be over ten times the cost of ordinary portland cement [6]. A fly ash replacement level of
30% by mass of cement increased corrosion initiation time by a factor of about 2 [9]. For
blast furnace slag, effective replacement levels to reduce the risk of corrosion are 50 to 70%
by mass of cement [18].
The benefits of using supplementary cementing materials can only be fully realized if the
concrete is treated properly during construction. With regard to silica-fume concrete, for
example, the typical low water content and lack of bleeding may give rise to plastic shrinkage
cracking, particularly in the hot and arid environment of the Arabian Gulf region. Techniques
364
Attiogbe and Al-Bahar
for dealing with this problem include starting the curing process early and using evaporationretarding materials. The implementation of good concrete construction practices that account
for the characteristics of the particular supplementary cementing material used are essential to
enhancing the resistance of concrete to reinforcement corrosion.
140
100 kg cement/cubic meter
250 kg cement/ cubic meter
400 kg cement/cubic meter
120
100
80
60
40
20
0
0
10
20
Addition of Condensed Silica Fume, %
Figure 2. Effect of silica fume on electrical resistivity of concrete
800
Plain Cement
10% Cement Replacement by Silica Fume
20% Cement Replacement by Silica Fume
600
400
200
0
0
2
4
6
8
10
12
14
16
C3A Content of Cement, %
Figure 3. Time to initiation of corrosion for replacement of cement by silica fume
Corrosion-Inhibiting Chemical Admixture
Calcium nitrite admixture is a commonly used chemical corrosion inhibitor for
reinforced concrete structures. It is added to the concrete mix during batching and enhances
the stability of the passivating layer on the surface of the reinforcing steel. Nitrite ions react
365
Corrosion in the Building Industry
with ferrous ions according to Eq. 1 [19] to produce a stable passive layer of ferric oxide
(Fe2O3):
2 Fe2 + + 2OH − + 2 NO2− → 2 NO + Fe2O3 + H2O
(1)
The chloride and nitrite ions compete for ferrous ions. If the chloride ion concentration
is greater, the corrosion process will start. If, on the other hand, the nitrite ion concentration
is greater, passive layer will form to close off the iron surface. Equation 1 shows that during
the reaction between the nitrite and ferrous ions, the supply of nitrite is depleted. The
effectiveness of the calcium nitrite admixture, therefore, is dependent on an accurate
prediction of the chloride loading of the structure over its expected design life and, hence, on
the selection of an appropriate dosage of the admixture. In the hot and chloride-laden
environment of the Arabian Gulf region, selection of an appropriate amount of calcium nitrite
should be carefully considered for cost-effective use of the admixture.
CONCRETE PLACEMENT AND CURING
Concrete placement, such as placement with bucket or by pumping, should be carried out
in such a manner as to ensure that the concrete remains cohesive. Inadequate consolidation of
the in-place concrete contributes significantly to the corrosion-induced deterioration of
concrete structures. Poor placement and consolidation is manifested in the form of large
voids, extensive honeycombing, rock pockets and bugholes, resulting in concrete of high
permeability. The effect of consolidation on corrosion initiation was evaluated by
Rasheeduzzafar et al. [7]. The results showed that the time to initiation of corrosion increases
with the consolidation effort, which was taken as the length of time for consolidating concrete
specimens on a vibrating table. Full or 100% consolidation was reached when the emission
of air bubbles ceased and the concrete surface was covered with a thin layer of cement paste.
With 40, 60 and 70% of full consolidation effort, the time to initiation of corrosion is,
respectively, 60, 76 and 95% of the corrosion initiation time in specimens where full
consolidation is achieved.
Curing is essential to producing high-quality concrete, and even more so in the hot-arid
climatic conditions of the Arabian Gulf region where extremely rapid, excessive evaporation
of moisture occurs from concrete surfaces. Water lost by evaporative drying can seriously
hamper the cement hydration reactions and the filling of capillary pores by the hydration
products. Hydration can take place only when the vapor pressure in the capillaries is
sufficiently high (i.e., about 80% of the saturation pressure). Below 30% saturation pressure,
there is only negligible hydration [20]. As such, the initiation stage and duration of curing are
crucial to producing durable, corrosion-resistant concrete under the climatic conditions in the
Gulf region. The beneficial effect of moist curing on the time to initiation of corrosion was
demonstrated by Rasheeduzzafar et al. [7]. Figure 4 shows that the time to initiation of
corrosion increases with the length of the curing period.
Placement and curing of concrete under the high ambient temperature conditions in the
Gulf region can cause visible thermal cracking of the newly placed concrete. Also, concrete
exposed to such high ambient temperatures at early ages develops cracks within its
microstructure, as shown in Fig. 5. This early-age cracking can lead to ingress of chlorides
366
Attiogbe and Al-Bahar
either from the salty groundwater or from the salt-laden atmosphere in the coastal zones of
the region. In addition, studies have shown that the higher the curing temperature, the higher
the porosity of concrete [21-23]. In a study of the microstructure of concrete cured to the
same degree of hydration at temperatures of 5, 20 and 50oC [21], the volume of hydration
products was found to decrease and the porosity was found to increase with an increase in
curing temperature (Table 2). This effect of the curing temperature on the microstructure
leads to increased chloride penetration of concrete cured at higher ambient temperatures. For
concretes of different compositions, the rate of chloride diffusion increased by a factor of at
least 3 when the curing temperature was increased from 23 to 70oC [22]. Hot weather
concreting measures as outlined by ACI [24] and CIRIA [11] need to be carefully followed to
minimize the detrimental effects of high ambient temperatures. In addition, supplementary
cementing materials may be used to mitigate the effects of elevated temperature curing.
120
100
80
60
40
20
0
0
5
10
15
20
25
30
Curing Period, days
Figure 4. Effect of curing period on time to initiation of corrosion
367
Corrosion in the Building Industry
Figure 5. Cracking within concrete microstructure
Table 2. Effect of Curing Temperature on Concrete Microstructure
Curing Temperature (oC)
5
20
50
Hydration Product (% vol.)
84.6
78.4
74.7
Porosity (% vol.)
4.3
10.9
15.1
CONCLUDING REMARKS
Based on the discussions in this paper, the following comments are offered as
recommendations to reduce the risk of corrosion in reinforced concrete structures:
1. Adequate steps must be taken to minimize both chlorides and sulfates in
concrete aggregates,
2. Concrete mix water should be of drinking quality,
3. Concrete permeability characteristics should be considered along with strength
requirements when specifying concrete quality to match specific
environmental conditions,
4. Low w/c ratio concrete used in conjunction with supplementary cementing
materials is effective in enhancing resistance to reinforcement corrosion,
5. Adequate placement and consolidation is necessary to produce concrete with
low permeability to the ingress of chloride solutions,
6. Moist curing of sufficient duration is highly beneficial in delaying the onset of
corrosion, and
7. Measures necessary for hot weather concreting should be carefully
implemented to minimize the detrimental effects of high ambient temperatures
on concrete durability.
REFERENCES
1. Rasheeduzzafar; F.H. Dakhil and A.S. Al-Gahtani, Deterioration of concrete
structures in the environment of the Middle East," ACI Journal 81, 1, 1984, pp. 13-20.
2. H.M. Shalaby, Case studies of corrosion and deterioration of reinforced concrete
structures in the State of Kuwait, First International Conference on Deterioration and
Repair of Reinforced Concrete in the Arabian Gulf, Bahrain, 1985, pp. 309-320.
3. Z.G. Matta, Chlorides and corrosion in the Arabian Gulf environment, Concrete
International 14, 5, 1992, pp. 47-48.
4. A.A. Hamid, Improving structural concrete durability in the Arabian Gulf, Concrete
International 17, 7, 1995, pp. 32-35.
5. S. Al-Bahar and E. Attiogbe, Corrosion-induced deterioration of reinforced concrete
structures in Kuwait, Proceedings CONSEC ‘95, International Conference, K. Sakai,
N. Banthia and O. E. Gjorv, Eds, Sapporo, Japan, Vol.1, August 1995, pp. 564-573.
368
Attiogbe and Al-Bahar
6. Rasheeduzzafar, S.S. Al-Saadoun and A.S. Al-Gahtani, Reinforcement corrosionresisting characteristics of silica-fume blended-cement concrete, ACI Materials
Journal 89, 4, 1992, pp. 337-344.
7. Rasheeduzzafar, A.S. Al-Gahtani and S.S. Al-Saadoun, Influence of construction
practices on concrete durability,” ACI Materials Journal 86, 6, 1989, pp. 566-575.
8. Rasheeduzzafar, Influence of cement composition on concrete durability, ACI
Materials Journal, 89, 6, 1992, pp. 574-586.
9. S.E. Hussain and Rasheeduzzafar, Corrosion resistance performance of fly ash
blended cement concrete, ACI Materials Journal 91, 3, 1994, pp. 264-272.
10. O.S.B. Al-Amoudi, Durability of reinforced concrete in aggressive sabkha
environments, ACI Materials Journal 92, 3, 1995, pp. 236-245.
11. CIRIA, Guide to Concrete Construction in the Gulf Region, CIRIA Special
Publication 31, 1984.
12. Uniform Building Code, International Conference of Building Officials (ICBO),
California, 1991.
13. K. Treadway, Corrosion period, in: Corrosion of Steel in Concrete, RILEM Report of
Technical Committee 60 CSC, P. Schiessl Ed., Chapman & Hall, London, 1988.
14. W.G. Hime, The corrosion of steel: random thoughts and wishful thinking, Concrete
International 15, 10, 1993, pp. 54-57.
15. O.E. Gjorv and O. Vennesland, Diffusion of chloride ions from seawater into
concrete, Cement and Concrete Research 9, 1979, pp. 229-238.
16. O.E. Gjorv, K. Tan and M. Zhang, Diffusivity of chlorides from seawater into highstrength lightweight concrete, ACI Materials Journal 91, 5, 1994, pp. 447-452.
17. O.E. Gjorv, Effect of condensed silica fume on steel corrosion in concrete, ACI
Materials Journal 92, 6, 1995, pp. 591-598.
18. J. Geiseler, H. Kollo and E. Lang, Influence of blast furnace cements on durability of
concrete structures, ACI Materials Journal 92, 3, 1995, pp. 252-257.
19. J.M. Gaidis and A.M. Rosenberg, The inhibition of chloride-induced corrosion in
reinforced concrete by calcium nitrite, Cement, Concrete and Aggregates 9, 1, pp. 3033.
20. T.C. Powers, A discussion of cement hydration in relation to the curing of concrete,
Proceedings, Highway Research Board, Vol. 27, 1947, pp. 178-188.
21. K.O. Kjellsen, R.J. Detwiler and O.E. Gjorv, Development of microstructure in plain
cement pastes hydrated at different temperatures, Cement and Concrete Research 21,
1991, pp. 179-189.
22. R.J. Detwiler, C.A. Fapohunda and J. Natale, Use of supplementary cementing
materials to increase the resistance to chloride ion penetration of concretes cured at
elevated temperatures, ACI Materials Journal 91, 1, 1994, pp. 63-66.
23. H.H. Patel, C.H. Bland and A.B. Poole, The microstructure of concrete cured at
elevated temperatures, Cement and Concrete Research 25, 1995, pp. 485-490.
24. ACI Committee 305, Hot Weather Concreting, American Concrete Institute, Detroit,
1991.
369
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
THE EFFECT OF THE TYPE OF COPPER ON ITS CORROSION
BEHAVIOR IN KUWAIT’S SOFT TAP WATER
H.M. Shalaby1 and F.M. Al-Kharafi2
1
Kuwait Institute for Scientific Research, P.O. Box 24885, 13109-Safat, Kuwait.
2
Kuwait University, Faculty of Science, P.O. Box 5969, 13060-Safat, Kuwait.
ABSTRACT
This paper presents the results obtained from a study of the corrosion behavior of annealed, halfhard and hard-drawn copper pipes in Kuwait’s soft tap water at room temperature. Accelerated
electrochemical experiments and long-term immersion tests were used to evaluate the performance of
the copper pipes under stagnant and flow conditions. The free corrosion potentials reached during the
long-term immersion tests were close to each other regardless of the tempering state of the copper.
The copper pipes, however, suffered from mild general corrosion which was more apparent in the
hard-drawn material.
This result was further confirmed using inductively-coupled plasma
spectrometry (ICPS). Potential-step polarization experiments indicated that the hard-drawn copper
exhibited the highest anodic current densities and the least noble breakdown potential among the
investigated tempering states. At low frequency, the impedance of the as-received half-hard copper
was lower than that of the annealed and hard-drawn copper. However, the impedance measurements
showed deviations from capacitive behavior when pitting occurred on polished copper. The half-hard
copper exhibited the largest pits in terms of depth and diameter, followed by the annealed and harddrawn copper. The differences observed were explained in terms of the nature of the oxide film, the
tempering state of the copper and its purity.
Key Words: Copper pipes, tap water, general corrosion, pitting corrosion, polarization
behavior, impedance.
INTRODUCTION
Generally, copper plumbing tubes have a service life of several decades. In some waters,
however, pitting can occur. Pitting of copper was first recognized in the UK in the late 1940s
when annealed and half-hard tubes were reported to be perforated in cold, fresh, hard water
within one or two years of service [1]. The pits were hemispherical in shape and contained
cuprous chloride and oxide. They were covered with copper carbonate corrosion product.
This type of pitting has been termed type 1 in the literature. Susceptibility to this type of
pitting has been attributed to the formation of a carbonaceous film during annealing.
In the late 1960s, pitting of another type (termed type 2) was identified in Sweden [2] on
hard-drawn tubes exposed to hot, soft water with pH values of 5-7. The pits were deep and
narrow in shape. They were covered with basic copper sulfate corrosion product. According
to case studies performed in Japan [3], the susceptibility to type 2 pitting is highly dependent
on water quality, but is seemingly independent of the presence of carbon films.
371
Corrosion in the Building Industry
A case study of the failure of copper tubes in the soft tap water of Kuwait revealed
characteristics which were different from those of types 1 and 2 [4]. The morphological
features of the pits were similar to those of type 1, while the water composition and the
corrosion products were those of type 2. Moreover, pitting occurred in cold and hot water
irrespective of the tempering state of the copper. These discrepancies promoted the present
authors to evaluate the corrosion behavior of annealed, half-hard, and hard-drawn copper
pipes in Kuwait's soft tap water.
EXPERIMENTAL PROCEDURE
The study was conducted on seamless copper water pipes made to ASTM B88-83a, type
L (annealed and hard-drawn) and to BS 2871 Part 1, type Y (half-hard). The pipes made to
ASTM B88-83a measured 28.57 mm in outer diameter and 1.27 mm in wall thickness, while
the pipes made to BS 2871 measured 28 mm in outer diameter and 1.2 mm in thickness.
Microscopic examinations of the as-received copper pipes revealed attack of grain boundaries
in the annealed copper; presence of a thick and porous film on the half-hard copper; and a
rough surface with longitudinal grooves on the hard-drawn copper (Fig. 1). The amount of
carbon present on the internal surfaces of the as-manufactured pipes was found to be 15.5,
6.75 and 3.15 mg/dm2 for the annealed, half-hard and hard-drawn pipes, respectively.
The pipes were tested under stagnant and flow conditions in laboratory tap water at room
temperature. Table 1 shows the chemical composition of the water used. Pipe sections,
measuring 10 cm each, were used in their as-received condition to study the corrosion
behavior of copper in stagnant tap water. The pipes were tested in horizontal position and
were connected with PVC pipes to form a U shape. The pipes were welded on the outside to
copper wires for measurement of the open-circuit potential and for conducting direct current
(DC) potential-step polarization experiments. At the end of the immersion period, the tap
water was chemically analyzed for its copper content using inductively-coupled plasma
spectrometry (ICPS).
A test rig was used to measure the open-circuit potentials of the as-received copper pipes
in flowing tap water. The rig was in the form of flow loops made of PVC pipes. The rig was
composed of a fiber glass tank (160 l capacity), a circulation pump, and six horizontal copper
pipes measuring 30 cm in length each. The pipes were connected in parallel with common
water inlet and outlet. The water tank contained a ball valve for maintaining the water level.
The time of water flow in the rig was automatically controlled using a timer switch. The
timer switch was adjusted to allow water to flow for 0.5 hours followed by a period of 5.5
hours of stagnation. The water flowed to waste after passing through the tubes. The flow
velocity in the rig was adjusted to 0.4 and 1.0 m/s using valves connected to the inlet and
outlet of the rig. The details of the test rig have been described elsewhere [5]
The potential measurements were performed against a saturated calomel electrode (SCE)
for a period of about 180 days. In the potential-step polarization tests, the pipes were left in
the stagnant water until their potential reached a quasi steady state before the application of
potential. This usually took a period of about 2 months. Polarization started from the
corrosion potential in the anodic direction in steps of 20 mV. The current was recorded after
10 minutes at each step of polarization. A scanning potentiostat and a strip-chart
potentiometric recorder were used, whereas graphite rods served as counter electrodes.
372
Shalaby and Al-Kharafi
(a)
(b)
(c)
Figure 1.
SEM micrographs showing the surface appearance the as-received
copper before testing: (a) annealed copper; (b) half-hard copper; and
(c) hard-drawn copper
DC potentiostatic polarization and alternating current (AC) electrochemical impedance
measurements were also made on copper specimens in the form of discs, measuring 2.4 cm in
diameter. These tests were carried out on as-received and polished copper in 1-l glass cells
containing stagnant water. Polishing was achieved with diamond paste of grades 15 and 6 μ
m. The specimens were mounted in holders made of Teflon so that only one side (1.5 cm2 in
area) was exposed to the water. Again, a SCE served as the reference, whereas platinum
sheet was used as the counter electrode. The potentiostatic polarization tests were made at a
single potential value of 200 mV for 4 and 6 days. The impedance measurements were taken
after separate potential measurements and DC potentiostatic polarization experiments. In
both cases, the potential of the specimens was left to reach a quasi steady state before the
measurements were taken at an amplitude of 10 mV.
Table 1. Chemical Composition of Kuwait's Soft Tap Water
pH
Conductivity (μScm-1)
M alkalinity (CaCO3, mgl-1)
HCO3- (mgl-1)
Total hardness (MgCO3, mgl-1)
Ca hardness (CaCO3, mgl-1)
Mg hardness (MgCO3, mgl-1)
Cl- (mgl-1)
7.6
540.0
30.0
37.0
135.3
86.3
49.0
80.2
373
Corrosion in the Building Industry
SO42- (mgl-1)
SiO2 (mgl-1)
HCO3-/SO42-
118.0
2.1
0.31
After the termination of the tests, the specimens were left to dry in a desiccator. Optical
and scanning electron microscopy (SEM) were used to examine the surface. The depth and
diameter of the pits were measured using a micrometer attached to the optical microscope.
RESULTS AND DISCUSSION
Potential and Polarization Behaviors
Figure 2 shows the potential-time behavior of as-received copper pipes of different
tempering states during testing in stagnant tap water. It can be seen from the figure that the
potential was initially negative, but rapidly shifted to the positive direction within the first
few days of testing. It is apparent from the figure that the potentials became more or less
steady (with occasional fluctuations) after about 100 days of testing. The potentials reached
were close to each other regardless of the tempering state, being on the order of 5-15 mV.
The copper concentrations in the stagnant water were 0.7, 0.8 and 1.5 mg/l for the as-received
annealed, half-hard and hard-drawn copper, respectively. It is worth noting that the
concentration of copper in the water was higher for the hard-drawn material.
The same potential-time behavior as above was obtained during testing under a
flow/stagnation sequence at 0.4 and 1.0 m/s. Again, the tempering state of the metal did not
significantly affect the potential values at the steady state. However, the increase in flow
velocity appeared to slightly shift the steady state potentials to more positive values.
Optical microscopic examinations of the copper pipes used in the above measurements
revealed that the flow/stagnation sequence caused more precipitation of corrosion products
than full stagnation. The copper pipes tested in stagnant water experienced very little
corrosion, mostly in the form of bluish-green stains. On the other hand, a significant amount
of bluish-green corrosion products were observed on the internal surfaces of the copper pipes
tested under flow/stagnation sequence of tap water. The precipitated corrosion products were
more pronounced on the lower halves of the pipes. They covered a large area of the surface
and were not in the form of mounds as is usually observed during pitting of copper. The
amount of precipitated corrosion products was observed to be more in the hard-drawn pipes
than in the half-hard and annealed ones.
374
Shalaby and Al-Kharafi
Figure 2. Potential-vs-time curves for as-received copper pipes tested in stagnant tap water
Figure 3 shows potential-vs-current curves obtained during potential-step polarization
tests carried out on as-received copper pipes in stagnant tap water. The hard-drawn copper
exhibited the highest anodic current densities and the least noble breakdown potential among
the investigated tempering states. The hard-drawn copper was followed by the annealed and
half-hard copper in terms of the noble shift in the breakdown potential. Optical examinations
showed the presence of scattered and agglomerated small cubic crystals on the surfaces of the
annealed and hard-drawn materials. The oxide layer that was observed on the as-received
half-hard material before the start of the test became less porous and more diffuse after the
test, suggesting that the generated current was consumed in repairing the highly defective
oxide layer. No pitting corrosion was found in the as-received materials.
The present results suggest that hard-drawn copper is somewhat more prone to corrosion
than half-hard or annealed copper in Kuwait's tap water. Since general corrosion was only
observed during the testing period, the attained results do not necessarily imply that harddrawn tubes are more susceptible to pitting than the other tubes in the soft tap water of
Kuwait. In the case of copper pipes tested in hard waters, Devroey and Depommier presented
evidence to show that hard-drawn tubes were more resistant to pitting than half-hard tubes,
which were in turn more resistant than annealed tubes (unpublished data). Cornwell et al. [6]
attributed these differences to carbon deposits on the surface of the annealed tubes which
were formed during bright annealing operations. The superior performance of the hard-drawn
tubes, on the other hand, was attributed to a protective effect exerted by the drawing lubricant
in the pores of the hard-drawn tubes. In the case of soft water, such as in Kuwait, Sato et al.
[3] indicated that the susceptibility to pitting is highly independent of the presence of carbon
film. In such cases, it is possible that the stored energy within the metal due to the
manufacturing process of the hard-drawn tubes, becomes an activating factor for corrosion.
375
Corrosion in the Building Industry
Figure 3. Potential-step anodic polarization curves for as-received copper pipes
tested in stagnant tap water
Lucey [7] and Rossum [8] stated that failures of copper pipes usually occur at the bottom
of horizontal copper pipes because only rarely should debris cling to vertical or steep
surfaces. The presence of a large amount of corrosion products on the bottom surfaces of the
copper pipes may pose a danger by creating differential aeration or concentration cells at
longer exposure duration, leading to the pitting type of corrosion attack.
An interesting finding in this work is that a relatively well defined breakdown potential
is exhibited in the potential-step polarization diagrams of as-received copper (Fig. 3). This
breakdown potential is usually termed critical pitting potential to signify the onset of pitting.
Critical pitting potentials have been reported for copper in different waters [6,9,10]. In our
study, the as-received copper suffered from general corrosion. Thus, the breakdown potential
of copper cannot always be considered to be a critical pitting potential as in the case of a
stainless steel immersed in a chloride-containing environment.
Impedance Response
The impedance response of as-received annealed, half-hard, and hard-drawn copper is
given in Fig. 4. At low frequencies, the impedance of the annealed material is higher than
that of the hard-drawn, and in order, of the half-hard copper. This result clearly expresses the
nature of the oxide films formed on these materials. The low impedance of the half-hard
copper is, therefore, possibly due to the presence of a thick, porous crystalline scale.
Figure 4. Bode-plots for as-received copper of different tempering states in
stagnant tap water.
376
Shalaby and Al-Kharafi
An attempt was made to study the kinetic behavior of the pitting of copper using
impedance spectroscopy. Figure 5 shows the impedance spectra in Bode format obtained for
polished annealed copper after free corrosion and potentiostatic polarization in stagnant tap
water at 200 mV for 4 and 6 days. The overall behavior of the impedance-vs-frequency
curves was similar in all cases with the exception of a slight change for the specimen tested
under free corrosion condition in the frequency range of 1 < f < 1000 Hz. On the other hand,
a significant change is noted in the phase angle-vs-frequency curves. Potentistatic
polarization appeared to cause a decrease in the phase angle peak which is associated with a
shift towards higher frequencies. Furthermore, the low frequency part of the spectra changed
in a manner which suggested that pitting had occurred. A comparison of Figs. 4 and 5
indicates that when the air-formed oxide film was removed by polishing, the impedance of
the material decreased significantly.
The AC impedance technique has been found to be successful in investigating the
kinetics of general corrosion [11]. Recently, however, some investigators [12,13] have shown
that the technique can be used to detect the initiation and growth of pitting. The present
results also indicate that the impedance technique can be a useful tool for the study of pitting
of copper in tap water. The results showed that a dominant capacitive behavior existed in the
case of polished copper immersed in tap water under free corrosion conditions. This suggests
that the surface of copper is to some extent blocked by the surface film under these
conditions. On the other hand, deviations from the capacitive behavior are noted in the case
of polarized copper specimens (Fig. 5). These deviations can signify the presence of pitting.
They are possibly caused by a diffusion process through the porous oxide scale which was
elevated from the metal surface.
Figure 5. Bode-plots for polished annealed copper tested in stagnant tap water
after different testing conditions.
377
Corrosion in the Building Industry
Pit Depth and Diameter
SEM examination of polished copper specimens which underwent potentiostatic
polarization in tap water at 200 mV for 4 days showed a reddish-brown scale covering the
surface regardless of the tempering condition of the tested specimens. The scale was totally
disbonded from the metal surface by an average distance of about 8.5 μm. When the
corrosion products were removed with a soft tissue, numerous pits were found in the metal
underneath the scale. Figure 6 shows the percentage of pits against the pit depth or diameter
for the different copper after potentiostatic polarization in tap water at 200 mV for 4 days. It
can be clearly seen from Fig. 6 that the half-hard copper exhibited the largest depth and
diameter of pits when compared with the annealed and hard-drawn copper. The majority of
the pits in the hard-drawn copper had about the same depth as in the annealed. On the other
hand, the pit diameter was clearly smaller in the hard-drawn copper than in the annealed, as
can be seen in Fig. 6b. In fact, the average pit depths were 0.022, 0.059 and 0.021 mm, while
the corresponding average pit diameters was 0.092, 0.126 and 0.05 mm for the annealed, halfhard and hard-drawn copper, respectively. The number of pits after 4 days of polarization
was almost the same (7-10 pits/cm2) for all tempering states. This number remained almost
unchanged when the potentiostatic polarization time was increased from 4 to 6 days.
The average pit depth for the annealed and half-hard copper after 6 days of polarization
was 0.035 and 0.065 mm, respectively. These results indicate an increase in average depth
for the annealed copper of about 60% and for the half-hard copper of 10%. However,
increasing the polarization time increased the maximum pit depth of about 20% of the pits in
the annealed copper to 0.05 mm (a 25% increase in the maximum depth) and of 17% of the
pits in the half-hard to 0.17 mm (a 90% increase in the maximum depth).
(a)
378
(b)
Shalaby and Al-Kharafi
Figure 6. Percentages of pits of: (a) different depths, and (b) diameters, grown on
polished annealed, half-hard and hard-drawn copper during
potentiostatic polarization in tap water at 200 mV for 4 days.
The fact that the hard-drawn material exhibited smaller diameter pits is readily
explainable in terms of the large deformation of the grains. On the other hand, it is rather
difficult to explain the behavior of the half-hard copper. The differences in behavior could
possibly be due to differences in the alloying elements and not the carbon content. The purity
of the phosphorus de-oxidized annealed and hard-drawn copper (high residual phosphorus)
was 99.9%, whereas the purity of the half-hard copper was 99.85. Pipes made of the former
only contain phosphorus and silver as residual elements, whereas pipes made of the latter
contain phosphorus, arsenic, antimony, bismuth, iron, lead, etc. The present work also
showed that the number of pits did not increase with increases in the polarization time,
suggesting that the number of pitting sites is more or less the same regardless of the
differences in manufacturing procedure and carbon content. These results lend support to the
observation made by Sato et al. [3] that pitting corrosion of copper in soft tap water appears
to be independent of the presence of a carbon film on the metal surface.
In a previous work [9], the present authors showed that pitting corrosion of copper starts
at the grain boundaries. Since there are no data available in the scientific literature with
regard to the diameter or depth of pits developed in copper in tap water, the authors assumed
that the cross-section of a pit grown along the grain boundaries of hard-drawn copper would
appear deep and narrow. Thus, the authors postulated that the shape of the pits depends on
the deformation of the grains. The present results provided the needed experimental proof
that the shape of the pits depends on the tempering state of the copper and not on the type of
water as was assumed previously [14].
For copper, which is rarely passivated, the mechanisms for pitting and general corrosion
seem to be different than those of passivable metals, such as stainless steels. In the presence
of the oxide film formed during manufacturing, as-received copper was found to experience
corrosion of a general nature while under anodic polarization. Thus, the protective oxide film
on copper functions not by making the corrosion rate minimal as in the case of passivable
metals, but by scattering the dissolution process to become through large amount of pores. In
the case of the polished material, a reddish-brown scale was formed. This scale separated
from the metal surface, creating the occluded cell required for pitting initiation. Thus, the
disbondment of the oxide film creates the requirements necessary for pitting corrosion to
occur.
CONCLUSIONS
1. During long-term immersion tests, the free corrosion potentials of the as-received
annealed, half-hard and hard-drawn copper were close to each other regardless of the
tempering state of copper. The copper pipes, however, suffered from mild general
corrosion which was more apparent in the hard-drawn material. This result was
further confirmed using ICPS.
2. At low frequencies, the impedance of the as-received half-hard copper was lower than
that of the annealed and hard-drawn copper. When pitting occurred on polished
379
Corrosion in the Building Industry
copper, the impedance measurements showed deviations from the capacitive behavior.
These deviations were attributed to a mass transport phenomenon through the porous
oxide scale. The AC impedance technique was found to be useful in studying the
pitting of copper.
3. The half-hard copper exhibited the largest depth and diameter of pits followed by the
annealed and hard-drawn copper. The number of pits was about the same (7-10
pits/cm2) for all types of copper and remained unchanged with increases in
polarization time. The difference in the size of pits was ascribed to the degree of
purity of the copper.
ACKNOWLEDGMENT
The authors of this paper would like to acknowledge the Kuwait Foundation for the
Advancement of Science for its financial support of this work through contract project No.
87-08-06.
REFERENCES
1. H.S. Campbell, Journal Institute of Metals 77, 1950, p. 345.
2. E. Mattsson and A.M. Fredriksson, British Corrosion Journal 3, 1968, p. 246.
3. S. Sato, T. Minamoto, K. Seki, H. Yamamoto, Y. Takizawa, S. Okada, S. Yamauchi,
Y. Hisamatsu, I. Suzuki, T. Fujii, T. Kodama, H. Baba and K. Nawara, Proceedings
International Symposium on Corrosion of Copper and Copper Alloys in Building,
Tokyo, Japan, 1982, p. 17.
4. F. Al-Kharafi, H.M. Shalaby and V.K. Gouda; Proceedings 10th International
Congress on Metallic Corrosion, Madras, India, November 1987, p. 767.
5. F.M. Al-Kharafi and H.M. Shalaby, Corrosion 51, 1995, p. 469.
6. F.G. Cornwell, G. Wildsmith and P.T. Gilbert, British Corrosion Journal 8, 1973, p.
202.
7. V.F. Lucey, British Corrosion Journal 2, 1967, p. 175.
8. J.R. Rossum, Journal American Water Association 77, 1985, p. 70.
9. H.M. Shalaby, F.M. Al-Kharafi and V.K. Gouda, Corrosion 45, 1989, p. 536.
10. M. Pourbaix, Corrosion 25, 1969, p. 267.
11. D.C. Silverman and J.E. Carrico, Corrosion 44, 1988, p. 280.
12. M. Keddam and R. Oltra, Materials Science Forum 8, 1986, p. 167.
13. F. Mansfeld and H. Shih, Journal Electrochemical Society 135, 1988, p. 1,171.
14. E. Mattsson, Corrosion Australasia 6, 1981, p. 4.
380
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
CORROSION BEHAVIOR OF VANADIUM IN AQUEOUS SOLUTIONS
W.A. Badawy, F.M. AI-Kharafi and M.H. Fath-Allah
Department of Chemistry, Faculty of Science
Kuwait University, P.O. Box 5969 Safat, 13060 Kuwait
ABSTRACT
Vanadium is an important transition metal used in various industrial applications; especially as a
structural material in metallurgical applications. The electrochemical behavior of vanadium is an
important subject due to its application in redox flow batteries. The behavior of the metal in aqueous
solutions of different pH in oxygen-saturated and oxygen-free electrolytes was investigated. The
behavior of vanadium electrodes in alkaline solutions (pH > 8) differed from its behavior in acidic or
neutral solutions (pH 1-8). Open-circuit potential measurements revealed that the steady-state
potential was a linear function of the solution’s pH. The slope of the linear relation changed from 0.039 V/decade for solutions of pH < 8 to -0.058 V/decade for solutions of pH > 8.
Electrechemical impedance spectroscopy, which represent an effective method for studying
corrosion phenomena, has shown that the metal undergoes active dissolution in aqueous media. In
solutions of pH > 8, the corrosion behavior of the metal can be simulated to a Randles equivalent
circuit model. In oxygen-saturated solutions, the electrode’s surface is covered with a thin oxide film.
The interaction of this film with the ambient electrolyte depends on the solution’s pH. Polarization
measurements have shown that the rate of corrosion in acidic solutions is not affected by the
prevailing gas. In alkaline solutions, the removal of air or oxygen from the ambient electrolyte leads
to a decrease in the corrosion rate.
Key Words: Corrosion, electrochemistry, impedance, passivity, polarization, vanadium
INTRODUCTION
Vanadium is an important transition metal due to its use and the use of its alloys as
structural materials not only in metallurgical applications but also in nuclear reactors [l]. The
metal ions are used extensively in redox flow batteries [2-5]. Unlike many transition metals,
vanadium shows active behavior [6-10]. The electrode potential of vanadium and the effect
of oxygen on this potential were the subjects of the very early studies concerning this metal
[11,12]. The active behavior of vanadium and its anodic dissolution in aqueous solutions was
investigated [13,14]. It was found that the rate of metal dissolution is independent of the
hydrogen ion concentration or the nature of the anions present. The rate determining step was
+
a monovalent adsorbed intermediate (V ads) [14]. Investigations of the anodic behavior in
acidic solutions containing different anions and cations have shown that the metal undergoes
2+
active dissolution in all acidic solutions except those containing Ba at a pH > 4, and that the
active dissolution of the metal occurs through a monovalent intermediate [15]. The behavior
of the metal in both acidic and basic media was found to obey the Tafel approximation of the
Butler-Volmer equation over a wide range of potentials [16]. Extensive studies have been
383
Fundamental Aspects
carried out in glacial acetic acid and have shown oxide film formation in the presence of
sodium borate and water [17,18,19].
The present investigation was aimed at throwing more light on the electrochemical
behavior of vanadium in aqueous solutions and on the effect of solution pH and oxygen on
the corrosion and passivation processes occurring at the electrode/electrolyte interface.
EXPERIMENTAL PROCEDURE
Massive, cylindrical, spectroscopically pure, vanadium rods (Alderich-Chemie) were
used as working electrodes; they were mounted in glass tubes of appropriate internal diameter
2
with an epoxy resin leaving a front surface area of 0.302 cm to contact the electrolyte. An
all glass three electrode cell with a large surface area, Pt, counter electrode and a Ag/AgCl/Cl
(3 M KCl) reference electrode was used. The solutions were prepared from analytical grade
reagents and triple distilled water. The buffer solutions covering a pH range of 1-13 were
prepared according to Clark and Lub's series [20]. The pH of each solution was controlled
before each experiment, and the electrodes were mechanically polished with successive
grades of emery papers down to 1200 grit, then wiped with a smooth cloth and washed with
triple distilled water. In this way, the electrodes acquired reproducible, bright silvery
surfaces. After polishing and rinsing, the electrodes were immersed directly in the test
solutions. Electrochemical impedance spectroscopy (EIS) investigations were carried out
using the IM5d-AMOS system (Zahner Elektrik GmbH & Co., Kronach, Germany). The
5
input signal was usually 10 mV peak to peak in the frequency domain of 0.1-10 Hz.
Frequencies down to 0.1 MHz were also investigated. Polarization measurements were
performed using an EG&G (Princeton Applied Research) model 273A
Potentiostat/Galvanostat
interfaced to an IBM PS3 computer.
The potentiostatic
measurements were traced using programs that enable ohmic drop compensation. The
steady-state potential, ESS, measurements were controlled separately using a high impedance
value voltmeter (Keithley type 197A Autoranging Multiplier, England). All measurements
were carried out at constant room temperature of 22oC. The potentials were measured against
the Ag/AgCl/Cl- reference, and then refered to the normal hydrogen electrode (nhe). The
gases used for deaeration or solution saturation were purified and dried before being bubbled
in the electrolytic cell. The gas was bubbled for at least 20 min. in the test solution prior to
each experiment. The details of the experimental procedures were as described elsewhere
[10,21].
RESULTS AND DISCUSSION
Steady-State Potential Measurements
The potential of the vanadium electrode was traced over a period of 180 min. in naturally
aerated aqueous solutions covering a pH range from 1 to 13. In all solutions, the electrode
potential became more positive with time. A typical example of the variation of the electrode
potential with immersion time in the buffer solutions is presented in Fig. 1A. In this figure,
the results in solutions of pH’s 2, 7 and 12 are presented. The steady-state potential, Ess ,
measured after 180 min. of electrode immersion in each solution is plotted against the pH of
the solution, and is presented in Fig. 1B. The results reveal that ESS is pH-dependent over the
384
Badawy et al.
whole pH range. The Ess versus pH relation is linear and can be presented by an empirical
equation of the form:
Ess = a - b pH
(1)
where a is a constant representing the value of ESS extrapolated from the linear relation at pH
= 0. As can be seen from Fig. 1B, the ESS versus pH relation has an inflection between pH =
8 and pH = 9 which means that the slope, b, changes. In the basic medium, i.e., pH > 8, the
slope of the linear relation is very close to the value of 0.059 V/pH, which was calculated
from the Nernst equation for a pH indicator electrode with one electron electrode process at
o
25 C according to:
o
E = E - 0.059 pH
n
(2)
o
E is the standard pH-independent electrode potential and n is the number of electrons
involved in the electrochemical process. Therefore, the value of n for basic solutions (pH >
8) is equal to 1. In acidic and neutral solutions (pH < 8), on the other hand, the slope of the
Ess versus pH relation is < 0.059 V/pH. A value of 0.039 V/pH was calculated. This can be
+
explained by the interaction of both H and OH ions with the electrode’s surface which is
essentially covered with a passive film or an adsorbed O2 film as will be discussed later. In
regions where the OH ions are inaccessible to the electrode surface, solvation of the
electrode occurs, and deviation from the n = 1 process is observed [22]. The vanadium
electrode can be considered to be a pH indicator electrode taking into account the inflection
in the ESS versus pH relation and the values of a and b in each pH in the range 1-8
Ess = 0.180 - 0.039 pH
(3)
Ess = 0.380 - 0.058 pH
(4)
and in solutions of pH’s > 8
The steady-state potential values obtained according to Eqs. 3 and 4 are quite different
2+
from the standard electrode potentials assigned for the systems V/V (Eo = -1.19 V) and
2+ 3+
o
V /V (E = - 0.26 V) [22]. This supports the notion that the electrode process cannot be
represented by a simple equilibrium relationship such as that given for the simple redox
2+
2+ 3+
equilibria of the metal and its ions, e.g., the V/V or V /V Systems. This can be
understood on the basis that the solution does not contain significant concentrations of those
ionic species, and that the electrode’s surface is covered by a thin oxide film in aqueous
solutions. Consistent with this is the dependence of the steady-state potential on the
prevailing gas and the stirring conditions of the solution.
Effect of Oxygen on the Steady-State Potential
385
Fundamental Aspects
The electrode’s potential was traced in naturally aerated, oxygen-saturated and oxygenfree solutions of different pHs. Oxygen was removed from the solution by bubbling N2 or H2
at least 20 min. before electrode immersion. A typical example of the results in solutions of
pH = 12 is presented in Fig. 2. The results reveal that the electrode’s potential is sensitive to
the oxygen concentration in the solution. Under all conditions, the potential became more
positive with immersion time until it reached a steady-state value. The use of either N2 or H2
to remove oxygen from the solution did not produce any remarkable difference; the steadystate potential lay, in both cases, in approximately the same range (≈ -385 mV (nhe) in pH =
12). In oxygen-saturated solutions, the Ess shifted in the positive direction by ≈ 100 mV.
This shift can be attributed to the presence of a thin oxide film on the electrode’s surface.
The same trend was observed over the whole pH range from 1 to 13.
Figure 1. (A) Variation of the electrode potential Figure 2.
with time of the vanadium electrodes in
naturally aerated solutions of different pHs
(o) pH = 2 (∗) pH = 7
(Δ) pH = 12
(B) Steady-state potential (Ess) vs. pH for the
vanadium electrode in naturally aerated
buffer solutions
Effect of the prevailing gas on
the electrode potential of
vanadium in solutions of pH 12
(o) naturally aerated
(∗) O2-saturated
(Δ) N2-saturated
(Δ) H2-saturated
The sensitivity of metals with active/passive transitions towards oxygen is well known,
especially for those which do not show active dissolution like niobium, tantalum and titanium
[10,21,24]. The passivation of vanadium in aqueous solutions was discussed very early by
386
Badawy et al.
Muthman and Frauenberger [11]. Later, it was suggested that the passivity is due to the
presence of a gaseous film [12]. Unlike many transition metals, vanadium has an active
corrosion behavior with a limited tendency for passivation. The passivation behavior of the
metal was explained earlier by the presence of a chemisorbed oxygen film [25]. Dry or moist
o
oxygen did not tarnish the polished metal’s surface, and even hot water (60-85 C) had no
effect on the surface’s brightness. The decrease of the steady-state potential on oxygen
removal by bubbling of N2 or H2 in the test solution can be considered as an indication of the
presence of a thin oxide film on the electrode’s surface. The oxide film formed represents a
non-stoichiometric oxidation state of the metal which is responsible for the observed behavior
of the metal in each solution. Removal of oxygen by bubbling of N2 or H2 in the solution
leads to dissolution of the formed oxide, and hence, a decrease of its thickness leading to a
shift of the steady-state potential in the negative direction (Fig. 2), and the following
equilibrium state is shifted to the left:
2V + xO2 → 2VOx
(5)
The value of x determines the stoichiometric factor n of Eq. 2.
In acidic solutions, where an excess of hydrogen ions are present, the interaction
between the non-stoichiometric oxide film and the hydrogen ion takes place and the
electrode’s potential is determined by the hydrogen ion concentration according to:
+
-
V-Ox + 2x H + 2xe → V + x H2O
(6)
The value of x in this case, and hence, the value of n of Eq. 2, can be calculated from the
slope of the first segment of the steady-state potential/pH relation (Fig. 1B), i.e., in the pH
range of 1-8.
slope = -0.039 V = - 0.059
n
where n = 2x i.e., n = 1.5
In basic solutions, the interaction between the electrode’s surface and the solution occurs
through OH ions according to:
-
V- Ox + 2OH → VO1+X + H2O + 2 e
-
(7)
The lower oxides of vanadium are basic and very unstable [25]; therefore, they cannot
protect the metal from corrosion as in the case of the valve metals with active/passive
transitions.
Open-Circuit Impedance Measurements
EIS is a powerful tool for investigating electrochemical and corrosion systems, since it is
essentially a steady-state technique that is capable of accessing relaxation phenomena with
387
Fundamental Aspects
relaxation times that vary over several orders of magnitude and permits single averaging,
within a single experiment to obtain highly precise levels. The open-circuit impedance of
vanadium electrodes was traced for 180 min. after electrode immersion in the test solutions.
Typical data for pH’s of 2, 7 and 12 are presented as Bode plots in Fig. 3.
(A) in solution of pH 2
(B) in solution of pH 7
(C) in solution of pH 12
Figure 3. Bode plots of the vanadium electrodes in naturally aerated solutions at different
time intervals from electrode immersion
(⎯) 15 min. (…) 60 min. (----) 130 min.
Bode plots are recommended as standard impedance plots since the phase angle, θ, is a
sensitive parameter for indicating the presence of additional time constants in the impedance
spectra [10,26-28]. It employs frequency as an independent variable, so that a more precise
comparison between experimental and calculated impedance spectra can be made [29-31].
The use of the log versus log format enables equal representation of all experimental data
over the whole frequency domain.
388
Badawy et al.
The EIS spectra in Fig. 3 contain only one capacitive contribution represented by the
linear variation of the electrode’s impedance, Z, with the frequency, f, [26-28]. In acidic and
neutral solutions (pH = 1-8), there is a part of the spectrum where the phase angle is
independent of frequency (Fig. 3a and b at f < 1 Hz). Such behavior is explained by the
assumption of frequency dispersion and surface inhomogeneity [29]. For such systems, the
electrode impedance is given by:
Rp
Z = ⎯⎯⎯⎯⎯
α
I + (sCRP)
(8)
where RP is the polarization resistance which is considered to be a pure charge transfer
resistance, C is the electrode capacitance and α is a fit parameter ( 0 < α < 1 ) that is
correlated to the angle of rotation of the center of the capacitive semicircle, φ, below the real
axis:
φ = (1- α) π/2
s = j ϖ where j =
(9)
−1 and ω = 2 π f.
The value of the fit exponent α corresponds to the extent of dispersion and is attributed
to surface inhomogeneity [29,30]. A nonlinear concentration of metal ions will occur, since
preferential charge transfer takes place at active sites. The impedance spectra of Fig. 3 show
that there is an active dissolution of the metal, as can be identified by the decrease of the
polarization resistance with immersion time in each solution.
The rate of dissolution is limited and is dependent of the solution’s pH. The impedance
spectra of the electrodes in solutions of pHs of 2, 7 and 12 after 180 rnin. of electrode
immersion are collectively presented in Fig. 4. The results show clearly that the behavior of
vanadium in basic solutions is different from its behavior in acidic or neutral solutions. In
solutions of pH 12, a phase maximum at a frequency of ≈1 Hz was observed. To be sure that
there was no second time constant at lower frequencies, impedance measurements down to
frequencies of 1 MHz were taken. A typical example of these measurements for solutions of
pH 12 is presented in Fig. 5. The behavior of the electrode in basic solutions is very similar
to an ideally corroding system with a single time constant corresponding to the corrosion
reaction rate determining step. This behavior can be simulated by a simple equivalent circuit
model of the Randles type which consists of a parallel combination of a resistor, RP ,
representing the polarization or charge transfer resistance and a capacitor, C , representing the
capacitance of the electrode/electrolyte interface. This parallel combination is in series to a
small resistor, Rs , equivalent to the electrolyte resistance.
The data in Fig. 5 were subjected to a procedure of data fitting [31] to fit the
experimental data in this figure to the electronic model described. Good agreement was
obtained between the experimental and theoretical data for values of RP = 10.22 kΩ, C =
-2
170.4 μF cm , and Rs = 42 Ω. The fitting procedure for the results of solutions of pH of 12
389
Fundamental Aspects
0
had around 2% mean error in the absolute impedance and ≈1.2 mean error in the phase
angle. Deviations of the phase angle from ideal behavior were found to be related to the
polishing of the electrode’s surface. The impedance data presented show that the vanadium
electrode, although it showed active dissolution, it had a great tendency towards oxygen and
the surface was covered with a thin film of non-stoichiometric oxide or a mixture of oxides of
varying valences. The instability of such oxides explain the corrodibility of vanadium’s
surface. The increased rate of corrosion with increases in the pH of the solution was varied
by measuring the corrosion currents and polarization resistance in each solution. The values
of these parameters in pHs 2, 7 and 12, are presented in Table 1.
Figure 4. Bode plots of vanadium electrode after Figure 5.
130 min. immersion in naturally aerated
solutions of pH 2
(…), pH 7 (----) and pH 12 (⎯)
Bode plot of the vanadium
electrode in naturally aerated
solution of pH = 12 in the
-3
5
frequency range 10 to 10 Hz
Table 1. Values of the Polarization Resistance, RP , Corrosion Current, icorr, and
Corrosion Potential Ecorr, of the Vanadium Electrode in Naturally Areated
Solutions of Different pHs
pH
2
7
12
RP
2
(kΩcm )
5.012
2.880
1.652
icorr
-2
(μAcm )
0.407
1.239
3.930
Ecorr
(mV)
-215
-407
-597
The effect of the prevailing gas on the impedance behavior of the vanadium electrode
was also investigated. An example of these measurements in solutions of pH 12 is presented
in Fig. 6. After 3 hrs of electrode immersion, the impedance behavior in H2 or N2 saturated
solutions is quite similar. In air or oxygen-saturated solutions, the measured polarization
resistance is lower. Polarization measurements have shown that the corrosion current
390
Badawy et al.
decreases in the presence of an inert gas in basic solutions. In acidic solutions, on the other
hand, the corrosion rate is not much affected by the inert gases. These results are summarized
in Table 2. The results of polarization experiments are in good agreement with the
+
explanation based on Eqs. 6 and 7. In acidic solutions, the interaction of the excess H ions
with the surface film leads to the removal of this film without any appreciable change in the
rate of corrosion of the metal by changing the gas. In basic solutions, the formation of the
non-stoichiometric, unstable, basic oxides is responsible for the increased rate of corrosion of
the metal. The formation of the oxide film and its stoichiometry is dependent of the presence
of air or O2 in the solution. Removal of air or oxygen from a basic solution shifts the
equilibrium of Eq. 5 to the left, and hence, decreases the rate of corrosion as can be seen from
the values of the corrosion currents presented in Table 2. The presence of the surface film is
confirmed by capacitance measurements. In all solutions, the electrode capacitance showed
an approximately constant value within a wide potential range (-100 - +100 mV from the
-2
steady-state potential). In acidic solutions, an average capacitance value of 25 μFcm was
measured. This value is higher than the reported value of the Helmholtz capacitance (17
-2
μFcm ) [32]. The higher capacitance value can be attributed to the presence of adsorbed
electroactive species on the surface film. The concentration of these electroactive species at
the electrode’s surface is constant in the potential range where the electrode capacitance is
potential-independent.
Figure 6. Effect of the prevailing gas on the impedance characteristics of the vanadium
electrode in solutions of pH 12
(- -) oxygen, (⎯) nitrogen, (----) naturally aerated, (....) hydrogen
Table 2. Corrosion Currents of Vanadium in Solutions of pHs 2, 7 and 12 Saturated with
Different Gases
Gas
-2
icorr (μA-cm )
2
Rp (kΩcm )
391
Fundamental Aspects
Air
Oxygen
Nitrogen
Hydrogen
pH 2
0.407
1.057
1.578
1.046
pH 7
1.239
2.939
1.967
1.082
pH 12
3.930
17.50
1.710
0.928
pH 12
1.652
1.596
2.352
2.072
Effect of Temperature on the Corrosion Behavior of Vanadium
To study the effect of temperature on the corrosion behavior of vanadium, an all glass,
double walled cell was used with the same arrangement of counter, reference and working
electrodes. The measurements were made in naturally aerated solutions of pHs 2, 7 and 12.
In all solutions, the general trend was an increase in the rate of corrosion with increasing
temperature. Polarization measurements were taken at each temperature, and the
corresponding corrosion current, iCorr , which represents the rate of corrosion, was obtained.
A plot of log icorr versus 1/T obeys the familiar Arrhenius equation [33].
d logicorr = Ea
2
dT
RT
(10)
where Ea is the activation energy which is given by
Ea = NA εa
(11)
εa is the energy relative to the ground state energy which an atom or molecule must have in
order to react, i.e., the activation energy per molecule, whereas Ea is the molar activation
energy. Figure 7 presents the Arrhenius plots obtained in solutions of pH 2, 7 and 12.
In the solutions investigated, almost parallel Arrhenius plots were obtained, which means
that the activation energy of the corrosion process lies in the same range without regard to the
solution’s pH. Calculation of the activation energy of the corrosion process in each solution
gave the values presented in Table 3. The values given in Table 3 show that the activation
-1
energy of the corrosion process is less than 40 kJmol , which supports the view that the
dissolution of the metal is a one-electron charge transfer process [34]. This supports the
mechanism suggested by Armstrong and Henderson [14], and the corrosion reaction may be
presented by
V(s)
⎯
⎯→
V(I) + e
2+
-
(12)
+
-
H2O + V(I)ads. ⎯fast
(13)
⎯
⎯→ VO +2H +3e
This reaction is enhanced in basic media, and hence, the rate of corrosion increases. In
accordance with this, the calculated activation energy in solutions of pH 12 is slightly lower.
Table 3. Activation Energy of the
392
Badawy et al.
Corrosion of Vanadium in
Naturally Aerated Solutions of
pHs 2, 7, and 12
pH
2
7
12
-1
Ea, (kJ mol )
34.8
36.4
30.6
Figure 7. Log icorr vs 1/T relations for the corrosion behavior
of vanadium in naturally aerated solutions of pH 2
(o), pH 7 (∗) and pH 12 (Δ)
CONCLUSIONS
The steady-state potential of vanadium is sensitive to the solution’s pH and can be used
for pH calculations. The rate of corrosion of the metal in basic media decreases with the
removal of air or oxygen. In acidic or neutral solutions, the prevailing gas has no significant
effect on the rate of the corrosion process. Activation energy calculations support a oneelectron transfer step as the rate-determining corrosion process.
ACKNOWLEDGMENT
This work has been supported by Kuwait University, Research Grant No. SCO59. The
financial support of the research administration is gratefully acknowledged.
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394
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
THE EFFECT OF UV-IRRADIATION ON PASSIVE FILMS FORMED ON
TYPE 304 AND 316 STAINLESS STEELS
M.S. Al-Rifaie, C.B. Breslin, D.D. Macdonald and E. Sikora
Center for Advanced Materials, The Pennsylvania State University,
517 Deike Building, University Park, Pennsylvania, 16802, USA
ABSTRACT
The effect of monochromatic ultraviolet (UV) light on the passive films formed on Types 304
and 316 stainless steels (SS) is described. Under UV irradiation 304SS and 316SS specimens, in
neutral and acidic solutions, exhibited an increased resistance to localized corrosion (pitting). This
resistance to localized corrosion was gauged by an increase in induction time, an increase in
breakdown potential, and some significant changes in the current noise at constant potentials. All
these changes indicate that shining UV light on a metal specimen can sometimes decrease its
susceptibility to pitting.
It was observed that the extent of photoinhibition of localized attack (PILA) depends on the
nature of the passive film, the period of illumination, and the incident photon energy. It was also
observed that the PILA effect can, in some cases, last for over 200 hours after the illumination has
been removed. Increased pitting resistance was observed with higher energy incident photons. The
minimum apparent incident photon energy corresponds to the band gap of the metal specimen (375
nm for SS). The optimum illumination period observed in these experiments was approximately 5
hours. In alkaline solutions a much decreased PILA effect was observed, this decrease in PILA was
attributed to the formation of a precipitate layer. This precipitate layer in turn interfered with the
incident photon interaction with the barrier layer. Finally, a possible explanation of PILA is given
within the framework of the Point Defect Model (PDM).
Key Words: Photoinhibition, stainless steel, localized corrosion, pitting, UV, PDM, passive films
INTRODUCTION
The initiation and propagation of pits and the corresponding breakdown of the passive
film that forms on metals and alloys is of great fundamental interest in electrochemistry and
corrosion science. Numerous efforts, ranging from the addition of inhibitors [1] to the
alloying of base metals [2-5], have been made in an attempt to prolong the effective life of the
passive film. Recently, it has been shown that irradiating immersed electrodes with
ultraviolet (UV) light can inhibit localized corrosion. The first observation of this kind was
made when polycrystalline nickel, in chloride-containing solution, was irradiated with white
light [6]. Other observations supporting photoinhibition of localized attack (PILA) include
similar effects on 304 stainless steel (SS) [7] and even pure iron under UV irradiation [8-9].
The purpose of this study is to examine the effects of UV irradiation on the photoinhibition of
304SS and 316SS in chloride-containing solutions. These results could help shed light on
395
Fundamental Aspects
how PILA affects SSs and lay the groundwork for photoinhibition as a new form of corrosion
control.
EXPERIMENTAL PROCEDURE
Test specimens were prepared from Types 304 and 316 SS rods, which were covered
with lacquer, mounted in a PVC holder and then embedded in a two-component epoxy resin.
The exposed surface, approximately 0.8 cm2 in area, was polished mechanically with
successively finer grades of SiC paper and 0.05 μ alumina powder to a mirror finish. The
chemical composition of the SS samples used is shown clearly in Table 1.
Table 1. Chemical Composition of 304 SS and 316 SS (in Wt %)
SS Type
C
304
0.08
316
0.08
Mn
2
2
Cr
18
16
Ni
8
12
P
S
0.04 0.003
0.04 0.003
Si
1
1
Mo
Trace
2
Fe
Bal.
Bal.
The electrochemical cell consisted of a three-electrode PTFE cell equipped with a quartz
window to allow irradiation of test electrodes. A saturated calomel electrode (SCE) was used
as the reference electrode, and a platinum wire, coiled inside the cell, was used as the
auxilliary electrode. All test solutions were prepared from Analar-grade reagents and
deionized water, and were deoxygenated with nitrogen. The pH of the solution was adjusted
to 7.5 with NaOH, or alternatively, buffered to pH 7.5 with a 0.15 mol dm-3 H3BO3/0.007 mol
dm-3 Na2B4O7 solution.
B
The working electrodes were irradiated at wavelengths between 300 and 425 nm using a
150 W UV-enhanced Xe lamp (Oriel Model 6254) and a 1/8 monochromator (Oriel Model
77250). The incident power density at 300 nm was 0.4 mW cm-2, giving a photon flux of
6.04 x 1014 cm-2. The photon flux was maintained at approximately this value at each
wavelength by adjusting the light intensity at the surface.
Electrochemical tests were carried out using a Solartron/Schlumberger Electrochemical
Interface (Model 1286). In potentiodynamic polarization tests, the working electrodes were
polarized at a rate of 0.1 mV s-1 in the anodic direction up to the breakdown potential. In
illumination experiments, the electrodes were illuminated continuously throughout the
potential scan. The breakdown potential was recorded as the potential at which the current
exceeded 80 μA cm-2. In current-time measurements, the electrodes were initially polarized
at a potential in the passive region for a 30-minute period, and then the potential was stepped
to an appropriate point where metastable pitting could be observed for the non-illuminated
specimens. The current transients were then recorded as a function of time, using a Keithley
Model 576 data acquisition unit at a sampling rate of 90 mS.
Additional experiments involved polarizing the working electrodes, under illumination
for periods of up to 15 hours, and then determining the breakdown potential using the
potentiodynamic polarization method. The exact same polarization periods were used for the
illuminated and non-illuminated electrodes.
396
Al-Rifaie et al.
RESULTS AND DISCUSSION
Figure 1 shows typical anodic polarization curves for type 304 SS in a neutral 0.5 mol
dm-3 NaCl solution (unbuffered) under conditions of illumination and non-illumination.
Figure 1 clearly shows that when the sample was illuminated, an increase in pitting resistance
resulted. This increase in pitting resistance was indicated by a shift of both the breakdown
potential and the initial metastable pitting potential towards the more noble direction.
Figure 1. Potentiodynamic polarization curves for type 304 SS in neutral 0.5 mol dm-3 NaCl
under: (a) non-illumination; and (b) illumination at 300 nm
The effect of illumination on the breakdown potentials of 304 SS and 316 SS can
perhaps be seen more clearly from the data shown in Table 2, which shows averages of the
breakdown potentials for 304 SS and 316 SS as a function of chloride concentration. In each
case, an average increase of about 60 ± 40 mV in the breakdown potential can be observed
upon illumination. Another observation that can be made from the data is that the breakdown
potentials of the illuminated 304 SS specimens approach those of the 316 SS specimens in the
dark. This observation could suggest a comparable degree of passivity enhancement between
alloying and illumination under these conditions.
An even greater increase in the breakdown potential, approximately 150 ± 50 mV, was
observed on prior illumination of the specimens at 300 nm for periods exceeding 5 hours. In
these experiments, the electrodes were polarized in a 0.1 mol dm-3, NaCl buffered solution at
+250 mV (SCE) for various periods of time under illumination and non-illumination. The
specimens were then polarized from +250 mV (SCE) in the anodic direction at a rate of 0.1
397
Fundamental Aspects
mV s-1 up to the breakdown potential. The displacement in the breakdown potential was
calculated by subtracting the average breakdown potential for the specimens polarized in the
dark from the average breakdown potential for the specimens polarized in the light. Each
experiment was repeated three times. This data are shown graphically in Fig. 2, where the
average displacement in the breakdown potential, ΔEb is shown as a function of the prior
illumination period.
Table 2. Breakdown Potential Values for 304 SS and 316 SS Under Light and Dark
Conditions
[Cl-] mol dm-3 304 SS Eb (dark) 304 SS Eb (light) 316 SS Eb (dark) 316 SS Eb (light)
0.025
355 ± 7 mV (3)
420 ± 10 mV (3)
430 ± 10 mV (3)
490 ± 11 mV (4)
0.5
275 ± 18 mV (20) 350 ± 20 mV (22) 330 ± 20 mV (17) 395 ± 20 mV (17)
2.0
160 ± 8 mV (4)
210 ± 15 mV (5)
230 ± 12 mV (5)
290 ± 10 mV (4)
Eb in mV vs. SCE
Light = 300 nm
The numbers in parentheses indicate the number of times the experiment was repeated
Figure 2. Displacement in the breakdown potential, ΔEb, as a function of the illumination
period for 316 SS at +250 mV (SCE) in a buffered 0.1 mol dm-3 NaCl solution
398
Al-Rifaie et al.
Further evidence for photoinhibition of pitting attack was obtained from current-time
measurements where the current decay transients were monitored as a function of time for
illuminated (300 nm) and non-illuminated 316 SS. These data are shown in Fig. 3.
It is evident from Fig. 3 that illumination causes a delay in the onset of metastable
pitting, even after approximately 110 minutes, the illuminated film was still able to
repassivate while the dark film had started breaking up after only 25 minutes. Therefore,
illumination seemed to postpone metastable pitting by at least a factor of 3, which in itself is
quite astounding.
One of the questions that this study aims to answer is how prior illumination affects the
passive film on 304 SS and 316SS. In order to quantify the permanent nature of the
photoinhibition effect, the pitting susceptibility of 316SS was studied at various periods of
time with prior illumination. The specimens were illuminated at 300 nm for 80 minutes under
polarizing conditions in a neutral 0.5 mol dm-3 NaCl solution. The specimens were then
immersed under open-circuit conditions (dark) in a borate buffer solution (pH of 7.5), and
removed at selected intervals. The breakdown potential was determined in a neutral 0.5 mol
dm-3 NaCl solution using the potential scan method. Identical experiments were carried out
in the dark; the specimens were polarized in the chloride solution (dark) for 80 minutes,
removed and immersed in the borate solution. The breakdown potential was determined at
selected intervals. Data collected in this manner for periods up to 350 hours are shown in Fig.
4, where the breakdown potential is plotted against the immersion period following
illumination or polarization.
399
Fundamental Aspects
Figure 3. Current-time decay profiles for 316 SS polarized at 285 mV (SCE) in a neutral
0.025 mol dm-3 NaCl solution under: (a) dark, and (b) light (300 nm) conditions
Figure 4. Breakdown potential of illuminated and non-illuminated 316SS in 0.5 mol dm-3
NaCl as a function of the immersion period in a borate buffer solution (dark)
following polarization or polarization and illumination at 300 nm.
(a) on a linear scale; and (b) on a semi-log scale
A clear difference between the breakdown potentials measured for the illuminated and
non-illuminated specimens can be seen for immersion periods up to about 220 hours,
indicating that the photoinhibition effect persists over this period of time. The gradual
increase in breakdown potential (i.e., a shift towards the more noble direction) may be
attributed to a crystallization process or chromium-enrichment in the passive film. A similar
trend was observed for 304SS.
The influence of solution pH, and thus the nature of the passive film, on the extent of
photoinhibition was studied by polarizing and illuminating the electrodes in solutions of
varying acidity. A 0.5 mol dm-3 NaCl solution was used as the test solution; the pH was
adjusted to the desired value by the addition of NaOH or HCl. All irradiation experiments
were carried out at 300 nm. The breakdown potentials were determined from polarization
measurements for specimens polarized in the dark and under conditions of continuous
400
Al-Rifaie et al.
illumination. Each experiment was carried out at least three times. The displacement in the
breakdown potential, ΔEb, was calculated as the difference between the light and dark
breakdown potentials. The average displacements in the breakdown potentials for 304SS and
316SS are shown as a function of pH in Fig. 5; the degree of scatter in the average
displacements was ± 30 mV. An essentially constant increase in the breakdown potentials, of
approximately 60 mV, was observed on illumination, except for those specimens polarized in
the alkaline solutions, where no apparent photoinhibition effect was detected. However, it
was found that the photoinhibition effect was partially restored under these alkaline
conditions (pH of 10) by the addition of a 0.01 mol dm-3 EDTA solution to the test solution.
The pH in this region was adjusted with NaOH and maintained at 10 on addition of the
complexing EDTA agent. The average displacements in the breakdown potentials for both
304SS and 316SS at a pH of 10 on addition of the EDTA were 45 and 50 mV, respectively.
The presence of EDTA at other pH values did not enhance the photoinhibition effect. This
seems to suggest that the precipitated layers formed in alkaline environments are photoelectrochemically inactive, but that the addition of a chelating agent hinders the formation of
this layer, allowing the photons to reach the barrier layer.
Figure 5. Displacement in the breakdown potential, ΔEb, as a function of the pH of a 0.5 mol
dm3 NaCl solution, on illumination of: (a) 304SS; and (b) 316SS at 300 nm
401
Fundamental Aspects
The effects of variations in the photon energy on the breakdown potential displacement
(a measure of photoinhibition) are shown graphically in Fig. 6. A constant photon flux was
maintained at each wavelength. A neutral 0.5 mol dm-3 NaCl solution was used as the test
solution. A total of twenty experiments were carried out for each of the SSs under conditions
of non-illumination in order to obtain adequate reference breakdown potentials. The amount
of scatter in the breakdown potentials, under these conditions, was on the order of ± 20 mV.
The mean value of the breakdown potential calculated for 304SS in the dark was 272 mV
(SCE), while that for 316SS in the dark was 327 mV (SCE). Displacements in the breakdown
potential, ΔEb, on illumination were calculated as the difference between the breakdown
potentials in the light and the dark. It can be seen from Fig. 6 that the degree of
photoinhibition depended on the photon energy, with the photoinhibition effect decreasing
with wavelengths exceeding 375nm.
Figure 6. Displacement in the breakdown potential, ΔEb, measured in a 0.5 mol dm-3 NaCl
solution, as a function of the incident light wavelength on illumination of: (a)
304SS; and (b) 316SS
The induction periods for 316SS specimens polarized at +285 mV (SCE) in a 0.025 mol
dm NaCl solution can be plotted as a function of varying photon energy as shown in Fig. 7.
Figure 7 also shows the induction periods measured for identical experiments carried out in
the dark. These data points are plotted at each wavelength so that the increase in the
induction period on illumination is evident. The induction periods were measured as the time
-3
402
Al-Rifaie et al.
between the application of the polarizing potential and the first metastable pitting events in
which the current exceeded 500 nA. It is clear from this figure that the induction time was
reduced slightly with decreasing photon energy.
Figure 7. Measured induction periods as a function of the incident light wavelength for
316SS polarized at +285 mV (SCE) in 0.025 mol dm-3 NaCl. Induction periods
for non-illuminated specimens are also shown for reference (i.e., dark triangles at
each wavelength)
In previous papers [6,7], PILA has been interpreted using the PDM for the growth and
breakdown of passive films [10] as the photo-quenching of the electric field within the barrier
layer. The PDM has already been developed to give theoretical expressions for the
breakdown potential, Eb, and the induction time, tind. It is proposed that on illumination,
incident photons (with energies in excess of the band gap of the metal) generate electron-hole
pairs that are separated by a steep potential gradient in a manner which quenches the electric
field. Once the electric field is decreased, the theoretical expressions predict a higher
breakdown potential and a larger induction time. Therefore, Figs. 1-7 are indeed consistent
with a PDM interpretation of the photoinhibition effect since they predict that incident photon
energy increases pitting resistance. It is not clear, at least from the experimental evidence
shown here, how the electric field itself is quenched.
The data presented in Fig. 5 suggests that the formation of a precipitate layer on the
metal specimens polarized in the alkaline solutions screens the barrier layer from the incident
403
Fundamental Aspects
photons, and thus inhibits the generation of electron-hole pairs (and subsequent quenching of
the electric field). This is supported by the experiments in which the PILA effect was
partially restored on addition of the complexing agent, EDTA. This complexing agent may
hinder the formation of the precipitate layer, and hence, facilitate the photon-barrier layer
interaction. Also, the fact that the PILA effect remains unaffected (in neutral and acidic
solutions), or increases (in alkaline solutions), in the presence of EDTA suggests that it is
unlikely that photo-induced reactions, involving oxidized iron, at the film/solution interface
account for the photoinhibition effect.
The semipermanent nature of PILA is evident in this study (as shown in Figs. 4 and 7)
and other studies [8,9]. The PILA effect remaining after the irradiation has been removed
seems to indicate that the photoinhibition effect cannot be explained solely by the changing of
the electronic structure of the film. Consequently, it is postulated that the suppression of the
electric field strength also modifies the vacancy distribution. If we assume that the thickness
of the barrier layer (L) is 3 nm, and the cation vacancy diffusivity is on the order of 10-19
cm2s-1 [12], we can calculate an approximate relaxation time (t=L2/D). The relaxation time of
the vacancy structure is thus estimated to be around 9 x 105 s. If we then compare this value
with the period of 220 hours (7.92 x 105 s, from Fig. 4), there is good agreement between the
two values, suggesting that the photoinhibition effect persists until the vacancy structure
relaxes. Figure 2 also supports this idea in the sense that longer illumination periods lead to
an increased PILA effect by altering the vacancy structure more dramatically, and therefore,
requiring a larger relaxation time with increasing illumination time. Another possible
explanation of PILA is that the UV irradiation leads to a chromium enrichment of the passive
film. Since the PILA effect on pure iron [8,9] and nickel [6] cannot be explained by
chromium enrichment of the passive film; a more likely possibility is one in which both the
electric field is quenched and the passive film is enriched in chromium simultaneously.
CONCLUSIONS
The results of this work show that photoinhibition of pitting corrosion can be achieved
for SSs on illumination with UV light. Increases in both the breakdown potential and
induction period, and a decrease in the frequency of metastable pitting events were observed
upon irradiation. It was also found that this PILA effect depended on the photon energy (with
photons having energies above the band gap of the specimen being more effective), the
illumination period, the pH of the test solution, and the nature of the resulting passive film.
It appeared that the precipitate layer formed on passivation of 304SS and 316SS in
alkaline solutions (pH > 10) adversely affected the interaction of the incident photons with
the barrier layer. The addition of EDTA, a complexing agent, partially restored the PILA
effect, which seems to support the hypothesis that the precipitate layer is indeed hindering the
photon-barrier layer interaction. These observations can be explained within the framework
of the PDM in terms of the generation of electron-hole pairs and consequent photo-quenching
of the electric field. This in turn modifies the vacancy structure, leading to an enhancement
in the pitting resistance of the specimens, that remains effective for some 220 hours.
ACKNOWLEDGEMENT
404
Al-Rifaie et al.
The authors gratefully acknowledge the support of this work by the Electric Power
Research Institute under Contract No. RP8041-07, and by the US Department of
Energy/Basic Sciences Division through Grant No. DE-FG02-91ER45461.
REFERENCES
1. M. Ohi, H. Nishihara and K. Aramaki, Corrosion 50, 1994, p. 226.
2. R.G. Wendt, W.C. Moshier, B. Shaw, P. Miller and D.L. Olson, Corrosion 50, 1994, p.
819.
3. B.A. Shaw, G.D. Davis, T.L. Fritz, B.J. Rees and W.C. Moshier, Journal of the
Electrochemical Society 138, 1991, p. 3288.
4. A.J. Sedriks, Corrosion of Stainless Steels, The Electrochemical Society, Princeton, New
Jersey, 1979.
5. P.M. Natishan, E. McCafferty and G.K. Hubler, Journal of the Electrochemical Society
135, 1988, p. 321.
6. S.J. Lenhart, M. Urquidi-Macdonald and D.D. Macdonald, Electrochimical Acta 32,
1987, p. 1739.
7. E. Sikora, M.W. Balmas, D.D. Macdonald and R.C. Alkira, Corrosion Science, in press.
8. P. Schmuki and H. Bohni, Journal of the Electrochemical Society 139, 1992, p. 1908.
9. P. Schmuki and H. Bohni, Electrochimical Acta 40, 1995, p. 775.
10. D.D. Macdonald, Journal of the Electrochemical Society 139, 1992, p. 3434.
405
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
KINETICS OF HIGH TEMPERATURE CORROSION OF A LOW Cr-Mo STEEL IN
AQUEOUS NaCl SOLUTION
W.A. Ghanem1, F.M. Bayyoumi1 and B.G. Ateya2
1
2
Central Metallurgical Research and Development Institute, Helwan, Egypt.
Corresponding Author Department of Chemistry, Faculty of Science, Cairo University,
Cairo, Egypt.
ABSTRACT
The kinetics of corrosion of a low Cr-Mo steel alloy were studied over a temperature range of 75250 C in 1 m NaCl in the absence and in presence of various levels of contamination with CuCl2. We
measured corrosion rates, weights of corrosion product (magnetite) film and total (integral) weight
loss of the alloy over exposure times of 1-480 hours. The corrosion rate decreased rapidly with time,
before it leveled off at longer time periods, indicating the formation of a protective corrosion product
film. The ability of the alloy to retain an adherent corrosion product (magnetite) film was expressed in
terms of a retention coefficient. This increased with temperature and exposure time, and decreased
with the level of contamination with CuCl2. The effect of temperature was attributed to the
improvement of the crystallinity of the corrosion product. On the other hand, the effect of the CuCl2
was attributed to the electro-deposition of Cu and its impregnation within the corrosion product, which
became less adherent. The free corrosion potential was affected by the presence of the CuCl2 in a
fashion compatible with the Wagner-Traud theory of mixed potential.
0
Key Words: Kinetics, corrosion, steel, high temperature, sodium chloride solution, copper chloride.
INTRODUCTION
The corrosion of steel in high temperature aqueous environments is encountered in many
industrial applications, e.g., boiling water reactors [1], desalination plants [2], high
temperature aqueous fuel cells [3], and steam generators [4]. In such environments,
contaminants, which are present in the aqueous media at trace levels, are concentrated by
several orders of magnitude to form highly corrosive solutions [5]. Due to the active nature
of iron, it corrodes in high temperature deaerated water and steam giving rise to the formation
of ferrous (Fe2+) species which change to ferrous hydroxide, Fe(OH)2, and magnetite, Fe3O4
[6-10]: i.e.,
3 Fe + 6 H2O = 3Fe(OH)2 + 3 H2
(1)
3 Fe(OH)2 = Fe3O4 + H2O + H2
(2)
The overall reaction is represented by:
3 Fe + 4 H2O = Fe3O4 + 4 H2
(3)
407
Fundamental Aspects
Some of the resulting magnetite adheres to the surface in the form of a film which affects the
kinetics of any subsequent corrosion of the steel. The rest of the resulting magnetite spalls off
the surface into the electrolyte. The qualities of the adherent magnetite film depend on the
temperature, composition of the environment and exposure time.
The objectives of this paper are to study the kinetics of corrosion of a low Cr-Mo steel in
high temperature NaCl solution, and the mechanism of protective film formation during this
process. Particular attention is given to the effect of contamination of the electrolyte with
CuCl2 on the integrity of the protective film, and hence, on the subsequent corrosion. The
effects of CuCl2 concentration, temperature and exposure time on the adherence of the
magnetite film were also determined.
EXPERIMENTAL PROCEDURE
All measurements were performed in an autoclave fabricated from 316 type stainless
steel. The autoclave consisted of two parts. A (Teflon) PTFE cell was machined to fit tightly
into the autoclave, to accommodate the electrolyte. Further details are given elsewhere [11].
The cell was preheated for about 5 hours to obtain thermal stability [12,13]. The autoclave
was placed in a furnace connected with the temperature regulator to the heating source. A NiCr thermocouple was used to regulate the temperature. It was inserted into a stainless steel
tube coated with a layer of PTFE. The corrosion rate measurements were taken on coupons
(2 x 5 x 0.2 cm) made of a low Cr-Mo steel of the following composition (wt%): 2.3 Cr; 1.0
Mo, 0.46 Mn, 0.2 Si, 0.13 C, 0.015 P, 0.007 S and Fe balance. They were annealed at 900oC
for 1 hour in an argon atmosphere and furnace-cooled. Their microstructure revealed fine
dispersed carbide in a matrix of ferrite. Before use, they were polished successively down to
600 grit using silicon carbide paper, rinsed with ethyl alcohol and distilled water, and then
dried in the air. Three electrolytes were used: (I) 1 molal (m) NaCl, (II) 0.999 m NaCl + 5 x
10-4 m CuCl2 and (III) 0.9 m NaCl + 5 x 10-2 m CuCl2. They were deaerated by boiling the
electrolyte under reflux for 15 miutes to give an oxygen content of < 0.1 ppm (as estimated
polarographically). The volume of the electrolyte in each test was 200 ml. The tests lasted
for various durations, i.e., 1, 3, 6, 12, 24, 48, 96 and 480 hours and were performed in
triplicate. Two of the three specimens were subjected to successive descaling by immersing
for 20 minutes in a 20% ammonium citrate solution at 80oC [14] to dissolve the corrosion
products. They were then rinsed with distilled water, dried and weighed until a constant
weight was obtained. We calculated the integral weight (ΔWi ) i.e., the weight of the alloy
which dissolved up to a particular time, and the weight of the corrosion product film (ΔWf )
which remained adherent per square centimeter of the area after a particular time. After each
test, the solution was found to contain some solid (spalled) corrosion products.
The working electrode consisted of a rectangular sheet about 9 cm long, 0.5 cm wide
and 0.2 cm thick. It was insulated with PTFE in such a way that an area of 1 cm2 was
exposed at its end. The other end was threaded and connected to a stainless steel rod of 0.3
cm diameter through a stainless steel connector. A graphite rod of 0.5 cm diameter and about
10 cm length was used as a counter electrode, and an Ag/AgCl was used as a reference
electrode [12,15].
RESULTS AND DISCUSSION
408
Ghanem et al.
Corrosion Rate
Figure 1 (a-c) illustrates the variation of the corrosion rate with time of immersion at
various temperatures in electrolyte I, electrolyte II and electrolyte III. Figure 1d compares the
behavior in the three electrolytes at 2500C. They clearly reveal same significant features.
During short time periods, the corrosion rate decreased rapidly with the time of immersion
before it tended to level off after longer time periods. This behavior is characteristic of
protective film formation [16]. There was a strong detrimental effect of CuCl2 on the ability
of the magnetite film to protect the substrate alloy. As the temperature and/or exposure time
increased, the corrosion rate decreased.
The present work reveals that the mechanism of corrosion changes after a transition time,
τ, the magnitude of which, generally, decreases as the temperature of the test increases. At
and beyond this transition time, an adherent layer of the corrosion product was shown to
protect the substrate alloy by acting as a diffusion barrier [11,17], thus reducing the rate of
corrosion. Before this transition time, the alloy corrodes more freely with a higher rate of
corrosion. It was found that, at a given temperature, increasing the concentration of CuCl2
increased the transition time, τ [17].
Retention Coefficient
The retention coefficient is introduced here to give a quantitative expression of the
ability of the alloy to retain an adherent corrosion product film on its surface under a
corrosive environment. It is defined as the ratio of the weight of the adherent (magnetite)
film, ΔWf (adh.), at a particular time of immersion to the total weight of the (magnetite) film
which would form if the integral weight loss of the alloy were to be totally consumed in
forming the film material (magnetite), i.e., ΔWf (total). The later value is related to the integral
weight loss ΔWi by a chemical factor (CF), which in the present case is given by the ratio of
the molecular weights of Fe3O4 and 3 Fe i.e., CF = 232/168 = 1.38. Thus, the retention
coefficient is given by
Retention coefficient ϕ = ΔWf (adh.) /ΔWf
(total)
(4)
The retention coefficient was determined at various temperatures, CuCl2 concentrations and
time intervals. Figure 2 (a-d) illustrates the variation of the retention coefficient, ϕ , with the
time of immersion for the three electrolytes at various temperatures. The curves clearly
reveal that ϕ increased as the temperature increased, and decreased as the concentration of
CuCl2 increased. Reaction 2, which is called the Schikorr reaction [18], has been extensively
studied [19,20]. Ferrous hydroxide, Fe(OH)2, decomposes rapidly above 1000C [21] , but
relatively slowly at lower temperatures. Robertson [22] stated that the corrosion of steel in
hot water is controlled by the dehydration of the hydroxide phase (Eq. 2), which proceeds
when the metal/solution interface becomes saturated with Fe(OH)2. Consequently, two
factors affecting reaction 2,
• The saturation of Fe(OH)2 at the metal/solution interface which is time
dependent, and
• The temperature, which enhances the reaction in the forward direction and
affects the solubility [23] and crystallinity of the magnetite film [17].
409
Fundamental Aspects
This explains the higher ϕ values obtained at higher temperatures, that longer exposure times,
and hence, the increased efficiency of the film in retarding corrosion. On the other hand, the
presence of CuCl2 in the electrolyte decreased ϕ . This has previously been shown [17] by xray diffraction and scanning electron microscopy to be due to the electro-deposition of Cu
and its impregnation within the magnetite film, which then becomes less adherent.
Figure 1. Variation of the corrosion rate with the time of immersion
at different temperatures in: (a) sol. I, 1 m NaCl; (b) sol. II, 0.999 m
NaCl + 0.0005 m CuCl2; (c) sol. III, 0.9 m NaCl + 0.05 m CuCl2.; and
(d) at 250oC in different solutions
410
Ghanem et al.
Figure 2. Variation of the retention coefficient, ϕ , with the time of immersion
at different temperatures in: (a) sol. I, 1 m NaCl; (b) sol. II, 0.999 m
NaCl + 0.0005 m CuCl2; (c) sol. III, 0.9 m NaCl + 0.05 m CuCl2; and
(d) at 250oC in different solutions
Comparing the results in Figs. 1 and 2, it can be concluded that, in most cases, as the
retention coefficient decreases, the corrosion rate increases. In other cases, both the retention
coefficient and the corrosion rate increase in the same direction. This result indicates that the
film formed, though retained on the alloy surface, is unable to protect it from subsequent
corrosion.
Potential-Time Curves
Figure 3 illustrates the time variation of the free corrosion potential, Ecor , of the alloy at
different temperatures in electrolytes I, II and III. It is seen that Ecor shifts toward the noble
direction as the concentration of CuCl2 increases. This is in agreement with the results of Lin
et al [23]. The increase in temperature above 750C shifted the values of Ecor in electrolytes I
and II closer to each other than they were at 750C. In electrolyte III, increasing the
temperature shifted the free corrosion potential to more noble values. A comparison of these
free corrosion potential values with the equilibrium potentials of the hydrogen evolution
(H2O/H2) and copper reduction (Cu/Cu++) reactions is in order to identify the cathodic half
reactions. Table 1 lists the values of the equilibrium potentials of both systems at various
temperatures in electrolytes I, II and III. Note that these Ecor values are considerably negative
(cathodic) with respect to the reversible equilibrium potentials of the Cu/Cu++ or the H2O/H2
electrode systems [6]. The approximate values for Cu/Cu++ system in electrolytes II and III
are calculated at various temperatures using the Nernst equation, i.e.,
Cu2+ + 2e
Cu
E = E0 Cu/CuCl2 + 2.303 RT/2F log [Cu2+]
(5)
(6)
The values of E0 were obtained from Latimer [24]; the activity of the Cu2+ species was taken
equal to its concentration. The values of E (H2O/H2) at different temperatures were taken
from Pourbaix diagrams [6, 25,26].
Consequently, under the potentials shown in Fig. 4, the cathodic half cell reaction
involves both the reduction of water i.e., reaction 7
2 H2O + 2 e →
H2 + 2 OH-
(7)
and the electro-deposition of Cu according to reaction 5, while the anodic reaction involves
the dissolution of the iron, i.e., reaction 8
Fe + H2O
→
Fe(OH)+ + H+ + 2e
(8)
Since the concentration of Cu2+ is rather small in electrolyte II, the time behavior of Ecor is not
significantly different from that in electrolyte I at the higher temperatures i.e., 125, 175 and
411
Fundamental Aspects
2500C. Alternatively, in presence of higher concentration of CuCl2, the rate of reaction 5 is
greatly enhanced leading to an increase in the corrosion rate.
The results of Figure 3 can be explained within the domain of the Wagner-Traud theory
of mixed potential [27,28], shown schematically in Fig. 4, which illustrates the effect of a
significant increase in the rate of the cathodic reaction on the corrosion rate (Icor) and the
corrosion potential (Ecor). For the sake of simplicity, we neglect the changes in the anodic
polarization curves of reaction 5 brought about by adding CuCl2. Upon changing the cathodic
half cell reaction from reaction 7 to reaction 5, Fig. 4 shows a significant increase in the
corrosion current, Icor , and a significant shift in the mixed (free corrosion) potential towards
more noble values. Both phenomena were confirmed by the experimental measurements
shown in Figs. 1 and 3.
412
Ghanem et al.
Figure 3. Time variation of the free corrosion potential for the alloy at
different temperatures in: electrolyte I (1 m NaCl); II (0.999 m NaCl +
0.0005 m CuCl2), and III (0.9 m NaCl + 0.05 m CuCl2)
Table 1. Approximate Values of the Electrode Potential, V (NHE) of the Cu/Cu2+
Calculated at Various Temperatures Using the Nernst Equation in
Electrolytes II and III and of the H2O/H2 Systems in Electrolytes I, II and
III [6, 25,26]
Temperature
750C
1250C
1750C
2500C
Electrolyte I
H2O/H+
- 0.442
- 0.551
Electrolyte II
Cu/Cu2+
H2O/H+
+ 0.312
- 0.405
+ 0.309
+ 0.305
+ 0.299
- 0.484
Electrolyte III
Cu/Cu2+
H2O/H+
+ 0.329
- 0.361
+ 0.327
+ 0.326
+ 0.325
- 0.415
Figure 4. Schematic representation of the effect of CuCl2 on the
mixed potential of iron in the corrosive medium
CONCLUSIONS
Inspection of the results presented reveals the following conclusions:
413
Fundamental Aspects
1.
At short times, the corrosion rate decreases rapidly with the time of immersion before
it tends to level off at longer times. This behavior is characteristic of protective film
formation
2.
The retention coefficient is introduced to give a quantitative expression for the ability
of the alloy to produce an adherent corrosion product film under the corrosive
environment. Increasing the temperature enhances the ability of the alloy’s surface
to retain the film. The concentration of CuCl2 has an opposite effect.
3.
The results of potential-time curves measured in different electrolytes reveal that Ecor
shifts in the noble direction to an extent that increases with the concentration of
CuCl2. This is compatible with the Wagner-Traud theory of mixed potential. The
increase in temperature above 750C shifts the values of Ecor in electrolyte I and II
closer to each other than they are at 750C.
ACKNOWLEDGMENT
The authors express their warm gratitude to Prof. A.A. Abdul Azim, the former chairman
of CMRDI, for valuable discussions.
REFERENCES
1. B.C. Syrett, Materials Performance 30, 8, 1992, p. 52.
2. O. Osborn and F.H. Coley, in High Temperature High Pressure Electrochemistry in
Aqueous Solutions, Houston, NACE 4, 1976, p. 7.
3. F.T. Bacon, in High Temperature High Pressure Electrochemistry in Aqueous
Solutions, Houston, NACE 4, 1976, p. 24
4. M.J. Wootten, G. Economy, A.R. Pebler and W.T. Linsay. Jr., Materials Performance
17, 2, 1978, p.30.
5. C.B. Ashmore, M.H. Hurdus, A.P. Mead, P.J B. Silver, L.Tomlinson and D.J.
Finnigan, Corrosion 44, 1988, p. 334.
6. M. Pourbaix, in Atlas of Electrochemical Equilibria in Aqueous Solutions, London,
Pergamon Press, 1974, p. 305.
7. J.E. Castle and G.M.W. Mann, Corrosion Science 6, 1966, p. 253.
8. J. Robertson, Corrosion Science 29, 1989, p. 1275.
9. J. Jelinek, P. Neufield, Corrosion 38, 1982, p. 98.
10. G. Butler, H.C.K. Ison and A.D. Mercer, British Corrosion Journal 6, 1971, p. 23.
11. F.M. Bayyoumi, M.Sc. Thesis, Cairo University, 1995.
12. M.H. Lietzke, R.S. Greeley, W.T. Smith and R.W. Stoughton, Journal of Physical
Chemistry 64, 1960, p. 652.
13. D.D.G. Jones and H.G. Masterson, in Advances in Corrosion Science and
Technology, Vol. 1, New York, Plenum Press, 1970, p. 1.
14. F.A. Champion, in Corrosion Testing Procedures, London, Chapman and Hall, 1964,
p. 192.
15. D.D. Macdonald, A.C. Scott and P. Wentrcek, Journal of the Electrochemical Society
126, 1979, p. 908
16. U.R. Evans, in The Corrosion and Oxidation of Metals: Scientific Principles and
Practical Applications, New York, St. Martin’s Press , 1960, p. 819.
17. W.A. Ghanem, F.M. Bayyoumi and B.G. Ateya, Corrosion Science, in press, 1996.
414
Ghanem et al.
18. G. Schikorr, Z. Anorg. Allg. Chem. 212, 1933, p. 533.
19. U. R. Evans and J.N. Wanklyn, Nature 162, 1948, p. 27.
20. B. McEnaney and D.C. Smith., Corrosion Science 18, 1978, p. 591.
21. F.J. Shipko and D.L. Douglas, Journal of Physical Chemistry 60, 1956, p. 1519.
22. J. Robertson, Corrosion Science 32, 1991, p. 443.
23. C.C. Lin, F.R. Smith, N. Ichikawa and M. Itow, Corrosion 48, 1992, p. 16.
24. W.M. Latimer, in The Oxidation States of the Elements and Their Potentials in
Aqueous Solutions, 2nd ed., New York, Prentice Hall, 1961.
25. H.E. Townsend, Jr., Corrosion Science 10, 1970, p. 343.
26. V. Ashworth and P.J. Boden, Corrosion Science, 10, 1970, p. 709.
27. C. Wagner and W. Traud, Z. Elektrochem. 44, 1938, p. 391.
28. D.D. Macdonald, Corrosion 48, 1992, p. 194.
415
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
CORROSION AND PASSIVATION BEHAVIOUR OF ALUMINIUM AND
ALUMINIUM ALLOYS
MECHANISM OF THE CORROSION PROCESS
F.M. Al-Kharafi, W.A. Badawy and A.S. El-Azab
Department of Chemistry, Faculty of Science
Kuwait University, P.O. Box 5969 Safat, 13060 Kuwait
ABSTRACT
Aluminum and Al-alloys represent technologically and industrially important materials. The
electrochemical behavior of these materials in different solutions represents a major subject of
investigation. The corrosion characteristics of naturally passivated Al, Al-Cu, Al-6061 and Al-7075
were studied in nitric acid and nitric acid containing chloride solutions. The effect of the
concentration of anions on the corrosion behavior of these materials was traced.
Electrochemical impedance spectroscopy (EIS) is a powerful tool in studying corrosion and
passivation problems. Besides polarization techniques, the method has been applied successfully to
investigate the corrosion behavior of Al and Al-alloys. The Al-6061 alloy was found to be the most
corrosion resistant. In all cases, the naturally occurring passive film was too thin to impart complete
passivity. Equilibrium occurred between barrier film dissolution and surface passivation especially in
dilute solutions (< 0.1 M HNO3). The electrode/electrolyte interface was fitted to a parallel
resistor/capacitor combination. The barrier film formed on Al or Al-6061 behaved like a perfect
dielectric whereas that formed on Al-Cu or Al-7075 alloys deviated from the ideal capacitor behavior.
X-ray photoelectron spectroscopy (XPS) experiments have shown that Al-Cu alloys contain
remarkable amounts of Cu on the material surface. Scanning electron microscopy (SEM)
investigations have shown that the presence of Cu on the alloy surface initiates flawed regions which
are responsible for the increased corrosion rate of the Cu-containing alloys.
Key Words: Aluminium, aIuminium alloys, corrosion, electrochemistry, impedance, passivation
INTRODUCTION
Due to the technological importance of Al and Al-Alloys and their increased industrial
applications, the electrochemical behavior of these materials represents an important subject
for many investigators [1-10]. Investigations have been conducted to optimize the anodic
polarization and passive film growth [11-15]. The effect of anions like Cl and the
mechanism of their attack on the metallic surface has been a major research subject
[5,10,16,17]. The corrosion and passivation behavior of aluminum and its alloys has been
subjected to intensive investigation [18-21]. The use of nitric acid and nitric acid/phosphoric
acid mixtures in the surface finishing of aluminum and aluminum alloys, especially in the
household industry, required detailed information about the electrochemical behavior of these
materials in this medium [17,22-25].
417
Fundamental Aspects
Several techniques have been used to study the corrosion and passivation behavior of
metals and alloys. Electrochemical impedance spectroscopy (EIS) is becoming a well
established method for investigating electrochemical systems in which the solid/electrolyte
interface plays the main role [26]. One of its important aspects is the direct matching that
often exists between experimental impedance data and data obtained from discrete electrical
components which represent physical processes taking place in the system under investigation
[27]. Such electronic components are usually termed electronic models or equivalent circuits.
EIS and other polarization techniques are very useful in studying the corrosion and
passivation behavior of Al and its alloys. The polarization resistance, RP , which represents
the major indication of kinetic facility [28], and electrode capacitance, C , which represents
the main source for calculating the barrier layer thickness and its dielectric properties [29] are
the main factors used to describe and control the corrosion and passivation of these materials
[22-25].
In the present investigation, the electrochemical behavior of mechanically polished,
naturally passivated Al, Al-Cu, Al-6061 and Al-7075 was investigated in nitric acid and nitric
acid containing chloride solutions. An electronic model for the barrier layer/electrolyte
interface was described. The effect of chloride ion concentration on the corrosion and
passivation behavior of the metal and its alloys in nitric acid solutions was studied. X-ray
photoelectron spectroscopy (XPS) and scanning electron microscopy (SEM) were used to
investigate the material’s surface.
EXPERIMENTAL PROCEDURE
Commercial-grade aluminum and aluminum alloys (Al-Cu, 6061 and 7075) were used as
electrodes. The mass spectroscopic analysis of these materials is presented in Table 1.
Electrodes in the form of cylindrical rods were mounted into glass tubes of appropriate
internal diameter with an epoxy resin leaving an exposed surface area of 0.50 , 0.21 , 0.20 and
0.21 cm2 for Al, Al-Cu, Al-6061 and Al-7075, respectively, to contact the test solution. The
electrolytic cell was an all glass, three electrode cell with a large surface area Pt counter
electrode and Ag/AgCl/Cl- (3 M KCl) reference electrode. The electrolytic solutions were
prepared using analytical grade reagents and triply distilled water. All measurements were
carried out at a constant room temperature of 25oC. The potentials were measured against the
Ag/AgCl/Cl- (3M KCI) reference {Eo = 0.1970 V(nhe)}. Before each experiment, the
electrode was mechanically polished with successive grades of emery paper down to 3/0, and
then with a smooth cloth and washed with triply distilled water. In this way, the electrode’s
surface acquired a reproducibly bright appearance. For comparison, some experiments were
carried out after chemical etching of the electrode’s surface to be sure that the mechanical
polishing had no effect on the alloy’s structure. The electrodes were chemically etched in a
80oC heated mixture of phosphoric, acetic and nitric acids for 5 minutes [22,23]. The
impedance data obtained for both mechanically polished and chemically etched electrodes
showed almost the same trend with slightly higher impedance values (5-10% higher) for the
chemically etched surface at different time intervals of electrode immersion in the test
solution.
EIS measurements were performed using the IM5d-AMOS system (Zahner Elektric
GmbH & Co., Kronach, Germany).
All experiments involved single frequency
measurements in the frequency domain of 0.1-105 Hz. To check the presence of another time
418
Al-Kharafi et al.
constant at lower frequencies, some experiments were conducted over a bandwidth of 1 mHz
- 105 Hz. The input signal’s amplitude was usually 10 mV peak to peak. Polarization
measurements were carried out using an EG&G (Princeton Applied Research) Model 273A
potentiostat/galvanostat interfaced to an IBM PS/3 computer. The XPS experiments were
carried out using an ESCA-Lab 200 (VG instruments). The surface was etched as required by
argon ion bombardment. In each spectrum, the XPS peaks of C 1S, O 1S, Al 2P and Cu 2P1
and 2P3 were traced. The electrode’s surface was examined by SEM before and after
immersion in the test solution. The details of experimental procedures were as described
elsewhere [22-25].
Table1. Mass Spectrometric Analysis of the Different Electrode Materials in Mass %
Alloy
Al
Cu
Mg
Al
Al6061
Al-Cu
Al7075
99.23
97.09
93.43
90.93
0.043
0.201
4.80
1.17
0.217
1.40
0.229
2.21
Si
0.038
0.601
0.047
0.272
Fe
Mn
Ni
Zn
pb
Sn
Ti
Cr
0.164
0.193
0.499
0.124
0.001
0.012
0.024
0.067
0.010
0.010
0.012
0.007
0.027
0.029
0.025
4.95
0.001
0.000
0.721
0.000
0.003
0.000
0.006
0.000
0.006
0.016
0.015
0.024
0.001
0.248
0.001
0.046
RESULTS AND DISCUSSION
Corrosion Behaviour in Nitric Acid Solution, Equivalent Circuit for the
Electrode/Electrolyte Interface
The impedance behavior of the different electrodes was investigated in 0.1 M HNO3.
The mechanically polished electrodes were left in 0.1 M HNO3 until a steady state was
reached, and then the impedance data were recorded. Although Bode plots for impedance
data presentation are always recommended as standard impedance plots [26], they sometimes
lead to no indication of features hidden at high frequency [10]. In such cases, the Nyquist
plot format is more favorable. For data fitting procedures, Bode plots are always used since
all experimental data are equally represented, and the phase angle is very sensitive for
indicating the presence of additional time constants in the impedance spectra. In our
experiments, both formats were used.
Typical Nyquist plots of Al-7075 alloy electrodes taken at different time intervals from
the steady state are presented in Fig. 1. Bode plots as a function of immersion time in the test
electrolyte of Al-6061 are presented in Fig. 2. For all electrodes, the impedance Nyquist plot
at any time interval consists of two semicircles (two phase maxima in the Bode plot). A high
frequency semicircle, which is due to the interaction between the electrode surface and the
electrolyte, is associated with a high field conduction mechanism through the oxide film and
its thickness, and a low frequency loop which is concerned with the relaxation processes
occurring in the barrier layer either in the bulk or at the surface, which is typical of passivated
surfaces. Below the assigned low frequency of the experiment (i.e., 0.1 Hz), no reproducible
data could be obtained. In the very low frequency range (0.1-100 mHz), the structural
changes of the interfacial region were faster than the measurements and no reliable data could
be obtained [30-33]. The diameter of the high frequency semicircle changed with the time of
immersion in the electrolyte. For all alloy electrodes, the diameter decreased with immersion
time which reflects a decrease in the polarization resistance of the barrier layer, Rp, and its
419
Fundamental Aspects
thickness, δ. Pure Al electrodes showed continuous increases of diameter with immersion
time in nitric acid solutions (Fig. 3). The increase in the diameter with time indicates oxide
film thickening which means continuous passivation of Al in nitric acid solutions as was
observed before [22]. The barrier layer thickness was calculated from the impedance data
according to
C = I/ 2 π f Zim
(1)
C = A ε ε0 / δ
(2)
where C is the electrode’s capacitance, f is the frequency, Zim is the electrode’s impedance, A
is the electrode’s area, ε0 is the permittivity of free space (8.85 x 10-14 Fcm-l), δ is the barrier
layer thickness and ε is the oxide film dielectric constant taken as 8.4 [33] considering that
the barrier layer consists mainly of A1-203. The calculated barrier layer thickness after a
long period of electrode immersion (≈ 4 hours) in the test solution ranges between 0.2 and 0.6
nm for all electrodes irrespective the barrier layer thickening or thinning that occurred at the
electrode/electrolyte interface. This thickness is about one-tenth of the thickness of the
barrier layer occurring on Al or Al alloys in neutral solutions (pH = 7) [6,34]. The presence
of such a barrier layer on Al or its alloys after long immersion times in nitric acid solution (≈
4 hours in 0.1 M solution) indicates the remarkable passive behavior of these materials in
these electrolytes.
Nyquist plots of Al-7075 electrode
at different time intervals of
immersion in 0.1 M HNO3 (steady
state potential = -258 mV vs.
Ag/AgCl/Cl- [3 M KCl])
(⎯) 45 min (steady state), (…) 75
min, (---) 135 min, (- - -) 250 min.
420
Bode plots of Al-6061 electrode at
different time intervals of
immersion in 0.1 M HNO3 (steady
state potential = -532 mV vs.
Ag/AgCl/Cl- [3 M KCl])
(⎯) 45 min (steady state), (…) 75
min, (---) 135 min, (- - -) 250 min.
Al-Kharafi et al.
The polarization and impedance data are in good agreement. Figure 4 presents the Tafel
polarization curves of the four materials investigated after reaching the steady state in 0.1 M
HNO3. The values of the polarization resistance, Rp , corrosion current, icorr , and steady state
potential, Ecorr , for the different electrodes as obtained from these measurements are
presented in Table 2. Figure 5 presents the impedance data of the electrodes after 250
minutes of electrode immersion under the same conditions. Taking into consideration the
polarization resistance and corrosion current data given in Table 2, the stability order of the
investigated materials after reaching the steady state in 0.1 M HNO3 (i.e., 45 minutes of
electrode immersion) follows the sequence:
Al-6061 > Al > Al-7075 > Al-Cu
After a long immersion time in the test solution (i.e., 4 hours), Al attained of comparable
stability or became even more passive than the Al-6061 alloy and the order changed to:
Al > Al-6061 > Al-7075 > Al-Cu
Nyquist plots of aluminium
electrodes at different time
intervals of immersion in 0.1 M
HNO3 (steady state potential = 650 mV vs. Ag/AgCl/Cl- [3 M
KCl]) (⎯) 45 min (steady state),
(…) 75 min, (---) 135 min, (- - -)
250 min.
Tafel polarization curves of Al-Cu
(1), Al-7075 (2), Al-6061 (3) and
Al (4) after 45 min of electrode
immersion in 0.1 M HNO3
This is clearly reflected in the Nyquist plots of Fig. 5. The change in the order of
stability of Al and Al-6061 after long immersion times in the nitric acid solution is due to the
421
Fundamental Aspects
observed passivation of aluminum from the moment of immersion in nitric acid (Fig. 3). The
polarization resistance of the Al-electrode increased from 230 Ωcm2 after 45 minutes of
electrode immersion in 0.1 M HNO3 solution at 25oC, to 372 Ωcm2 after 250 minutes of
electrode immersion in the same solution under the same conditions.
Table 2. Values of RP , icorr and Ecorr (vs. Ag/AgCl/3M Cl-) for the Different Electrodes
Measured in 0.1 M HNO3 at the Steady State ( ≈ 45 min from electrode immersion)
Electrode
Al-6061
Al
Al-7075
Al-Cu
2
Rp, (Ω cm )
502
374
276
269
icorr, (μA cm2)
113.5
121.2
169.6
196.1
Ecorr, (mV)
-500
-778
-271
-111
The data presented in this section show that Al-6061, either from the moment of
immersion in nitric acid or after a long period of immersion, represents the most stable alloy
in this solution of the alloys investigated. The results reveal that the presence of the small
amount of Mg (1.4%) improves the passivation behavior of the aluminum alloy. The mass
spectrometric investigation of the alloy showed that it contains 1.40% Mg and 0.60% Si.
Such a combination in a heat-treatable wrought alloy leads to the formation of a Mg2Si phase,
which is the basis for precipitation hardening. Either in solid solution or as submicroscopic
precipitate, Mg2Si has a negligible effect on electrode potential. The alloy is normally used
in a heat-treated form; therefore, no detrimental effects derive from the major alloying
element or from the minor components like Cr and/or Zn which are usually added to control
the grain structure. Copper additions which increase strength in the alloy are limited to very
small amounts 0.2% in this alloy (Al-6061), to minimize its effects on corrosion resistance
[35]. Increasing the copper content decreases the corrosion resistance of the alloy, as can be
seen for Al-7075 (1.17% Cu) and Al-Cu (4.80% Cu) in Table 2.
The electrochemical system can be represented by a theoretical model consisting of a
parallel combination of resistor, Rp , and capacitor, C , in series with the electrolyte
resistance, RS, [22,23]. Other equivalent circuit models including capacitive features and
inductive features are successful in describing the electrochemical behavior of Al or its alloys
in the very low frequency regions, (i.e., f < 0.1 Hz) [10,19,36]. Since it is necessary to
compare the electrochemical behavior of Al and the investigated alloys, it is useful to reduce
the theoretical model to the least number of components which can describe the dielectric
properties of the oxide film. The capacitor/parallel resistor model, investigated in the high
frequency region (f > 0.1 Hz) is suitable for such investigation. At high frequencies the
resistance of the inductive features becomes included in the polarization resistance, RP ,
which is equivalent to the corrosion resistance, Rcorr , of the material. The capacitive features
of the high frequency semicircle are related to the barrier layer itself [30,33]. The impedance
data of the different electrodes were correlated to the model described above.
A procedure of data fitting with minimum error was used in which a fitting program was
applied to fit the experimental data to the computer-generated data. The program used
422
Al-Kharafi et al.
enables data fitting in the required range of frequency. For the measurements presented, it
was necessary to fit experimental data of the high frequency semicircle to the computergenerated data of the proposed model. For data fitting procedures, Bode plots are always
recommended as standard impedance plots [26,37]. Figure 6 presents the experimental Bode
plots for Al-6061 after ≈ 4 hours of electrode immersion in 0.1 M HNO3 (dotted line)
correlated to the computer-generated data of RS = 25.2 Ω, Rp = 1.39 kΩ and C = 2.36 μF
according to the data fitting program. The data fitting of Al-6061 gives a mean error in the
absolute impedance of 1.4% and a mean deviation in the phase angle, θ, of 1.00. The
procedure of data fitting was applied to other electrode materials and also to impedance
spectra taken at different time intervals of electrode immersion in the test solution. The Al
and Al-6061 impedance data represent the best fitt to the theoretical model after reaching the
steady state, whereas Al-Cu showed the largest deviation. The absolute impedance and phase
angle deviation for Al-Cu electrodes after 4 hours of electrode immersion in 0.1 M HNO3
solution are 3.1% and 1.6o, respectively. The small deviation of the absolute impedance
values obtained with Al-6061 and pure Al indicate that the barrier layer on these materials
approaches ideal capacitor behavior.
Nyquist plots of Al-Cu (⎯), Al7075 (…), Al (----) and Al-6061 (- -) after 250 min of electrode
immersion in 0.1 M HNO3. The
values of Rp for each electrode in
2
Ωcm are 122, 142, 372 and 234,
respectively
Computer fitted data of RS = 25.2
Ω , Rp = 1.39 kΩ and C = 2.3 μF
(⎯) to experimental Bode plot of
Al-6061 after 250 min of electrode
immersion in 0.1 M HNO3 (••)
Effect of Chloride Ion Concentration
In this series of experiments the effect of chloride ion concentration on the corrosion and
passivation behavior of Al, Al-Cu, Al-6061 and Al-7075 was investigated. The electrodes
423
Fundamental Aspects
were mechanically polished and investigated in 0.1 M HNO3 solutions containing different
concentrations of Cl- ions ranging between 3.5 mM and 0.35 M. A typical example of the
data from these investigations is presented as Nyquist plots for the Al-Cu alloy in Fig. 7. For
all investigated materials in all measurements, two semicircles were recorded. A high
frequency semicircle and a low frequency inductive loop. The diameter of the high frequency
semicircle depended on the concentration of Cl- (Fig. 7). It decreased as the concentration of
Cl- ions increased. This means that the natural passivity of Al or Al alloys decreases in the
presence of chloride ions, as was reported for aluminum. Chloride ions attack the base metal
by dissolving the passive film at defective areas [16,37,38]. The decrease of polarization
resistance and passive layer thickness with increasing Cl- ion concentrations means that the
native barrier layer is too thin to impart complete passivity. As the concentration of Clincreases, the extent of surface attack increases and Cl- spreads laterally beneath the original
native film leading to the loss of its passive characteristics with the formation of a
nonprotective, oxyhalide layer on the metallic surface [16,24,25]. At very low concentrations
of Cl- (i.e., 3.5 mM), the rate of barrier layer removal is very low and is exceeded by the rate
of passive film repair that occurs in nitric acid solutions; hence, no remarkable attack can be
observed. In this case, oxide film thickening occurs as was indicated by the increase of 1/C
versus t relations [17,24,25]. At higher concentrations of Cl- (i.e., > 35 mM), the rate of
barrier film removal is compensated for by the rate of passive film formation, and equilibrium
is attained in which an approximate constant value of δ can be calculated. The remaining
barrier layer thickness and its polarization resistance depends on the electrode material.
Figure 8 presents the effect of Cl- ions with the same concentration (i.e., 35 mM) on the
different electrode materials. As can be seen from the Nyquist plots of Fig. 8, Al-6061 has
better corrosion resistance against Cl- than the other alloys investigated and even better
resistance than aluminum itself. The Al-Cu and Al-7075 alloys are much affected by the
presence of chloride ions. A remarkable decrease in the corrosion resistance of both alloys
with the increase of the concentration of Cl- was recorded. This behavior can be attributed to
the alloy’s constituents. The presence of Mg in the Al-6061 alloy improved the corrosion
characteristics of the alloy in chloride media, whereas the presence of Cu increased the
corrosion rate of the alloy. The values of corrosion resistance of the different materials after
reaching the steady state in 0.1M HNO3 containing 35 mM Cl- took the order, Rp Al-6061 >
Rp Al > Rp Al-7075 > Rp Al-Cu. The values in kΩ-cm2 for the sequence are 0.754, 0.382,
0.157 and 0.097 for Al-6061, Al, Al-7075 and Al-Cu, respectively.
The presence of Cu on the alloy’s surface was confirmed by XPS measurements. Figure
9 presents the XPS spectra for Al, Al-Cu and Al-6061. In all spectra, the characteristic peaks
of aluminum (Al 2P at 75.5 eV and Al 2S at 120.0 eV), oxygen (O 1S at 532.5 eV) and
carbon (C 1S at 285.5 eV) were recorded. The XPS of Al-Cu (Fig. 9b) contains additional
copper peaks (Cu 2P3 at 932.5 eV and Cu 2Pl at 952.5 eV) which indicate the presence of Cu
on the electrode’s surface even after 3 hours of electrode immersion in the test solution. The
XPS of Al-6061 did not show a pronounced Mg peak, i.e., the characteristic sharp Mg XPS
peak (Mg 1S at 1305 eV) is not present [39]. This means that Mg is present more likely in
the form of Mg2Si in the Al-6061 bulk and not on the alloy’s surface. The similarity between
the XPS spectra for Al and Al-6061 (compare Fig. 9a and b) explains the close corrosion
behavior of both materials and supports the conclusion that the barrier layer on both materials
consists of a stable Al2O3 film. The presence of Cu on the Al-Cu surface is responsible for
424
Al-Kharafi et al.
the higher rates of corrosion recorded for this alloy. It initiates cathodic areas or flawed
regions which leads to the observed decrease in the corrosion resistance after long immersion
times in the test electrolyte. The presence of flawed regions on Al-Cu was confirmed by
SEM. Figure 10 presents the scanning electron micrographs of Al and Al-Cu surfaces before
immersion in the test solution (Fig. 10a and b, respectively) and of an Al-Cu surface after 3
hours of immersion in 0.1 M HNO3 containing 0.35 M Cl-.
Nyquist plots of Al-Cu electrodes
after 250 min of electrode
immersion in 0.1 M HNO3
containing different concentrations
of chloride ions
(⎯)
0.35 M Cl , (…) 0.035 M Cl and
(---) 3.5 mM Cl-
Nyquist plots of Al-Cu (⎯), Al7075 (----), Al (- - -) and Al-6061
(…) electrodes after 250 min of
immersion in 0.1 M HNO3 + 35
mM Cl- solution. The values of Rp
2
for each electrode in Ωcm are 91,
114, 630 and 336, respectively
Mechanism of the Corrosion/Passivation Process
The mechanism of corrosion of Al and its alloys is based on the dissolution of Al atoms
from the active sites or flawed regions on the naturally passivated material. The dissolved
atoms are gradually removed through the formation of hydroxide with increased coordination
from 1 to 3 to form Al(OH)3. The formed hydroxide sticks to the surface in neutral solutions,
and hence, a decrease in the corrosion rate takes place which gives the remarkable passive
behavior of Al and its alloys in neutral solutions [40]. In acid and alkaline solutions, the
formed Al(OH)3 reacts in a purely chemical manner to form soluble species which go in
solution leaving bare active sites which in turn lead to the observed increase in the corrosion
rate in these media [25,411. The mechanism of the corrosion process represents an
+
irreversible coupled reaction, the anodic part of which is the reduction of H , water or
dissolved oxygen leading to hydrogen evolution and OH- formation in the vicinity of the
active regions according to:
425
Fundamental Aspects
- cathodic reaction:
H+ + e- → H
H2O + e- → H + OH-
(3)
(3’)
H + H2O + e- → H2 + OH-
(4)
1/2 O2 + H2O → OHads. + OH-
(5)
OHads..+ e- → OH-
(6)
Al + OH- → Al(OH)ads + e-
(7)
Al(OH)ads. + OH- → Al(OH)2ads. +e-
(8)
Al (OH)2ads. + OH- → AI(OH)3ads. + e-
(9)
- anodic reaction:
In the presence of Cl- ions, metal dissolution occurs through the attack of Cl- according to
Al + Cl- → AlClads. + e-
(10)
AlClads.. + Cl- → AlCl2ads. + e-
(11)
AlCI2ads.. + Cl- → AlCl3ads. + e-
(12)
and the most appropriate cathodic counterpart is reaction (3) followed by reaction (4). The
presence of cathodic areas enhances the corrosion process which was observed with the Cucontaining Al-alloys (i.e., Al-Cu or Al-7075). The presence of Cu on the material surface
increases the ratio of cathodic/anodic areas leading to an increase in the corrosion rate. The
natural tendency of Cu or Zn to form oxyhalide complexes is also an additional effect which
causes the loss of the protective properties of the naturally occurring barrier layer in the
chloride solution. In chloride-free nitric acid solutions, the formed Al(OH)3 can be oxidized
to the stable A12O3 passive film. The oxidation power of nitric acid depends on its
concentration [22-24]. This explains the passive behavior of Al and Al-6061 in nitric acid
solutions. The presence of Cu on the surface of the Cu-containing alloys increases the
cathodic areas, and hence, increases the tendency for galvanic corrosion to occur, which
explains the comparatively high rates of corrosion of these alloys after long immersion time
in nitric acid solutions (Table 2).
426
Al-Kharafi et al.
X-ray photoelectron survey spectra
of naturally passivated Al(a), AlCu(b) and Al-6061 (c) after 3 h of
electrode immersion in 0.1 M
HNO3
. SEM micrographs of
mechanically polished Al(A), AlCu (B) and Al-6061 (C) after 3 h
of immersion in 0.1 M HNO3
solution containing 0.35 M Cl
CONCLUSIONS
The corrosion and passivation behavior of Al, Al-6061 Al-7075 and Al-Cu, is dependent
of both the alloying element and the corrosive medium. In nitric acid and nitric acid
containing Cl- ions solutions, Al-6061 has the highest corrosion resistance. This alloy with its
1.40% of Mg. and 0.60% of Si behaves like a perfect dielectric. The corrosion behavior of
which is most likely similar to aluminum, especially after long periods of immersion in nitric
acid or nitric acid containing chloride solutions. The presence of Cu on the surface of Cucontaining alloys initiates flawed regions in the barrier layer which are responsible for the
higher corrosion rates of these alloys.
ACKNOWLEDGEMENT
427
Fundamental Aspects
The financial support of Kuwait University, Research Grant No. SC060, is gratefully
acknowledged.
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1986, p. 342.
6. W.A. Badawy, M.S. El- Basiouny and M.M. Ibrahim, Ind. J. Technol. 24, 1986, ,p. 1.
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9. C.B. Breslin and W.M. Carroll, Corros.Sci. 34, 1993, p. 327.
10. C.M.A. Brett, I.A.R. Gomes, J.P.S. Martins, J. Appl. Electrochem. 24, 1994, p. 1158.
11. M. Elboujdaini, E. Ghall, R.G. Barradas and M. Glrgis, Corros. Sci. 30, 1990, p. 855.
12. N. Khalil and J.S.L. Leach, Electrochim. Acta 31, 1986, p. 1279.
13. V. Surganov, P. Morgan, J.G. Nielsen, G. Gorokh and A. Mozalev, Electrochim. Acta 32,
1987, p. 1125.
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Shershulskii, Electrochim.Acta 35, 1990, p. 955.
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429
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
THE SUSCEPTIBILITY OF MOLYBEDNUM AND VANADIUM-BEARING
AUSTENITIC STAINLESS STEEL WELDMENTS TO INTERGRANULAR
CORROSION
M.K. Karfoul
College of Chemical Engineering and Petroleum
Ba’ath University, Syria
ABSTRACT
The development of the chemical, fertilizer, petrochemical, refining and energy industries
depends, in most cases, on resolving the problems associated with the use and maintenance of stainless
steels. The most important problem faced with the use of stainless steels is intergranular corrosion
(IGC). This subject has attracted the attention of many research organizations for several years and
still concerns many to date.
Key Words: Austentic stainless steel, weld metal, intergranular corrosion
INTRODUCTION
Nowadays, the metal manufacturing industries produce chromium and nickel-bearing
austenitic stainless steels which are highly resistant to intergranular corrosion (IGC). This
level of resistance is achieved by stabilizing the characteristics of the steels by adding
titanium or niobium, or by decreasing the percentage of carbon in the steel to a very low level
or both [1]. In most cases, these stainless steels are used in welded mechanical equipment
which are also repaired by welding. Such uses, however, might lead to IGC (sensitization),
especially when the steel remains at temperatures of 500-700°C during the welding process.
Within this temperature range, depletion of chromium can occur at the grain boundaries in the
heat affected zone. Welding or repeated welding promotes such a phenomenon. Therefore, it
is deemed desirable to have a stainless steel alloy that resists sensitization during welding in
order to prevent IGC. One of the basic methods used to stabilize the austenitic chromiumnickel stainless steel weldments is through the addition of titanium or niobium.
After welding, steel weldments passes through the temperature range of 900-1000°C for
a period of time. This allows the formation of carbide stabilizers such as Ti and Nb carbides
[2-4], resulting in lowering the carbon content in the solid solution. However, when such an
alloy reaches the critical temperature range for sensitization, i.e., 500-750°C, chromium
carbide is formed (i.e., Cr23C6) along the grain boundaries. This results in dispersed and unconnected areas depleted in chromium. Therefore, the grain boundaries become less
resistance to IGC than those that are stabilized with Ti or Nb. The percentage of these latter
alloying elements in the steel should depend on the amount of carbon present in the steel.
431
Fundamental Aspects
This matter appears simple, but many complications, however, occur upon the addition
of Ti or Nb. Titanium has a higher affinity for reaction with carbon than does chromium. Ti
is also more reactive to oxygen than chromium or the residual elements such as Mn and Si
are. Therefore, during electric arc welding, a titanium-stabilized steel electrode tends to react
with oxygen and be consumed completely. To avoid such complications, inert gas is used
during welding of such electrodes. However, this increases the cost of the components to be
welded. Alternatively, titanium can be replaced by niobium with its lower oxidation rate as a
stabilizer for the steel weld metal.
When mechanical components are operated under severe corrosive conditions and/or are
exposed for an extended period to a temperature in the range of 500-600°C, it is
recommended that the amount of Nb to be 10-12 times greater than the percentage of carbon
in the steel. In addition to the amount of Nb which is required in such weld metal, the amount
of azote and carbon need to be calculated [4,5]. Therefore, more Nb is required than is
mentioned above, and this leads to a loss of steel toughness and the appearance of hot short
cracks in the weld metal during the welding process [6].
The sensitivity of welding materials to IGC increases with increases in carbon, titanium
and niobium. When using a welding electrode with a maximum carbon percentage of 0.06,
the sensitivity of weld metal may be increased. However, this does not negate the weld
metals sensitivity to IGC, especially during solid welding processes or applied on different
layers. This is related to the presence of large percentage of carbon in the welding material
the migration of carbon from the base metal or the degradation of carbon containing
materials that cover the welding electrode [7].
The chromium-nickel steels with niobium added had the tendency to crack during
cooling. This action appears more than in castings, solid welding joints and during the
welding of thick plates. This damage increases with the increase in acidity of the slag during
welding [8] (especially for electrode materials containing cilium) which produces hot short
cracks. Getting rid of compounds containing cilium in materials covering electrodes is
impossible.
The negative effect of niobium on hot short cracking is associated with the small
dissociation of niobium in iron especially at the low eutectic melting temperature of Fe-Nb
[9,10]. This eutectic effect cannot be avoided in practice unless it is associated with the
amount of niobium and carbon. Niobium usually increases the ferrite phase in steels,
especially if it is present in steel in a ratio of 1:10 with respect to carbon. Therefore, due to
the inhomogeneous concentrations in weld metal, microscopic cracks appear in pure
austenitic areas. This is expected since niobium activate cracking in weldments of pure
austenite. One may thus conclude that stabilizing the chromium-nickel steels with niobium is
associated with many technological difficulties.
Molybdenum is a more positive stabilizing additive with this respect. Mo is known to
posses a large degree of dissociation in iron, chromium and nickel. The eutectic melting
point of Mo is not that different from the melting point of the original metals of Fe, Ni and
Cr. Metals that contain Mo do not exhibit hot short cracks. Molybdenum is also known to
promote the ferrite phase and has the affinity to react with carbon, but to a lesser degree than
Nb with carbon. Therefore, Mo aids in raising the resistance of weld metals of chromiumnickel stainless steels to hot cracks.
432
Karfoul
Molybdenum plays a role in softening the microstructure of weld metals of pure
austenite to hot cracks [11]. Molybdenum also increases the resistance of weld metals to
corrosion, and increases its surface negativity [12.3]. The positive effect of Mo is exhibited
in improving the resistance of weld metal to hot cracks and improving its technological
soundness [14] or increasing the IGC resistance of weldments [15] in the presence of special
electrodes for welding the austenitic stainless steel with the use of the common stabilizing
elements.
Vanadium possesses a great ability to dissociate in iron. It is strongly reactive with
carbon, vanadium carbides are formed such as V4C3 and remain stable at higher temperatures.
Therefore, vanadium is considered to be a stabilizer for carbon as carbides in the grains of the
chromium-nickel stainless steel. Vanadium also plays a major role in promoting ferrite
formation [16]; thus increasing the resistance of weldments to hot cracks.
The effect of vanadium in reducing the IGC of chromium-nickel steels is not clear. The
information available with respect to the subject is lacking. Therefore, the objective of this
research was to evaluate the addition of vanadium to steel weldments with molybdenum,
rather than of titanium and niobium, as to its resistance to IGC.
EXPERIMENTAL PROCEDURE
In this research, stainless steel specimens were prepared by electric arc welding of two
sheets, 500 mm in length and 5 mm thick, made of chromium-nickel steels stabilized with
titanium. These sheets were composed of: 16.5% Cr, 9.6% Ni, 0.66% Ti, 0.1%C and the
balance was Fe.
The welding system was chosen in a way to ensure minimum interaction with the base
metal. The operational angle of the welded surfaces was 90°. The distance between the two
plates was 2 mm, and the welding current and potential were 160 A and 25 V, respectively.
The welding speed was 0.17 cm/s, and the diameter of the welding electrodes was 5 mm.
Using this welding system, a single pass was applied. the base metal was protected by a
copper sheet that was placed underneath the two welded plates.
Stainless steel welding rods were used. They were composed of 0.06-0.08 wt.% C, 18.4
wt.% Cr, 11.00-11.38 wt.% Ni, 2.17-2.37 wt.% Mo, 1.35 wt.% Ti, 0.35 wt.% Si, 0.029 wt.%
P, and 0.014 wt.% S.
The proper chemical composition of the weld metals was achieved by adding the
necessary elements to the substance covering the welding electrode.
The chemical and phase compositions of the weldments [15] was based on physical and
mathematical methods. The chemical and phase compositions of the weldments are shown in
Table 1.
Table 1 shows that the first three components of the weldment had a fixed composition
except for their Mo content. The percentage of Mo was increased to 2, 3, and 6% to study the
effect of the alloy on the weld metal’s resistance to IGC. In the forth composition, half of the
Mo was replaced with vanadium, the percentage of Mo was 3 wt%. The percentage of Cr
was kept at 17 in order to allow for the formation of a ferrite phase in the weld metal. In
addition, a fifth composition was also tried with the same Mo and V percentages as in the
fourth composition, but with 20 wt% Cr to maximize the ferrite phase in the weld metal.
433
Fundamental Aspects
Table 1a. The Chosen Chemical and Phase Compositions of Weldments
Chemical Composition (wt%)
Composition No C
Mn Cr Ni
Mo V
1
0.1 2.0
20 10.5 2
2
0.1 2.0
20 10.5 3
3
0.1 2.0
20 10.5 6
4
0.1 2.0
17 10.5 3
3
5
0.1 2.0
20 10.5 3
3
Phase Composition
S(%) Grain Diameter (μm)
4
20
5
20
10
20
4
20
10
20
Table 1b. Actual Experimental Chemical Compositions of Weldments
Composition No.
1
2
3
4
5
C
0.11
0.11
0.11
0.11
0.11
P
0.03
0.028
0.027
0.033
0.03
Chemical Composition (wt %)
S
Si
Mn
Cr
Ni
Mo
0.018 0.13 1.83 20.8 10.55 1.9
0.018 0.16 1.74 20.0 10.96 3.1
0.018 0.20 0.81 19.7
9.84 6.38
0.018 0.20 1.79 16.6 10.50 2.95
0.018 0.20 1.80 19.5 10.85 2.77
V
2.85
1.86
Table 1c. Chemical and Phase Compositions of Weldments
Composition No.
1
2
3
4
5
δ (%)
4.7
4.4
8.2
6.0
9.0
Phase Composition
Grain Diameter (μm)
21
17
19
17
17
The amount of the ferrite phase on the weld metal specimens was determined
magnetically by an α-phase meter with ± 5% error, as is shown in Fig. 1. These same
specimens were also used to study the weld metals resistance to IGC.
RESULTS
A metallurgical microscope was used to measure the austenite grain diameter of the weld
metals, as shown in Fig. 2. After being welded, the specimens were prepared (Fig. 1) and
heat-treated at a temperature range of 500-800°C for different time periods as shown in Table
2. Then, the specimens were quenched with water. To determine the susceptibility of the
specimens to IGC after the heat treatment, they were all immersed in Strauss solution for 24
434
Karfoul
hrs. Strauss solution is made up of 100 gr of CuSO4 5H2O + 0.1l H2SO4 + 1l of distilled
water + thin copper sheets.
Figure 1. The Location of the points where the δ ferrite phase was measured
in the welded metal
Composition No. 1
Composition No. 2
Composition No. 4
Composition No. 3
Composition No. 5
435
Fundamental Aspects
Figure 2. Microstructures of the weld metals for the different studied compositions
Table 2. Time in Minutes of the Different Heat-Treatment Temperatures Studied
550
600
650
700
750
800
1
Heat Treatment
Temperature (°C)
-
25
50
3
5
10
25
10
25
10
25
3
500
55
110
300
500
5
30
50
10
300
10
30
50
110
300
50
10
25
50
100
300
500
25
50
100
300
500
0.5
1
3
5
10
5
10
30
50
110
300
500
5
10
2
1
3
5
10
100
5
30
50
300
10
30
50
110
300
500
10
25
50
100
300
500
25
50
110
300
500
5
10
30
50
90
-
5
10
30
50
100
-
4
50
100
300
500
5
50
100
300
500
10
25
50
100
300
60
25
50
110
300
500
50
110
110
300
After being boiled for 24 hours. in Strauss solution, the specimens were bent to a 180°
angle. If no cracking was observed in the bent samples, they were considered to be resistant
to IGC.
DISCUSSION
The aqueous solution of copper sulfate and sulfuric acid was chosen because it attacks
only the regions of the specimens with chromium contents of < 12%. This negates the use of
weight-loss or other methods to check whether or not such specimens are susceptible to IGC.
On the other hand, results obtained using acetic acid may require the determination of the
weight loss of specimens because IGC might not be well defined. This is because in this
case, the acid attacks regions containing more than 12% Cr and removes grains from the
436
Karfoul
metal surface. Therefore, weight -loss measurements are not representative of intergranular
corrosion.
It can be seen in Figs. 3, 4 and 5 that the addition of 6 wt% Mo to the weldments
increased greatly the IGC resistance of the austenitic stainless steel weld metal. It also
increased annealing of the microstructure of the weld metal, resulting in increasing the
austenitic grain boundary surface area and thus not allowing the precipitation of Cr23C6 and
preventing the depletion of Cr around the grains. The addition of Mo also aids the formation
of δ ferrite phase in the microstructure of the weld metal and thus decreases the IGC
susceptibility of the metal.
It was also observed from Figs. 6 and 7 that the replacement of half of the Mo by V in
the weld metal of the composition No 3 did not change the resistance of the alloy to IGC,
indicating similarity of V and Cr. However, the increase in the amount of δ ferrite in the
microstructure of the weld metal tended to increase the resistance of Cr-Ni weld metals to
IGC, as shown in Fig. 7.
Figure 3. Response of composition No. 1
Figure 4. Response of composition No. 2
Figure 5. Response of composition No. 3
Figure 6. Response of composition No. 4
437
Fundamental Aspects
Figure 7. Response of composition No. 5
CONCLUSIONS
The results of tests conducted on the Cr-Ni austenitic stainless steel showed that the
addition of either Mo or V to the weld metal possess similar effects on the sensitivity of the
weld metal to I.G.C.
The addition of Mo and V tends to shift the Rolason curves down and to the right,
indicating an increase in the resistance of the weld metal to IGC.
REFERENCES
1.
E.C. Bain, R.H. Aborn and J.J.B Rutherford, The nature and prevention of intergranular
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I.Z. Kogan, Niobium in welding electrodes, Heat resisting, No. 4, 1951, pp. 1-3.
H.J. Rocha, Die sensibilisiertes austenisches stahle durch chromkarbide, Zeitchrift fur
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B.I. Medovar, Welding of heat resisting austenitic stainless steel and alloys, Moscow,
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3.
4.
5.
6.
7.
8.
9.
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Karfoul
10. M.Ck. Sharshorov, Hot cracks during the welding of heat resistant alloys, Moscow,
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molybdenum, The Sixth Middle East Corrosion Conference, Conference Proceeding,
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Houdermont,
Handbuch
der
Soonder
Stahlkunde,
Springer-verlag,
Berlin/Gohingen/Heidd, 1956, Russian Translation, Moscow, 1960, p. 1064.
439
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
EFFECT OF CRYSTALLIZATION ON THE CORROSION BEHAVIOR OF
AMORPHOUS FeCr9P6C3Si0.2 ALLOY IN 1M H2SO4
F. Hajji1,2, S. Kertit1, J. Aride1 and M. Ferhat2
1
Laboratoire de Physico-Chimie des Matériaux associé l'AUPELF.UREF (LAF502),
Ecole Normale Supérieure de Takaddoum, B.P. 5118, Rabat, Morocco
2
Laboratoire de Chimie-Physique Générale, Faculté de Sciences, Rabat, Morocco
ABSTRACT
The crystallization of amorphous FeCr9P6C3Si0.2 alloy was investigated by x-ray diffraction
(XRD), differential thermal analysis (DTA) and scanning electron microscopy (SEM). Specific
differential heat curves of the FeCr9P6C3Si0.2 alloy exhibited three exothermic peaks indicating that
the crystallization of the amorphous alloy occurred through the formation of three kinds of metastable
crystalline phases. The deterioration of corrosion resistance by the crystallization of the amorphous
FeCr9P6C3Si0.2 alloy was studied by electrochemical methods to correlate the corrosion behavior
with the increase in the heterogeneity of the alloy. As soon as the stage formed in the amorphous
matrix, the anodic current density increased. The current density in the active and passive regions
increased continuously in sulfuric acid solution during the crystallization processes. Chemical
heterogeneity, based on the formation of precipitation segregation and other compositional
fluctuations, seemed to be responsible for the deterioration of the corrosion resistance.
Keys Words: Amorphous alloys, corrosion, crystallization, potentiodynamic measurements
INTRODUCTION
The corrosion behavior of amorphous alloys was first studied in 1974 [1]. It was then
extensively reviewed [2-6] for simple binary alloys such as Fe80B20 as well as commercial,
multicomponent systems. Amorphous iron-based alloys prepared by quenching from the
liquid state contain a large amount of various metalloid elements that stabilize the amorphous
structure in the solid state. Among the metalloid elements, phosphorus is the most effective
at concentrating chromium in the passive film [7], and hence, the passive film formed on the
amorphous alloys contained phosphorus and a small amount of chromium consisting
exclusively of hydrated chromium oxyhydroxide. This is partly responsible for the extremely
high corrosion resistance of these amorphous alloys [8-9]. Amorphous alloys containing P
were almost always more corrosion resistant than alloys containing B, Si or C. However, the
alloys containing P became much less corrosion resistant [10,11] after treatment at the
crystallization temperature, at which point P migrated into the grain boundaries and caused
severe intergranular corrosion. The chemical heterogeneity seemed to form a high density of
weak points in the passive film with respect to corrosion as well as localized corrosion attack.
The heat treatment of amorphous alloys gives rise to the formation of various metastable
crystalline phases (MS) in the amorphous matrix before the formation of stable crystalline
441
Fundamental Aspects
phases [12]. In this work, amorphous FeCr9P6C3Si0.2 alloy was used. Mainly
potentiodynamic polarization experiments were conducted to understand the electrochemical
corrosion of the alloy. The effect of heat treatment on the behavior of the amorphous alloy in
aerated sulfuric acid media was investigated.
EXPERIMENTAL PROCEDURE
Amorphous FeCr9P6C3Si0.2 alloy ribbons 2 mm in width and 10 μm in thickness were
produced by a rapid quenching technique (melt-spinning). The number attached to a
respective element in the alloy formula denotes the nominal content in atomic percentage.
After isothermal heat treatment of the alloy in an evacuated quartz tube at various constant
temperatures at a gas pressure of 3.10-6 torr, diffractometric measurements were made using
a diffractometer, with Co Kα radiation at a scanning speed of 8o/minute.
Differential thermal analysis (DTA) at a heating rate of 10 K/minute was carried out to
confirm the multistage crystallization which produced multiple exothermic peaks. The
crystallization process was also examined by a scanning electron microscopy (SEM) after
undergoing isothermal heat treatment.
The electrochemical experiments were performed with an Amel potentiostat system.
Anodic polarization curves for the alloy were measured potentiodynamically with a potential
sweep rate of 1 mV/s. The electrochemical measurements were conducted in unstirred,
aerated 1 M H2SO4 solution which was prepared using a reagent. The alloy was not
mechanically polished. Both sides of the ribbons were immersed in the test solutions. All the
experiments were carried out at room temperature. Polarization studies were conducted in a
simple electrolytic cell at three electrode using platinum counter electrodes and a saturated
calomel (reference) electrode (SCE).
RESULTS
Differential Thermal Analysis (DTA) Measurements
The DTA curves for the amorphous FeCr9P6C3SiO.2 alloy exhibited three exothermic
peaks as shown in Fig. 1. The first peak was at about 448°C, the second peak was at 489°C
and the final peak was at 578°C. This indicates that the crystallization of the amorphous
alloy occurred through the formation of three kinds of MS. In the case of heat treatment at
300 and 400°C, the pace of the thermograph was not modified. However, the DTA curves for
the alloy treated at 500°C for 1, 2 and 3 hours exhibited only one exothermic peak. In this
case, the crystallization was only partial. But, for the alloy treated at 500°C for 8 hours or at
600, 700 and 800°C for 1 hour, crystallization was total.
X-Ray Diffraction (XRD) Measurements
The x-ray diffraction (XRD) patterns for the amorphous FeCr9P6C3SiO.2 alloy and a
series of isochronal heat treatments are shown in Fig. 2. It is evident that the as-quenched
state was typical of the amorphous state, and no crystalline phases were observed. With
annealing at 300 and 400°C the alloy stayed amorphous. When the temperature was
increased, the x-ray pattern became totally crystalline. As shown in Fig. 2, XRD patterns
revealed that the crystallized FeCr9P6C3SiO.2 alloy consisted of many phases. XRD
confirmed the presence of α-Fe as an fcc phase with a = 2,866 Å in the ribbons annealed at
442
Hajji et al.
500°C. Because of the very complicated nature of the diffraction patterns and the intense
diffraction lines of the α-Fe phase, an accurate analysis of the lattice constants of the other
phases was not possible.
The crystallization of various amorphous metal-metalloid alloys has been studied by
many authors. When the metalloid concentration is not exceedingly high, such as at 28
atomic percentage, the crystallization of amorphous alloys generally takes place as follows:
the heat treatment of the amorphous alloy gives rise to the precipitation of a MS in the
amorphous matrix. This phase contains a large amount of the main metallic component of the
alloy and thus it has the same crystal structure as the main metallic component. The
amorphous phase then disappears by the formation of two or three MS, through
transformation diffusion of various elements and by recrystallization and/or decomposition of
the metastable phases.
Figure1. Differential thermal analysis curves Figure
of amorphous FeCr9P6C3Si0,2
alloy before and after isothermal
heat treatment with a heating rate of
10 K/minute
2.
X-ray diffraction patterns
FeCr9P6C3Si0,2 alloy before
and after isothermal annealing
at different temperatures and
times
443
Fundamental Aspects
Figure 3. Breaking faces of the alloy FeCr9P6C3Si0,2 thermally treated at 300°C for 1 hour
(a), 500°C for 1 hour (b) and 800°C for 1 hour (c), Micrograph of the shining face
of the alloy FeCr9P6C3Si0,2 after annealing at 500°C for 1 hour (d), 800°C for 3
hours (e) and mate face at 800°C for 3 hours (f). Micrograph of the alloy
FeCr9P6C3Si0,2 after annealing at 500°C for 1 hour (g). x-ray cartography of the
shining face of the alloy FeCr9P6C3Si0,2 after annealing at 800°C for 1 hour (h)
Scanning Electron Microscopy (SEM)
Scanning electron microscopy (SEM) images obtained on breaking faces that have been
thermally treated are shown on Fig. 3. For samples annealed at 300 or 500°C for 1 hour,
these alloys formed in a disordered grain stacking. They seemed to have a spherical form,
and their size increased with increasing temperature (Fig. 3a and 3b). On the other hand,
when the temperature and the annealing time increased, the alloy presented a mixture of two
aspects: a granular aspect and a column-like aspect (Fig. 3c). The observed crystallization
was probably due to an intrinsic heating up of the alloy during the annealing process.
Annealing of the alloy did not change the state of the external surface of the alloy. Indeed the
state of the surface stayed amorphous, as shown in the SEM micrograph (Fig. 3d). In the case
of the alloy annealed at 800°C for 3 hours, the state of the surface of the shining face (Fig. 3e)
seemed to be formed by a granular stacking of a quasi-spherical form, even if the mate face
444
Hajji et al.
always stayed amorphous (Fig. 3f). In order to clarify this phenomenon, note that the
micrograph in Fig. 3g, obtained for a sample annealed at 500°C for 1hour, presents a mixture
of two different crystallographic aspects: an amorphous aspect (alloy surface) and a crystal
aspect (inside the alloy).
An x-ray cartograph giving the distribution of the Fe, Cr, P and C elements of a sample
annealed at 800°C for 1 hour showed good homogeneity of the shining surface (Fig. 3). This
confirms that the surface of the sample was not affected by the temperature. These results
could be explained by the fact that increases in the annealing temperature and time led to an
increase in the molecule mobility inside the alloys before their freezing into a crystal state.
Therefore, the grain size increased with annealing temperature and time. Nevertheless, this
increased mobility did not lead to a perfect atomic order which would have filled the entire
volume of the alloy, mainly the external surface as shown by the SEM results. The inside of
the alloy was then formed by a stacking of micro-crystallites separated by a granular
boundary which are probably disordered. Furthermore during the growth of such microcrystals, depending on the annealing temperature and time, it is possible that some impurities
formed around the grains with compositions quite different from those of the crystallites.
Figure 4. Change in the anodic polarization curves of amorphous FeCr9P6C3Si0,2 alloy
before and after annealing in 1 M H2SO4 with 1 hour of heat treatment at 300°C,
400, 500, 600, 700 and 800°C
Electrochemical Measurements
445
Fundamental Aspects
Figure 4 shows changes in the anodic polarization curves of the alloy measured in 1 M
H2SO4 according to the temperature of heat treatment at different time periods. The anodic
polarization curve of the untreated alloy is also shown in Fig. 4 for comparison. The various
anodic parameters determined from these curves are given in Table 1. The curve for the
amorphous alloy exhibits a typical active-passive transition (Fig. 4), and passivation in 1 M
H2SO4 solution. When the alloy was amorphous with the annealing (300 and 400°C), the
activation current density began to increase slowly. However, as soon as the first crystalline
phase was formed in the amorphous matrix, the speed of the activation current density began
to increase. The corrosion current density for activation (Ic) and for passivation (Ip) was
dependent on the annealing temperature. Before annealing, the activation current density (Ia)
was 34 μA/cm², and after annealing (Ic) became 40, 64, 726, 13793, 32307 and 47984 μ
A/cm² at 300, 400, 500, 600, 700 and 800°C for 1 hour, respectively. The current density in
the active and passive regions continuously increased during all the stages of nucleation. The
crystallized alloy also exhibited a wide passivity range. The superior corrosion resistance of
the amorphous alloy decreased with the crystallization of the alloy in 1 M H2SO4 solution.
Table 1.
Electrochemical Parameters of the Amorphous FeCr9P6C3Si0,2 Alloy
Before and After Isothermal Heat Treatment at 1 Hour in 1 M H2SO4
Solution
Amorphous 300°C 400°C 500°C
Ecor(mVvsS.C.E) -310
-310
-315
-365
34
40
64
726
Ia (μA/cm²)
15
13.2
27
222
Ip (μA/cm²)
600°C 700°C
-375
-382
13793 32307
1034
462
800°C
-380
47984
161
DISCUSSION
The present results revealed clearly that the formation of a crystalline phase in the
amorphous matrix increased the anodic current density. The anodic current density of
FeCr9P6C3Si0.2 alloy in the active and passive regions continuously increased during the
growth of metastable phases crystallites. Therefore, the crystallization of the amorphous
alloys led to the appearance of chemical heterogeneity with a subsequent increase in the
current density in the active region. The appearance of chemical heterogeneity also increased
the current density in the passive region. It has been assumed [13-16] that the passive film
was not essentially uniform but contained weak points (micropores) which were responsible
for the apparent passive current density in aggressive solutions. The micropores could be
formed on heterogeneous sites of the underlying alloy surface as well as on the phases which
are relatively difficult to passivate. Accordingly, the formation of chemically heterogeneous
sites in the alloys by heat treatment increased the passive current density. The question raised
is whether or not the chemical heterogeneity, in comparison with structural heterogeneity, is a
dominant factor in decreasing the corrosion resistance. The present authors [17] have shown
that the rapidly quenched, single phase alloys showed significantly high corrosion resistance
in comparison with the corresponding ordinary crystalline alloys. Naka et al. [11] have
reported that the passive current density of the Fe-10Cr-13P-7C alloy increased by two orders
of magnitude due to the formation of the solid solution phase in the amorphous matrix.
446
Hajji et al.
Therefore, rapidly quenched single phase alloys show significantly high corrosion
resistance. In contrast, heat treatment inevitably induces solid state diffusion and hence
results in various compositional fluctuations such as precipitation, segregation, and other
composition gradients. These compositional fluctuations may act as dominant active surface
sites with respect to corrosion.
CONCLUSIONS
It can be concluded that crystallization of this amorphous alloy FeCr9P6C3Si,0.2 did not
considerably alter its excellent corrosion resistance as long as the alloy remained a single
phase solid solution. This suggests that structure may not be a dominant factor in
determining the corrosion resistance of this amorphous alloy. When the crystallization was
complete, the corrosion resistance of the alloy deteriorated significantly.
REFERENCES
1. M. Naka, K. Hashimoto and T. Masumoto, J. Japan. Inst. Metals 38, 1974, p. 835.
2. Y. Waseda and K.T. Aust. J. Mater. Sci. 16, 1981, 2337.
3. R.B. Diegle, N.R. Sorensen, T. Tsuru and R.M. Latanision, in Treatise on Materials
Science and Technology (Edited by J. Scully), Vol. 23, p. 63, Academic Press, London
(1983).
4. K. Hashimoto, in Amorphous Metallic Alloy (Edited by T.E. Luborsky), p.471,
Butterworth, London (1983).
5. M.D. Archer, C.C. Corke and B. H. Harji, Electrochim. Acta 32, 1987, p. 13
6. P.C. Searson, P.V. Nagarkar and R.M. Latanision, in Modern Aspects of Electrochemistry
(Edited by R.E. White, J.O. Bockris and B.E. Conway), No. 21, pp.121-161. Plenum
Press, New York (1990).
7. K. Hashimoto, M. Naka, K. Asami and T. Masumoto, Corros. Eng. 27, 1978, p. 279.
8. K. Asami, K. Hashimoto, T. Masumoto and S. Shimodaira, Corros. Sci. 16, 1976, p. 909.
9. K. Hashimoto, K. Asami, K. Asami, and T. Masumoto, Corros. Eng. 28, 1979, p. 271.
10. R.B. Diegle and D.M. Lineman, Interim Technical Report No. 0NR-00014-77-C-0488-3
to the Office of Naval Research.
11. M. Naka, K. Hashimoto and T. Masumoto, Corrosion 36, 1980, p. 679.
12. T. Masumoto and R. Maddin, Acta Metall. 19, 1971, p. 725.
13. K. Hashimoto, K. Asami and K. Teramoto, Corro. Sci. 19, 1979, p. 3
14. K. Hashimoto and K. Asami, Corros. Sci. 19, 1979, p. 251.
15. K. Hashimoto and K. Asami, Passivity of Metals, Proceedings, 4th Intern. Symp. on
Passivity of Metals, (Edited by R.P. Frankenthal and J. Kruger), the Electrochemical
Society, Princeton, New Jersey, p.749 (1978).
16. K. Sugimoto and Y. Sawada, Corros. Sci. 17, 1977, p. 425.
17. M.Naka, K. Asami, K. Hashimoto and T. Masumoto, Proceedings, 4th International
Conference on Titanium (1980).
447
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
EXPERIENCE WITH VOC-COMPLIANT WATERBORNE AND HIGH
SOLIDS COATINGS IN CORROSIVE ENVIRONMENTS
P Kronborg Nielsen
HEMPEL Coatings, Lyngby, Denmark
ABSTRACT
Waterborne and high solids paints-VOC compliant paints are becoming more important in coating
specifications for environmental reasons. They contain between 0% and 35% organic solvents (VOC)
per litre, compared to standard paints, where VOC often contributes 60% or more. High solids paints,
such as low solvent epoxies, have been used successfully in, for example, submerged areas. However,
petrochemical installations are still often coated with standard paint types.
In recent years, official tests, practical demonstrations and case histories have shown that
waterborne and low solvent coating systems protect steel in aggressive environments on a par with
standard paints, especially in cases when high solid epoxy primers are combined with waterborne acrylic
top coats-the hybrid system.
Calculated per square meter applied and considering the longer maintenance free intervals, the cost
of high solids and hybrid systems is not excessive. Painting contractors and specifiers can thereby meet
upcoming VOC-legislation on sound economic and technical bases.
Key Words:
Waterborne coatings, high solids coatings, corrosion, VOC, coating specification,
petrochemical industry
INTRODUCTION
For environmental reasons, coating specifiers and painting contractors are today
constrained to using low volatile organic content (VOC) coatings for an increasing number of
painting jobs. The low VOC coatings are paints with a reduced content of organic solvents, and
they are present in the market in the form of waterborne paints and coatings with high solids.
In less corrosive environments, the long-term performance of the low VOC coatings is now
recognized to be fine. Their performance is often even above that of comparable standard
coatings such as alkyd, acrylic and epoxy [1]. However, in corrosive environments, experience
especially with waterborne coatings, is still limited. Lately, though, the offshore market and
independent testing laboratories have shown that coating systems with low VOC and
waterborne paints are well suited for severe corrosive environments. These systems may reduce
VOC emission by 70% or more, and the cost on an applied-square-meter basis is not excessive.
Thus, the new coating technology is in full accordance with the principle of BATNEEC, the
Best Available Technology Not Entailing Excessive Cost, which is statutory for all
environmentally directed developments.
EFFECT OF VOC's
449
Corrosion Protection and Monitoring
The VOC in paints are organic solvents, and solvents are necessary to facilitate production
and application. But once the paints are applied, solvents are only a nuisance. They are
inflammable, and have a negative influence both on man and nature. A predisposed painter's
long-term continuous exposure to organic solvents will have a negative effect on his
•
•
•
•
Respiratory system,
Nervous system,
Capacity for reproduction, and
Skin.
Various governments and civil councils have, therefore, introduced health and safety measures
to protect painters, such as the Control of Substances Hazardous to Health (COSHH)
regulations in force in the UK.
When organic solvents evaporate, they are decomposed by the ultraviolet radiation from
the sun. The decomposed molecules are highly reactive and easily form compounds with the
exhaust from automobiles and industrial air pollution. These chemical reaction products will
affect the local environment, and eventually cause smog and reduce metabolism in human
beings, animals, and plants.
In Europe and the USA, VOC emission is being addressed by various legislative measures,
e.g., organic solvents from painting processes by an European Union Directive [2] to be
implemented into a law. The directive requests a solvent management and reduction plan, and
sets limits on emissions, but it is aimed at the user of the paint, i.e., the painting shops.
Upcoming British and existing American laws address the solvent content in the paint itself.
The aim of both types of regulation is to reduce the overall VOC emissions. It is expected
that the actions laid down in the European Union Directive-once introduced-will reduce the
solvent emissions in European Union member states by 30% in 1999 (compared with 1985).
Paint manufacturers today have products in their assortment that can meet these
regulations, among which are the waterborne and the high solids (low VOC) products for the
protective coating of steel structures.
The user-the painting contractor-may address an emission directive by installing an
abatement system in his plant, but this could be a costly solution. A better way is to modify
working procedures and adjust the equipment to handle waterborne and/or high solids paints,
and to train the applicators accordingly.
COATING SYSTEMS WITH REDUCED SOLVENT EMISSIONS
The satisfactory performance of paint coatings for the protection of steel structures against
corrosion is determined by
• The choice and formulation of the products used in differently classified
environments, and
• The standard of workmanship and execution of the contract.
Agreement between the client and the contractor as to the specifications to be applied is
essential to the satisfactory execution of the work. The paint producer may also introduce
specifications, or they can be made in accordance with national standards, such as BS 5493,
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Kronborg Nielsen
DS/R 454, or DIN 55928/5.
There has, however, been a long-felt wish to standardize coating specifications and
workmanship. Therefore, at the time of writing (1996), the secretariat of the International
Organization for Standardization (ISO) has established working groups which have already
presented a series of working drafts on, among other topics, surface preparation, classification
of environments and protective paint systems (ISO 12944 [3]). The various corrosiveenvironments classifications in ISO 12944-2 [3] are given in Table 1. Industrial and coastal
areas are in Category 4, and Category 5M covers marine and aggressive areas.
ISO/CD#1 12944-5 (Table 2) contains a wide selection of paint systems suited for each of
these environments. Their performance has been confirmed after long experience and/or series
of successful trials. In order to make the ISO standard suitable for the future, systems with low
solvent emissions are also included. The paint systems are water dilutable, contain high solids,
or may be combined (hybrids).
In Figs. 1 and 2 are typical examples of the emission from the three types of systems
included in Categorys 4 and 5M. They are compared with normal standard systems for the
same environments.
Table 1. ISO 12944-2 Classification of Environments
Corrosivity
Category
C1: Very low
C2: Low
C3: Medium
C4: High
Examples of Typical Environments in a Moderate Climate
Exterior
-
Atmospheres with low pollution and
dry climate.
Mostly rural areas.
Urban and industrial atmospheres,
moderate sulphur dioxide pollution.
Moderate coastal climate.
Industrial and coastal areas.
Interior
Inside, heated buildings with
neutral atmospheres, and relative
humidity below 60%, e.g., offices,
shops, schools and hotels.
Unheated buildings where
condensation may occur, e.g.,
depots and sports halls.
Production rooms with high
humidity and some air pollution,
e.g., food processing plants,
laundries, breweries, and dairies.
Chemical processing plants, and
boat yards over seawater.
-
C5: Very high
Industry and areas high humidity
(industry)
and aggressive atmosphere.
C5M: Very high Marine coastal, offshore, areas with
(marine)
high salinity.
Table 2. Selected ISO/CD#1 12944-5 Coating Systems for Corrosion Categories
C4 and C5M, Aggressive Industrial and Marine Areas
Corrosivity
Category &
Paint System
Dry Film
Thickness
Solvent
Emission
451
Corrosion Protection and Monitoring
Number
(g/sqm)
(μm)
66
Chlorinated rubber primer
60
66
Chlorinated rubber primer
60
72
C4 - 5
Acrylic intermediate
60
60
72
Acrylic Top Coat
240
276
Standard solvent-borne system
Ethyl zinc silicate primer
67
80
C4 - 16
Waterborne epoxy intermediate
5
60
60
5
Waterborne epoxy top coat
77
200
Hybrid system
13
High solids epoxy primer
80
13
C4 - 7
High solids epoxy intermediate
80
80
13
High solids epoxy top coat
240
39
High solids system
2
Waterborne zinc epoxy
40
3
Waterborne acrylic intermediate
60
0.5
C4 - 10
Waterborne acrylic top coat
50
50
0.5
Waterborne acrylic top coat
200.
6
Waterborne system
67
Ethyl zinc silicate primer
80
52
Epoxy intermediate
60
43
C5 - 10x 1)
Epoxy intermediate
50
50
40
Polyurethane top coat
240
175
Standard solvent-borne system
40
21
High solids zinc epoxy primer
150
24
C5M - 7x
High solids intermediate
50
0.5
Waterborne acrylic top coat 2)
Hybrid system
240
∼46
C5M - 3
High solids epoxy primer
24
150
150
24
High solids epoxy top coat
300
48
High solids systems
40
2
Waterborne zinc epoxy
70
6
Waterborne epoxy intermediate
70
6
C5M - 6x
Waterborne epoxy intermediate
60
0.6
Waterborne acrylic top coat 3)
Waterborne system
240
∼14
Corrosivity Categories:
C4: Industrial and coastal areas
(ISO/WD 12944-2)
C5M: Marine areas with high salinity and corrosive areas
All systems are claimed in ISO/CD#1 12944-5 to have an expected
medium durability in the respective corrosion categories.
Notes:
1). In ISO/CD#1 12944-5 this system is a 5 coat system
2). In ISO/CD#1 12944-5 the top coat is chlorinated rubber
3). In ISO/CD#1 12944-5 the top coat is polyurethane
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Kronborg Nielsen
Figure 1. Solvent emission from coating systems for industrial and coastal areas
Figure 2. Solvent emission from coating systems for marine and aggressive areas
The paint systems mentioned are only examples of many possible combinations having the
same performance. However, a considerable decrease in emissions is possible when one of the
waterborne or high solids/low VOC systems is selected.
453
Corrosion Protection and Monitoring
EXPERIENCE WITH WATERBORNE AND LOW VOC COATINGS
Application
High solids paints are normally applied without problems if the application is carried out
with heavy-duty airless spraying equipment, e.g., pumping at least 45:1. On the other hand,
waterborne coatings are more delicate because of their nature and their weather window, i.e.,
limited by temperature and relative humidity during application. The most frequently observed
mistakes are as follows
• Mixing waterborne paint with thinners which results in clogging spray equipment.
Preventive action includes cleaning spraying equipment carefully with thinner followed
by fresh water before the waterborne paint is introduced.
• Applying waterborne paint on cold steel and/or at low temperatures which results in
insufficient curing and poor resistance. Preventive action includes painting indoors in
ventilated, heated facilities (ambient temperatures of 5°C are sufficient for waterborne
acrylics), or making covers with heating, if possible. If not possible, low VOC systems
should be used.
• Exposing newly waterborne painted objects to frosty weather which results in the
cracking of coating films. Preventive action includes keeping coated objects away from
frost for at least 24 hours after the application of waterborne paints.
• Using waterborne paint indoors in areas without ventilation which results in runners and
slow drying. Prevention action includes allowing sufficient ventilation to extract the
water liberated during the application; for 20 l of paint more than 10 l of water have to
be removed in the form of vapor. The ventilation requirement is at least 75 m3 air/l paint
at 20°C and 50% relative humidity.
The above errors can be overcome by changing the painting procedures and by training the
applicators.
Performance
Paint manufacturers, specifiers and societies use a number of accelerated test methods to
predict the lifetime performance properties of coatings. In particular, the corrosion resistance of
coating systems is important. Some of the test methods used are
• Salt spray test (ISO 7253, ASTM B-117),
• Continuous condensation test (ISO 6270), and
• Prohesion chamber cyclic test (No ISO yet).
The salt spray test has been used for a number of years. However, its inability to
demonstrate a direct relationship between the resistance of organic coatings to the action of salt
spray and resistance to corrosion in other (natural) exterior environments is now acknowledged.
Actually, for a number of years, the salt spray test averted the introduction of waterborne
coatings. The low dry-film thicknesses of systems with these coatings (dft 50 -100 μm) fail
quickly in the salt spray test, but perform well at exterior exposure sites, as demonstrated by
Andrews et al. [4]. Salt spray tests are, however, valuable in comparison situations for high dft,
high solids, and solvent-borne systems.
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Kronborg Nielsen
In the prohesion chamber cyclic test, the coated panel is dried between the salt spray and
the ultraviolet radiation cycles. Thereby more reliable test results, i.e., results with correlations
to genuine exterior situations, are obtained.
Waterborne and low VOC coating systems have performed satisfactorily in a number of
accelerated cyclic tests in comparison with standard high solvent systems. The results have
been confirmed by exterior exposure.
PERFORMANCE IN CORROSION CATEGORY 4: INDUSTRIAL AND COASTAL
AREAS
Waterborne coatings have been tested and applied since 1984 on dry cargo containers for
carriage by sea, and the obtained experience has only been positive [1]. Today more than 3000
containers are in service, either with waterborne systems only (zinc epoxy/epoxy/acrylics), or
with a hybrid system made of solvent-borne (zinc) epoxy prime coat(s) followed by waterborne
acrylic top coats.
In Kuwait, the Kuwait Institute for Scientific Research has compared 11 coating systems in
two laboratory cyclic tests and at five exposure sites in Kuwait [5]. The sites selected for the
outdoor panel exposure had different environmental parameters of the industrial area and were
at varying distances from the Arabian Gulf. Inorganic zinc silicate (IOZ)/epoxy/polyurethane
systems and waterborne acrylic systems performed better than other systems, both in the
laboratory (3000-hour test) and at the sites (two year exposure).
Also in the Gulf area, the petrochemical market has discovered that the exterior of storage
tanks can be advantageously finished with an acrylic waterborne top coat. The traditional
system of zinc epoxy/epoxy/polyurethane is occasionally being replaced by zinc epoxy/high
solids epoxy/waterborne acrylic on tank farms. Although the gloss retention of the
polyurethane top coat is slightly superior, the new system has three distinct advantages:
• Cost. Calculated per square meter, the IOZ/high solids epoxy/waterborne coating
system is 10-15% cheaper than an IOZ/epoxy/polyurethane system (dft 225 μm).
• Application. A one-pack product like a waterborne acrylic is easier to apply than a
two pack polyurethane. Additionally, pot-life problems are avoided.
• Lower VOC-emission. VOC emission is reduced by 70% (Fig. 3).
The change from polyurethanes to waterborne acrylics has also been introduced in the UK.
On the exterior of tanks at a major oil terminal near the coast of Southhampton, a test program
has been concluded, with the result that a number of tank externals were primed in 1995 with
high solids (low VOC) epoxies and finished with waterborne acrylics [6].
Also in the UK, waterborne acrylics have been used, among others, for
• British Gas' gasholders. Many of the gasholders seen near the major cities are now
maintained with waterborne coatings with excellent performance results.
• Maintenance of ships' interiors. Maintenance of bulkheads, deckheads, engine
rooms, etc., is normally carried out with alkyd paint onboard or during dry-docking.
However, the Royal Fleet Auxiliary (RFA) recognized in 1990 that a switch to
455
Corrosion Protection and Monitoring
waterborne coatings has several advantages over solvent-borne systems. First there
is no fire risk. Fire and explosion risks from paint are eliminated, during both
application and storage. Additionally, waterborne paints have low flame spread
characteristics once they are applied. Second, no thinner is required, water is the
diluent and is also used for the cleaning of application tools. Third, there is less odor
and less inconvenience. Waterborne coatings are popular for use indoors because
they allow other trades to work in the immediate vicinity, so painters do not have to
work night shifts. Fourth, waterborne coatings have good protective properties. The
performance of the coatings in respect to gloss and color retention, and protection is
highly satisfactory.
The number of RFA-ships using waterborne coatings is now nearly 25.
Figure 3. Solvent emission from coating systems for marine and aggressive areas
CORROSION CATEGORY 5: MARINE AND AGGRESSIVE AREAS
In 1992, the three major Norwegian petrochemical and offshore operators, Statoil, Saga
Petroleum and Norsk Hydro, introduced a prequalification test for coatings used on offshore
structures. The prequalification test [7] included, among other things, coating systems for
structural steel, exteriors of vessels and tanks, piping (not insulated), valves, and steel in
noncorrosive areas; all decks; and submerged steel.
The tests used for the prequalification and the acceptance criteria are listed in Table 3.
High solids/low VOC coating types like reinforced polyester and solvent-free epoxies have
been specified for offshore decks and have performed very satisfactorily over the years. The
testing confirmed their good performance. Similarly, the good experience with solvent-free or
low solvent epoxies in submerged areas has been verified.
456
Kronborg Nielsen
Table 3. Prequalification Testing by Statoil R-SP-630
Test
Salt spray
Method
ISO 7253
Duration
8000 h
Acceptance Criteria
Max disbonding 5 mm
(ISO 4628).
Blistering: not visible
(ISO 4628).
Remarks
Adhesion: 2,0 MPa
(ISO 4624) and
maximum 50%
reduction from
original value.
Condensation
chamber
Weatherometer
Cathodic
disbonding
ISO 6270
8000 h
ASTM 623-89
2000 h
ASTM G8
28 d
Max disbonding 5 mm
Overcoatable
without
mechanical
pretreatment.
Only for non
submerged
coatings.
Only for
submerged
coatings.
Among systems tested for topsides are those mentioned in Table 4.
In general all three systems performed equally overall in the prequalification tests. The
IOZ/vinyl topside system in Table 4 is a system that has been used since the 1960s, especially
by American-owned offshore operators, and is still performing well in, for example, the Gulf of
Mexico and the North Sea.
During the 1980s the IOZ/HB epoxy/PU system gradually took over. Important reasons
were the lower cost of the paint, and the possibility of applying coatings in higher film
thicknesses, thereby reducing the number of coats and, as a result, the application costs. The
change was not caused by environmental pressure.
Table 4. Systems Tested for Topsides
Standard IOZ/Vinyl System
Zinc silicate primer
Vinyl tie coat
IOZ/HB Epoxy/PU System
60 m Zinc silicate primer
25 m Epoxy tie-coat
IOZ/HB Epoxy/WB Acrylic
System
60 m Zinc silicate primer
60 m
25 m Epoxy tie-coat
25 m
457
Corrosion Protection and Monitoring
3x Vinyl interm/finish 215 m HB epoxy, LTC*
165 m
Polyurethane
50 m
Total
300 m Total
300 m
Solvent emission: 450 g/m2
Solvent emission: 257 g/m2
Paint price/m2, index: 100
Paint price/m2, index: 86
*LTC: Low temperature curing (-10°C - 20°C)
HB epoxy, LTC*
165 m
WBorne acrylic
50 m
Total
300 m
Solvent emission: 217 g/m2
Paint price/m2, index: 83
In the 1990s, with greater focus on environmental issues, systems employing low VOC
products are becoming more important. Also the isocyanate curing agent in the polyurethanes is
being put under surveillance (officially and unofficially) in, for example, Great Britain and
Norway. Therefore, the IOZ/HB epoxy/WB acrylic described in Table 4 is a step towards top
side systems formulated with respect for environmental concerns both for the applicator and his
surroundings. It is worth noticing that the IOZ/HB epoxy/WB acrylic top-coat system has both
the lowest emission and the lowest cost.
Since 1992 a major Norwegian offshore operator, Amoco Norway, has employed this zinc
silicate/HB epoxy (amide type)/waterborne acrylic system both for offshore maintenance
painting and for new construction in manufacturing units. Occasionally, zinc-rich epoxies are
also used as prime coats. The operator decided to change to the low VOC system after a
thorough evaluation of the anticorrosive properties, and the reduced solvent, low molecular
amine and isocyanate exposure of the applicators. After the initial adjustment of procedures, the
result has been very positive from environmental, operational, and economical points of view
[8].
The low VOC, high solids epoxy mastic has for a long time been used for maintenance on
ships' topsides, superstructures and exposed steel structures, especially on power-tool cleaned
surfaces. They have replaced traditional chlorinated rubber and alkyd systems, and their
success is again due to the possibility of applying high film systems in a few coats; a
comparably lower cost per square meter when applied at the same dry film thickness with
recognized better protective performance [9]. The epoxy mastic may be top coated with
waterborne acrylic coatings to obtain gloss and better color retention while keeping the solvent
emission down.
An exceptional coating test object is situated on the Thames near the Tower Bridge: The
HMS Belfast. This World War II warship was painted with waterborne acrylics in 1993, and
the surface condition is excellent [10,11].
SUMMARY
Legislation is gradually forcing painting contractors and shipyards to employ paint systems
with low solvent emissions (i.e., low VOC). Among the low VOC paints on the market are high
solids epoxies and waterborne acrylics. The introduction of these coatings in less aggressive
environments is already in place. However, the use of waterborne coatings in particular has
been more cautious in highly corrosive environments, e.g., ships' topsides, petrochemical
installations, and offshore. The reluctance is mainly originating from limited experience with
their long-term performance.
However, official tests, practical demonstrations and case histories in aggressive areas have
458
Kronborg Nielsen
now shown that waterborne acrylic finishes combined with high solids/low VOC prime coats
are as fully resistant to corrosion as any alternative. Similarly, high solids epoxies,
polyurethanes and polyesters have also demonstrated their performance in corrosive
environments. Therefore, on sound economic and environmental bases, coating contractors,
specifiers and shipyards can meet the upcoming VOC legislation with environmentally
acceptable and resistant coating systems.
REFERENCES
1. S. Nysteen, Surface Coatings International, July 1994, p. 311.
2. European Union, Proposal for a council directive (EEC) on the limitations of the emissions
of organic compounds due to the use of organic solvents in certain processes and industrial
installations, April 1994.
3. ISO 12944. Secretariat of ISO/TC 35/SC 14, N76, N37 and N94.
4. J. Andrews et al., Cleveland Society for Coatings Technology Technical Committee, Journal
of Coatings Technology, October 1994, p. 49.
5. J. Carew et al., Materials Performance, December 1994, p. 24.
6. I. Walker, Petroleum Review, November 1994, p. 520.
7. Statoil, Norway, Specification for purchase, Surface Preparation and Protective Coating
Doc. no. R-SP-630, 1992.
8. T.M. Ege and H. Erikstein, Maling offshore: Bruk av vanntynnbare toppstr ksmalinger i
Nordsj en. Hvilke krav og erfaringer har man? (Painting offshore: Demands and experience
with waterborne top coats in the North Sea), Overflatedagar 95, Paper A-5, Oslo, Norway,
Teknologisk Institutt, November 1995 (In Norwegian).
9. P.K. Nielsen, J.H. Hansen, Ecology and economy in the development and use of heavy-duty
protective coatings for Steel, Corrosion Asia/94, Paper No. 1130, Singapore, National
Association of Corrosion Engineers, September 1994.
10. Lloyds List, 29 September 1994, p. 14.
11. D. Woodyard, Lloyds List, 19 April 1991, p. 5.
459
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
ANTICORROSIVE FILM-FORMING NONPOLLUTING PRODUCTS
ACHIEVED IN ROMANIA
R. Serban, N. Moga and E. Stockel
Anticorrosive Protection, Paints and Varnishes Research Institute
(ICEPALV), 49A Theodor Pallady Av. 74585, Bucharest, Romania
ABSTRACT
Present law stipulations concerning environmental protection require that modern film-forming
products should produce an as little pollution as possible. To achieve this, water should be used as a
thinner to produce waterborne anticorrosive paints for metal and emulsion paints for anticorrosive
protection of concrete. As concerns waterborne products, ICEPALV is researching products applied
by electrophoresis, and offering licenses for anaphoretic primers; and has developed air-drying
waterborne primers and paints, based on epoxy and alkyd resins. As concerns emulsion paints,
ICEPALV has researched and developed acrylic, acrylo-styrene, vinyl-acrylic, and vinylic as well as
decorative plasters for concrete protection. ICEPALV is also researching anticorrosive emulsion
paints, based on meth(acrylic) copolymer latices with vinyl esters of heavily monocarboxylic acid
containing 10 carbon atoms (VeoVa 10).
Key Words: Metal, concrete, anticorrosive protection, waterborne products, emulsion paints
INTRODUCTION
Since 1980, the ICEPALV research in the field of film-forming materials has been
directed by the need to protect the environment. The use of water as a thinner in the
development of waterborne products and emulsion paints is one of the ways to achieve this
purpose.
Water Thinnable Products
ICEPALV is concerned with developing this important group of products used especially
in the building machines industries (i.e., bodies, cases, wheels and other accessories), but also
in other industries, as for example, for the anticorrosive protection of metallic buildings.
As concerns the application of these products, they are used simultaneously in both older
application systems (i.e., immersion, flow-coating, and spraying) as well as in newer ones
(i.e., anionic and cationic electrodeposition, and autophoresis) [1]. The main waterborne
products developed by ICEPALV in Romania are presented in Table 1.
Emulsion Paints
Concrete corrosion is a complex physicochemical process. Corrosion may be spoken
about as an electrochemical phenomenon only in the case of reinforced concrete; otherwise, it
is only about the support isolation from the corrosive medium. Therefore, nowadays
especially, the outdoor painting of buildings is a necessity, having both an aesthetic and an
461
Corrosion Protection and Monitoring
anticorrosive protection function. Emulsion paints are more and more often used for these
purposes, and their advantages are well known.
A whole series of putties, primers, decorative plaster and paints based on acrylic, acrylostyrene, acrylo-vinyl, vinyl, etc. latices for concrete protection have been developed and
launched on the market by ICEPALV. Among these the most important are those mentioned
in Table 2.
Table 1. The Main Waterborne Products Developed by ICEPALV in Romania
Product
Series
Grey primer
Brown primer
Grey primer
Grey, red A-B primer
Colourless, black,
khaki primer
Grey anticorrosive
primer
Dark-black primer
7100
7101
7003
7004
7004
alkyd-phenol
alkyd-phenol
alkyd-phenol
alkyd-phenol
alkyd-phenol
Application
Method
flow-coating
spraying
immersion
electrophoresis
electrophoresis
7003
epoxy-ester
immersion
7206
acrylic
flow-coating,
immersion
7005
polybutadiene electrophoresis
7210
acrylic
Anticorrosive grey
primer
Black primer
Binder
Recommended Uses
Building-machines industry
Building-machines industry
Building-machines industry
Automotives industry
Building-machines industry
Automotives and
electrodomestic industry
The protection of gasoline
tanks and of the other accessories in the automotive
industry
Automotives industry
immersion,
Phosphatized plates
spraying
protection
Air drying primer
7352 waterborne
by brush, roller Building-machines industry
alkyd
or spraying
Air drying primer
7702 waterborne
by brush, roller Building-machines industry
alkyd
or spraying
Air drying primer
7752 waterborne
by brush, roller Building-machines industry
alkyd
or spraying
Beige primer
7385 epoxy-ester
electrophoresis Automotives industry
Black primer
7325 epoxy-ester
electrophoresis Automotives industry
Air-drying paint
7351 epoxy
by brush, roller Building-machines industry
or spraying
Air-drying paint
7751 epoxy
by brush, roller Building-machines industry
or spraying
Colourless,
8771 vinyl-polyby brush or roller Protection of paint spray
removable varnish
acetate latex
booth glass surfaces
White, removable
8772 vinyl-polyby brush or roller Protection of paint spray
paint
acetate latex
booth metal surfaces
Table 2. The Main Emulsion Products Developed by ICEPALV in Romania
462
Serban et al.
Product
Soaking primer
Series
8440
Putty for concrete
8640
LUCICRIL-half-glossy
paint for outdoors
DASIROM-half-matt
paint for outdoors
Half-matt paint for
outdoors
VEPAROM-matt paint
for indoors
VEPATIM-matt paint
for indoors
Matt paint for indoors
8513
Matt paint for indoors
8213
STROP- decorative
plaster
8411
8415
8427
8430
8630
8426
Latex
Acrylo-styrene
hydrosol
Recommended Uses
To fill in the pores and to
increase adherence (concrete,
masonry)
Vinyl-maleic
To putty the concrete before
copolymer
finishing
Pure acrylic
Outdoor finishing of buildings
copolymer
(wood, concrete, masonry)
Acrylo-styrene
Outdoor finishing of buildings
copolymer
(concrete, masonry)
Acrylo-styrene-maleic Outdoor finishing of buildings
copolymer
(concrete, masonry)
Acrylo-styrene
Indoor finishing of buildings
copolymer
(concrete, masonry)
Vinyl-maleic
Indoor finishing of buildings
copolymer
(concrete, masonry)
Acrylo-styrene-maleic Indoor finishing of buildings
copolymer
(concrete, masonry)
Vinyl-polyacetate
Indoor finishing of buildings
homopolymer
(concrete, masonry)
Acrylo-styrene
Decorative finishings for outdoor
copolymer
buildings (concrete, masonry)
NEW RESEARCH TRENDS
Waterborne Epoxy-Ester Primer Series 7301
ICEPALV has lately ended the research concerning waterborne epoxy-ester primer series
7301. It is used for the anticorrosive protection of wheels and some automotive accessories,
and it is applied as a first coat over zinc phosphate pretreated iron plate by Bonder 125
technology. Application is carried out by anionic type electrophoresis.
EXPERIMENTAL PROCEDURE
The epoxy-ester, which represents the binder of this primer, is produced from an epoxy
resin of diglycidil ether of the A bisphenol type, with a molecular weight of 900-1000, which
in the first stage is partially fatty acids reaction, with some of the hydroxyl groups left are
esterified with COOH groups from the tricarboxyl adduct and hydrolyzed according to the
following scheme:
463
Corrosion Protection and Monitoring
O
OH
unsaturated fatty
OH
acids
OH
OH
OH
245 °C
catalyst NaOH
O
COOH
HOOC
COO
CH
OOC
CH
HOOC
OH
120-150°C
OH
OH
HOOC
O OC
HOOC
COO
Epoxy resin
M=900-1000
COOH
HOOC
COO
COOH
COOH
HOOC
Tricarboxyl adduct
COOH
COO
The tricarboxyl adduct is previously achieved from fatty acids, maleic anhydride and
water, according to the reaction:
COOH
COOH
CH
COOH
HC
C
CH2
C
COOH
O
H2O
O
O
HC
C
HC
C
O
O
COOH
O
HOOC
Unsaturated
fatty acids
Hydrolized succinic adduct
125°C
Succinic adduct
220°C
COOH
O
HC
C
HC
C
O
CH
COOH
CH
COOH
Hydrolized Diels-Alder adduct
O
Diels-Alder adduct
Finally, the epoxy-ester is so reactioned to provide the achievement of an epoxy-ester
with free functional groups: hydroxyl and carboxyl, which exhibit a good water solubility,
increase the system’s stability, as well as the salt spray resistance by increasing film
adherence in an alkaline medium (due to the hydroxyl polar groups and being inert to
alkalies).
The primer also contains cosolvents (for spreading and adherence) and various additives:
wetting and dispersing agents, antifoaming agents, antioxidants, antibacterial, etc., which
provide high quality films. The pigments and extenders were so selected to resist the alkaline
medium, and to provide a high corrosion resistance, a good hiding power for the support and
a migration speed in an electric field similar to that of the film-forming.
464
Serban et al.
RESULTS AND DISCUSSION
The electrophoretic primer developed is a slightly thixotropic, grey fluid, with a medium
viscosity (under 100 P), with a nonvolatile matter content of about 40%. It is neutralized with
an alkaline base. It has an alkaline pH, being water soluble and sensitive to low temperatures
(under 10°C). It is not flammable and presents low toxicological hazards compared to the
classical products.
Characteristics of the Electrodeposition Bath
By diluting the primer with demineralized water, an electrodeposition bath is achieved
having a nonvolatile matters content of about 12.5%. The following parameters should be
kept constant: pH, conductivity, content of cosolvents and free fatty acids, degree of
neutralization and free acidity, and pigment/binder ratio. The formulation is so balanced to
provide physical stability for the system expressed by an adequate settling curve, as can be
seen in Fig. 1.
0
-10
Settling degee (%)
-20
-30
-40
-50
-60
-70
-80
-90
-100
0
2
4
6
8 10 12 14 16 18 20 22 24
Time (hours)
Figure 1. Evolution of settling degree over time
Application Conditions
•
•
•
•
•
Electric voltage (V)
Medium density of anode current (A/m2)
Application time (sec)
Bath temperature (°C)
Film curing(drying) is carried out in the oven
180°C
140 - 220
max. 20
60 -360
25-28
for 30 minutes at a temperature of
From Fig. 2, the variation of film thickness according to application time, at the
application voltage (180 V), as well as at breakdown voltage (50 V) can be seen.
Film Characterization
465
Corrosion Protection and Monitoring
Between 20 and 30 µm films are achieved with a uniform appearance, free of surface
faults (i.e., pinholing, cratering). The mechanical characteristics are very good:
•
•
•
•
Cross-cut adherence (mm)
Erichsen elasticity (mm)
Impact resistance (1 kg/cm)
Flexibility (mm)
1
4 min.
30 min.
1
The films corresponded from the point of view of corrosion resistance, so:
• Salt spray resistance (hours)
192
6
4
192
absent
21
absent
Blistering (note), (min.)
Rust spreading (mm, max.)
• Water resistance by immersion (hours)
Blistering
• Humidity resistance (days)
Surface alteration
The throwing power was determined on 24-cm samples, and their values are presented in Fig.
3.
In conclusion, from the short presentation of the epoxy-ester electrophoretic primer
series 7301, it may be noticed that this product, used for the anticorrosive protection of some
parts and units in the automotive industry, corresponds to the present requirements of the
Romanian industry.
35
Thickness (um)
30
25
20
15
50 V (max. 25 minutes)
10
180 V (max. 5 minutes)
5
0
0
5
10
15
20
25
30
Time (minutes)
Figure 2. Variation of film thickness with the application time
466
Serban et al.
Throwing power (um)
30
20
10
0
0
3
6
9
12
15
18
21
24
Distance (cm)
10 20 30 40 50 60 70 80 90 100
Figure 3. Evolution of throwing power with distance
Anticorrosive Primer Paints Based on VEOVA 10 (Meth) Acrylate Latices
ICEPALV has been researching anticorrosive primer paints based on VeoVa 10
(meth)acrylate latices. Various studies [2,3] demonstrated that the incorporation of VeoVa
monomers in (meth)acrylic copolymers improves the chemical and, especially, the water
resistance of the latex films. The bulky aliphatic entity gives the copolymer a high
hydrofobicity, an excellent UV resistance and also a good alkali resistance by protecting it
from saponification [4].
EXPERIMENTAL PROCEDURE
The latices were obtained by the copolymerisation of VeoVa 10 and 2 ethylhexyl
acrylate (which also contributes to good water repellency) as soft monomers with
methylmethacrylate as a hard monomer.
Performing the polymerization essentially in the absence of colloids and in the presence
of a minimum quantity of surfactant with phosphate groups (e.g., organic ester phosphate:
REWOPHAT E 1027 - REWO, Germany), whilst at the same time carefully adding the paint
formulation additives (i.e., coalescing agents, thickener, and dispersant) and pigments and
mineral fillers, films possessing intrinsically good barrier properties can be achieved. The
best results were obtained with monomer compositions falling in the shaded area of Fig. 4.
In the primer paint formulations zinc, phosphate was introduced. It proved to be an
anticorrosive and nontoxic pigment. It appears to pack in the film in a manner which presents
a high resistance to the passage of water molecules and salts and an anticorrosive efficiency
similar to zinc chromate, in long-term exposure tests. The primer paints were formulated at
two different PVC: 20% and 30%, at basic pH. The evolution of the viscosities was
essentially unchanged after 8 months of storage, thus demonstrating the good stability of
these primers. They are being kept under observation.
467
Corrosion Protection and Monitoring
RESULTS AND DISCUSSION
Some Mechanical Properties
Adhesion was evaluated both on concrete (by a pull-off test) and on metal (by a cross-cut
test). The improvement of adhesion property is directly related to the amount of VeoVa 10
from the copolymer composition indifferent to the PVC (Fig. 5).
Due to the good metal adherence, Erichsen elasticity could be evaluated. It showed the
same evolution according to the quantity of VeoVa 10 from the latex as the former case. The
paints with lower PVC, have higher elasticities (Fig. 6). The thickness of the analyzed primer
paint films was about 100µm.
Water Resistance of Latices
Concerning the water-spot resistance test (100 µm latex films, 24 hours), Fig. 7
demonstrates clearly, the positive contribution of the increase of proportion of VeoVa 10
from the copolymer composition at Tg = constant = 15°C. It can be noticed from Fig. 7, that
the VeoVa latices have superior values when compared to acrylo-styrene latices, which have
about 2.
Water Vapor Permeability of Primer Paints
The studies of Geelhaar and Melan [5] have shown that an emulsion film absorbs 10 to
100 times greater amounts of moisture than normal solvent cast films from alkyds, and
reactive and crosslinked polymers. Continuous films produced by solvent-borne coatings,
absorb water in the polymer matrix by a very slow diffusion process. The water absorption of
emulsion films is via microcapillaires between coalesced particles.
From Fig. 8, it can be noticed that primer paints based on VeoVa 10 (meth)acrylate
latices have lower water vapor permeabilities than acrylo-styrene emulsion paints, at the same
PVC level (e.g., 2.9 g/100cm2/100µm/day at PVC = 20% and 3.5g/100cm2/100µm/day at
PVC = 30%). It can be also noticed that an increase of the VeoVa quantity from the latex up
to 55-60% leads to a great decrease in permeability. As was expected, if the PVC increases,
water vapor permeabilities also increases.
Corrosion Resistance
The primer paints studied were applied on panels of concrete with a thickness of about
100µm, and after 7 days of drying they, were salt-spray tested in order to simulate a marine
atmosphere, one of the harshest climates. After 400 hours of exposure, the films with 50-60%
VeoVa 10 in latex were unchanged. The others exhibited about 10% blistering. The test is
continuing. The same type of panels were also exposed outdoors in an industrial climate, and
after 8 months of exposure, the films were unchanged.
In conclusion, the primer paints based on VeoVa 10 (meth)acrylate latices, containing at
least 50-60% VeoVa 10 monomer and stabilized by a phosphate based surfactant and without
colloids, provide good qualities as an anticorrosive protection for concrete even in a marine
environment. In addition, they exhibit resistance to flash rusting and early rusting as well as
good metal adherence, qualities which convinced ICEPALV to keep on testing these primer
paints for the anticorrosive protection of metals, too.
468
Serban et al.
Adhesion
10
8
6
PVC= 20 %
4
PVC= 30 %
2
0
20
30
40
50
60
70
80
VeoVa 10 concentration (%)
Figure 4. Optium copolymer compositions
Erichsen
elasticity
(mm)
Tg = 15 oC; 0 - bad,
10 - excellent
Figure 5. Influence of VeoVa 10
concentration
on primer paint adhesion
10
Water
spot
rezistence
8
7
8
6
6
5
PVC= 20 %
4
4
PVC= 30 %
2
3
0
2
20
30
40
50
60
70
VeoVa 10 concentration (%)
Tg = 15°C
Figure 6. Influence of VeoVa 10
concentration on primer paint
elasticity on metal
80
20
30
40
50
60
70
80
VeoVa 10 concentration (%)
Tg = 15 oC; 0 - film completely white,
10 - film unaffected
Figure 7. Water spot resistance of
latices
469
Corrosion Protection and Monitoring
W a te r v a p o u r
p e rm e a b ility
(g / 0 .1 m m /
100cm / day)
2
P VC = 2 0 %
P VC = 3 0 %
1 .6
1 .2
0 .8
0 .4
0
10
20
30
40
50
60
70
Ve o Va 1 0 c o n c e n tra tio n (% )
Tg = 15 oC
Figure 8. The effect of VeoVa 10 concentration on the
water vapour permeability of primer paints
CONCLUSIONS
ICPALV is concerned with developing some new electrophoretic products of the latest
generation (i.e., cataphoresis) which together with other measures (e.g., the use of galvanized
sheet, waxes and protection products, and good and severe services) will increase the storage
life of the new types of automotive and metallic surfaces, in general.
As concerns emulsion paints, ICEPALV is concerned with developing new products
with higher durability for concrete protection and new anticorrosive primer and emulsion
paints for metal protection.
REFERENCES
1.
2.
3.
4.
5.
470
C. Robu, XV FATIPEC Congress, Amsterdam, June 1980.
M. Slincks and M.F. Daniel, PPCJ April, 1995, pp. 28-29
M. Slincks and M.S. Sonderman, XXI FATIPEC Congress, Amsterdam, 14-18 June 1992.
C. Bondy and M.M. Coleman, JOCCA, 1970, 53, p. 555.
H. Geelhaar and M. Melan, 13e AFTPV Congressbook 147, 1979.
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
CATHODIC PROTECTION UNDER DISBONDED COATINGS
OF 56 INCH GAS PIPELINE ALONG THE KANGAN-SHIRAZ
M. Pakshir
Department of Materials Science and Engineering
Shiraz University, Shiraz, Iran
ABSTRACT
The present study investigates the disbonding phenomenon according to the British Gas Method
PS/CW6 and the ASTM-G8 standard. This investigation showed that the best protection for
disbonded buried pipelines under cathodic protection would be achieved if the applied potential was
kept -950 to -1000 mV with respect to the Cu/CuSO4 half-cell rather than the usual value of -850 mV
(Cu/CuSO4) which is currently used.
Key Words: Pipeline, coating, soil, disbonding, cathodic protection
INTRODUCTION
Underground corrosion of buried steel is a major problem in the oil and gas industries
[1]. A good practice in modern pipeline corrosion control work comprises the use of good
coatings, in combination with cathodic protection as the main lines of defence [2], and better
current distribution is accomplished by using an insulating coating [3]. Therefore, attempts to
control pipeline corrosion rely on the use of coating materials with the reasoning that if the
pipeline metal is isolated from contact with the scratches from soil, no corrosion could occur.
However, the insulating coatings must be free of any defects such as scratches or pinholes.
Pitting associated with such defects in the coating, i.e., Holidays, and subsequent disbonding
have been observed on pipelines which are nominally cathodically protected [4,5].
Fessler [6] suggested that stress corrosion cracking on buried, coated pipeline tends to
occur under disbonded coatings near small pinholes or holidays. It was deduced that a
disbonded coating acted as a shield to the cathodic protection current and caused potential
gradients under the coatings [7]. Aqueous displacement as a possible mechanism of the
cathodic disbondment of protective organic coatings was suggested by Evans [8] in terms of
the ability of an alkaline solution to creep over the metal surface and displace the organic
coating. Bolger and Micheals [9] argued that displacement of coatings from a metal’s surface
is promoted by pH values far removed from the isoelectric point of the surface oxide so that
the oxide has a greater affinity for the water than for the organic materials.
EXPERIMENTAL PROCEDURE
471
Corrosion Protection and Monitoring
Soil Analysis
Since the types of soil differ along the buried pipeline, four types of soil were taken from
different sites as follows:
•
•
•
•
Site 1 - 76 km Kangan - Shiraz
Site 2 - 298 km
Site 3 - 151 km
Site 4 - 36/900
The soil samples obtained from each site were taken from areas adjacent to the pipeline
and from the same depth as the pipeline was buried, to evaluate the chemical characteristics
of the soil. A soil-distilled water slurry was obtained by mixing 50 g of soil with l00 ml of
distilled water. The pH of the soil was determined using a special pH electrode cup for soil.
The slurry was placed in the electrode cup and the pH value was read directly from the meter.
Conductivity tests were conducted on the slurry using the Wenner four-electrode
method, a procedure similar to ASTM G57-78. The concentration of soluble materials was
evaluated from the slurry filtrate by saturation and flame photometric techniques.
Table 1 shows the soil analysis for different sites along the 56 in. pipeline. Soil Type 4
with a pH of 8.44 and an Ec of 1 1.52 was chosen as a saline alkaline-type soil which is
representative of the southern part of Iran. The results were compared with soil taken from
the Northern part of the country, which is acidic. It was decided to make an acidic soildistilled water slurry using the same procedure (Table 2).
Table 1. Characteristics of Saline Alkaline-Type Soil
NO3
CO3
HCO3
SO4
Cl
Na
K
Mg
Ec
Ca
pH
Soil
Type
(mmhos/cm)
33
0
146.5
71
1036
15.5
18.5
68
516
2.15
7.56
1
26
0
148.2
106.5
1728
75.5
20.5
49.5
560
2.70
7.43
2
37
0
341.6
461.5
960
980
17
113.7
320
4.87
7.79
3
6.4
0
1753
770.5
1536
1006.5
105
121.5
528
11.52
8.44
4
Ec (mmhos/cm)
pH
Soil
Table 2. Characteristics of Acidic-Type Soil
NO3
CO3
HCO3
Cl
SO4
Na
K
Mg
Ca
Type
188
0
0
Tested Material
472
2912.5
4467
1398.6
868
754
540
28.42
4
acidic
Pakshir
Test specimens were cut and flattened, and then they were stamped with an identifying
number from the existing new pipeline used in the Shiraz gas industry according to ASTMA53 grade B having the following composition:
0.30% C, 1.20 % Mn, 0.05 % P, 0.06 % S
Each specimen was hand brushed with a wire brush, and soaked in a solvent to remove
the cutting oil.
Test Procedure
In order to investigate the potential gradient under the disbonded coatings
experimentally, special suitable cells were constructed ( Fig. 1).
Steel samples 200 x 160 mm were coated with a cold tape used in Shiraz gas industries
called NITO. The tape consisted of a polyethelene backing and a thermoplastic adhesive.
For better adhesion, NITO primer was used. The technique for measuring the potential under
the disbonded coatings was based on British gas method PS/CW6 in which backside
electrodes were prepared by inserting a 500 mm length of a 5 mm outside diameter PVC
tubing filled with agar and KCI. They were then filled with a sintered glass plug through four
predrilled holes in the back of the steel plate. The holes were positioned in a line so as to
give eight positions 5 mm apart at distances from 10 to 40 mm from the center of the holiday.
Figure 1. Experimental apparatus
On the coated side of the specimen, a PVC pot (150 mm in diameter and 160 mm high)
was attached by silicon rubber to the coating. This pot contained 1.5 l of prepared simulated
473
Corrosion Protection and Monitoring
solution. A platinum wire was used as the counter electrode. The steel plate was connected
to a voltage supply by a wire tapped into it and insulated. The holiday potentials were
controlled by a voltage regulator. The voltage was measured using a reference electrode and
a digital multimeter. The holiday potentials were set at values of -780, -920 and - 1200 mV
(SCE).
In order to extract solutions from beneath the disbonded coating, a piece of polyethylene
tubing was inserted through the steel in a fation similar to that used for the electrodes. A
thinner length of tubing was then be inserted through the first hole to allow solution to be
extracted with a syringe.
The pH of the extracted solution was measured using a pH paper, and the chloride
solution was measured using a Ag/AgCl microprobe. However, at distances further than 10
cm from the holiday, little success was achieved in extracting any crevice solution due to the
thinness of the layer of electrolyte. Along the crevice, most of this liquid was absorbed into
the adhesive. In this region, the pH values of the solution were measured directly from
moisture that was found on the steel underneath the backing after the coating had been
removed.
Polarization Study
In order to obtain the potential-pH diagram, a polarization study was carried out.
Specimens 1 x 1 cm in dimensions were cut from the original pipe and mounted in a
specimen holder so that 1 cm2 of the steel was exposed, and the polished specimen was
placed in a corrosion cell with a platinum counter electrode and a lugging probe. The cell
was filled with an already made. simulated solution of pH 8.44. To simulate the alkaline
environments found beneath the disbonded coating, the pH of the solution was increased by
adding various amounts of NaOH.
RESULTS AND DISCUSSION
Using a backside electrode as a special technique, the potentials under the disbonded
coating were measured as a function of distance away from the holiday (Figs. 2 and 3). Also,
using a catheter arrangement, the pH and the concentration of the chloride solution under the
disbonded coating could be determined (Tables 3 and 4).
In this investigation, three holiday potentials were chosen: the potential of -780 mV
(SCE) was chosen since it represents the minimum cathodic current density in order to
polarize the pipe to 850 mV Cu/CuSO4, and the holiday potential of -1500 mV (SCE) since it
represents the overprotection potential. A test temperature of 40°C was used to simulate the
conditions of the hottest part of the soil. As can be seen from Tables 3 and 4, the
concentration of chloride ions in the crevice did not vary significantly from the bulk solution,
and also, the concentration of chloride ions in the crevice was not a function of the holiday
potential.
Since the solution pH beneath the disbonded coating was thought to ncrease, polarization
tests were carried out in the alkaline range of pH. Two distinct points can be seen in Figs. 4
and 5: the interaction of polarization current with the potential axis. i.e., where the current
density was zero and was representative of the corrosion potential (Ecorr); and the start of the
decrease in current density while the potential increased and was representative of the
474
Pakshir
protection potential (Ep). Therefore, in the potential-pH diagram, the immunity-general
corrosion boundary represented by the corrosion potential points and the general corrosioncomplete passivation boundary could be represented by protection potential points (Figs. 6
and 7).
Potential vs. distance variation
for alkaline-saline soil
Potential vs. distance variation
for acidic type soil
Table 3. Experimental Results for Alkaline-Saline Soil and the Holiday Potentials
of -0.780, -0.920 and -1.5 V
Cl
(ppm)
Crevice
Tip/pH
Time
(hour)
805.2
798.5
752.1
794.3
10.48
11.02
10.63
10.83
40
-0.604
-0.603
-0.604
-0.707
30
-0.606
-0.605
-0.709
-0.714
20
-0.603
-0.710
-0.716
-0.719
10
-0.715
-0.724
-0.731
-0.736
0
-0.780
-0.780
-0.780
-0.780
513
862
1103
1223
791
812.3
783.4
800.8
751.5
818.1
759.5
791.2
11.08
10.54
10.96
10.61
10.41
10.52
11.02
10.75
-0.611
-0.611
-0.610
-0.673
-0.612
-0.611
-0.610
-0.743
-0.608
-0.607
-0.674
-0.680
-0.605
-0.607
-0.722
-0.747
-0.616
-0.675
-0.685
-0.691
-0.609
-0.717
-0.785
-0.828
-0.704
-0.728
-0.751
-0.778
-0.725
-0.823
-0.907
-0.991
-0.920
-0.920
-0.920
-0.920
-1.500
-1.500
-1.500
-1.500
363
605
770
863
143
325
297
335
Distance from Holiday (mm)
Holiday
Potential
(Volt)
-0.780
-0.920
-1.580
Table 4. Experimental Results for an Acidic-Type Soil and the Holiday Potential
of -0.780, -0.920 and -1.5 V (cm/CuSO4)
475
Corrosion Protection and Monitoring
Cl
(ppm)
Crevice
Tip/pH
2850.5
2920.2
2860.4
2898.7
29.20.2
2870.8
2932.7
2946.5
2873.6
2950.2
2873.8
2932.6
10.33
10.28
10.47
10.86
10.52
10.63
10.21
10.44
10.73
10.96
10.85
10.91
40
-0.605
-0.605
-0.604
-0.655
-0.600
-0.601
-0.600
-0.628
-0.611
-0.612
-0.611
-0.686
30
-0.609
-0.607
-0.665
-0.673
-0.603
-0.604
-0.632
-0.641
-0.608
-0.608
-0.687
-0.701
Polarization curve for alkalinesaline soil
476
Time
(hour)
Distance from Holiday (mm)
20
-0.608
-0.676
-0.681
-0.692
-0.610
-0.630
-0.647
-0.652
-0.614
-0.695
-0.724
-0.750
10
-0.690
-0.704
-0.711
-0.715
-0.680
-0.708
-0.729
-0.744
-0.701
-0.807
-0.891
-0.945
0
-0.780
-0.780
-0.780
-0.780
-0.940
-0.920
-0.920
-0.920
-1.500
-1.500
-1.500
-1.500.
623
1038
1320
1482
421
719
911
1016
215
359
455
505
Holiday
Potential
(Volt)
-0.780
-0.920
-1.500
Figure 5. Polarization curve for acidic soil
Pakshir
. Pourbiax diagram extracted from
polarization curve for alkalinesaline
urbiax diagram extracted from
larization curve for acidic soil
As can be seen from Figs. 4 and 5, a significant shift of the crevice potentials to more
positive values occurred as the distance from the holiday along the crevice increased. Also,
the largest potential drop occurred near the holiday. The potential gradient appeared to
decrease at distances further along the holiday. The major differences between the three
potentials applied is the disbondment time, i.e., as the holiday potential became negative, the
disbondment time decreased. For example, for an alkaline-saline soil, it took two months to
disbond the coating at a holiday potential of -780 mV (SCE), while at a potential of -1500
mV (SCE), it only took 14 days. The crevice tip potential for the three applied holiday
potentials varied between -673 and -734 mV (SCE), and the pH of the tip of the crevice
varied between 10.41 and 11.08. Therefore, neither the crevice potential nor the crevice tip
pH were a function of the applied holiday potentials.
Polarization tests enabled an experimental potential-pH diagram to be constructed for the
steel exposed to conditions which simulated those formed beneath the disbonded coatings
(Figs. 6 and 7). Superimposed on these diagrams are the crevice tip environments which
were determined from the corresponding cathodic disbondment test. Hence, if the -crevicetip potential lay between -673 and -734 mV (SCE), and the crevice-tip pH lay between 10.41
and 11.08, then according to Fig. 6, only a small part of the blackened area lay within the
general corrosion area, whereas a large part of it was in the complete passivation area.
Hence, maintaining the steel exposed along the crevice in a passive state depended only on a
continued high PH. Consequently, the condition of the crevice tip determined the corrosion
behavior of the metal beneath the disbonded coating. In another words, by knowing the
condition of the crevice tip on the potential-pH diagram, one can estimate the corrosion and
noncorrosion condition beneath the disbonded coating.
477
Corrosion Protection and Monitoring
Figures 8 and 9 show that the condition required for the occurrence of corrosion beneath
a disbonded coating is for the crevice potentials to be within the general corrosion range of
about -695 to -845 mV (SCE). Therefore, as cathodic disbondment occurs, the crevice
potentials place the steel at some distance along the crevice into a region of general corrosion.
This can be seen by the surface morphology of the specimens after the disbonding test ( Figs.
10-15).
As can be seen from Fig. 10, all the disbonded surfaces were in the region of general
corrosion, but when the holiday potential was kept at -920 mV (SCE), the first 5 mm of the
disbonded surface was in the immunity region, from 5 to 20 mm was in the general corrosion
region and at distances > 20 mm from the holiday, the surface was in the passive state. These
situations can also be predicted by Figs. 8 and 9.
. Potential distance curve for
alkaline-saline soil
478
. Potential distance curve for
acidic soil
Pakshir
Figures 10-12. Surface morphology after disbondment for alkaline-saline type soil at various
holiday potentials
479
Corrosion Protection and Monitoring
Figures 13-15. Surface morphology after disbondment for acidic type soil at various holiday
potentials
From Fig. 11, it can be seen that in the first 20 mm from the holiday, the disbonded
surface was covered by an oxide layer which represents the general corrosion and suggests
that the metal underneath the disbonded coating was in the passive state. When the holiday
potential was kept at 1500 mV (SCE), the potential-pH diagram predicted that up to 18 mm of
the disbonded surface would be immune to corrosion, and then would show general
corrosion. In Fig. 9, the curve predicts the regions of corrosion, passivation and immunity
under the disbonded coating, i.e., for an acidic soil it predicts that at a holiday potential of 780 mV (SCE), the surface would be in the general corrosion region, at a holiday potential of
480
Pakshir
-920 mV (SCE), the first 4 mm of the disbonded surface would be in the immunity region,
and between 4 and 20 mm, the metal would be in the general corrosion region and then in the
passivated region. This prediction can be shown by the surface morphology examination
shown in Figs. 13, 14, and 15.
6. Distance from holiday vs. time
curve for alkaline-saline soil
7. Distance from holiday vs. time
curve for acidic soil
CONCLUSIONS
For an alkaline-saline soil, the occurrence of corrosion beneath a disbonded coating
requires that the crevice potentials be within the general corrosion range of -695 to -845 mV
(SCE). The crevice potentials for a holiday potential of -920 mV lie solely within the perfect
passivation region beyond a distance of about 5 mm from the holiday. However, during the
growth of the disbondment area and the consequent movement of the crevice potentials to
their final values, the crevice potentials must be in the region of general corrosion. For the
crevice potential obtained for a holiday potential of -1500 mV (SCE), the potentials must lie
within the general corrosion region from about 18 to 28 mm from the holiday.
Important conclusions which can be drawn from this investigation is that the applied
potential of -780 mV (SCE) is not satisfactory, and a holiday potential of -1500 mV (SCE) is
not representative of an instantaneous off potential which suggests that overprotection is
undesirable. However, when the applied potential is -920 mV (SCE), only small parts around
the holiday are in the general corrosion region while the rest of the surface is in the protected
state. Therefore, this potential is recommended for buried pipeline.
Also from Figs. 16 and 17, it can be concluded that for both types of soil, the disbonded
mechanisms follow similar patterns, although the disbonding rate is much slower for an
acidic soil.
REFERENCES
481
Corrosion Protection and Monitoring
2.
3.
4.
5.
6.
7.
8.
9.
1. H. Azad, M.Sc. thesis, School of Engineering, Shiraz University, 1994.
A.W. Peabody, NACE, Control of Pipe Line Corrosion, 1976.
H.H. Uhlig, Corrosion and Corrosion Control, John Wiley and Sons Inc, 1971.
C.G. Manger and R.C. Robinson, Materials Performance 20, 7, 1981, p. 46.
B.W. Cherry and A.N. Gould, Pitting Corrosion of Nominally Protected Land Base
Pipelines, Materials Performance, Aug. 1990.
R.R. Fessler, Oil and Gas Industry 74, 7, 1976, p. 81.
R.P. Fessler, Sixth Symposium on Pipe Line Research, American Gas Association,
Arlington, Virginia, P-R-1, 1960.
V.R. Evans, Corrosion and Oxidation of Metals, St. Martin's Press, New York, 1960.
J.C. Bolger and A.S. Micheals, Interface Conversion for Polymer Coatings, Weiss and
Cheever, eds., Elsivier, New York, 1968.
482
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
SYNERGISTIC EFFECT EXISTING BETWEEN AND AMONG
A PHOSPHONATE, Zn2+, AND MOLYBDATE ON THE INHIBITION OF
CORROSION OF MILD STEEL IN A NEUTRAL AQUEOUS ENVIRONMENT
S. Rajendran1, B.V. Apparao2 and N. Palaniswamy3
1
2
Department of Chemistry, G.T.N. Arts College,
Dindigul - 624 001, Tamil Nadu , India
Department of Chemistry, Regional Engineering College,
Warangal - 506 004, Andhra Pradesh, India
3
Corrosion Science and Engineering Division,
Central Electrochemical Research Institute,
Karaikudi - 630 006, Tamil Nadu, India
ABSTRACT
The synergistic effect existing between and among the sodium salt of ethyl phosphonic acid
(EPA), Zn2+, and molybdate on the inhibition of corrosion of mild steel in a neutral aqueous
environment containing 60 ppm Cl- was evaluated by the classical weight-loss method. The
formulation consisting of 300 ppm EPA, 300 ppm MoO42- and 50 ppm Zn2+ had 99% inhibition
efficiency. The mechanistic aspects of corrosion inhibition are discussed, in a holistic way, based on
the results obtained from a potentiostatic polarization study, the x-ray diffraction (XRD) technique,
and UV-visible diffused reflectance, FTIR and luminescence spectra.
Key Words: Mild steel, neutral environment, corrosion inhibition, synergistic effect, ethyl
phosphonate-zinc-molybdate
INTRODUCTION
Molybdates are among the most broadly applied inhibitors, chiefly because of their
efficacy towards both ferrous and nonferrous metals and their very low order of toxicity [1].
Molybdate can be used as corrosion inhibitor alone or in combination with other synergists
like nitrite [2], metallic cations like Ni2+, Mn2+, Zn2+ [3], azoles like benzotriazole and
tolyltriazole [4], chromate [5], amine phosphonates [6], inorganic phosphates [7], citrate and
calcium [8]. Even though several papers [1-14] have discussed the use of molybdate as
corrosion inhibitor, the mechanistic aspects of corrosion inhibition have not been studied in
detail. The present work evaluates the synergistic effect existing between and among
molybdate, Zn2+ and ethyl phosphonate by the weight-loss method. The mechanistic aspects
of corrosion inhibition were studied, in a holistic way, based on the results obtained from a
potentiostatic polarization study, the x-ray diffraction (XRD) technique, UV-visible diffused
reflectance, FTIR and luminescence spectra.
EXPERIMENTAL PROCEDURE
483
Corrosion Protection and Monitoring
Preparations of the Specimens
Mild steel specimens (0.02 to 0.03% S, 0.03 to 0.08% P, 0.4 to 0.5%Mn, 0.1 to 0.2% C
and the rest iron) of the dimensions 1 x 4 x 0.2 cm were polished to a mirror finish and
degreased with trichloroethylene for use in the weight-loss method and surface examination
studies. For the potentiostatic polarization studies, a mild steel rod encapsulated in Teflon
with an exposed cross section 0.5 cm in diameter was used as the working electrode.
Weight-Loss Method
Mild steel specimens, in triplicate, were immersed in 100 ml of the solutions containing
various concentrations of the inhibitors for a period of seven days. The weights of the
specimens before and after immersion were determined using a Mettler balance, AE-240.
Potentiostatic Polarization Study
This study was carried out in a three-electrode cell assembly connected to a bioanalytical
system (BAS-100 A) electrochemical analyzer, provided with an IR compensation facility,
using mild steel as the working electrode, platinum as the counter electrode and a saturated
calomel electrode as the reference electrode.
Surface Examination Study
The mild steel specimens were immersed in various test solutions. After two days, the
specimens were taken out and dried. The nature of the film formed on the surface of the
metal specimens was analyzed by various surface analysis techniques.
FTIR Spectroscopic Study
The FTIR spectra were recorded using a Perkin-Elmer 1600 FTIR spectrophotometer.
UV-Visible Diffused Reflectance Spectroscopy
The UV-visible diffused reflectance spectra were recorded using a Hitachi U-3400
spectrophotometer.
X-Ray Diffraction Technique
The XRD patterns were recorded using a computer-controlled x-ray powder
diffractometer, JEOL JDX 8030, with CuKα (Ni-filtered) radiation (λ = 1.5418 A).
Luminescence Spectroscopy
The luminescence spectra were recorded by Hitachi 650-10 S fluorescence
spectrophotometer equipped with a 150 W xenon lamp and a Hamamatsu R 928 F
photomultiplier tube.
RESULTS AND DISCUSSION
Analysis of the Results of the Weight-Loss Method
The corrosion rates of mild steel in a neutral aqueous environment containing 60 ppm Clin the absence and presence of inhibitors at various concentrations, obtained by the weight
loss method are given in Table 1. The corrosion inhibition efficiencies of various systems are
also given in Table 1.
484
Rajendran et al.
Table 1. Corrosion Rates of Mild Steel in a Neutral Aqueous Environment (Cl- = 60 ppm) in
the Absence and Presence of Inhibitors, and the Inhibition Efficiencies Obtained by
the Weight-Loss Method
2+
SI. EPA
Zn
MoO4
No. (ppm)
(ppm)
(ppm)
1
2
300
3
50
4
300
50
5
300
50
50
6
300
50
100
7
300
50
200
8
300
50
300
9
300
300
10
50
300
11
300
Inhibitor system: EPA + Zn2+ + MoO42-
2-
Corrosion
rate
(mdd)
15.54
15.39
19.11
6.22
13.98
11.66
9.32
0.16
0.16
1.58
3.11
Inhibition
Efficiency
(%)
1
-23
60
10
25
40
99
99
90
80
It is evident from Table 1 that ethyl phosphonic acid (EPA) by itself is not a good
inhibitor and Zn2+ is corrosive. Interestingly the formulation consisting of 300 ppm EPA and
50 ppm Zn2+ had a 60% inhibition efficiency. This indicates the synergistic effect between
EPA and Zn2+. When various concentrations of molybdate were added to the above system,
the inhibition efficiency increased at 300 ppm MoO42-. The formulation consisting of 300
ppm EPA, 50 ppm Zn2+ and 300 ppm MoO42- had a 99% efficiency. It was found that the
formulations consisting of 300 ppm EPA and 300 ppm MoO42-, and also MoO42- (300ppm)Zn2+ (50 ppm) showed a synergistic effect.
Analysis of the Potentiostatic Polarization Curves
The potentiostatic polarization curves of mild steel immersed in various environments
are given in Fig. 1. It is observed that, when molybdate was added to chloride or EPA or Zn2+
or EPA-Zn2+, the corrosion potential shifted to the anodic side. The formulation consisting of
300 ppm EPA, 300 ppm MoO42- and 50 ppm Zn2+ shifted the corrosion potential to -390 mV
vs. SCE. This indicates that this formulation acts as a mixed inhibitor. This is further
supported by the fact that the anodic and cathodic Tafel slopes shifted almost equally (28
mV/decade).
Analysis of the FTIR Spectra
The FTIR spectrum of pure EPA (KBr) is given in Fig. 2a. The FTIR spectrum (by the
multiple internal reflection (MIR) technique) of the film formed on the surface of the metal
specimen immersed in an environment consisting of 60 ppm Cl-, 300 ppm MoO42- and 300
ppm EPA is given in Fig. 2b. It is found that the P-O stretching frequency [15-17] of the
phosphonic acid decreased from 1071.7 cm-1 to 1018.6 cm-1. This suggests that the oxygen
atom of the phosphonic acid coordinated to Fe2+ on the metal surface [18,19].
485
Corrosion Protection and Monitoring
The FTIR spectrum (MIR) of the film due to the environment consisting of 60 ppm Cl-,
300 ppm EPA, 300 ppm MoO42- and 50 ppm Zn2+ reveals that the P-O stretching frequency
decreased from 1071.7 cm-1 to 1018.4 cm-1. This suggests that in this case also, the oxygen
atom of the phosphonic acid coordinated to Fe2+ on the metal surface. Furthermore, the peak
at 1456 cm-1 was due to ZnO2 [20]. This may be explained by the fact that Zn(OH)2 formed
on the cathodic sites [19] converts into ZnO2.
Analysis of the UV-Visible Reflectance Spectra
The UV-visible reflectance spectra of the films formed on the surface of metal specimens
immersed in various test solutions are given in Fig. 3. The spectrum of the film due to the
environment containing 60 ppm Cl- and 300 ppm MoO42- shows a peak at 320 nm (Fig. 3a).
This may be due to a complex formed between the iron and molybdate.
The spectrum of the film formed on the surface of the metal immersed in the
environment, consisting of 60 ppm Cl-, 300 ppm MoO42- and 50 ppm Zn2+ is given in Fig. 3b.
The wavelength transition at 550 nm indicates the presence of oxides of iron (band gap =
1.239/0.55 = 2.25 eV) on the metal surface [21] having semiconducting property [22]. The
peak at 320 nm may be due to an iron-molybdate complex.
Figure 1. Potentiostatic polarization curves
(a) Cl- 60 ppm
(e) Cl- 60 ppm + EPA 300 ppm
(b) Cl- 60 ppm + Zn2+ 50 ppm
(f) Cl- 60 ppm + EPA 300 ppm + Zn2+ 50 ppm
(c) Cl- 60 ppm + MoO42- 300 ppm
(g) Cl- 60 ppm + EPA 300 ppm + MoO42300 ppm
(d) Cl- 60 ppm + MoO42- 300 ppm
+ Zn2+ 50 ppm
(h) Cl- 60 ppm + EPA 300 ppm + MoO42300 ppm + Zn2+ 50 ppm
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Rajendran et al.
Figure 2. FTIR spectra
Figure 3. UV-visible reflectance spectra
(a) Pure EPA
(a) Cl- 60 ppm + MoO42- 300 ppm
(b) Cl- 60 ppm + EPA 300 ppm
+ MoO42- 300 ppm
(b) Cl- 60 ppm + MoO42- 300 ppm + Zn2+ 50
ppm
(c) Cl- 60 ppm + EPA 300 ppm
+ MoO42- 300 ppm + Zn2+ 50 ppm
(c) Cl- 60 ppm + EPA 300 ppm + MoO42- 300
ppm
(d) Cl- 60 ppm + EPA 300 ppm + MoO42300 ppm + Zn2+ 50 ppm
The spectrum of the film due to the environment consisting of 60 ppm Cl-, 300 ppm
MoO42- and 300 ppm EPA (Fig. 3c) does not show any wavelength transition at 550 nm
indicating the absence of any oxides of iron on the metal surface. The peak at 320 nm is due
to an iron-molybdate complex formed on the metal surface.
The reflectance spectrum of the film due to the environment containing 60 ppm Cl-, 300
ppm MoO42-,300 ppm EPA and 50 ppm Zn2+ (Fig. 3d) has a peak at 320 nm due to an ironmolybdate complex. Absence of a wavelength transition at 550 nm indicates the absence of
oxides of iron on the metal surface.
Analysis of the X-Ray Diffraction Patterns
The XRD patterns of the film formed on the surface of the metal specimens immersed in
various test solutions are given in Fig. 4. The XRD pattern of the film due to the environment
consisting of 60 ppm Cl- and 300 ppm MoO42- is given in Fig. 4a. The film consisted of
Fe2(MoO4)3 (2θ = 14.1°, 22.6°, 30.63° and 31.88°) [23]. The peaks due to iron appear at 2θ =
44.5°, 64.8° and 82.2°.
The film formed on the surface of the metal specimen immersed in the environment
containing 60 ppm Cl-, 300 ppm MoO42- and 50 ppm Zn2+ contained Fe2(MoO4)3 (2θ = 22.0°,
26.7°, 45.5°, 66.1°) [23], ZnO2 (2θ = 45.5° and 66.1°) [24] and γ-FeOOH (2θ = 60.9°) [25].
The iron peaks appear at 2θ = 44.6°, 65.0° and 82.3°.
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Corrosion Protection and Monitoring
Figure 4. XRD patterns
Figure 5. Luminescence spectra
(a) Cl- 60 ppm + MoO42- 300 ppm
(a) Cl- 60 ppm + EPA 300 ppm + MoO42- 300 ppm
(b) Cl- 60 ppm + MoO42- 300 ppm
+ Zn2+ 50 ppm
(b) Cl- 60 ppm + EPA 300 ppm + MoO42- 300
ppm + Zn2+ 50 ppm
(c) Cl- 60 ppm + EPA 300 ppm
+ MoO42- 300 ppm
(d) Cl- 60 ppm + EPA 300 ppm
+ MoO42- 300 ppm + Zn2+ 50 ppm
The film due to the environment containing 60 ppm Cl-, 300 ppm MoO42- and 300 ppm
EPA consisted of Fe2MoO4 (2 θ = 34.9°) [26]. The peaks due to iron appear at 2θ = 44.5°,
64.9° and 82.2°.
The film formed on the surface of the metal immersed in the environment containing 60
ppm Cl-, 300 ppm MoO42-, 300ppm EPA and 50 ppm Zn2+ consisted of Fe2(MoO4)3 (2θ =
22.7°) [23], ZnMoO4 (2θ = 30.5°) [27] and ZnO2 (2θ = 41.1°) [24]. The iron peaks appear at
44.6°, 65.0° and 82.4°.
Analysis of the Luminescence Spectra
The emission spectrum (λex = 300 nm) of the film formed on the surface of the metal
immersed in the environment containing 60 ppm Cl-, 300 ppm EPA and 300 ppm MoO42- is
given in Fig. 5a. This spectrum may be due to an Fe2+-EPA complex and Fe2MoO4.
The emission spectrum (λex = 300 nm) of the film due to the environment consisting of
60 ppm Cl- , 300 ppm EPA, 300 ppm MoO42- and 50 ppm Zn2+ is given in Fig. 5b. This
spectrum may be due to an Fe2+-EPA complex and Fe2(MoO4)3 in the presence of ZnMoO4
and ZnO2.
Mechanism of the Inhibition of Corrosion
The results of the weight loss method show that the formulation consisting of 300 ppm
EPA, 300 ppm MoO42- and 50 ppm Zn2+ had an inhibition efficiency of 99%. The
polarization study revealed that this system acts as a mixed inhibitor. The FTIR spectra
indicate that the protective film consisted of an Fe2+-EPA complex and ZnO2. The UVvisible reflectance spectra show that the film did not contain any oxides of iron. The XRD
patterns show that the protective film consisted of Fe2(MoO4)3, ZnMoO4 and ZnO2. The film
was found to be luminescent. In order to explain these observations in a holistic way, the
following mechanism of inhibition of corrosion is proposed.
1. When the environment consisting of 60 ppm Cl-, 300 ppm EPA, 300 ppm MoO42- and 50
ppm Zn2+ is prepared, there is formation of an Zn2+-EPA complex and a Zn2+-MoO42complex in solution.
2. When the metal is immersed in this environment, the Zn2+-EPA complex and the Zn2+MoO42- complex diffuse from the bulk of the solution to the surface of the metal.
488
Rajendran et al.
3. On the surface of the metal, the Zn2+-EPA complex is converted into an Fe2+-EPA
complex in the local anodic sites, since the latter is more stable than the former.
Zn2+-EPA + Fe2+ ---> Fe2+ -EPA + Zn2+
(1)
4. Similarly, the Zn2+-MoO42- complex is converted into an iron-molybdate complex,
namely, Fe2(MoO4)3
Zn2+-MoO42- + 2 Fe3+ ---> Fe2(MoO4)3 + 3 Zn 2+
(2)
(Formation of an Fe3+-EPA complex and an Fe2+-MoO42-complex to some extent cannot be
ruled out)
5. The released Zn2+ on the metal surface forms Zn(OH)2 in the local cathodic regions.
Zn2+ + 2 OH- ---> Zn(OH)2
(3)
This may be converted into ZnO2
6. ZnMoO4 also forms on the metal surface.
CONCLUSIONS
1. A synergistic effect was noticed between MoO42- and Zn2+; MoO42- and EPA; and MoO42, EPA and Zn2+.
2. Molybdate shifted the corrosion potential of the Zn2+, EPA or EPA-Zn2+ system to the
anodic side.
3. The formulation consisting of 300 ppm EPA and 300 ppm MoO42- had a 99% inhibition
efficiency. The protective film consisted of an Fe2+-EPA complex and Fe2MoO4. This
film was found to be luminescent.
4. The formulation consisting of 300 ppm EPA, 300 ppm MoO42- and 50 ppm Zn2+ had 99%
inhibition efficiency.
The protective film consisted of an Fe2+-EPA complex,
Fe2(MoO4)3, ZnMoO4 and ZnO2. This film was found to be luminescent.
ACKNOWLEDGEMENT
S. Rajendran wishes to thank the University Grants Commission, India, for awarding him
a fellowship; and Mr. Ranjit Soundararajan, the Correspondent, Prof. S. Ramakrishnan, the
Principal, and Prof. P. Jayaram, HOD, Chemistry Department, GTN Arts College, Dindigul,
for their encouragement.
REFERENCES
1. M.S. Vukasovich and J.P.G. Farr, Materials Performance, May 1986, p. 9.
2. A.Y. Al-Borno, R.A. Haleem, A. Al-Shatti, A. Abdulla and T.H. Mustafa, Technical
Report No. 2132, Kuwait Institute for Scientific Research, Kuwait, 1986.
3. M.S. Vukasovich and D.R. Robitaille, J. Less-Common Metals 54, 1977, p. 437.
489
Corrosion Protection and Monitoring
4. C. O'Neal, Jr., R.N. Borger, Materials Performance 15, 1976, p. 9.
5. J.I. Bregman, US Patent 3,024,201, 1962.
6. T.C. Breske, Materials Performance 16, 1977, p. 17.
7. H. Leidheiser , Jr., Corrosion 36, 1980, p. 339.
8. J.P.G. Farr and M. Saremi, Surface Technology 17, 1982, p. 19.
9. D.B. Alexander and A.A. Moccari, Corrosion 49, 1993, p. 921.
10. M.R. Reda and J.N. Alhajji, Journal of the University of Kuwait (Science) 20, 1993, p.
171.
11. A. Vonkoepper, G.A. Emerle, K. Nishio and B.A. Metz, Materials Protection
Performance 12, 1973, p. 23.
12. Y.J. Qian and S. Turgoose, British Corrosion Journal 22, 1987, p. 268.
13. A. Hussain, K. Habib and R. Jarman, Proceedings 7th European Symposium on Corrosion
Inhibitors, Ferrara, Italy, 1, 1990, p. 621.
14. A. Al-Borno, Proceedings 7th European Symposium on Corrosion Inhibitors, Ferrara,
Italy, 1, 1990, p. 583.
15. R.M. Silverstein, G.C. Bassler and T.C. Morrill, Spectrometric Identification of Organic
Compounds, New York, John Wiley and Sons, 1981.
16. K. Nakamoto, Infrared and Raman Spectra of Inorganic and Coordination Compounds,
New York, Wiley-Interscience, 1986.
17. A.D. Cross, Introduction to Practical Infrared Spectroscopy, London, Butterworths
Scientific Publication, 1960.
18. L. Horner and C.L. Horner, Werkstoff und Korrosion 27, 1976, p. 223.
19. S. Rajendran, B.V. Apparao and N. Palaniswamy, Proceedings 8th European Symposium
on Corrosion Inhibitors, Ferrara, Italy, 1, 1995, p. 465.
20. R.A.Nyquist and R.O. Kadel, Infrared Spectra of Inorganic Compounds, New York,
Acadamic Press, 1971.
21. C. Sanchez, K.D. Sieber and G.A. Somorjai, Journal Electroanalytical Chemistry 252,
1988, p. 269.
22. S.M. Wilhelm and N. Hackerman, Journal Electrochemical Society 128, 1981, p. 1668.
23. JCPDS Nr. 200526.
24. JCPDS Nr. 130 311
25. M. Favre and D. Landolt, Corrosion Science 34, 1993, 1481.
26. JCPDS Nr. 251403.
27. JCPDS Nr. 251024.
490
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
EVALUATION OF CORROSION INHIBITORS FOR CARBON STEEL, MONEL 400
AND STAINLESS STEEL 321 IN A MONOETHANOLAMINE ENVIRONMENT
UNDER STAGNANT AND HYDRODYNAMIC CONDITIONS
J. Carew, H. Al-Sumait, A. Abdullah and A. Al-Hashem
Materials Application Department
Kuwait Institute for Scientific Research
P.O.Box 24885, Safat, 13109, Kuwait
ABSTRACT
Four organic based corrosion inhibitors have been evaluated for carbon steel, Monel 400 (UNS
No 400) and stainless steel 321 (UNS No. 32100) in fresh monoethanolamine (MEA) environments
saturated with an H2/CO2 gas mixture at 40oC. The test solutions were prepared from fresh MEA
solutions with and without the addition of 250 ppm of each inhibitor. Initial screening of the
inhibitors was performed using the wheel test to determine the corrosion rates of the three alloys with
and without inhibitors. The rotating disc electrode (RDE) method was used to determine the
effectiveness of the four organic inhibitors under hydrodynamic conditions. It was found that flow
conditions tended to increase the effectiveness of some corrosion inhibitors with respect to stagnant
conditions. The weight-loss and electrochemical tests conducted under hydrodynamic conditions
indicated that the quaternary ammonium-based inhibitor was the most effective of the three alloys in
the different MEA solutions.
Key Words: Corrosion inhibitors, carbon steel, monel 400, stainless steel 321, monoethanolamine,
weight-loss, rotating disc electrode
INTRODUCTION
A large variety of corrosive conditions is encountered in the different industries. The
costs of corrosion, and correspondingly, the savings gained through the use of appropriate
corrosion mitigation techniques is considerable. Corrosion inhibitors are one of the main
methods used to reduce corrosion problems in metallic installations all over the world.
One of the most expensive and corrodible installations in chemical plants and refineries
is the gas purification system. The most economical and effective method of protection is the
addition of inhibitors to the closed circulation circuit of the system. Because carbon steel,
UNS No 400 and stainless steel 321 (SS321) are the major alloys of construction in
monoethanolamine (MEA) gas treating systems, both inorganic- and organic-based inhibitors
have been utilized to reduce the corrosion rate in such an environment. However, due to
environmental regulations and toxicity considerations, the use of inorganic inhibitors is
declining and that of organic-based inhibitors is rising. The ability to evaluate and screen
corrosion inhibitors for this type of application is important in order to choose an appropriate
inhibitor.
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Corrosion Protection and Monitoring
It is well known that flow velocity exerts a great influence on the corrosion rate of
metallic materials [1,2]. Despite this knowledge the importance of the flow rate is not
sufficiently considered in evaluation tests of corrosion inhibitors. The use of circulating
loops to enable the examination of a wide range of flow rates (laminar and turbulent) is
sometimes considered too expensive and time consuming because of the long sequence of
operations needed [3]. Several authors [2,4,5] suggest simulating the turbulent flow present
in many systems by using rotating disc electrodes (RDEs). In this technique, cylindrical
electrodes are rotated at different speeds to measure the effect of flow at various velocities.
This work was undertaken to screen four corrosion inhibitors in fresh monoethanolamine
saturated with 85% CO2/15% H2 using the weight-loss and RDE methods. The aim of this
investigation was to determine which corrosion inhibitor produces the lowest corrosion rate
for carbon steel, UNS No 400 and SS321 in MEA solution by correlating the data obtained by
the weight-loss and RDE methods.
EXPERIMENTAL PROCEDURE
Materials
The materials tested were carbon steel (ASTM A283 grade B) Monel 400 and SS321.
All the alloys were supplied by the Kuwait National Petroleum Company (KNPC) and
represent the materials of construction of the plant.
Inhibitors Tested
Four corrosion inhibitors were tested under stagnant conditions by the weight-loss test,
and under flow conditions by the RDE technique. The inhibitors were labeled A, B, C and D
for simplicity and are shown in Table 1.
Table 1. Type of Commercial Inhibitors Tested
Inhibitor Type
A
B
C
D
Chemical Family
Amides
Aklylamino acid and ethylene glycol
Amines
Quaternary ammonium compound
Weight-Loss Method
The procedure for the weight-loss method is essentially like that given in ASTM
standard G-31 (1990). The specimens were in the form of coupons in dimensions of 40 x 20
x 1 mm. Before exposure to the test environment, the surfaces of the coupons were
successively ground with 180, 400 and 600 grit silicon carbide papers, washed with detergent
and dried. The weight of each specimen was determined accurately on an electronic balance.
Prior to immersion in the test medium, the coupons were decreased in acetone and dried by
hot air and hung with nylon thread. The test vessels were 800 ml Pyrex glass fitted with a gas
inlet and immersed in thermostatically controlled water baths. The volume of the test
solution was 400 ml. Each cell contained 2 coupons of carbon steel. During the course of the
494
Carew et al.
test, coupons were removed after 2 and 4 weeks. After weight-loss determination, the
corrosion rate was calculated from the weight loss data as (mpy) from the following formula
according to ASTM G-31 (1990).
Corrosion rate = (K x W)/(A x T x D)
(1)
where
K = constant = 3.45 X 106, W = mass loss in g, A = area in cm2, T = time of exposure
in hours, and D = density in g/cm3.
Rotating Disc Electrode (RDE)
For the RDE measurements, the system Model 616 RDE, by EG&G PARC, was used,
mounting and rotating the cylindrical electrodes at 200, 1000, and 3000 rpm. Some tests
were conducted under stagnant conditions for comparison. The corrosion rates for the steel
electrodes were determined at each speed selected, by using a potentiostat/galvanostat
M000odel 273 A by EG&G PARC through the linear polarization resistance (LPR)
measurement method.
The electrochemical cell used consisted of the working electrode, a graphite counter and
a saturated calomel electrode (SCE) as a reference electrode. The test was conducted with 400
ml of fresh MEA solution. Tests were conducted in 4 uninhibited MEA solutions as well as
solutions inhibited ones with addition of 250 ppm of each inhibitor. The solution was
continuously purged with a gas mixture of 85% CO2 and 15% H2 ,and the temperature was
maintained at 40°C.
RESULTS
Weight-Loss Method
Carbon Steel. Figure 1 shows the corrosion rate of this alloy in fresh (18% H2O) MEA
solution in the presence of four organic-based inhibitors at a temperature of 40°C for a period
of 4 weeks. To determine the most effective inhibitor in the MEA solution for this alloy, it
was decided to rank the inhibitors by comparing the corrosion rates of the alloy in the
absence and presence of inhibitors. In other words, the corrosion inhibitor that reduced the
corrosion rate of carbon steel to the lowest value would be ranked as the best, and the one
with the highest corrosion rate would be ranked as the worst. The corrosion rate of carbon
steel in fresh MEA and in the absence of any corrosion inhibitor was considered to be the
reference point (blank conditions). Therefore, the ranking of the inhibitors in terms of their
corrosion performance for carbon steel in fresh MEA solution was as follows (Fig. 1):
D>C>A>B
Monel 400. Figure 2 illustrates the corrosion rates of this alloy in fresh MEA solution in
the absence and pressure of the four inhibitors at 40°C and after four weeks of immersion.
The ranking of these inhibitors in terms of their corrosion protection to Monel 400 in the
fresh MEA-solutions is as follows:
A>D>B>C
495
Corrosion Protection and Monitoring
Figure 1. Corrosion rate of carbon steel in fresh MEA solution as a function of
inhibitor type at a temperature of 40°C
Figure 2. Corrosion rate of Monel 400 in fresh MEA solution as a function of
inhibitor type at a temperature of 40°C
Stainless steel 321. Figure 3 shows the corrosion rate of this alloy with and without
inhibitors at 40°C for 4 weeks. The ranking of those inhibitors for this alloys was as follows:
D > C> A>B
496
Carew et al.
Figure 3. Corrosion rate of stainless steel in fresh MEA solution as a function of
inhibitor type at a temperature of 40°C
Rotating Disc Electrode Technique
Carbon Steel. Figure 4 shows the corrosion rate of carbon steel in fresh MEA solution
with and without corrosion inhibitors under hydrodynamic conditions at 40°C. At a rotational
speed of 200 rpm the inhibitors were ranked as follows:
D > A> B> C
However, at high speeds, the inhibitors were not as effective as under blank conditions.
Monel 400. Figure 5 shows the corrosion rate of Monel 400 in fresh MEA with and
without corrosion inhibitors under three different hydrodynamic velocities. The corrosion
protection performance varied from one inhibitor to another, as well as from one speed to
another. None of the inhibitors were effective for Monel 400 under hydrodynamic
conditions.
Stainless Steel 321. Figure 6 shows the corrosion rate of this alloy in fresh MEA with
and without corrosion inhibitors under hydrodynamic conditions at 40°C. The inhibitors
were ranked in terms of their corrosion protection as follows:
C>B>A
DISCUSSION
This investigation was carried out to evaluate the relative performance of 4 organicbased inhibitors for the CO2 removal system of one of the refineries in Kuwait using MEA
solution. The 3 main alloys that comprise such a system are carbon steel, Monel 400 and
SS321. The two methods used in the evaluation process were the weight-loss and RDE
methods representing stagnant and flow conditions, respectively. The 4 inhibitors were
studied to asses their effect on the general corrosion of carbon steel, Monel 400 and SS321 in
CO2 saturated fresh MEA solutions under stagnant and hydrodynamic conditions.
Under stagnant conditions, the corrosion rates of carbon steel in fresh MEA (Fig. 1) was
slightly more than 0.2 mpy for blank conditions. The addition of inhibitors A, C and D
tended to lower the corrosion rate of carbon steel to an acceptable level. However, inhibitor
B was found to enhance the corrosion rate of carbon steel to an acceptable level. This
behavior might be attributed to the nonuniform distribution of the inhibitor film on the
surface of the carbon steel specimen.
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Corrosion Protection and Monitoring
Figure 4. Corrosion rate of carbon steel in fresh MEA solution as a function of
inhibitor type and rotational speeds at a temperature of 40°C
Figure 5. Corrosion rate of Monel 400 in fresh MEA solution as a function of
inhibitor type and rotational speeds at a temperature of 40°C
Figure 6. Corrosion rate of stainless steel in fresh MEA solution as a function of
inhibition type and rotational speeds at a temperature of 40°C
The corrosion rate of Monel 400, as shown in Fig. 2, was surprisingly high for such type
of nickel-based alloy under stagnant blank conditions. The addition of any of the four
inhibitors at the recommended dosage level reduced the corrosion rate quite dramatically
especially for inhibitors A and D.
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Carew et al.
Figure 3 shows the corrosion rate of SS321 under stagnant conditions which was quite
low in the uninhibited media. The addition of inhibitors C and D reduced the corrosion rate
of SS321 to very low levels. However, the addition of inhibitors A and B increased the
corrosion rate of this alloy. This behavior was observed for SS304 and SS316 in identical
media in previous studies [6,7].
Under hydrodynamic conditions, the corrosion rate of carbon steel, as shown in Fig. 4,
indicates that under blank conditions, the flow velocity tended to decrease the corrosion rate
of this alloy, in comparison to stagnant condition. This observation has been reported by
many authors [1,2,4,5]. However, the dissolution rate of the steel cylinders in the inhibited
MEA solutions were some what independent of the rotation velocity. This behavior may be
interpreted by assuming the formation of a thick surface layer on the carbon steel electode.
This layer strongly hindered either the anodic or the cathodic reaction or both on the surface
of the steel electrode. Such a phenomenon was also reported by Zucchi et al.[2].
The corrosion rate of Monel 400 under flow conditions (Fig. 5) was lower under
uninhibited conditions for the three different velocities than for samples with inhibitors A, B,
C, and D. According to Fig. 5, Monel 400 did not seem to be affected by the different
rotational speeds in the uninhibited MEA solutions. In other words, the passive oxide layer
on the surface of Monel 400 was sufficient to resist destruction at up to 3000 rpm. The
increase in the corrosion rate of this alloy upon the addition of the 4 inhibitors could be
attributed to the removal or nonuniform formation of inhibitor film under hydrodynamic
conditions.
Figure 6 shows the corrosion rate of SS321 under hydrodynamic conditions in the
uninhibited and inhibited MEA solutions. The addition of inhibitors A, B and C tended to
lower the corrosion rate of SS321 under flow conditions. Inhibitor C seemed to be the most
effective one in reducing the corrosion rate of SS321.
Based on the results obtained by the weight-loss and RDE methods, the inhibitor that
seemed to be the most effective in reducing the overall general corrosion rate for the three
alloys was inhibitor D (quaternary ammonium compound). Generally, the inhibition
mechanism of this compound is due to its ability to adsorb into the metal or alloy surface to
form a protective film. Organic inhibitors can adsorb on to a metal surface by hydrogen
bonding, by electron donation from the nitrogen atom, or by interaction of the dipole with the
surface charge [8, 9, 10].
CONCLUSIONS
The most effective corrosion inhibitor that may be used for carbon steel, Monel 400 and
SS321 as determined under laboratory conditions, is the quaternary ammonium compound.
The corrosion rate of the three alloys in the inhibited fresh MEA solutions varied with respect
to flow conditions. The corrosion rate of carbon steel under such conditions was independent
of the rotational speed, slightly increased for Monel 400 and decreased with respect to SS321.
ACKNOWLEDGMENT
The authors would like to acknowledge the in-kind support of KNPC’s - Shuaiba
Refinery of this research work.
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Corrosion Protection and Monitoring
REFERENCES
1. E. Heitz, Corrosion ‘90, Paper No. 1, NACE, Houston, Texas, USA., 1990.
2. F. Zucchi, G. Trabanelli, and G. Brunoro, Effect of flow velocity on corrosion inhibition
of steel in HCl, in Progress in the Understanding and Prevention of Corrosion, Vol.2, J.
M. Costa and A. D. Mercer, eds.,The Institute of Materials, p. 845, 1993.
3. J.L. Dawson, C.C. Shih, R.G. Miller and J.W. Palmer, Corrosion 90, Paper No. 14,
NACE, Houston, Texas, USA., 1990.
4. A. Mazanek and H. Bala, Corrosion Science 28, 5, 1988, p. 513.
5. D.C. Silverman, Corrosion prediction from circuit models application to evaluation of
corrosion inhibitors, in electrochemical impedance - Analysis and Interpretation, J.
Scully, D. Silverman, and M. Kendig Editors, ASTM Publication, Philadelphia,
Pensylvania,1993, p. 192.
6. M. Islam, A. Abdullah, W. Riad and G. Mansi, An investigation of corrosion and its
control in the PIC monoethanolamine carbon dioxide removal unit”. Kuwait Institute for
Scientific Research, Report No. KISR 1446, Kuwait 1984.
7. M. Islam, A. Abdullah, W. Riad , R. Al-Taib and G. Mansi, An investigation of corrosion
and its control in the PIC MEA carbon dioxide removal unit laboratory studies”, Kuwait
Institute for Scientific Research, Report No. KISR 1603, Kuwait 1984.
8. A. Moccari and D. D. Macdonald, Corrosion 41, 5, 1985, p. 263,
9. J.S. Robinson, Corrosion Inhibitors, 1979.
10. C.C. Nathan, Corrosion Inhibitors, NACE publications, Houston Texas, USA, 1973.
500
Industrial Corrosion and Corrosion Control Technology
Shalaby, H.M. et al. (Editors)
1996 Kuwait Institute for Scientific Research. Printed in Kuwait
LABORATORY EVALUATION OF THE EFFECTS OF OZONE ON
CORROSION RATES AND PITTING OF ENGINEERING ALLOYS
S. Nasrazadani
Department of Materials Engineering
Isfahan University of Technology, Isfahan, 84156, Iran
ABSTRACT
Cyclic polarization experiments were performed on 1018 steel, yellow brass, and 6061-T6 aluminum
in ozonated and non-ozonated tap water under stagnant conditions to evaluate pitting and corrosion
tendencies of these engineering metals. Results show that the corrosion rate of 1018 carbon steel in tap
water under stagnant conditions increased about three-fold with the injection of ozone (0.05-0.1 ppm). No
considerable changes in the corrosion rate occurred for yellow brass and aluminum when ozone was added
under similar conditions. Study of the pitting behavior of the materials also demonstrated that ozone did
not increase pitting when injected into the test solution for yellow brass. But in the case of 1018 carbon
steel and aluminum, pitting was more pronounced.
Key Words:
Ozone, pitting, corrosion rates, electrochemical testing, engineering alloys, laboratory
evaluation
INTRODUCTION
It is now a well known fact that the injection of ozone into the cooling water circulated in
cooling towers can help to prevent biofouling due to