Copyright © 2008 ASM International®. All rights reserved.
Stainless Steels for Design Engineers (#05231G)
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STAINLESS STEELS
FOR DESIGN ENGINEERS
MICHAEL MCGUIRE
ASM International®
Materials Park, Ohio 44073-0002
www.asminternational.org
Copyright © 2008
by
ASM International®
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First printing, December 2008
Great care is taken in the compilation and production of this book, but it should be made clear that NO WARRANTIES, EXPRESS OR IMPLIED, INCLUDING, WITHOUT LIMITATION, WARRANTIES OF MERCHANTABILITY OR FITNESS FOR A PARTICULAR PURPOSE, ARE GIVEN IN CONNECTION WITH
THIS PUBLICATION. Although this information is believed to be accurate by ASM, ASM cannot guarantee
that favorable results will be obtained from the use of this publication alone. This publication is intended for use
by persons having technical skill, at their sole discretion and risk. Since the conditions of product or material use
are outside of ASM’s control, ASM assumes no liability or obligation in connection with any use of this information. No claim of any kind, whether as to products or information in this publication, and whether or not
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FROM THE NEGLIGENCE OF SUCH PARTY. As with any material, evaluation of the material under end-use
conditions prior to specification is essential. Therefore, specific testing under actual conditions is recommended.
Nothing contained in this book shall be construed as a grant of any right of manufacture, sale, use, or reproduction, in connection with any method, process, apparatus, product, composition, or system, whether or not covered by letters patent, copyright, or trademark, and nothing contained in this book shall be construed as a defense against any alleged infringement of letters patent, copyright, or trademark, or as a defense against liability
for such infringement.
Comments, criticisms, and suggestions are invited, and should be forwarded to ASM International.
Prepared under the direction of the ASM International Technical Book Committee (2007–2008), Lichun L.
Chen, Chair.
ASM International staff who worked on this project include Scott Henry, Senior Manager of Product and Service
Development; Steven R. Lampman, Technical Editor; Eileen De Guire, Associate Editor; Ann Britton, Editorial
Assistant; Bonnie Sanders, Manager of Production; Madrid Tramble, Senior Production Coordinator; Diane
Grubbs, Production Coordinator; Patty Conti, Production Coordinator; and Kathryn Muldoon, Production Assistant
Library of Congress Control Number: 2008934669
ISBN-13: 978-0-87170-717-8
ISBN-10: 0-87170-717-9
SAN: 204-7586
ASM International®
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Printed in the United States of America
Contents
Preface . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . v
METALLURGY
Chapter 1
Metallurgy........................................................................................................1
CORROSION AND OXIDATION
Chapter 2
Corrosion Theory............................................................................................11
Chapter 3
Corrosion Kinetics..........................................................................................19
Chapter 4
Corrosion Types..............................................................................................27
Chapter 5
Oxidation.......................................................................................................57
STAINLESS STEEL ALLOYS
Chapter 6
Austenitic Stainless Steels ..............................................................................69
Chapter 7
Duplex Stainless Steels...................................................................................91
Chapter 8
Ferritic Stainless Steels .................................................................................109
Chapter 9
Martensitic Stainless Steels ..........................................................................123
Chapter 10
Precipitation-Hardening Stainless Steels ......................................................137
PROCESSING
Chapter 11
Casting Alloys...............................................................................................147
Chapter 12
Melting, Casting, and Hot Processing...........................................................155
Chapter 13
Thermal Processing ......................................................................................161
Chapter 14
Forming........................................................................................................173
Chapter 15
Machining ....................................................................................................181
Chapter 16
Surface Finishing ..........................................................................................193
Chapter 17
Welding........................................................................................................201
Chapter 18
Architecture and Construction.....................................................................213
Chapter 19
Automotive and Transportation Applications................................................225
Chapter 20
Commercial and Residential Applications ....................................................233
Chapter 21
Marine Systems Applications........................................................................243
Chapter 22
Petroleum Industry Applications ..................................................................247
Chapter 23
Chemical and Process Industry Applications ................................................257
Chapter 24
Pulp-and-Paper Industry Applications ..........................................................265
APPLICATIONS
APPENDIXES
Appendix 1
Compositions...............................................................................................269
Appendix 2
Physical and Mechanical Properties of Select Alloys....................................279
Appendix 3
Introduction to Thermo-Calc and Instructions for
Accessing Free Demonstration ....................................................................281
Index .................................................................................................................................285
iv
Preface
The rate of growth of stainless steel has outpaced that of other metals and alloys, and by 2010 may
surpass aluminum as the second most widely used metal after carbon steel. The 2007 world production of stainless steel was approximately 30,000,000 tons and has nearly doubled in the last ten years.
This growth is occurring at the same time that the production of stainless steel continues to become
more consolidated. One result of this is a more widespread need to understand stainless steel with
fewer resources to provide that information. The concurrent technical evolution in stainless steel and
increasing volatility of raw material prices has made it more important for the engineers and designers who use stainless steel to make sound technical judgments about which stainless steels to use and
how to use them.
This book provides design engineers with an up-to-date source of information at a level useful for
both metallurgists and other engineers and technicians. It seeks to bridge the gap between the internet where much current, but raw information is available and scholarly books and journals that provide theory that is difficult to put into practice. The content of the book is selected for utility for the
user of stainless steel. The first section gives elementary metallurgy and identification of constituents
of stainless, the effects of alloying elements and a significant section on corrosion. A second section
is oriented toward processes important to users of stainless steel. The third section is about each family of stainless alloys and includes the most recent additions that have come to the market. The fourth
section deals in some depth with the major applications for stainless steel. This last part is presented
without the promotional bias which is found in many steel producers’, alloy producers’, and trade associations’ literature. While a number of steel producers have provided assistance to the author, there
has been no attempt to unfairly bias information in their favor. To the contrary, those producers responsible for generating factual, useful data for the user community are those who should benefit the
most by books such as this. The author is particularly indebted to Allegheny Ludlum and John
Grubb, and his many colleagues who assisted him, for technical assistance throughout the writing
and to Carnegie Mellon University for their support. The author also wishes to thank Professor Sridhar Seetharaman at Carnegie Mellon University for his help in writing the corrosion chapter and others who helped: Roy Matway of CMU, Vittorio Boneschi of Centro-Inox; Paul Mason of ThermoCalc; Bob Drab of Schmolz Bichenbach; Elisabeth Torsner and Chuck Turack Outukumpu, USA;
Scott Balliett of Latrobe Steel; Jim Halliday and Fred Deuschle of Contrarian Metals Resources; Professors Tony DeArdo of Pitt and Gerhard Welsch of CWRU; the staffs of Centro-Inox, Euro-Inox,
SSNA, The Nickel Institute; and the editorial staff at ASM International, Scott Henry, Eileen
DeGuire, Charlie Moosbrugger and Steve Lampman. I would also like to thank the many members
of my forum at Eng-tips.com who have contributed much collective knowledge and perspective to
this book.
ASM International is the society for materials
engineers and scientists, a worldwide network
dedicated to advancing industry, technology, and
applications of metals and materials.
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information is believed to be accurate by ASM, ASM cannot guarantee that favorable results will be
obtained from the use of this publication alone. This publication is intended for use by persons having
technical skill, at their sole discretion and risk. Since the conditions of product or material use are
outside of ASM's control, ASM assumes no liability or obligation in connection with any use of this
information. As with any material, evaluation of the material under end-use conditions prior to
specification is essential. Therefore, specific testing under actual conditions is recommended.
Nothing contained in this publication shall be construed as a grant of any right of manufacture, sale,
use, or reproduction, in connection with any method, process, apparatus, product, composition, or
system, whether or not covered by letters patent, copyright, or trademark, and nothing contained in this
publication shall be construed as a defense against any alleged infringement of letters patent,
copyright, or trademark, or as a defense against liability for such infringement.
ASM International is the society for materials
engineers and scientists, a worldwide network
dedicated to advancing industry, technology, and
applications of metals and materials.
ASM International, Materials Park, Ohio, USA
www.asminternational.org
This publication is copyright © ASM International®. All rights reserved.
Publication title
Product code
Stainless Steels for Design Engineers
#05231G
To order products from ASM International:
Online Visit www.asminternational.org/bookstore
Telephone 1-800-336-5152 (US) or 1-440-338-5151 (Outside US)
Fax 1-440-338-4634
Mail
Customer Service, ASM International
9639 Kinsman Rd, Materials Park, Ohio 44073-0002, USA
Email CustomerService@asminternational.org
American Technical Publishers Ltd.
27-29 Knowl Piece, Wilbury Way, Hitchin Hertfordshire SG4 0SX,
In Europe United Kingdom
Telephone: 01462 437933 (account holders), 01462 431525 (credit card)
www.ameritech.co.uk
Neutrino Inc.
In Japan Takahashi Bldg., 44-3 Fuda 1-chome, Chofu-Shi, Tokyo 182 Japan
Telephone: 81 (0) 424 84 5550
Terms of Use. This publication is being made available in PDF format as a benefit to members and
customers of ASM International. You may download and print a copy of this publication for your
personal use only. Other use and distribution is prohibited without the express written permission of
ASM International.
No warranties, express or implied, including, without limitation, warranties of merchantability or
fitness for a particular purpose, are given in connection with this publication. Although this
information is believed to be accurate by ASM, ASM cannot guarantee that favorable results will be
obtained from the use of this publication alone. This publication is intended for use by persons having
technical skill, at their sole discretion and risk. Since the conditions of product or material use are
outside of ASM's control, ASM assumes no liability or obligation in connection with any use of this
information. As with any material, evaluation of the material under end-use conditions prior to
specification is essential. Therefore, specific testing under actual conditions is recommended.
Nothing contained in this publication shall be construed as a grant of any right of manufacture, sale,
use, or reproduction, in connection with any method, process, apparatus, product, composition, or
system, whether or not covered by letters patent, copyright, or trademark, and nothing contained in this
publication shall be construed as a defense against any alleged infringement of letters patent,
copyright, or trademark, or as a defense against liability for such infringement.
Stainless Steels for Design Engineers
Michael F. McGuire, p 1-10
DOI: 10.1361/ssde2008p001
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
CHAPTER 1
Metallurgy
Summary
COMPARED TO ALLOY STEELS, stainless
steels are chemically complex. The large number
of alloying elements makes possible a larger
range of possible phases or basic crystal structures. The large amount of the alloying elements
makes the deviation from the behavior of pure
iron greater; consequently, the calculations that
predict which phases will exist are more difficult.
The three basic phases of stainless steels are ferrite, austenite, and martensite. The wide variety
of alloys that exist is based on:
• Combinations of these phases
• Altering the composition of these phases
• Adding secondary phases for particular
purposes
Metallurgy, as discussed in this chapter, focuses on phases normally encountered in stainless steels and their characteristics. In subsequent
chapters on types of stainless steel, there are
more detailed treatments of the alloys made of
these phases and their properties.
Introduction
Most widely used alloy systems, such as carbon steels, alloy steels, and aluminum alloys, are
relatively dilute solutions of several elements in
the parent matrix. Carbon and alloy steels, with
very few exceptions, are principally of the magnetic body-centered cubic (bcc) phase or a
slightly distorted version of it. Aluminum alloys
share the face-centered cubic (fcc) structure of
pure aluminum. A given structure, which can
have a certain range of compositions, is what is
meant by a phase, just as a gas or liquid is a
phase. In solid metals, there can be a number of
phases coexisting simultaneously. Stainless
steel is an exceptional alloy system in that it is
not a dilute solution. Alloy steels may contain
several percent of alloying elements, such as
carbon, manganese, nickel, molybdenum,
chromium, and silicon, in addition to the impurities sulfur, oxygen, and phosphorus. Alloy
steels typically contain very small amounts of
titanium, niobium, and aluminum. The total
amount* of these alloying elements seldom exceeds 5%. The same is true for most aluminum
alloys. In contrast, stainless steels contain no
less than about 11% chromium alone. Most
stainless alloys have manganese, silicon, carbon, and nickel in thermodynamically meaningful amounts as well as large concentrations
of nickel and/or molybdenum.
The result of the large number of alloying elements in relatively high concentrations is that
stainless steel can have many stable phases concurrently. In almost every case, having phases
other than the principal one or two phases for
which the alloy was designed is undesirable because of the possibility of undesirable variations
in mechanical or corrosion performance. The
producer of stainless steel controls the chemical
composition and thermomechanical processing,
so that when the processor or end user receives
the product it is usually in the correct condition.
However, subsequent processing or service conditions may alter the carefully established phase
structure. Therefore, it is necessary to discuss
the phases that can exist in stainless steel and
the conditions under which they form so that the
enlightened user will know which phases to
avoid and how to avoid them.
It is possible to use thermodynamics to calculate which phases may exist at a given tempera* All compositions are given in weight percent unless
stated otherwise.
2 / Stainless Steels for Design Engineers
ture for a given composition. It is not remotely
feasible, however, to give an adequate treatment
of the thermodynamics required to do this. The
topic alone requires a book. The necessary
knowledge has been embedded in proprietary
computer programs that will be used instead.
Thermodynamics of Stainless Steel
Pure metals, from a practical viewpoint, are
either liquid or solid depending on temperature,
with the possibility of some trivial small gas
vapor pressure. A law of thermodynamics is that
the number of possible condensed (i.e., solid)
phases equals the number of elemental constituents plus one. The solid has a crystallographic structure that may vary with temperature. Many metals have a less-dense bcc
structure at high temperature and transform to a
denser fcc structure at lower temperatures. Iron
does this. Iron has the curious characteristic of
transforming from fcc back to the low-density
bcc at still lower temperatures. This is a result of
the unpaired 3d orbital electrons (those that give
rise to ferromagnetism) that are not given up as
valence electrons, causing repulsive forces between atoms and requiring a more widely spaced
structure.
All thermodynamic properties are based on
interatomic attractions. In metals, the metal
atoms give up valence electrons to the entire
mass. These electrons are of varying energy
states and highly mobile. They are responsible
for the ability of metals to conduct heat and
electricity well. The attraction, the strength of
the bond, is proportional to the charge difference and distance. The attraction determines
such macroscopic properties as melting temperature, density, and elastic modulus.
In this book, the main concern with thermodynamics is predicting which phases are present
both at equilibrium and in the quite frequent
metastable state. The prediction involves calculating the free energy of the various possible
phases. The phase with the lowest energy is
most favored, but others may have free energies
that permit them to exist. The difference between these two is that the equilibrium state,
that of the lowest free energy, may require
atomic rearrangements to occur for equilibrium
compositions to be reached on an atomic scale.
If diffusion is too sluggish for these rearrangements to take place, the structure may retain the
prior metastable structure indefinitely. This is
not a small, pedantic point. Most stainless steels
are used in the metastable condition. For example, the common alloy 304 (also called 18-8) is
normally used in the fully austenitic condition.
It would “rather” be partly ferritic, but the substitutional diffusion of chromium in austenite
that is required to form a ferrite phase of a separate composition is so slow that it cannot occur
in terrestrial time frames. However, if energy is
applied by mechanical shear, the austenite can
transform without diffusion to the lower freeenergy martensite phase, a quasi-bcc structure
of lower free energy.
The calculation of which phases exist under
equilibrium conditions proves to be extraordinarily difficult in complicated alloy systems.
This is because thermodynamic values can be
measured accurately only in the liquid state, so
the values for the solid state are extrapolations.
Also, the interaction between elements is very
important in nondilute alloys such as stainless
steel. Consequently, most published phase diagrams are experimentally derived. To determine
which phases exist at a given composition and
temperature, a sample is made, equilibrated at
the appropriate temperature, and quenched to
room temperature. It is assumed that the characteristic equilibrium phases have been frozen and
are then identified by various techniques for
structure, composition, and the like. This important work is obviously tedious and susceptible
to experimental error and applies only to specific compositions. Any “what if” extrapolation
to a different alloy composition carries the risk
of error.
A practical tool has been developed that permits phase diagrams to be calculated for arbitrary
compositions. These are computer simulated,
mathematical models that can perform the complex thermodynamic calculations. To do this with
accuracy requires databases of thermodynamic
values. These values must be derived from computer analysis of experimental phase equilibrium
diagrams. They are expensive to derive and validate, and only a few exist. Hence, they are proprietary. In Appendix 3, a license to one such
program, Thermo-Calc, can be found. The version has a reduced three-element capability but
uses the same proprietary thermodynamic database of the full version. The program allows determination of which phases can exist for any
composition and temperature. Whether the
phases will form depends also on kinetic factors.
First, however, it is good to become familiar with
the principal phases found in stainless steel.
Chapter 1: Metallurgy / 3
Phases
Ferrite
The basis of stainless alloys is, of course,
iron. Iron, as stated, solidifies as a bcc alloy before transforming to the denser fcc austenite at
lower temperatures. At still lower temperatures,
it reverts to the bcc structure. It is accurate to
surmise from this that the free energy of both
structures is close. Alloying elements that promote one structure over the other can therefore
change which one predominates. The element
that produces the ability to form the passive film
that makes stainless corrosion resistant,
chromium, has the characteristic of stabilizing
the bcc structure. As chromium is added to iron,
the temperature range over which austenite is
stable grows smaller until, at about 12%
chromium, ferrite is stable at all temperatures.
This is, coincidentally, the approximate level of
chromium needed to keep alloys from rusting
under ambient conditions, but this effect is not
related to whether the structure is bcc or fcc.
The iron-chromium phase diagram (Fig. 1)
shows the composition and temperature regions
where ferrite (a), martensite (α' ), austenite (γ),
and sigma phase (σ) are stable.
While chromium is the principal ferrite-promoting alloying element, other elements have
similar effects, but none produces the quality of
stainlessness. Silicon, aluminum, molybdenum,
tungsten, niobium, and titanium all favor ferrite.
Carbon, nitrogen, manganese, nickel, and copper do not and expand the temperature range
over which austenite exists. Elements that are
insoluble in iron at austenite-forming temperatures, such as the impurities phosphorus, sulfur,
Fig. 1 The iron chromium phase diagram. Courtesy of
Thermo-Calc Software
and oxygen, have no influence on which phase
is favored. Again, it must be emphasized that
the influence of an alloying element on structure
has zero bearing on its influence on corrosion
resistance.
The elements that promote ferrite over austenite also have the effect, at still lower temperatures, of promoting intermetallic compounds
generally composed of iron, chromium, and
some of those alloying elements. These are discussed separately.
Metals are effective solvents in both the liquid and solid states. An important part of steelmaking is refining the molten metal to remove
the undesired impurities dissolved in it. The
normal technique is to add elements that react
selectively with the targeted impurities to form
an immiscible reactant that can become part of
the slag and physically separated from the refined alloy. This is done for the primary impurities oxygen and sulfur. A third common impurity, phosphorus, is not so easily removed and
must be excluded from raw materials to be kept
under control.
In stainless steel, carbon and nitrogen can be
detrimental impurities. Both are quite soluble
in molten iron-chromium alloys and are fairly
soluble in ferrite at high temperatures. This solubility decreases exponentially with temperature so that it is essentially zero at room temperature. These elements have small atomic
sizes compared to iron and chromium and,
when dissolved, squeeze into interstitial sites
within the bcc matrix. Such interstitial solute
atoms profoundly distort the structure. They are
much more soluble in the fcc structure, which,
while denser, has roomier interstitial spaces, so
they stabilize that structure. To preserve the ferrite structure, carbon and nitrogen must be
eliminated.
There are additional reasons to eliminate carbon and nitrogen. During cooling as these elements become less and less soluble, they must
precipitate. The most thermodynamically favorable form in which they can precipitate is as a
compound of chromium, with which they are
very reactive. This occurs at the grain boundaries, where nucleation is favored, and depletes
those regions of chromium, rendering them less
corrosion resistant. A second effect is a loss of
toughness due to these precipitates. The diffusion
rates of carbon and nitrogen in ferrite are too
high to prevent this precipitation by quenching.
Modern refining methods can reduce carbon
plus nitrogen to under 0.020%, but even this is
4 / Stainless Steels for Design Engineers
too high. So, to avoid the detrimental effects of
chromium carbide and nitride formation in ferrite, other benign carbides and nitrides such as
those of titanium or niobium are allowed to form
preferentially. This approach is called stabilization and is used for most ferritic alloys today. The
older approach, as characterized by alloy 430, is
to permit chromium carbides and nitrides to form
but then to perform a subcritical anneal to rehomogenize the chromium and coarsen the precipitates so that they have only a small negative
effect on mechanical properties.
Hydrogen and boron are other elements that
can be interstitially dissolved in ferrite. Boron is
normally found at levels of around 5 to 10 ppm.
At higher levels, boron substitutes for carbon in
carbides. Hydrogen is soluble to several parts
per million by weight. It does not cause hydrogen embrittlement in annealed ferrite. If the ferrite is cold worked, the solubility of hydrogen
increases as the defect structure accommodates
hydrogen atoms. In this condition, ferrite may
be embrittled by hydrogen, especially if it enters the metal through corrosion processes like
pitting. This is one explanation of, and the most
likely explanation for, stress corrosion cracking.
While hydrogen is easily removed by argon
oxygen decarburization (AOD), assuming absolutely dry blowing gases and additions are
used, it can be picked up during pickling, welding, or annealing as well as by corrosion.
All stainless alloys rely on having a uniform
level of chromium and the other element,
molybdenum, which assists in corrosion resistance, distributed throughout the matrix. If there
are locally low levels of these elements, localized resistance to corrosion is reduced, and localized corrosion can occur. This can occur by
the precipitation of any phase that is richer in
chromium or other corrosion-resisting elements. Because chromium is a reactive element, its success depends to a great degree on
maintaining the homogeneity required for
proper corrosion-resistant performance. Incorrect thermal processing is the main way homogeneity can be lost. Stabilizing makes it much
easier to keep chromium from segregating in
ferritic alloys.
A by-product of stabilization with titanium is
that oxygen and sulfur are also eliminated as
compounds of titanium along with carbon and
nitrogen. These impurity elements would otherwise also precipitate as compounds containing
some chromium, potentially depleting chromium
in the vicinity of their precipitation.
The bcc structure of ferrite allows more
rapid diffusion than does the fcc structure of
austenite. This is true for both the interstitial
diffusion of the elements helium, boron, carbon, nitrogen, and oxygen and the substitutional diffusion of all other elements. The rate
of diffusion of all elements, both interstitial
and substitutional, in ferrite is about two or
three orders of magnitude higher than in
austenite. The practical implication of this is
that precipitation reactions generally cannot be
suppressed by quenching in ferrite if they involve interstitial elements, whereas they can
be in austenite. Intermetallic phases can form
more rapidly in ferrite. This becomes an issue
only when total chromium plus molybdenum
exceeds about 20%, above which the sigma
phase appears. This is thus only an issue for
superferritic (high chromium content) alloys or
for the ferrite phase of duplex (ferrite-austenite) alloys.
The mechanical properties of the ferrite phase
are discussed extensively in Chapter 8, “Ferritic
Stainless Steels.” Here, it is only necessary to
note that ferrite in stainless steel closely resembles low-carbon steel in mechanical behavior. It
shares the following characteristics:
• A toughness transition that occurs around
room temperature
• Notch sensitivity
• A yield point phenomenon
• Pronounced crystallographic anisotropy of
mechanical properties
• High stacking fault energies and low workhardening rates
These issues are dealt with in the same way as
in carbon steel when these characteristics become an issue. The first two are controlled by
reduction of interstitial levels and refining of
grain size. The yield point is eliminated by
slight elongation by temper rolling or elimination of interstitial carbon and nitrogen, whose
interaction with dislocations causes the yield
point. The anisotropy is either utilized to advantage by maximizing it, as in the case of deepdrawing alloys, or minimized by refining grain
size and randomizing grain orientation by special thermomechanical processing.
Ferrite has a greater thermal conductivity
and lower thermal expansion than austenite. Its
strength decreases with temperature more than
that of austenite, but the good match in thermal
expansion between the ferrite and its oxide
still makes it an excellent high-temperature
Chapter 1: Metallurgy / 5
material. Ferrite has very nearly the same corrosion resistance as austenite, but since ferrite
can hold no nitrogen in solution, it cannot benefit from this element. In duplex alloys, the
ferrite is generally the more corrosion resistant
phase because it is richer in chromium and
molybdenum.
Austenite
The second major constituent phase of the
stainless steel alloy system is austenite. Austenite has an fcc atomic structure. The fcc structure
is common in many transition metals to the right
of iron in the periodic table. As stated, the fcc
structure should be considered normal for metals
well below their melting temperature as it is a
denser structure. The presence of the bcc structure relates to the unpaired 3d electrons, which
provide ferromagnetism. Adding elements to
iron that causes pairing of the 3d electrons diminishes ferromagnetism and promotes the fcc
structure. Nickel and manganese are the most
prominent alloying elements that do this, but the
interstitials carbon and nitrogen are the most
powerful austenite stabilizers on a percentage
basis. Their use is limited by their solubility and
their tendency to form precipitating compounds
with chromium. Manganese acts largely through
its ability to promote nitrogen solubility. Superaustenitic stainless steels, such as S34565, use 4
to 6 % manganese to permit nitrogen levels of
0.4 to 0.6% to be achieved, resulting in higher
pitting corrosion resistance.
Since all stainless steels contain principally
iron and chromium, the addition of a substantial
amount of austenitizing elements is necessary to
transform the structure to austenite. As a rule of
thumb, iron alloys require about 17% chromium
and 11% nickel (or its equivalents) to remain
austenitic at room temperature. One percent
nickel can be replaced by about 2% manganese
as long as nitrogen is present to maintain the
same phase stability. The omnipresent carbon
and nitrogen have an effect 30 times that of
nickel, so even in the small amounts in which
they are normally present, they have a significant effect. These stabilizing factors are mapped
in the Schaeffler diagram of Fig. 2 (Ref 1),
whose purpose is to predict the phase makeup
of weld metal. Since welds solidify relatively
rapidly, no carbides or intermetallic phases
Fig. 2 Schaeffler-Delong constitution diagram showing phases present in as-solidified stainless steels at room emperature as a function of composition demonstrating carbon and nitrogen contributions to nickel effects. Adapted from A.L. Schaeffler,
Constitution Diagram for Stainless Steel Weld Metal, Met. Prog., Vol 56, Nov 1949, p 680–688; and W.T. Delong, A Modified Phases
Diagram for Stainless Steel Weld Metals, Met. Prog., Vol 77, Feb 1960, p 98
6 / Stainless Steels for Design Engineers
form, and only ferrite, austenite, and martensite
will be present. Thus, they provide useful information about the compositional effects on phase
development in nonequilibrium situations. The
nickel equivalent (vertical axis) summarizes
how nitrogen, carbon, and other elements combine to create a nickel-like effect. The horizontal axis does the same for chromium and those
elements that have a similar effect.
In most common stainless steels, austenite is
normally present in the metastable state, for example, the retained austenite in alloy steels. Those
with carbon above 0.02% would eventually break
down into austenite plus carbides, and those with
less than about 30% chromium plus nickel will
form martensite if deformed sufficiently. But in
the annealed state, the austenite in standard
austenitic stainless steels will remain indefinitely
as fully austenitic without precipitates unless
heated above 400 °C (750 °C) for protracted periods of time or deformed extensively.
Interstitial elements are much more soluble in
austenite than in ferrite. Of these, only nitrogen
is considered a beneficial alloying element. It
both strengthens and improves the pitting corrosion resistance of austenite. Carbon has a parallel effect, but its tendency to form chromium
carbides limits its use and in fact leads to its
minimization in most alloys. Before the AOD
was developed and carbon levels in stainless
steels were higher, austenitic stainless steels
were sometimes stabilized by titanium or niobium to counter the effects of carbon. Both carbon and nitrogen stabilize the austenite phase,
permitting lower levels of nickel to be used in
austenitic alloys.
Interstitial atoms of carbon and nitrogen distort the fcc lattice, causing it to expand about
1% linearly per 1 wt% of solute (Fig. 3) (Ref 2).
This produces solid solution hardening of the
austenite. The work hardening of austenite is increased by nitrogen. A third interstitial solute,
hydrogen, produces the same effect but to a
lesser degree. Austenite is not embrittled by hydrogen to the extent ferrite or martensite is, but
hydrogen does raise its flow stress and hardness
while lowering its work-hardening rate.
Sulfur and oxygen are considered impurities
because they form inclusions, usually chrome/
manganese silicates and sulfides. If present in
sufficient amounts, sulfur and oxygen precipitate
as primary inclusions before or during solidification. In most austenitic stainless alloys, the remainder of these elements are near saturation in
the as-solidified ferrite at very high temperatures
and then frozen in a state of supersaturation in
the austenite when it forms on cooling. The sulfur and oxygen then precipitate during cooling or
subsequent hot working as isolated inclusions.
The interface between these inclusions and the
matrix is the locus of corrosion pit initiation,
quite probably because of chromium depletion
occurring during and as a result of inclusion
growth. When an alloy solidifies as austenite,
sulfur immediately segregates to the grain
boundaries because of its low solubility in
austenite, and it forms a low-strength film with a
low melting temperature. This causes poor hot
workability and hot cracking of welds.
The diffusion rates in austenite are quite low
compared to ferrite, so even interstitial elements
cannot move quickly enough to precipitate below
about 400 °C (750 °F). This permits carbon and
nitrogen to exist in very high degrees of supersaturation if introduced below this temperature, as
is done by various proprietary processes. The
low diffusion rates restrict such colossally supersaturated zones to thin surface layers, but they
can reach phenomenal hardness of over Rc 70.
The austenite structure does not discourage the
formation of intermetallic compounds such as
sigma, but it does, fortunately, make their formation very sluggish, as seen in Fig. 4. The difference of three orders of magnitude for carbide
formation reflects the difference between the diffusion of carbon and that of substitutional elements. The formation of sigma in ferrite is about
100 times faster than in austenite. Sigma is almost never seen in commercial 316 alloys.
Fig. 3 Lattice expansions due to carbon. Source: Ref 2
Chapter 1: Metallurgy / 7
The mechanical properties of austenite are
quite different from those of ferrite. Austenite is
characterized by:
• Low stacking fault energies leading to high
work-hardening rates
• Good toughness even at very low temperatures
• Low notch sensitivity
• Lack of a sharp elastic limit
• Good high-temperature strength
• Fairly isotropic mechanical properties
While there is not a great deal of difference
in the yield strengths of austenitic and ferritic
alloys of similar alloy levels, austenitic alloys
are more ductile, have high work-hardening
rates, and therefore have higher tensile
strengths. Austenite can be cold worked to extremely high strengths, around a maximum of
2000 MPa (290 ksi). Chapter 3, “Austenitic
Stainless Steels,” gives a more thorough and
quantitative treatment of the mechanical properties of austenite.
In duplex stainless steels, a secondary austenite, γ2, can form from ferrite below 650 °C
(1200 °F). At this temperature, it has the same
composition as the ferrite from which it forms
and is called type 1. In the 650 to 800 °C (1200
to 1470 °F) range, a range that can be encountered in the heat-affected zone (HAZ) at γ/δ
boundaries during welding, another type forms.
This so-called secondary austenite, γ2, type 2, is
somewhat enriched in nickel over the ferrite
from which it forms but poorer in nitrogen than
the primary austenite, giving it poorer corrosion
resistance. Secondary austenite can also coform
with sigma as γ/δ grain boundaries are depleted
of chromium. This secondary austenite is called
type 3 and is also poor in chromium.
Fig. 4 Precipitation kinetics in 316 stainless steel. Source: Ref 3
The physical properties of austenite compared to ferrite include lower thermal and electrical conductivity and greater thermal expansion. It is also, of course, nonmagnetic.
Martensite
Martensite is a phase that forms from the diffusionless shear of austenite to a distorted cubic
or hexagonal structure. This transformation can
occur spontaneously on cooling or isothermally
with externally applied deformation. It is essentially ferrite that has been formed with a supersaturation of carbon. The resulting structure is
very fine and highly faulted, making it quite
hard. As in carbon steel, the hardness of the
martensite increases dramatically with interstitial content because of the huge strain interstitial elements impose on the bcc lattice, distorting it into tetragonality.
Martensite in stainless steels is restricted to
alloy levels at which austenite can form at
higher temperatures but at which the austenite is
unstable at ambient temperatures. This gives
martensite a fairly narrow composition range.
The lowest alloy level is that of the basic 12%
chromium steels with 0.1 to 0.2% carbon. The
most highly alloyed martensites are found in the
precipitation-hardening grades. Thus, martensitic stainless steels are inherently limited in
corrosion resistance to a level no better than a
17 or 18% chromium alloy and often barely
qualify as stainless after the chromium tied up
as chromium carbide is recognized as not contributing to the corrosion resistance.
The as-formed martensite to the degree it has
significant carbon content is hard and requires
tempering to give it adequate toughness. The
tempering reaction is the precipitation of carbon in the form of carbides with the concurrent
loss of internal strain in the martensite lattice.
The complexities of tempering require its discussion in detail to be found in Chapter 3,
“Martensitic Stainless Steels.” It is worth noting,
however, that all tempering involves carbide formation, thus losing some corrosion-fighting
chromium.
There are two forms of martensite, the ε, epsilon, and the α', alpha prime. Epsilon is formed
in steels with low stacking fault energy, which
are primarily the leaner austenitic alloys. Thus, it
forms at cryogenic temperatures or by cold
working. It appears in martensitic alloys of the
precipitation-hardening type. It is nonmagnetic,
has a hexagonal close-packed (hcp) structure,
8 / Stainless Steels for Design Engineers
and is very difficult to identify microscopically.
The a' martensite is the familiar magnetic variety known in alloy steels that forms both by
quenching and by deformation.
The mechanical properties of stainless
martensite are parallel to those of alloy steels.
The high quantity of alloying elements in
stainless give an extreme depth of hardening,
so there is no concern with ancillary phases
such as bainite. The physical properties are
very close to those of ferrite of the same composition.
Intermetallic Phases
The number of phases that can coexist in an
alloy is proportional to the number of alloying
elements in the alloy. Table 1 lists data on the
more common precipitates found in stainless
steel. It is not surprising that stainless steel with
iron, chromium, nickel, manganese, silicon, and
often molybdenum, titanium, and niobium
should have numerous ancillary phases. Intermetallic phases are normally hard and brittle.
They can render the bulk alloy brittle when they
form along grain boundaries. The other concern
arising from intermetallic phase formation is the
depletion from the surrounding matrix of
Table 1
steels
Precipitated phases found in stainless
Precipitate
Structure
Parameter, A
Composition
NbC
fcc(a)
a = 4.47
NbC
NbN
fcc
a = 4.40
NbN
TiC
fcc
a = 4.33
TiC
TiN
fcc
a = 4.24
TiN
Z-phase
Tetragonal
a = 3.037 c = 7.391
CrNbN
M23C6
fcc
a = 10.57–10.68
Cr16Fe5Mo2C (e.g.)
M23(C,B)6 fcc
a = 10.57–10.68
Cr23(C,B)6
M6C
a = 10.62–11.28
(FeCr)21Mo3C;
Fe3Nb3C; M5SiC
Diamond
cubic
M2N
Hexagonal
a = 2.8 c = 4.4
Cr2N
MN
Cubic
a = 4.13–4.18
CrN
Gamma
prime
fcc
a = 3.59
Ni3(Al,Ti)
Sigma
Tetragonal
a = 8.80 c = 4.54
Fe, Ni, Cr, Mo
Laves
phase
Hexagonal
a = 4.73 c = 7.72
Fe2Mo, Fe2Nb
Chi phase bcc(b)
a = 8.807–8.878
Fe36Cr12Mo10
G-phase
a = 11.2
Ni16Nb6Si7,
Ni16Ti6Si7
fcc
(a) fcc, face-centered cubic. (b) bcc, body-centered cubic.
chromium or molybdenum, causing localized
lower corrosion resistance. Intermetallic phases
form by diffusion of substitutional alloying elements, which makes their precipitation slower
than that of carbides, but they can form in a
matter of minutes in alloy-rich grades. Deformation, which enhances substitutional diffusion, accelerates their formation. The principal
intermetallic phases are described next.
Alpha Prime. Not to be confused with
martensite, alpha prime is an ordered ironchromium phase (i.e., iron and chromium atoms
occupy specific, rather than random, sites on
two intersecting superlattices). This structure is
quite brittle. It forms at relatively low temperatures, between 300 and 525 °C (570 and 980
°F). Before its true nature was understood, its
presence was known through its causing the
phenomenon called 475 embrittlement, originally called 885 °F embrittlement. This is sometimes confused with temper embrittlement,
which occurs in the same temperature range but
is caused by phosphide precipitation on prior
austenite grain boundaries of martensite. Alpha
prime precipitation can cause 475 embrittlement in ferritic or duplex stainless steels and
limits their use in this temperature range but not
at higher temperatures, at which the phase dissolves. This phase forms at chromium contents
as low as 15%, but fortunately it takes a relatively long time to form, on the order of hours,
so it will not occur inadvertently during thermal
processing such as welding or annealing.
Sigma. Sigma is a brittle tetragonal phase
richer in chromium and molybdenum than either the ferrite or austenite matrix around it. It
forms preferentially at ferrite-austenite boundaries in the temperature range 600 to 1000 °C
(1110 to 1470 °F) in alloys with more than
about 18% chromium plus molybdenum. Its
composition is sometimes given as (CrMo)35
(FeNi)65, but examination of the iron-chromium
phase diagram shows that it is archetypically an
equiatomic iron chromium compound. It is
strongly promoted by silicon and suppressed by
nitrogen. Stabilized alloy grades show more
rapid sigma formation than unstabilized alloy
grades (e.g., 347 versus 304). In unstabilized alloys the prior precipitation of carbides destabilizes austenite, leading to subsequent sigma formation. This makes alloys like 310H,
essentially 25Cr-20Ni, especially prone to
sigma formation.
Sigma forms much more rapidly from ferrite
than from austenite because of the 100-fold
Chapter 1: Metallurgy / 9
higher diffusion rate of alloy elements in ferrite.
This makes it a much larger issue in superferritic and duplex alloys, which have high
chromium and/or molybdenum levels. Chapter
7, “Duplex Stainless Steels,” contains an indepth discussion of sigma.
Chi. Chi, χ, is similar to sigma except it contains more molybdenum and less chromium and
has a cubic structure. It can coexist with sigma
and forms in the same temperature range. It also
precipitates at ferrite-austenite boundaries and
has the same deleterious effects.
Laves Phase. The laves phase has the structure A2B where A is iron or chromium and B is
molybdenum, niobium, titanium, or silicon. It
forms at 550 to 650 °C (1020 to 1200 °F) over
the course of hours. Thus, although its effect
would be deleterious, it seldom becomes a practical problem. It is possible for it to form at temperatures below sigma and above alpha prime,
but the long times for formation make it rare.
Carbides, Nitrides, Precipitation Hardening,
and Inclusions
Carbon and nitrogen are very important in all
steels, but they take on a special significance in
stainless steel because chromium, the essential
alloying element of stainless steel, reacts more
vigorously with carbon and nitrogen than iron
does. Except for its role in hardening martensite
and strengthening austenite at high temperatures, carbon is almost universally a detrimental
impurity from a corrosion point of view and is
minimized. Its beneficial effect on corrosion resistance when it is in solution is negligible
because so little of it can be held in solution.
Nitrogen has a lesser tendency to form compounds with chromium, so it is considered a
beneficial alloying element in austenite but not
in ferrite, in which it has essentially zero solubility. Common carbide and nitride precipitating
phases are also listed in Table 1.
Carbides. M23C6 is the main carbide found in
stainless steel. Its structure is orthorhombic, and
it contains both iron and chromium. It can form
at any temperature at which the host austenite or
ferrite becomes saturated with carbon. It is
mainly chromium carbide, but iron can substitute
for chromium up to about 50%. Other elements,
such as tungsten, vanadium, and molybdenum,
can also dissolve in this carbide. The ratio of
chromium to iron in the carbide increases with
time and temperature, as chromium diffusion
permits, up to a maximum of 4 or 5 to 1.
The precipitation of the carbide from ferrite
occurs at grain boundaries, is extremely rapid,
and cannot be suppressed by quenching. Less
than 20 ppm carbon content is required to prevent its precipitation from ferrite, although up
to 50 ppm can be effectively kept in solution by
very vigorous quenching. From austenite, carbide precipitation occurs below about 900 °C
(1650 °F) for carbon levels under 0.10% and at
650 °C (1200 °F) for carbon levels below
0.03%. For practical purposes, precipitation
ceases below 500 °C (930 °F) due to the slowing diffusion of carbon. While carbon is essentially insoluble in austenite at room temperature, quenching can easily preserve up to 0.10%
in supersaturation, as is commonly seen in type
301 stainless.
The carbide precipitation occurs first at grain
boundaries. The chromium that combines with
the carbon comes from the matrix in the immediate vicinity and therefore decreases the
chromium content of that region, giving rise to
the phenomenon of sensitization, which comes
from the original phrase “sensitization to intergranular corrosion.” Nickel and molybdenum
decrease the solubility of carbon and thus accelerate the precipitation. Nitrogen retards precipitation. Cold work accelerates precipitation. The
carbide has a hardness of about Rc 72. This
makes the phase a useful constituent in wear
resistance in martensitic alloys.
In higher carbon grades such as the martensitic stainless alloys, additional, more carbonrich, carbides may form. These include M7C3
and M3C. The latter carbide forms during the
low-temperature tempering of martensite, while
the former precipitates at higher temperatures.
Stabilizing carbides are those that are formed
by the intentional addition of elements such as
titanium and niobium. These elements form carbides of the type MC (metal carbide). The carbon in these compounds may be replaced by nitrogen or, in the case of titanium, sulfur. These
carbides form preferentially over chromium carbides and thus prevent sensitization. They precipitate in both the liquid and solid states. In the
solid state, the precipitate normally forms within
grains. The Ti(CN) appears as a cube of gold
TiN surrounded by gray TiC. The Nb(C,N) is
less regularly shaped. They affect mechanical
properties in ferrite both by their influence on
recrystallization and by their ability to act as nucleation sites for brittle fracture
Nitrides. At low levels, nitrogen can substitute for carbon in M23C6. At higher nitrogen
10 / Stainless Steels for Design Engineers
levels, Cr2N can form. This can occur in duplex
alloys if they are heated to a solution annealing
temperature at which the alloy has high solubility for nitrogen. Cooling from these temperatures can cause the excess nitrogen to precipitate as needles of Cr2N. Another nitride CrN
can form in HAZs of welds.
Precipitation-Hardening Phases. Phases
that have a very similar lattice match to the parent phase can precipitate coherently, that is,
without changing the continuity of the crystal
lattice. In these cases, the slight mismatch
causes a strain that can significantly restrict dislocation movement and thereby strengthen the
matrix. One such precipitate is gamma prime,
an intermetallic, ordered, fcc phase with the
composition Ni3(AlTi). Copper forms the epsilon phase, essentially pure copper, which
causes precipitation hardening. The secondary
hardening of martensite due to the precipitation
of molybdenum nitride or carbide is also a precipitation-hardening reaction.
Inclusions. Inclusions are principally oxides
and sulfides that form in the melt (type I), at the
end of solidification (type II), or in the solid
(type III). Type I inclusions are the largest and
are globular. Except when they are deliberately
kept to improve machinability, they are physically removed by various steelmaking practices.
Type II inclusions form in interdendritic spaces
as the solubility of oxygen and sulfur drop on
solidification. Type III inclusions precipitate the
remaining oxygen and sulfur, up to 100 ppm for
normal manganese-silicon killed stainless
steels, in the solid state either on preexisting inclusions or as micron-size particles. Inclusions
are mainly oxides and sulfides of silicon and
manganese. If more reactive elements, such as
aluminum or titanium, are present, their oxides
and sulfides can also be present.
Sulfides and oxysulfides can be beneficial for
machining as solid-state lubricants and chip
breakers. Otherwise, their presence is detrimen-
tal as inclusions have been shown to be the initiation sites for corrosion pits, which have been
linked to both their sulfur ions disrupting the
passive layer and their chromium content causing slight local chromium depletion.
Properties of Stainless Steels
Physical and mechanical properties of representative stainless steel alloys are summarized
in Appendix 2. Properties are also discussed in
chapters specific to each alloy family. The
reader is referred to primary sources, such as
company web sites, such as Ref 4 and 5.
REFERENCES
1. D.J. Kotecki, Welding of Stainless Steels,
Welding, Brazing, and Soldering, Vol 6,
ASM Handbook, ASM International, 1993,
p 677–707
2. G.E. Totten, M. Narazaki, R.R. Blackwood, and L.M. Jarvis, Failures Related to
Heat Treating Operations, Vol 11 ASM
Handbook, ASM International, 2002, p
192–223
3. High Performance Stainless Steels, Reference Book Series 11 021, Nickel Development Institute, p 16
4. ASM Handbook, Vol 1, Properties and
Selection, ASM International, 1990
5. ASM Speciality Handbook, Stainless Steels,
ASM International, 1996
SELECTED REFERENCES
• D.J. Kotecki and T.A. Siewert, WRC 1992
Constitution Diagram, Welding Journal, Vol
5, 1992, p 171s–178s
Stainless Steels for Design Engineers
Michael F. McGuire, p 11-18
DOI: 10.1361/ssde2008p011
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
CHAPTER 2
Corrosion Theory
Summary
Electrochemical Reactions
THIS CHAPTER INTRODUCES THE fundamentals of electrochemical theory as it pertains
to corrosion. Topics covered include an overview
of electrochemical reactions, Faraday’s law, the
Nernst equation, galvanic versus electrochemical
cells, and Pourbaix diagrams. The examples provided relate these fundamentals to the corrosion
resistance of stainless steels.
In electrochemical reactions, charge is transferred across interfaces of species of different
chemistry. Consider, for example, the reaction:
Introduction
Corrosion—the environmental degradation
of materials through electrochemical reactions—is a key subject for more or less all
classes of alloys that fall within the broad definition of stainless steels because these alloys
were developed with the intention of preventing corrosion. This chapter aims first to provide
an introduction to the fundamentals of electrochemical theory as it pertains to corrosion.
Thermodynamics are presented in light of electrochemical potentials as opposed to purely
chemical ones. Chapter 3 introduces the formal
terms needed to describe electrode reaction kinetics. Chapter 4 describes the various forms of
corrosion and how they are related to alloy
metallurgy, chemistry, and structure. Chapter 5
focuses on oxidation. For an in-depth study of
electrochemical kinetics and electroanalytical
methods, Ref 1 is recommended. For a broader
study of corrosion, the reader is referred to
texts by Jones (Ref 2), Uhlig and Revie (Ref 3),
and Fontana (Ref 4) and to ASM Handbook,
Volume 13A (Ref 5).
2Fe (s) + O 2 (g) + 2H 2 O → 2Fe 2+ + 4 OH −
(Eq 1)
An inspection of this reaction suggests that
three phases must be present for the reaction to
proceed: an ion-conducting phase (water-based
solution), a metallic phase (iron), and a gas
phase O2(g). Second, electrons have been transferred from the metallic phase, iron to O2 +
H2O. Figure 1(a) shows the arrangement of an
experimental setup in which Reaction 1 could
proceed.
On the left, iron is allowed to dissolve
according to:
2Fe (s) → 2Fe 2+ + 4 e −
(Eq 2)
resulting in Fe2+ ions that dissolve in the
water-based solution and electrons that are carried to the right side, where they participate in
the reaction:
O 2 + 2H 2 O + 4 e − → 4 OH −
(Eq 3)
Inside the water-based solution, ions (Fe2+,
OH⫺, H+, or any others) migrate, thereby constituting a so-called ionic current. This current
together with Reactions 2 and 3 and the transport of electrons from left to right form a closed
circuit called an electrochemical cell. The cell is
made up of four parts: the two electrodes where
the charge transfer Reactions 2 and 3 take place
12 / Stainless Steels for Design Engineers
often are described as half cells, for example,
Fe / O 2 / OH − and Fe / Fe 2+ .
Faraday’s Law
If the cell in Fig. 1(a) was allowed to proceed
and thermodynamics favored to proceed according to the direction in Reaction 1, then a
current i will flow from the anode to the cathode, and the amount of charge passed per unit
time as a result of this current will be linked to
the amount of iron dissolved per unit time or the
amount of oxygen reacted per unit time by
virtue of Eq 2 and 3. This is given by Faraday’s
law:
nNF = it
(Eq 4)
Here, i * t is the charge passed (in coulombs);
N is the moles of consumed/produced specie
(e.g., moles consumed iron in Reaction 2); n is
the ratio of electrons to consumed/produced
species, which in the case of Reaction 2 will be
2; and F is Faraday’s constant, which is essentially the charge in coulombs corresponding to
1 mole of electrons.
Fig. 1 Schematic illustration of (a) a differential aeration cell
involving iron dissolution and (b) the same cell with a
variable resistor and voltmeter
(the anode and cathode, respectively), an electrolyte, and an electron pathway. It should be
noted that electrodes are interfaces that require
several phases to be in contact. Oxidation, Reaction 2, occurs at the anode and reduction, Reaction 3, occurs at the cathode. The electrolyte
is the medium through which the ions migrate;
in the case of corrosion reactions, this is most
commonly a water-based solution, but at high
temperatures it could be a solid oxide. The final
constituent of the electrochemical cell is a pathway through which electrons can migrate from
the anode to the cathode.
As a shorthand notation, electrochemical cells
are written by separating components within a
phase by a comma and separating phases by a
slash; gaseous species are written next to their
conducting electrode. For example, the cell described in Fig. 1(a) would be recorded as
Fe / O 2 / OH − , Fe 2+ / Fe . This cell is an example
of a differential aeration corrosion cell, which is
discussed later. Processes at a single electrode
The Nernst Equation
Electrochemical reactions require a transfer
of charge; hence, there is a coupling between
chemical and electrical energy. Consider the hypothetical setup in Fig. 1(a) with the addition of
a variable resistor and a voltmeter, resulting in
the arrangement shown in Fig. 1(b). Thermodynamically, the Gibbs free energy of the cell is
that of Reaction 1:
ΔG = ΔH − T ΔS
(a ) (a )
= ΔG + RT ln
4
0
OH −
Fe 2+
aH O PO
2
2
(Eq 5)
2
where ⌬G is the Gibbs free energy, H is the enthalpy, S is the entropy, R is the gas constant,
and T is the absolute temperature. If this is negative, the reaction would be expected to proceed
spontaneously as written in Reaction 1. Let us
assume that this is the case. The thermal heat
produced by the system can be divided into two
parts: the thermal heat produced by the cell Qt
and the heat produced at the resistor QRes. QRes
Chapter 2: Corrosion Theory / 13
in this case is heat, but in essence it represents
the available energy or work, which in the case
of a resistance is given by the product of charge
passed times potential difference. If the resistance approached infinity ( R→ ), Reaction 1
would proceed through infinitesimal steps and
can be considered thermodynamically reversible.
In this case, the thermal heat produced by the
cell is minimized and according to thermodynamics is given as Qt = Qrev = T⌬S1. On the
other hand, the net work gained QRes is maximized and constitutes the rest of the free energy:
QRes = ΔG = ΔH − T ΔS
(Eq 6)
As mentioned, the energy dissipated through
the resistance is charge passed times potential
difference, and in this case the potential difference is the reversible potential difference E;
thus, in an absolute sense:
Here, pi and ri are the concentrations of reactant and products, respectively, and αi and βi are
the numbers that are needed to balance the reaction stoichiometrically. In the case of Reaction
1, Eq 10 would be:
0
Erxn = Erxn
−
(Eq 7)
Here, n is the number of electrons passed per
atom of iron reacted, and F = 96,485 C per mole
electrons, is Faraday’s constant. The reversible
potential difference E represents the potential
difference between the two electrode reactions
(cathode and anode), and as such they are associated with Reaction 1 rather than a physical
cell. The potential difference is referred to as
the electromotive force (emf) of the cell. It is
also referred to as the open circuit potential because it is the potential measured by the voltmeter in Fig. 1(b) when a negligible current
flows. It is defined here as Erxn. By convention,
this potential is positive for a spontaneous reaction (as opposed to the chemical free energy,
which is negative); hence, Eq 7 becomes:
ΔG = − nFErxn
(Eq 8)
and if all elements have unit activities:
0
ΔG 0 = − nFErxn
(Eq 9)
Equation 8 is the Nernst equation. By virtue
of Eq 8 and 9 and the expression for Gibbs free
energy of a reaction (e.g., Eq 5), an expression
for Erxn is obtained:
RT ⎛ p1 1 p2 2 p3 3 ... ⎞
ln ⎜
⎟
nF ⎝ r1β1 r2β2 r3β3 ... ⎠
α
0
Erxn = Erxn
−
α
α
(Eq 10)
)( )
(Eq 11)
If the emf according to Eq 11 is positive, this
means that the free energy is negative (according
to the Nernst equation); hence, the net reaction is
thermodynamically favored as it is written in
Reaction 1. By inspection of Eq 11, it can be
seen that it is the difference between two hypothetical half reactions, ( Erxn = EO /OH− − EFe 2+ /Fe )
2
defined as:
)
(
4
⎛
⎞
RT ⎜ aOH− ⎟
EO /OH− = EO /OH− −
ln
2
2
4 F ⎜ aH O PO ⎟
⎝ 2 2⎠
0
ΔG = nFE
(
4
2
⎛
⎞
RT ⎜ aOH− aFe 2+ ⎟
ln
⎟
4F ⎜
aH O PO
2
2
⎠
⎝
(Eq 12)
which corresponds to the reduction Reaction 3
and:
)
4
RT ⎛ ( aFe ⎞
⎟
ln ⎜
EFe 2+ /Fe = EFe 2+ /Fe −
4 F ⎜⎝ aFe 2+ ⎟⎠
0
(Eq 13)
which corresponds to the reverse of Reaction
2, that is, if it was a reduction reaction. The
potentials as written in Eq 12 and 13 are
called reduction potentials, and because
Erxn = EO /OH− − EFe 2+ /Fe has to be positive for the
2
reaction to be thermodynamically favored as
written in Eq 1, the reduction potential E O 2 /OH−
has to be larger than E Fe 2+ /Fe. If it was not, then
Reaction 1 would proceed in the reverse direction, which means that the electrode Reactions
2 and 3 would be reversed and thus so would
the anode and cathode of the cell.
It is useful to list reduction potentials for halfcell reactions, just as it is useful to list free energy data. However, half-cell potentials (like
any electrical potentials) cannot be measured in
an absolute sense; only potential differences can
be measured. ( Erxn = EO /OH – − EFe 2+ /Fe can be
2
measured because it is a difference.) Therefore,
half-cell potentials are measured with respect to
a reference electrode. Reference electrodes are
constructed such that they have a stable potential;
this is discussed further in Chapter 3. A common reference electrode in aqueous solutions is
14 / Stainless Steels for Design Engineers
the normal hydrogen electrode (NHE), also
known as the standard hydrogen electrode
(SHE), with a potential set (arbitrarily) as zero
at all temperatures. The NHE is schematically
shown in Fig. 2. In shorthand notation, it is:
Pt / H 2 (a = 1)/ H + (a = 1) , and the half-cell reaction is:
2 H + + 2 e− = H 2
(Eq 14)
Table 1 (Ref 6) lists half-cell reduction stan0
dard potentials ( EOx/Re) versus NHE that are a
result of the emf of the following types of cells
(for Reaction 2, as an example):
Pt / H 2 (a = 1)/ H + (a = 1), Fe 2+ (a = 1) / Fe
Galvanic versus Electrochemical Cells
When reactions in a cell occur spontaneously
in the direction dictated by the open-circuit potential of a cell that is positive ( Erxn > 0) , a current flows as shown in Fig. 3(a). This is the case
in environmentally caused electrochemical corrosion reactions. It also is the case in fuel cells
and batteries (under discharge), in which the
current is used as electricity. These types of
cells are called galvanic cells, in which chemical energy is converted to electrical energy.
Most of the discussion in the following chapters
concerns these types of cells. In electrolytic
cells (Fig. 3b), an imposed electrical potential
counters the “natural” cell potential to drive a
reaction in a desired direction. These types of
cells are used for many metallurgical processes,
such as electroplating, electrorefining and electroextraction (e.g., the Hall-Heroult aluminum
smelting cell), and for other applications, such
as charging batteries. In the case of corrosion,
the principle is used for protection against corrosion. In electrolytic cells, electrical energy is
converted to chemical energy.
Table 1 Standard half-cell reduction potentials
versus the normal hydrogen electrode
Reaction
Fig. 2 The normal hydrogen electrode (NHE)
Fig. 3 Schematic of (a) galvanic cell and (b) electrolysis cell
Standard half-cell reduction
potential vs. NHE(a) (V)
Fe 3+ + e− = Fe 2 +
0.771
O 2 + 2H 2 O + 4e − = 4 OH – (pH = 14)
0.401
2 H + + 2 e− = H 2
0.000
Ni 2 + + 2 e− = Ni
–0.250
Fe 2+ + 2e− = Fe
–0.447
Cr 3+ + 3e− = Cr
–0.744
2 H 2 O + 2 e− = H 2 + 2 OH − (pH = 14 )
–0.828
(a) NHE, normal hydrogen electrode. Source: Ref 6
Chapter 2: Corrosion Theory / 15
Corrosion Tendency
The tendency to corrode, that is, whether a
system consisting of anode, cathode, and electrolyte can react thermodynamically, is determined by evaluating Erxn. If this is positive, then
there is thermodynamically a possibility for corrosion. The rate of corrosion, which is in most
cases determined by corrosion kinetics, is discussed in Chapter 3. Consider, for example, a
case of iron in aerated water. Figure 1 (with
electrode Reactions 2 and 3) can be viewed as
an idealized equivalent cell for this situation. It
should be noted, however, that the locations of
anode(s) and cathode(s) on the iron surface cannot be identified with ease. At room temperature
(298 K), 1 atm oxygen partial pressure, and
using Table 1, Eq 12 can be written by assuming
unit activity for water and unit activity coefficient for OH−:
E O / OH – = 0.401 + 0.059 pOH
EFe 2+ /Fe = −0.447 − 0.0295 log(aFe 2+)
= –0.624 V (vs. NHE)
(Eq 15)
Here, the following definition of pH has been
used: pH = –log CH+, pOH = –log COH− and pH
+ pOH = 14. Similarly, the iron dissolution Reaction 2 will have a reduction potential accord-
(Eq 16)
Figure 4(a) shows a schematic plot of the two
reduction potentials (Eq 15 and 16) versus pH.
Because a spontaneous reaction requires Erxn
to be positive, if the only pertinent reactions
were Eq 2 and 3, this means that corrosion (due
to iron dissolution to Fe2+ and oxygen reduction) is possible when the line representing
EO2/OH− (Eq 15) lies above the line representing
EFe / Fe (Eq 16). This is indicated by the region
shaded in gray in Fig. 4(a). Hydrogen reduction
is another possible cathode reaction in water:
2+
2 H + + 2 e− → H 2
2
= 0.401 + 0.059(14 − pH )
= 1..227 − 0.059 pH V (vs. NHE)
ing to Eq 13, which, assuming a Fe2+ activity of
10−6 (this is an arbitrary value but is usually
taken to represent a low ion concentration), becomes at room temperature (using Table 1 for
the standard potential):
(Eq 17)
and its reduction potential is (using the definition of pH):
EH+ / H = EH0 + / H −
2
2
RT PH2
ln
nF aH2 +
= 0 − 0.059 pH V vs. NHE
Fig. 4 Reduction potential versus pH for iron and (a) oxygen gas reduction and (b) hydrogen ion reduction
(Eq 18)
16 / Stainless Steels for Design Engineers
Figure 4(b) shows the condition in which corrosion under deaerated conditions (due to iron dissolution to Fe2+ and hydrogen ion reduction) is
possible as a gray shaded region. In Fig. 4(a)
and (b), the regions where iron is stable are
denoted as immunity (corresponding to immunity from corrosion). When comparing these
two figures, it is noteworthy that hydrogen ions
are able to cause corrosion only under relatively
low pH conditions, whereas oxygen gas is able
to corrode iron in the entire pH range.
The Construction of Pourbaix Diagrams
Figures 4(a) and (b) are types of phase diagrams that show the stable phases in an area
bounded by pH and potential. In reality, several
electrochemical and chemical reactions need to
be considered when constructing these types of
diagrams. Each reaction is represented by a line.
In the case of iron, the following chemical reactions will have to be considered (the pH dependency of these reactions is listed next to
them [Ref 7] and since they are not electrochemical, they are evaluated from the equilibrium constants):
Fe 2 + + 2 H 2 O = Fe(OH )2 + 2 H + ,
(
pH = 6.65 − 0.5 log aFe 2 +
)
(Eq 19a)
Fe(OH)2 = HFeO 2− +H + ,
(
pH = 14.30 + log aHFeO−
2
)
(Eq 19b)
Fe 3+ + 3H 2 O = Fe(OH)3 +3H + ,
pH = 1.613 – (1/3) log(a Fe3+ )
(Eq 19c)
Since these are independent of potential, they
will appear as vertical lines (see lines 19a to 19c
in Fig. 5a).
The following pH-independent electrochemical reactions need to be considered, and they
will result in horizontal lines (Fig. 5a):
Fe 2+ + 2 e− → Fe,
(
EFe 2+ /Fe = −0.447 + 0.0295 log aFe 2+
)
HFeO 2− + H 2 O = Fe ( OH 3 + 2e− ,
(
)
(Eq 20a)
EFe(OH) / HFeO2– = −0.810 − 0.0591 log aHFeO–
3
2
)
(Eq 20b)
The following electrochemical reactions will
depend on pH and thus will be sloped depending on this dependence (Fig. 5a).
Fig. 5 Pourbaix diagram for iron. (a) Schematic matching Eq 19 to 21 in text to lines. (b) Actual complete diagram. Source: Ref 7
Chapter 2: Corrosion Theory / 17
)
Fe + 2H 2 O = Fe ( OH 2 + 2H + + 2e− ,
EFe(OH) / Fe = −0.0470 − 0.0591 pH
(Eq 21a)
2
Fe + 2H 2 O = HFeO −2 + 3H + + 2e− ,
EHFeO- /Fe = 0.495 − 0.0886 pH
2
(
+ 0.0295 log aHFeO–
2
)
(Eq 21b)
)
Fe 2+ + 3H 2 O = Fe ( OH 3 + 3H + + e− ,
EFe(OH) / Fe 2+ = 1.057 − 0.1773 pH
3
( )
− 0.0591 log aFe 2+
(Eq 21c)
Fe(OH)2 + H 2 O = Fe(OH)3 + H + + e− ,
EFe(OH) /Fe(OH) = 0.271 − 0.0591 pH
3
Fig. 6 Pourbaix diagram for chromium in water. Source: Ref 8
(Eq 21d)
2
For the pH-dependent reactions (chemical
and electrochemical), one can readily label the
regions depending on what iron species increasing pH favors. If iron would be an anode and
the tendency to corrode were to be evaluated,
then the reduction potential for a possible cathode reaction would be placed on this diagram. If
this point were to be, for example, in A in Fig.
5(a), this means that the reduction potential for
this assumed cathode lies below any reduction
potential of iron, and hence under these conditions iron is immune (since Erxn is negative). In
fact any Fe2+ ions present could plate as iron.
On the other hand, if the reduction potential of
the assumed cathode reaction were to lie in
point B, then there is a tendency to dissolve iron
to Fe2+ since Erxn is positive. Finally, if the reduction potential of the assumed cathode was at
point C, corrosion would occur, resulting in
Fe(OH)3, but when oxides or hydroxides are
formed there is a possibility that this product
could form a solid protective layer that kinetically hinders further corrosion. These types of
diagrams are called Pourbaix diagrams. Figure
5(b) shows the Pourbaix diagram for iron overlaid with the common cathode reactions in
water, Eq 15 and 18 (Ref 8). The ionic activity
was previously arbitrarily set at 10–6, but from
the Pourbaix diagram it can be seen that changes
in ion activity do not have dramatic effects on
the boundaries. It can be seen that both the
oxygen gas reduction reaction and hydrogen ion
reduction are able to cause corrosion through
the entire pH region. Unfortunately, Fe-OH corrosion products are generally not passivating.
Iron or carbon steel alloys are therefore not particularly corrosion resistant in water solutions.
Figure 6 shows the Pourbaix diagram for
chromium (Ref 8). While chromium oxidizes
even more readily than iron, it forms Cr2O3 over
a significantly large region that is of relevance
to pH values in water solutions. Since Cr2O3 is
protective, it prevents further corrosion. When
chromium is added to iron as an alloying element, it corrodes selectively due to its low reduction potential, but this means that it also protects the iron alloy due to the properties of
Cr2O3. This is the basic design principle behind
iron-chromium-based stainless steels.
REFERENCES
1. A.J. Bard and L.R. Faulkner, Electrochemical Methods: Fundamentals and Applications, 2nd ed., Wiley, 2001
2. D.A. Jones, Principles and Prevention of
Corrosion, 2nd ed., Prentice Hall, 1996
3. H.H. Uhlig and R.W. Revie, Corrosion and
Corrosion Control: An Introduction to Corrosion Science and Engineering, 3rd ed.,
Wiley, 1985
4. M.G. Fontana, Corrosion Engineering, 3rd
ed., McGraw-Hill, 1986
18 / Stainless Steels for Design Engineers
5. ASM Handbook, Vol 13A, Corrosion: Fundamentals, Testing, and Protection, S.D.
Cramer and B.S. Covino Jr., Ed., ASM
International, 2003
6. Handbook of Chemistry and Physics, 71st
ed., CRC Press, 1991
7. D.A. Jones, Principles and Prevention of
Corrosion, 2nd ed., Prentice Hall, 1996, p. 59
8. S.A. Bradford, Corrosion Control, 2nd ed.,
CASTI Publishing, Inc., 2001, p 41
SELECTED REFERENCE
• M. Pourbaix, Atlas of Electrochemical Equilibria in Aqueous Solutions, NACE, 1974
Stainless Steels for Design Engineers
Michael F. McGuire, p 19-25
DOI: 10.1361/ssde2008p019
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
CHAPTER 3
Corrosion Kinetics
Summary
the penetration due to dissolution of element i
becomes:
CORROSION INVOLVES chemical reactions with equilibrium that is known through
thermodynamics. In practice, the rate at which
corrosion reactions occur is the most important
consideration. This chapter deals with corrosion kinetics, which allows us to understand the
rates of corrosion.
Introduction
Consider the differential aeration cell discussed in the Chapter 2 on corrosion theory, Fe/
O2/OH–, Fe2+/Fe. If the thermodynamic conditions favor electrochemical corrosion of iron,
that is, Erxn = EO / OH – EFe / Fe is positive, then
a net corrosion current i will flow, resulting in
iron dissolution and consumption of oxygen gas
according to the net reaction, 2Fe (S) + O2 +
–
2H2O → 2Fe2+ + 4OH . The magnitude of this
current will determine the rate or iron dissolution according to Faraday’s law, which was
introduced in Chapter 2: nNF = it. Because n =
2 and F = 95,485 C per mole electrons, the
moles of dissolved iron are given as a function
of time as N = i* t/(2* 95,485). Practically, this
can be readily converted to lost mass m, which
in the case of iron loss becomes m = N* MFe =
MFe* i * t/(2* 95,485), or thickness reduction r,
which in the case of iron becomes r = MFe* i*
t/(2* 95,485* A* ␳Fe). Here, MFe and ␳Fe are
molar mass and density of iron, respectively. It
is often the thickness loss (referred to as
penetration per unit time) that is useful;
therefore, i/A is often replaced by j, which is
defined as current density and has the units
amperes/square meters. A general equation of
–
r= j
t ⋅ Mi
ni ⋅ F ⋅ ρ
(Eq 1)
The penetration rates for iron and various
stainless steels are listed in Table 1 (Ref 1) in
units of mils (0.001 in.) per year, or mpy. In the
case of alloys, the ratio of Mi /ni is computed as
an equivalent weight (EW) according to:
EW =
1
fn
∑ Mi i
i
(Eq 2)
2+
2
where fi, ni, and Mi are the weight fraction, valence, and molar mass of element i,
respectively.
As stated, the amount of corroded (dissolved) iron is determined by the current i, and
the magnitude of this current is determined by
Table 1 Penetration rates for a current of 1
µA/cm2 (mpy)
Alloy
Element/oxidation state
Density,
g/cm3
EW(a)
Penetration
rate, mpy
Fe
304
321
309
316
430
446
20Cb3
Fe/2
Fe/2,Cr/3,Ni/2
Fe/2,Cr/3,Ni/2
Fe/2,Cr/3,Ni/2
Fe/2,Cr/3,Ni/2,Mo/3
Fe/2,Cr/3
Fe/2,Cr/3
Fe/2,Cr/3,Mo/3,Cu/1
7.87
7.9
7.9
7.9
8.0
7.7
7.6
7.97
27.92
25.12
25.13
24.62
25.50
25.30
24.22
23.98
0.46
0.41
0.41
0.41
0.41
0.42
0.41
0.39
(a) Equivalent weight. Source: Ref 1
20 / Stainless Steels for Design Engineers
EFe 2+ /Fe = EFe0 2+ /Fe −
Fig. 1 Schematic illustration of a differential aeration cell
involving iron dissolution. Kinetic steps: (1) electrode
reactions, (2) ion conduction, (3) electron conduction
the corrosion potential. The corrosion potential
is determined by the reaction potential (which
was discussed in Chapter 2) and the kinetics of
the various steps involved in completing the
electrochemical circuit depicted in Fig. 1.
These involve: (a) electrode reactions at the
cathode and anode, (b) conduction of ions in
the electrolyte, and (c) conduction of electrons
from the anode to the cathode. The conduction
of electrons is generally not a problem in stainless steels because the corroding metal (iron)
and scale (Cr2O3) provide an easy path for electrons. The other two kinetic processes are discussed briefly in this chapter.
4
RT ⎛ ( aFe ) ⎞
ln ⎜
⎟
4 F ⎝ aFe 2+ ⎠
When a cell is not under open circuit (i.e., a
net current passes through it), the cathode and
anode potentials deviate from the half-cell potentials, and the electrode states are then defined
as being polarized. The polarization is quantified
as overpotentials ␩, which are defined by the deviation from the equilibrium half-cell potentials,
that is, for the cathode, ηc = Ecathode – EO2 / OH and
for the anode, ηa = Eanode – EFe 2 + Fe. Effectively, the
overpotential reduces the activation energy for
the electrode Reactions 3 and 4. In the case of a
reduction reaction at a cathode, such as Reaction 3, the overpotential is negative, and driving
an electrode toward a lower potential drives
electrons from the electrode into the solution,
resulting in a net cathodic current ic at this electrode. Similarly, at the anode the overpotential
is positive, which results in electrons that are favored to be removed from the solution and
transferred into the electrode, thus producing a
net anodic current ia.
If the magnitudes of cathode and anode polarization are large, as would be expected in a galvanic cell, the relation between each electrode
current/current density and overpotential is
given by the following equations (for a thorough derivation of the current overpotential
equation, Ref 2 is recommended):
The Butler-Volmer Equation
For the case study 2Fe (S) + O2 + 2H2O
→ 2Fe2+ + 4OH–, the cathode and anode reactions are:
O 2 + 2H 2 O + 4e − → 4 OH −
jc =
(Eq 4)
The Nernst equation predicts an open circuit
potential of Erxn = EO / OH – EFe / Fe
where
–
2+
2
(
)
4
⎛
⎞
RT ⎜ aOH – ⎟
EO /OH – = EO /OH – −
ln
2
2
4 F ⎜ aH O PO ⎟
⎝ 2 2⎠
0
and
⎛ (1 − α )nF ηa ⎞
ia
= j0 ,a exp ⎜
⎟⎠
Aa
RT
⎝
(Eq 6)
and
ja =
2 Fe ( s ) → 2 Fe 2 + + 4 e−
(Eq 5)
2
(Eq 3)
and
CO (0, t )
⎛ αnF ηc ⎞
ic
exp ⎜ −
= j 0 ,c 2
Ac
C O*
RT ⎟⎠
⎝
where the jo,i terms are the exchange current
densities and represent the equally large cathode and anode currents at equilibrium (zero
overpotential) at the electrodes. The exchange
current densities are a measure of the electrocatalytic ability of the surface to promote/demote
the electrode charge transfer reactions; as such,
they can vary over many orders of magnitude
depending on the surface chemistry and structure and on electrode reaction. The α-terms are
fractions that define the amount to which the
Chapter 3: Corrosion Kinetics / 21
activation energies are lowered. They do not
have to be the same for the anode and the cathode, but due to the uncertainty in evaluating
them, they are often taken as 0.5. The concentration terms represent the ratios between the reactant concentration at the electrode/electrolyte
interface and bulk, which could deviate from
unity as a result of consumption/production of
species at the interface. In an iron-based alloy,
this ratio for the anode would be close to unity
because the reactant is iron itself, and no
concentration gradient would be expected as a
result of the corrosion reactions. When a net
corrosion current flows, icorr = ia = ic. If the cathode and anode areas are assumed to be equal,
then jcorr = ja = jc and Eq 5 and 6 can be rewritten (using βc = 2.3RT/(αnF) and βa = 2.3RT/
((1 ⫺ α)nF) as:
ηc = β c log
j 0 ,c
jcorr
+ β c log
CO (0, t )
2
C O*
electrode, and the rate of cathode reaction will
depend on how rapidly oxygen molecules diffuse to the electrode/electrolyte interface. As a
limiting case, when the oxygen concentration is
actually zero at the interface, the corrosion current can, through Faraday’s law, be coupled to
the steady-state flux of diffusive oxygen supply
through a boundary layer δ. This limiting case
current is called the limiting current (iL or jL)
and can be expressed as:
jcorr = jl =
DO nFCO*
2
In a nonlimiting case, the corresponding
equation would be:
jcorr =
DO nF (CO* − CO (0, t ))
2
2
2
δ
(Eq 7)
2
(Eq 11)
2
δ
(Eq 12)
Combining Eq 11 and 12, one obtains:
and
CO (0, t )
j
ηa = β a log corr
j0 ,a
2
(Eq 8)
Tafel Regime: Electrode-Kinetics Control.
If the electrode charge transfer reactions are
rate limiting, the supply of oxygen to the reaction site would be rapid enough to maintain a
concentration at the electrode close to that of
the bulk. In this case, Eq 7 and 8 would both result in a linear dependence of the overpotentials
versus log jcorr:
ηc = β c log j0 ,c − β c log jcorr
CO *
= 1−
2
jcorr
jl
(Eq 13)
Thus, Eq 7 can be written:
ηc = β c log
⎛
j ⎞
+ β c log ⎜ 1 − corr ⎟
jcorr
ji ⎠
⎝
j 0 ,c
(Eq 14)
The slower the diffusion (small d), the lower
the limiting current and thus a larger contribution from the mass-transfer-dependent second
term on the overpotential.
(Eq 9)
Migration and Ionic Diffusion
and
ηa = β a log jcorr − β a log j0 ,a
(Eq 10)
The ionic transport in the electrolyte phase,
the flux of an ion i under an electric field φ
across a distance L, can be shown to be:
Mass Transfer Control. In Eq 7, the term:
CO (0, t )
2
CO*
2
stands for the ratio of oxygen gas concentration
at the electrode/electrolyte interface and the
concentration in the bulk, sufficiently far away
from the interface. If the electrode reaction kinetics are very fast, the depletion of oxygen will
lead toward a zero oxygen concentration at the
J i = − Di
∂ci zi F
Δφ
−
Dc
∂x RT i i L
(Eq 15)
In an electrolyte with many different ions, an
ion current through an area A can be computed
by multiplying Eq 15 with zi * A and summing
the contribution from all ions:
i = FA∑ zi Di +
∂Ci + ( x )
∂x
+
F2A
RT
22 / Stainless Steels for Design Engineers
× ∑ zi 2Ci Di Δφ / L
(Eq 16)
Because the first term is important only at the
regions near the electrodes (where consumption/creation of species occur), the current in
the majority region of the electrolyte can be estimated as:
2
i≅
F A
∑ zi 2Ci Di Δφ / L
RT
(Eq 17)
Using Ohm’s law (R = U/i), the electrolyte resistance can be computed as:
⎛ F2A
⎞
Relectrolyte = L / ⎜
∑ zi 2Ci Di ⎠⎟
⎝ RT
(Eq 18)
The resistivities of some test solutions are
shown in Table 2.
ηc = β c log
⎛
j ⎞
+ β c log ⎜ 1 − corr ⎟
jcorr
ji ⎠
⎝
j 0 ,c
Viewing the electrochemical cell as an electrical circuit, Kirchoff’s law can be used to
design a so-called polarization diagram. Consider, as a case study, a steel corroding under
deaerated conditions, in a water solution, as
shown in Fig. 2. Assume that the pH is such that
a passive layer does not form (see the discussion of Pourbaix diagrams in Chapter 2). The
cathode and anode reactions, respectively, are:
2 H + + 2 e− → H 2
(Eq 19)
Fe ( s ) → Fe 2 + + 2 e−
(Eq 20)
j
ηa = β a log corr
j0 ,a
(Eq 22)
Now, the potentials of anode and cathode
when current is flowing are in each case the
equilibrium potential plus overpotential, that is:
Ecathode = EH0 + / H −
2
RT PH2
ln
+ ηc
nF aH2 +
= 0 − 0.059pH + β c log
j0 ,c
jcorr
(Eq 23)
( )
Eanode = −0.447 − 0.0295 log aFe 2+ + ηa
j
= −0.624 + β a log corr
j0 ,a
Natural Seawater
Fresh (tap) water adjusted
with seawater
Fresh (tap) water adjusted
with seawater
Fresh (tap) water adjusted
with seawater
Deionized water adjusted
with fresh (tap) water
Ratio by
volume
Resistivity,
ohm-cm
...
28:1
25
500
68:l
1,000
950:1
3,000
21:10
10,000
(Eq 24)
A polarization diagram is now constructed by
plotting the anode and cathode potentials versus
log jcorr. Strictly speaking, to close the circuit,
Test solution resistivity
Test solution
(Eq 21)
and
⎛
j ⎞
+ β c log ⎜ 1 − corr ⎟
ji ⎠
⎝
Mixed Potential Theory and
Polarization Diagrams
Table 2
and the respective equations describing the
overpotentials will be:
Fig. 2 Schematic polarization diagram
Chapter 3: Corrosion Kinetics / 23
the potential drop across the electrolyte needs to
be included, which simply equals icorr * Relectrolyte
(the electrolyte resistance is evaluated from Eq
18); however, in many cases, this term can be
neglected. A schematic polarization diagram is
shown in Fig. 3. The anode polarization is linear
with decade current as predicted by Eq 24 because the overpotential has only a Tafel regime
and no mass transfer dependence. On the other
hand, the cathode polarization deviates from the
Tafel behavior as a result of the effect of the
mass transfer (hydrogen ion supply), dependent
on the limiting current in Eq 23. It is noteworthy that the cell shown in Fig. 1 does not have a
macroscopic anode and cathode. Different microscopic regions on the surface are assumed to
act as cathodes and anodes, and in the lack of
more detailed knowledge, the cathode and
anode areas are assumed to be equal. The overall mixed potential of the surface would be at a
corrosion potential Ecorr, defined in Fig. 2.
In effect, the corrosion current resulting from
the cell depends on the equilibrium half-cell potentials (Ecathode and Eanode), the Tafel slopes (βc
and βa), the exchange current densities (jo,c and
jo,a), and any limiting current density (jl). Figure
3 shows schematically how decreasing any of the
Tafel slopes and increasing an exchange current
density increases the corrosion rate. The effect of
the electrolyte resistance has been ignored; that
is, corrosion current is where the two polarization curves intersect. Figure 4 shows the effect of
increased mass transfer, which would result in an
increase in the limiting current. In an active (nonpassive) alloy, this results in an increased corrosion current up to a point.
Passivation
Theory. In Chapter 2, it was identified
through the Pourbaix diagrams that there were
conditions under which an alloy could be
passive. In the case of stainless steels, the range
of pH and other conditions under which this
would occur has been increased thanks to the
chromium content, which readily forms a Cr2O3
scale. In general, a passive layer constituted of
adsorbed molecules or thin oxide/hydroxide
layers decreases the corrosion current. Researchers (Ref 3) have reported that the constituents of the passive film are alpha Cr2O3 and
Cr(OH)3nH2O. The structure is reported to be a
nanocrystalline spinel, epitaxial to the surface.
The grain size may decrease with increasing
chromium content. This protection by
chromium requires a threshold level of 11 to
12% chromium.
Effect on Polarization Diagrams. The polarization diagram for a passive alloy is quite different from those discussed for active alloys. A
schematic of a typical polarization curve is
shown in Fig. 5. When a passive alloy is anodically polarized, it initially behaves like an active
alloy (i.e., with a Tafel slope, etc., as the passive layer is building up). The building up is
actually a selective dissolution of iron, which
causes a greater remaining surface concentration of chromium and other alloying elements.
Once the passive layer is formed and offers
protection against further dissolution, the potential-decade current relation drops to lower
currents. This happens at potentials beyond the
passivation potential Epp. At some high enough
Fig. 3 Corrosion rate and the effect of (a) Tafel slope and (b) exchange current density
24 / Stainless Steels for Design Engineers
polarization level, the passive layer breaks
down, and the metal becomes active again; this
region is called the transpassive regime. The
design of a structure involving a passive metal
should aim at forming a corrosion cell in which
the cathode polarization curve intersects the anodic one in the passive regime.
Consider, for example, an alloy that exhibits
the behavior shown in Fig. 6 in deaerated
acidic solutions with different pH. If mass
transfer limitations due to hydrogen ion supply
are neglected, then the cathode polarization is
given by:
Ecathode = EH0 + / H −
2
RT PH2
ln
+ ηc
nF aH2 +
= 0 − 0.059pH + β c log
j0 ,c
jcorr
This will result in a straight line as shown in
Fig. 6, which will be shifted vertically depending on the pH. The dashed circles indicate the
intersection between anode and cathode polarization curves that would yield the corrosion
current. At a sufficiently high pH (= pH1), the
alloy is clearly not optimal because intersection
occurs in the active regime, and the passive
properties are not utilized. This is what occurs
when a reducing acid is too strong for a given
stainless steel, such as with concentrated hydrochloric acid. At pH = pH2, on the other
hand, a low-corrosion current is obtained as a
result of intersection at the passive regime.
This is the benevolent case when stainless steel
is correctly matched to the environment, and
low rates of uniform corrosion occur. Finally, at
pH = pH3, the resulting corrosion current is
again high as a result of intersection occurring
at the transpassive region. This could occur
with some stainless steels exposed to a very
strong alkali solution. Using a similar argument, the readers can themselves deduce the
effects of cathode exchange current density and
Tafel slopes. In the discussion of active anode
polarization, it was found that increasing the
transport rate of cathode reactants through
agitation, for example, would increase the corrosion rate up to a point but beyond that have
no further effect (see Fig. 4). In the case of a
passive/active behavior, the effect of mass
Fig. 4 Effect of increasing the limiting current by, for example,
increased agitation in the electrolyte. Beyond the
dashed line, increasing the limiting current would have no further effect
Fig. 5 Schematic of a passive anode polarization curve
(Eq 25)
Fig. 6 Effect of cathode polarization
Chapter 3: Corrosion Kinetics / 25
Fig. 7 Effect of mass transport
transport is somewhat different, as shown
schematically in Fig. 7.
Increasing mass transport, such that the limiting current increases, results initially in an increased corrosion current (e.g., increasing jl
from 1 to 2). It should be noted that there are
several intercepts possible (both in the active
and passive regime), but assuming there are
defects present, it is likely that there will be corrosion corresponding to the higher current. Increasing the limiting current beyond the knee
corresponding to Epp, however, results in a drop
in the current because now the only corrosion
potential possible is at the intersection in the
passive regime. This is the case for jL3.
In the normal use of stainless steel, achieving
passivity takes on several forms. What is often
called passivation is actually a cleaning process
in which contaminants, such as tramp iron, are
removed from the surface. Dilute nitric acid is
an excellent vehicle to achieve this. This
medium has the additional benefit of forming a
passive film on an active stainless surface. This
is the actual passivation; the iron removal is
really a chemical cleaning operation, which happens to be called passivation. During the production of stainless steel, after a final anneal another
version of passivation is carried out. The oxide
from annealing in air is dissolved by a strong
mixture of nitric and hydrofluoric acids, which
does not allow passivation. This treatment,
called pickling, removes by dissolution both the
oxide layer and the chromium-depleted layer
below the oxide formed during annealing. The
depleted layer can extend a number of microns
in depth and would seriously degrade corrosion
resistance if not removed (Ref 4). This is then
followed by a straight nitric acid immersion,
which ensures complete passivity. This is the
procedure that should be performed on the oxides formed during welding if full corrosion resistance is to be restored. Simply removing the
oxide through mechanical means leaves a
chromium-depleted layer that corrodes more
readily than is expected of the alloy.
REFERENCES
1. D.A. Jones, Principles and Prevention of
Corrosion, 2nd ed., Prentice Hall, 1996
2. A.J. Bard and L.R. Faulkner, Electrochemical Methods: Fundamentals and Applications, 2nd ed., Wiley, 2001
3. M.P. Ryan et al., Critical Factors in Focalized Corrosion, Proc. Electrochem. Soc.,
Vol 150, 2003, p 583–594
4. J. Grubb and J. Maurer, “Corrosion of the
Microstructure of a 6% Molybdenum Stainless Steel with Performance in a Highly Aggressive Test Medium,” paper 300 presented
at
Corrosion
95,
NACE
International, 1995
Stainless Steels for Design Engineers
Michael F. McGuire, p 27-56
DOI: 10.1361/ssde2008p027
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
CHAPTER 4
Corrosion Types
Summary
STAINLESS STEEL is unusual among alloy
systems in that its corrosion resistance derives
from the passivating ability of a minor constituent, chromium. Thus, while stainless steels
can be made to be essentially immune to corrosion in many environments, it can also experience various debilitating forms of localized corrosion, which stem from the failure of this
passive film. This chapter explores the behavior
of stainless steel in media that promote uniform
corrosion and the various mechanisms of localized corrosion, such as pitting and crevice corrosion.
plex in their behavior because the influence of
processing and alloying variables changes the
ability of this layer to form and remain stable in
the face of aggressive environments. The behavior of stainless steel is further affected by its
microstructural complexity. Stainless steel alloys may have many constituent elements and
many thermodynamically possible phases, and
none of these are necessarily uniform in their
composition. Yet, it is the composition of the
alloy in contact with the specific environment at
any microscopic point that determines the corrosion resistance of that particular point.
Uniform Corrosion
Introduction
To most designers, the most recognized characteristic of stainless steel is corrosion resistance. Stainless, unlike noble metals such as
gold, does not obtain its excellent corrosion resistance from inertness. Instead, it is the reactivity of chromium that allows the surface layer of
corrosion product to become sufficiently adherent and impenetrable, which effectively stops
further corrosion by isolating the base material
from the environment. This resistance to corrosion is called passive behavior or passivity.
Other metals, such as aluminum and titanium,
form similar layers and also exhibit passivity.
The important difference in the case of stainless
steel is that chromium is still a minor constituent, never more than 30% by weight, sometimes little more than 10%. How much
chromium there is and how uniformly it is distributed have a profound effect on corrosion resistance by virtue of its ability to concentrate
into the surface film. Stainless steels are com-
When all parts of a corroding surface have
equal access to the corroding atmosphere and
the structure of the corroding metal is relatively
uniform, a uniform thinning of the material is
expected. Stainless steels are materials of choice
because, by virtue of their passive behavior, they
show very low rates of uniform corrosion in
many environments. The metallurgy and processing of a particular grade are designed to provide passivity in a given environment. The environment can be too aggressive to allow passivity
to be maintained either by being too reducing, as
with some acid media, so that passivating
species cannot form or by being too oxidizing so
that the oxidized species that normally affect
passivity are no longer stable. The former is
called dissolution in the active state, while the
latter is termed transpassive dissolution.
Intelligent design and knowledge of the environmental variables for a stainless steel component ensure that the alloy is used in the passive
state, at which uniform corrosion occurs at a
very low rate.
28 / Stainless Steels for Design Engineers
Among the important media with which we
encounter uniform, but acceptably controlled,
corrosion in stainless steel are atmospheric and
marine environments and chemical environments such as sulfuric acid, phosphoric acid, nitric acid, strong bases, and organic acids, such
as acetic and formic.
Pickling is an example of controlled, accelerated uniform corrosion. This is typically done
with 10 to 20% hot sulfuric acid or a mixture of
hydrofluoric and nitric acids.
Environmental Variables Influencing
Uniform Corrosion
The corrosion of stainless steels is usually the
result of contact with an electrolyte, allowing a
complex set of partial electrochemical reactions, which may occur sequentially or concurrently. The corrosion rate depends on the
current exchanged between the negative and
positive electrode (anode and cathode). These
may be on a macroscopic or microscopic level.
The main consideration is normally ionic transpor tthrough the passive film, which after all
is what makes stainless so effective against
corrosion.
The chemical parameters that influence the
media with respect to uniform corrosion rate are
the acidity and the oxidation-reduction (redox)
potential of the electrolytic medium, both of
which act through their influence on the stability of the passive film, rendering it active, pas-
sive, or transpassive; Fig. 1 illustrates the effect
of redox potential on a solution.
Certain anions have strong effects in media
through their well-known, if not well understood, disruption of the passive film. Halides are
well known for this effect, but sulfides are
active. These anions seem to intervene in the
adsorption of the hydroxyl ions. In acid media,
these anions accelerate uniform corrosion,
while in neutral media they may result in localized corrosion. Anions that form soluble complexes with elements in stainless, such as
amines, formates, or acetates, can also disrupt
the stability of the passive film and thus promote active corrosion.
Of the physical variables of the environment,
it should be obvious that temperature is paramount since all the reactions are thermally activated. Increasing temperature may speed the
formation of the passive film when thermodynamic conditions are favorable, but in general
one expects increasing temperature to increase
corrosive attack. Access to passivating species,
such as oxygen, is important in establishing and
maintaining passivity.
Increased diffusion of reacting species in the
liquid will normally accelerate the partial reactions, but if passivity is stable, the rate-limiting
transport through the passive film will not be affected. Therefore, increasing the flow rate of a
corrosive fluid does not automatically accelerate corrosion. The reduction of concentration
gradients can be beneficial against localized
Fig. 1 Reduction potential versus pH for iron and (a) oxygen gas reduction and (b) hydrogen ionreduction
Chapter 4: Corrosion Types / 29
corrosion, and flow can bring to the surface an
increased supply of passivating species.
Increased flow rate in a fluid medium is deleterious if it induces mechanical damage to the
passive film by erosion, abrasion, or cavitation.
These are complex mechanisms, but it should
be apparent that the success of a stainless steel
to a given flow condition will depend mainly on
its ability to form and re-form its passive film.
Somewhat counterintuitively, thinner passive
films are more protective than thicker films
among stainless alloys. The tenacity of the thin
passive films on stainless (and titanium) make
these alloys quite resistant to flow-accelerated
corrosion, as contrasted to copper and aluminum alloys, which have soft, thick corrosion
product films.
Material Variables
Stainless steels have a great variety of alloying elements and microstructures. As a generalization, we can say that corrosion resistance is a
function of composition rather than structure.
Then, we must quickly add the qualifiers to this
statement. On an undisrupted, stress-free surface, local composition does quite precisely determine corrosion resistance. But, stainless
steels are seldom homogeneous or at thermodynamic equilibrium. Impurities such as oxygen
and sulfur are usually present, mainly as inclusions since they have diminishingly small solubility at room temperature. At high temperatures after solidification, as in welds, they can
be present in supersaturation, ready to precipitate as inclusions that alter local composition.
The tendency of carbon and nitrogen to form
precipitates is controlled by diffusion rates,
which if elevated by increasing temperature can
cause debilitating, composition-altering precipitation. The even more slowly diffusing substitutional alloying elements, such as chromium,
molybdenum, and nickel, have strong tendencies to form phases that disturb their uniformity
in the austenite or ferrite matrix in which they
are intended to work. So, any discussion of the
influence of alloying element on corrosion resistance of a phase like austenite or ferrite must
recognize that alloying elements exert their effect when they are in solution in that phase. The
same element may under some conditions not
be in solution and have a contrary effect. An example is molybdenum, which is obviously a
great enhancer of corrosion resistance when in
solid solution. When it precipitates as a con-
stituent of sigma phase, which it promotes,
however, it combines with chromium. If this
happens at relatively low temperatures, the surrounding matrix is depleted of both chromium
and molybdenum, and the corrosion resistance
in that region is diminished. Nitrogen also is effective when in solid solution in austenite but
can precipitate as a chromium nitride under certain conditions and cause depletion of the remaining matrix. Local structure and composition are paramount. This becomes more
important to localized corrosion, as discussed
later, but it should be remembered in examining
uniform corrosion because corrosion will cease
to be uniform when composition becomes
nonuniform.
The compositional material variables that influence uniform corrosion are not exactly the
same as those that will be seen to influence localized corrosion. The foremost element is, of
course, chromium. Researchers (Ref 1) have reported that the constituents of the passive film
are alpha Cr2O3 and Cr(OH)3nH2O. The structure is reported to be a nanocrystalline and epitaxial to the surface. The grain size may decrease with increasing chromium content. This
protection by chromium requires a threshold
level of 11 to 12% chromium. This threshold
has been attributed most convincingly to the
minimum chromium content that permits
chromium atoms on surface sites to be linked by
adsorbed oxygen atoms (Ref 2). In any event,
the mechanism by which this thin, severalnanometer-thick, film forms is the subject of ongoing debate, but we do know that it is enriched
in chromium, and that it is thinner for higher
chromium alloys. The critical current density j,
as measured during polarization, is also smaller
as chromium content increases (Fig. 2). This is
consistent with the lower dissolution of noncontributing elements required to achieve a critical
surface chromium concentration. Increases in
chromium can also be seen to lower the current
density in the passive region. This is manifest in
alloy performance as a reduction in the uniform
corrosion rate in a given medium. From an electrochemical point of view, this is explained as a
manifestation of the stability of the Cr(OH)3
nH2O.
The role of molybdenum is less clear. The observed action of molybdenum is to greatly reduce the critical current density required for
passivation. This is also seen as accelerating the
formation of the passive films and as increasing
the resistance of the alloy to depassivation at
30 / Stainless Steels for Design Engineers
Fig. 3 Influence of alloying element on corrosion rate as explained by the effect on polarization.Source: Ref 6
Fig. 2 Schematic illustration of polarization behavior for a
passive alloy with and without pitting occurring
lower pH. The role of molybdenum is not to enrich in the passive film itself, although it can be
found in the film. Its potency is far more than its
presence can take into account. Pure molybdenum is itself not passive. Its action does not appear to be via a product of reaction. Instead, it
seems to reduce the dissolution rate of elements
other than iron, which would promote a surface
richer in chromium (Ref 3). The action of
molybdenum as an alloying element is complicated by the fact that molybdate ions are known
to impede pit growth as a separate effect from
their action within the alloy matrix (Ref 4).
Copper has a similarly complicated effect, with
copper ions gettering sulfide ions and redepositing as metallic copper (Ref 5).
Nickel also lowers the critical current density
for passivation without contributing directly to
the passive film’s stability. This also may be the
result of the stronger bond between nickel and
chromium reducing the anodic dissolution rate
of the alloy by permitting the anodic enriching
of the surface by selection iron dissolution.
Nickel does not actively help passive film formation and can actually hinder film stability in
highly acidic/oxidizing environments.
Nitrogen, however, appears to be more like
molybdenum in its effect. While nickel and copper provide no benefit to the stability of the passive film once it is formed, both nitrogen and
molybdenum do, and to a degree that cannot be
explained by their presence in the film. This
may then relate to their thermodynamic action
within the alloy itself. Molybdenum and nitrogen act both to enhance the enrichment of
Fig. 4 Influence of alloying elements on uniform corrosion
rate in 20% sodium chloride solution with carbon
dioxide pressure of 20 MPa. Source: Ref 7
chromium in the passive layer and to decrease
active dissolution of noniron alloying elements,
thereby promoting both the formation and stability of the passive film. A summary of the
known major alloying effects in acidic chloride
media is shown in Fig. 3 in acidic chlorides. Alloying elements provide benefits in the part of
the chart where they appear (Ref 6). From this,
it can be seen that chromium, molybdenum,
nickel, copper, and nitrogen all assist in the active region, while chromium, molybdenum, and
nitrogen expand the region of passivity and diminish the corrosion current.
An example of the influence of these alloying
elements on the uniform corrosion rate of stainless steels in a sodium chloride/carbon dioxide
environment is shown in Fig. 4 (Ref 7). Note
the alloying composition is measured by a
crevice corrosion index (CCI), which is discussed in the section Localized Corrosion.
Chapter 4: Corrosion Types / 31
Fig. 5 Corrosion table for stainless steels and titanium in sulfuric acid plus copper sulfate. Corrosion rate legend: 0, < 0.1 mm/yr
(corrosion resistant); 1, 0.1–1 mm/yr (useful in certain circumstances); 2, > 1.0 mm/yr (material not recommended). Source:
Ref 8; see source for interpretation of data. Courtesy of Outukumpu Stainless
Unfortunately, it cannot be assumed that this
relationship is true for other environments, although other empirical relationships exist or can
be generated. Because the influence of alloying
element varies with environment, we need to
discuss some of the more commonly encountered severe environments.
Corrosion in Acids and Bases
The examples discussed in Chapter 3, “Corrosion Kinetics,” refer mostly to corrosion in
these aqueous solutions, in which the slow thinning rate of the chosen alloy can be determined
through the mixed potential theory and polarization diagrams. In the case of stainless steels,
the alloy chemistry is chosen such that the
passive-active behavior favors corrosion in the
passive regime. The corrosion rate of the various stainless steels in the myriad possible environments has been measured in probably all
practical cases. These data can be obtained from
a number of sources, such as the National Association of Corrosion Engineers (NACE) and
ASM Handbook volumes. None is more accessible than the Web site of Outukumpu, which
contains a “Steel Professional Tool,” a lookup
table in which the corrosion rate of many stain-
less steels in a great number of environments
can be obtained. Figure 5 shows an example of
one such table. Many of the isocorrosion charts
in this book are reprinted from this source,
http://www.outokumpu.com/applications/documents/start.asp (Ref 8).
These tables are supplemented by isocorrosion diagrams such as that shown in Fig. 6.
These diagrams show constant corrosion behavior under varying environmental conditions such
as temperature and solution composition. This
information is available to guide the designer in
selecting appropriate steels for various environments, and it is highly recommended that it be
used. Free sites tend to promote proprietary alloys, as these charts suggest. The serious engineer will consult multiple sources and unbiased
sources before making alloy decisions.
The influence of alloying element is by no
means the same in all environments. So, while it
is useful and necessary to have these experimental data, it is also helpful to understand the
peculiarities of some of the major alloy-environment pairings.
Sulfuric Acid. Stainless steels require more
than a minimum amount of alloying to resist sulfuric acid. Straight 16% chromium grades such
as 430 fare poorly, while the nickel-containing
32 / Stainless Steels for Design Engineers
Fig. 6 Isocorrosion curves for 17-12-2.5 stainless steel and titanium in sulfuric acid plus copper sulfate.Source: Ref 8. Courtesy of
Outukumpu Stainless
304 has more than an order of magnitude better
corrosion rate in either dilute or concentrated
sulfuric acid at ambient temperatures. Figure 7
(Ref 9) shows the isocorrosion rate curves for
several common alloys. Alloying with molybdenum is also very effective, as is alloying with
copper. If passivity cannot be established, increasing chromium content actually increases
corrosion rate.
The corrosion behavior of sulfuric acid varies
greatly with concentration. At low concentrations, sulfuric is a classic reducing acid. It dissociates in water to create hydrated hydrogen
ions (H3O+) that release hydrogen gas bubbles
during the corrosion reaction. As the acid concentration increases, the solutions become more
corrosive, and progressively more highly alloyed stainless steels are required to provide adequate corrosion resistance. At about 50% acid,
only the most highly alloyed stainless alloys
(alloy 20, AL-6XN, C-276, etc.) can provide acceptable corrosion rates, and even these alloys
are restricted to use at near ambient temperatures. As acid concentration increases beyond
50%, the solution begins to show oxidizing behavior. At acid concentrations above 80%,
nickel-molybdenum-copper-bearing stainless
steels begin to exhibit useful corrosion resistance. In the 93 to 98% sulfuric acid concentration range, carbon steel can be used to hold sulfuric acid at ambient temperatures, although
stainless steels provide better performance at elevated temperatures or if flow-erosion can occur.
In the 96 to 100% sulfuric acid concentration
range, at elevated temperatures, the oxidizing
character is quite pronounced, and oxidation-
Fig. 7 Isocorrosion rates of various stainless steels in sulfuric
acid. Source: Ref 9
resistant high-chromium (type 310S) and highsilicon (MECS ZeCor UNS S38815 and Sandvik SX S32615) stainless steels are frequently
used, especially in sulfuric acid-manufacturing
Chapter 4: Corrosion Types / 33
Fig. 8 Influence of alloying element on corrosion rate in contaminated sulfuric acid. Source: Ref 11
Fig. 9 Isocorrosion curves for various alloys in sulfuric acid
equipment. Sulfuric acid-containing dissolved
sulfur trioxide is called oleum, and such solutions are often identified as sulfuric acid of
greater than 100% concentration. Highchromium stainless steels (i.e., type 310S) are
among the very few materials that exhibit corrosion resistance in oleum. (See MTI Materials
Selector Volume 3—Sulfuric Acid at www.mtiglobal.org for more information.)
Aeration has a major influence on corrosion
rates because oxygen stabilizes the passive film.
Molybdenum-alloyed stainless has dramatically
lower corrosion rates in aerated solutions than
non-molybdenum-bearing alloys. Their superiority in deaerated solutions is much less
marked. Studies (Ref 10) have shown that in
sulfuric acid molybdenum is highly enriched in
the passive film, and when molybdenum is an
alloy, chromium also is enriched. This is a manifestation of selective dissolution of other elements in the matrix.
Oxidizing impurities, such as ferrous ions, act
like aeration to diminish the corrosive attack,
but reducing impurities such as halides have an
extremely negative effect, as the corrosion tables will show. These effects are not linear and
underscore the value of these tables.
The uniform corrosion rate in contaminated
sulfuric acid may be more important than in
pure acid since this represents a potentially
likely failure mode because contamination is a
constant hazard.
Figure 8 shows the corrosion rate of various
alloys in sulfuric acid contaminated with chlorides and iron. These researchers (Ref 11) found
that the resistance to attack correlated to the
alloy content by the formula shown.
Figures 9 and 10 show how isocorrosion rates
vary with alloy and contamination level.
Hydrochloric acid is very destructive of the
passive film on stainless. An alloy like 304 is
not suitable even in a deaerated 1% HCl solution at room temperature. Chromium additions
34 / Stainless Steels for Design Engineers
Fig. 12 Isocorrosion curves for austenitic AL-6XN (UNS
Fig. 10 Isocorrosion curves for various alloys in sulfuric acid
N08367) and 904L (UNS N08394) stainless steels
in hydrochloric acid. Source: Ref 12
with chlorides
Fig. 11 Isocorrosion curves for various stainless steels in
hydrochloric acid. Source: Ref 8. Courtesy of Outokumpu Stainless
are only modestly helpful, while nickel, copper,
and molybdenum are more beneficial. Stainless
steels are not good materials for contact with
hydrochloric acid. Figures 11 (Ref 8) and 12
(Ref 12) show how even the most highly alloyed grades can withstand only dilute concentrations and low temperatures.
While a stainless steel vessel may not be intended to be used for hydrochloric acid, resistance to lesser amounts of chlorides is important
because of the possibility that an acidic environment may be contaminated with chlorides. When
this is a possibility, then proper alloy selection
must guard against it. Figure 13 shows a correlation between alloy content and resistance to the
general corrosion (GI) by sulfuric acid contaminated with hydrochloric acid (Ref 11).
Nitric acid is strongly oxidizing. This actually promotes the passive film formation; consequently, even low-chromium alloys remain
passive at all concentrations at ambient temperature (see Fig. 14) (Ref 8).
The addition of molybdenum, which is so
generally helpful, is deleterious in this case because it forms soluble compounds. It is useful
to keep carbon, silicon, and phosphorus as low
as possible.
Silicon is unusual in that normal levels (0.4 to
1.0 %) are worst, with very low (0.05%) or very
high levels (4.0%) beneficial. The low levels of
silicon contents of these alloys are useful for
their action in minimizing grain boundary segregation, which is the usual locus of attack.
High silicon levels contribute to a general protective silica surface layer in concentrated acid,
which augments the true passive layer. This superiority appears above the azeotropic composition of about 67%, which is a common commercial concentration, as shown in Fig. 15 for
high-silicon austenitic stainless steels (Ref 13).
Phosphoric Acid. This oxidizing acid behaves more like sulfuric acid in that simple
iron-chromium alloys have only moderate resistance to uniform corrosion in them, while alloying with molybdenum and copper produces
major improvements. This can be seen in Fig. 16,
in which alloys with increasing nickel (18-10)
show clear benefits over a chromium-molybdenum alloy (18-2), and the added molybdenum
in 317 (17-14-4) is better, while 904L with
nickel, molybdenum, and copper is even better
(Ref 8).
In the commercial production of phosphoric
acid, halide impurities may be present, in which
Chapter 4: Corrosion Types / 35
Fig. 13 Influence of alloy content on corrosion rate in hydrochloric acid
Fig. 14 Isocorrosion curve for nitric acid. Courtesy of
Outokumpu Stainless
Fig. 15 Corrosion behavior of high-silicon alloys in concentrated nitric acid. Courtesy of Outokumpu Stainless
case alloys with higher molybdenum, chromium,
copper, and nitrogen may be required.
Organic Acids. The weakly dissociating organic acids are normally not aggressive against
stainless steels. The exceptional dangerous environments are those that include high temperature and the presence of chloride contamination.
It should be noted that in formic acid, which
does dissociate more strongly, nickel is detrimental. This phenomenon is also seen in the
production of urea via the intermediary ammonium carbamate. The difficulty lies in the hightemperature solubility of nickel complexes and
is best addressed by the use of ferritic or duplex
alloys.
Alloying with molybdenum seems to provide
the greatest resistance to uniform corrosion in
strong organic acids, as illustrated in Fig. 17
(Ref 8). If halides are present in organic acids
and liberated by contact with water, then pH
and chloride concentration will govern the corrosive attack, which could then become nonuniform.
Strong Bases. In strong bases, the stainless
steels are generally quite resistant to uniform
corrosion. Straight chromium (17%) alloys are
usable at any concentration up to 50 °C. Adding
molybdenum and nickel does little to further
improve performance as the underlying resistance is due to chromium. Increasing chromium
levels provide increased resistance. Attack
when it does occur can be manifested as grain
boundary attack. Figure 18 shows isocorrosion
curves for sodium hydroxide (Ref 8).
36 / Stainless Steels for Design Engineers
Fig. 16 Isocorrosion curves in phosphoric acid: (a) 0.1 mm/yr for various stainless steels; (b) 0.1 mm/yr for titanium and 17-12-2.5
stainless steel. Courtesy of Outokumpu Stainless
Fig. 17 Isocorrosion curves in organic acids: (a) acetic acid; (b) formic acid. Source: Ref 8. Courtesy of Outokumpu Stainless
In the pulp-and-paper industry, chemical
pulping is called the kraft or sulfate process. In
the presence of sulfur, nickel can be quite detrimental, and ferritic or duplex alloys are
preferred. This again is caused by the solubility
of nickel complexes formed in the presence of
sulfur-containing compounds. This can be seen
in Fig. 19, which shows a 26-1 (chromiummolybdenum) alloy significantly outperforming
higher alloys that contain nickel and molybdenum (Ref 12).
Atmospheric Corrosion
Fig. 18 Isocorrosion curves for various materials in sodium
hydroxide. SCC, stress corrosion cracking. Courtesy
of Outokumpu Stainless
Atmospheric corrosion is an example of uniform corrosion that occurs when a thin layer of
water condenses on a metal surface and as such
Chapter 4: Corrosion Types / 37
Fig. 19 Corrosion rates of various alloys in simulated evaporator liquid. Source: Ref 12
depends on humidity, temperature, and other atmospheric conditions. The rate of corrosion
measured as defined in Chapter 3, “Corrosion
Kinetics,” as dissolution r of element i as
r= j
t ⋅ Mi
ni ⋅ F ⋅ρ
where j is current density, t is time in seconds,
M is molar mass, n is valence, F is the faraday
constant, and ρ is the density), and therefore has
two contradicting effects of temperature. In
general, temperature increases the exchange
current density and transport properties and thus
the kinetic rates involved in corrosion. On the
other hand, increasing temperatures may reduce the concentration of dissolved oxygen in
the electrolyte and eventually will dry the surface and thus limit the electrochemical corrosion due to the access to an electrolyte. In steels,
the corrosion products are (a) an outermost
layer of porous rust (FeOOH) characterized by
low water content but easy access to oxygen
and (b) an inner layer of magnetite (Fe2O3) in
which pores are filled with water. The access of
oxygen to the bare metal limits the cathode reduction reaction rate, and in relatively pure atmospheres, the corrosion rate is negligible due
the protective nature of the oxide. However, sulfur dioxide impurities in the atmosphere react
with water to form sulfuric acid, which tends to
dissolve the protective oxide. In the case of
stainless steels, the passive region is extended
due to chromium, to a wide enough region in
terms of pH (see Pourbaix diagrams in Chapter
2) that atmospheric corrosion can in effect be
prevented.
The most deleterious impurity in the atmosphere for stainless is the chloride ion. Chlorides
are pervasive. Borne from oceans by normal climatological processes, they are found far inland.
In many colder climates, they are also seen in
high concentrations from road salts. Without
washing or the natural rinsing by rain, surface
chloride concentrations can become very high.
Thus, the rules of corrosion of aerated aqueous
solutions are followed by stainless with respect
to atmospheric corrosion. The difficulty is accurately estimating the solution that constitutes the
aqueous solution. Much experience has shown
that if coastal and road salt effects are minimal,
then 18% chromium alloys such as 304 experience such negligible visible corrosion that they
can be used for exposed, unrinsed architectural
purposes. If the same alloy is used in an unrinsed
coastal environment, red rust stain will occur.
This is the corrosion product from metastable
and possibly stable pitting. In Japan, where
coastal conditions prevail throughout, much research has been done that has shown that a pitting resistance equivalent number (PREN) of 25
is necessary for freedom from corrosion (i.e.,
zero pitting) (Ref 14). This contrasts to a requirement of about PREN 35 to resist pitting in
seawater. Pitting is a form of localized corrosion, and PREN is an index to pitting resistance.
These concepts are examined in the next section.
Localized Corrosion
Localized corrosion is in general more
damaging from a structural integrity point of
view than uniform corrosion since the corrosion current is limited to a small area and the
penetration distance is large. Often, localized
corrosion involves a large-area cathode and a
small-area anode, which means that for a given
corrosion current, the corrosion current density
at the anode is very large. In localized corrosion, unlike uniform corrosion, the anode and
cathode are clearly identifiable locations, and
the reason that certain structural features assume the roles of cathode and anode can be
used to categorize and exemplify different
cases. Interestingly, the cathode and anodes,
while identifiable, can vary across scales, that
is, from distinct macroscopic components or
parts to microstructural features.
In Chapter 2, the tendency for corrosion
was introduced as a positive value for an
38 / Stainless Steels for Design Engineers
electrochemical cell potential (Erxn) corresponding to a spontaneous electrochemical reaction
forming a galvanic cell. Erxn is obtained as:
0
Erxn = Erxn
−
α
α
α
RT ⎛ p1 1 p2 2 p3 3 ... ⎞
ln ⎜ β β β
⎟
nF ⎝ r1 1 r2 2 r3 3 ... ⎠
(Eq 1)
Here, pi and ri are the concentrations of reactant and products, respectively, and αi and βi are
the numbers that are needed to balance the reaction stoichiometrically. Any time that Erxn is
positive, there is thermodynamically a tendency
for an electrochemical reaction, in our case a
corrosion reaction. The rate of corrosion, as discussed in Chapter 3, is dependent on the polarization behavior.
Dissimilar Metals and Differential
Aeration Cells
The case of dissimilar metals and differential
aeration cells is perhaps more important in active alloys than for stainless steels, which are
generally passive, and occurs when two
metals/alloys are in contact that have elements
in them that are dissimilar in the electromotive
force (emf) series (see Chapter 2) and there is an
electrolyte present. For example, if nickel and
iron pipes are connected and water flows though
them containing some traces of Ni2+ ions, then:
0
Erxn = ENi
− EFe0 2 +/ Fe −
2+
/ Ni
RT ⎛ aFe 2 + ⎞
ln ⎜
⎟
2 F ⎝ aNi 2 + ⎠
(Eq 2)
In this case, the corrosion tendency is primarily caused by the first two terms on the right
side of Eq 2, the dissimilarity in the standard
half-cell reduction potentials:
0
ENi
− EFe0 2 + / Fe = − 0.250 + 0.447 = 0.197 V
2+
/ Ni
This tendency is caused by the galvanic dissimilarity between the metals. This is normally
important for alloys joined to stainless that are
themselves less noble. Less-noble alloys, such
as carbon steel, can fail rapidly if coupled to
stainless. A classic example is the use of carbon
steel fasteners for joining stainless sheets. Different stainless steel alloys have minor differences when passive, but if the environment is
such that one alloy is active while another is
passive, then the galvanic differential could be
harmfully large. In many cases, a situation
arises in which access of oxygen is not the same
to different areas of a sample. In effect, this results in that the cathode reaction:
O 2 + 2H 2 O + 4e − → 4 OH −
(Eq 3)
is limited from proceeding in some areas but not
others. This gives rise to a differential aeration
cell. For example, consider Fig. 20(a), in which
a metal is partially immersed in water. Transport distance of oxygen increases with depth;
the limiting current would then vary with depth,
such as at locations 1 and 2, and cathode polarization curves as a result of this are schematically plotted in Fig. 20(b). Near the surface of
the water, where oxygen is readily replenished,
passivation is likely to be fast, and thus anodic
iron dissolution is slow. This region assumes the
role of the cathode, and reaction 1 occurs. Sufficiently far away from the surface, if there are
regions where passivation is incomplete (e.g.,
surface defects or scratches) or has broken
down as a result of, for example, Cl− (see section on pitting), repassivation does not readily
occur since oxygen transport is too slow. These
regions become anodes where the following reaction occurs:
2Fe (s) → 2Fe 2+ + 4e −
(Eq 4)
The distance at which this occurs is balanced
by being large enough to limit the rate of
oxygen transport but not too long to be
strongly influenced by ion transport that is
needed to complete the electrochemical cell.
Resulting corrosion currents are shown in Fig.
20b. This type of degradation is called waterline corrosion.
Crevice Corrosion. In stainless, the more
significant occurrence of this type of cell occurs
when a crevice, from whatever cause, exists,
and reactions within the crevice or pit cause the
accumulation of iron ions by:
2Fe (s) → 2Fe 2+ + 4e −
(Eq 5)
The regions adjacent to the drop that maintained their passive layer and have access to
oxygen act as cathodes where the oxygen reduction reaction takes place:
O 2 + 2H 2 O + 4e − → 4 OH −
(Eq 6)
This reaction maintains an alkali solution. As
a result of the geometry, Fe2+ ions remain and
Chapter 4: Corrosion Types / 39
Fig. 20 Schematic illustration of (a) sample partially immersed in water; (b) resulting polarization behavior for two different passivating alloys (A and B polarization curves)
enrich in the water-filled pit; to maintain charge
neutrality, Cl– migrates into the pit. This causes
the following reaction:
)
Fe 2+ + 2H 2 O + 2Cl − → Fe ( OH 2 + 2HCl
(Eq 7)
which has several consequences: (a) Hydrochloric acid further acidifies the pit and
increases the rate of iron dissolution since decreasing pH increases cathode half-cell potential, which increases corrosion rate (see polarization diagram construction in Chapter 3,
“Corrosion Kinetics.”) (b) The formation of
porous Fe(OH)2 further helps to isolate the pit,
thereby separating anode and cathode regions in
the differential aeration cell. (c) The presence of
Cl– prevents repassivation.
As a result of increased acidification, the dissolution rate becomes autocatalytic, and as a result the pit grows in depth.
At the outside, the reaction:
)
)
2Fe ( OH 2 + O 2 + H 2 O → 2Fe ( OH 3
(Eq 8)
further consolidates the isolation of the pit and
impedes the ingress of oxygen.
Pitting Corrosion
Pitting corrosion is the most intensely studied
and debated form of corrosion of stainless steel.
Pitting corrosion is important to designers because it is corrosion under conditions at which
corrosion may not have been anticipated. Thus,
it is both a materials selection and an environmental control problem. Its consequences may
be only cosmetic, such as on a building or appliance facade, or potentially catastrophic, such as
if leaks of toxic materials were to result from
perforation. Stainless steels are designed to be
passive, and localized corrosion is the local loss
of passivity. Whether the consequences are
major or not, it is always undesirable, and good
design allows it to be avoided.
What do we know for certain about pitting?
We know quite a lot, really. Experts now conclude that since the early 1970s the local chemistry of pitting has been understood (Ref15).
The greatest contributions to this field have
been electrochemical studies. The tools of electrochemistry have been especially successful in
elucidating the mechanism involved in pit
growth and pit stability (Ref 16). The local environment within pits has been sufficiently
measured and correlated with cavity geometry
that some experts can say, “In a sense, all pitting
is crevice corrosion” (Ref 15). This is to say that
the electrochemistry of cavities such as pits and
crevices is quite similar and has been well modeled. These same tools, however, have been
much less successful in clarifying the mechanism of pit initiation, which is still the subject
40 / Stainless Steels for Design Engineers
Inclusions. The question of what causes the
initial dissolution that causes both stable and
metastable pits focuses on inclusions, which
most authorities (Ref 18) have concluded are
associated in some way with pit initiation. In
the absence of inclusions, metastable pitting
events are not noted, and the potential at which
pitting occurs is the beginning of the transpassive regime.
What are the typical inclusions in stainless
steel? Inclusions in steel are normally the
residue of normal deoxidation and desulfurization taken during steel refining usually done in
an argon oxygen decarburization (AOD). After
removal of the carbon, the subsequent objective
is to remove or render less harmful the dissolved oxygen and sulfur, which if left in solution would later precipitate as low-meltingpoint iron compounds that would make the steel
fragile and unworkable at high temperatures.
Inclusions in stainless steel are typically oxides
and sulfides. A key point to understand when
considering inclusions as initiation sites for pitting is that inclusions are not simply inert debris
but precipitates that are seeking thermodynamic
equilibrium with the steel in which they have
previously been dissolved. The reactions in
stainless steel differ thermodynamically from
those in carbon steel because of the presence of
high chromium concentrations. This lowers the
activity of oxygen and sulfur, making them
more soluble, as Table 1 indicates (Ref 19). It
also alters the efficiency of deoxidizing elements. Aluminum is a powerful deoxidant in
carbon steel but is less effective in stainless,
while titanium becomes a stronger deoxidizer in
stainless. Their effect on sulfur is similar to that
on oxygen.
The bottom line is that oxygen and sulfur are
generally removed by silicon/manganese deoxidation, but that this process occurs in both the
liquid and solid states. That it carries over significantly into the solid state means that diffusion has a major role in determining if equilib-
of debate, possibly indicating that the root
causes are more metallurgical than electrochemical.
Figure 21 depicts a polarization curve for
stainless steel in a chloride-containing solution.
Pitting occurs in the zone in which passivity is
expected. As potential increases, small spikes in
corrosion current occur. These spikes measure
local dissolution, called metastable pitting.
Some such sites complete their dissolution and
repassivate, while others continue to grow as
stable pits. The potential at which stable pitting
occurs is the pitting potential, while metastable
pitting can occur at much lower potentials. Pitting events, stable or not, cause the generation
of iron ions and local pH reduction. To the extent these remain concentrated in a small volume, they will affect subsequent events. The
dissolution during metastable pitting is located
at the matrix-inclusion interface. Different researchers assume dissolution of the inclusion,
while others assume dissolution of the matrix.
The dissolution parameters, as measured by
current transients, depend on variables not of
the inclusion chemistry but of the matrix composition, notably molybdenum and nitrogen
levels (Ref 17), which is in keeping with the reduction in dissolution of the matrix that these
alloying elements confer.
Fig. 21 Schematic of a passive anode polarization curve
Table 1 Typical values of activities and activity coefficients in liquid steels: activities in the 1 mass %
solution: ai = fi . %i
Metal
Carbon steel,
1600 °C
Stainless steel,
1600 °C
%i
fi
ai
%i
fi
ai
Al
C
Mn
P
S
Si
Ti
H
N
O
Cr
Ni
...
1.05
...
...
3.6
...
0.05
1.06
0.053
0.05
0.49
0.025
0.45
1.0
0.45
0.45
1.0
0.45
0.02
1.1
0.022
0.02
0.32
0.006
0.01
1.0
0.01
0.01
0.66
0.007
0.3
1.15
0.345
0.3
1.24
0.372
0.05
0.93
0.046
0.05
9.4
0.47
...
1.0
...
...
0.93
...
...
0.97
...
...
0.17
...
...
0.85
...
...
0.21
...
...
...
...
...
...
...
18
0.97
17.5
8
1.0
8.0
Chapter 4: Corrosion Types / 41
rium reactions occur and whether they go to
completion. We will see that they do not.
Oxide inclusions also are common. They are
formed as the products of the reactions of silicon and manganese with dissolved oxygen. The
thermodynamics of the reactions determine at
any time how much oxygen can be dissolved in
the steel at equilibrium. That equilibrium is easily achieved in the molten state, in which diffusion is very rapid, but achieved more slowly
once the material has solidified. The inclusions
in the solid state grow by the diffusion of oxygen to inclusion sites, where it precipitates as an
oxide of silicon or manganese to the extent that
these are locally present or of chromium when
its local concentration (or more properly, its activity) makes it more favorable. These oxides
are often the nucleation sites for manganese sulfide inclusions.
Sulfur is a very surface active impurity that assists in weld penetration in stainless by virtue of
its effect on weld pool circulation. Otherwise, it
is a detrimental impurity, forming low-melting
oxysulfides that diminish hot workability. Manganese is a strong sulfide former, and it is the
main agent used to tie up sulfur. Manganese sulfide precipitates as an inclusion as a function of
manganese and sulfur concentrations and temperature. Inclusions form not only in the molten
metal but also in the solidified metal. The solubility, which is high in the liquid state, decreases on
solidification, as seen in Fig. 22. Only resulfurized free-machining stainless steels have sufficient sulfur to precipitate manganese sulfide in
Fig. 22 Pseudo-binary-phase diagram for iron and sulfur at
1.8% manganese and 18% chromium
the liquid. At high sulfur and manganese concentrations, some manganese sulfides can precipitate
during solidification interdendritically, while normal alloys with less than 100 ppm of sulfur form
their inclusions after solidification. The distinction is important because precipitation in the liquid state permits rapid diffusion, which results in
the most thermodynamically favorable species,
manganese sulfide, to form. It may, and often
does, nucleate on a preexisting inclusion, such as
silicate present from the deoxidation process. In
austenitic steels, manganese is generally present
at a level of around 1.5% as a deoxidant and as a
substitute for some nickel. An inclusion formed
in the molten metal does not cause alloy depletion around it. One that forms or grows in the
solid state does cause depletion of the elements
that are precipitating, causing its growth.
If manganese is lowered to very low levels,
the supersaturation of sulfides is pushed to a
lower temperature, at which lower diffusion
rates hinder or prevent the precipitation. Thus,
low-manganese alloys can be free of manganese
sulfide inclusions even at somewhat high sulfur
levels. Such alloys have elevated resistance to
pit initiation. Lower manganese levels also thermodynamically reduce the chromium sulfide coprecipitation in inclusions, lowering chromium
depletion around manganese sulfide/chromium
sulfide inclusions.
Elements more effective than silicon and
manganese are now in use for deoxidation and
desulfurization. These include aluminum, calcium, cerium, and other rare earth metals
(REMs), and titanium. The action of calcium is
notable. In a well-deoxidized and well-stirred
melt and with a basic slag, calcium dissolved in
the metal will react with dissolved sulfur to
form calcium sulfide, which will be incorporated into the slag phase. Aluminum, while a
potent deoxidizer, is less effective directly in
desulfurization, but it can act indirectly by reducing a small amount of Ca2+ in the slag, allowing the formation of calcium sulfide. Titanium can sequester some sulfur as titanium
carbosulfide precipitates. The greatest amount
of sulfur removal is obtained by the addition of
cerium or other REMs, usually in the form of
the alloy mischmetal. These reactive elements
typically form oxysulfide particles in the melt
that may be trapped in the slag before metal
solidification.
Oxygen is normally dissolved in solidifying
stainless steel, also at amounts in the neighborhood of 100 ppm depending on deoxidation
42 / Stainless Steels for Design Engineers
methods. Inclusions based on oxygen and sulfur
formed in the liquid or during solidification are
relatively large, greater than 1μ. As the alloy
cools after solidification, precipitation continues
since sulfur and oxygen are decreasingly soluble with temperature, to virtually nil at room
temperature. This causes existing inclusions to
grow and new ones to nucleate. This precipitation is similar to that which carbon undergoes in
stainless, except carbon is generally not supersaturated until below 1200 °C at the highest in
most alloys, whereas sulfur and oxygen are normally near saturation even at freezing or almost
always when the solidifying ferrite transforms
to austenite. Thus, inclusions grow via diffusion
of oxygen and sulfur, which, as interstitials, diffuse much more rapidly than the silicon or manganese with which they have the greatest thermodynamic affinity. But precipitate they must,
even if the silicon and manganese in the vicinity
of their inclusion are exhausted. Thus, inclusions can grow with chromium substituting for
either silicon or manganese as the precipitating
partner for oxygen and sulfur. The inclusion
growth necessarily depletes the surrounding region of reactants, silicon, manganese, and
chromium (Ref 20). Inclusions thus formed are
nonequilibrium in nature, and thermal cycles of
steel production are rarely sufficient for the
equilibrium to be attained. The chromium enrichment of such inclusions and corresponding
chromium depletion of surrounding regions has
been measured (Ref 21) and corresponds to the
depletion seen next to chromium carbide precipitates at grain boundaries in sensitized alloys. These zones are altered in size and shape
by thermomechanical processing in wrought alloys but exist fairly undistorted in welds. Hot
rolling and cold rolling followed by annealing
elongate manganese sulfide inclusions and flatten them, allowing depleted zones around the
inclusion in the reduced dimension to be more
rapidly homogenized during annealing. Thus,
wrought material has better pitting resistance
than cast or welded material. Inclusions that
precipitate from the liquid, as is more the case
for alloys solidifying in an austenitic mode, are
at equilibrium with the surrounding matrix by
virtue of the faster diffusion in liquids, do little
to diminish the chromium content around them,
and have a small effect on lowering pitting resistance. Pitting resistance is still affected to a
degree by alloy depletion due to solidification
segregation. However, if the alloy solidifies in a
ferritic mode (FA, i.e., ferrite forming first on
solidification as opposed to austenite first, AF),
as is almost always the case with commercial
alloys, more sulfide precipitation happens in the
solid state, pitting resistance is lowered proportionately to the sulfur level (Ref 22), and there
is little negative effect from solidification segregation, as is shown in Fig. 23 and 24 (Ref 23).
Solidification can also occur in a mixed ferriticaustenitic mode, in which case each microstructural component behaves according to the chart.
The ratio of chromium and chromium-like elements molybdenum and silicon to nickel and
nickel-like elements carbon, nitrogen, manganese determines the mode of solidification. It
Fig. 23 Influence of sulfur level on pitting resistance of
unannealed welds for different solidification modes.
Source: Ref 23
Fig. 24 Influence of sulfur level on pitting resistance of welds
without homogenizing anneal. FA, ferrite forming
first on solidification as opposed to austenite first, AF. Source:
Ref 23
Chapter 4: Corrosion Types / 43
can also be altered by freezing rate. Faster cooling favors austenitic solidification.
Long-term annealing of welds has shown
that sufficient time and temperature to achieve
some rehomogenization the alloy result in better pitting resistance (Ref 24), approaching that
of the wrought alloy. Examination of the decreasing solubility of sulfur in stainless in Fig.
22 indicates that the precipitation of sulfides
that cause chromium depletion occurs in delta
ferrite on freezing when sulfur exceeds 0.007%
and in austenite when sulfur exceeds 0.003%.
Oxygen behaves in a parallel manner and is
usually present in sufficient quantities, about
0.01% in manganese/silicon deoxidized steels,
to cause the same phenomenon. This fundamentally is due to the high ratios of the diffusivities of oxygen and sulfur to chromium,
which are about 10,000 and 680, respectively.
Whenever fast-diffusing elements such as oxygen, sulfur, carbon, and nitrogen, which have a
strong affinity for chromium and a solubility
that decreases strongly with temperature, are
present in steel, their precipitation will result in
some degree of chromium depletion around the
precipitation site because chromium diffuses
too slowly to be replenished.
The low chromium around inclusions is a sufficient condition for the local dissolution measured as metastable pitting, and if the depletion
zone shape and size are favorable, then stable
pitting would ensue.
Certain other types of inclusions/precipitates
are less harmful in this regard. Titanium, for instance, which is often added to form carbides
and nitrides, also forms sulfides and oxides
more strongly than manganese and therefore
does so at higher temperatures. Such precipitates have a much lower tendency to allow
chromium to join in the precipitation since the
higher the temperature of precipitation the more
that diffusion allows the more favorable reaction to occur. Rare earths also behave the same
way. Metastable pitting is diminished by the
presence of these elements.
The initiation of pitting is also affected by
stress and inclusion orientation (Ref 25), which
the researchers correlated to the dimensions of
the inclusion-derived cavity being able to sustain a sufficiently low pH due to iron dissolution to maintain stable pitting. The influence of
stress was to cause cracking at otherwise unfavorably shaped inclusions, which then provided
a crevice capable of sustaining stable pitting.
This will be relevant to later discussions of
stress corrosion cracking (SCC). There have
been numerous proposed mechanisms for the
breakdown of a passive film in chloride-containing media; these have been summarized in
other publications (Ref 23). These hypotheses
deal with how a passive film on a homogeneous
surface could break down. They include:
• Adsorption of chloride ions
• Penetration of the passive film by chloride
ions
• Film breakdown by electrostriction
• Formation of stable metallic chlorides
• Coalescence of cationic vacancies
• Random localized thinning of the passive
film
• Local variations in the composition of the
corrosive medium
By and large, these mechanisms presuppose
a stainless steel surface that is homogeneously
passive and try to explain the observed inhomogeneous behavior of the passive film. However, since it is clear that the surface is not homogeneous, especially with regard to the
passive film, these hypotheses are not necessary to explain the behavior of everyday stainless steels, which unfortunately have abundant
inclusions and chemical inhomogeneities capable of locally diminishing the integrity of the
passive film. More research in understanding
the exact nature of the inhomogeneity of stainless steel surfaces is necessary for a complete
understanding of pit nucleation and therefore
prevention.
Pitting Resistance. Pitting has been extensively correlated with environment and compositional variables. The most well-known and
useful correlations are between the PREN and
the critical pitting temperature (CPT) and by extension to the pitting potential.
For austenitic alloys:
PREN = % Cr + 3.3 % Mo + 30 % N
(Eq 9)
For ferritic alloys, which hold no nitrogen in
solution:
PREN = % Cr + 3.3 % Mo
(Eq 10)
For duplex alloys, which have two phases,
neither of which matches the bulk composition:
PREN = % Cr + 3.3 % Mo + 16 % N
(Eq 11)
These equations are useful, if approximate,
and their correlation is shown in Fig. 25 (Ref
26) They do not include tungsten, which, if
44 / Stainless Steels for Design Engineers
Fig. 25 Variation of critical pitting temperature with pitting resistance equivalent number (PREN) of
austenitic steels in water plus 6% FeCle. Source: Ref 26
Fig. 26 Differential variation of critical pitting temperature
of several stainless steel alloys for unwelded
wrought and welded material. Source: Ref 13
present, has half the effectiveness of molybdenum. They neglect carbon, which seldom varies
enough to have a visible effect but has been
shown when in colossal supersaturation to have
a factor of about 10, not unlike nitrogen, another interstitial that it resembles in solution
thermodynamically (Ref 27). It also does not include the negative influence of elements such as
sulfur. Likewise, the equations cannot deal with
inhomogeneity issues, so welded alloys have
different CPTs for the same PREN (Fig. 26)
(Ref 13). These equations are all-other-thingsbeing-equal equations and are useful for gross
alloy behavior predictions. It is noteworthy that
the elements copper and nickel, which are beneficial against uniform corrosion and which slow
the growth of pits by this same action, do not
contribute to increasing the resistance to the
onset of pitting. This is another manifestation of
pitting initiated by the local stability of the passive film, which is primarily a function of local
chromium content. Molybdenum and nickel
thus seem to bolster local chromium content in
the passive film. Nitrogen seems to act by concentrating at the passive film-alloy interface
rather than by buffering the solution by ammonia formation, which has been proposed (Ref
36). Research on very pure sputtered films of
iron-chromium alloys have demonstrated that
both titanium and niobium in solution diminish
active dissolution, assist repassivation, and improve pitting resistance (Ref 29). In most practical cases, these elements are not found in solution because of their affinity for oxygen, sulfur,
carbon, and nitrogen, with which they form
compounds.
It should also be noted that the critical PREN
values vary with crystallographic structure. Ferritic alloys require somewhat lower PREN values to exhibit similar pitting resistance as
austenitic alloys of somewhat higher PREN.
While pitting is of great theoretical and practical interest, there are significant problems in
actually conducting good tests. Monitoring of
the electrochemical potential during the test is
considered mandatory by most researchers.
How is a metallic sample suspended in a solution without creating any crevices and without
exposure at the liquid-gas interface? The development of the flooded gasket technique (used in
ASTM G150) was a milestone, but it also has
some problems—most notably the potential for
dilution of the test solution, especially during
prolonged testing. FeCl3 testing benefits from
the fact that the solution creates a reproducible
positive potential.
While the PREN approximates the pitting resistance of an alloy, there is a standard test by
which the CPT is measured. Pitting in a given
medium capable of causing pitting does not
occur below a temperature that is characteristic
of the medium and the material, with the myriad
exceptions of stress state, surface finish, microstructure, etc. The most commonly used test
media are the unacidified 10% FeCl3, which is
used in the ASTM G 48 practice B, and the 3.5%
NaCl solution of the ASTM G 150. The latter, if
modified to 0.1N NaCl, allows the ECPT, the
Chapter 4: Corrosion Types / 45
electrochemical pitting potential, of lower alloys
such as 304 to be measured (Ref 30).
Crevice Corrosion
In the case of pitting, the geometry that
makes up the pit is essential in creating the differential aeration cell and to cause the autocatalytic dissolution process. In many cases, a
geometry that retains and acidifies water is already present in crevices in different types of
structures such as gaskets, under faulted coatings, under bolt or screw heads, etc. Crevice
corrosion occurs because zones have restricted
access of reactants and restricted exit of corrosion products. It is especially the inhibition of
the cathodic reaction inside the crevice by the
dearth of oxygen, which sets up a more aggressive environment within the crevice than without. The interior reactions become increasingly
anodic, and the aggressiveness of the environment can reach a threshold at which active corrosion occurs, while the situation exterior to the
crevice is safely passive. Crevice corrosion occurs at lower temperature than pitting in the
same environments, so it is a greater danger in
that sense.
The relationship between the alloy content,
given as the crevice corrosion resistance equivalent number (CCREN), and critical crevice
corrosion temperature (CCT), shown in Fig. 27
(Ref 11), is similar to that of PREN (PI) to CPT
except for the molybdenum factor being more
important:
Cl = %Cr + 4.1%Mo + 27%N
(Eq 12)
Since a crevice has a preexisting favorable
geometry for pit growth, any pitting event,
metastable or stable, can initiate ongoing
crevice corrosion. Crevices are thus incubators
for corrosion triggered by metastable pitting
events. The dissolution of iron during passivation itself as well as the differential oxygen cell
created by the crevice contribute to the process.
It is logical to think that alloying the elements
that contribute to lowering the critical current
density for passivation and the uniform corrosion rate, such as nickel, would reduce the creation of the reactants that start the crevice corrosion process, but this presumed effect is not
strong enough to be reflected in this actual behavior Eq 12 represents, although it is generally
acknowledged that austenitic steels perform better than ferritic steels in the absence of molybdenum. Materials are characterized as having a
critical depassivation pH. If crevice conditions
are such that the reactions over time allow the
pH to be reduced to this level, then active corrosion will begin within the crevice. Thus, passive
film stability seems to be the critical factor
rather than corrosion rate after initiation.
Preventing Crevice Corrosion. The countermeasures against crevice corrosion are cathodic protection, design, maintenance, and, of
course, alloy selection. Designing to avoid
crevices should include maximizing the volume
of unavoidable crevices, engineering flow to enhance transport in and out of crevices, and
Fig. 27 Variation of critical crevice corrosion temperature with alloy content
46 / Stainless Steels for Design Engineers
avoiding stagnation. Any maintenance or design
procedure that prevents formation of deposits is
beneficial. Welds are particularly vulnerable
surface sites, so any combination of welds and
crevices or crevices caused by poor weld geometry must be avoided. S32205 is a benchmark
alloy of sorts. It has just sufficient alloying to
resist pitting in seawater, but it is susceptible to
crevice corrosion.
As a practical matter, crevices are almost impossible to eliminate. Threaded fasteners and
joints represent severe crevices and should be
avoided in aggressive environments if possible.
Gasketed joints are another severe crevice location, and their usage should be curtailed to the
minimum practical extent. In these situations,
judicious use of very expensive, highly corrosion
resistant materials is justified. The use of smooth
welded joints is thus generally preferred. In a
more general consideration, deposition and
fouling create crevice sites, and design and operational controls to preclude the formation of
deposits and the prompt removal of sludge and
the like are necessary. But in some situations,
such as marine exposures, biofouling will create
crevice sites. This fouling may be macroscopic,
such as from shellfish and barnacles, or it may
be microscopic. Microscopic biofouling causes
the special form of crevice corrosion called microbiologically influenced corrosion (MIC) discussed in a separate section).
Sensitization/Grain Boundary Corrosion
The maintenance of a passive layer in a wide
range of pH conditions in stainless steels is dependent on the alloying elements, primarily
Fig. 28 Schematic illustration of sensitization due to
chromium-rich precipitates that deplete adjacent
regions of chromium. GB, grain boundary
chromium. In the various grades of stainless
steels, there are many intermetallic phases that are
thermodynamically stable but kinetically slow to
precipitate that are enriched in chromium. An example of such a phase is chromium carbide (Fe,
Cr)23C6. These phases tend to form at grain
boundaries where nucleation is favored, resulting
in a depletion of chromium in the adjacent regions, as shown in Fig. 28. Thus, the chromiumdepleted regions near the grain boundaries are
sensitized in that they behave as active anodes
compared to the larger interior of the grains that
are still passive. In an aerated corrosive environment, the smaller chromium-depleted nonpassive
anodes dissolve, whereas the larger cathodes reduce oxygen, resulting in a localized corrosion
along grain boundaries. Any heat-treating or
welding procedure of stainless steels should thus
be tailored to avoid sensitization.
When a stainless steel is heat treated, there is
a risk that the unwanted phases may form, depending on the time-temperature history and
precipitation kinetics of the unwanted phase.
Figure 29 shows schematically the temperature
versus time due to welding and the resulting
sensitization. Figure 29 shows a TTT (timetemperature-transformation) curve for precipitation of the unwanted phase. Near the weld
(A), the time spent in the temperature region
where precipitation occurs is too short, whereas
far away from the weld (C) the temperature experienced is too low. At location B, there is,
however, a risk for sensitization.
Austenitic. Sensitization can occur at any
temperature at which carbon is supersaturated
in an alloy. Current austenitic stainless steels
have carbon levels of under 0.10% normally
Fig. 29 Schematic illustration of how a heat treatment re-
lates to sensitization due to precipitation kinetics.
TTT, time-temperature-transformation
Chapter 4: Corrosion Types / 47
and under 0.03% for low-carbon L grades.
Thus, normal grades sensitize below around
800 °C. The supersaturation increases with
decreasing temperature, but below about 500 °C
diffusion of carbon is too slow for carbon to
move to grain boundaries and cause the damaging combination with chromium that causes
sensitization. Low-carbon grades avoid sensitization because they are not sufficiently supersaturated at temperatures at which carbon is
mobile enough to diffuse to grain boundaries.
Ferritic. Another situation exists in ferritic
stainless steels, in which carbon is much less
soluble but is much more mobile. Annealing
over 900 °C can put enough carbon in solution to
cause sensitization even at the lowest carbon levels attainable in an AOD and even at the fastest
possible quench rates. The damaging chromium
depletion caused by this very rapid precipitation
can be undone by a simple rehomogenization
anneal of the remaining chromium. This is theoretically possible with austenitic alloys also, but
the diffusion rates of chromium in austenite as
so slow that it is impractical in most real cases.
Duplex steels have a subtle near immunity to
carbide sensitization. While they are typically
low carbon anyway, the carbides that do form
do so at ferrite-austenite grain boundaries. Here,
chromium is consumed from the chromium-rich
ferrite phase, leaving the austenite intact. Their
large grain boundary area keeps carbide concentration per unit area low, and the fast diffusion in the ferrite keeps austenite from becoming depleted. However, the rapid formation of
intermetallic phases at the ferrite-austenite interfaces can lead to a rapid loss of corrosion resistance and a severe loss of toughness if exposure to temperatures within the intermetallic
precipitation range is not controlled.
Martensitic steels are quenched as austenite
to and through the Ms temperature without time
for carbon to precipitate in austenite. The carbon in the martensite can precipitate and cause
sensitization if reheated to the 300 to 700 °C region. Fortunately, heating to above 700 °C rehomogenizes the chromium and eliminates sensitization.
Effect of Alloying. Besides determining
basic phase structure, alloying plays a role in
susceptibility to sensitization. Those elements
that reduce the tendency of chromium carbides
to form also reduce the susceptibility to sensitization. This is a purely thermodynamic effect.
Molybdenum, silicon, and nickel promote carbide formation by increasing the thermodynamic
activity of carbon and make alloys more susceptible. Nitrogen lowers the tendency for carbide
formation and slows sensitization. Nonthermodynamic effects are those of austenite grain size
and prior cold work. Decreasing grain size and
therefore increasing grain boundary surface
area decreases the amount of precipitate per unit
area of grain boundary and therefore the amount
of chromium depletion per unit area. Cold work
accelerates diffusion and makes precipitation
more rapid, thus aggravating sensitization.
The thermodynamic affinity tool can be used
to prevent chromium carbide formation in another way. Introducing alloying elements that
combine with carbon more strongly and rapidly
than chromium can exhaust the supply of carbon available to precipitate as chromium carbide. There are a number of candidate elements,
zirconium, vanadium, tantalum, niobium, and
titanium, most prominently. Of these, the diffusivity and affinity for carbon of niobium and titanium make them the best for this purpose.
Each forms stable carbides at much higher temperatures than chromium, starving chromium of
sufficient carbon to form damaging precipitates.
The caveat with titanium is that it forms oxides,
sulfides, and nitrides preferentially to carbides.
Therefore, sufficient quantities must be used to
accommodate the prior formation of these
phases. Niobium tends more toward carbide
than nitride formation but is a weaker carbide
former than titanium. The solubility products of
these precipitation reactions are:
log [Ti ][C ] = 2.97 −
6780
T
(Eq 13)
9350
T
(Eq 14)
log[ Nb][C ] = 4.55 −
These equations follow the form of the general equation for precipitation reactions:
log[ M ][ X ] = A − H / RT
(Eq 15)
in which A is a constant, H is the heat of dissolution, R is the gas constant, and T is the absolute temperature. If the amount of titanium or
niobium is stoichiometrically sufficient, no carbon will form chromium carbides under equilibrium conditions. It is possible to defeat the stabilization reactions by quenching the alloys
from temperatures at which titanium carbide or
niobium carbide is dissociated. If free carbon is
48 / Stainless Steels for Design Engineers
left free in the matrix by quenching, then on reheating it may form carbides with the most locally accessible favorable element, such as
chromium, rather than the most thermodynamically favorable element, which would be titanium or niobium. This can occur when a stabilized alloy such as 321 is welded. A zone away
from the weld may experience a high enough
temperature to put carbon into solution and then
cool just rapidly enough to not form only the
equilibrium titanium carbide but also Cr23C6 at
grain boundaries, causing the type of sensitization called knife-line attack. This problem has
nearly ceased to exist as modern 321 has low
levels of carbon and nitrogen for economic reasons; this effectively precludes this chromium
carbide precipitation in most cases.
Welding. Many of the most severe problems
of sensitization arise when stainless steels are
welded to carbon or low-alloy steels. In these
situations, construction code rules usually require that the carbon steel component be given a
stress relief annealing (SRA) treatment. Such
SRA treatments are typically in the sensitization
temperature range for austenitic stainless steels.
Use of low-carbon or stabilized grades is necessary in such cases. Even then, use of the lowest
allowable temperature SRA treatment for the
shortest allowable time is preferred.
Corrosion Combined with
Fatigue or Fracture
Environmentally induced failure occurs when
brittle failure under tensile mechanical loading
occurs at a lower stress when a material is subjected to a corrosive environment than what
would happen in a noncorrosive environment.
This introduces us to what is perhaps the most
controversial technical subject in all of stainless
steel research, SCC.
layer occur, and region II, where the protective
layer is not fully developed, suggesting an appreciable electrochemical effect. The latter is a
zone that exists in alloys that have zones of
chromium depletion.
Stress corrosion cracking has always been
among the most controversial subjects among
metallurgists and electrochemists. The debate
centers on whether the critical mechanism is
dissolution or fracture, and if a fracture, by what
mechanism. Is the cracking zone locally softened, locally hardened, transformed, to a more
brittle phase or embrittled by hydrogen? As of
this writing, there is no general agreement on
which type mechanism is the fundamental
cause, but there is room for convergence. Obviously, elements of many may come into play. It
is likely, as in most prolonged arguments, that
no hypothesis is completely correct. We will try
to fairly set out what is known and agreed on as
fact and then present researchers’ views in an
unbiased manner, but since we concern ourselves only with stainless steel, no attempt is
made to address an all-encompassing theory.
Crack Initiation. In stainless steels, cracks
can be seen to initiate at surface defects and irregularities. In stainless steel, it must be agreed
by all that the preponderant initiation site is a
corrosion pit or, in some cases, a crevice. Intergranular corrosion sites, as are seen in sensitized material, can also provide the conditions
for SCC initiation. The interrelationship between pits and SCC cracks has been studied
(Ref 25). Stress lowers the anodic potential at
which pitting occurs and permits metastable pits
Stress Corrosion Cracking
The key cause for SCC is the cooperating effects of tensile stress and a corrosive environment. Such cases can be identified in most alloy
systems, and even pure metals, which were
thought to be more or less immune, also have
had cases of SCC reported.
In passive metals, two sensitive potential regions for the occurrence of SCC have been
identified and are shown in Fig. 30: region I,
where pitting and breakdown of the passive
Fig. 30 Zones of susceptibility to stress corrosion cracking
Chapter 4: Corrosion Types / 49
to become stable via the generation of cracks.
Cracks, once formed, presumably have favorable geometry to duplicate pit internal chemical
reactions and must be considered to be described by the models that apply to pits and
crevices.
The stress at which SCC initiates has a
threshold, which has been reported as between
25 and 50% of the yield strength in austenitic
stainless steel.
The temperature at which SCC is initiated
ranges from ambient to under 100 °C for
martensitic materials, while austenitic alloys
begin their sensitivity above room temperature
and increase in susceptibility with increasing
temperature. The ferritic steels, while considered nearly immune to SCC, have their maximum susceptibility in the same range as martensitic steels. In environments of mixed chlorides
and sulfides, however, SCC can occur in all
types of stainless at room temperature. This has
been seen in the SCC of austenitic stainless
steel in swimming pool environments, in which
chloride ions can condense on the stressed steel
and cause pitting and SCC.
Cracks propagate very slowly below specific certain stress intensity levels, but once that
intensity is reached, they have a plateau rate
that is fairly constant until the stress level at
which catastrophic failure occurs at very high
propagation rates. Rates of crack propagation
are exponentially increased by increasing temperature. The crack propagation rate has been
seen across a range of alloys to be linearly proportional to the average current density that
alloy experiences when its surface is strained,
indicating that reactions at the crack tip are
strain sensitive, and overall rate limiting, but
not necessarily the mechanism of cracking.
Crack growth is discontinuous with individual
steps of growth many times the average rate,
which is similar to that seen with gaseous hydrogen embrittlement (HE). The crack growth
gives off acoustic emissions as cracking steps
occur. These steps of growth are brittle and are
seen as facets on fractographs with cleavages
corresponding to crystallographic planes. The
crack facets match with high perfection, showing almost no evidence of plastic deformation
or dissolution.
The propagation path may be intergranular or
transgranular. Grain boundary propagation in
stainless steels usually corresponds to conditions under which grain boundaries are less corrosion resistant because of either material or
environmental variables. The most common example is that of sensitized 304 in high-temperature water or caustic media. The relevance of
this to the normal case of stainless steels must
be questioned since, by definition, the sensitized
grain boundaries themselves can be depleted of
chromium to a degree they are not stainless and
have a much less stable austenitic structure,
having their martensite start temperature Ms,
raised by the loss of chromium.
Material Variables. Martensitic stainless
steels and martensitic precipitation hardened
stainless steels are quite susceptible to SCC.
This susceptibility increases with hardness,
yield strength, and embrittling heat treatments.
They will crack at threshold stresses equal to
50% of yield strength. These alloys can be tempered at sufficiently high temperatures that they
become soft and tough enough to have very
good resistance.
Ferritic stainless steels of low and medium
chromium are generally not susceptible to SCC.
Ferritic alloys, which can have a martensitic
structure, should be considered martensitic for
SCC purposes. If purely ferritic alloys are alloyed with copper, molybdenum, and nickel,
they can become susceptible. The presence of
α'( or high-temperature embrittlement also increases susceptibility, as does cold work.
Despite the controversy surrounding the
mechanism of SCC in austenitic stainless steels,
there is almost complete agreement that SCC of
body-centered cubic (bcc) stainless steels,
martensitic, ferritic, and pH is simply a manifestation of HE, with hydrogen provided by either anodic (e.g., active corrosion within a pit)
or cathodic reactions.
Duplex stainless steels have low susceptibility to SCC. Their dual-phase microstructure ensures that under conditions that crack austenite,
ferrite remains as a crack-arresting phase, while
under conditions that cause SCC in highly alloyed ferrite, the austenite is a crack arrester.
A second explanation of the resistance of duplex alloys to SCC is that their two phases have
different corrosion potentials, and that the
mixed potential that arises because they are in
intimate contact is outside the potential range
for SCC on either phase. This fits with the resistance to SCC of wrought alloys with a lamellar structure and the lesser resistance of cast alloys that lack that structure.
Austenitic stainless steels are the type of stainless steel generally associated with SCC, and
they vary in their degree of susceptibility to
50 / Stainless Steels for Design Engineers
SCC. All other things being equal, alloying
elements that delay or prevent localized corrosion do the same to delay SCC. This is simply
the delay of initiation. However, if pitting can be
delayed indefinitely, then SCC can also, assuming, of course, more harmful localized corrosion,
such as that due to intergranular chromium depletion, is not occurring. Molybdenum, which
we already know helps prevent pitting and
crevice corrosion, also increases the threshold
stress for SCC, as shown in Fig. 31 (Ref 31).
But, if metastable or stable pitting is occurring,
the threshold stress has been reached, and the
temperature is sufficient, then SCC will proceed.
It is mitigated by material variables such as cold
work and by alloying elements that increase
austenite stability. Many publications cite nickel
as beneficial in enhancing resistance to SCC,
often referring to the data from Fig. 32. However,
its role seems mainly to be as an austenite stabilizer and as a retarder of active corrosion. The
minimum in the curve corresponds to the nickel
level at which the structure is entirely austenitic,
but least stably so. Lower nickel levels produce
better immunity through the duplex structure,
while higher levels promote austenite stability
and correspond to alloys having more alloying
elements, such as chromium and molybdenum,
The cases of polythionate and thiosulfate solutions are industrially important but can be adequately explained as simply the stress-assisted
intergranular corrosion of sensitized material.
Hot caustic solutions are aggressive against
stainless steels. Certain combinations of concentrations, temperature, impurity, and dissolved
oxygen can cause SCC as well as other undesirable corrosive attack. Resistance to general corrosion is proportional to nickel content, but ferritics and duplex alloys are less prone to SCC.
304 has been reported to have no meaningful
threshold stress for SCC in hot caustic solutions,
leading one to question whether such a failure
should even be classified with SCC of the typical chloride-induced type or belong with the previous polythionate and thiosulfate solutions.
Chloride containing environments are the
main ones that induce SCC. Water can cause
Fig. 31 Influence of molybdenum on resistance to stress
Fig. 32 Variation of resistance to stress corrosion cracking
corrosion cracking (SCC) in austenitic steels
which are probably the greater cause of resistance to SCC.
Environmental Variables. There are three
key types of environments in which SCC occurs
in stainless:
• Chloride-containing solutions
• Caustic solutions
• Polythionate and thiosulfate solutions
with nickel (and other) content and structure
Chapter 4: Corrosion Types / 51
SCC at sufficiently high temperatures (i.e.,
above 100 °C) if there are even very low combined concentrations of chloride (greater than
0.1 ppm) and oxygen (greater than 0.1 ppm)
dissolved (see Fig. 33) (Ref 32).
Failure times decrease exponentially with decreasing chloride content. Crack growth rate increases by a factor of ten with each 30 °C rise in
temperature. Decreasing pH lowers the temperature at which SCC occurs in a given time.
Mechanisms. There have been many mechanisms proposed for SCC in stainless steel. We
focus only on those that address the failure in
chloride-containing media, the main concern for
users of stainless steel.
The models that have found some support
are:
• Slip dissolution
• Adsorption-enhanced plasticity
• Adsorption-induced brittleness
• Hydrogen embrittlement
Slip dissolution (anodic dissolution) was the
earliest proposed model for SCC. It simply proposes that at a crack tip a passive film forms, and
after time it fractures by an unspecified mechanism. The fresh active surface may or may not
repassivate, after which the process repeats itself.
The strength of this model is that it actually
does describe what is happening. The crack
does advance discontinuously, and after each
advance there is fresh surface, which comes
into equilibrium with the solution within the
crack. So, any experiment, such as that shown
in Fig. 34 (Ref 33), that tests crack propagation
against electrochemical events will absolutely
support this model.
It is axiomatic that films must rupture and reform as cracks advance discontinuously. The
weakness of this model is that it does not provide a mechanism for brittle fracture, and the
very brittle features of transgranular SCC fracture surfaces do not show any supporting evidence of dissolution. Research (Ref 34) showing that metal dissolution at the crack tip is
isotropic rather than crystallographically oriented make dissolution models incapable of
being reconciled with the crystallographic fracture surface facets.
Adsorption-induced brittleness, also known
as stress-sorption, looks to the parallels between liquid metal embrittlement and SCC to
explain the mechanism of SCC as the action of
adsorbed species weakening atomic bonds on
the crack tip surface. If the action is on the surface, however, the mechanism cannot produce
the observed discontinuous, brittle cracks which
characterize SCC. Only in alloys such as Fe-3Si
are steps small enough to make this mechanism
plausible.
Fig. 33 Variation of susceptibility to stress corrosion cracking (SCC) with media oxygen and chloride content for 304 stainless steel.
Source: Ref 32
52 / Stainless Steels for Design Engineers
Fig. 34 Crack propagation rates of various metals plotted versus current density. Source: Ref 33
Adsorption-enhanced plasticity/hydrogen embrittlement encompasses a number of models
that observe that adsorbed species enter the lattice in the vicinity of the crack tip and then
cause failure by one of several mechanisms:
• Dealloying and porosity
• Adsorption-induced brittleness
• Coalescence of voids formed by cross slip
enhanced by the adsorbed species
Since hydrogen is the only species that is produced in quantity and is capable of diffusing into
the lattice, HE is implicit in all these models.
All of the models have support in that they
have some experimental observations that show
that the phenomena they propose as causal
actually take place, but none is specific enough
to have been tested by critical experiments to
prove or disprove it.
It has been demonstrated that hydrogen is absorbed into the material at the crack tip. The
main question is whether it causes damage by
creating porosity, altering dislocation mobility,
or causing lattice decohesion. There is support
for each.
It has been observed that where SCC occurs
there is a large concentration of vacancies. This
has led to speculation that porosity is a weakening mechanism responsible for SCC (Ref 35). It
has been proposed and supported by calculations that hydrogen lowers the energy required
or vacancy formation. The lowest energy
Chapter 4: Corrosion Types / 53
configuration is calculated as two hydrogen
atoms per vacancy. This pairing of hydrogen
solute atoms to dislocations is very reasonable
given the major distortion the interstitial hydrogen causes to the lattice, so there is no basis to
challenge the enhanced vacancy formation.
Whether the effect is large enough to cause failures has not been demonstrated. The largest
measurable effect of hydrogen has been a slight
acceleration of stress relaxation in martensite.
The relevance of hydrogen-induced vacancy agglomeration as the principal cause of failure must
be considered questionable until some further
critical experiments link the vacancies to the observed instances of failure quantitatively, and
more important, to show how this mechanism
could account for the temperature and stress dependence observed. The major influence of vacancy formation due to hydrogen may be to enhance the volume expansion due to hydrogen.
Adsorption-induced brittleness, also known
as stress sorption, looks to the parallels between liquid metal embrittlement and SCC to
explain the mechanism of SCC as the action of
adsorbed species weakening atomic bonds on
the crack tip surface. If the action is on the surface, however, the mechanism cannot produce
the observed discontinuous, brittle cracks that
characterize SCC. Only in alloys such as Fe3Si are steps small enough to make this mechanism plausible. In stainless steels, there
seems to be nothing to support this proposed
mechanism.
Adsorption-enhanced plasticity has become
known recently as HELP or hydrogen-enhanced
localized plasticity. The underlying mechanism
at work in this model is the hydrogen-induced
shielding between microstructural defects. This
has been observed distinctly in single crystals of
austenitic stainless alloys. The Cottrell atmosphere of hydrogen around dislocations causes
mutual repulsion, causing strain to be localized
on certain slip systems. This has been observed
to occur and has caused deformation to become
concentrated in Luders bands in austenitic alloys, which of course do not show such behavior without hydrogen (Ref 36). This also produces ε-martensite in austenitic alloys, which
would be considered stable without hydrogen
and deformation.
This theory encounters a problem, however,
with the fact that the same studies showed that
hydrogen actually strengthens the matrix by
solid solution hardening. It acts in much the
same way as carbon and nitrogen do, as shown
Fig. 35 Stress-strain curve for single crystals of stable
austenitic stainless steel with and without hydrogen.
Source: Ref 36
in Fig. 35 (Ref 36). All these interstitials strain
the lattice and therefore harden in proportion to
their atomic size. Hydrogen, as the smallest of
them, has about half the distorting effect and
half the hardening effect. But, its small size
makes it mobile at ambient temperatures, so it
can diffuse to sites where it can alter mechanical properties.
But, while hydrogen causes dislocation motion and lower work hardening, it does not
weaken austenite, so this theory by itself cannot
account for the role of hydrogen in SCC and, by
inference, in HE in the more general case.
The quandary of hydrogen finally having
been shown to have a clear effect on mechanical
properties but having that not account for either
SCC or HE may be put to rest by the additional
observations of hydrogen’s role as a lattice distorter (Ref 37). While not formalized as a proposed hypothesis for SCC, the role of hydrogen
as a generator of very high stresses has been
pointed out as a factor that cannot be neglected
when evaluating other proposed mechanisms.
Hydrogen has been shown to distort the lattice
in proportion to its concentration. The effect is
not small, accounting for about 1% strain per
0.1% concentration by weight, as shown in
Fig. 36 (Ref 38). At hydrogen levels of over
1000 ppm, which are thought to exist around
growing SCC crack tips, there could therefore
be hydrogen concentration gradients capable of
producing additional tri- or biaxial stresses on
the matrix ahead of the crack tip that may approach the yield stress and account for some or
all of the difference between the normal fracture
toughness KI and the KISCC, that for SCC. This
also precludes the necessity of hypothesizing
hydrogen-induced phase changes, although
54 / Stainless Steels for Design Engineers
Fig. 36 Dilation of austenite due to hydrogen in solution. Source: Ref 38
were they to exist, they would result in the same
lattice expansion. In both cases, the failure
would occur at a region ahead of the crack tip
and beyond the highest hydrogen concentration,
which is what is observed to occur. The growth
of this stress over time with increasing hydrogen-producing corrosion would account for the
observed kinetics, locus, and stress dependence
of SCC.
If nothing else, the main models for SCC and
the experimental results on which they are
based should be reexamined in view of the fact
that the stresses induced by hydrogen are not
negligible and, in fact, may account for much of
the observed SCC behavior of stainless steels.
The next few years may finally see the resolution of the lengthy debate over the causes of
SCC. If it comes, it will be from critical experiments, which can quantitatively differentiate
among the above effects and measure the contribution of each.
Hydrogen Embrittlement
Like SCC, there has been debate about HE
that has produced more heat than light. This in-
volved distinguishing among the same mechanisms, namely:
• Decohesion
• Enhanced local plasticity
• Adsorption embrittlement
• Void coalescence
The identification of the operative mechanism
for HE involves again distinguishing what role
each of the above contributes to HE in a given
situation since all are known to be real metallurgical phenomena.
The main difference between HE and SCC in
stainless steel is that HE is limited to ferrite,
which is hardened by cold work or alloying, and
martensite. Austenite is somewhat diminished
in ductility by hydrogen, but not subject to the
completely brittle, discontinuous cracking of
bcc stainless. The observations that make a
given model plausible as a mechanism for SCC
lack traction for the same materials in HE. It
hard to envision enhanced plasticity involved in
the completely brittle fracture of high-strength
martensitic stainless steels, whereas void coalescence by vacancy creation seems more likely
to account for the observed behavior.
Chapter 4: Corrosion Types / 55
The resolution of mechanism here also must
account for the contribution of hydrogen-induced stress as well as hydrogen effects on mechanical processes, especially since the observed susceptibility to HE is proportional
hardness, therefore to the amount of hydrogen a
given material can hold both in normal interstitial solution and the amount it can trap at lattice
defects (Ref 39), especially the dislocations
within the plastic zone at the crack tip, which
provide enhanced hydrogen solubility where it
can aggravate the applied crack opening with a
wedge effect from hydrogen dilation.
Corrosion Fatigue
Just like SCC, corrosion fatigue causes brittle
failure under a combined environment of corrosion and a tensile stress component. The stress,
however, is cyclic and in a test of stress versus
number of cycles (S vs. N), failure will occur at a
lower N under the corrosive environment. The
cracks are transgranular, and the collaborative effect of corrosion and fatigue is that corrosion accelerates the plastic deformation that accompanies the evolution of extrusions and intrusions.
In corrosion fatigue, an obvious pit corrosion
site may not be necessary because of the combined action of cyclic stresses and the environment. However, an initiation site that is the
weakest link in a combined mechanical and
metallurgical sense will be the initiation point
after which conditions that may not cause SCC
can help propagate fatigue cracking at lower
stresses than would be expected in more benign
environments and in environments that may not
cause SCC or pitting under static loads.
The importance of the environmental interaction is reflected in the sensitivity to frequency of
stress application. High-frequency loading
gives less time for corrosive attack and brings
crack propagation rates down closer to those in
air. In some materials, crack propagation rates
are elevated above those in air at all stress levels, while in others a threshold stress intensity
must be reached before an acceleration is noted.
Some materials show a combination of both.
The first case seems to be merely fatigue assisted by corrosion, while the last two seem to
indicate an SCC–type behavior.
The same uncertainties that cloud our understanding of SCC necessarily disguise the precise mechanism of corrosion fatigue, which
must be viewed as a combination of SCC and
fatigue.
Biocorrosion and Microbiologically
Induced Corrosion
There are many cases for which biological organisms contribute to initiating or enhancing
rates of corrosion. This can occur in natural environments such as ground or seawater as well
as domestic and industrial environments such as
the nuclear and chemical processing industries,
for example. This is called biocorrosion or
MIC, microbiologically induced corrosion.
The bacteria that are known to influence corrosion can be sorted as aerobic bacteria that lie
in aerated water and anaerobic bacteria. Among
the anaerobic bacteria that are known to [16]
(Ref 40) affect stainless steels can be counted:
Desulfibrio and Desulfotomaculum. Both of
these are so-called sulfate-reducing bacteria
(SRB), which means that they promote the reaction:
SO 24− → S2− + 4 O
(Eq 16)
which in turn accelerates the cathode reaction:
2H = + 2e − → 2H
(Eq 17)
Aerobic bacteria flourish under oxygen (Ref
40). Examples are the iron-oxidizing Gallionella and Sphaerotilus, which increase the
anode dissolution reaction:
Fe → Fe 2+ = 2e −
(Eq 18)
by converting the ferrous iron-ion product
(Fe2+) to less soluble ferric (Fe3+). Due to this,
macroscopic so-called tubercules form that can
cause crevice-type shelters where differential
aeration and pit initiation can occur.
Countering MIC with biocides can cause
problems in manganese-containing waters. Oxidizing biocides, such as ozone, chlorine, or peroxide, can cause manganese to be oxidized to
manganese dioxide. The precipitated deposits of
manganese dioxide can accelerate pitting corrosion even in low-chloride waters in which alloys such as 316 would otherwise be safe from
pitting attack.
Biocorrosion is most commonly encountered
in ambient aqueous environments, which are the
environments in which most microorganisms
have evolved to thrive. So, it tends to be a problem for the medium-alloyed steels such as 304
and 316, which are used in these environments.
56 / Stainless Steels for Design Engineers
The more chemically or thermally hostile environments in which higher alloyed grades are
used are also hostile to bioorganisms and thus
minimize the problem.
The development of microbiological consortia
allow anaerobes to flourish under biofilms that
form in an aerated environment. These represent
a differential aeration cell that acts just like a severe crevice. Also, the action of microbes in
raising the corrosion potential is key to understanding why natural seawater is so much more
corrosive than sterile sodium chloride or synthetic seawater solutions. And, macrofouling organisms are important. They create crevices and
sites where microfouling can start early. At the
same time, they are sources of turbulence in
flowing systems, and this turbulence can cause
flow erosion in copper materials, making use of
stainless steels more attractive.
REFERENCES
1. M.P. Ryan et al., Critical Factors in Localized Corrosion, Proc. Electrochem Soc., Vol
150, 2003, p 284–294
2. K. Sieradski and R.C. Newman, J. Electrochem. Soc., Vol 133, 1986, p 1980
3. L. Brewer, Science, Vol 161, 1968, p 115
4. W.J. Tobler and S.Virtanen, Critical Factors
in Localized Corrosion, Proc. Electrochem
Soc., 2003, p 583–594
5. B. Baroux et al., Corros. Sci., Vol 47 (No.
5), 2005, p 1097–1117
6. http://www.alleghenyludlum.com/pages/
products/xq/asp/T.1/qx/productCategory.
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7. K. Kimura et al., High Cr Stainless OCTG
with High Strength and Superior Corrosion
Resistance, JFE Technical Report 7, Jan 2006
8. http://www.outokumpu.com/applications/
documents/start.asp
9. J.E. Truman, Corrosion: Metal/Environment Interaction, Vol 1, Newness-Butterworths, 1976, p 352
10. J.P. Audouard, Stainless Steels, Les Editions
de physique, 1993, p 268
11. H. Mimura et al., Nippon Steel Tech. Report
90, July 2004, p 94–99
12. http://www.alleghenyludlum.com/ludlum/
Documents/AL-6XN_sourcebook.pdf
13. http://www.alleghenyludlum.com/ludlum/
Documents/al610_611.pdf
14. F. Tagashi et al., Kawasaki Technical Report 31, 1994
15. R.C. Newman, Corrosion, Dec 2001,
p 1030–1041
16. N.J. Laycock and R.C. Newman, Corros.
Sci., Vol 39, 1997, p 1771
17. Y. Kobyashi, S.Virtanen, and H. Bohni,
Proc. Electrochem. Soc., 1999, p 533–540
18. Z. Szlarska-Smialowska, Pitting Corrosion
of Metals, NACE, Houston, TX, 1986
19. E.T. Turkdogan, Fundamentals of Steelmaking, Institute of Materials, 1996
20. H.S. Kim and H. Lee, Met. Trans. A,
Vol 32A, June 2001, p 1519
21. M.P. Ryan, D.E. Williams, et al., Nature,
Vol 415, Feb 2002, p 770–777
22. A.J. Grekula et al., Corrosion, 40, 1984,
p 569
23. Stainless Steels Les Editions de physique, 1993
24. N. Suutala and M.Kurkela, Stainless Steel
‘84, Metals Institute, 1985, p 240–247
25. T.Suter, E.G. Webb, H. Bohni, and
R.C. Alkire, J. Electrochem. Soc., Vol 148
(No. 5), 2001, B174–B185
26. M.O. Spiedel, Stainless Steels ‘87, Institute
of Metals, London, 1988, p 247–252
27. Y. Cao, F. Ernst, and G.M. Michal, Acta
Mater., Vol 51, 2003, p 4171.
28. G. Lothongkum et al., Corros. Sci., Vol 48,
2006, p 137–153
29. S. Fujimoto , Sci. Technol. Adv. Mater., Vol
5, 2004, p 195–200
30. J.D. Fritz, J.F. Grubb, B.W. Parks, and
C.P. Stinner, Stainless Steel World, KCI,
P01488, 2001
31. M.O. Spiedel, Met. Trans. A, Vol 12A,
1981, p 779
32. A.J. Sedricks, Corrosion of Stainless Steels,
Wiley, 1979, p 158
33. R.N. Parkins, Br. Corros. J., Vol 14, 1979, p 5
34. S. Tahtinen, H. Hahhinen, and T. Hakkarainen,
Stainless ‘84, Metals Institute, 1985, p 143–148
35. M. Nagumo et al., Met. Trans. A, Vol 32A,
Feb 2001, p 332
36. H. Hanninen et al., Hydrogen Effects on
Materials Behavior, TMS, 2003, p 201–210
37. V.J. Gadgil, Scr. Metal., Vol 28, 1993,
p 1489–1494
38. M. Hoelzel et al., Mater. Sci. Eng. A,
Vol 384, 2004, p 255–261
39. B.G. Pound, Hydrogen Effects on Materials
Behavior, TMS, 2003, p 93–103
40. S.C. Dexter, Microbiologically Influenced
Corrosion, Corrosion: Fundamentals, Testing, and Protection, Vol 13A, ASM Handbooks, 2003, p 398–413
Stainless Steels for Design Engineers
Michael F. McGuire, p 57-68
DOI: 10.1361/ssde2008p057
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
CHAPTER 5
Oxidation
Summary
STAINLESS STEEL, often considered mainly
as a corrosion-resisting material, plays an important role as a heat-resisting material. This is
partly due to its ability to retain strength at higher
temperatures at which many otherwise useful alloying systems, such as aluminum, copper, and
even titanium, soften. Stainless steel retains
strength and has excellent oxidation resistance
from room temperature to nearly 1000 °C, at
which other economical alternatives are lacking.
Introduction
High-temperature oxidation is a form of environmental degradation of metals and alloys that
results from the following chemical reaction in
which metal atoms M react with gaseous oxidants:
M (s) + 0.5 yX 2 ( g) → MX y
(Eq 1)
Due to the high temperatures involved, these
reactions are generally rapid and thus are a concern for high-temperature applications such as
components for power generation. The electronegative gaseous oxidant X could be sulfur,
chlorine, etc., but the discussion here mainly is
limited to oxidation by oxygen or water vapor
(in the latter case, hydrogen would be added as
a product in Eq 1. For a thorough study of oxidation, referred to Ref 1 to 3).
Thermodynamics of Oxidation
As discussed in Chapter 2, Corrosion Theory, a
reaction will be possible when the net free energy
is negative. In Eq 1, the free energy G is decreased
by a lower nobility of the metal (or a higher activity a of a metallic alloying element), a lower temperature T, and a higher partial pressure P of the
oxidizing gas according to:
⎛ aMeX ⎞
y
ΔG = ΔG 0 + RT ln ⎜
⎟
⎜⎝ aM PX0.5 y ⎟⎠
2
(Eq 2)
In the case of alloy oxidation, for which temperatures are high enough to form mixed oxides
or spinels, the activities of the oxide species
also need to be considered.
The standard Gibbs free energy ΔG0 is often
presented in Richardson-Jeffes (Gibbs free
energy-temperature) diagrams such as the one
shown in Fig. 1 (Ref 4).
It is evident from Fig. 1 that the major alloying
element in stainless steels, chromium, forms a
thermodynamically significantly more stable
oxide, Cr2O3, than those of the base alloy iron
(FeO, Fe3O4, and Fe2O3) or the major ternary element nickel (NiO), and to a great extent, the
chromium content determines the oxidation behavior of stainless steels.
The Effect of Chromium. The oxidation of
multicomponent alloys is a complex process
from both thermodynamic and kinetic points of
view. A range of oxides may form with various
degrees of thermodynamic stabilities and stoichiometries (including complex ones with different cations), and there might be degrees of
solubilities of oxides in one another. Kinetics of
their growth is complex because metal solute
diffusion in the metal phases varies, as do metal
and oxygen ion mobilities in the different oxide
phases.
Birks, Meier, and Pettit distinguished between
two basic types of behavior: (a) a noble matrix
58 / Stainless Steels for Design Engineers
Fig. 1 Standard Gibbs free energy of formation of some metal oxides as a function of temperature. Source: Ref 4
metal with less-noble alloying elements and (b)
both matrix element and alloying elements are
nonnoble. The concept of nobility is decided by
the thermodynamic conditions of Eq 2; that is,
an element for which the free energy defined by
Eq 2 is negative is nonnoble. We first discuss the
more common case (b), in which oxidation takes
place under significantly oxidizing conditions,
such as air. In such a situation, it can be seen
from Fig. 1, that the matrix iron is nonnoble and
so are many of the solutes (chromium, molybdenum, aluminum, silicon, manganese, etc.).
A high-temperature Fe-Cr-O phase diagram is
shown in Fig. 2. It can be seen here that Fe2O3
and Cr2O3 are soluble in each other, and that the
spinels Fe1.5Cr1.5O4 (with a solid solubility with
Fe3O4) and FeCr2O4 form. The progressive
change in oxidation behavior as chromium is
added to iron has been described in the literature (Ref 1).
At lower chromium contents and above a minimum temperature, an iron-chromium alloy
would behave as pure iron, where FeO would
form next to the metal, then gradually Fe3O4 and
Fe2O3 would form toward the gas as oxygen potential increases. Isolated pockets of spinel may
form within the FeO layer. The oxidation of iron
proceeds predominantly due to the rapid ionic
diffusion of Fe2+ cations on the FeO layer,
which leads to growth of this layer. If chromium
Chapter 5: Oxidation / 59
Fig. 3 Parabolic rate constants for the growth of several oxides.
Source: Ref 6
Fig. 2 The iron-chromium-oxygen phase diagram at 1300 °C.
Source: Ref 5
in the base alloy is increased, the spinel pockets
increase, and the mobility of Fe2+ decreases. As
chromium content is increased further, a mixedspinel scale is formed. Iron diffusion through the
mixed spinel is significant, and thus the scale is
not yet protective. As chromium content is increased further, an outer layer of Cr2O3 is
formed, and the oxidation behavior becomes
similar to that of chromium. A chromium limit
of roughly 20% is needed to achieve a permanent Cr2O3 scale. This amount decreases if
nickel is added. While the thermodynamic driving force is important, the chromium content of
the alloy can override because the supply of
chromium to the interface becomes dominant.
At chromium contents less than about 16 wt%,
the oxidation rate is influenced by the rate of
supply of chromium from the alloys beneath the
oxide. Above 16%, the supply of chromium is
fast enough that chromium gradients are low
enough that instead transport in the oxide layer
controls the rate. The rate of oxidation then follows a so-called parabolic law (this is explained
in the next section), by which the mass change
per unit area due to oxidation (incorporation of
oxygen) is given by:
( m / A) = k t
2
m
(Eq 3)
Here, m is the added mass, A is the area
exposed to the oxidizing atmosphere, t is the
time exposed, and km is the parabolic rate constant. The subscript “m” is added here to denote
that the reaction is measured as added mass (it
can also be defined for oxide thickness X). As
shown in Fig. 3, the parabolic rate constant for
the oxides of chromium, silicon, and aluminum
are low compared to others, and this is the reason that these elements are used as alloying elements to reduce oxidation rates for alloys in
high-temperature applications.
The solubility of Fe3O4 in the spinel will
eventually result in continued iron oxide formation, as the iron-chromium system is not an
optimal basis for high-temperature oxidation
resistance, although it might be an option from
an economical standpoint compared to other
alloy systems such as superalloys. The high
mobility of both iron and manganese in the
spinel structure is also an important factor.
The major stainless steels used for oxidation
resistance fall into two categories: the ferritic
stainless steels and the austenitic. Table 1 lists
some of the more significant alloys commonly
encountered in applications for which oxidation resistance is paramount. The value of ferritic alloys (such as 409, 439, and 446) is that
they are relatively inexpensive, and that they
have a thermal expansion coefficient that is
closer to that of the oxide than do austenitic alloys (such as 302B, 309, and 310). This gives
them an advantage in cyclic oxidation applications even though their strength at high temperatures does not rival that of austenitic alloys.
The ferritic stainless steels are the most widely
used alloys based on their low cost, which has
60 / Stainless Steels for Design Engineers
Table 1
Oxidation-resisting grades of stainless steel in common use
UNS
Name
S40900
409
0.08
S43935
11Cr-Cb(a)
12SR(a)
439
0.01
0.02
0.07
S44600
S30215
S30415
S38150
S30900
S31000
18Cr-Cb(a)
18SR(a)
4742(a)
446
302B
153MA
253MA
309
310
0.02
0.015
0.08
0.2
0.15
0.04–0.06
0.05–0.10
0.2
0.25
Composition, %
C
N
Cr
Ni
Mn
10.5–11.75
0.5
1
0.015
0.015
0.04
11.35
12
17.0–19.0
0.2
...
0.5
0.25
...
1
...
...
...
0.25
...
0.12–0.18
0.14–0.20
...
...
18
17.3
18
23.0–27.0
17.0–19.0
18.0–19.0
20.0–22.0
22.0–24.0
24.0–26.0
...
0.25
...
0.6
8.0–10.0
9.0–10.0
10.0–12.0
12.0–15.0
19.0–22.0
0.3
0.3
0.7
1.5
2
0.8
0.8
2
2
Si
Ti
6X(C + N) to
1.10
1.3
...
...
0.3
1
0.20 + 4X(C +
N) to 0.75
0.45
0.25
...
0.25
...
...
1
...
2.0–3.0
...
1.0–2.0
...
1.4–2.0
...
0.75
...
1
...
Nb
Other
...
...
0.35
0.6
...
...
1.2 Al
...
0.55
...
...
...
...
...
...
...
...
...
1.7 Al
1.0 Al
...
...
0.04 Ce
0.04 Ce
...
...
1
Note: All compositions include Fe as balance. Single values are maximum, unless otherwise specific (a) Indicates typical analysis
made them the standard alloys for automotive
exhaust systems.
Transient Oxidation
The oxidation of a clean metal surface on exposure to an oxidizing environment will initially lead to all the nonnoble components of the
alloy being oxidized together, forming mixed
oxides having composition similar to the base
alloy. In stainless steels, these initial oxides are
typically Fe-Cr-Ni-Mn mixed oxides. This is
called transient oxidation. As these oxides
thicken, the partial pressure of oxygen at the
scale-metal interface falls until only the most
reactive element present in high concentration
can be oxidized. For stainless steels, this means
that a layer of Cr2O3 is eventually established in
contact with the alloy.
scale-gas surface. Similarly, the electric charge
can be carried by either n-type (electrons) or ptype (electron holes) electronic defects. The
case will be determined by the equilibrium defect structure of the oxide, which depends on
temperature and oxygen partial pressure. In the
case of Cr2O3, chromium cations are the predominantly mobile defects (Fig. 4b) as a result
of a very small degree of deviation from stoichiometry in the cation lattice, that is, Cr2-5O3,
leading to metal deficiency. The contribution of
chromia grain boundary diffusion is large and
probably dominates the process at temperatures
of interest. The defect can be described as an interaction with oxygen, at high oxygen potentials, through Kroger-Vink notations as:
3
3
O ( g) = VCr''' + OOx + 3h •
2 2
2
−3
The Electrochemical Nature of
Oxidation
Once an inner scale of Cr2O3 is formed, as
shown in Fig. 4(a), the oxidizing gas is reduced
at the gas-scale interface, and the chromium is
oxidized at the metal-scale interface. The Cr2O3
scale serves as both electrolyte, through which
ions are transported, and electron lead, through
which electronic defects are transported. In
principle, either or both metal or oxygen ions
can migrate. If oxygen ion mobility dominates,
then the scale would continue to grow at the
oxide-metal interface, whereas if chromium ion
mobility dominates, the oxide will grow at the
3
⎡VCr''' ⎤ p3 = K 2 pO3/ 4 → ⎡VCr''' ⎤ = K12 K 2 Po4
⎣ ⎦
⎣ ⎦
2
2
(Eq 4)
The electron holes that form as charge-compensating defects serve as the “electron lead” in
the electrochemical cell in Fig. 4b. The free energy of Reaction 4 determines the concentration
of mobile defects and thus the diffusion coefficient and electrochemical mobility (Be) of the
cation according to:
⎛ ΔGm ⎞
DCr 3+ = ⎡⎣VCr 3+ ⎤⎦ γΛ 2 υ ⋅ exp ⎜ −
⎝ RT ⎟⎠
⎛ ΔH m ⎞
= ⎡⎣VCr 3+ ⎤⎦ ⋅ cons tan t ⋅ exp ⎜ −
⎝ RT ⎟⎠
(Eq 5)
Chapter 5: Oxidation / 61
Fig. 4 Metal with oxide scale. (a) A protective scale that prevents gas access. (b) Schematic of electrochemical oxidation through a
protective oxide scale that serves as electrolyte and electron lead. The case is for mobile cations
e
BCr
3+ =
3FDCr 3+
RT
(Eq 6)
where zi is the ion charge, F is Faraday’s constant (96,457 C.eq–1). Compared to wustite
(FeO), the equilibrium constant of Eq 4 is quite
low, resulting in a low degree of nonstoichiometry in Cr2–δO3 compared to Fe1–δO (where δ can
be as large as 0.05), and thus the transport of
Cr3+ through its scale is much slower than the
transport of Fe2+ through FeO and thus the difference in parabolic rate constants in Fig. 3.
Kinetics and Oxidation Rates: Wagner’s
Theory
The parabolic oxidation rate was introduced
without explanation in Eq 3. It was first described in terms of oxide defect structure and resulting transport properties by Wagner (Ref 6),
and the theory is explained in most of the monographs on oxidation, such as Chapter 3 in Ref 3
and Chapter 4 in Ref 1. This treatment follows
the derivation in Ref 1. Consider a general case,
as shown in Fig. 4 under the assumptions that
(a) the scale is compact and adherent, (b) electrode reactions are rapid enough to be in equilibrium at the interface and surface, (c) nonstoichiometry is small and uniform throughout the
scale (i.e., defects are in thermal equilibrium
throughout the scale), and (d) double-layer
effects are ignored (i.e., scales are relatively
thick compared to range of space charge effects). As a case study, let us assume that the
mobile ion defect is cations due to metal vacancies in the scale.
The molar flux J (moles/m.s) of a particle i in
an electrolyte subjected to an electrochemical
potential gradient was shown to be:
Ji = − ci Bi
∂ ( μ i + zi Fφ
∂x
)
(Eq 7)
where zi is the ion charge, F is Faraday’s constant (96,457 C/gram equivalent), and φ is the
electric field (V). The electronic or ionic conductivity κ in an electrolyte can be computed
through:
κ = F 2 ∑ zi Bi ci
2
(Eq 8)
The contribution of a given ion specie type or
electron defect type to this conductivity is denoted as the partial conductivity and computed
as:
2
κ i = F 2 zi Bi ci
(Eq 9)
Inserting Eq 9 into Eq 7 yields:
Ji = −
∂ ( μ i + zi Fφ
κi
2
F zi
2
∂x
)
(Eq 10)
62 / Stainless Steels for Design Engineers
Now, if the mobile particles are a single type
of metal cations (e.g., Cr3+) and electrons, then
two fluxes are present:
Jc = −
∂ ( μ c + zc Fφ
κc
2
F zc
∂x
2
)
(Eq 11)
and
Je = −
∂ ( μ e + ze Fφ
κe
2
F ze
∂x
2
)
(E q 12)
dx
k
=
dt X (t )
(Eq 18)
The constant k is the parabolic rate constant.
A mass balance can be written where the flux of
cations for a period of time dt is equated to the
amount of metal being accumulated as cations
inside the scale of thickness dx:
Jc ⋅ dt = CC dx
Electrical neutrality requires that:
Jc Z c + Je Z e = 0
assumed, that is, a linear drop across X, at all
times:
(Eq 13)
And at the oxide-scale/metal interface, the
anode reaction is in equilibrium, that is:
z +
M = M + zce − and therefore,
c
(Eq 19)
By combining Eq 18 and 19 and inserting Eq
16 for the flux, an expression for the parabolic
rate constant is obtained:
μ'
M
κ cκ e
1
k= 2 2
dμ
∫
zc F CC μ '' κ c + κ e M
(Eq 20)
M
μ M = μ c + zc μ e
(Eq 14)
Combining Eq 11 to 14, the potential gradient
is eliminated, and the chemical potential gradients can be replaced by the metal (M) potential
gradient, and the following equation results:
Jc = −
κ cκ e
zc2 F 2 ( κ c + κ e
)
∂μ M
∂x
μ"
(Eq 21)
Since from diffusion theory we know that
Dc = BcRT, and inserting this in Eq 9, one
obtains
κ c = F 2 zc
2
Dc
c
RT c
and inserting this into Eq 21 results in:
μ'
(Eq 16)
M
1 M
k=
D dμ
RT μ∫'' c M
(Eq 22)
M
Now, if the growth of the oxide scale is controlled by the flux of cations:
∂C
dx
∝ JC = − DC C
∂x
dt
μ'
M
1
k= 2 2
κ dμ
zc F CC μ∫M c M
(Eq 15)
To obtain an explicit function for the cation
flux, Eq 15 needs to be integrated after variable
separation, keeping in mind that conductivities
and metal chemical potential may vary within
the scale. Integrating from the gas-scale surface
(x = 0, μM = μ′′M) to the scale-metal interface
(x = X, μM = μ′′M) one obtains.
M
κκ
1
Jc = − 2 2 ∫ c e dμ M
κ
zc F x μ ' c + κ e
If the mobility and thus partial conductivity
of electrons is significantly higher than that of
the ions (a reasonable assumption), then Eq 20
can be simplified as:
(Eq 17)
The concentration drop across the scale is
constant since interface and surface reactions
are at equilibrium. If quasi steady state is
It is often more convenient to express Eq 22
in terms of oxygen potentials rather than the
metal potentials. It was assumed at the onset of
this analysis that the deviation from stoichiometry is small and constant throughout the scale.
Therefore, the oxide potential is constant and:
zc
μ + μ M = μ MO = constant
zC / 2
4 O2
(Eq 23)
Chapter 5: Oxidation / 63
And thus, dμ O = dRT ln PO ∝ − dμ Me. Therefore,
2
2
Eq 22 can be written:
ln PO''
PO''
Dc
dPO
2
P
O
P'
2
2
k ∝ ∫ Dc d ln PO = ∫
2
ln PO'
O2
2
(Eq 24)
2
In the case of Cr2O3, if bulk diffusion is dominating, Eq 4 and 5 inserted into the diffusion
coefficient in Eq 24 for C = Cr3+, results in:
PO''
PO''
2
P 3/ 4
k∝ ∫
dPO = ∫ P −1/ 4 dPO
2
2
PO
P'
P'
2
O2
2
O2
=
( ) ( )
3/ 4
3 ⎡ '' 3/ 4
⎤
PO
− PO'
⎢
⎥⎦
2
2
4⎣
(Eq 25)
Thus, the parabolic rate constant would be predicted to vary with the power of three-quarters of
the external oxygen partial pressure.
Grain boundary diffusion has however been
identified to be important in the case of Cr3+
transport (Ref 8). The observed growth rate of
Cr2O3 polycrystalline films is far too fast to be
accounted for by bulk diffusion of chromium
ions; instead, grain boundary diffusion would
be expected to dominate (Ref 9).
The Volatile Nature of Cr2O3
At high enough temperatures and high
enough oxygen partial pressures, the formation
of a gaseous hexavalent chromium oxide CrO3*
could lead to thinning of the Cr2O3 scale according to the following reaction:
3
Cr2 O3 (s) + O 2 ( g) = 2CrO3 ( g)
2
(Eq 26)
Figure 5 shows the vapor pressure of the superoxide as a function of temperature and partial pressure of O2.
The effect of this reaction on the oxidation kinetics can be described as follows: The thickness change described through the parabolic
rate constant in Eq 18 is corrected for by the
* Hexavalent chromium is now considered a human carcinogen and is rigorously regulated by both the Occupational Safety and Health Administration (OSHA) and the
U.S. Environmental Protection Agency (EPA).
thickness loss due to evaporation, which is described by a first-order reaction kinetics expression with rate constant ke. Thus, the thickness
change becomes:
dx
k
=
−k
dt X (t ) e
(Eq 27)
This results in a so-called paralinear (as opposed to parabolic) rate for the oxide thickening
(Fig. 6), and at a critical oxide scale thickness
X( the rate of thinning due to evaporation equals
the rate of thickening due to oxidation. In Eq
27, this means that dx/dt = 0; consequently, X =
k/ke. While at first this seems to suggest that it
does not affect the oxidation process in that the
rate of oxidation does not increase, the formation and evaporation of chromium oxides results in greater chromium consumption in the
alloy compared to what would be the case if
evaporation did not occur. As a result of evaporation losses, stainless steels that depend on a
protective chromium oxide layer are limited in
use to temperatures up to 900 to 1000 °C.
The presence of water vapor promotes the
formation of even more volatile oxyhydroxides
(e.g., CrO2(OH)2 ) (Ref 11, 12).
Spalling and Cracking of the Scale
At elevated temperatures or during temperature cycling, there are multiple ways in which
stresses can develop that may crack and blister
the scale, rendering it nonprotective. The different causes of stress generation are described in
Chapter 5 in Ref 1. So-called growth stresses
arise due to changes caused by the oxidation
process itself. These include differences in lattice mismatch, alloy depletion in the metal,
point-defect gradients in scales containing oxides such as FeO, with large deviation from stoichiometry, recrystallization, and volume differences between the oxide and metal. The last is
perhaps the most commonly mentioned and is
characterized by the Pilling-Bedworth ratio, abbreviated as PBR (Ref 13).
PBR =
VOxide
VmOxide
= Metal
VMetal Vm * ν
(Eq 28)
Here, the subscript m stands for molar volume,
and υ is the number of metal atoms needed to
form a stoichiometric unit of the oxide (in the
64 / Stainless Steels for Design Engineers
Fig. 5 Chromium-oxygen system species volatility as a function of temperature and oxygen pressure. Source: Ref 10
case of Cr2O3, υ is 2, and in the case of FeO, it
is 1). When PBR is greater than 1, then the
oxide is expected to be in compression and is
likely to be protective, whereas if it is less than
1, the oxide is in tension and thus nonprotective.
There are, however, many exceptions to this,
partly because the stress state often depends
more on the mechanisms and conditions of the
oxidation process rather than the properties of
metal and oxides.
Thermal, stresses are caused by differences in
thermal expansion between the oxide and metal,
and the stresses generated in oxide scales can be
estimated:
σ Oxide =
)
− EOxide ( α Oxide − α Metal ΔT
⎛
⎞
t E
1 − ν p ⎜ 1 + 2 Oxide Oxide ⎟
t Metal E Metal ⎠
⎝
(
)
(Eq 29)
The equation is written for a case shown
schematically in Fig. 7, where both sides on a
metal undergo oxidation. Here, σ is the stress,
υp is Poisson’s ratio (it has been assumed that
there is no mismatch), α is the coefficient of
thermal expansion, E is the modulus of elastic-
Fig. 6 Schematic of paralinear oxidation as a result of evaporation of chromium superoxide
ity, t is thickness, and ΔT is the temperature
change. In general, α is larger for the metal than
the oxide; thus, during cooling the stresses are
expected to be compressive and during heating
tensile. Thermal stresses can cause spalling of
the protective oxide layer, and it is most severe
under cyclic conditions.
Effect of Silicon, Aluminum, and Molybdenum. Due to concerns about the cost of
chromium and its (former) classification as a
Chapter 5: Oxidation / 65
tox
tm
tox
Oxide
Metal
Oxide
Fig. 7 Schematic of a cross section of oxidized sample indi-
cating dimensions in Eq 29 for predicting thermal
stresses
strategic material, there were efforts to try to
substitute less-expensive elements such as aluminum and silicon that also are known to form
protective layers (Ref 14), even though they
have significant metallurgical and mechanical
drawbacks.
Silicon additions of 1.5 wt% or more have the
effect of forming a continuous amorphous subsurface layer in iron-silicon and Fe-Cr-Si alloys
that is relatively impervious to transport of ions.
The mechanism for the evolution of such a layer
is as follows: (1) The more readily available
iron or chromium first forms a surface layer,
and this causes an enrichment of silicon at the
oxide-metal interface. (2) As sufficient silicon is
enriched, the SiO2 layer is formed. It has been
reported that alloys with chromium content as
low as 6 wt% and silicon content of 1.5 wt%
perform in terms of oxidation as well as commercial stainless steels. Also, an addition of 4
wt% Si to a Fe14wt%Cr14wt%Ni alloy resulted
in a 200-fold reduction in weight gain at 900 °C
However, this SiO2 layer seems to promote
oxide spalling, especially in cyclic service.
Aluminum forms a very stable thin outer
layer of Al2O3 that initially reduces the oxidation rate. Alumina is among the most stable and
defect-free oxides, giving it an extremely low
diffusion rate. In consequence, if sufficient aluminum is present to maintain the protective alumina scale, the aluminum-bearing alloys provide the greatest oxidation resistance attainable
in engineering alloys. If, however, the aluminum content is not sufficient to force alumina
scale formation at the scale-metal interface, an
internal Cr2O3 layer forms and thickens. Eventually, the surface aluminum-oxide layer flakes
off. Also, due to the low oxygen potential
needed to form Al2O3, internal oxidation may
result below the metal-scale interface in alloys
in which formation of a continuous alumina
scale film does not occur.
Molybdenum is suggested to strain the lattice due to its larger size and consequently increase the rate of bulk diffusion of elements
(Ref 11), which can enhance the rate of initial
Cr2O3 formation. Molybdenum is usually considered detrimental for oxidation resistance.
Molybdenum normally forms MoO2 oxide, but
this can oxidize further to form the low-melting and volatile MoO3. If the MoO3 evaporates, there is little problem, but if its
volatilization is inhibited by low atmosphere
circulation, liquid MoO3 can accumulate and
dissolve the protective Cr2O3 scale, leading to
catastrophic oxidation.
Effect of Rare Earth Additions. Cerium, lanthanum, and yttrium additions are known to improve oxidation resistance of high temperature
nickel- and iron-based alloys (Ref 6). Rare earth
additions have been suggested to have a multitude of beneficial effects, such as reducing the
growth kinetics of Cr2O3 scales, stabilizing
Cr2O3 scales at lower chromium levels, increasing adhesion, and preventing spalling of the
oxide scale during thermal cycling. The explanation for any of this does not seem clear, but
some hypotheses have been suggested. The effect on the growth kinetics could be because the
reactive element ions collect at grain boundaries
and block fast path diffusion.
The improved adherence could be because
these elements getter tramp elements such as
sulfur and suppress void formation at the interface. Furthermore, they might form so-called
oxide pegs at the interface (Ref 16). The precise role of the rare earth additions to Cr2O3
oxide protection and the mechanism by which
they are incorporated into the scale during the
surface treatment processes remain unknown.
An understanding of these fundamental issues
would help to develop optimum alloy
chemistries for selected high-temperature and pressure applications and to further develop the
surface infusion process. The lack of fundamental understanding of how rare earths improve oxidation resistance has not stopped the
development of several alloys that benefit from
the effect.
66 / Stainless Steels for Design Engineers
Oxidation Under Less-Oxidizing
Atmospheres
quasi-steady-state situation as shown in Fig. 9,
the flux can be written:
The two types of alloy oxidation behaviors,
(a) a noble matrix metal with less-noble alloying elements and (b) both matrix element and
alloying elements are nonnoble, were mentioned. When designing against oxidizing environments, case b is perhaps the most relevant,
and most of the discussion has been devoted to
this. However, during annealing for microstructural control, steels are exposed to furnace gases
at high temperatures that have relatively low
oxygen or steam contents, for which case a will
apply, that is the atmosphere does not cause iron
(or nickel) to oxidize but chromium (and aluminum, silicon, molybdenum, etc.) does.
For simplicity, assume a binary system A-B
of “noble” iron and “reactive” chromium. In
this case, depending on the concentration of
the reactive element and atmosphere, the
oxide of the reactive element can, in principle,
form either on the surface (as has been discussed so far) or internally as discrete oxide
particles in a metal matrix through oxygen
diffusion into the metal. Both cases are shown
schematically in Fig. 8. Let us discuss the
conditions that promote one or the other of
these by starting with a situation in which (a)
no surface oxide exists and (b) the oxygen atmosphere is such that the solubility of oxygen
within a distance X is enough to thermodynamically render Cr2O3 stable according to Eq 2
but none of the iron oxides. The derivation is
done in terms of both a generic system A-B
causing an oxide BOν and for iron-chromium
causing Cr2O3. Assume for a start that the DO
>> DCr, and thus while oxygen diffuses into
the alloy, chromium does not counterdiffuse.
The flux of oxygen inward into the metal is
then the cause of increased mass. Assuming a
JO =
Fig. 8 Schematic of two cases in a less-oxidizing atmosphere.
(a) Adsorption of oxygen leading to internal oxidation
and (b) external oxidation as the B element migrates.
N X − N OS
NS
dm
= − DO O
= DO O
dt
Vm X
Vm X
(Eq 30)
Within the depth X, the oxygen solubility is
such that Eq 2 is negative enough that Cr2O3
forms. Beyond X, it is not. At the distance X, the
oxygen concentration is negligibly low compared to the surface composition, which is in
equilibrium with the gas phase. The molar volume Vm is used to obtain the flux in units of
moles per square meter. Within the layer 0 < x <
X, all the chromium is assumed to be oxidized;
therefore, the accumulated mass due to oxygen
addition:
o 3
N B0 υX N Cr 2 X
m=
=
Vm
Vm
(Eq 31)
Differentiating Eq 31 with time, equating to
Eq 3, and separating variables results in:
XdX =
N OS DO
dt =
O
υN B
N OS DO
3 0
N
2 Cr
(Eq 32)
Integrating Eq 32 from x = 0 to x = X results in
an expression for the internal oxidation depth X:
⎛ 2N S D ⎞
X = ⎜ O 0 O t⎟
⎝ υN B ⎠
1/ 2
⎛
⎞
⎜ 2 N OS DO ⎟
=⎜
t⎟
0
⎟
⎜ 3 N Cr
⎠
⎝ 2
1/ 2
(Eq 33)
Fig. 9 Quasi-steady-state approximation of the moving
boundary problem of internal oxidation. Counterdiffusion of B is assumed to be negligible
Chapter 5: Oxidation / 67
This expression predicts a parabolic dependence of X with time, just as Eq 3 did for the external oxidation. For a more rigorous derivation
(without assuming quasi steady state), refer to
Appendix B in Ref 1. Now, let us see what
causes this to transition into an external scale. It
was assumed in the derivation of Eq 33 that
counterdiffusion of chromium does not occur.
When considering Eq 33, is clear that the rate of
penetration of the internal oxidation front will
decrease with (a) increasing NB0, (b) decreasing
NOS, and (c) decreasing DO. If DCr was not negligible, there would be a gradual change in oxidation morphology if the ratio (NB0DB)/(NOSDO).
Gradually, if the ratio were increased there
would be a slowing of the inward penetration of
the internal oxide front and an enrichment of
BOυ in the internally oxidized zone. Wagner developed a model (Ref 16), based on that at some
point, when the volume fraction of BOυ versus
volume metal in the internally oxidized zone
reaches a critical value g*, there is a transition
from internal to external oxidation; specifically,
this happens when:
1/ 2
⎡ πg* S DOVM ⎤
N >⎢
NO
⎥
DBVOx ⎦
⎣ 2ν
0
B
Fig. 10 Temperature dependence of metal dusting of iron.
Source: Ref 18
trogen absorption, but if the oxygen is depleted
before all surfaces are oxidized, the remaining
material can be rapidly nitrided by the residual,
essentially pure, nitrogen atmosphere.
REFERENCES
(Eq 34)
If more than one reactive elements were present (such as is the case in stainless steels in
which aluminum, silicon, molybdenum, niobium, etc., may be present), this will decrease
the inward flux of oxygen, and thus the transition to external may occur at a lower solute
(CCr0) concentration than what is predicted by
Eq 34.
Metal Dusting. Under reducing conditions,
in products of combustion atmospheres, oxidation and carburization may occur simultaneously and at a higher rate than exhibited in pure
oxidation. Under even more reducing conditions, the condition called metal dusting may
occur. Metal dusting is often characterized by
the generation of large, smooth pits that look as
if metal had been scooped from the surface. The
underlying phenomenon is the formation of
metal carbides, which manifests itself as the
breakup of bulk metal to metal powder. This occurs at temperatures at which the carbide is
most stable (Fig. 10).
During oxidation in air, if large surface area is
present under conditions of restricted air supply,
oxygen can be depleted to the point that oxidation essentially ceases. Oxide films inhibit ni-
1. N. Birks, G.H. Meier, and F.S. Pettit, Introduction to the High-Temperature Oxidation
of Metals, 2nd ed., 2006, Cambridge University Press, New York
2. P. Kofstad, High Temperature Corrosion,
1988, Elsevier Applied Science, London
3. K. Hauffe, Oxidation of Metals, Plenum
Press, New York, 1965
4. F.D. Richardson and J.H.E. Jeffes, J. Iron
Steel Inst., Vol 160, 1948, p 261
5. C. Wagner and K. Grünewald, Z. Phys.
Chem., Vol 40B, 1938, p 455
6. J.H. Park, W.E. King, N.L. Peterson, and
S.J. Rothman, The Effect of Reactive Element on Self-Diffusion in Cr2O3, Norman
L. Peterson Memorial Symposium, Oxidation of Metals and Associated Mass Transport, edited by M.A. Dayananda, S.I. Rothman, and W.E. King, AIME, Warrendale,
PA, 1998, p 103–107
7. C. Wagner and K. Grünewald, Z. Phys.
Chem., Vol 40B, 1938, p 455
8. D. Caplan and G.I. Sproule, Effect of Oxide
Grain Structure on the High Temperature
Oxidation of Cr, Oxid. Met., Vol 9, 1975,
p 459–472
9. B.B. Ebbinghaus, Combust. Flame, Vol 93,
1993, p 119–137
68 / Stainless Steels for Design Engineers
10. K. Hilpert et al., JECS, Vol 143/11, 1996, p
3642–3647
11. N.B. Pilling and R.E. Bedworth, J. Inst.
Met., Vol 29, 1923, p 529
12. J.K. Tien and J.M. Davidson, Oxide Spallation Mechanisms, Stress Effects and the Oxidation of Metals, ed. J.V. Cathcart, AIME,
New York, 1975, p 200
13. J. Rawers, Understanding the Oxidation
Protection of Fe-Cr-Si Alloys, Norman L.
Peterson Memorial Symposium, Oxidation
of Metals and Associated Mass Transport,
edited by M.A. Dayananda, S.I. Rothman,
and W.E. King, AIME, Warrendale, PA,
1998, p 323–340
14. E.J. Felten, J. Electrochem. Soc., Vol 108,
1961, p 490
15. C. Wagner, J. Electrochem. Soc., Vol 103,
1956, p 571
16. C.M. Chun, J.D. Mumford, and T.A. Ramanarayanan, J. Electrochem. Soc., Vol 147,
2000, p 3680
Stainless Steels for Design Engineers
Michael F. McGuire, p 69-90
DOI: 10.1361/ssde2008p069
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
CHAPTER 6
Austenitic Stainless Steels
Summary
AUSTENITIC STAINLESS STEELS are the
most common and familiar types of stainless
steel. They are most easily recognized as nonmagnetic. They are extremely formable and
weldable, and they can be successfully used
from cryogenic temperatures to the red-hot temperatures of furnaces and jet engines. They contain between about 16 and 25% chromium, and
they can also contain nitrogen in solution, both
of which contribute to their high corrosion resistance. Were it not for the cost of the nickel
that helps stabilize their austenitic structure,
these alloys would be used even more widely.
because their greater thermal expansion coefficient tends to cause the protective oxide
coating to spall.
2. They can experience stress corrosion cracking (SCC) if used in an environment to which
they have insufficient corrosion resistance.
3. The fatigue endurance limit is only about
30% of the tensile strength (vs. ~50 to 60%
for ferritic stainless steels). This, combined
with their high thermal expansion coefficients, makes them especially susceptible to
thermal fatigue.
However, the risks of these limitations can be
avoidable by taking proper precautions.
Introduction
Alloy Families in Perspective
Austenitic stainless steels have many advantages from a metallurgical point of view. They
can be made soft enough (i.e., with a yield
strength about 200 MPa) to be easily formed by
the same tools that work with carbon steel, but
they can also be made incredibly strong by cold
work, up to yield strengths of over 2000 MPa
(290 ksi). Their austenitic (fcc, face-centered
cubic) structure is very tough and ductile down
to absolute zero. They also do not lose their
strength at elevated temperatures as rapidly as
ferritic (bcc, body-centered cubic) iron base alloys. The least corrosion-resistant versions can
withstand the normal corrosive attack of the
everyday environment that people experience,
while the most corrosion-resistant grades can
even withstand boiling seawater.
If these alloys were to have any relative
weaknesses, they would be:
The fundamental criterion in the selection of
a stainless steel is generally that it can survive
with virtually no corrosion in the environment
in which it is to be used. Good engineering
practice sometimes requires that materials be
selected for sufficient, but finite, service life.
This is especially true for high-temperature
service, for which creep and oxidation lead to
limited life for all materials. The choice among
the stainless steels that can be used in that environment is then based on the alloy from which
the component can be produced at the lowest
cost, including maintenance, over the intended
service life. The ferritic stainless steels are less
expensive for the same corrosion resistance but
sometimes are found lacking because of:
• Lack of toughness, as is the case at subambient temperatures or in thicknesses greater
than about 1.5 mm
• Lack of great ductility, specifically if more
than about 30% elongation is needed
1. Austenitic stainless steels are less resistant
to cyclic oxidation than are ferritic grades
70 / Stainless Steels for Design Engineers
• Susceptibility to high-temperature embrittling phases when moderately alloyed
The less-expensive martensitic grades are
used instead of austenitic when high strength
and hardness are better achieved by heat treating rather than by cold work, and mechanical
properties are more important than corrosion resistance. This is also the case for the more expensive PH grades, which can achieve corrosion
resistance only matching the least corrosion resistant of the austenitic alloys.
Duplex grades match austenitic grades in corrosion resistance and have higher strength in the
annealed condition but present the designer
with challenges with regard to embrittling
phases that can form with prolonged exposure
to elevated temperatures and only moderate
ductility like the ferritic alloys.
So, the austenitic grades are the most commonly used grades of stainless mainly because,
in many instances, they provide very predictable
levels of corrosion resistance with excellent mechanical properties. Using them wisely can save
the design engineer significant costs in his or her
product. They are a user-friendly metal alloy
with life-cycle cost of fully manufactured products lower than many other materials.
The austenitic alloys can have compositions
anywhere in the portion of the Delong diagram
labeled austenite shown in Fig. 1 (Ref 1). This
diagram was designed to show which phases are
present in alloys in the as-solidified condition,
such as found in welds. Thus it also applies to
castings and continuously cast products. As a
practical matter of castability, the composition
of most commercial alloys falls along the zone
of several percent ferrite as cast. The salient feature of austenitic alloys is that as chromium and
molybdenum are increased to increase specific
properties, usually corrosion resistance, nickel
or other austenite stabilizers must be added if
the austenitic structure is to be preserved.
The traditional way of displaying the
austenitic stainless steels is to present 302 as a
base. Figure 2 shows one such diagram. Diagrams such as these treat alloys as an evolutionary family tree and subtly mislead. Many alloys
were pushed toward obsolescence because of
advances in processing. For instance, 321 was
developed as an alloy in which the detrimental
effects of carbon were negated by addition of titanium. The widespread adoption of the argon
oxygen decarburization (AOD) in the 1970s
made this alloy unnecessary, except for special
circumstances, since carbon could be cheaply
Fig. 1 Schaeffler-Delong stainless steels constitution diagram. Adapted from Ref 1, 2
Chapter 6: Austenitic Stainless Steels / 71
Fig. 2 The austenitic stainless family. Source: Ref 3
removed routinely. Likewise, 302 gave way to
the lower-carbon 304, for which the even lowercarbon 304L is commonly substituted and dually certified to qualify as either grade. While
low carbon prevents sensitization, stabilized
grades may still be preferred for special applications such as type 321 in aerospace and type
347 in refinery service. Similar inertia keeps the
higher-nickel 300 series as the de facto standard
when the more cost-efficient high-manganese
200 series is the logical basic grade. The relevant types of austenitic alloys can nonetheless
be rationalized with this diagram.
As chromium is added, oxidation resistance
and corrosion resistance increase. Because
nickel equivalents (manganese, nitrogen, carbon,
etc.) must also be added in matching amounts,
austenite stability is also increased. If molybdenum, a chromium equivalent, is added, corrosion resistance but not oxidation resistance is
enhanced. And, if nitrogen is the austenite stabilizer added to balance increases chromium or
molybdenum, then corrosion resistance is also
increased. With small exceptions, that is the
rationale of austenitic grade design. Silicon
is used as an alloy to promote oxidation resistance and resistance to corrosion by oxidizing
acids. Copper is used to promote resistance to
sulfuric acid. Rare earths make a more stably
oxidation-resisting scale. Niobium increases
72 / Stainless Steels for Design Engineers
Lean Alloys
creep resistance. Sulfur and selenium increase
machinability.
In this chapter, austenitic alloys are classified
into three groups:
• Lean alloys, such as 201 and 301, are generally used when high strength or high formability is the main objective since the lower,
yet tailorable, austenite stability of these alloys gives a great range of work-hardening
rates and great ductility. Richer alloys, such
as 305, with minimal work hardening are the
high-alloy, lowest work-hardening rate
grade for this purpose. The general-purpose
alloy 304 is within this group.
• Chromium nickel alloys when the objective
is high temperature oxidation resistance.
This can be enhanced by silicon and rare
earths. If the application requires high-temperature strength, carbon, nitrogen, niobium,
and molybdenum can be added. 302B, 309,
310, 347, and various proprietary alloys are
found in this group.
• Chromium, molybdenum, nickel, and nitrogen alloys when corrosion resistance is the
main objective. Alloys such as silicon and
copper are added for resistance to specific
environments. This group includes 316L,
317L, 904L, and many proprietary grades.
Lean austenitic alloys constitute the largest
portion of all stainless steel produced. These are
principally 201, 301, and 304. Alloys with less
than 20% chromium and 14% nickel fall into
this unofficial category. Since they are stainless,
it is generally taken for granted that these alloys
will not corrode, and these alloys have sufficient corrosion resistance to be used in any indoor or outdoor environment, excluding coastal.
These grades are easily weldable and formable
and can be given many attractive and useful surface finishes, so they are very much generalpurpose alloys. Table 1 lists some typical compositions of the most commonly used lean
austenitic alloys. These typical compositions
vary with end use, raw material cost factors, and
the preference of a given manufacturer. The
compositions of standard alloys are often finetuned to the intended end use. In this table, the
word drawing indicates higher nickel for lower
work hardening, while tubing indicates alloys
with higher sulfur to facilitate gas tungsten arc
welding (GTAW) penetration. Tensile indicates
lower alloy levels to increase the work-hardening rate for material that is intended to be used
in the cold-worked, high-strength condition.
316L is included in its most common tubing end
use chemistry even though it is a corrosion-resisting alloy because it is so pervasively used as
a service center sheet item.
The main difference among the lean
austenitic alloys lies in their work-hardening
rate: the leaner the alloy, the lower the austenite
stability. As unstable alloys are deformed, they
transform from austenite to the much harder
martensite. This increases the work-hardening
rate and enhances ductility since it delays the
onset of necking since greater localized
Wrought alloys generally have cast counterparts that differ primarily in silicon content.
Versions that require enhanced machinability
have a high content of controlled inclusions,
sulfides, or oxysulfides, which improve machinability at the expense of corrosion resistance.
Carbon is kept below 0.03% and designated an
L grade when prolonged heating due to multipass welding of heavy section (greater than
about 2 mm) or when welds requiring a postweld stress relief are anticipated.
Table 1 Typical compositions of the most commonly used lean austenitic alloys
Alloy
Designation
C
N
Cr
Ni
Mo
Mn
Si
Other
Other
Other
201
201 drawing
201LN
301 tensile
301 drawing
303
304
304 drawing
304 extra drawing
304L tubing
305
321
316L
S20100
S220100
S20153
S30100
S30100
S30300
S30400
S30400
S30400
S30403
S30500
S32100
S31603
0.08
0.08
0.02
0.08
0.08
...
0.05
0.05
0.06
0.02
0.05
0.05
0.02
0.07
0.07
0.13
0.4
0.04
...
0.05
0.04
0.04
0.09
0.02
0.01
0.0
16.3
16.9
16.3
16.6
17.4
...
18.3
18.4
18.3
18.3
18.8
17.7
16.4
4.5
5.4
4.5
6.8
7.4
...
8.1
8.6
9.1
8.1
12.1
9.1
10.5
0.2
0.02
0.2
0.2
0.02
...
0.3
0.3
0.3
0.3
0.2
0.03
2.1
7.1
7.1
7.1
1.0
1.7
...
1.8
1.8
1.8
1.8
0.8
1.0
1.8
0.45
0.5
0.45
0.45
0.45
...
0.45
0.45
0.45
0.45
0.60
0.45
0.50
0.001 S
0.001 S
0.001 S
0.001 S
0.007 S
...
0.001 S
0.001 S
0.001 S
0.013 S
0.001 S
0.001 S
0.010 S
0.03 P
0.30 P
0.03 P
0.03 P
0.03 P
...
0.03 P
0.03 P
0.030 P
0.030 P
0.02 P
0.03 P
0.03 P
0.2 Cu
0.6 Cu
0.5 Cu
0.3 Cu
0.6 Cu
...
0.3 Cu
0.3 Cu
0.4 Cu
0.4 Ci
0.2 Cu
0.4 Ti
0.4 Cu
Chapter 6: Austenitic Stainless Steels / 73
deformation is more than offset by greater localized strain hardening.
These grades are best viewed as a continuum
with a lower boundary at 16%Cr-6%Ni and an
upper boundary at 19%Cr-12%Ni. This represents the range from minimum to maximum
austenite stability. Since that is the main distinction within this grade family, let us examine its
basis.
Martensite and Austenite. Stability. The
formation of martensite at room temperature
may be thermodynamically possible, but the
driving force for its formation may be insufficient for it to form spontaneously. However,
since martensite forms from unstable austenite
by a diffusionless shear mechanism, it can occur
if that shear is provided mechanically by external forces. This happens during deformation,
and the degree to which it occurs varies with
composition according to (Ref 4):
such as occurs when they are sensitized or when
solute segregation occurs, as from welding, then
the equation applies on a microscopic scale.
Sensitized zones (i.e., the regions near grain
boundaries where chromium carbides have precipitated) will have a much higher tendency to
transform to martensite. Figures 3(a) and (b)
show the changes in phase structure as a function of composition over ranges that encompass
these alloys.
Martensite can be present in two different
forms. The α′-form is the bcc magnetic form,
while ε is a nonmagnetic, hcp (hexagonal closepacked) version. The formation of ε versus α′ is
related to the stacking fault energy of the alloy,
which is given by (Ref 6):
Md30 (°C) = 551 – 462(%C + %N)
– 9.2(%Si) – 8.1(%Mn) – 13.7(%Cr)
– 29(%Ni + Cu) – 18.5(%Mo)
– 68(%Nb) – 1.42 (GS – 8)
(Eq 1)
Y300SF (mJ m-2) = Y0SF + 1.59Ni – 1.34Mn
+ 0.06Mn2 – 1.75Cr + 0.01Cr2
+ 15.21Mo – 5.59Si
–60.69(C + 1.2N)1/2
+ 26.27(C + 1.2N)
(Cr + Mn + Mo)1/2
+ 0.61[Ni•(Cr + Mn)]1/2
(Eq 2)
This is the temperature at which 50% of the
austenite transforms to martensite with 30%
true strain (Ref 5). It should be noted that even
elements that are chromium equivalents in promoting ferrite are austenite stabilizers in that
they impede martensite formation. This temperature is the common index of austenite stability.
This regression analysis was generated for homogeneous alloys. If alloys are inhomogeneous,
Epsilon martensite formation is favored in
alloys of lower stacking fault energy. The fcc
structures deform by slip between (111)
planes. Viewed from these planes, the structure
is a series of ABCABC atom arrangements.
Slip between planes can result in an
ABCA/CAB structure. This so-called stacking
fault generates an hcp structure. With lower
stacking fault energies, these are more readily
Fig. 3 (a) Iron-chromium phase diagram at 8% nickel; (b) iron-nickel phase diagram at 18% chromium
74 / Stainless Steels for Design Engineers
Fig. 4 Variation of martensite formation with temperature and true strain for 304. Source: Ref 7
formed, and ε predominates. The stacking fault
can also be viewed as two partial dislocations
with the material between them faulted. These
partial dislocations, when generated in abundance, cannot readily slip past one another
and thus pile up, increasing work-hardening
rates.
As in carbon and alloy steels, the martensite
transformation can take place simply by cooling, but in the lean austenitic alloys the temperatures are well below ambient. The more stable
alloys do not transform even with cryogenic
treatment. Figure 4 shows the variation of
martensite formation with temperature and true
strain for 304. Martensite formed in these alloys
is quite stable and does not revert until heated
well above the temperatures (Fig. 5) at which it
was formed. The carbon levels of austenitic
stainless steels are always relatively low, so
strain-induced martensite is self-tempering and
not brittle.
Martensite has been found to form in unstable
austenite due to the electrochemically induced
supersaturation by hydrogen (Ref 9). Under
conditions of cathodic charging, superficial layers were found to transform to ε under conditions of intense hydrostatic compression. During subsequent outgassing, α′ was found to
form due to reversals in the stress state. Martensite thus formed is, of course, susceptible to
hydrogen embrittlement.
Mechanical Properties. The tensile properties in the annealed state not surprisingly relate
well to composition. The 0.2% yield strength
and tensile strength, respectively, are reported
(Ref 10) to follow the equations:
YS( MPa) = 15.4[4.4 + 23(%C) + 32(%N)
+ 0.24(%Cr ) + 0.94(%Mo)
+ 1.3(%Si ) + 1.2(%V)
+ 0.29(%W ) + 2.6(%Nb)
+ 1.7(%Ti ) + 0.82(%Al )
+ 0.16(%Ferrite )
+ 0.46(d −1/1/ 2 )
(Eq 3)
TS ( MPa) = 15.4[29 + 35(%C) + 55(%N)
+ 2.4(%Si ) + 0.11((%Ni ) + 1.2(%Mo)
+ 5.0(%Nb) + 3.0(%Ti ) + 1.2(%Al )
+ 0.14(%Ferrite ) + 0.82(d −1/ 2 )
(Eq 4)
In each case, d is the grain diameter in millimeters.
Another researcher (Ref 11) gave the relationships as:
YS ( MPa ) = 120 + 210 N + 0.02 + 2 Mn + 2Cr
+ 14 Mo + 10 Cu + (6.15 − 0.054δ)δ
+ (7 + 35(N + 0.2))d −1/ 2
(Eq 5)
TS = 470 + 600( N + 0.2)
+ 14 Mo + 1.5δ + 8d −1/ 2
(Eq 6)
Again, d is grain diameter in millimeters, and
δ is percent ferrite. The claimed accuracy for
the latter set of equations is 20 MPa and is said
Chapter 6: Austenitic Stainless Steels / 75
Fig. 6 Variation of impact strength with temperature for (a)
austenitic, (b) duplex, and (c) ferritic stainless steels
Fig. 5 Reversion of martensite formed by cold work. Source:
Ref 8
to apply to both austenitics and duplex stainless
steels, but clearly the tensile strength relationship must break down for leaner alloys, such as
301, in which tensile strength increases with decreasing alloy content because of the effect of
increasing alloying causing less transformation
to martensite, which inarguably produces higher
tensile strengths in austenitic stainless steels.
Equation 3 must also be favored over Eq 5 in
that it accounts for carbon explicitly.
One other hardening mechanism is possible
in austenitic stainless steels, and that is precipitation hardening. Most precipitation-hardening
stainless steels are unstable austenite, which is
transformed to martensite before the precipitation hardening takes place. One commercial
alloy, A-286, is entirely austenitic and employs
the precipitation within the austenite matrix of
Ni3 (titanium, aluminum) for strengthening.
This is dealt with in a separate section.
Austenitic stainless steels do not have a clear
yield point but can begin to deform at as little as
40% of the yield strength. As a rule of thumb,
behavior at less than half the yield strength is
considered fully elastic and stresses below twothirds of the yield strength produce negligible
plastic deformation. This quasi-elastic behavior
is a consequence of the many active slip systems in the fcc structure. Even highly coldworked material exhibits this phenomenon, although stress-relieving cold-worked material
will cause dislocations to “lock in place” and
form more stable dislocation arrays that break
loose at a higher and distinct yield point.
The tensile properties of austenitic stainless
steels with unstable austenite, that is, those with
Md30 temperatures (Eq 1) near room temperature, are very strain rate dependent. This is simply due to the influence of adiabatic heating
during testing increasing the stability of the
austenite. Tests run under constant temperature
conditions, either by slow strain rates or use of
heat sinks, produce lower tensile strengths.
Thus, reported tensile strengths should not be
taken as an absolute value but a result that can
be significantly changed by changes in testing
procedure, even with accepted norms and
standards.
Highly cold-worked austenitic stainless
steels are often used for their robust mechanical
properties. Few metallic materials can match
the very high strengths they can achieve. Very
lean 301 can be cold worked to yield strengths
on the order of 2000 MPa because of its unstable austenite transforming to martensite. When
cold worked to lower degrees, it can provide
very high strength while keeping impressive
ductility.
Austenitic stainless steels have exceptional
toughness. The ambient temperature impact
strength of austenitic stainless steels is quite
high. This is not surprising in view of their high
tensile strengths and high elongations. What is
most remarkable is the absence of a transition
temperature, which characterizes ferritic and
martensitic materials. Figure 6 shows impact
strength of the various stainless steel types versus temperature. This again is due to the multiplicity of slip systems in the fcc structure and
the fact that they do not require thermal activation. This makes the austenitic stainless steels,
especially the 200 series, the optimal cryogenic
76 / Stainless Steels for Design Engineers
material, surpassing the 9% nickel martensitic
steels in cost, toughness, and, of course, corrosion resistance.
Precipitation of Carbides and Nitrides.
Carbon is normally considered as an undesirable impurity in austenitic stainless steel. While
it stabilizes the austenite structure, it has a great
thermodynamic affinity for chromium. Because
of this affinity, chromium carbides, M23C6, form
whenever carbon reaches levels of supersatura-
Fig. 7 Carbon solubility in 18–10 austenitic stainless. Source:
Ref 13
tion in austenite, and diffusion rates are sufficient for carbon and chromium to segregate into
precipitates. The solubility of carbon in austenite is over 0.4% at solidification but decreases
greatly with decreasing temperature. The solubility is given by (Ref 12):
log (C ppm ) = 7771 −
6272
T (°K )
(Eq 7)
The equilibrium diagram for carbon in a basic
18%Cr10%Ni alloy is shown in Fig. 7. At room
temperature, very little carbon is soluble in
austenite; even the 0.03% of L grades is mostly
in a supersaturated solution. The absence of carbides in austenitic stainless is due to the slow
diffusion of carbon and the even slower diffusion of chromium in austenite. At a carbon level
of 0.06%, which is found in most 304, supersaturation is reached below about 850 °C. Below
this temperature, supersaturation increases exponentially, while diffusion decreases exponentially. This results in precipitation rates that vary
with temperature and carbon level as shown in
Fig. 8. At these temperatures, grain boundary
diffusion is much more rapid than bulk diffusion, and grain boundaries provide excellent nucleation sites, so precipitation occurs along
grain boundaries. And, because carbon diffuses
several orders of magnitude more rapidly than
chromium, carbon diffuses to and combines
with chromium essentially in situ, depleting the
grain boundaries of chromium in solution.
Fig. 8 The precipitation rates for Cr23C6 as a function of carbon content
Chapter 6: Austenitic Stainless Steels / 77
Fig. 10 Variation of carbide precipitation locus with time.
Source: Ref 16
Fig. 9 Depletion of chromium from the austenite near grain
boundaries due to chromium carbide precipitation.
Source: Ref 14
Figure 9 shows that the local chromium depletion is such that the chromium level can become low enough that it has not even enough to
be stainless and certainly much lower corrosion
resistance than the surrounding area. This zone,
because it is lower in chromium, also has very
unstable austenite and is quite prone to martensite formation. Figure 10 shows how the locus
of precipitation changes with time and temperature. Carbon relatively far from grain boundaries in the interior of grains remains in supersaturation until much longer times and much
greater supersaturation since bulk diffusion is
required for the nucleation and growth of these
precipitates.
The key observation is that any solid-state
precipitation of a chromium-rich precipitate
will necessarily cause local chromium depletion
and a resulting loss of corrosion resistance.
Much longer term heat treatment is required to
eliminate these depleted zones by rehomogenization of slowly diffusing chromium than the
short time required to form them. This is very
evident for carbides, but also true for oxides.
Underneath chromium-rich oxide scales is a
layer depleted in chromium and lower in corrosion resistance. This is why not only scale from
welding must be removed, but also the underlying chromium-depleted zone.
Other precipitation processes that give rise to
chromium depletion are α and χ and the solidstate precipitation of oxides, nitrides, and sulfides. Chromium precipitates that form in the
liquid alloy do not cause depletion of chromium
locally because no chromium gradients are set
up around them during precipitation as diffusion in the liquid is very rapid. Thus, primary
carbides, oxides, and sulfides are not per se
harmful to corrosion resistance. But, if the same
compounds form and grow in the solid state,
chromium depletion occurs (Ref 15).
Alloying elements can have a major influence
on carbide precipitation by their influence on the
solubility of carbon in austenite. Molybdenum
and nickel accelerate the precipitation by diminishing the solubility of carbon. Chromium and
nitrogen increase the solubility of carbon and
thus retard and diminish precipitation. Nitrogen
is especially useful in this regard (Fig. 11).
Increasing austenite grain size accelerates
precipitation, as does cold work, especially in
the interior of grains, where diffusion is enhanced by increased defect density.
Nitrogen is much more soluble than carbon
and does not give rise to sensitization phenomena as does carbon even though Cr2N can be a
78 / Stainless Steels for Design Engineers
Fig. 11 Delay in carbide precipitation induced by nitrogen
level. Source: Ref 17
stable phase when the solubility limit is exceeded. The solubility is over 0.15% in austenite, so its precipitation seldom has the possibility of occurring, but it does become an issue in
ferritic stainless steels in this regard, for which
solubility is much lower. Manganese and
chromium increase the solubility of nitrogen in
austenite.
Stabilization. Before carbon was easily lowered to harmless levels, it was found that adding
more powerful carbide formers than chromium
could preclude the precipitation of chromium
carbides. Titanium and niobium are the most
useful elements in this regard. They form carbides with solubility that follows the following
equation type:
log [ M] [ X] = + A − H
RT
(Eq 8)
For titanium carbide and niobium carbide, the
respective solubilities are:
log [Ti ] [C] = 2.97 −
6780
T
(Eq 9)
9350
T
(Eq 10)
log [ Nb] [C] = 4.55 −
Oxides and sulfides are more energetically favorable than are carbides and nitrides of these
metals. Thus, any additions made to form carbides must be sufficient to account for the prior
formation of these compounds. Nitrogen also
competes with carbon for available titanium or
niobium. Thus, for successful gettering of all
carbon, there must be sufficient titanium or niobium to combine stoichiometrically with all
these species present.
This requires in rough terms that titanium exceed four times the carbon plus nitrogen, or that
niobium exceed eight times, after accounting
for the oxygen and sulfur. It would be a mistake
Fig. 12 Variation of hardness with depth and therefore carbon content in colossal supersaturation
to ignore the titanium-consuming capacity of
oxygen and sulfur unless they have been minimized by refining, which can be done quite
readily.
Even if sufficient titanium or niobium is present to combine with all carbon, kinetic considerations may result in that not occurring. High
temperatures, such as encountered in welding,
dissociate carbides. If quenched from this state,
carbon can be free to form Cr23C6 if it is reheated to temperatures above 500 °C.
Carbon has always been considered totally
undesirable from a corrosion point of view because of its tendency to form chromium carbides. Recently, however, new processes have
been developed to supersaturate carbon in
austenite below the temperatures at which it has
sufficient mobility to form carbides. This socalled colossally supersaturated austenite results in very high hardness (Fig. 12) and corrosion resistance over limited depths. From this,
however, we can see that carbon, like nitrogen,
is actually beneficial to corrosion resistance in
solid solution, although this is not observed at
normal concentrations. It is possible to see that
if it could be kept in solution it would be appropriate to give it a factor of around 10 in the pitting resistance equivalent number (PREN)
equation:
PREN = %Cr + 3.3(%Mo)
+ 30(%N) + 10(%C)
(Eq 11)
This is consistent with the similar thermodynamic interaction coefficients that carbon and
nitrogen share with regard to chromium.
Chapter 6: Austenitic Stainless Steels / 79
High-Temperature Alloys
The austenitic stainless steels can have an
exceptional combination of strength and corrosion resistance at temperatures above 500 °C.
They are often called on to resist attack by
oxygen, sulfur, carburizing, nitriding, halogens,
and molten salts. Austenitic stainless steels are
the most creep resistant of the stainless steels.
Alloying with carbon, nitrogen, and niobium
produces the greatest strength at elevated temperatures. Refer to the properties database for
comparisons among the grades. This discussion concentrates on their resistance to hightemperature environments, which is their salient
characteristic.
Oxidation Resistance. Resistance to oxidation comes from the protective Cr2O3 scale that
forms on the surface of the material. Above
about 18% chromium, a continuous scale forms.
The scale acts as a barrier to oxygen and greatly
slows further oxidation of metal below the
scale. Below the composition at which complete Cr2O3 coverage occurs, the film will also
contain the less-protective spinel FeCr2O4. The
Cr2O3 scale is more protective because it better
restricts the diffusion of oxygen to the interface
between the scale and the base metal, which is
where the oxidation reaction occurs. As the
oxide grows, the path to the interface lengthens,
and the rate of oxidation slows. This generates
the parabolic-type oxide growth rate that characterizes these alloys. The rate of oxide growth
is expressed simply as:
R = kt
(Eq 12)
in which the rate is in units of mass gained per
unit of surface area and time. This rate is a
strong function primarily of chromium level, as
can be seen in Fig. 13.
The rate increases exponentially with temperature since diffusion governs this phenomenon.
The rate drops dramatically as chromium
reaches the concentration necessary to generate
the protective Cr2O3 layer. Above this sufficiency level, further increases in chromium are
not as beneficial; they mainly provide a reservoir of chromium to re-form the Cr2O3.
As long as chromium can diffuse to the interface at a sufficient rate to satisfy the incoming
flux of oxygen, the parabolic rate is maintained.
If there is insufficient chromium flux, then the
oxygen penetrates beyond the interface and
forms Cr2O3 in situ. The oxide will change to
Fig. 13 Variation of parabolic oxidation rate with chromium
level and temperature. Source: Ref 18
the less-protective FeCr2O4, and the scale
growth rate will increase beyond the parabolic
relationship, leading to breakaway oxidation.
The breakaway temperature increases with increasing chromium level.
The austenitic alloys benefit over the ferritic
alloys from the presence of nickel. For a given
chromium level, oxidation rates decrease with
increasing nickel content. Figures 14, 15, and
16 display this relationship. The optimal range
for the iron base stainless steels, shown in Fig.
14, is reached by the commercial alloy 310,
with 25Cr-20Ni composition.
Alloying can alter the oxidation-resisting performance of the austenitic stainless steels. Some
elements form more protective oxide layers
than Cr2O3. Aluminum and silicon are most useful in this regard. Aluminum forms a layer of
Al2O3 that is more restrictive to oxygen diffusion than is Cr2O3, as does silicon through the
formation of SiO2. The alloys 302B, 153MA,
and 253MA all use elevated silicon levels. Aluminum’s strong ferrite-promoting tendency restricts its utility in austenitic grades, however.
While the gains from using under 3% silicon
are impressive, rare earths can yield even
greater benefits from mere trace additions.
153MA (UNS S30415) is a variation on 304
using silicon and cerium. Cerium appears to act
at the metal-scale interface such that the oxides
80 / Stainless Steels for Design Engineers
Fig. 14 Influence of nickel on oxidation of iron-chromium
alloys. Source: Ref 19
Fig. 16 Corrosion rates for various stainless steels and nickel
base alloys. Source: Ref 20
Fig. 15 Isooxidation curves. Source: Ref 20
formed are thinner, tougher, more adherent, and
more protective, although there is no consensus
on the mechanism. Figure 17 shows the improvement quantitatively (Ref 21).
Because austenitic stainless steels have a
greater thermal expansion coefficient than fer-
ritic alloys, they stress their scale more during
thermal cycling. This can fracture the scale,
causing spalling and rapid subsequent oxidation
of the underlying metal. This serious performance flaw also is remedied by rare earths, as
shown in Fig. 18 (Ref 21).
Chapter 6: Austenitic Stainless Steels / 81
Fig. 17 Comparison of rare earth-alloyed stainless alloys to conventional stainless alloys, 4833 = 309S, 4845 = 310S.
Source: Ref 21
Fig. 18 310S (4845) versus rare earth-modified 253MA in
cyclic oxidation. Source: Ref 21
While the mechanism by which rare earths
make the scale tougher and more adherent are
vague, their effect in making austenitic alloys
much better at resisting high-temperature oxidation, especially cyclic oxidation, are undeniable.
Alloying elements can also be detrimental to
oxidation resistance. Manganese, an even more
potent oxide former than chromium, forms a
manganese-chromium spinel that is less protective than the Cr2O3. Molybdenum and tungsten,
which are refractory metals and are beneficial to
aqueous corrosion resistance, form volatile, lowmelting oxides (MoO3 and WO3) that promote
catastrophic oxidation (Ref 22, 23). Vanadium
also forms an oxide, V2O5, which melts at 660
°C and can also cause catastrophic oxidation.
The formation of an oxide on stainless steel
should be understood to imply de facto the depletion of chromium from the underlying
metal surface. Whether the scale is formed in
service, during heat treating, or during welding
the surface, once the oxide is removed, there
will be less chromium than the bulk alloy,
often by a very significant amount, and therefore the corrosion resistant will be less. This is
why welds must have not only their heat tint
removed, but also the underlying metal which
is depleted in chromium, to a depth on the
order of 10 μ (Ref 24).
Microstructure can also affect oxidation resistance. As a generalization, it can be said that
changes that promote the diffusion of chromium
assist in the formation of a protective scale and
improve oxidation resistance. Thus, cold work
and finer grain size are positive factors via their
enhancement of chromium diffusion.
At the very highest temperatures, Cr2O3 can be
further oxidized to CrO3, which has significant
vapor pressure above about 1100 °C. The compositions of some of the main high-temperature
austenitic alloys mentioned here are given in
Table 2.
Other Environments. The most common addition species that aggravates high-temperature
oxidation is water vapor. At 10%, water vapor
will increase oxidation by a factor of ten. It acts
by increasing the porosity of the oxide scale and
by promoting formation of the volatile CrO2
82 / Stainless Steels for Design Engineers
(OH)2 species. As a rule of thumb, maximum
service temperatures should be reduced by 50 to
100 °C in the presence of steam.
Halogens can attack oxide scales and cause
their degradation or volatilization.
Carburization and nitriding are best prevented
by an oxide layer that forms at very low oxygen
partial pressures as chromium and silicon contents are increased. Austenitic alloys have no
advantage over ferritic in this regard.
Intermetallic Phases. Transition elements
may combine to form intermetallic phases in
which the formula can vary from B4A to A4B.
Table 3 lists the most common secondary
phases encountered in austenitic stainless steels,
i.e., apart from austenite and ferrite. Sigma formation is retarded by nitrogen, so alloys such as
153MA are less prone to it. Lower chromium
and higher nickel are beneficial. Silicon and
aluminum are detrimental, as is molybdenum.
The most relevant is σ. It can contain as little
as four iron to one chromium or molybdenum in
a tetragonal structure. Thus, it can exist in many
conventional austenitic alloys. Other relevant
phases are χ and Laves. The greatest risk from
these phases is the loss of room temperature
toughness, followed by some loss of corrosion
resistance.
Table 2
In lean austenitic alloys used in high-temperature, 600 to 1000 °C service, formation times
are relatively long, on the order of 100 h or
more. In richer alloys, such as 310, times can be
as short as several hours. If the temperature at
which the alloy is to be used is one in this temperature range, then some σ is a foregone conclusion, and while σ will have little detrimental
effect on short-term properties at these temperatures, long-term properties such as creep, stress
rupture, and especially rupture ductility are degraded by σ. For alloys, σ is an even greater
concern as these are prone to its formation and
can inadvertently form some during processing.
If such alloys are intended for use near room
temperature, then their toughness will be seriously reduced by the brittle σ, which will form
first at triple points and then throughout grain
boundaries. Because of this morphology, just a
few percent intermetallic phase can cause
toughness to decrease by an order of magnitude.
High-Temperature Mechanical Properties.
Above about 500 °C, yield strength is less appropriate than creep strength in assessing the
adequacy of an austenitic stainless steel for
structural purposes. The resistance of a material
to creep is generally measured by the creep rupture strength, which is the stress that causes
Notable high-temperature austenitic alloys
Alloy
Designation
C
N
Cr
Ni
Mo
Mn
Si
Other
Max temp, °C
302B
304H
321H
153MA
309S
309Si
253MA
310S
353MA
330
332Mo
S30215
S30409
S32109
S30415
S30909
DIN 1.4828
S30815
S31008
S35315
S33000
S35125
0.15
0.08
0.06
0.05
0.08
0.08
0.08
0.05
0.05
0.06
0.04
0.07
0.08
0.03
0.15
0.07
0.07
0.17
0.03
0.15
...
0.04
17.8
18.8
17.8
18.5
23.0
19.8
21.0
24.6
25.0
18.0
21.0
8.1
8.1
9.1
9.5
12.2
11.1
10.5
19.2
35.0
35.0
34.5
...
...
...
...
...
...
...
...
...
...
2.4
1.8
1.8
1.8
0.6
1.7
1.8
0.6
1.6
1.5
1.7
1.1
2.5
0.50
0.50
1.3
0.50
2.0
1.5
0.60
...
0.90
0.40
...
...
950
820
820
1000
1040
1040
1100
1090
1200+
1200
1200
Table 3
Secondary phases in austenitic stainless steel
Precipitate
Structure
NbC
NbN
TiC
TiN
Z-phase
M23C6
M6C
σ-phase
Laves phase
χ-phase
G-phase
fcc(a)
fcc
fcc
fcc
Tetragonal
fcc
Diamond cubic
Tetragonal
Hexagonal
bcc(b)
fcc
(a) fcc, face-centered cubic; (b) body-centered cubic
Parameter, (Å)
a = 4.47
a = 4.40
a = 4.33
a = 4.24
a = 3.037, c = 7.391
a = 10.57–10.68
a = 10.62–11.28
a = 8.80, c = 4.54
a = 4.73, c = 7.72
a = 8.807–8.878
a = 11.2
0.6 Ti
0.05 Ce
...
...
1.0 Al, 0.05 Ce
...
0.05 Ce
...
0.40 Nb
Composition
NbC
NbN
TiC
TiN
CrNbN
Cr16Fe5 Mo2C (e.g.)
(FeCr)21Mo3 C; Fe3Nb3C; M5SiC
Fe,Ni,Cr,Mo
Fe2Mo, Fe2Nb
Fe36Cr12 Mo10
Ni16Nb6 Si7, Ni16Ti6 Si7
Chapter 6: Austenitic Stainless Steels / 83
rupture after 10,000 or 100,000 h. If deformation is a greater concern, however, the creep deformation strength, that is, the stress that results
in a strain of 1% after 10,000 or 100,000 h, can
be used as a basis for design.
Cold work and precipitates tend to be ineffective strengtheners at temperatures that produce
solution annealing and precipitate coarsening
(overaging). Thus, solid solution hardening is
Fig. 19 Charpy V toughness after 200 hr aging
Fig. 20 Relative 100,000-h creep strength
preferred. Substitutional elements have limited
effect, but interstitial solid solution elements,
such as carbon and nitrogen, are quite useful.
Nitrogen is the better addition in this regard,
plus it has the collateral benefit of strongly
retarding intermetallic phase precipitation.
Figures 19 to 22 compare mechanical properties
of the major high-temperature austenitic alloys
(Ref 25).
84 / Stainless Steels for Design Engineers
Fig. 21 100,000-h creep rupture strength
Fig. 22 High-temperature, short-time yield strength
Corrosion-Resistant Austenitic Alloys
Stainless steels are almost always chosen at
least in part for their corrosion resistance. In
normal atmospheric conditions, alloys with
more than 10.5% chromium do not rust.
Austenitic alloys require higher levels than
this to stabilize the austenitic structure at room
temperature, thus giving the common perception
that they are superior in corrosion resistance.
The main advantage austenitic alloys have is
their ability to utilize the powerful and inexpensive alloying element nitrogen. That is the key
aspect of the more modern austenitic stainless
steels that have come into use in the last two
decades.
Chapter 6: Austenitic Stainless Steels / 85
The ion that is most aggressive against stainless steel is one of the most pervasive in our
environment. The chloride ion is found in
abundance over the entire earth. It is, of
course, in seawater, but also in the rain, on
roads, in food, and even in our bodies. Chlorides destabilize the passive film. If the conditions of chloride concentration, temperature,
and acidity are sufficiently aggressive to break
down the film, then active corrosion ensues. If
this is highly localized because of a local
weakness in the passive film, then pitting occurs. Such a pit may be unstable and repassivate, or it may grow without limit. Other
halides have the same effect, but they are less
ubiquitous.
Because of the specific virulence of the chloride ion and because of its universal presence,
corrosion-resistant austenitic stainless steels all
look like they were designed to resist chloridepitting attack. Pitting in stainless steels is in
most instances the threshold level of corrosion.
Crevice corrosion is, however, more severe and
usually design limiting vis-à-vis corrosion.
Crevices can exist not only in deliberate joints
but also under environmental deposits, paint
films, weld splatter, etc. There are other environments in which the resistance follows different rules, such as oxidizing acids, bases, and organic acids, but these are best regarded as
exceptions.
The main factors in the resistance of austenitic
alloys to pitting attack are generally given by:
PREN = %Cr + 3.3(%Mo) + 30(%N)
(Eq 13)
Pitting resistance is measured by ASTM G 48
(practice C) in which the lowest temperature at
which pitting occurs in a 6% FeCl3 is measured.
This is the critical pitting temperature, CPT. The
relationship between PREN and CPT is shown
in Fig. 23 (Ref 26).
The function of chromium in the passive film
is intuitively clear. As the chromium content of
an alloy increases, the ready reservoir of
chromium to form the chromium-rich layer is facilitated. The roles of molybdenum and nitrogen
are subtler and are still subject to controversy. It
is the subject of much research, which has been
summarized in reviews. The obvious paradox is
how can elements that are not active in the passive film be so effective in maintaining its integrity. We do know that the essential chromium
in the matrix of stainless steels is quite reactive
and will form compounds with carbon, oxygen,
sulfur, and other transition elements. When it
does, it is no longer effective as a passive film
former. The regions from which the chromium
diffused to form the chromium-rich phase will
be poor in chromium unless subjected to a
lengthy homogenizing anneal. Most theories of
pitting founder at the start because they assume
Fig. 23 Critical pitting temperature versus pitting resistance equivalent number (PREN); SUS 329J4L = S31260, YUS 270 = S31254.
Source: Ref 26
86 / Stainless Steels for Design Engineers
a homogeneous passive film, which is an impossible goal in reality. The chapter on corrosion
deals with this topic in more depth.
The passive layer is extremely thin compared
to oxide layers. It is on the order of 1 to 10 nm
thick. Its formation does not cause chromium
depletion beneath it, as oxide layers do. As the
alloy content of chromium and molybdenum increase, the film is thinner, and the current density required to form the film is correspondingly
reduced.
The corrosion-resistant austenitic stainless
steel grades range from 316 to the various highmolybdenum, high-nitrogen alloys commercialized in the last ten years, the most notable of
which are listed in Table 4 with their typical
analyses.
Early grades were based on alloying with
chromium and molybdenum with sufficient
nickel added to preserve the austenitic structure.
Each of these elements facilitates the formation
of the passive film and reduces the corrosion
rate in the active state. Further experimentation
showed that molybdenum was not beneficial
under highly oxidizing conditions, but that silicon was helpful under such conditions. Copper
was beneficial against sulfuric acid.
Alloy development came in stages. First, 317
was the most significant corrosion-resistant
alloy. Then, more chromium and molybdenum
were added, and the class of alloys known collectively as the 6%Mo alloys was commercialized. Allegheny’s AL-6X™ is representative of
this group. With PRENs of around 40, these alloys were resistant to seawater at ambient but
not at elevated temperatures. This left a great
deal wanting in corrosion resistance. These alloys were also very difficult to process, at least
in part because they rapidly formed brittle grain
boundary σ-phases in as little as several minutes
at some temperatures. This limited chromium
and molybdenum levels to a total of about 30%.
Table 4
The discovery that nitrogen was beneficial
against corrosion permitted a breakthrough in
alloy development by the 1980s. Nitrogen was
increased to around 0.20% from nominal levels
of 0.05%. This was found to increase PREN by
another five units, but more importantly, also
suppressed σ formation to times that permitted
thicker sections to be welded without embrittlement. Research into the thermodynamics of
nitrogen in austenite showed that manganese increased the solubility of nitrogen appreciably.
This permitted even higher levels of total alloying to be achieved. This was exploited in the
alloys UNS S34565 and S32654, which contain
3 to 6% manganese and about 0.50% nitrogen.
The PRENs of these alloys are over 50, which
gives them a critical pitting temperature around
100 °C. Table 5 lists the performance of the various popular corrosion-resistant austenitic stainless steels.
The advances are quite appreciable and made
stainless steel a viable material for many applications for which previously nickel base or titanium alloys had been required.
The obvious question in view of the success
of the use of high manganese levels in conjunction with high nitrogen levels in the most highly
alloyed austenitic stainless steels is when this
approach will be used for the medium-alloyed
austenitics to make alloys superior to 316, 317,
and 904 without the high levels of nickel and
molybdenum that render these alloys so expensive. It does not take much imagination to envision alloys equal to 904L in PREN with less
nickel and molybdenum than 316L that would
be almost totally resistant to intermetallic phase
precipitation and have much greater resistance
to SCC because of higher austenite stability.
The same case could be made for a 317-equivalent alloy in corrosion resistance with less than
7% nickel and 0.5% molybdenum, essentially a
301 in alloy cost. In view of these trends in
Typical compositions of corrosion-resistant austenitic stainless steels
Alloy
Designation
C
N
Cr
Ni
Mo
Mn
Si
Other
Other
316L
316Ti
317L
317LM
904L
JS700
254SMO
4565
654SMO
AL6-XN
Al6-XN Plus
S31603
S31635
S31703
S31725
N80904
N08700
S31254
S34565
S32654
N08367
N08367
0.02
0.02
0.02
0.02
0.02
0.02
0.02
0.01
0.01
0.02
0.02
0.03
0.03
0.06
0.06
0.06
0.06
0.20
0.45
0.50
0.22
0.24
16.4
16.4
18.4
18.4
19.5
19.5
20.0
24.0
24.0
20.5
21.8
10.5
10.5
12.5
13.7
24.0
25.0
18.0
18.0
22.0
24.0
25.3
2.1
2.1
3.1
4.1
4.1
4.4
6.1
4.5
7.2
6.2
6.7
1.8
1.8
1.7
1.7
1.7
1.7
0.8
6.0
3.0
0.4
0.3
0.5
0.5
0.5
0.5
0.5
0.5
0.4
...
...
0.4
0.4
...
0.40
...
...
1.3 Cu
0.4 Cu
0.8 Cu
...
0.5 Cu
0.2 Cu
0.2 Cu
...
...
...
...
...
0.3 Nb
...
...
...
...
...
Chapter 6: Austenitic Stainless Steels / 87
alloy development, the future use of 316 and
317 should be numbered. There is no justification for the use of scarce and expensive resources such as nickel and molybdenum when
cheap, abundant replacements like manganese
and nitrogen are available. The use of 316 as a
standard alloy should in the future be eroded by
more cost-effective alloys such as the lean duplex alloys like 2003.
The same environments that cause pitting
corrosion also cause crevice corrosion. A
crevice is a volume in and out of which diffusion is restricted to a degree that corrosion products accumulate and cause the contained environment to become increasingly aggressive in
pH and [Cl–]. Conditions that are below the
threshold for pitting can cause crevice corrosion. The critical temperature for crevice corrosion is also measured in 6% FeCl3 (ASTM G-48
B or D). It is the lowest temperature at which
crevice corrosion occurs. This temperature, the
CCT (critical crevice temperature), is lower
than the CPT. GTAW a wrought alloy also lowers the CPT to about the level of the CCT. The
reason for this lowering of resistance to localized attack has been thought to be related to
alloy depletion caused either by dendrite coring
during solidification or chromium depletion
around inclusions. The relation to crevices
would thus seem to be that surfaces contain numerous flaws with respect to corrosion resistance, which, while not capable of sustaining pitting, can in a crevice dissolve and alter the
environment sufficiently that the new harsher
environment can generally destabilize the passive film and proceed autocatalytically.
Stress corrosion cracking is the bane of
austenitic stainless steels. SCC occurs when
there is both a tensile stress of a sufficient magnitude and a sufficiently aggressive environment. The threshold stress varies with alloy and
thermomechanical history. As a rule of thumb,
the environment to initiate SCC must be sufficiently severe to cause localized corrosive attack. The most dangerous situation is one in
which the expectation of pitting is marginal.
The mechanism of SCC has been a subject of
intense academic controversy for many years.
Because of this, much of the literature has focused on arguing a case rather than clarifying
the phenomenon. What can be said about SCC
in austenitic stainless steels with consensus?
• Risk of SCC is low at room temperature and
increases exponentially with temperature.
• SCC is preceded by localized corrosive attack, which has an incubation time, and then
proceeds in a discontinuous manner.
• Fracture may be transgranular, intergranular,
or both. It is almost entirely lacking in plastic deformation with little, if any, metal loss.
• Alloying or treatments that delay localized
attack or stabilize austenite can delay SCC
up to the point of virtual immunity.
• SCC is aggravated by increased chloride
concentration and acidity, but also exists in
caustic environments.
• Stress must exceed a threshold for a given
set of conditions for SCC to occur.
• Anodic or cathodic polarization may prevent
SCC under conditions at which it would otherwise be expected, or it may cause SCC in
environments in which it would not otherwise occur.
Austenitic stainless steels are not alone in their
susceptibility to SCC. All stainless steels suffer
from SCC under some set of conditions of environment and material thermomechanical history.
The key is to choose an alloy that is resistant
under the conditions of use. To a first approximation, this means using an alloy that will not pit
under the conditions of use, then designing its
Table 5 Corrosion resistance ratings of various austenitic stainless steels, using 30 factor for
nitrogen
Alloy
Designation
PREN(a)
CPT oC
CCT oC
316L
316Ti
317L
317LMN
904L
JS700
254SMO
4565
654SMO
AL6-XN
AL-6XN Plus
S31603
S31635
S31703
S31726
N80904
N08700
S31254
S34565
S32654
N08367
24
23
30
34
35
36
46
53
63
50
50 min
15
15
25
30
40
43
73
90
105
78
95
–3
–3
0
4
15
15
38
50
75
43
60
(a) PREN, pitting resistance equivalent number.
88 / Stainless Steels for Design Engineers
use to be below the threshold stress completes a
sound design approach if residual stresses can be
accurately known. Otherwise, assuming that the
metal will have residual stresses equal to 100%
of the yield strength is the prudent engineering
approach. Figure 24 shows threshold stress for a
number of alloys.
Special Corrosive Environments. Knowledge of the ability of the various stainless steels
to resist specific environments is essential to the
design process. This information is extensive
since it must correlate many environments and
temperatures for many materials. Hence, refer
to the publications of organizations such as the
National Association of Corrosion Engineers
(NACE) or to the Web sites of companies such
as Allegheny Ludlum or Outukumpu, where
such information is available freely. The more
reputable producers will give assistance on specific questions. Engineering forums on the Internet, such as www.eng-tips.com, should also
be considered a resource. The following discussion presents just the principles of the resistance
of austenitic stainless steels to specific, more
common environments.
Sulfuric acid is common, aggressive, and must
be contained. Figure 25 shows the isocorrosion
curves for several alloys in pure sulfuric acid.
Alloy 20, 904L, and alloy 825 were developed specifically for sulfuric acid service. Each
contains several percent copper that, while not
beneficial against pitting, concentrates in the
passive film and diminishes general corrosion.
Molybdenum and tungsten are also very beneficial for resistance to sulfuric acid. Phosphoric
acid is similar to sulfuric acid in its effect on
austenitics but somewhat less aggressive.
Fig. 24 Threshold stress for stress corrosion cracking (SCC)
for various alloys. Source: Ref 27
Nitric acid is not particularly aggressive
against stainless steels. Resistance to it is proportional to chromium content. So, attack,
should it take place, is preferably at grain
boundaries, where segregation of elements such
as carbon, phosphorus, and silicon can lower
chromium locally. These elements are kept low
for nitric acid service. Pitting is not a risk. Standard usage is:
• Below 50% concentrations and below 100
°C, 304L and 17% Cr ferritics are used.
• Around the 65% aziotrope, 310 is most resistant, especially a low carbon version, but
304 NAG with low carbon, phosphorus, and
silicon is more often used.
• For 98% solutions or for lower concentrations that contain other stronger oxidizers,
alloys with 4% Si, 18% Cr, and 15% Ni or
5% Si, 17% Cr, and 17% Ni (UNS S30600
and S30601) have been developed.
Hydrochloric acid, not surprisingly, is quite
aggressive against stainless steel. It is very effective in destabilizing the passive film. Thus,
resistance to hydrochloric acid is simply an extreme case of resistance to pitting in chlorides
with resistance given by Eq 12. Only the most
highly alloyed austenitic alloys, such as AL6XN®, should be considered and then under
conditions that are tolerable, such as those
shown in Fig. 26.
Strong bases such as NaOH and KOH are not
especially aggressive against stainless. The 17%
chromium alloys can be used up to 50 °C, while
304L can be used to 90 °C. As is the case with
nitric acid, chromium and nickel are beneficial,
while molybdenum is counterproductive. The
25% chromium alloys such as 310 or an equal
Fig. 25 Isocorrosion in pure sulfuric acid. Source: Ref 28
Chapter 6: Austenitic Stainless Steels / 89
Fig. 26 Resistance to hydrochloric acid. Source: Ref 28
chromium duplex can be used to 150 °C, above
which temperature nickel base alloys are required. High-chromium ferritic stainless steels
are also very good choices.
Organic acids are generally less aggressive
against stainless than are mineral acids since
they are less dissociated in solution. They become hazardous when they contain chloride
ions, at high temperatures, or when they dissociate strongly, such as with formic acid.
Because of the large number of organic compounds that may be considered, refer to the various corrosion tables.
Surface Finish. The corrosion resistance of
austenitic stainless steels is quite dependent on
surface condition, as are other stainless steels.
Treatments that enhance surface concentrations
of beneficial elements or remove detrimental
constituents can greatly alter performance.
Oxide formation depletes surface chromium, so
strong pickling or electropolishing of the
descaled surface is especially important. Studies
have shown chromium depletion of a maximum
of 6% extending 10 μ before reaching bulk
chromium levels. This is equivalent to the depletion seen in sensitization. The increase in attack rate from this depletion is huge. A 1000fold increase in weight loss in the ASTM G 48
B test has been seen by a superficial loss of 6%
chromium.
Likewise, surface abrasion, especially coarse
abrasion, has a major detrimental effect. The
120 grit #3 finish often seen on stainless reduces
pitting resistance by as much as the equivalent
of 5 PREN, that is, equal to a reduction in
chromium content of 5%. Rolled finishes are
much preferred. The mechanism for this has not
been clearly established; exposure of MnS inclusions, the microcrevices abrasion produces,
and residual stress have been cited as possible
contributing causes. Powder injection-molded
stainless components often have porosity that is
generally spherical. When exposed to the surface, such pores act as crevices and lower the
pitting potential also. All of these factors are operative and can act in unison.
Very fine abrasive polishing causes little
residual stress and has very minimal crevice
creation. Thus, mirror-type polished finishes do
not degrade corrosion resistance, but they do
not enhance corrosion resistance as does electropolished mirror finishes, which remove
chromium-depleted sites, which can initiate
pitting.
REFERENCES
1. A.L. Schaeffler, Constitution Diagram for
Stainless Steel Weld Metal, Met. Prog.,
Vol 56, Nov 1949, p 680–688
2. W.T. Delong, A Modified Phases Diagram
for Stainless Steel Weld Metals, Met. Prog.,
Vol 77, Feb 1960, p 98
3. Design Guidelines for Selection and Use of
Stainless Steel, SSINA,1998, p 3
4. K.-J. Blom, “Press Formability of
Stainless Steels,” paper presented at Stainless steels ‘77
5. F.B. Pickering, “Physical Metallurgical Developments in Stainless Steel,” paper presented at Stainless Steel ‘84, Goteborg
6. Q.-X. Dai et al., Chin. Phys., Vol 11, 2002, p
596–600, doi:10.1088/1009-1963/11/6/315
7. Aciers Inoxidables, Les Editions de
Physique, Les Ulis, Paris, 1993, p 564
8. Aciers Inoxidables, Les Editions de
Physique, Les Ulis, Paris, 1993, p 565
9. P. Marshall, Austenitic Stainless Steels, Microstructure and Mechanical Properties,
Elsevier, 1984
10. Aciers Inoxidables, Les Editions de
Physique, Les Ulis, Paris, 1993, p 579
11. H. Nordberg, Mechanical Properties of
Austenitic and Duplex Stainless Steels, Innovation in Stainless Steels ‘93 (Firenze),
1993, p 2.217
12. Aciers Inoxidables, Les Editions de
Physique, Les Ulis, Paris, 1993, p 566
13. S.J. Rosenberg and C.R. Irish, J Res. Nat.
Bar. Stand., Vol 48, 1952, p 40
14. Aciers Inoxidables, Les Editions de
Physique, Les Ulis, Paris, 1993, p 410
15. M. McGuire, “A Diffusion Model for the
Influence of Oxygen and Sulfur on the NonEquilibrium Distribution of Chromium in
90 / Stainless Steels for Design Engineers
Austenitic Stainless Steel Welds and Slabs,”
paper presented at Proceedings MS&T
‘04,2004
16. R. Stickler and A. Vinckier, Trans. ASM,
Vol 54, 1961, p 362
17. Aciers Inoxidables, Les Editions de Physique, Les Ulis, Paris, 1993, p 568
18. Aciers Inoxidables, Les Editions de
Physique, Les Ulis, Paris, 1993, p 448
19. Aciers Inoxidables, Les Editions de Physique, Les Ulis, Paris, 1993., p 453
20. Aciers Inoxidables, Les Editions de
Physique, Les Ulis, Paris, 1993, p 454
21. www.outukumpu.com
22. W.C. Leslie, Mechanism of Rapid Oxidation at High Temperature, Trans. ASM, Vol
41, 1958, p 1213–1219
23. N. J. Grant, Accelerated Oxidation of Metals at High Temperature, Trans. ASM, Vol
44, 1961, p 128–137
24. J.F. Grubb, paper 04291 presented at Corrosion 2004, NACE, 2004, p 1–15
25. ACOM Files, High Temperature Stainless
Steels, www.outukumpu.com
26. J. Okamoto et al., A Super-Austenitic Stainless Steel for Tubing and Piping Applications, Nippon Steel Technical Report 90,
July 2004
27. ACOM Files, High Temperature Stainless
Steels, www.outukumpu.com
28. Allegheny
Technologies,
“AL6-XN®
Alloy”
Stainless Steels for Design Engineers
Michael F. McGuire, p 91-107
DOI: 10.1361/ssde2008p091
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
CHAPTER 7
Duplex Stainless Steels
Summary
THE NEWEST FAMILY of stainless steels is
the duplex alloys. The mixture of ferrite and
austenite in their structure gives them higher
strength than either phase by itself. Duplex alloys have at least 20% chromium, so they are
considered as highly corrosion-resistant alloys
but not high-temperature alloys because of embrittling phases. Their low nickel content makes
them more economical than austenitic alloys of
the same level of corrosion resistance, especially when their greater strength can be utilized
to reduce the amount of material required. They
should largely replace alloys such as 316L
and317L in the future.
Introduction
Duplex stainless steels are the newest and
fastest-growing alloy group in the stainless steel
family. They are called duplex because at room
temperature they consist of two phases, ferrite
and austenite. Discovered in the 1920s, they
languished in a suboptimized and underutilized
state until recently. They possess excellent
strength, toughness, and corrosion resistance.
They also display exceptional resistance to
stress corrosion cracking (SCC) and corrosion
fatigue. The leaner grades, such as 2304, correspond to 316L in corrosion resistance but have
double the yield strength, while the higher alloy
grades like 2507 compete with the 6% molybdenum superaustenitics in corrosion resistance
while still possessing much greater strength.
Their limitations lie in their lack of cryogenic
toughness and their inability to withstand temperatures much above 300 °C without forming
embrittling phases. But between –100 and 300 °C
they are exceptional materials. Whether these
duplex alloys will grow to the full extent of
their potential depends on several factors:
• Will high nickel and molybdenum prices be
sufficient motivation to drive designers to
explore alternatives to traditional austenitic
grades?
• Will producers overcome their inhibition to
aggressively market these grades through
their cost-saving potential?
• Will producers perfect the techniques to produce these grades reliably so that their availability is unquestioned?
• Will design codes change to correctly reflect
the duplex materials’ higher ratio of yield
strength to tensile strength?
Why are there such issues with a family of alloys that has been successfully used for 20 years?
The concept of duplex stainless steels is simple:
islands of austenite in a continuous matrix of
highly alloyed ferrite, as seen in Fig. 1. This
Fig. 1 Wrought 2205 duplex microstructure
92 / Stainless Steels for Design Engineers
combination in principle offers high strength because of the possibility of refining the dual-phase
grain structure and thereby raising yield strength
according to the Hall-Petch relationship as well
as by solid solution hardening, especially with
nitrogen. In addition, the absence of a continuous
austenite phase provides relief from SCC by having any propagation of cracks in austenite arrested by the ferrite phase.
The optimization of the alloy system had to
wait for two events, both related to nitrogen.
First, the control of nitrogen in the refining by
the argon oxygen decarburization (AOD) process
allowed the control nitrogen content up to the
solubility limit. Second, the understanding of
the thermodynamics of the alloy system became
understood and reduced to a computer model.
At this point, the alloys developed over the first
50 years of development became obsolete, and
new grades with higher nitrogen vastly improved
the performance and user friendliness. Why this
was so important can be seen by studying the
structure of these alloys.
Structure and Alloy Design
The ideal structure of a duplex grade would
be a stable 50-to-50 ratio of austenite to ferrite
at all temperatures at which it is to be used without other phases. The austenite would be islands
in the ferrite matrix, and each phase would have
equal corrosion resistance despite having different compositions. It took a long time for that to
be accomplished.
Figure 2 shows a simple Fe-Cr-Ni constitutional diagram (Ref 1). The salient points are
that the typical successful alloys nearly bisect
the two-phase field for austenite and ferrite. It is
also obvious that the composition of the austenite and the ferrite must be quite different.
Ferrite contains a great deal more chromium
than austenite; hence, its pitting corrosion resistance contribution from chromium is much
greater than the resistance of the austenite because in duplex grades:
PREN = %Cr + 3.3 × %Mo + 16 × %N
(Eq 1)
If one were to add molybdenum to increase
pitting resistance, it would preferentially partition to the ferrite, further exacerbating the differential between the two phases.
This is where nitrogen saves the day. Additions of nitrogen concentrate nearly entirely in
the austenite. This lowers the activity of
chromium and thereby effectively attracts more
chromium to the austenite phase than would
otherwise be present. This stabilizes the austenite, keeping the ratio of ferrite to austenite more
nearly constant with temperature. The pitting
resistance of the austenite increases signifi-
Fig. 2 The Fe-Cr-Ni phase diagrams. The shaded area results from nitrogen additions
Chapter 7: Duplex Stainless Steels / 93
cantly to approximately that of the ferrite. In addition, the nitrogen solid solution strengthens
the austenite and retards the formation of intermetallic phases, which is not bad for an element
that costs nothing.
ThermoCalc, developed by the Swedish Royal
Academy, has been an especially valuable tool in
helping us understand and design better duplex
stainless steels. Without being able to computer
model the thermodynamics of the system, it
would be impossible to project the partitioning of
the various potential alloying elements. Figure 3
shows isopleth diagrams for a basic 2205 composition in which nickel is varied.
The 2205 is the workhorse grade of duplex. It
has a pitting resistance equivalent number
(PREN) of about 35 and fills a niche in corrosion
resistance where austenitics and ferritics are
lacking, greater than 317L stainless, PREN = 30,
and below the 6% molybdenum grades, such as
AL-6XN alloy, with PRENs of around 45. Ferritics have a gap between 442 (18Cr-2Mo) and
the super ferritics (28Cr-4Mo). By varying the
chromium, nickel, and molybdenum, leaner alloys can be devised that save cost based on reduced molybdenum and nickel. Conversely,
more corrosion-resistant alloys with higher
PRENs can also be mapped, such as Fig. 4, with
the same diagrams varying nickel for a higher
molybdenum level. This composition includes
the important 2507 alloy.
Partitioning of elements (Fig. 5) between
austenite and ferrite is an important issue. The
partitioning tendency is a strong function of
temperature. Figure 6 shows that as temperature
Fig. 3 The iron-nickel diagram for 22% Cr, 3% Mo, 0.15% N
Fig. 5 Partitioning tendencies of various elements between
ferrite and austenite. Source: Ref 2
Fig. 4 The iron-nickel diagram for 25% Cr, 4% Mo, and
0.25% N: N is a nitride, χ is chi, σ is sigma, α is ferrite, and γ is austenite
Fig. 6 Variation of partitioning ratio with temperature.
Source: Ref 2
94 / Stainless Steels for Design Engineers
increases, the partitioning diminishes until at
just above 1300 °C it approaches unity for all
normal substitution-alloying elements (Ref 2).
For nitrogen, however, the tendency is to increasingly segregate to austenite as temperature
increases (Ref 2).
A danger in these alloys is that austenite
formed from ferrite on heating, such as during
welding or annealing, will contain only the low
amount of nitrogen that was in the ferrite from
which it was formed, until diffusion can restore
equilibrium. If the heating time does not permit
this, this so-called secondary austenite will
have low nitrogen and therefore low pitting
corrosion resistance, as shown in Fig. 7(e). Nitrogen alters the phase stability, making austenite stable to higher temperatures. This helps
keep welds from becoming excessively ferritic
and disturbing the desirable 50-to-50 ratio of
austenite to ferrite. Secondary austenite with
low nitrogen is remedied by diffusion if the
phase forms at higher temperatures at which
diffusion of nitrogen can rehomogenize the nitrogen level.
A crucial aspect of alloy design in the duplex
alloys involves the avoidance of unwanted
phases. The duplex stainless steels have all the
potential problems with embrittling phases of
the ferritic and austenitic stainless steels combined since they contain both as phases. Ferrite
forms two main embrittling phases, α′ and σ.
The α′ is generally believed to be a result of
the miscibility gap that exists in the ironchromium system, by which ferrite undergoes
spinodal decomposition into the iron-rich α,
normal ferrite, and the chromium-rich α′,
which is a brittle ordered alloy. Higher levels
of chromium or the presence of copper or
molybdenum exacerbate this reaction, which
has a formation that follows an Arrhenius-type
curve with a maximum at around 400 °C. Figure 8 shows the α′ formation kinetics for five
duplex alloys (Ref 2). While duplex grades
have good oxidation resistance and high-temperature strength, the α′ problem restricts their
use to below about 315 °C. Ferrite and austenite both form intermetallic phases, of which
the most prominent and dangerous is σ, a
tetragonal phase richer in chromium and
molybdenum than the ferrite from which it
forms. It is brittle and forms at grain boundaries, so its precipitation has the immediate effect of lowering toughness. Cold work accelerates the precipitation process by up to an order
of magnitude by virtue of its dual effect on nu-
cleation and diffusion. The areas around the
newly formed σ are naturally somewhat diminished in chromium and molybdenum, so
the alloy’s resistance to localized corrosion is
compromised also.
Figure 9 shows the TTT (time-temperaturetransformation) diagram for various high-alloy
stainless steel, including austenitic, ferritic, and
duplex. Alloys of all structures, ferritic,
austenitic, and duplex, with high chromium and
molybdenum encounter the σ problem fairly
equally and in proportion to their alloy content
(Ref 2). This is the reason that the use of nitrogen instead of molybdenum is so beneficial to
the leaner alloys, not just in cost, but for the
major reduction in rate of formation of sigma.
Figure 10 shows the large reduction in σ formation enjoyed by the lean alloy AL 2003™ material compared to the higher molybdenum 2205
alloy (Allegheny Ludlum).
There are other intermetallic phases in addition to σ. They include χ, R, π, and τ. These are
of more research than practical interest because
σ, with its bad consequences, forms sooner and
in greater quantity under the same conditions
compared to the others.
Carbides and nitrides can also form in duplex
alloys. The nitride Cr2N can form when saturated ferrite is quenched from a high temperature, as can occur in the welding process. It is
possible that this would result in nearby
chromium depletion and a decrease in corrosion
resistance. Carbide formation does not as easily
cause chromium depletion in duplex alloys because the precipitation at the ferrite-austenite
grain boundary does not deplete the austenite as
greatly in chromium locally because of the
neighboring ferrite having a much higher diffusivity for chromium. The point is generally
moot since all modern duplex grades contain
less than 0.030% carbon.
Table 1 lists the duplex grades currently
available commercially. Figure 7 shows a series
of duplex photomicrographs.
Mechanical Properties
In many ways, the duplex stainless alloys
represent a best of both worlds in combining
traits from the austenitic and ferritic alloys.
They offer high as-annealed strength with
good toughness and ductility. Table 1 lists
the major grades of duplex stainless steels;
Chapter 7: Duplex Stainless Steels / 95
Fig. 7 (a) As-cast duplex structure, austenite in a ferrite matrix. (b) 2205 annealed; austenite phase contains twins.
(c) 2507 as-welded; weld is highly ferritic because of rapid cooling rate. (d) Same weld as (c) after homogenization anneal.
(e) 7-Mo Plus with ( (dark areas) that has induced the formation of secondary austenite (arrows)
Table 2 lists typical and minimum properties
for the major duplex alloys and those of some
comparable ferritic and austenitic alloys for
comparison.
The most striking and unexpected characteristic of the duplex grades is their high yield
strength, more than double that of comparable
austenitic grades.
96 / Stainless Steels for Design Engineers
Fig. 8 Kinetics of ( formation
The strength of the duplex grades is driven by
the strength of the continuous ferrite phase. It
owes its strengthening primarily to:
• Solid solution hardening by nickel, molybdenum, chromium, copper, and manganese
• Interstitial solid solution hardening by carbon and nitrogen
• Strengthening by grain refinement
These components have been related to the
mechanical properties by the following equations (Ref 3):
Fig. 9 Sigma formation kinetics at various alloy levels
Rp0.2 = 120 + 210 N + 0.02
+2( Mn + Cr ) + 14 Mo + 10 Cu
+ (6.15 − 0.54δ)δ
+ (7 + 35( N + 0.02))d −1/ 2
Rp1.0 = Rp0.2 + 40 ± 9
(Eq 2)
(Eq 3)
Rm = 470 + 600( N + 0.02)
+14 Mo + 1.5δ + 8d −1/ 2
Fig. 10 Delay in ( precipitation in lean duplex 2003
(Eq 4)
where δ is the ferrite content in percent, d is the
lamellar spacing, and results are in megapascals.
The influence of nitrogen is interesting in that
at lower levels (e.g., below 0.1% nitrogen)
austenite is the weaker phase, but additional ni-
Chapter 7: Duplex Stainless Steels / 97
Table 1
UNS
S32900
S31200
S31260
S31500
S31830
S32001
S32003
S32101
S32205
S32304
S32520
S32550
S32750
S32760
S32906
S32950
S39274
S39277
Table 2
Duplex compositions
Name
C
N
Cr
Ni
Mo
Mn
P
S
329
44LN
DP3
3RE60
2205(old)
19 D
2003
2101
2205
2304
Uranus
52N+
255
2507
Zeron 100
2906
7–Mo Plus
DP3W
AF 918
0.08
0.03
0.03
0.30
0.03
0.03
0.03
0.04
0.03
0.03
0.03
...
0.14–0.20
0.10–0.30
0.05–0.10
0.08–0.20
0.05–0.17
0.14–0.20
0.20–0.25
0.14–0.20
0.05–0.20
0.20–0.35
23.0–28.0
24.0–26.0
24.0–26.0
18.0–19.0
21.0–23.0
19.5–21.5
19.5–21.0
21.0–22.0
22.0–23.0
21.5–23.5
24.0–26.0
2.5–5.0
5.5–6.0
5.5–7.5
4.25–5.25
2.5–3.5
1.0–3.0
3.0–4.0
1.35–1.70
4.5–6.5
3.0–5.0
5.5–8.0
1.0–2.0
1.2–2.0
2.5–3.5
2.5–3.0
2.5–3.5
3.0–5.0
1.0
2.0
1.0
1.2–2.0
2.0
4.0–6.0
2.0
4.0–6.0
1.0
2.5
1.5
0.75
...
...
1.0
...
...
0.75
0.2–0.8 0.1–0.5
1.4–2.0 . . .
...
1.0
...
...
1.0
...
...
1.0
...
...
1.0
0.1–0.8 . . .
2.0
...
...
1.0
0.05–0.6 . . .
0.8
0.5–3.0 . . .
0.040
0.040
0.030
0.030
0.030
0.040
0.040
0.040
0.030
0.040
0.035
0.030
0.030
0.030
0.030
0.020
0.030
0.030
0.030
0.020
0.040
0.020
0.04
0.03
0.03
0.03
0.03
0.03
0.025
0.10–0.25
0.20–0.30
0.20–0.30
0.30–0.40
0.15–0.35
0.24–0.32
0.23–0.33
24.0–27.0
24.0–26.0
24.0–26.0
28.0–30.0
26.0–29.0
24.0–26.0
24.0–26.0
6.0–8.0
6.0–8.0
6.0–8.0
5.8–7.5
3.5–5.2
6.0–8.0–
6.5–8.0
2.9–3.9
3.0–5.0
3.0–5.0
1.5–2.6
1.0–2.5
2.5–3.5
3.0–4.0
1.5
1.2
1.0
0.8–1.5
2.0
1.0
0.8
1.0
0.8
1.0
0.5
0.6
0.8
0.8
1.5–3.0 . . .
0.5
...
0.5–1.0 0.5–1.0
0.8
...
...
...
0.2–0.8 1.5–2.5
1.2–2.0 0.8–1.2
0.040
0.035
0.030
0.030
0.035
0.030
0.030
0.030
0.020
0.010
0.030
0.010
0.020
0.020
1.5–2.0
0.1–0.8
3.0–3.5
Si
Cu
W
Duplex mechanical properties
Grade
Name
Rp0.2
Rm
A5
HB
RC
Charpy-V
–40 °C, J
S31200
44LN
450
690
25
293
31
...
S31260
DP3
485
690
20
290
31
...
S31830
2205(old)
450
62
25
293
31
...
S32003
2003
450
620
25
290
30
40
S32001
19D
450
640
25
290
25
...
S32101
2101
450
650
25
290
32
40
S32205
2205(new)
460
640
25
290
32
40
S32304
2304
400
600
25
290
31
40
S32520
Uranus 52N+
550
770
25
...
28
...
S32550
Ferralium
550
760
15
302
32
...
S32750
2507
550
795
15
310
32
40
S32760
Zeron 100
550
750
25
270
...
...
S32960
7-Mo Plus
485
690
15
293
32
...
trogen strengthens the austenite so that above
0.2% nitrogen, the austenite becomes the
stronger phase.
The two phases are elongated parallel to the
major strain axis of working such as from hot or
cold rolling. As working increases, the microstructure and properties become increasingly
anisotropic, with the austenite taking on a (110)
[223] texture and the ferrite (100) [011] to (211)
[011] (Ref 2).
Because the ferrite phase controls mechanical properties, the dependence of these properties on temperature is significant since flow in
body-centered cubic (bcc) structures is thermally activated. Figure 11 shows the variation
of yield and tensile strengths of various grades
along with that of austenite and ferrite of similar composition. Because use of these alloys
above 300 °C is not recommended, no highertemperature properties are shown. Since they
have a ductile-to-brittle transition, they also are
not well suited to cryogenic use.
Impact Strength. Toughness is a significant
consideration when using duplex alloys to replace the extremely tough austenitic alloys.
Duplex alloy low-temperature toughness is intermediate to that of ferritic and austenitic alloys. This having been said, it should be noted
that the duplex alloys can have excellent toughness levels, such as 100 J at –100 °C in the solution-annealed condition. As would be expected, toughness improves with decreasing
grain size and deteriorates with cold work. The
most deleterious effect on toughness comes
from the precipitation of intermetallic phases,
such as α′ and σ, which cause a sharp decrease
98 / Stainless Steels for Design Engineers
Fig. 11 Variation of ferrite, austenite, and duplex with temperature. Source: Ref 4
in toughness level and a concurrent increase in
transition temperature. The combined effect of
cold work and α′ can be seen in Fig. 12. Lean
alloys such as 2001, 2003, and 2101 have a
much slower rate of formation of α′ and are
much less at risk for loss of toughness from exposure in the 300 to 600 °C range, as was
shown in Fig. 11.
Fatigue. Fatigue tests on duplex stainless
steels indicate that they possess a fatigue limit
of about 50% of the yield strength when tested
in air (Ref 4). The ratio of the fatigue strength in
a hostile environment to that in air is a useful
measure of the complementary strong points of
the duplex grades (i.e., strength and corrosion
resistance). Figure 13 shows that ratio for various alloys plotted versus their PREN. As an
alloy’s resistance to corrosive attack increases,
its fatigue limit in a given environment approaches that in air, indicating simply that corrosion plays an increasingly small role in fatigue crack propagation as corrosion resistance
increases. While this is intuitively reasonable, it
is not diminished because the duplex reward the
user with a higher level of yield strength and fatigue strength in air, so the net useful strength
under cyclic loading is much greater than that of
equivalent-PREN austenitic alloys.
Fig. 12 Increase in transition temperature with α′ formation
with aging for (a) annealed 2705 and (b) coldworked 2205. Source: Ref 4
Chapter 7: Duplex Stainless Steels / 99
Fig. 13 Influence of pitting resistance equivalent number (PREN) to fatigue strength in NaCl solution versus in air. Source: Ref 2
Forming and Machining
The higher strength and lower ductility of the
duplex grades compared to austenitics gives
them correspondingly less ability to be cold
formed. Duplex alloys have sufficient ductility
to be cold drawn; they behave like ferritics or
austenitics of similar alloy level. This, however,
is an alloy level at which excellent formability
is seldom expected. Nevertheless, duplex alloys
can be cold formed like austenitic alloys. Operations such as bending, drawing, and pressing
can readily be performed. Bend radii should be
at least twice sheet thickness. Tubing can be expanded into tube sheets, but care must be taken
to produce tight roller-expanded joints. Tubing
bend radii should be at least twice tubing outside diameter (OD).
Heavily formed sections should be fully
annealed, not just stress relieved, whenever
there is a potential for SCC in the service
environment.
Corrosion Resistance
Because duplex alloys are made up of two
phases, ferrite and austenite, each must carry its
own weight in resisting corrosion. Early alloys
that were lacking in nitrogen generally had a ferrite phase that, because of the greater partitioning of the chromium and molybdenum to the ferrite, had higher corrosion resistance than the
austenite. As nitrogen is added, it enriches the
austenite phase preferentially until the corrosion
resistance of the austenite phase reaches that of
100 / Stainless Steels for Design Engineers
the ferrite. This approach is common to all more
recently developed alloys starting with the revision of 2205 from UNS S31803 to S32205,
which has primarily higher nitrogen. The net result is a type of alloy that has most of the highly
desirable corrosion resistance characteristics of
superferritic grades without their limiting lack of
mechanical properties, mainly toughness.
The duplex alloys offer important advantages
in performance over the austenitic grades in a
number of significant aggressive media, including sulfuric acid, hydrochloric acid, sodium hydroxide, phosphoric acid, and organic acids.
This performance extends to situations in which
the aggressiveness of these media is enhanced
by contamination.
Sulfuric Acid. Figure 14 shows the behavior
of S32304 compared to 304 and 316 in sulfuric
acid. Figure 15 shows additional, more highly
alloyed duplex grades. The use of copper as an
alloying element in S32550 (1.5%) and S32760
(0.5%) gives them much better performance
than the otherwise similar S32750.
In real-life situations, such as seen in flue gas
desulfurization, sulfuric acid can be contaminated with chlorides. While this contamination
is deadly to 316 and 317, it has only a minor effect on the copper-alloyed duplexes (Fig. 16).
Hydrochloric Acid. Historically stainless
steels have had their poorest performance when
confronted by hydrochloric acid. Here again, the
Fig. 14 The 0.1 mm isocorrosion curves. Source: Ref 5
Fig. 15 The 0.1 mm isocorrosion curves. Source: Ref 5
General Corrosion
Fig. 16 Isocorrosion (0.1 mm/yr) performances of several austenitic and duplex alloys. Source: Ref 6
Chapter 7: Duplex Stainless Steels / 101
copper/tungsten-alloyed duplexes show exceptionally good performance, as seen in Fig. 17.
This extends the usefulness of stainless steels to
an environment that had previously been off
limits. Indeed, the duplex stainless steels in general can be said to be relatively indifferent to the
pH of chloride solutions and are affected rather
by the chloride concentration and temperature.
Nitric Acid. It is fairly well known and accepted that resistance to nitric acid, which was
one of the first uses of stainless steel, depends
almost entirely on the chromium content.
Molybdenum, in all other instances a very beneficial alloying element, has a strongly negative
influence on resistance to this highly oxidizing
acid. Consequently, only the leanest-molybdenum duplex alloys, such as S32304, should be
considered for use with nitric acid, and even
then no advantage can be claimed.
Sodium Hydroxide. Much of the older published data on the behavior of stainless steel has
seemed to promote the notion that higher nickel
levels were beneficial in strong bases. There
seems now to be little to support that notion.
Figures 18 and 19 clearly indicate, respectively,
that the duplex alloys with their relatively low
nickel levels significantly outperform the higher
nickel 304L and 316L, with performance improving with increasing chromium content. The
advantage is magnified when the environment is
contaminated with chlorides, as is the case of
the white liquors of kraft digesters.
Phosphoric Acid. While pure phosphoric
acid is not a very corrosive medium for stainless
Fig. 17 Isocorrosion (0.1 mm/yr) performance of duplex in HCl compared to 316L. Source: Ref 6
Fig. 18 Corrosion rates in boiling NaOH. Source: Ref 7
Fig. 19 Corrosion rates in white liquors plus chlorides.
Source: Ref 8
102 / Stainless Steels for Design Engineers
Fig. 20 Minimum temperatures for wet phosphoric acid (WPA) with an isocorrosion rate of 0.127 mm/yr. Source: Ref 9
In combinations of acetic and formic acid, the
superiority of duplex alloys is quite evident, as
seen in Fig. 22. S32750 shows virtual immunity, while in mixtures contaminated with
halides its performance ranks very closely to
expensive nickel-based superalloys such as
N06625 and N06455. Even the lower alloyed
S32205 can offer an order of magnitude improvement over S31703 in hot contaminated
acetic acid.
Pitting Corrosion
Fig. 21 Isocorrosion (0.1 mm/yr) performances of various
alloys. Source: Ref 9
steel, contaminants again can render it so.
Halides are particularly common and aggressive
contaminants. Figure 20 shows the substantial
improvement in performance of the duplex alloys over 316L when contaminants are present.
Performance again improves with increasing
chromium, molybdenum, and nitrogen levels.
Organic Acids. Duplex alloys perform particularly well in organic acids and have an excellent
record in industrial plants. In acetic acid, 304L
handles lower temperatures and concentrations.
Alloys such as S32205 perform well. In
formic acid, the most aggressive organic acid,
S32750 is resistant at all concentrations almost
to the boiling point, outperforming even titanium (see Fig. 21).
The different analysis of the two main phases
in duplex alloys means that each has its own pitting resistance equivalent number, PREN. The
ferrite phase has the relationship common to
ferritic grades:
PREN + %Cr + 3.3%Mo
(Eq 5)
while the austenite obeys the more familiar:
PREN + %Cr + 3.3%Mo + 30%N
(Eq 6)
The duplex grades partition these critical elements in such a way that the overall PREN of
most alloys comes out to be approximately Eq 1.
If one has the actual analysis of each phase,
then the proper relationship to use is Eq 2. These
relationships are incomplete in that they only address the major alloying elements. Tungsten has
half the value of molybdenum and is frequently
included:
PREN = %Cr + 3.3(%Mo + 0.5 × %W) + 16%N
(Eq 7)
Chapter 7: Duplex Stainless Steels / 103
Fig. 22 Corrosion rates for various alloys in 50% acetic plus formic acid, boiling. Source: Ref 10
If nonwrought material is involved, as in ascast and welded alloys, these relationships
greatly overstate PREN. This is because nonequilibrium-diminished
concentrations
of
chromium are often found around precipitates,
especially (manganese, chromium) S inclusions
(Ref 11, 12) and because of lower alloy content
locally due to solidification segregation, principally of molybdenum. This is most significant
in welded tubing, which can have higher sulfur
levels to increase weld penetration. Tube welds
can be reequilibrated by high-temperature annealing, but field girth welds will show diminished corrosion resistance if unannealed. So,
untreated welds can have PREN’s 5 to 15 lower
than the parent alloy, which equates to the localized lowering of chromium levels. The critical
pitting temperature (CPT) of welds often decreases to near the critical crevice corrosion
temperature (CCT) of the parent metal.
The precipitation of chromium- or molybdenum-rich second (third, in this case) phases,
such as σ or α′ inevitably results in diminishment of these key alloying elements in the region surrounding the precipitate, which will
make it more prone to localized corrosion.
This can also occur when secondary austenite forms during the heating of alloys to high
temperatures. This austenite, which forms from
ferrite, has very little nitrogen, which clearly
lowers its pitting corrosion resistance.
The duplex alloys stand up very well in comparison to corresponding superaustenitic alloys.
Figure 23 shows how CPT varies with PREN.
This ranking is not always linear, as Fig. 24
shows, with pitting potential dropping fairly
rapidly with temperature and at different rates
for different alloys. Figures 25 and 26 show the
influence of pH and chloride concentration, respectively. In 3% NaCl (Fig. 26), the rankings
show a minor variation with pH and a rational
relationship to alloy content. The influence of
chloride concentration is strong over a wide
range of concentrations.
These tests are best for judging relative performance of alloys and must be used cautiously
when extrapolating lab results to service performance. The degree to which short-term tests,
whether potentiostatic or strictly immersion, reflect long-term performance has not been well
established.
Crevice Corrosion
Crevices exist both by design and inadvertently. Crevices are occluded volumes of liquid
104 / Stainless Steels for Design Engineers
Fig. 23 Critical pitting temperature in seawater measured potentiostatically versus pitting resistance equivalent number (PREN).
Source: Ref 13
Fig. 24 Variation of pitting potential with temperature. SCE, saturated calomel electrode. Source: Ref 14
in which oxygen and corrosion products reach
levels quite different from the exterior environment and become highly corrosive. Thus, the
tighter the crevice is, the greater the restriction
of diffusion between the crevice and the bulk
and therefore the greater the chance of crevice
corrosion occurring. An alloy’s susceptibility to
crevice corrosion is proportional to its resistance to pitting corrosion under the same conditions. The CCT is lower than the CPT by about
10 to 30 °C.
The difference increases with total alloy content, as can be seen in Fig. 27. Interestingly, the
difference is approximately the same as is the
difference in CPT between the wrought alloy
and the welded alloy.
Stress Corrosion Cracking
Stress corrosion cracking (SCC) has long
been the Achilles’ heel of stainless steels. Only
soft ferritic stainless steels are immune to it. It
Chapter 7: Duplex Stainless Steels / 105
Fig. 25 Variation of critical pitting temperature (CPT) with
pH. Source: Ref 14
Fig. 26 Critical pitting temperature (CPT) as a function
of NaCl concentration. SCE, saturated calomel
electrode. Source: Ref 5
Fig. 27 Critical crevice temperature (CCT) and critical pitting temperature (CPT). Source: Ref 15
occurs at temperatures and in environments
where stainless would be a perfect material if
only it did not stress corrosion crack. The
arrival of duplex stainless steels has to a very
large degree ameliorated, if not solved, that
problem.
SCC is unfortunately poorly understood. Like
pitting, whose initiation mechanism has not
been identified, SCC has both its initiation and
propagation mechanisms still open to debate.
But the duplex alloys have good strength
mainly through fine grain size and solid solution
hardening, which seems to avoid the hydrogentrapping dislocation types that seem to be associated with hydrogen failures. So, while we cannot state the mechanism for SCC, we can map
out the conditions under which duplex alloys
are susceptible to SCC. The major environmental factors that affect SCC are chloride concentration and temperature. Figure 28 shows the
remarkable advantage the duplex alloys have
over the comparable austenitic alloys with regard to the temperatures at which they may be
used without SCC. The duplex alloys in this regard are governed in their behavior by their ferrite matrix, through which cracks must propagate (Ref 16). Ferritic stainless steels are known
for their resistance to SCC in the annealed condition. The advantage of the duplex lies in their
composite-type microstructure with the crackarresting austenite phase and the toughening
fine grain structure. The duplex alloys show a
higher threshold stress for SCC as a percentage
of their yield strength (Fig. 29) than austenitic
alloys. This is in spite of their higher yield
strength, again giving these alloys more usable
strength.
In ferrite, SCC susceptibility is a maximum
below 100 °C, while in austenite susceptibility
appears to begin around 50 °C and increase
monotonically with temperature. The temperature at which SCC occurs at the fastest rate increases with nickel content. This is also characteristic of ferritic and martensitic materials
and mirrors their hydrogen embrittlement behavior. H2S also accelerates failure in chloride
environments (Fig. 30), and cold work accelerates failure and lowers threshold stress values.
While duplex alloys behave in many regards
like ferritic alloys in their SCC or hydrogen
embrittlement response, they do not have the
same relationship between susceptibility and
bulk hardness. Other ferritic and martensitic
alloys display pronounced susceptibility to
these failures modes when their hardness
106 / Stainless Steels for Design Engineers
Fig. 28 Stress corrosion cracking (SCC) in neutral aerated NaCl. Testing duration 1000 hr. Source: Ref 5
Fig. 30 Suggested chloride and pH limits for cold-worked
duplex alloys. Source: Ref 17
Fig. 29 Constant load stress corrosion cracking (SCC) tests in
aerated MgCl2 at 150 °C. Source: Ref 5
exceeds Rc 22. The duplex alloys have annealed hardness over Rc 30 without being in
danger. This probably simply indicates that
hardness as a measure of susceptibility is valid
only insofar as it reflects a certain yield
strength threshold as it does in tempered
martensite and is not valid for ferrite/austenite
composite structures. Thus, it is very important
to understand duplex SCC behavior as a separate study and not interpret it in terms of
austenitic or martensitic SCC.
Chapter 7: Duplex Stainless Steels / 107
REFERENCES
1. P. Lacombe, B. Baroux, and G. Beranger,
Stainless Steels, Les Editions de Physique,
2003
2. R.N. Gunn, Duplex Stainless Steels, Abington Publishing, 1997, p 28
3. H. Nordberg H, Innovation of Stainless
Steel, Conf. Proc., AIM, Florence, 1994, p
2.217–2.229
4. Charles, Duplex Stainless Steels ’91, Vol 1,
Beaune, Les Editions de Physique, 1991, p
3–48
5. S. Bernhardsson, Duplex Stainless Steels
’91, Vol 1, Beaune, Les Editions de
Physique, 1991, p 137–150
6. J. Nichols J, 12th International Corrosion
Congress, Houston, NACE, p 1237
7. E.-M. Horn, Werkstoffe und Korrosion, Vol
42, 1991, p 511–519
8. J.P. Audouard, Stainless Steel Europe, April
1992, p 45
9. Avesta Sheffield, Corrosion Handbook for
Stainless Steels, 1994
10. B. Walden et al., Stainless Steel ’93,
Florence, AIM, 1993, p 3.47
11. M.F. McGuire, MS&T Conf. Proc., 2004, p
831–846
12. M. Ryan, D. Williams, R. Chater, B.
Hutton, and D. McPhail, Why Stainless
Steel Corrodes, Nature, Vol 412, p 770
13. C.V. Roscoe et al., Duplex Stainless Steels
’86, The Hague, Nederlands Instituut voor
Lasteckniek, 1986, p 126–135
14. J.M. Drugli et al., Paper 270 presented at
Corrosion ’90, Las Vegas, NACE, 1990
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Steel ’87, The Institute of Metals, 1988, p
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Stainless Steels for Design Engineers
Michael F. McGuire, p 109-122
DOI: 10.1361/ssde2008p109
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
CHAPTER 8
Ferritic Stainless Steels
Summary
THE FERRITIC STAINLESS STEELS are
the lowest-cost highly corrosion- and oxidationresisting alloys in existence. They are useful
mainly as light-gauge sheet since their toughness drops off rapidly for heavier sections. Even
as they have grown in use more than any other
type of stainless, they could still economically
displace the popular but expensive 304 for
many routine applications.
Introduction
Ferritic stainless steels are simplest, lowestcost stainless steels. In their minimal form, they
contain simply enough chromium to overcome
their inherent level of carbon impurity and hit
the 11% chromium in solution required for
“stainlessness.” Early in the 20th century, 430
came into being, and the attainable levels of
carbon removal required 16% chromium for
this to occur. So much extra chromium was required because during annealing, to develop the
fully ferritic structure, carbon combines with
chromium, rendering it useless as a corrosion
fighter. In October 1967, the first commercial
use of argon oxygen decarburization (AOD)
changed the world for ferritic stainless steel.
This process, in which argon and oxygen are
blown through the molten metal to selectively
remove carbon without removing chromium
(described in detail elsewhere in this book), reduced the carbon plus nitrogen levels sufficiently that their effect could be nearly negated
by small additions of titanium or niobium,
which combine strongly with carbon and nitrogen and effectively remove them from solution.
This process is called stabilization, and the
technology was documented long before AOD
was invented (Ref 1). It was not until carbon
and nitrogen levels were brought down to AOD
levels that it became truly practical for ferritic
alloys. The level of carbon plus nitrogen was
lowered from around 0.10% to around 0.04%,
and less-expensive high-carbon ferrochromium
could be used instead of expensive low-carbon
versions. Thus, there exist two types of ferritics:
the early high-carbon types such as 430, 434,
436, and 446 and the more modern stabilized alloys led by 409 and 439.
The older, unstabilized grades are not always
fully ferritic. Their carbon levels cause them to
form some high-temperature austenite, which
transforms to martensite if quenched. This
makes their welds brittle. To be used, they are
normally in the annealed condition, which requires a lengthy subcritical anneal to avoid
martensite and to evenly distribute chromium
after all carbides have stably formed. The newer
stabilized alloys behave as if they are interstitial
free. They are ferritic at all temperatures (excluding for the moment the possibility of extraneous phases such as (α' and σ) and can be easily welded without fear of unwanted phases.
Stabilization does not preclude excessive grain
growth in the fusion or heat-affected zone
(HAZ) of welds, which can render them brittle.
The mechanical properties of ferritic stainless
steels appear similar to austenitics strengthwise,
but they lack the ductility of austenitics, and
they are limited at low temperatures by brittleness and at high temperatures by softness.
The lower thermal expansion coefficient of
ferritics makes their scale more compatible with
the base alloy and provides them with a lesser
tendency to spall. This makes them excellent
for high-temperature applications with thermal
cycles, provided their strength is adequate.
110 / Stainless Steels for Design Engineers
The corrosion resistance of ferritics is hampered by their inability to utilize nitrogen. The
absence of nickel, which characterizes these alloys, is not a problem since nickel adds little to
corrosion resistance. The titanium stabilization
of the modern alloys has quite a beneficial effect
since titanium is a powerful deoxidizer and
desulfurizer, both of which can cause local
chromium depletion and pitting. Ferritics, moreover, are essentially free from stress-corrosion
cracking (SCC) since they are below the threshold hardness for hydrogen embrittlement in
body-centered cubic (bcc) ferrous alloys. There
are a few exceptions.
The main attraction of ferritic stainless steels
over austenitics is their cost. The old comparison of 430 versus 304 is a bit unfair since 304 is
richer in chromium. A fair comparison might be
between 439 and 304. The corrosion resistance
of these two alloys is barely distinguishable
under normal ambient conditions. They are both
very formable and weldable. The vast majority
of the objects made commercially from 304
could be switched to 439 with no adverse consequence. But, if nickel is selling for $7 per
pound, then the total cost of 304 versus 439 is
doubled by its presence. No design engineer can
afford to ignore this level of incentive to learn
to use ferritic stainless steels.
Ferritic Stainless Alloys
The ferritic stainless alloys generally group in
low (10.5 to 12.0%), medium (16 to 19%), and
high (greater than 25%) chromium. They can be
stabilized or not. These distinctions are somewhat imposed after the fact. Rather than giving
them an order that they truly do not possess, the
most significant alloys are all listed in Table 1
with their compositions.
The low-chromium ferritic stainless steels
began with the development of MF-1, the predecessor of 409, in the 1960s. Its excellent corrosion resistance, compared to carbon steel; relatively low cost; good welding; and formability
permitted it to replace aluminized carbon steel
and cast iron in automotive exhaust systems,
opening up what eventually became the largest
single market for stainless steel. It was made
possible by the very low carbon plus nitrogen
levels the AOD process provided and the use of
stabilization. Thus, 409 was an improvement on
405 in which aluminum performed a quasi stabilization, and low carbon suppressed martensite.
A similar predecessor was 410S, a low-carbon
version of 410 to which some understabilizing
amount of titanium is added but that still requires annealing for full ferritic properties. The
key issue of the 11% chromium ferritics is how
to deal with carbon and nitrogen. The 405 and
410S take the approach of minimizing it and live
with annealing. The 409 uses full titanium stabilization. The hidden problem with using only titanium is that unless nitrogen levels are made
very low, the amount of titanium required to
combine with it can reach levels at which the
first TiN precipitates in the molten metal. This
slaggy precipitate agglomerates, causing casting
problems and surface defects. This gave 409 a
reputation as a grade unsuitable for applications
that required good appearance because the titanium streaks were difficult to avoid and greatly
highlighted by polishing. This has largely been
overcome by better refining techniques to reduce
carbon plus nitrogen to levels below 0.02% and
the use of dual stabilization by titanium and niobium; 468 (UNS S40930) is such an alloy.
The historical archetype of ferritic stainless
steels was 430, which has existed since the
1920s and is still widely used. Its drawbacks are
lack of weldability, relatively poor corrosion resistance because so much of its chromium is
tied up as carbides, and modest formability. The
new archetype for this medium-chromium level
is 439. With 17% chromium and single (439) or
dual stabilization (468), this alloy overcomes
the problems of 430 and can readily replace 304
in most applications with significant cost savings. In North America, 439 is mainly used as a
higher-temperature automotive exhaust alloy,
but in Europe 430Ti is used extensively in more
visually challenging applications, such as appliances. There, it is generally used instead of 439
whenever the part can be designed to be formed
from it.
Now, 434 and 436 are little used as their historical application in automotive trim finds little
place in today’s automotive styling. A modern
offshoot of these alloys, which are basically
molybdenum enhanced 430, is 444. This alloy
has roughly the corrosion resistance of 316L but
is fully resistant to SCC in the welded or annealed condition. This makes it especially useful
for applications such as hot water heaters, heat
exchangers, and food- and beverage-processing
equipment.
Both the nominally 11 and 18% chromium
alloys are sometimes modified to enhance
their high-temperature strength or oxidation
0.01
...
0.07
0.03
0.01
0.03
S40910
S40920
AK alloy
S40930
S40940
S40975
AK alloy
typical
AK alloy
typical
ATI alloy
typical
ATI alloy
typical
Outukumpu
typical
S42900
S43000
S43020
S43023
S43036
S43035
S43932
ATI, AK
alloys
S46800
AK alloy
typical
AK alloy
typical
Outukumpu
typical
409
409
409 ultraform
466
409Cb
409Ni
11 Cr-Cb
430Se
430Ti
439
439LT
439 HP
439 ultraform
468
18 Cr-Cb
4742
18SR
429
430
430F
4724
Alfa II
Alfa I
12 SR
0.08
S40900
409
0.02
0.02
...
...
0.08
...
0.015
0.02
0.1
0.04
0.03
...
0.04
0.12
...
...
...
0.12
0.12
0.12
...
...
0.025
0.08
18
17.3
18
18.0–20.0
17.5
17.0–19.0
17.0–19
16.0–19.5
16.0–18.0
14.0–16.0
16.0–18.0
16.0–18.0
13.5
13
13
12
0.015
...
11.35
10.5–11.7
10.5–11.75
10.5–11.7
10.5–11.7
10.5–11.75
10.5–11.7
10.5–11.75
11.5–14.5
12.0–13.0
Cr
0.015
0.03
0.025
0.02
0.01
0.03
0.02
...
0.03
0.03
0.02
0.06
0.03
0.03
...
...
...
0.08
0.05
S40500
AK alloy
405
400
N
C
Designation
Ferritic stainless compositions
Alloy
Table 1
...
0.25
0.7
0.3
1
0.3
...
0.35
1
1
1
1.25
1
1
1.25
0.7
0.035
0.035
...
0.25
1
0.5
0.2
0.5
0.5
1
...
0.75
0.75
...
...
...
...
...
0.2
0.5–1
1
1
0.75
0.5
0.05
0.5
1
0.5
1
1
0.5
1
1
Mn
0.5
0.6
...
Ni
...
1.3
...
0.45
(continued)
...
...
...
...
...
0.45
1
...
...
...
...
...
...
0.04
...
0.25
...
...
0.25
Ti + Nb: 0.20 + 4x(C + N) to 1.10
0.35
0.20 + 4x(C+N) to 1.10
0.20 + 4x(C+N) to 0.75 Ti+Nb
...
...
0.55
...
...
...
...
...
...
...
0.20 + 4x(C+N) to 1.10
...
...
...
...
...
...
...
...
...
0.4
...
0.4
0.35
0.6
...
...
...
...
...
...
0.17
...
...
...
Nb
0.3
...
0.03
0.001
0.03
0.03
0.04
0.04
0.02
0.03
0.03
0.3
0.15
min
0.06
...
...
...
...
...
...
0.8 + 8x(C+N) Ti + Nb
10xC to 0.75 Nb
0.01
0.04
0.03
8x(C + N)
8x(C + N) to 0.15–0.50
6x(C + N) to 0.5
6x(C + N) to 0.75
...
...
Ti
0.01
0.02
0.02
0.45
0.03
0.03
S
0.04
0.06
0.04
0.04
0.06
...
...
...
...
...
...
...
...
...
0.04
0.04
0.04
0.04
0.04
0.04
0.45
0.04
0.03
P
...
...
...
...
...
...
...
...
...
...
...
...
...
Mo
1
1
1
1
1
1
1
1
0.03
0.03
...
1.3
1
1
1
1
1
1
1
1
1
Si
1.7
Al
1
Al
...
...
...
...
0.15 Al
0.15 Al
0.15 Se
...
...
...
1 Al
4 Al
3 Al
1.2 Al
...
...
...
...
...
...
...
...
0.10–0.30 Al
0.25 Al
Other
Chapter 8: Ferritic Stainless Steels / 111
S43400
S43600
S44100
S44200
ATI alloy
typical
S44400
442
436S
S44627
S44635
S44660
S44735
S44600
Cast alloy
4762
E-Brite,
26-1
Monit
Sea-cure
29-4C
446
CC-50
453
ATI alloy
typical
Outukumpu
typical
ATI alloy
typical
433
444, YUS
190-EM
Designation
434
436
441,
4509,
430J1L
0.025
0.025
0.025
0.2
0.5
0.01
0.03
0.08
0.01
0.025
0.2
0.01
0.12
0.12
0.03
C
0.035
0.035
...
0.25
...
0.015
...
...
...
0.035
...
0.015
...
...
...
N
Ferritic
stainless compositions
(continued)
Alloy
Table 1
Cr
24.5–26.0
25.0–27.0
28.0–30
23.0–27.0
26.0–30
25.0–27.5
22
24
20
17.5–19.5
18.0–23.0
17.3
16.0–18.0
16.0–18.8
17.5–18.5
Ni
3.5–4.5
1.5–3.5
0.5
0.6
4
0.5
0.3
...
0.25
1
0.6
0.3
...
...
...
Mn
1
1
1
1.5
1
0.4
0.3
0.7
0.3
1
1
0.2
1
1
1
Si
0.75
1
0.75
1
1.5
0.4
0.3
1.4
0.4
1
1
0.4
1
1
1
Mo
3.5–4.5
2.5–3.5
3.5–4.5
...
...
0.75–1.25
...
...
...
0.75–1.25
...
1.2
0.75–1.25
0.75–1.25
...
P
0.04
0.04
0.04
0.04
...
0.02
0.02
...
0.02
0.04
0.04
0.02
0.04
0.04
0.04
S
...
...
Ti + Nb: 0.20 + 4x(C + N) to 0.80
Ti + Nb: 0.20 + 4x(C + N) to 0.80
Ti + Nb: 0.20 + 4x(C + N) to 0.80
...
...
0.03
0.03
0.03
0.03
...
0.02
...
...
Ti + Nb: 0.20 + 4x(C + N) to 0.80
8x(C + N) min
0.1-0.6
Ti
...
...
0.02
0.03
...
0.001
0.03
0.04
0.001
0.03
0.03
0.015
...
...
...
...
...
...
...
...
0.5–0.20
...
...
Other
1.5
Al
0.60 Al
0.1
REM
0.2 Cu
0.5
Cu + Ni
...
...
...
...
...
10x(C + N) . . .
...
...
...
...
...
Nb + Ta 5xC 0.7 . . .
9xC 0.3–1.0 . . .
Nb
112 / Stainless Steels for Design Engineers
Chapter 8: Ferritic Stainless Steels / 113
resistance. Again, the driving force has been
the requirements of the hot end of exhaust
systems (e.g., exhaust manifolds). Alloying
with niobium and molybdenum adds to the
high-temperature strength, while additions of
chromium, silicon, and aluminum increase oxidation resistance. There exists an array of
proprietary alloys as shown in Table 1; these
are usually developed for specific automotive
needs and employ all or some of these alloying variations. The use of silicon and aluminum decreases formability and can accelerate (formation, so their use involves
trade-offs.
Alloys with more than 20% chromium are
used specifically for high-oxidation or corrosion
resistance. Despite the relative lack of hightemperature strength, these alloys are particularly useful because of their high-oxidation resistance, which they derive from the tight
adherence of their oxide scale. The close match
between the thermal expansion coefficient of
the scale and the alloy prevents spallation of the
oxide, which would lead to breakaway oxidation. This was the purpose of the earliest highchromium ferritic stainless, 446. The performance of 446 has been exceeded by lower alloyed
grades, such as the aluminum-alloyed ferritics.
A prime example of the state of the art is 453,
which has not only 22% chromium and 0.6%
aluminum but also rare earths in trace amounts
(i.e., 0.1%).
As in austenitic alloys, rare earths act as very
powerful oxide and sulfide formers that concentrate at the metal-oxide interface and stabilize it,
again preventing spallation. This type of alloys
finds use in high-temperature applications such
as planar oxide fuel cells.
The high-chromium alloys, when used for
corrosion resistance, are usually called superferritics. In the 1960s, E-Brite® was developed.
To obtain high toughness, it was vacuum refined
to very low carbon plus nitrogen levels. It was
followed by the more capable 29-4®. Later, this
alloy was stabilized and became the still-popular AL 29-4C®. (E-Brite now has a new life as a
fuel cell material based on its oxidation resistance and very low thermal expansion coefficient.) These alloys saw success as replacements for 316L when SCC was a problem. This
alloy and its close neighbor SeaCure® are used
primarily in tubing where corrosion resistance
is most important. It was developed for welded
condenser tubing where seawater or brackish
water is involved. It is also used in heat ex-
changers and extensively in condensing portions of high-efficiency residential furnaces.
The lower-alloyed Seacure had a slight toughness advantage that permitted it to be used at
wall thicknesses of 1/16 in. when AL29-4C®
was too brittle. As with other ferritics, these alloys are generally only suitably tough when
used in thin section size (i.e., less than several
millimeters).
It is difficult to say ferritic stainless steels are
underutilized since they account for about half
the world’s production of stainless, but there are
many applications in which more expensive
austenitic stainless steels are used needlessly.
Ferritic stainless steels are a viable alternative
to nickel-bearing austenitics when thickness is 2
mm or less and drawing and bending instead of
stretch forming is permitted. There are many
applications where the longer corrosion life of
low-chromium ferritics should economically replace carbon steel, as they have in automotive
exhaust systems. There are no technical barriers
to obtaining these savings; design engineers
need to learn how to use these alloys.
Metallurgy of Ferritic Stainless Steels
Chromium stabilizes the ferritic structure at
high temperatures. Thus, above about 11%
chromium, austenite does not exist at any temperature in pure iron chromium alloys, as seen
in Fig. 1.
However, iron-chromium alloys devoid of
carbon are not practical, so early metallurgists
saw the diagram shown in Fig. 2 with the level
Fig. 1 Iron-chromium phase diagram from Thermocalc
114 / Stainless Steels for Design Engineers
of carbon at 0.20%, which represented the purity level attainable in arc furnace refining. Carbon is essentially insoluble in ferrite at ambient
temperatures, and carbides of chromium and
iron will form to the extent carbon is available.
Since carbon diffuses interstitially much more
rapidly than chromium can substitutionally,
chromium is combined in situ, especially along
grain boundaries, which are fast-diffusion
paths. This locally depletes chromium, and the
alloy is sensitized. This can be eliminated by a
Fig. 2 Iron-chromium phase diagram at 0.20% carbon
Fig. 3 Iron-chromium diagram at low carbon levels Source: Ref 2
sufficiently long homogenization anneal at a
low enough temperature that carbon and nitrogen have very little solubility. This is standard
in the processing of unstabilized ferritic stainless steels, such as 430. Rapid cooling of unstabilized alloys causes carbon and nitrogen to precipitate within grains. This severely embrittles
the material and does not avoid sensitization.
This is called high-temperature embrittlement
because it comes from putting carbon and nitrogen into solution at a high temperature and then
causing it to precipitate in a harmful manner.
These alloys were only ferritic at room temperature if they were given a subcritical anneal
to transform austenite to ferrite. Otherwise, at
room temperature they would be ferrite plus
martensite. There are alloys that are intended to
use a mixed ferrite/martensite structure, but
they are treated later as a variation from the normal ferritic alloys.
The introduction of AOD refining permitted
much lower levels of carbon, as seen in Fig. 3,
opening the door for fully ferritic stainless steels.
Carbon and nitrogen added together produce
about the same effect as carbon alone. So, unstabilized fully ferritic alloys are not feasible
below 20% chromium without extreme refining
techniques, such as electron beam refining,
which are not commercially viable for low-cost
alloys. Thus, nearly all modern ferritic alloys
Chapter 8: Ferritic Stainless Steels / 115
are “stabilized.” This means that a strong carbide former such as titanium or niobium is
added in sufficient quantity to combine with all
the carbon plus nitrogen, removing them from
solution.
These reactions are simply:
Ti + C = TiC
(Eq 1)
Ti + N + TiN
(Eq 2)
Nb + C = NbC
(Eq 3)
Nb + N = NbN
(Eq 4)
Titanium is the stronger getter for carbon and
nitrogen. The thermodynamic driving force for
carbide and nitride formation is given by
(Ti )(C) =
−7700
+ 2.75
T
(Eq 5)
(Ti )( N) =
−15790
+ 5.40
T
(Eq 6)
It must be noted that titanium has an even
higher affinity for oxygen and sulfur than for
carbon, so that the removal of carbon from solution is preceded by the removal of oxygen,
nitrogen, and sulfur in that order. This will be
seen to have a major influence on corrosion resistance as the MnS inclusions generally associated with the initiation of pitting are not
found in titanium-stabilized grades of normally low sulfur. In practice, the removal of
oxygen begins in the molten state with the formation of titanium sulfide and nitride and next
in the molten or solid state, depending on concentrations. It is desirable to keep sulfur and
nitrogen low enough that precipitation is in the
solid state so that precipitates do not agglomerate and cause large primary inclusions that
become unsightly surface defects. TiCS forms
in the solid state if sulfur is present; if not, TiC
forms. Essentially all carbon is removed from
solution below 1250 °C if carbon and nitrogen
are kept as low as possible and a stoichiometric amount of titanium is available (i.e., greater
than about four times the carbon plus nitrogen).
The stabilization formula in various specifications is more than four times the carbon plus
nitrogen because experimentally it has been
found that sometimes understabilization occurs. This is due to the influence of oxygen and
sulfur having prior compound formation with
the titanium and less importantly that kinetic
factors prevented TiC formation. The latter effect was real in early austenitic alloys, such as
321, leading to knife-line corrosion attack after
welding, but does not exist in low interstitial
ferritic alloys, which have much greater diffusion rates than austenitic alloys. But, since carbon mobility is quite high, it is not practical to
quench alloys quickly enough to prevent carbide precipitation as is possible in austenitics
(detailed in the Chapter 6, “Austenitic Stainless
Steels”). Figure 4 shows the time-temperaturetransformation (TTT) curve for an unstabilized
430-type alloy with carbon plus nitrogen of
0.08% (Ref 3).
Stabilization causes nonchromium carbides to
form at high temperatures, precluding chromium
carbide precipitation. The net effect is that
modern stabilized ferritic alloys behave as interstitial free and can be mapped using the pure
iron-chromium diagram shown in Fig. 1.
The rate of diffusion of carbon in ferrite is
around 100 times greater than that of carbon in
austenite. The solubility of carbon in ferrite is
vastly lower than it is in austenite. Because of
these factors, the heat treatments to avoid sensitization are essentially reversed. Carbon in
austenite can be retained in supersaturation for
extended periods of time. This is why austenitic
L grades do not sensitize even though they are
slightly supersaturated. Sensitization occurs
at higher levels of carbon by prolonged heating
at 600 to 850 °C. In ferritics, carbon cannot
be kept in supersaturation even by the most
rapid quenching, and sensitization is alleviated
by prolonged heating in the 600 to 850 °C
range to allow chromium to equalize where
carbide precipitation has previously made it
inhomogeneous.
Fig. 4 430 time-temperature-transformation (TTT) curve. K,
carbideSource: Ref 3
116 / Stainless Steels for Design Engineers
Ferritic alloys, like austenitic alloys, can form
intermetallic phases. The most prominent is σ,
which can be seen to form in higher-chromium
stainless steels (i.e., those with chromium plus
molybdenum of 20% or more). Formation of σ
occurs when such alloys are held between 500
and 800 °C; it is a hard, brittle tetragonal phase
with equal parts iron and chromium. Thus, its
formation causes chromium depletion of the adjoining ferrite. Formation requires substitutional
diffusion of chromium so is slower to form than
carbides, minutes rather than seconds. Since
cold work enhances substitutional diffusion, it
accelerates σ formation. The σ forms preferentially along grain boundaries for diffusion reasons, and this causes it to have a major embrittling effect. The σ may be redissolved by
solution annealing, but regaining full homogeneity is not immediate.
Another embrittling phenomenon is the formation of α'. This was named 885 °F or 475 °C
embrittlement before its cause was understood.
Before the nature of α' was known, it was confused with temper embrittlement, which occurs
in martensitic alloys at the same temperature.
Temper embrittlement is the segregation of
phosphorus to prior austenitic grain boundaries
and does not occur in fully ferritic alloys. The α'
is the ordered equiatomic chromium iron phase
that forms by spinodal decomposition; it has the
same composition as σ but exists at a lower
Fig. 5 Influence of α' formation on hardnessSource: Ref 4
temperature with the same structure as ferrite
but with the chromium and iron atoms in an ordered bcc matrix in which iron and chromium
occupy sites equivalent to two interlocking simple cubic matrices. Because the lattice so
closely matches that of ferrite, the precipitate is
coherent and causes hardening. The α' embrittlement causes an extreme loss of toughness as
well as hardening. It also causes a loss in corrosion resistance via the chromium depletion of
that part of the matrix that surrenders chromium
to the α'-phase.
Figures 5 and 6 show the hardening effect of
α' and the resulting loss of toughness, respectively (Ref 6).
Mechanical Behavior
Ferritic stainless steels are quite similar in
their mechanical behavior to carbon steel. The
main influence of chromium is to produce some
solid solution hardening. Let us review the
strengthening mechanisms of bcc iron. Pure
iron is an extremely soft material with a yield
strength well under 10,000 psi. This softness is
not seen in practice because steel is never pure.
Carbon has an extremely powerful effect on
hardening, as does nitrogen.
The influence of substitutional alloying elements is also quite significant. According to
Chapter 8: Ferritic Stainless Steels / 117
Fig. 6 Influence of α' formation on toughness Source: Ref 5
Paxton (Ref 7), the misfit of solute atoms causes
lattice strains proportional to the amount dissolved and provides strengthening through the
lattice friction term. This mechanism also increases the impact transition temperature unfavorably. Elements that produce a refining of
grain size are the exception to this general rule in
carbon steel, but the lack of an austenite-to-ferrite transformation in stabilized ferritic stainless
steels negates this benefit for them. Figure 7
shows that fairly common ingredients and impurities have strong hardening effects (Ref 6). Manganese and silicon are normally deoxidizers, but
in titanium-stabilized alloys, titanium takes over
the deoxidizing role so their presence can be limited. Phosphorus is virtually impossible to refine
from stainless steel, so its presence at around
0.02% is normally a given unless low-phosphorus raw materials are used as a starting point.
The worst toughness-inhibiting effects come
from interstitial elements to grain boundaries:
oxygen, carbon, and nitrogen. The effect of carbon plus nitrogen on transition temperature is
profound, as seen in Fig. 8 (Ref 8).
Stabilizing removes the interstitial carbon and
nitrogen, along with oxygen and sulfur, from
solution. This does not produce a major softening, however, because the precipitate itself has a
hardening effect.
Fig. 7 Influence of substitutional elements on hardness of
iron alloys
The softest ferritic stainless alloys are the
409 variations made for highly formed exhaust
system components. They contain as little
manganese, silicon, nickel, and other substitutional elements as possible and have a minimum of carbon plus nitrogen, so that the resulting precipitate fraction after titanium
addition is as low as possible. To maximize
softness and formability, titanium and niobium
in excess of that required for stabilization must
also be minimized as they will cause solid solution hardening.
118 / Stainless Steels for Design Engineers
Fig. 9 Corrosion of titanium-stabilized 29% Cr plus 4% Mo
alloys in ASTM A 763 Y test. Source: Ref 11
Fig. 8 Influence of interstitial carbon and nitrogen on toughness transition temperature Source: Ref 8
Stabilization
Stabilization is essential to ferritic stainless
steels to avoid the precipitation of grain boundary carbides. Combined carbon plus nitrogen
levels below 100 ppm are necessary to avoid
both sensitization and embrittlement, but without proper heat treatment even alloys of this purity can incur debilitating loss of toughness due
to carbide and nitride precipitates (Ref 9). These
levels are not economically attainable for commercial alloys, so stabilization is the correct engineering answer.
Stabilization is generally considered as the
simple gettering of carbon and nitrogen by a
suitable carbide and nitride former. It was not
known until about 1980 just what the mechanisms of embrittlement were in the ferritic
stainless steels, however. The distinguishing of
α' from those related to interstitials and their
stabilizers (Ref 10) permitted stabilizing elements to be optimized.
Titanium combines with carbon and nitrogen
stoichiometrically by:
Ti = 4 × C + 3.4 × N
(Eq 7)
Niobium requires a greater weight percentage:
Nb = 7.7 × C + 6.6 × N
(Eq 8)
Fig. 10 Corrosion of niobium-stabilized 29% Cr plus 4% Mo
alloys in ASTM A 763 Y test. Source: Ref 11
As titanium and niobium are added to alloys,
their corrosion resistance is improved (Figs. 9
and 10) (Ref 11). Maximum improvement in
corrosion resistance levels off once full stabilization is reached. Excess amounts of the stabilizing elements have negligible effect, but titanium-stabilized alloys have a lower rate of
corrosion than niobium-stabilized alloys. This is
probably due to titanium’s ability to eliminate
sulfur and oxygen from solution.
Toughness improves for niobium-stabilized
alloys up through full stabilization and then begins to decline. This is a result of excess stabilizing alloy acting as a solid solution hardener
and therefore a toughness reducer. This toughness reduction is more pronounced with titanium, which is a stronger solid solution hardener (Figs. 11 and 12) (Ref 11).
The upshot of this understanding was the introduction of dual stabilization, through which
both weld and base metal toughness and corrosion resistance are optimized. The same study
recommended that dual stabilization follow the
following formula:
Chapter 8: Ferritic Stainless Steels / 119
(Ti + Nb) ≥ 6 × (C + N)
Fig. 11 Charpy V-notch impact ductile to brittle transition
temperature (DBTT) of titanium-stabilized 29%Cr
plus 4%Mo alloys test. Source: Ref 11
Fig. 12 Charpy V-notch impact ductile to brittle transition
temperature (DBTT) of niobium-stabilized 29%Cr
plus 4%Mo alloys test. Source: Ref 11
(Eq 9)
The toughness of these alloys has a broad optimum that takes advantage of the corrosion-resisting benefits of titanium (Fig. 13) (Allegheny
Ludlum).
Other strong carbide formers such as zirconium and vanadium are ineffective stabilizers
because their mobility at the temperatures at
which they are thermodynamically capable of
forming sufficiently large percentages of carbides and nitrides is too low to rid the matrix of
these elements. They also have too great a tendency to form intermetallic compounds.
Toughness in ferritic stainless steels is a
major consideration. If ferritic alloys enjoyed
the same toughness as austenitic alloys, there
would be few instances when the use of the
much more expensive nickel-bearing grades
would be justified. Because stabilized alloys
are ferritic at all temperatures, there is no
automatic grain-refining transformation as exists in carbon steel. If grains grow large from
annealing at high temperatures or welding, then
the transition temperature increases. Section
size also has an effect. Stabilized ferritic stainless steels are seldom used in thicknesses of
over several millimeters because of decreasing
toughness. Figure 14 shows how transition
Fig. 13 Toughness of dual-stabilized low-alloy ferritic stainless. AL 466 is recognized as S40930
120 / Stainless Steels for Design Engineers
tion. But, the anisotropy does result in remarkable drawing characteristics, with ferritic
stainless steels with elongations in tensile tests
in the mid-30% range being nearly equal to
austenitic stainless steels with over 50% elongation.
The measure of anisotropy is the Lankford
ratio. It is expressed as:
r + 2r45 + r90
R= 0
4
Fig. 14 Change in transition temperature with thickness for
29Cr-4Mo-2Ni alloy. Source: Ref 12
temperature can increase with thickness (Ref
12). This effect is due simply to stress states
transitioning from biaxial to the more embrittling triaxial with increasing thickness.
Texture and Anisotropy
The deformation of ferritic bcc materials is
characterized by limited slip systems, high
stacking fault energy, and lattice anisotropy.
So, when ferritic stainless are deformed, dislocations tend not to dissociate as they do in
austenitic stainless steels. The lack of dissociation of dislocations encourages cross slip.
This minimizes dislocation tangles and work
hardening.
When ferritic stainless steels are deformed,
certain crystallographic slip systems predominate, so that large deformations mechanically
bring different grains via rotation into closer
crystallographic alignment. This preferred deformation along easier slip planes results
macroscopically in overall mechanical properties varying with direction with respect to the
prior deformation. Thus, ferritic stainless
steels, like low-carbon steels, have pronounced
mechanical anisotropy. This is manifest in
their deep drawing characteristics. Heavily
cold-rolled and annealed ferritic stainless
steels draw quite well. They resist thinning.
When elongated, they contract in the width direction while keeping virtually the same thickness. This same phenomenon means that they
cannot be stretch formed since plain strain
quickly results in fracture because of the resistance to deformation in the thickness direc-
(Eq 10)
When this expression equals 1, then a material is isotropic. As the value increases from 1,
the drawability increases, as measured by the
limiting drawing ratio (LDR), the ratio of the
diameter of a disk to that of the deepest cylinder
into which it can be drawn. The ferritic stainless
steels in sheet form have LDRs of around 2.2
compared to 2.0 for 304.
The good formability of ferritic stainless
steels has some drawbacks. They are subject to
ridging, which is the formation of visible ridges
parallel to the direction of elongation. This is an
artifact of texture in the material. A combination
of careful chemistry design and thermomechanical processing is required to keep it under control. The approach centers on variables that increase stored energy from deformation to
promote recrystallization over recovery during
annealing.
The ferritic stainless steels even carry forward some of the preferred grain orientation
that come from initial solidification when
growth of dendrites is along preferred crystallographic directions. Hot working merely reorients these similarly oriented grains en masse.
Without phase changes or enough stored energy
to provoke full recrystallization, randomness of
grain orientation is never achieved.
Titanium-stabilized steels show more texture
and recovery versus recrystallization than do
niobium-stabilized alloys. This is because titanium carbides and nitrides form at higher temperature and are therefore coarser. They thus
present less obstruction to dislocation motion
than finer niobium precipitates. Furthermore,
niobium precipitates tend to dissociate to a
greater degree than those of titanium. This puts
niobium in solution during hot working where it
can interact with dislocations. Thus, alloys at
least partially stabilized with niobium can
achieve greater recrystallization, which can
translate to finer grain size and less anisotropy.
Chapter 8: Ferritic Stainless Steels / 121
Boron additions to ferritic stainless steels
result in the formation of grain boundary carbides, M23(C, B)6. If added to titanium-stabilized steels, the carbides form on preexisting
TiN particles and result in coarser overall precipitate arrays since finer, lower-temperature
precipitating TiC or TiCS precipitates are at
least partially precluded. The net result is
coarser grain size and no major improvement
in mechanical properties over the use of titanium alone (Ref 13). Additions of boron to
niobium-stabilized steels does cause finer precipitates and grain size than would niobium
alone (Ref 14).
High-Temperature Properties
High-temperature mechanical properties of
ferritic stainless steels are often important to
their successful use because their oxidation resistance is excellent and better than austenitics,
but their high-temperature strength is lower
than that of austenitics. This has led to considerable development of high-temperature properties, primarily for the automotive market. Research has determined that high-temperature
strength and creep resistance are best served by
stabilizing grain size and having niobium in
solid solution. Adding titanium to niobium-stabilized steels stabilizes the type of carbide, especially preventing the formation of the coarse
M6C, whose growth decreases strength. The
relatively high insolubility of TiC causes this.
Niobium is concurrently made available for
high-temperature solid solution strengthening.
Corrosion and Oxidation Resistance
Corrosion resistance is chemistry dependent
rather than structure dependent, so ferritic stainless steels behave just as do other stainless
steels of the same crucial alloy content.
The main alloying elements that provide resistance to localized corrosion, general corrosion, and crevice corrosion are chromium,
molybdenum, and nitrogen. Since nitrogen is
essentially insoluble in ferrite, it cannot contribute to the corrosion resistance of ferritic
stainless steels as it can in austenite. Other alloying elements, such as copper and nickel, can
add to corrosion resistance in special cases, but
they are of secondary importance compared to
chromium and molybdenum. Likewise, other
elements can have a negative effect. Any element that can combine with chromium or
molybdenum can detract from corrosion resistance by their removal of these essential elements from solution. The most notorious of
these is carbon, whose tendency to form
chromium carbides causes areas around such
carbides to be partially depleted of chromium.
However, nitrogen, oxygen, and sulfur can also
form chromium compounds and cause localized
loss of corrosion resistance. Manganese sulfides, for instance, are almost always seen to be
the locus of pitting corrosion (Ref 15). More
careful examination has shown that such sulfides grow in the solid state as chromium/manganese sulfides and deplete their very close surroundings of chromium, inviting corrosion to
begin at the inclusion-matrix interface, where
chromium levels in solution are reduced (Ref
16).
Other factors that lead to loss of localized corrosion resistance are the formation of
chromium-rich phases such as α' and σ. Either
of these with about 50% chromium will cause
adjoining ferrite to have lower chromium levels.
Because ferrite has a non-close-packed structure, diffusion rates, both substitutional and interstitial, are about two orders of magnitude
higher than in austenite. That means that any
deleterious chromium-depleting reaction can
happen more rapidly. Alloys cannot be
quenched rapidly enough to forestall sensitization, the precipitation of chromium carbides
that depletes grain boundary regions of
chromium. Instead, carbon must be neutralized
by stabilization, or the chromium depletion
must be removed by homogenization in longbox anneals. Note that the latter technique is
also possible in austenitics but would require
annealing for excessively long times, 102 h or
so.
The ferritic stainless steels are valued for
their resistance to SCC. Even in environments
that cause pitting, the normal initiation step for
SCC, annealed ferritic stainless steels do not
undergo SCC as long as alloying elements such
as nickel, copper, and cobalt are kept below
0.5% in aggregate. Cold work sufficient to raise
their hardness above Rc 20 to 22 can make them
susceptible to both SCC and its cousin, hydrogen embrittlement. The more highly alloyed superferritic alloys are even susceptible to hydrogen embrittlement in the annealed condition
(Ref 17). As with martensitic stainless steels,
122 / Stainless Steels for Design Engineers
this susceptibility is a maximum near room temperature and declines with increasing temperature, as opposed to austenitics, which see their
maximum susceptibility above room temperature. This limits these alloys’ ability to employ
cathodic protection safely to –0.80 Vsce, at
which point corrosion in seawater is, if not
eliminated, reduced to very low levels (Ref 18).
REFERENCES
1. F.M. Beckett and R. Franks, Trans AIME,
Vol 113, 1934, p 126–143
2. Stainless Steel, Les Editions de Physiques,
1992, p 483
3. Stainless Steels, Les Editions de Physique,
2003
4. H.D. Newell, High Chromium Irons, Met.
Prog., April 1947, p 617–626
5. P.J. Grobner, The 885 °C (475 °C) Embrittlement of Ferritic Stainless Steels, Metall.
Trans., Vol 4, 1973, p 251–260
6. Handbook of Stainless Steels, Peckner and
Bernstein, McGraw Hill, 1977, p 5–9, 5–12
7. H.W. Paxton, Alloying, ASM, 1998, p 213
8. H. Abo et al., Stainless Steel ‘77
9. J. Grubb and R. Wright, The Role of C and
N in the Brittle Fracture of Fe-26Cr, Met.
Trans. A, Vol 10A, Sept 1979, p 1247–1255
10. J. Grubb, R. Wright, and P. Farrar, “Micromechanisms of Brittle Fracture in TitaniumStabilized Stainless Steels,” Special Publication 706, ASTM, 1980
11. J. Grubb, Stabilization of High-Chromium
Ferritic Stainless Steels, Proc. Int. Conf.
Stainless Steels, ISIJ, Chiba, 1991
12. M.A. Streicher, Stainless Steel ‘77, p 27
13. E. El-Kashif, K. Asakura, T. Koseki, and
K. Shibata, ISIJ Int., Vol 44, 2004, p
1568–1575
14. N. Fujita, K. Ohmura, E. Sato, and A.
Yamamoto, Nippon Technical Report 71,
Oct 1996
15. T. Suter, E. Webb, H. Bohni, and R. Alkire,
Pit Initiation in I M NaCl With and Without
Mechanical Stress, J. Electrochem. Soc.,
Vol 148 (No. 5), 2001, B174
16. M. Ryan, D. Williams, R. Chater, B. Hutton,
and D. McPhail, Why Stainless Steel Corrodes, Nature, Vol 412, 2002, p 770
17. J. Grubb, “Hydrogen Embrittlement of Superferritic Stainless Steels,” paper presented
at 1984 ASM Int’l Conference on New Developments in Stainless Steel Technology,
Detroit, September 1984
18. J. Grubb and J. Maurer, “Use of Cathodic
Protection With Superferritic Stainless
Steels in Seawater,” paper presented at Corrosion 84, New Orleans, April 1984
Stainless Steels for Design Engineers
Michael F. McGuire, p 123-135
DOI: 10.1361/ssde2008p123
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
CHAPTER 9
Martensitic Stainless Steels
Summary
THE SMALLEST CATEGORY of stainless
steels in usage volume is the martensitic stainless steels. This is mainly because these alloys
are limited in corrosion resistance because of
the necessity of keeping alloy levels low to produce the martensite structure. Even so, they fill
an important niche as a strong, hard, and tough
alloy of fairly good corrosion resistance and as
a strong, stable, high-temperature alloy.
Introduction
Nearly 100 years ago cutlery was first sold in
Great Britain with a composition of 13%
chromium and 0.25% carbon. This was the first
commercial use of stainless steel and cutlery
with the same basic analysis is still sold today.
The useful alloys of martensitic stainless steel
contain from roughly 11 to 18% chromium and
up to 1.0% carbon. Relatively small amounts of
nickel, molybdenum, tungsten, vanadium, and
niobium are also added at times for specific purposes explained in this chapter. Those martensitic stainless steels in which elements such as
copper and titanium are added to produce additional hardening through precipitation are discussed in Chapter 4, “Corrosion Types.”
The designers and engineers already familiar
with martensitic carbon and alloy steels will
find nothing confusing about martensitic stainless steels. There is no aspect of martensitic
steels that does not apply directly to stainless
martensitic steels. The additional concerns one
must have with stainless martensite relate
mainly to those that are due to the strong ferritizing influence of chromium. Chromium
strongly promotes the formation of ferrite,
which restricts the temperature and composition
ranges over which it is possible to obtain a fully
austenitic structure from which to form martensite. The presence of ferrite in a martensitic
structure is detrimental to strength, hardness,
and toughness. Ferrite can appear in the as-cast
structure and be formed during austenitizing or
tempering. All the usual concerns inherent in
any martensitic alloys are still present; temper
embrittlement, retained austenite, etc.
Martensitic stainless steels are the most marginally corrosion resistant of all the stainless alloys. The requirement that they be fully austenitizable limits the amount of corrosion-resisting
chromium and molybdenum they can contain.
Much of the carbon in them detracts from the effective chromium content by forming chromium
carbides. In addition they are always susceptible
to stress corrosion cracking (SCC) when their
hardness exceeds about Rc 22. These limitations
combine to make their excellent properties usable in only mild environments compared to
other stainless steels. Their high strength and
hardness for their relatively low cost ensure their
place as a very useful engineering material.
Table 1 lists the most significant of the
martensitic stainless steel alloys. The reader
should be aware that some alloys which are
quite similar are discussed primarily in other
chapters dealing with specifically PH stainless
steels or primarily ferritic stainless steels. The
distinction between martensitic stainless steels
and some other stainless alloy families is sometimes vague. Nearly all the precipitation-hardening stainless steels are used in the martensitic
state, but their special hardening mechanism of
precipitation within a martensitic matrix causes
them to be categorized separately somewhat arbitrarily. By this conventional logic, some of the
martensitic alloys containing molybdenum or
S41400
S41425
S41500
S41600
S41623
S41800
Wrought
S42000
Wrought
DIN
Wrought
1.4116 Nominal
S42020
Wrought
S42023
Wrought
S42200
Wrought
S42400
Wrought
Wrought
S42500
...
Wrought
...
Wrought
414
414 mod
415
416
416Se
418
420
4116
NT-CRS
HP13Cr-2
HP13Cr-1
JFE
Nominal
JFE
Nominal
Nippon
Nominal
S41003
412
420F
420FSe
422
424
425
425mod
Trinamet
S41040
410Cb
Wrought
Wrought
Wrought
Wrought
Wrought
Wrought
Wrought
Wrought
Wrought
Wrought
Wrought
S41008
410
Form
Wrought
Wrought
Wrought
S40300
S41000
S41003
UNS
0.03
0.025
0.025
0.15 min
0.15 min
0.20-0.25
0.06 max
0.08-0.20
0.50-0.55
0.30 max
0.15-0.20
0.15 min
0.5
0.05 max
0.15 max
0.15 max
0.15 max
0.05
0.030 max
0.18 max
0.08
0.15 max
0.15 max
0.03
C
Mn
...
1.45
0.45
0.45
1.25
1.25
1
0.50-1.00
1
1
1
0.5
1
0.50-1.00
1.25
1.25
1
0.5-1.0
1.5
1
1.5
1
1
1
Compositions (wt%) of martensitic stainless steels
403
410
410S
Alloy
Table 1
S
...
...
...
...
...
...
Si
...
...
...
1
1
0.75
0.30-0.60
1
1
1
0.5
1
0.6
1
1
1
0.6
1
1
1
0.5
1
(continued)
0.15 min
0.06
0.03
0.03
0.01
0.03
0.03
0.03
0.03
0.03
0.15-0.30
0.06
0.03
0.005
0.03
0.03
0.03
0.03
0.03
0.03
Cr
12.7
13
12.0-14.0
12.0-14.0
11.0-13.5
12.0-14.0
14.0-16.0
13.0-14.0
12.0
-14
13
11.5-13.5
11.5-13.5
10.512.5
11.513.5
11.513.5
10.512.5
11.5-13.5
12.015
11.5-14.0
12.0-14.0
12.014
12.0-14.0
12.0-14.0
14.5
...
...
1.4
2
0.6
0.6
0.75-1.25
0.30-0.70
0.30-0.70
0.80-1.20
1.00
-3
1
...
...
0.65
...
1.52
0.50-1.00
...
...
0.6
Mo
...
...
...
...
...
4.5
5
4
...
...
0.50-1.00
3.50-4.50
1.00-2.00
0.5
...
1.80-2.20
...
...
1.25-2.50
4.07
3.50-5.50
...
...
1.5
1.5
Ni
...
...
0.03
...
1.5 Cu
...
...
0.15 min Se
0.75-1.25 W
...
...
...
2.00
-3.00 Cu
...
0.05-0.30
Nb
0.030 max
N
...
0.06-0.12
N
...
...
0.15 min
Se
2.50-3.50 W
...
...
N
...
...
Other
0.040 N
...
...
...
...
0.15-0.30V
...
...
...
...
...
...
0.15 V
...
...
...
...
0.30 Cu
...
...
...
...
...
...
Other
124 / Stainless Steels for Design Engineers
KL-HP
12Cr
431
440A
440B
440C
440F
440FSe
BG-42
ATS-34
14-4
CrMo
154 CM
CPM
S30V
CPM
S60V
CPM
S90V
CA-15
CA 15M
CA-40
CA-40F
CB-6N
CB-6MN
CA-28MVW
KL-12Cr
NT-CRSS
Alloy
PM
PM
Nominal
Nominal
J91150
J91151
J91153
J91154
J91650
J91540
J91422
Wrought
PM
Nominal
Nominal
Cast
Cast
Cast
Cast
Cast
Cast
Cast
0.15 max
0.15 max
0.20-0.40
0.20-0.40
0.06 max
0.06 max
0.20-0.28
2.2
2.15
1.05
1.45
0.01
0.20 max
0.60-0.75
0.75-0.95
0.95-1.20
0.95-1.20
0.95-1.20
1.15
1.05
1.05
Wrought
Wrought
Wrought
Wrought
Wrought
Wrought
Wrought
Wrought
Wrought
Wrought
C
0.01
0.02
Wrought
Form
Wrought
Nippon
Nominal
JFE
Nominal
JFE
Nominal
S43100
S44002
S44003
S44004
S44020
S44023
Nominal
Nominal
Nominal
UNS
...
...
Mn
...
1
1
1
1
0.5
1
0.50-1.00
0.4
0.45
...
1
1
1
1
1.25
1.25
...
0.4
0.5
2
Table 1 (continued)
Compositions (wt%) of martensitic stainless steels
...
...
S
0.04
0.04
0.04
0.020-0.040
0.02
0.03
0.03
...
...
...
...
0.03
0.03
0.03
0.03
0.10-0.35
0.03
...
...
...
...
Si
...
...
...
1.5
0.65
1.5
1.5
1
1
1
0.3
1
1
1
1
1
1
0.3
0.35
0.3
...
...
...
Cr
11.5-14.0
11.5-14.0
11.5-14.0
11.5-14.0
10.50-12.50
11.50-14.00
11.00-12.50
13
17
14
14
15.0-17.0
16.0-18.0
16.0-18.0
16.0-18.0
16.0-18.0
16.0-18.0
14.5
14
14
12
11
12.3
...
Mo
0.40-1.00
0.90-1.25
0.5
0.15-1.00
0.5
0.5
1
0.4
4
2
...
0.75
0.75
0.75
0.40-0.60
0.6
4
4
4
2
2
Ni
1
1
1
1
1
3.50-4.50
0.50-1.00
...
...
...
...
1.25-2.50
...
...
...
0.75
0.75
...
...
...
5.5
2.4
5.8
Other
...
...
...
...
...
...
0.90-1.25 W
9.0 V
5.5 V
...
4.0 V
...
...
...
...
...
0.15 min Se
1.2 V
...
...
...
0.5 Cu
1.5 Cu
Other
...
...
...
...
...
...
0.20-0.30 V
...
...
...
...
...
...
...
...
...
...
...
...
0.010 N
0.010 N
0.015 N
Chapter 9: Martensitic Stainless Steels / 125
126 / Stainless Steels for Design Engineers
tungsten should also be considered precipitationhardening alloys, but they customarily are not
and will not be in this work.
The ferritic alloys often have compositions
that allow them to be partially martensitic under
some conditions. 430 (UNS S43000) and
3CR12 (UNS S41003) can contain some
martensite if their heat treatment is such that
austenite is allowed to form and is followed by
rapid cooling. Even 409 (UNS S409XX) can
form some austenite if chromium is at the high
end of its possible range and nickel and manganese residual levels are high. The martensitic
alloys themselves can be made to be partially
ferritic by forcing their carbon contents to low
levels as is customarily done with 410S (UNS
S41003). Not understanding these alloys can
lead to unexpected consequences in mechanical
properties or corrosion performance.
on temperature. The amount is given by the
Koistinen and Marburger equation (Ref 1):
1− Vα ′ = exp{β (Ms − T )}
(Eq 1)
The martensite is coherent with the parent
austenite and resembles the passage of slip dislocations through the crystal. The sum of many
such dislocations is shear, and this can be
macroscopically visible as in Fig. 1.
The formation of martensite is essentially mechanical (i.e., via deformation, not diffusion).
The shear and volume expansion, about 4%,
which accompanies the transformation, involves a great deal of strain energy that must be
taken into account. This is shown diagrammatically in Fig. 2 (Ref 2).
Martensite Formation
Martensite as a phenomenon deserves a brief
review. Martensite forms as result of the diffusionless transformation of austenite. The
austenite may be supersaturated with carbon or
nitrogen, but that is not necessary for the transformation. The driving force for the transformation is simply the much lower free energy of the
ferrite phase over the austenite phase, which
can be attributed largely to large mutual repulsion between iron atoms that possess unpaired
outer electrons with the same quantum number
and magnetic polarity. This free-energy differential increases with decreasing temperature. At
a certain temperature, the martensite start temperature Ms, the transformation occurs spontaneously via the coordinated movement of atoms
in a shearing-type mode at very high speeds approaching the speed of sound in the material.
The composition of the martensite is identical to
that of the parent austenite.
There is regularity to the relationship between
the parent austenite and the martensite.
Greninger and Troiano determined that the
close-packed planes of the austenite {111} varied from the {011} of the martensite by only
0.2°. Further, the direction of the <101 bar> of
the austenite was only 2.7° from the <1 bar 11
bar> of the martensite. These relationships define the habit plane that constitutes the austenite
martensite boundary.
Martensite forms essentially independent of
time and the fraction transformed depends only
Fig. 1 Martensite platelets emerging from the surface.
Source: Ref 2
Fig. 2 The martensite reaction ab contrasted to the nucleation
and growth-type transformation of austenite to ferrite, ac
Chapter 9: Martensitic Stainless Steels / 127
This energy differential between ferrite and
martensite is stored in the high-strain energy
matrix. Applied strains affect the transformation. Indeed, metastable austenite can readily be
transformed to martensite by deformation.
However, the untransformed austenite is hindered from transforming by the compression it
receives from the already-formed martensite.
Thus, some residual austenite is commonly
found between lathes of martensite.
At the Ms temperature, the body-centered
cubic (bcc) phase becomes preferable energetically, but this temperature is too low for diffusion transformation, and a slight shear in the
austenite lattice causes a rearrangement of the
atoms from a face-centered cubic (fcc) to a distorted bcc structure. The amount of distortion is
proportional to the amount of carbon in the interstices of the structure. These interstices are
considerably smaller in the bcc structure even
though it is expanded from the fcc. The octahedral sites change from 2.86 by 3.56 A to 2.86 by
2.86 A, as shown in Fig. 3 (Ref 3). The distortion is accommodated by accommodation from
site to site at low carbon levels, but above about
0.018% carbon this can no longer be accommodated and a tetragonal distortion occurs (Ref 3).
The carbon is in a state of supersaturation in the
as-formed martensite. When the martensite is
tempered, the carbon diffuses from these inter-
Fig. 3 Change in size of the octahedral interstitial site with
the change from face-centered cubic (fcc) to bodycentered cubic (bcc). Source: Ref 2
stitial sites and forms various carbides, leaving
the parent martensite less strained, softer, and
tougher. Figure 4 shows that the large strain energy in martensite varies with the carbon content, and Fig. 5 (Ref 4) shows how hardness
varies with carbon content.
Nitrogen behaves similarly to carbon in both
austenite and martensite, but its solubility is
lower, and it is less significant as an alloying element accordingly. Hydrogen and boron, as interstitials, also raise hardness.
Phase Structure
Figure 6(a) to (h) shows a series of photomicrographs of various martensitic alloys (Ref 5).
A stainless martensitic alloy should have the
following characteristics:
• It must have at least 10.5% chromium to
qualify as stainless and even more for better
corrosion resistance.
• It should be fully austenitic at some temperature.
• The temperature at which austenite forms on
heating should be sufficiently high to permit
tempering above the temper embrittlement
range.
These criteria are somewhat challenging. Figure 7(a) shows that at low-carbon (0.05%) levels austenite is stable up to about 12%
chromium, above which some δ-ferrite tends to
be stable at all temperatures below the melting
point. Increasing carbon slightly expands the
chromium level at which full austenitization can
occur (Fig. 7b) (Ref 3).
Fig. 4 Strain energy of martensite dependence on carbon
content. Source: Ref 2
128 / Stainless Steels for Design Engineers
Fig. 5 Variation in martensite hardness with carbon content
The interplay between chromium and carbon
is further explained in Fig. 8(a) and (b), in
which it becomes clear that for higherchromium alloys the range over which full
austenitization can occur is further restricted.
The variety of martensitic stainless steels
would be very limited if only chromium and
carbon were available as alloying elements, but
fortunately nickel again can make an important
contribution. Nickel greatly expands the
Chapter 9: Martensitic Stainless Steels / 129
Fig. 6 (a) Annealed 410 showing carbides within an equiaxed ferrite matrix. (b) 410 quenched and tem-
pered. (c) 416 quenched and tempered: white ferrite and gray sulfides in a martensite matrix. (d) 420
quenched and tempered showing fine carbides in a martensite matrix. (e) 420 quenched and tempered showing
surface decarburization. (f) 440A annealed displaying primary and smaller secondary carbides in a ferrite matrix. (g) 440B quenched and tempered displaying both primary and secondary carbides. (h) 440C quenched and
tempered displaying significant primary carbides plus finer secondary carbides in a martensite matrix
130 / Stainless Steels for Design Engineers
Fig. 7 Iron-chromium phase diagrams at two low-carbon levels
1800
1800
Liquid
1400
α+γ
Liquid
1600
Temperature, °C
Temperature, °C
1600
γ
1200
1000
L+α
1400
σ
γ
α+γ
1200
γ + carbide
1000
γ + carbide
800
800
α + carbide
α + carbide
600
600
0
0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0
Mass, %C
0
0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0
Mass, %C
Fig. 8 (a) Iron-chromium phase diagram at 12% chromium; (b) iron-chromium diagram at 17% chromium
chromium levels and temperatures at which
austenite is stable as is shown in Fig. 9 (Ref 3).
Table 2 quantifies the influences of the various possible alloying elements on the key properties of martensitic stainless steels.
It can be seen that the elements that promote
austenite, with the exception of cobalt, all depress the Ms temperature. This puts a limit on the
amount of total alloy that can be used and in the
end puts an upper limit on the ability of martensitic stainless steels to achieve high corrosion resistance. This is because as the main corrosion
fighters, chromium and molybdenum, which are
ferritizers, are increased, so must austenitizers
such as nickel. The coordinated increase in these
elements lowers the Ms to such a degree that the
Chapter 9: Martensitic Stainless Steels / 131
alloys become stably austenitic before much
higher corrosion resistance is obtained.
The ability to temper without austenite reversion is an important trait. Obviously, if transforming martensite to austenite during tempering caused subsequent untempered martensite
or other undesirable phases, this would limit
one’s ability to temper at a high enough temperature to achieve desired toughness. This limits
the use of nickel while encouraging the use of
elements like molybdenum.
Copper has become an important alloying element in martensitic stainless steels because it
greatly improves corrosion resistance in certain
environments without diminishing an alloy’s
ability to be tempered.
Thermal Processing
The main concerns with processing martensitic stainless steels are austenitizing, quenching, tempering/stress relieving, and annealing.
Fig. 9 The expansion of the range of austenite stability with
nickel content
Table 2
Austenitizing is complicated in martensitic
stainless steels because many grades contain
carbon at levels intended to produce carbides
for wear resistance purposes. Since carbon solubility varies strongly with temperature at
austenitizing temperatures (Fig. 7a and b), control of temperature is vital to have the correct
balance of carbon in solution versus carbon as
carbide since carbon in solution has such a
strong influence on ferrite content, Ms, and mechanical properties.
Austenitizing temperature also determines
austenite grain size. This affects Ms, but more
importantly it influences subsequent toughness.
Phosphorus precipitates at prior austenite grain
boundaries during tempering with a maximum
effect at 475 °C. This is the infamous temper
embrittlement. Figure 10 (Ref 3) shows the significant toughness change that occurs as increasing austenitizing temperature increases
austenite grain size and permits greater phosphorus concentrations at grain boundaries.
Refining phosphorus from any chromiumcontaining steel is quite challenging thermodynamically, so achieving low phosphorus levels
depends mainly on restrictions on raw materials
for melting. Because this is difficult or costly,
grain size control is the main tool for controlling temper embrittlement.
The higher-carbon grades, those above 0.20%
carbon, should be heated gradually through
stage heating to avoid cracking due to thermal
stresses. Soaking at 800 °C until uniform temperature is achieved minimizes this risk.
Another concern during austenitizing is superficial carbon loss, an example of which is
shown in Fig. 6(e). Heating in air to 1050 °C
can cause surface carbon to decrease by approximately 0.10% per hour, resulting in much
lower surface hardness. This loss increases with
base carbon level and austenitizing temperature.
Carbon or nitrogen pickup could also occur if
the atmosphere was rich in these elements. The
carbon potential of the furnace atmosphere must
be controlled to avoid potentially serious problems. If hydrogen atmospheres are used the
Influence of alloying elements on ferrite, Ms, and austenite start
Element
Lowering of %
ferrite per % element
Lowering of MS
per % element
Change of AC
per % element
N
C
Ni
Co
Cu
Mn
Si
Mo
Cr
V
Al
–220
–210
–20
–7
–7
–6
6
5
14
18
54
–475
–475
–17
0 to 10
–17
–30
–11
–21
–17
–46
...
0
0
0 to 280 0 to 250 –30 to –115
–25 to –66 25 to 73 25 to 60 0 to 35 50 to 290 30 to 750
132 / Stainless Steels for Design Engineers
+120
P: 0.047%
[ ]
P: 0.035%
+80
[ ]
Transition temperature, °C
P: 0.021%
+40
P: 0.003-0.004%
[ ]
0
[ ]
[ ]
[ ]
−40
−80
5
10
20
50
100
200
Austenite grain size, μm
Fig. 10 Influence of austenite grain size and phosphorus level on toughness
danger of embrittlement after quenching must
be recognized. Stress relief without delay would
be mandatory.
The high chromium content of these alloys
renders them very deep hardening. Air hardening is generally sufficient. Oil quenching which
is faster may be slowed by heating the oil.
Avoiding quench cracking and excessive
warpage is almost always a greater concern than
depth of hardening so air quenching is standard.
Because the quenching and the transformation it causes are inevitably accompanied by
residual stresses in a brittle material, stress relieving should be immediate to avoid cracking.
Higher-carbon grades should not even be allowed below room temperature before stress
relief. Pickling should never be done on asquenched material because this could easily result in hydrogen uptake and delayed cracking
by hydrogen embrittlement.
Heating as-quenched material to between 150
and 400 °C produces stress relieving. Besides
the normal flow on a microscopic scale, which
we understand as stress relieving, there is a
slight growth in the number of fine cementite
particles and a corresponding decrease in the
amount of carbon in solid solution. This results
in a slight decrease in hardness. At 400 °C, a
further precipitation of M2X and M7C3 as well
as the transformation of M3C into M7C3 can result in a secondary hardening, a true precipitation-hardening effect. In the presence of strong
carbide-forming alloying elements such as
molybdenum, vanadium, and tungsten, the M2X
carbide can become the more stable species and
be responsible for the secondary hardening. At
500 °C, coarser M23C6 and M7C3 begin to grow
at grain boundaries. This is accompanied by a
pronounced softening. The hardening reduction
with stress relief and tempering for a 12% Cr
alloy is shown in Fig. 11 (Ref 6).
Separately at the 475 °C range, the previously
mentioned phosphorus segregation to prior
austenitic grain boundaries occurs. This effect
begins to disappear above 550 °C. Thus true
tempering is conducted above this temperature.
The microstructural changes at these temperatures are the above-mentioned loss of carbon
from solid solution, carbide precipitation and
coarsening, and, of course, stress relief. The result is a pronounced softening and toughening.
If the material contains retained austenite, it
may decompose to ferrite and carbide with a
negative effect on toughness.
The molybdenum, vanadium, and tungstenalloyed grades will resist softening during tempering because of the strength of the secondary
hardening they undergo due to precipitation
Chapter 9: Martensitic Stainless Steels / 133
500
Initial hardness
300°C
450
300°C
350°C
350°C
400°C
400
400°C
450°C
450°C
500°C
550°C
Hardness
350
600°C
650°C
700°C
300
750°C
500°C
550°C
250
600°C
650°C
200
700°C
750°C
150
11
12
13
14
15
16
17
18
19
20
21
22
23
T (20 + LOG t) × 10−3
Fig. 11 Influence of tempering on hardness
hardening of carbides and nitrides. Nickel
seems to amplify this action by its influence on
diminishing the solubility of carbon in the matrix. Thus, the tempering of the higher-alloy
martensitic stainless steels can truly be considered a precipitation-hardening reaction.
The higher-carbon, higher-chromium grades
are typically only stress relieved because the removal of chromium from solution by carbide
formation at higher temperatures causes an unacceptable loss of corrosion resistance.
Applications
High-Temperature Use. The basic 12 % Cr
martensitic alloy has been the basis of alloying
improvements that were done to produce better
high-temperature performance, especially for
turbines.
The addition of vanadium and niobium, both
of which form much more stable carbides than
chromium, results in alloys that have vastly improved creep resistance in the 550 °C range, as
shown in Fig. 12 (Ref 3).
Tool and Cutlery Alloys. A high-profile use
of martensitic stainless steels is in cutlery. Hunting knives, sport knives, and chefs’ cutting tools
are highly valued items and contain some of the
most sophisticated martensitic stainless alloys.
While 420 is the common alloy and is quite
serviceable, much more wear- and corrosion-resistant alloys exist. At one time, 440C was the
maximum step up from 420; however, further
alloying with molybdenum for corrosion resistance and vanadium for hardness of the carbide
phase has led to improvements. The wear resistance of a blade is largely determined by the
hardness and amount of carbides while the
toughness is governed by the matrix properties.
These alloys are used at very high hardness
levels, so cleanliness is very important to toughness, which measures the ability to withstand
chipping in use. Electroslag remelting (ESR) or
vacuum induction melting-vacuum arc remelting (VIM-VAR) provides the cleanliness required, while powder metallurgy is optimal for
obtaining very fine carbide size and uniformity.
The nominal analyses of some prominent
grades are shown in Table 3.
The martensitic alloys have a tendency toward centerline segregation during solidification as well as toward the formation of primary
carbides. This has produced limitations in the
amount of highly wear-resistant constituents
such as vanadium carbide (hardness Rc 75),
which can be introduced into the matrix in conventional production. Powder metal techniques
134 / Stainless Steels for Design Engineers
4
50
2
45
Applied stress, kg/mm2
40
35
3
30
1
25
20
1 (0.2C-10.5Cr)
2 (0.2C-10.5Cr-0.1Nb)
3 (0.2C-10.5Cr-0.1V)
4 (0.2C-10.5Cr-0.1V-0.1Nb)
16
3
10
30
100
300
1000
3000
10,000
Rupture life, h
Fig. 12 Influence of vanadium and niobium on high-temperature properties
Table 3
Alloy
420
4116
440A
440C
BG-42
ATS-34
14-4
CrMo
154 CM
CPM
S30V
CPM
S60V
CPM
S90V
Tool and cutlery martensitic stainless steels alloy compositions
UNS
Form
C
Mn
Cr
Mo
Ni
Other
Other
...
12.0-14.0
14.5
...
0.65
...
...
...
...
...
0.15 V
...
0.4
0.5
0.03
0.03
...
...
...
1
1
0.3
0.35
0.3
16.0-18.0
16.0-18.0
14.5
14
14
0.75
0.75
4
4
4
...
...
...
...
...
...
...
1.2 V
...
...
...
...
...
...
...
1.05
1.45
0.45
...
...
...
0.3
...
14
14
4
2
...
...
...
4.0 V
...
...
PM
2.15
0.4
...
...
17
0.4
...
5.5 V
...
PM
2.2
...
...
...
13
1
...
9.0 V
...
S42000
Wrought
DIN 1.4116 Wrought
Nominal
S44002
Wrought
S44004
Wrought
Nominal
Wrought
Nominal
Wrought
Nominal
Wrought
0.15 min
0.5
1
0.60-0.75
0.95-1.20
1.15
1.05
1.05
1
1
Nominal
Nominal
Wrought
PM
Nominal
Nominal
...
S
Si
0.03
...
1
are not subject to the same limitations as continuous casters and have alloys the production of
alloys with high volume content of VC. One
such alloy is Crucible CPM 90V with 14% Cr,
9% V, 1% Mo, and 2.3% C. This alloy has equal
or better toughness and corrosion resistance as
440C but has ten times the wear resistance at
the same macrohardness.
Oil Country Tubular Good and Line Pipe.
The need for corrosion resistance in oil production tubulars has grown as the quality of petroleum deposits has become less optimal. Use of
stainless can eliminate for corrosion inhibitors
in H2S and CO2 environments. This has led to
the use of low-carbon martensitic stainless
steels. Low carbon and nitrogen levels give
good toughness without tempering and minimize the loss of chromium to carbides, maintaining it in solution for corrosion resistance.
The addition of nickel and molybdenum yields
full austenite and martensite transformation and
improves corrosion resistance. Table 4 lists several such alloys by JFE: the first two can be
made to meet L80 specifications and produced
as seamless. JFE reports production of over
100,000 tons per year of this product (Ref 7).
The third alloy is a near match for the precipitation-hardening stainless Custom 450 (UNS
S45000) (Ref 8). Like other precipitation-hardenable steels, it shows excellent resistance to
SCC at high strength levels. Figures 13 and 14
show the improvements in corrosion resistance
Chapter 9: Martensitic Stainless Steels / 135
Table 4
Alloy
JFE Steel/Nippon Steel oil country tubular goods and line pipe alloys
UNS
Form
C
Mn
S
Si
Cr
Mo
Ni
Other
Other
HP13Cr-! JFE
Nominal
HP13Cr-2 JFE
Nominal
NT-CRS Nippon
Nominal
NT-CRSS Nippon
Nominal
KL-12Cr JFE
Nominal
KL-HP
JFE
12Cr
Nominal
Wrought
0.025
0.45
...
...
13
1
4
...
...
Wrought
0.025
0.45
...
...
13
2
5
...
...
Wrought
0.03
1.45
...
...
12.7
1.4
4.5
1.5 Cu
0.040 N
Wrought
0.02
2
...
...
12.3
2
5.8
1.5 Cu
0.015 N
Wrought
0.01
...
...
...
11
...
2.4
0.5 Cu
0.010 N
Wrought
0.01
...
...
...
12
2
5.5
...
0.010 N
Fig. 13 Corrosion rates of stainless versus carbon steel
ronments to about Rc 22 to avoid SCC by the hydrogen embrittlement mechanism. The stainless
can resist this failure mode at higher strengths. As
specifying bodies such as the American Petroleum
Institute (API) approve the use of stainless tubulars at higher strength levels than carbon steel
tubulars are safely capable of handling, then the
strength improvement, coupled with the orders of
magnitude improvement in corrosion resistance,
will cause a great increase in their use.
Lower carbon levels permit the use of field
welds without tempering so that similar alloys
can be used for line pipe. These are corrosion
resistant and yet meet X70 and X80 class specifications. These alloys are the last two in Table
2. These uses of martensitic stainless steels for
oil production represent possibly the greatest
growth area for any kind of stainless steel in the
first decade of the 21st century.
REFERENCES
Fig. 14 Corrosion rates for stainless oil country tubular goods
(OCTG) alloys under severe operating conditions
over carbon steel L80 oil country tubular goods
(OCTG) under test conditions representative of
difficult real-use environments (Ref 7).
The improvements in martensitic steels for
these applications are hardly more than a thorough revisiting of the developments of the
1950s and 1960s. This does not diminish their
importance. Carbon steels are limited in sour envi-
1. D.R. Koistinen, R.E. Marburger, “A General
Equation Prescribing the Extent of the
Austenite/Martensite Transformation in
Pure Iron,” Acta Met, Vol 7, 1959, p 59
2. http://www.msm.cam.ac.uk/phase-trans/
2002/martensite.html
3. Bletton, Aciers Inoxidables, Les Editions de
Physique les Ulis, Paris, 1993, p 481
4. ASM Handbook Desk Edition, 1985, p 28–9
5. http://products.asminternational.org/mgo/
6. K.J. Irvine et al., JISI, Vol 195, ISIJ International 1960, p 386–405
7. S. Deshimaru et al., “Steels for Production,
Transportation and Storage of Energy, JFE
Technical Report (No. 2), March 2004
p 55–67
8. M. Kimura et al., “High CR Stainless
OCTG with High Strength and Superior
Corrosion Resistance,” JFE Technical Report (No. 7), Jan 2006, p 7–13
Stainless Steels for Design Engineers
Michael F. McGuire, p 137-146
DOI: 10.1361/ssde2008p137
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
CHAPTER 10
Precipitation-Hardening
Stainless Steels
Summary
Introduction
THE PRECIPITATION-HARDENABLE (PH)
grades are a highly specialized family of stainless steels whose existence derives from the
need for very high-strength materials with good
corrosion resistance. The workhorse alloys are
the martensitic PH grades, which are used in
many forms. Primarily used as forgings, bar,
and other hot-worked forms, they can also be
obtained in cold-rolled sheet and strip, although
not with the flatness expected from non-PH
stainless. The semiaustenitic alloys are more
amenable to production as sheet, strip, and wire
and are designed for applications that require
extensive forming before hardening. The fully
austenitic PH alloys fill a small niche where
high mechanical properties are required at temperatures above or below which the other PH
grades are found lacking, when a nonmagnetic
material is required, or when the higher thermal
expansive coefficient of an austenitic material is
desired. In no case is corrosion resistance better
than that of normal 304 found in PH stainless
steels. If enhanced strength and very high corrosion resistance are required, then the designer
should look to duplex stainless steels for the optimal material. If cost is a greater concern than
corrosion resistance or toughness, then martensitic stainless steels should be considered for
applications where strength and hardness over
that of annealed ferritic and austenitic stainless
is required.
The increased use of titanium alloys and advanced composite materials may occur at the
expense of the stainless PH alloys and at the
same time may create some new niche applications for them.
The PH stainless steels exploit the low austenite stability possible in the chromium/nickel
stainless steels by making the alloys so lean in
composition that they can be made to transform nearly entirely to martensite by thermal
or mechanical treatment. This martensite can
then be further hardened by the coherent precipitation of intermetallic compounds, elemental copper, nitrides, or even phosphides.
This precipitation hardening can also be made
to occur in a fully austenitic matrix, and this
also provides a commercial PH alloy. But, the
martensitic PH grades are by far the more
common. The border between the more highly
alloyed martensitic stainless steels, which undergo secondary hardening during tempering,
and the PH alloys is indeed vague. Some authors have astutely treated them as a single
group. Here, we treat them separately because
they are traditionally considered as separate
alloys.
The advantage of the PH alloys over the
strictly martensitic stainless steels is that they
attain great strength with higher toughness and
corrosion resistance than can be obtained
through the hardening of martensite through
carbon. In addition, they can be fabricated in a
relatively soft state and then hardened with very
little dimensional change.
The PH grades were developed at the beginning of World War II, with Stainless W (UNS
S17600) by U.S. Steel generally acknowledged
as the first. The later-developed grades are distinguished from the first by their more uniform,
and therefore tougher, microstructure through
the elimination of residual δ-ferrite and retained
138 / Stainless Steels for Design Engineers
austenite and by more astute alloy design and
chemistry control.
The mechanism of precipitation hardening is
parallel to that used to strengthen aluminum alloys in which the precipitation of a coherent second phase from a supersaturated solid solution is
produced by an aging heat treatment. The coherent precipitate strains the lattice and impedes the
motion of dislocations, producing strengthening.
Overaging causes the precipitates to lose coherency, and softening follows. The precipitate
that causes the hardening is normally nickel
(aluminum/titanium) (Ref 1). Figure 1 shows the
compounds that can form from the precipitation
of supersaturated aluminum and titanium in an
iron alloy matrix.
It is also possible to produce a hardening reaction by the precipitation of elemental copper
(Ref 2). In nitrogen-bearing alloys, a hardening
may be produced by the precipitation of Cr2N
(Ref 3). The precipitation begins with the diffusion of the precipitating species to sites on the
existing matrix. These enriched zones are called
Guinier-Preston (GP) zones. Close dimensional
matchup between the precipitating species and
the parent matrix is required. The differential
should be on the order of a percent. This allows
not only coherency but also strain. The coherent
precipitate is a effective barrier to dislocation
movement. As time and temperature of precipitation increase, the zones can grow to sizes that
cannot accommodate the small size differential;
coherency is lost, and with it the hardening effect diminishes. The precipitation has the dual
function of stress relieving the martensite while
further hardening the matrix through the precipitation of the coherent precipitate. The mechanical properties of the final microstructure depend
on the initial strength of the matrix before aging,
the amount of precipitate, and the coherency of
the precipitate. The ideal microstructure for the
initial matrix is 100% martensite. To the extent
there is δ-ferrite or retained austenite, properties,
especially yield strength and toughness in the
transverse direction, are compromised. The
aging temperatures can also be high enough that
reversion of martensite to austenite occurs,
which also lowers subsequent tensile properties.
While the presence of persistent, large bands
of either δ-ferrite or γ-austenite is undesirable,
but both also have benefits. The presence of
some fine bands of δ-ferrite promotes easier and
more reproducible precipitation of chrome carbides at the δ/γ interface during the “austenite
conditioning” or “trigger anneal” heat treatment
step for semiaustenitic alloys (17-7, AM350,
etc.). Although bands of stable austenite are
undesirable, it is the presence of residual interlath
γ + Ni(AlTi) + Ni3AlTi
Ni(AlTi)
Ni(AlTi) + Ni2AlTi
Ni2AlTi
4
Ni2(AlTi) + Ni3AlTi
Aluminum, wt%
Ni2AlTi
γ + Ni2AlTi
γ + Ni(AlTi)
Ni3Ti
σ
3
Limit of austenite
ferrite region in
solution treated
conditions
2
γ + Ni2(AlTi) + Ni3AlTi
γ + Ni3(AlTi)
1
γ + Ni3Ti + Ni3(AlTi)
γ
Cellular
precipitation
γ + Ni3Ti
0
1
2
3
Titanium, wt%
Fig. 1 Possible aluminum/titanium precipitates
4
Chapter 10: Precipitation-Hardening Stainless Steels / 139
austenite that provides the work-hardening ability in many of these PH alloys. It is this work
hardening that gives the PH alloys, especially the
semiaustenitic ones, their unusual combination
of high strength plus ductility and toughness in
the fully hardened state.
The complexity of PH steels comes from the
processing involved in producing the martensitic
structure in which the precipitation will occur.
The most straightforward alloys are the martensitic, also called the martensitic PH alloys. These
steels are supplied in the fully martensitic condition with hardness in the low Rc 30s. This is confusingly called the annealed condition, or condition A, even though the matrix is untempered
martensite. After the material is fabricated, it is
subjected to an aging treatment designated by
the aging temperature in Fahrenheit (e.g.,
H-950). These aging temperatures range from
950 °F (510 °C) to 1150 °F (620 °C).
A second major group of PH grades is the
semiaustenitic. These grades in the normally
furnished condition A are fully austenitic. This
is accomplished by adding elements that lower
the martensite start temperature, such as more
chromium, molybdenum, and nickel. The
austenite is more formable than martensite, and
it has the possibility of superior corrosion resistance because of higher chromium content. This
is balanced by the need to use either cold work,
cryogenic treatment, or a destabilizing anneal to
cause the matrix to become martensitic before
its precipitation aging treatment.
Last, if the austenite is made very stable by
further alloying additions, a precipitation reaction can still be made to occur by the same type
of aging treatment without martensite ever forming. The precipitation takes place in austenite
and therefore results in lower room temperature
Table 1
Alloy
strength than that of which the martensitic or
semiaustenitic alloys are capable. The austenitic
PH strength is better above 750 °F.
Martensitic Precipitation-Hardenable
Stainless Steels
The martensitic PH alloys are, as stated, fully
martensitic at room temperature. Their martensite is a relatively soft, low-carbon (less than
0.05%) martensite as opposed to the higher carbon found in the martensitic stainless steels. The
early alloys of this type, 17–7 PH and 17-4 PH,
contained up to 10% δ-ferrite stringers, which
caused poor through-thickness toughness. This
would be expected from the Schaeffler-Delong
diagram, but this is asking too much of the Schaeffler-Delong diagram, which was developed for
welds, to predict the phase composition of alloys
that have been homogenized by hot working.
The inaccuracy of the diagram for more complex systems was overcome, and alloys were designed that had minimal δ-ferrite and still transformed entirely to martensite, if not at room
temperature, at least at a reasonably attainable
subzero temperature. This was done first by trial
and error and more recently by using thermodynamic computer models, such as ThermoCalc, to
predict equilibrium phase composition. This development was very significant for making the
alloy family useful as a high-strength/hightoughness material for demanding applications
requiring high mechanical properties and corrosion resistance. The most advanced PH alloys
are martensitic PH grades by Cartech, Custom
465 and 475.
Table 1 shows the more significant of these alloys compared on a strength basis; Table 2 shows
Mechanical properties of martensitic precipitation-hardenable alloys
UNS
Condition
Yield, MPa
Tensile, MPa
Elongation, %
HRC
Stainless W
17-4 PH
15-5 PH
S17600
S17400
S15500
13-8 PH
S13800
1240
1210
1210
930
1450
1340
1310
1300
1025
1550
14
14
15
17.50
12
42
41
41
34
47
Custom 450
S45000
Custom 455
S45500
Custom 465
S46500
Custom 475
...
Ferrium S53
...
H-950 ( 510)
H-925 (495)
H-925 (495)
H1100 (595)
H-950 (510)
H1050 (565)
H-900 (480)
H1100 (595)
H-950 (510)
H1050 (565)
H-950 (510)
H1000 (535)
H-975 (525)
H 1100 (595)
...
1270
460
1515
1205
1650
1500
1855
1315
1565
1350
970
1585
1310
1765
1600
2005
1572
1985
14
23
10
14
11
13
5
13
14–16
42
...
48
40
49
48
54
48
54
Toughness,
CVN ft-lb
...
40
20
70
25
70
60
180
8
25
13
28
...
18
140 / Stainless Steels for Design Engineers
Table 2
Composition of martensitic precipitation-hardenable alloys
Alloy
Stainless W
17-4 PH
15-5 PH
13-8 PH
Custom 450
Custom 455
Custom 465
Custom 465 (275)
Custom 475
Ferrium S53
Designation
C
Mn
Si
Cr
Ni
Mo
Al
Cu
Ti
Other
S17600
S17400
S15500
S13800
S45000
S45500
S46500
...
...
...
0.1
0
0
0
0
0
0
0
0
0.2
0.5
0.6
0.6
0.1
0.3
0.3
0.2
0.2
0.4
0.1
0.5
0.6
0.6
0.1
0.3
0.3
0.2
0.2
0.4
0.1
17
16
15
13
15
12
12
12
11
10
6.3
4.3
4.3
8.5
6
8.5
11
11
8
5.5
...
...
...
0.2
...
...
1.1
...
...
...
...
1.2
0
...
3.2
3.2
...
1.5
2.5
...
...
...
...
...
...
1
...
...
...
...
1
2
2
...
...
P 0.3
...
...
...
0.3 Nb
0.3 Nb
...
0.2 Nb
8.0 Co
1W
0.3 V
14 Co
2
0.8
...
1
1
5
2
Fig. 2 Typical microstructures of precipitation-hardenable (PH) stainless steels: (a) 15-5PH as-quenched martensite; (b) 13-8 PH so-
lution treated and aged displaying fine martensite; (c) 17-7 PH displaying ferrite stringers in a martensite matrix; (d) 17-7 PH
showing residual ferrite stringers and inclusions
their nominal compositions. Figure 2 shows a series of photomicrographs of PH alloys.
The main advancement metallurgically in
these alloys from the top, and earliest, in Stain-
less W to the latest in Custom 475 besides the
elimination of δ-ferrite is in the volume fraction
of the precipitating phase and the elimination of
retained austenite. Alloy designers found that to
Chapter 10: Precipitation-Hardening Stainless Steels / 141
reduce δ-ferrite they also tend to stabilize
austenite, which also happens to reduce the
temperature at which martensite forms, MS.
Last, the higher levels of molybdenum reduce
the tendency to form secondary austenite during
aging. To minimize δ-ferrite requires reducing
chromium or molybdenum, which also reduces
corrosion resistance. As a result, as alloy
strength increases, corrosion resistance is compromised. The alloys with the greatest strength
potential, Custom 465 and 475, barely qualify
as stainless with around 11% chromium. But,
from a utility point of view, these alloys are designed to have maximum mechanical properties
with adequate corrosion resistance, so this is
viewed as an acceptable compromise.
The newest alloy is Ferrium S53, one of the
recent alloys designed by computer-assisted
thermodynamic calculations. It was designed to
replace 300M, 4340, and AerMet 100 on an
equal mechanical properties basis but also provide the corrosion resistance necessary to be
used in aircraft components without cadmium
plating. Its composition superficially seems deficient in chromium to provide “stainlessness,”
but the cobalt level raises the thermodynamic
activity of chromium sufficiently that the equivalent of 12% chromium in a non-cobalt-containing alloy is achieved. The precipitation hardening mechanism is the precipitation of Mo2C. It
has been established that this hardening mechanism optimizes resistance to stress corrosion
cracking (SCC) for a given strength level.
The alloying characteristics of these grades
are:
• Low carbon, nitrogen, silicon, and manganese because these elements lower MS
without contributing to age hardening
• Low chromium to suppress δ-ferrite
• Sufficient nickel to suppress δ-ferrite and
provide for precipitates without excessively
depressing MS
Table 3
• Molybdenum to offset loss of corrosion resistance by minimization of chromium, to
increase the temperature at which austenite
forms, and to form another hardening precipitate in the presence of cobalt
• Cobalt to stabilize austenite while raising
MS
• Aluminum or titanium to form intermetallic
precipitates with nickel or copper to precipitate as elemental copper
It is possible to quantify these various influences on phases. This is summarized in Table 3
in terms of the influence of the element on different factors measured in degrees Centigrade
for a 12% chromium alloy.
Rapid quenching of these alloys is not required. They are air hardenable. But, the cooling of these alloys must be completed expeditiously through the final stages of martensite
formation with minimal delay. During delays
after the start of martensite transformation has
occurred, the remaining austenite tends to stabilize, and full transformation to martensite does
not occur. When this happens, the higher levels
of austenite reduce subsequent mechanical
properties after aging.
As with any alloy used at such high strength
levels, microstructural cleanliness is essential,
but air melting and argon oxygen decarburization (AOD) refining are quite adequate.
Corrosion Resistance. The martensitic PH
stainless steels obey the same rules as other
stainless steels with regard to corrosion resistance. The martensite carries no nitrogen in solution, so the resistance to pitting is given by:
(Eq 1)
PREN = %Cr + 3.3%Mo
In the martensitic PH alloys, no chromium is
rendered ineffective by the formation of Cr23C6
since carbon is either held low or stabilized by
titanium or niobium. Thus, the corrosion
Influence of alloying elements on key transformations
Element
N
C
Ni
Co
Cu
Mn
Si
Mo
Cr
V
Al
–220
–210
–20
–7
–7
–6
6
5
14
18
54
Lowering of MS –475
per % element
Change of AC
0 to 280
per % element
–475
–17
0 to 10
–17
–30
–11
–21
–17
–46
0
0
25 to 60
0 to 35
Lowering
of % ferrite
per % element
0 to 250 –30 to –115
–25 to –66 25 to 73
50 to 290 30 to 750
142 / Stainless Steels for Design Engineers
resistance will equal that of stabilized ferritic alloys of the same pitting resistance equivalent
number (PREN) for which voluminous data are
available.
A greater concern is the risk of SCC in these
alloys. While the mechanism of SCC in
austenitic alloys is still debatable, it has long
been clear that, for martensitic alloys, SCC is
simply a manifestation of hydrogen embrittlement in which the hydrogen is provided by local
corrosion. The existence of pitting is a sufficient, if not necessary, condition for SCC to
occur if the temperature is within the range of
susceptibility and the material is inherently susceptible. The material susceptibility is largely a
function of resistance to crack propagation in
any given alloy that is measured by fracture
toughness. The martensitic PH grades have excellent toughness and low rates of crack propagation, but none should be considered immune
to SCC since their hardness is never below the
Rc 22 level, which is considered to be the
threshold hardness for susceptibility to SCC in
body-centered cubic (bcc) ferrous alloys.
The suitability of high-strength alloys for use
in potential SCC-provoking environments containing H2S is regulated in many locales by National Association of Corrosion Engineers
(NACE) International Standard MR01-75. In it,
the use of S17400 is permitted if it is double
tempered at 620 °C and its hardness is 33 HRC
or less, while S45000 can be used if it has been
aged at 620 °C for 4 h and its hardness is 31
HRC or less. These permissible hardness levels
are significantly higher than allowed in non-PH
martensitic alloys, 22 HRC, which reflects the
fact that the martensitic matrix has the toughness of a lower-hardness martensite.
In marine environments, the PH alloys are
susceptible to SCC if used at a high strength
level. S17400 aged at 480 °C with a yield
strength of 1240 MPa is susceptible to SCC,
while higher aging temperatures (above 540 °C)
producing lower strengths renders the materials
immune at stresses near the yield strength, approximately 1170 MPa. The threshold strength
for ordinary martensitic stainless steels would
be 1030 MPa (Ref 4).
Semiaustenitic PrecipitationHardenable Stainless Steels
If a martensitic stainless steel were alloyed
more strongly with austenite-stabilizing ele-
ments, the austenite could be made stable at
room temperature. This would make the alloy
softer and more fabricable and, most importantly, permit them to be manufactured as coldrolled sheet and strip. If the austenite could then
be transformed to martensite by cryogenic treatment, cold work, or special heat treatment, then
it could be age hardened just like the martensitic
PH grades. This has been accomplished for a
group of alloys called the semiaustenitic PH
grades. The “semi” signifies that the austenite
in these alloys is metastable rather than stable at
ambient temperatures. Also, it should be noted
that these semiaustenitic alloys usually contain
some δ-ferrite in their predominantly austenitic
microstructure after annealing.
These alloys are complex metallurgically because of the technique used to make the austenite
stable at room temperature after a full solution
anneal. The austenite is rendered stable by fairly
high levels of carbon, a powerful austenite stabilizer, in solution. The amount of carbon that can
be held in solution is a function of annealing temperature. The 1050 °C anneal of the condition A
mill anneal puts all the carbon in solution, giving
the austenite the stability of a normal 301-type
alloy. This permits extensive forming. The key is
to apply a subsequent lower-temperature anneal
so that less carbon goes into solution. Some of it
will thus form M23C6. This is in a sense deliberately sensitizing the alloy, but the sensitization
takes place at such a high temperature that
chromium deficits around precipitated carbides
are minimized by diffusion. This causes a higher
Ms temperature because of the lower amount of
carbon, and chromium, in solution in the austenite. Depending on the temperature at which the
anneal is done, the Ms temperature can be controlled so that a transformation to martensite can
be raised to either room temperature or some attainable cryogenic temperature. Figure 3 shows a
chart of these heat treatment options.
The lower strength levels achieved by the
condition T route in Fig. 3 reflect the lower carbon content of the martensite, while the highest
strength of the condition C route reflects the
compound influence of cold work and martensite hardness with similar subsequent contribution from age hardening.
The main alloys of this group are listed in
Table 4. Examination of the chemistries in this
table shows that the first two alloys rely on the
precipitation of Ni3Al for the hardening, while
the last two have no apparent precipitating
components. Their hardening is a more subtle
Chapter 10: Precipitation-Hardening Stainless Steels / 143
Fig. 3 Processing routes for S15700 Source: Ref 5
Table 4
Alloy
17-7 PH
15-7 PH
AM-350
AM-355
Compositions of semiaustenitic precipitation-hardenable alloys
Designation
C
Mn
Si
Cr
Ni
Mo
Al
N
S17700
S15700
S35000
S35500
0.1
0.1
0.1
0.1
0.5
0.5
0.8
0.9
0.3
0.3
0.4
0.4
17
15
17
16
7.1
7.1
4.3
4.3
...
2.2
2.8
2.8
1
1
...
...
0
0
0.1
0.1
secondary hardening from the tempering of
martensite rather than the classic precipitation
hardening via precipitation of intermetallic
compounds. In 17-7 and 15-7, aluminum rather
than titanium is the precipitating agent because
titanium would preferentially deplete the alloy
of carbon and nitrogen, precluding the action of
the conditioning heat treatment, which relies on
manipulating the amount of carbon in solution.
In AM350 and AM355, it is the precipitation of
144 / Stainless Steels for Design Engineers
Table 5
Mechanical properties of semiaustenitic precipitation-hardenable alloys
Alloy
UNS
Condition
Yield,
MPa
Tensile,
MPa
Elongation, %
HRC
17–7 PH
...
15–7 PH
...
AM-350
...
AM-355
...
TH 1050 (565)
RH 950 (510)
CH 900 (480)
TH 1050 (565)
RH 950 (510)
CH 900 (510)
SCT 850 (450)
SCT1000 (540)
SCT 850 (450)
SCT 1000(540)
1100
1380
1585
1380
1550
1720
1210
1020
1250
1035
1310
1520
1655
1450
1650
1790
1420
1165
1510
1124
10
9
2
7
6
2
12
15
13
22
42
46
49
45
48
50
46
40
48
38
(Cr,Fe)2N within the martensite phase that is
responsible for age hardening. In addition,
molybdenum produces a secondary hardening
in carbon-bearing martensite. These two early
alloys did not possess the hardening potential
that alloys employing copper-, titanium-, or
aluminum-based precipitates enjoy.
The mechanical properties of these alloys
are not greatly different from the martensitic alloys, as can be seen in Table 5. Their separate
existence is due to the need for alloys that are
more fabricable at room temperature than are the
alloys that are martensitic at room temperature.
This benefit is offset by the necessity to condition anneal before age hardening. If corrosion
resistance is a high concern, then the alloy and
heat treatment that yields the greatest amount of
chromium in solution should be chosen. Thus
condition CH is better than RH, which is better
than TH, the order of ascending solution annealing temperature and ascending Ms.
The martensitic PH grades have somewhat
better strength because they have a more uniformly martensitic structure, and they employ a
tougher, lower-carbon martensite. The fact that
the semiaustenitic alloys are typically sheet
products generally makes their service toughness requirements less onerous, so that their retained δ-ferrite is not a crippling drawback because of its detrimental effect on through
thickness toughness.
Corrosion Resistance. The semiaustenitic
PH alloys tend to higher values of PREN than
the martensitic alloys inherently since they are
alloyed to have lower Ms temperatures. The
thermal processing of these alloys causes a significant portion of the chromium to be removed
from solution as chromium carbide. This lowers
the corrosion resistance from what would be expected based on the bulk composition. From an
engineering point of view, it is best to assume
that all carbon is present as chromium carbide,
and that the chromium content is diminished by
that amount before applying Eq 1. In addition,
carbon can remove some molybdenum in the
form of carbides, and nitrogen can remove
chromium as a nitride. The effective corrosion
resistance of these alloys thus is similar to ferritic 430. The designer is thus advised to consult
with producers about corrosion resistance depending on the thermal processing that will be
used, especially if double aging is performed,
which can cause some degree of chromium depletion at grain boundaries.
While the general and pitting corrosion resistance of the semiaustenitic PH alloys are
never quite as good as most austenitic stainless,
they have very good resistance to SCC compared to ordinary martensitic stainless steels.
The δ-ferrite and the generally well-tempered
martensitic matrix provide a crack-arresting
feature and good inherent toughness that resist
SCC at higher strength levels than in straight
martensitic stainless steels. AM-355 in the SCC
(850 °F) condition can withstand stresses of
75% of 0.2% offset yield strength in salt spray
without SCC failure.
Austenitic Precipitation-Hardenable
Stainless Steels
The austinitic PH class consists of just one
important commercial alloy, A-286. The importance of the alloy is that it is entirely stable
austenite in both the solution-annealed and the
age-hardened condition. This means it is very
formable and nonmagnetic. And, because the
precipitation takes place in an austenite matrix,
the precipitation takes place at a higher temperature, around 700 °C. This gives the alloy the
Chapter 10: Precipitation-Hardening Stainless Steels / 145
Table 6
Alloy
A-286
Discalloy
Austenitic precipitation-hardenable composition
UNS
C
Mn
Si
Cr
Ni
Mo
Al
V
Ti
S66286
S66220
0.05
0.04
1.5
1.6
0.5
0.5
15
14
25.5
26
1.3
3
0.15
...
0.3
...
2.15
1.7
Fig. 4 A-286 properties as a function of test temperature. Source: Ref 5
potential to be used to near 700 °C without fear
of overaging. Thus, austenitic PH stainless represents a way to strengthen the austenite matrix,
which has the following advantages:
• High ductility and therefore high formability
in the soft, unaged condition
• High toughness at all temperatures and
strength levels
• Excellent creep and stress rupture properties
• Excellent oxidation, corrosion, and SCC resistance
This alloy is rightly considered an ironbased superalloy and is the root of the group
that succeeds it in properties, the nickel- and
cobalt-based superalloys. These last alloys can
attain greater strength and creep resistance
than A-286, effectively ending further development of austenitic PH alloys. The lower cost
of the A-286 alloy, compared to nickel-base
PH alloys, makes it attractive for a variety of
aerospace and nonaerospace uses.
The hardening mechanism of A-286 is the
precipitation of Ni3 (aluminum, titanium). Diffusion, even at the higher temperatures, is
slower in austenite, so aging treatments are typically 16 h. Table 6 gives the typical composition of A-286, and Fig. 4 shows some typical
properties as a function of temperature for a
standard 980 °C solution treatment followed by
a 720 °C, 16 h aging.
The mechanical properties can be greatly enhanced by cold working prior to aging, as is
shown in Fig. 5.
The toughness of the austenitic matrix is
abundant and quite temperature insensitive.
Charpy V-notch values of over 60 J are typical
from –200 to 800 °C.
The corrosion resistance of A-286 is comparable to that of 304 and 316. It has slightly
146 / Stainless Steels for Design Engineers
Fig. 5 The influence of cold work on aging response in A-286. DPH, diamond pyramid hardness. Source: Ref 5
better resistance to SCC despite its higher
strength level.
REFERENCES
1. F.B. Pickering, Physical Metallurgical Developments of Stainless Steel, Stainless ’84,
Goteborg, Sept 3–4, 1984, p 2–28.
2. M. Murayama, Y. Katayama, and K. Hono,
Microstructural Evolution in a 17-4 PH
Stainless Steel After Aging at 400 °C, Met-
allurgical and Materials Transactions A,
Vol 30A, Feb 1999.
3. G. Aggen, Ph.D. thesis, Carnegie Mellon
University
4. E.E. Denhard, “Stress Corrosion Cracking
of High Strength Stainless Steels in Atmospheric Environments”, paper presented at
the Twenty-fourth Meeting of the AGARD
Structures and Materials Panel (Turin,
Italy), April 17–20, 1967.
5. Allegheny Technology Blue Sheets
Stainless Steels for Design Engineers
Michael F. McGuire, p 147-154
DOI: 10.1361/ssde2008p147
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
CHAPTER 11
Casting Alloys
Summary
WITH TYPICAL ALLOY SYSTEMS, casting is often the most convenient method by
which to produce components. This is true for
stainless steels—both for corrosion-resisting
and for heat-resisting applications. This chapter
discusses primarily the alloys used for stainless
steel castings and their metallurgy. Foundry
methods are discussed to the degree they are
specific to the stainless alloys.
Stainless Steel Casting Alloys
Essentially any wrought stainless alloy composition can be modified to be made as a cast alloy.
The systemic difference between cast alloys and
their wrought equivalents is that cast alloys generally contain between 1.0 and 2.5% silicon. As
with other ferrous alloys, this is done to increase
the fluidity of the melt to make it cast more effectively. Silicon has strong metallurgical effects,
both beneficial and detrimental, which should be
understood by the user of cast stainless steels.
These are explained. A second general observation is that the stabilized ferritic stainless steel alloys, which constitute almost half the tonnage of
all stainless steel used, are notably absent from
the cast alloys. This is because these alloys are
single phase at all temperatures in the solid state
and because they have large as-cast grain sizes
that can only be refined by heavy cold work followed by annealing. This makes them quite lacking in toughness as cast. Since heavy cold work
defeats the purpose of casting to achieve a nearnet shape, stabilized ferritic stainless steels are
seldom used as castings. Also, the standard stabilizing alloy, titanium, is too readily oxidized for
normal foundry practice to avoid the loss of this
essential element. Thus, the casting alloys listed
in Tables 1 and 2 (Ref 1) are recognizable as approximate counterparts of the co-listed wrought
alloys (AISI grade). This cross reference to
wrought equivalents is helpful when looking for
data about an alloy that may be more easily
found for wrought alloys than for cast.
The High Alloy Product Group of the Steel
Founder’s Society of America employs a naming system (ACI, the Alloy Casting Institute)
for cast alloys that is significant; these designations are currently assigned by ASTM as grades
and are added to ASTM specifications. The first
letter, “C” or “H,” indicates corrosion resisting.
The second letter indicates the relative amount
of nickel, from a minimum of 0 to 1% for “A”
up to 30% nickel for “N” alloys. The number
following the hyphen for “C” alloys designates
the maximum carbon in hundredths of a percent.
The suffix letters designate additional alloying
elements, such as Cu for copper, M for molybdenum, N for nickel or nitrogen, F for free machining, and C for columbium (niobium). The
heat-resisting, “H,” alloys have generally only a
second letter designating relative nickel level on
the same scale as “C” alloys but going past stainless steels all the way to nickel-based alloys. The
inclusion of a number after the first two letters
indicates the center of the carbon range expressed in hundredths of a percent by weight.
To learn more about the influence of alloying
elements, refer to the chapters on the individual
alloy families; see Section 3. Here, we briefly
summarize:
• Pitting and crevice corrosion resistance, as
well as general corrosion resistance, are enhanced by chromium, molybdenum, tungsten, and nitrogen and carbon in solution.
• Localized corrosion is caused by chromium
depletion, which occurs when precipitates
148 / Stainless Steels for Design Engineers
Table 1
Compositions of cast stainless corrosion resisting alloys
Composition(a), wt%—maximum or range
ACI designation
Nearest
AISI grade
Chromium alloys
CA-15
410
CA-15M
CA-40
420
CA-40F
420F
CB-30
431,442
CC-50
446
Chromium-nickel alloys
CA-6N
CA-6NM
S41500
CA-28MWV
422
UNS
%C
%Mn
%Si
%Cr
%Ni
J91150
J91151
J91153
J91154
J91803
J92613
0.15
0.15
0.40
0.2–0.4
0.30
0.30
1.00
1.00
1.00
1.00
1.00
1.00
1.50
0.65
1.50
1.50
1.50
1.50
11.5–14.0
11.5–14.0
11.5–14.0
11.5–14.0
18.0–22.0
26.0–30.0
1.0
1.0
1.0
1.0
2.0
4.0
J91650
J91540
J91422
0.06
0.06
0.20–0.28
0.50
1.00
0.50–1.00
1.00
1.00
1.00
10.5–12.5
11.5–14.0
11.0–12.5
6.0–8.0
3.5–4.5
0.5–1.0
CB-7Cu-1
17-4PH (AISI 630) J92180
0.07
0.70
1.00
15.5–17.7
3.6–4.6
CB-7Cu-2
15-5 PH (XM-12)
J92110
0.07
0.70
1.00
14.0–15.5
4.5–5.5
CD-3MN
2205 (S32205)
J92205
0.03
1.50
1.00
21.0–23.5
4.5–6.5
CD-3MCuN
255 (S32550)
J93373
0.03
1.20
1.10
24.0-26.7
5.6–6.7
CD-3MWCuN (S32760)
J93380
0.03
1.00
1.00
24.0–26.0
6.5–8.5
CD-4MCu
CD-4MCuN
J93370
J93372
0.04
0.04
1.00
1.00
1.00
1.00
24.5–26.5
24.5–26.5
4.75–6.0
4.7–6.0
CD-6MN
J93371
0.06
1.00
1.00
24.0–27.0
4.0–6.0
J93404
J93345
J93423
J92500
J92800
J92700
J92600
J92710
J92900
J92590
J92901
J92971
J92972
J92701
J92602
J93790
0.03
0.08
0.30
0.03
0.03
0.03
0.08
0.08
0.08
0.04–0.10
0.04–0.10
0.10
0.10
0.12
0.16
0.20
0.06
1.50
1.00
1.50
1.50
1.50
1.50
1.50
1.50
1.50
1.50
1.50
1.50
7.00–9.00
1.50
1.50
1.50
4.00–6.00
1.00
1.50
2.00
2.00
2.00
1.50
2.00
2.00
2.00
2.00
1.50
1.50
3.50–4.50
2.00
2.00
2.00
1.00
24.0–26.0
22.5–25.5
26.0–30.0
17.0–21.0
17.0–21.0
17.0–21.0
18.0–21.0
18.0–21.0
18.0–21.0
18.0–21.0
18.0–21.0
15.0–18.0
16.0–18.0
18.0–21.0
18.0–21.0
18.0–21.0
20.5–23.5
6.0–8.0
8.0–11.0
8.0–11.0
8.0–12.0
8.0–12.0
9.0–13.0
8.0–11.0
9.0–12.0
9.0–12.0
8.0–11.0
9.0–12.0
13.0–16.0
8.0–9.0
9.0–12.0
9.0–12.0
8.0–11.0
11.5–13.5
CE-3MN
CE-8MN
CE-30
CF-3
CF-3M
CF-3MN
CF-8
CF-8C
CF-8M
CF-10
CF-10M
CF-10MC
CF-10SMnN
CF-12M
CF-16F
CF-20
CG-6MMN
2507 (S32750)
CG-8M
CG-12
CH-8
CH-10
CH-20
CK-3MCuN
317
308
309S
309H
309
254SMO™
J93000
J93001
J93400
J93401
J93402
J94653
0.08
0.12
0.08
0.04–0.10
0.20
0.025
1.50
1.50
1.50
1.50
1.50
1.20
1.50
2.00
1.50
2.00
2.00
1.00
18.0–21.0
20.0–23.0
22.0–26.0
22.0–26.0
22.0–26.0
19.5–20.5
9.0–13.0
10.0–13.0
12.0–15.0
12.0–15.0
12.0–15.0
17.5–19.5
CK-20
CN-3M
CN-3MN
CN-7M
CN-7MS
CT-15C
310
904L
AL-6XN®
320
J94202
J94652
J94651
N08007
J94650
N08151
0.20
0.03
0.03
0.07
0.07
0.05–0.15
2.00
2.00
2.00
1.50
1.50
0.15–1.50
2.00
1.00
1.00
1.50
3.50
0.50–1.50
23.0-27.0
20.0–22.0
20.0–22.0
19.0-22.0
18.0–20.0
19.0–21.0
19.0–22.0
23.0–27.0
23.0–27.0
27.5–30.0
22.0–25.0
31.0–34.0
312
304L
316L
316LN
304
347
316
304H
316H
316H
NITRONIC™ 60
316
303
302
NITRONIC™ 50
(a) Balance Fe for all compositions. Source: Ref 1
%Mo
%Other
0.50
0.15–1.00
0.50
0.20–0.40 Ss
0.4–1.0
0.9–1.25
0.9–1.25 W,
0.2–0.3 V
2.5–3.2 Cu,
0.2–0.35 Nb,
0.05 N max
2.5–3.2 Cu,
0.2–0.35 Nb,
0.05 N max
2.5–3.5
1.0 max Cu,
0.10–0.30 N
2.9–3.8
1.4–1.9 Cu,
0.22–0.33 N
3.0–4.0
0.5–1.0 Cu,
0.5–1.0 W,
0.20–0.30 N
1.75–2.25 2.75–3.25 Cu
1.75–2.25
2.75–3.25
Cu, 0.10–0.25 N
1.75–2.25 1.75–2.5 Cu,
0.15–0.25 N
4.0–5.0
0.10–0.30 N
3.0–4.5
0.10–0.30 N
2.0–3.0
2.0–3.0
0.10–0.20 N
Nb
2.0–3.0
2.0–3.0
1.75–2.25
2.0–3.0
1.5 max
(10xC)–1.2 Nb
0.08–0.18 N
0.2–0.35 Se
1.5–3.0
0.1–0.3 Nb,
0.1–0.3 V,
0.2–0.40 N
6.0–7.0
0.5–1.0 Cu,
0.18–0.24 N
4.5–5.5
6.0–7.0
2.0–3.0
2.5–3.0
0.18–0.24 N
3.0–4.0 Cu
1.5–2.0 Cu
0.5–1.5 Nb
Chapter 11: Casting Alloys / 149
Table 2
Compositions of cast heat-resistant stainless and nickel base alloys
Composition(a), wt%—maximum or range
ACI designation
HA
HC
HD
HE
HF
HH
HI
HK
HK-30
HK-40
HL
HN
HP
HP-50WZ
HT
HT-30
HU
HW
HX
Nearest AISI grade
504
446
327
312
302B
309
310
330
UNS
%C
%Cr
%Ni
%Si max
J82090
J92605
J93005
J93403
J92603
J93505
J94003
J94224
J94203
J94204
N08604
J94213
N08705
0.20 max
0.50 max
0.50 max
0.20–0.50
0.20–0.40
0.20–0.50
0.20–0.50
0.20–0.60
0.25–0.35
0.35–0.45
0.20–0.60
0.20–0.50
0.35–0.75
0.45–0.55
0.35–0.75
0.25–0.35
0.35–0.75
0.35–0.75
0.35–0.75
8–10
26–30
26–30
26–30
19–23
24–28
26–30
24–38
23–27
23–27
28–32
19–23
24–28
24–28
13–17
13–17
17–21
10–14
15–19
4–max
4–7
8–11
9–12
11–14
14–18
18–22
19–22
19–22
18–22
23–27
33–37
33–37
33–37
33–37
37–41
58–62
64–68
1.00
2.00
2.00
2.00
2.00
2.00
2.00
2.00
1.75
1.75
2.00
2.00
2.00
2.50
2.50
2.50
2.50
2.50
2.50
N08605
N08603
N08005
N08006
N06050
(a) Balance Fe for all compositions. Manganese content: 0.35–0.65% for HA, 1% for HC, 1.5% for HD, 2% for the other alloys. Phosphorus and sulfur contents: 0.04
(max) for all but HP-50WZ. Molybdenum is intentionally added only to HA: 0.90–1.2%. Maximum molybdenum for other alloys is 0.5%. HH contains 0.2% N
(max). HP-50WZ also contains 4–6% W, 0.1–1.0% Zr, and 0.035% S (max) and P (max). Source: Ref 1
form in the solid state. These precipitates are
carbides, oxides, and sulfides as well as intermetallic phases richer in chromium, molybdenum, or nitrogen than the matrix.
• General corrosion resistance follows the
above guidelines but is also helped by copper and nickel, which do not assist in pitting
resistance.
• High-temperature oxidation resistance is enhanced by increasing chromium and silicon.
Wrought alloys employ aluminum and rare
earths to help oxidation resistance, but the
difficulty of keeping these elements from
being oxidized requires special techniques
such as vacuum induction melting and inert
refractories for molds.
• Iron-chromium (ferritic) alloys have better
thermal fatigue resistance but poorer creep
resistance
than
iron-chromium-nickel
(austenitic) alloys.
The alloy designation system largely ignores
the wrought alloy distinctions by microstructure
(i.e., ferritic, austenitic, duplex, PH [precipitation hardening], and martensitic). One reason is
that the most widely used wrought-stabilized
ferritics (e.g., 409, 439) do not exist as common
casting alloys, and nominally austenitic alloys
in the cast form contain enough ferrite to be significantly magnetic. Thus, the distinctions based
on phase are not as well defined for casting alloys. The high ferrite content in the nominally
austenitic casting alloys is to avoid or at least
minimize solidification hot cracking or to allow
weld repair of cracks that do form. It has been
shown that the existence of ferrite can increase
the resistance to stress corrosion cracking.
Metallurgy of “C” Alloys
The corrosion-resistant “C” series have
wrought counterparts from which they differ essentially only in silicon content. This silicon has
no significant influence on corrosion resistance
or mechanical or physical properties, so an understanding of these alloys by approximating
them to their wrought counterparts is justified.
The main difference between the cast and
wrought product forms of these alloys is the
grain structure. In wrought grades, the grain
structure can be manipulated by deformation
and heat treatment. The use of deformation is
not an option in cast alloys; consequently, the
opportunities for grain refinement in cast alloys
are limited. To counteract the problems of lower
corrosion resistance of cast grades, a homogenizing solution heat treatment is necessary to
counteract the chromium depletion that occurs
due to solidification segregation and precipitating
phases. Representative mechanical properties
for “C” alloys are listed in Table 3 (Ref 1).
Martensitic. CA alloys are martensitic. The
metallurgy is straightforward and equivalent to
their wrought counterparts. The mechanical
Heat treatment(a)
>955 oC (1750 oF), AC, T
980 oC (1800 oF), AC, T
980 oC (1800 oF), AC, T
1040 oC (1900 oF), OQ, A
790 oC (1450 oF), AC
1040 oC (1900 oF), AC
1120 oC (2050 oF), FC to
1040 oC (1900 oF), WQ
1120 oC (2050 oF), FC to
1040 oC (1900 oF), A
1095 oC (2000 oF), WQ
>1040 oC (1900 oF), WQ
>1040 oC (1900 oF), WQ
>1040 oC (1900 oF), WQ
>1040 oC (1900 oF), WQ
>1095 oC (2000 oF), WQ
>1040 oC (1900 oF), WQ
>1040 oC (1900 oF), WQ
>1065 oC (1950 oF), WQ
>1065 oC (1950 oF), WQ
>1095 oC (2000 oF), WQ
>1040 oC (1900 oF), WQ
>1095 oC (2000 oF), WQ
1150 oC (2100 oF), WQ
1150 oC (2100 oF), WQ
1120 oC (2050 oF), WQ
669
531
600
531
586
531
552
621
552
531
531
565
607
524
770
476
97
77
87
77
85
77
80
90
80
77
77
82
88
76
112
69
130
896
ksi
120
115
150
190
95
97
108
827
793
1034
1310
655
669
745
MPa
Tensile strength
434
248
290
255
310
248
262
310
290
262
276
303
345
262
365
214
634
689
689
862
1172
414
448
558
MPa
24
22
10
14
15
18
25
20
18
60
50
55
50
50
55
45
50
39
52
45
38
37
50
48
92
63
36
42
37
45
36
38
45
42
38
40
44
50
38
53
31
Elongation %
in 50 mm, 2 in.
100
100
125
170
60
65
81
ksi
Yield strength, 0.2% offset
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
60
55
30
54
...
...
...
Reduction
in area, %
Room temperature mechanical properties of corrosion resisting cast stainless alloys
(a) AC, air cool; FC, furnace cool; OQ, oil quench; WQ, water quench; T, temper; A, age. Source: Ref 1
CE-30
CF-3
CF-3A
CF-8
CF-8A
CF-20
CF-3M
CF-3MA
CF-8M
CF-8C
CF-16F
CG-8M
CH-20
CK-20
CN-3MN
CN-7M
CA-6NM
CA-15
CA-40
CB-7Cu
CB-30
CC-50
CD-4MCu
Alloy
Table 3
190
140
160
140
156
163
150
170
170
149
150
176
190
144
185
130
305
269
225
310
400
195
210
253
Hardness,
HB
J
9.5
149.2
135.6
100.3
94.9
81.4
162.7
135.6
94.9
40.7
101.7
108.5
40.7
67.8
190
94.9
35.3
94.9
27.1
2.7
33.9
2.7
...
74.6
7
110
100
74
70
60
120
100
70
30
75
80
30
50
140
70
26
70
20
2
25
2
...
55
ft-lb
Specimen
Keyhole notch
V-notch
V-notch
Keyhole notch
Keyhole notch
Keyhole notch
V-notch
V-notch
Keyhole notch
Keyhole notch
Keyhole notch
V-notch
Keyhole notch
Izod V-notch
V-notch
Keyhole notch
V-notch
V-notch
Keyhole notch
Keyhole notch
V-notch
Keyhole notch
...
V-notch
Charpy toughness
150 / Stainless Steels for Design Engineers
Chapter 11: Casting Alloys / 151
properties are governed by the thermal processing, and strength, hardness, and toughness
can be varied over a wide range. The CB 30
and CC 50 alloys are ferritic and, as such, have
negligible toughness but effectively deliver
corrosion resistance. The toughness of CB 30
can be improved by balancing the chromium
and silicon to a lower part of the range and the
carbon and nickel to the higher end to render
the microstructure partly martensitic.
Precipitation Hardening. The cast PH alloys
include CB-7Cu-1, which behaves in a similar
way to 17-4PH, which has an overlapping composition range. Note that most other major
wrought PH grades rely on titanium and aluminum to form coherent strengthening precipitates and so do not have cast counterparts. Copper, which can harden ferrite but not austenite,
is thus the only strengthener available. There is
one cast PH alloy that has no wrought counterpart. It is CD-4MCu; however, it is rarely used
in the precipitation-hardened condition and is
most commonly classified as a duplex stainless
steel in which the nitrogen level is closely controlled. This is a highly alloyed duplex grade
that contains copper to precipitation harden the
ferrite phase. Oil field CO2 corrosion is resisted
by alloys that resemble the martensitic PH
grades. These alloys are discussed in Chapter
22, “Petroleum Industry Applications” and can
be considered castable alloys.
Duplex. The cast equivalents of alloys 2205
and 2507, J 92205, and J 93380 have similar
properties and corrosion resistance. Modern
wrought duplex alloys rely on nitrogen to partition the alloy with uniform corrosion resistance
in each phase and to suppress intermetallic phase
formation. Cast alloys are effectively limited to
0.25% nitrogen before gas porosity becomes excessive. Porosity can be reduced by replacing
some nickel with manganese, which increases nitrogen solubility. Doing so would expand the
most promising area of stainless steel development, lean duplex alloys such as 2101 and 2003,
to the cast grades. Alloy 2101 with 4 to 6% manganese provides the corrosion resistance of CF8M or 316L with total nickel plus molybdenum
of only 2% versus the 12% required for the
austenitic alloy. The duplex alloys also have
greater strength and are nearly immune to stress
corrosion cracking. These alloys represent significant cost-savings potential for the foundry and
for its customers. CE-30 is duplex steel, which is
fairly simple metallurgically and uses only
chromium for corrosion resistance. However, its
high level of nickel negates much of the potential
cost savings duplex alloys offer. All castings are
solution annealed and quenched to eliminate embrittling intermetallic phases.
The cast duplex alloys may offer a better engineering approach than the equivalent austenitic
cast alloys because they have greater strength
and lower alloy cost for the same level of corrosion resistance. They do not have the same problems of hot cracking that make casting austenitic
steels difficult. The poor hot workability of
duplex steel is not an issue for castings. It is important for designers to understand that cast duplex steels are totally compatible galvanically
with wrought or cast austenitic alloys of the
same corrosion resistance. Mixing components
with different microstructures does not create a
galvanic differential when corrosion resistance
levels are similar. Reluctance to mix alloys for
galvanic reasons can be an expensive error when
their similar corrosion resistances makes them
compatible, even if they are quite different
microstructurally.
Austenitic-Ferritic. The typical CF alloys,
which make up about two-thirds of U.S. stainless steel castings, are nominally austenitic but
always contain ferrite. This is not detrimental
and improves resistance to stress corrosion
cracking and sensitization. Homogenization annealing can reduce the amount of ferrite and result in lower yield and tensile strength and
higher elongation and toughness. The composition balance is the main determinant of ferrite
level. Increasing the nickel, nitrogen, manganese,
or carbon content decreases ferrite. Increasing
chromium, silicon, or molybdenum content increases ferrite. Increasing the solidification rate
will increase the ratio of austenite to ferrite in
duplex or austenitic-ferritic alloys. The predominantly austenitic matrix has a very high toughness even at cryogenic temperatures. Ferrite, if
continuous, decreases toughness. Fortunately, it
is seldom present as a continuous phase. The
loss of toughness associated with high ferrite
content can be aggravated by heating the ferrite
above 475 oC (885 oF) for a sufficient time for
the ferrite to decompose to the brittle α and α'.
At higher temperatures, development of the σ
phase would have a similar embrittling effect.
These phases thus formed are quickly redissolved and removed by annealing. Note that
sometimes copper is added to austenitic alloys
to improve corrosion resistance in sulfuric acid
environments. It has no precipitation hardening
effect in austenite, as it does in ferrite. When
152 / Stainless Steels for Design Engineers
used as a precipitating hardening agent, copper
does not increase corrosion resistance.
Virtually any non-titanium-bearing, corrosion-resistant, austenitic, wrought alloy can
have a cast counterpart. Curiously, the 2xx lownickel alloys are not found in most cast alloys
lists. If a specific wrought alloy cannot be found
to have a published cast counterpart, the designer should not avoid requesting a producer
to supply a version that the foundry is confident
of making. The designer must thoroughly understand the design of the alloy desired so that
any alterations to its composition necessary to
allow castability will not compromise expected
performance.
Metallurgy of “H” Alloys
The heat-resisting “H” alloys are principally
austenitic. Alloying elements and impurities
diffuse more slowly through the face-centered
cubic (fcc) austenitic structure than the bcc
ferrite structure, making the austenite more resistant to diffusion-controlled creep. Austenite
Table 4
Mechanical properties of heat-resistant cast stainless alloys at room temperature
Tensile strength
Alloy
Standard grades
HA
HC
HD
HE
HF
HH, type 1
HH, type 2
HI
HK
HL
HN
HP
HPNb(d)
HPNbTi(e)
HT
HU
HW
HX
has higher thermal expansion and lower thermal
conductivity than ferrite, which aggravates thermal fatigue and oxide spalling. Nevertheless,
the better high-temperature strength of austenite
generally is the predominant consideration, and
most “H” alloys are austenitic. Tables 4 and 5
list properties of “H” alloys (Ref 1 to 3).
Ferritic HA, HC, HD. Of the ferritic alloys
HA, HC, and HD, HA with less than 10%
chromium is not quite stainless but is useful to
650 oC (1200 oF) for petroleum refinery applications. HC and HD are very high chromium
ferritics that have very low toughness and creep
resistance but are quite oxidation and sulfidation resistant. They can be cost-effective materials when high-temperature strength is not an
overriding concern.
Austenitic HE-HP. The predominant hightemperature grades are the austenitic HE
through HP, after which come the nickel alloys,
which are not generally classified as stainless
steels because they contain less than 50% iron.
The high material cost of the nickel base alloys
restricts their use to those specific environments
where maximum carburization or nitriding
Yield strength
Condition
MPa
ksi
MPa
ksi
N + T(a)
As-cast
Aged(b)
As-cast
As-cast
Aged(b)
As-cast
Aged(b)
As-cast
Aged(b)
As-cast
Aged(b)
As-cast
Aged(b)
As-cast
Aged(c)
As-cast
As-cast
As-cast
As-cast
As-cast
As-cast
Aged(c)
As-cast
Aged(f)
As-cast
Aged(g)
As-cast
Aged(f)
738
760
790
585
655
620
635
690
585
595
550
635
550
620
515
585
565
470
490
450
450
485
515
485
505
470
580
450
505
107
110
115
85
95
90
92
100
85
86
80
92
80
90
75
85
82
68
71
558
515
550
330
310
380
310
345
345
380
275
310
310
450
345
345
360
260
275
220
220
275
310
275
295
250
360
250
305
81
75
80
48
45
55
45
50
50
55
40
45
45
65
50
50
52
38
40
70
75
70
73
68
84
65
73
40
45
40
43
36
52
36
44
Elongation, %
21
19
18
16
20
10
38
25
25
11
15
8
12
6
17
10
19
13
11
8
8
10
5
9
5
4
4
9
9
Hardness, HB
220
223
...
90
200
270
165
190
185
200
180
200
180
200
170
190
192
160
170
...
...
180
200
170
190
185
205
176
185
(a) Normalized and tempered at 675 °C (1250 °F). (b) Aging treatment: 24 h at 760 °C (1400 °F), furnace cool. (c) Aging treatment: 24 h at 760 °C (1400 °F), air cool.
(d) ISO 13583-2 specification minima. (e) ISO 13583-2 specification limits for microalloyed grade. (f) Aging treatment: 48 h at 980 °C (1800 °F), air cool. (g) Aging
treatment: 48 h at 980 °C (1800 °F), furnace cool. Source: Ref 1
Chapter 11: Casting Alloys / 153
Table 5
High-temperature mechanical properties of “H” alloys
Alloy
Temp
HA
1400 oF
1800 oF
1400 oF
1800 oF
1400 oF
1800 oF
1400 oF
1800 oF
1400 oF
1800 oF
1400 oF
1800 oF
1400 oF
1800 oF
1400 oF
1800 oF
1400 oF
1800 oF
1400 oF
1800 oF
1400 oF
1800 oF
1400 oF
1800 oF
800 oC
800 oC
1400 oF
1800 oF
1400 oF
1800 oF
1400 oF
1800 oF
1400 oF
1800 oF
HC
HD
HE
HF
HH
TYPE 1
HH
TYPE 2
HI
HK
HL
HN
HP
HPNb(a)
HPNbTi(a)
HT
HU
HW
HX
Yield strength
ksi, MPa
8.7
2.1
Tensile
strength
ksi, MPa
10.5
2.5
36
15
35
17
6.3
18
7
33
9
35
11
50
18.7
Creep rate
0.0001%/h
psi, MPa
1% in 100,000
h psi, MPa
Stress to
rupture
in 1000 h
16
27
1.3
3.6
3.5
1.0
3.5
1.0
6.0
1.6 (est)
3.0
1.1
7.0
2.1
6.6
1.9
6.8
2.7
7.0
2.8 (est)
1.3
0.6
7.0
2.5
11.0
2.5
9.1
4.4
2.0
6.3
0.9
8.0
1.6
8.5
2.6
12.0
2.8
15
5.2
2.4
...
...
26
8
6.2
23
8
19.5
6.9
...
...
35
11
40
10
32
10
42
10.7
2.1
51 MPa
54 MPa
8.0
2.0
8.5
2.2
6.0
1.4
6.4
1.6
Stress to
rupture
in 10,000 h
0.9
1.7
2.1
...
...
12
2.7
55 MPa
64 MPa
2.9
7.8
2.6
1.8
1.7
2.2
(a) Data from ISO 13583-2. Source: Ref 2, 3
resistance is mandatory. For oxidation and sulfidation resistance, the iron base alloys are preferred. These cast stainless steels derive their
oxidation resistance from their chromium level.
The chromium near the surface acts as a reservoir to replenish the protective iron/chromium
oxide scale as explained in Chapter 6 in the section on oxidation. Silicon, another stable oxide
former, assists in forming this protective scale
and resistance to carburization. Other typical
alloying elements do not aid in oxidation resistance. If it were possible to cast these alloys with
aluminum or rare earth additions without them
being lost to oxidation before solidification,
there could be some impressive benefits. Such
alloys exist in wrought form, for example,
153MA and 253MA. The metallurgical basis of
the benefits from aluminum and rare earth
alloying also are discussed in the oxidation section of Chapter 6.
The major problem that all producers of stainless steels face is that of transferring molten
metal from the furnace to the mold cavity. This
problem is heightened when the foundry makes
complex shapes. Methods developed to protect
the molten stream from exposure to air to prevent
reoxidation have shown great promise and have
been demonstrated by the wrought alloy producers who tend to produce much simpler shapes
than the foundry. Protection of the molten stream
could result in castings with much better hightemperature performance that could be used instead of some use of higher nickel alloys.
High-temperature strength is modestly improved by higher levels of chromium and
nickel. Molybdenum improves high-temperature
154 / Stainless Steels for Design Engineers
strength, but its detrimental effect on oxidation
resistance and its promotion of intermetallic
precipitation limits its use. Carbon is very effective for promoting high-temperature strength
and suppression of intermetallic phase formation. All “H” alloys, therefore, employ much
higher carbon levels than the “C” alloys. This
does, however, directly imply that the corrosion
resistance of “H” alloys, should it be an issue, is
significantly degraded over otherwise similar
“C” alloys.
The HP grades have undergone significant
development over the last 30 years. This development has come about through the addition of
niobium to increase creep and rupture properties. The use of microalloying additions has delivered creep and rupture properties some 30%
higher than the HP grade without niobium microalloying. It is unfortunate that these HP
grades have not at this time found their way into
ASTM standards; however, work is under way
to remedy this omission. Currently, the most upto-date collection of these grades can be found
in ISO 13583-2.
higher-performance grades. It is also possible to
use AOD-refined master melt stock to achieve
the same benefits as AOD refining while using
induction melting.
Welding of cast stainless alloys is a common
practice and does not present problems when
using approved weld procedures and qualified
welders. Chapter 17 describes joining methods
in detail. The same precautions about sensitization apply to castings. Welding of non-niobiumstabilized “C” alloys with carbon levels above
0.03% will require postweld annealing to redissolve chromium carbides, which will otherwise
make the alloy susceptible to corrosive attack in
the chromium-depleted regions of the heataffected zone.
Iron and nickel base “H” alloys that are fully
austenitic can suffer from hot shortness due to
sulfide films that precipitate along grain boundaries even at low bulk sulfur levels. This makes
them susceptible to hot cracking of welds. Alloys with some ferrite are less susceptible to hot
cracking, so most “C” alloys are highly resistant
to this problem.
Foundry Practice
REFERENCES
While the scope of this book does not extend
to the production of castings, certain aspects are
important to the user of castings. For the last 50
years, virtually all stainless has been refined in
argon oxidation decarburization (AOD) vessels
or versions thereof. For “C” alloys, this refining
method should be considered a basic requirement for good quality where the carbon levels
are restricted to low levels (e.g., CF3M). “H”
alloys are less refined inherently and can be
simply arc or induction melted; however it may
be necessary to use refining techniques for the
1. Cast Stainless Steels, Metals Handbook,
desk ed., J.R. Davis, Ed., ASM International, 1988, p 386–390
2. International Organization for Standardization, www.iso.org, ISO 13583-2
3. Steel Founders Society of America, online
documents: http://www.sfsa.org/sfsa/pubs/
index.html
SELECTED REFERENCE
• http://www.sfsa.org
Stainless Steels for Design Engineers
Michael F. McGuire, p 155-160
DOI: 10.1361/ssde2008p155
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
CHAPTER 12
Melting, Casting, and Hot Processing
Summary
THE PRIMARY PRODUCTION PROCESSES of melting, casting, and hot processing
are invisible to the end user. The vast majority of
stainless steel is made by arc furnace melting followed by argon oxygen decarburization (AOD)
refining and continuous casting. It is not normal,
and it is seldom beneficial for the end user to
specify processing paths. The end user should,
however, be knowledgeable and require the producer to document the process and the producer’s
control of it.
Introduction
The manner in which stainless steel is made
at the producing mill can have a great impact on
its final properties. These production methods
have undergone a major evolution over the last
50 years and are mainly responsible for stainless steels becoming the practical, widespread
engineering materials they are today. Traditional
carbon and alloy steel-making methods are not
suitable for stainless steels. The fundamental
difference is that the basic decarburization step,
which is common to all steel making, is thermodynamically very difficult in stainless steel because the essential element, chromium, reacts
more strongly with the purifying agent, oxygen,
than does carbon. Thus, early stainless steel making, done in an arc furnace, was a lengthy process
that necessarily involved high chromium losses
to the slag as carbon was removed. This process
was not only very expensive, the carbon levels
that could be achieved were not much below
0.10%, making most of today’s stainless steels,
whose carbon levels range from 0.010% in stabilized ferritic alloys to about 0.07% in normal
austenitic alloys, impossible to produce. The
advent of AOD, continuous casting, ladle metallurgy, and powerful hot rolling mills has led to
stainless steels of much higher quality produced
at lower cost. Ironically, the low processing cost
of stainless steel has spurred demand and made
some of its ingredients, such as molybdenum
and nickel, which are relatively scarce and expensive commodities, even more costly, forcing
the cost of many alloys to spike even higher
than in earlier years.
Melting and Refining
The arc furnace is nearly universally used for
the first step in the production of stainless steel.
The arc furnace is quite flexible in the types of
charge materials it can accept. Since the charge
materials for stainless steel are typically carbon
steel and stainless steel scrap, this flexibility allows scrap of all types to be used. The necessary
chromium is added as ferrochromium, whose
cost is inversely related to its carbon content.
The carbon content of the heat of steel is roughly
1.5 to 2.5% when it is melted and ready to
charge into the separate refining vessel.
It is this carbon whose removal is the primary
focus of refining. In the 1960s, Union Carbide
engineers perfected a method, the previously
mentioned AOD process, of removing nearly all
the carbon from molten stainless steel without
significant loss of chromium. This process is
based on the following chemical reaction:
Cr3O4 (Solid) + yC = yCO (gas) + Cr
(Eq 1)
The equilibrium for this reaction is:
Ln (K ) =
−ΔG
4575T
(Eq 2)
156 / Stainless Steels for Design Engineers
where K is the equilibrium constant, and G is
the Gibbs free energy.
Working through the thermodynamics yields
the relationship that summarizes the important
relationship among carbon, chromium, and CO
(Ref 1):
Log =
%Cr
13,800
=−
+ 8.76 − 0.925 log p CO
T
%C
(Eq 3)
Thus, increasing the temperature works to increase the elimination of carbon as CO, which
evolves from the melt. This is similar in principle
to the basic oxygen furnace (BOF) process for
carbon steel in which oxygen is injected into
molten steel to remove carbon by oxidizing it.
The key to the AOD process, though, is the injection of oxygen and argon into the bath to keep
the partial pressure of CO (pCO) very low. This is
done at a temperature consistent with economic
refractory life. The injection is done through
tubes called tuyeres in the bottom of the barrelshaped vessel. The injection and the reaction
cause extremely thorough mixing, which would
never happen in the flat, stagnant, arc furnace
bath. This mixing not only allows the CO-producing reaction to reach equilibrium, but also
the mixing of the slag and metal also permits
desulfurization. By increasing the ratio of argon
to oxygen in the injected gas as the refining proceeds, the carbon is selectively oxidized without concurrent chromium oxidation. A typical
starting ratio is 3 to 1 oxygen to argon/nitrogen
by volume. The ending ratio can be as low as 1
to 9, oxygen to argon/nitrogen. The choice of
which inert gas to use, argon or nitrogen, is based
on cost and final nitrogen content desired. Stabilized stainless steels require low carbon and nitrogen levels, for instance, so the more expensive
argon must be used.
It is possible to use a vacuum system to keep
the partial pressure of CO low when refining
with injected oxygen. This is the vacuum oxygen decarburization (VOD) process. The VOD
process can achieve slightly lower carbon levels
but does not achieve cleaner steel as some
believe.
In both processes, after final carbon content
has been achieved ferrosilicon is added to reduce
the chromium in the slag and have it return to the
molten steel. The excellent mixing of the slag
and metal in the AOD permits this to be done
efficiently. The silicon plus the manganese in the
steel combine to reduce the oxygen content of
the steel to around 100 ppm. This could be further reduced by aluminum, but aluminum-based
inclusions are generally undesirable. The thermodynamic activity of aluminum is considerably reduced in iron as chromium levels increase, so its role as a deoxidizer is less valuable
in stainless steels. Titanium, on the other hand, is
enhanced as a deoxidizer in chromium-iron alloys, and consequently small amounts of it are
sometimes used as a supplementary deoxidant in
alloys even though an alloy specification may
not call for any. Titanium is believed to reduce
hot working defects. More active deoxidants,
such as calcium and magnesium, can be used
when required. Also note that even if no intentional addition of metallic calcium is made,
strong deoxidation with aluminum or titanium
can reduce small amounts of calcium from the
CaO in the slag, producing measurable calcium
content in the metal.
Besides carbon and oxygen, other impurities
can be removed from the molten stainless. Once
the steel has been deoxidized, sulfur can be
readily removed by contact with a basic slag.
Sulfur can be reduced to less than 0.001% in
the AOD, and this excellent purity level is commercially furnished without additional price
premium. Sulfur, although a harmful impurity
from a corrosion standpoint, is often deliberately kept at moderate levels (0.008 to 0.015%)
for tungsten inert gas (TIG) welding penetration
(see Chapter 17) and at high levels (0.15%+) for
machinability (see Chapter 15). These trade-offs,
which are beneficial to processors, should be
viewed with skepticism by end users, whose
product integrity is compromised. There are
processing methods for which higher levels of
sulfur are not necessary that are preferable to
the end user while not compromising welding
or machining costs. For example, machinability
can be improved by calcium additions that produce malleable oxides to replace the deleterious
sulfides (see Chapter 15), and welding methods,
such as laser welding, can be used in many
cases to eliminate the need for the weld penetration enhancement of sulfur while increasing
welding speeds.
Phosphorus is an impurity for which no practical removal technology exists in stainless
steel. Any known process to remove it first removes chromium. Thus, it exists in almost all
stainless steel at levels close to its normal specification limit, about 0.030% in austenitic alloys
and 0.020% or less in ferritic alloys, which are
Chapter 12: Melting, Casting, and Hot Processing / 157
on interstitial solubility. The higher solubility
of carbon, nitrogen, and oxygen in stainless
steels is significant. A manganese/silicon deoxidized stainless steel will still have about 100
ppm of dissolved oxygen at the freezing temperature as opposed to the less than 10 ppm of
oxygen found in aluminum-killed carbon steel.
This oxygen precipitates as oxides in the solid
state.
Vacuum induction melting (VIM) is another
method of melting stainless steels. This is a
nearly slag-free process, and little refining is
possible. Melt purity is largely controlled by the
purity of the starting material, and use of AOD
master melt stock for VIM remelting is common. Limited decarburization is possible via injection of oxides such as Fe3O4 or SiO2 to create
CO evolution inside the vessel. Using this technique, very low carbon levels (less than 50
ppm) are achievable commercially. Use of VIM
is generally limited to high-value, high-purity,
or low-tonnage melts.
made from a higher percentage of low-phosphorus carbon steel scrap. The deleterious effects of
phosphorus on corrosion are not avoided unless
much lower levels are achieved. Consequently,
its presence is tolerated since it has no differential effect over the range in which it is found.
Heavy metals are eliminated by high-temperature AOD blowing, as is hydrogen. Care must
be taken not to reintroduce such impurities after
refining, which is a risk when using damp or
contaminated scrap for coolant.
Alloy adjustment can be done in the AOD or
preferably in a treatment-and-transfer ladle. The
tapped molten steel generally has excess heat
from the highly exothermic refining process.
This allows the composition to be measured and
adjusted before it must be cast. This can be done
very precisely by wire feeding of alloying elements through the slag into the heat, which can
be stirred by argon bubbling via porous plugs.
This technique is very effective for the finetuning of reactive elements such as titanium.
The refining treatments used for carbon steel
and stainless steel are very similar, but there are
subtle differences because of the difference in
the thermodynamics of dilute solutions like carbon steel and highly alloyed, nondilute solutions
like stainless steel. Table 1 shows the factors by
which additions of various elements to stainless
steel (j) alter the thermodynamic activity of
other alloying elements (i).
Equation 4 is used to calculate the activity of
elements in steel. The activity coefficient γ
varies with the concentration of alloying element x by:
n
RT lnγ i = RT lnγ i0 + ∑ RT
j =1... n
δ ln γ i
δxj
Remelting
Some stainless steels and related alloys are
remelted to refine composition or ingot structure. There are two principal remelt processes:
vacuum arc remelting (VAR) and electroslag
remelting (ESR).
In VAR, the material to be remelted is cast
into a cylindrical electrode and placed inside a
cylindrical water-cooled vacuum chamber. A
high-current direct current (dc) arc is established between the electrode and a starter plate
at the bottom of the chamber. The end of the
electrode is melted, and the molten drops fall
through the intervening vacuum. Volatile constituents escape from the molten drops, and the
purified drops collect to form a molten pool on
top of the starter plate. VAR parameters are adjusted to maintain a shallow pool, which solidifies in a bottom-up fashion. The shallowness of
the molten pool produces a refined grain
(Eq 4)
This calculation is best left to computer programs such as Thermo-Calc that have been perfected for these lengthy procedures. It should
be noted that chromium, which is always present in nondilute quantities, has a powerful effect
Table 1 Influence of alloying elements on the thermodynamic activity of carbon, nitrogen,
sulfur, and oxygen
J
i
Al
C
Cr
Mn
Mo
N
Ni
O
S
Si
Ti
W
O
N
S
0
.04
–.03
...
–.39
.14
.13
.11
–.45
–.02
–.05
–.01
–.04
–.01
–.02
–.03
–.02
–.01
–.01
.003
.003
.11
0.0
.01
.06
.01
.01
0.0
.006
–.34
.05
–.27
–.20
.05
.01
–.03
–.13
.08
.05
.06
–.13
...
–.53
–.07
–.6
–.005
–.001
.01
–.01
158 / Stainless Steels for Design Engineers
structure with less solidification segregation
than found in typical cast product.
In ESR, the material to be remelted is cast
into an electrode of similar shape, but slightly
smaller than the water-cooled mold. A gap between the electrode and a starter plate at the bottom of the mold is filled with a prepared slag.
Typically, this slag is calcium fluoride-based
with high lime (CaO) content. Additional ingredients control the basicity, fluidity, oxidizing
potential, and other properties of the slag. A
high current is used to melt the slag, which in
turn melts the end of the electrode, and the
molten drops fall through the slag. Reaction of
the molten drops with the slag removes sulfur
and some other impurities, and the purified
drops collect to form a molten pool on top of the
starter plate. ESR melting typically is done at a
higher rate than VAR, and the molten pool is
deeper. This deeper pool produces a grain structure between that of VAR and typical cast product, with commensurate intermediate segregation patterns.
Casting
Continuous slab, billet, and bloom casting
have become the standard methods of making
stainless steel primary products, replacing the
obsolete ingot method. There are some alloys
that cannot be continuously cast, but these represent a miniscule percentage of stainless production. Continuous casting produces slabs directly,
thus removing the costly soaking and slabrolling processes. In a well-executed continuous
casting operation, slabs are of sufficient quality
that they require no surface conditioning before
being hot rolled. Slabs range in thickness from
13 to 63 cm (5 to 15 in.). The segregation in continuous casters is less than in ingots because of
the smaller section size. It is not eliminated,
however, and certain alloying elements concentrate at the centerline, where they defy homogenization. Carbon and molybdenum are examples
of alloying elements with this tendency.
In properly executed continuous casting, the
ladle feeds by a slide gate, or preferably a stopper rod gate, into a ceramic tube into the large
tundish situated over the caster mold. The metal
in the tundish is covered with a protective slag
cover, and flow patterns within the tundish are
designed to minimize dead spots and encourage
removal of inclusions by impingement with the
slag cover. The metal feeds through another
ceramic tube, called the submerged entry
nozzle, into the mold, which is covered with a
consumable protective and lubricating slag
cover, called a mold powder. The mold powder,
which melts in the mold as it is added, contains
ceramics, fluxes, and carbon. The level of the
molten metal should be carefully controlled by
ultrasonic measurement, or other methods, to
prevent fluctuations in level that may entrap
slag in the slab surface. The entire water-cooled,
copper alloy mold oscillates in a precise pattern
as the solidifying strand of steel is withdrawn
from the mold bottom by pinch rolls and
sprayed with water to cool it. The pinch rolls
apply enough pressure to slightly deform the
slab. This deformation has a crucial, seldom-recognized effect. It causes a beneficial recrystallization that improves hot working characteristics of austenitic and duplex alloys. In ferritic
alloys, it can cause excessive grain growth,
which detracts from hot workability. The initial
portion of slab cast in a sequence is seldom
of adequate quality to be used because of
exogenous inclusions, entrapped mold powder,
and non-steady-state solidification structure.
The defective portion must be identified and
scrapped or diverted to low-quality requirement
end uses.
The strand is bent from an initial slightly
curved shape to flat and cut into slabs. More
than one heat of steel may be cast sequentially
without restarting the process. This is ideal economically and for quality reasons since initial
and final segments of a casting can contain
more inclusions and aberrant structure. Some
end users stipulate that no first slabs be applied
to their orders. Producers generally apply first
slabs to less-critical uses or discard suspect sections of them. If casting conditions are not optimal, the result can be slabs with poor surface
quality that must be surface ground.
Slabs are sometimes quenched to avoid precipitation of phases; however, they may be held
at high enough temperatures prior to hot rolling
to stay above the temperature range in which
embrittlement can occur or to stay above the
temperature at which an embrittled slab can
fracture. Ferritic and martensitic alloys are especially prone to these problems.
There has been great interest for decades in
producing stainless steel coils directly from
the melt in so-called strip casters. Elimination
of hot rolling could be quite valuable in stainless steel, whose hot rolling from slab can be
both expensive and problematic. There are a
Chapter 12: Melting, Casting, and Hot Processing / 159
number of such machines in pilot or limited
production. They have not had sufficient commercial or technical success to have become a
factor in the industry. Since their development
is only being undertaken by those large stainless steel producers who already have the hot
rolling assets that strip casting would replace,
it seems unlikely that strip casting will soon
become a major factor even if it is perfected
technically.
Another method of shortcutting the casting/
ingot step has been perfected: the powder metallurgy approach. In powder metallurgy, the refined molten metal is atomized by gas or liquid
and made to freeze into small particles. These
particles, having been quenched extremely
rapidly, are quite homogeneous. Powder technology methods allow for the design of alloys
that would otherwise freeze with too much
segregation and too coarse a structure with
conventional production methods. Traditional
powder metallurgy production methods are
used to make small near-net shape components, avoiding most of the costly machining
steps. More impressively, powder technology
is also used to produce massive components.
For example, very high carbon/vanadium
stainless tool steel components can be made by
encapsulating powder in an evacuated canister
in which it can be sintered and hot worked to
100% density and virtually complete homogeneity. Chapter 9 on martensitic alloys discusses these materials.
Hot Rolling
Hot rolling remains an essential process for
the vast majority of stainless steel used. Hot
rolling characteristics of stainless steels vary
greatly. Ferritic stainless steels are extremely
easy to hot roll since they have a soft, singlephase structure at hot rolling temperatures.
Martensitic stainless steels roll like their carbon
and alloy steel counterparts since their microstructure during hot rolling is a moderately
alloyed austenite similar to alloy steels. The microstructure during hot rolling is the crucial factor. Austenitic stainless steels have high strength
at hot rolling temperatures. Furthermore, the
low diffusion rates in austenite slow recrystallization so that the steel does not always soften
between stands in tandem mills. This increases
mill loads, and lower reductions must be taken
than for alloy steels. Powerful hot strip tandem
mills that routinely roll carbon steel to 1.5 mm
(0.06 in.) can struggle to attain 4.5-mm (0.18-in.)
thickness for 316 stainless.
The high separating forces on the hot rolling
mill stands also cause greater roll deflection and
compression, which if not countered by roll
bending or roll shifting schemes can lead to significant variation in thickness across the sheet,
as much as 0.25 mm (0.01 in.). This variation as
a percentage of thickness is not reduced by cold
rolling and is a major cause of tolerance loss in
sheet and strip. Hot-rolled bands vary in thickness along the length of the coil because the tail
end of the slab is colder and harder to roll. Coil
boxes (on reversing mills) address this problem
to a degree by permitting the semirolled coil to
equalize in temperature.
Hot strip tandem mills powerful enough to
successfully roll high-quality stainless steel hotrolled bands are massively expensive and are
seldom justified for the tonnage of stainless steel
rolling a given melt shop produces, although
rolling stainless on hot tandem mills used primarily for carbon steel can be an excellent production method.
Hot Steckel mills have become the favored
method of hot rolling stainless steel because
their throughput better matches stainless steel
melt shop production outputs. This permits the
melt shop and caster to be adjacent to the hot
mill, which permits energy-saving hot charging
of slabs. In hot Steckel mills, typically a fourhigh reversing rougher rolls slabs to about 3-cm
(1.2-in.) thick. Then, the transfer band is rolled
to final gauge on a separate reversing four-high
finishing mill with coil boxes to preserve temperature. The economy of having only two mill
stands makes these mills ideal for typical stainless production quantities and permits the cost
of sophisticated mill capabilities, such as roll
shifting, roll crossing, or roll bending, not to
have to be duplicated among many stands. This
is the same justification for using Sendzimir
mills to cold roll stainless. In both cases, the
logic applies more to austenitic alloys than to
the easily rolled ferritic stainless alloys.
In either case, the hot-rolled band carries a
heavy, embedded scale that must be removed
from the surface before further processing in
most cases. Some alloys can be cold rolled in the
“black band” state at a cost of coarser surface
finish and greater rolling mill roll wear. If normal cold rolling or use as hot-rolled coil is foreseen, the hot-rolled band must then be annealed
and pickled since the as-rolled hot-rolled band
160 / Stainless Steels for Design Engineers
has poor corrosion resistance, poor mechanical
properties, residual cold work and hardness variations, as well as a heavy oxide layer.
Defects
Stainless steel hot-rolled bands can contain
many types of defects. These are seldom seen
by the end user because they are removed
when they are not prevented. They do have
repercussions on delivery. The major categories are:
• Hot mill defects
• Inclusion-related defects
• Hot ductility-related defects
Stainless steel is less forgiving of hot mill
faults because its surface is not removed by oxidation to the degree carbon steel’s surface is.
Thus, a skid mark from a slab-heating furnace
will remain through the hot rolling, annealing,
and cold rolling processes. This is true of all hot
mill scratches, gouges, digs, etc. Rolling stainless requires a different mindset than rolling
carbon steel, which argues against the benefits
of rolling stainless on a mill built and primarily
used for carbon steel.
Inclusion-related defects are all essentially
avoidable by using state-of-the-art technology.
Protecting metal from reoxidation and keeping
precise mold-level control in the continuous
caster prevents all inclusions of a size that can
produce a defect.
Hot ductility defects are more subtle. They
arise from many causes and are manifest prima-
rily as edge cracks and slivers. Edge cracks are
simply a lack of ductility at the colder strip edge.
Stainless hot ductility often has a narrow temperature window, and many factors can affect
the size of that window depending on alloy type.
The most inherently challenging alloys for hot
working are the duplex alloys and the alloys that
solidify in the fully austenitic state. The former
has a mixed-phase structure, and the phases can
exhibit mechanical incompatibility at certain
temperatures. The latter alloys reject sulfur and
oxygen during solidification and slab reheating
to the grain boundaries, where they form very
weak films. But, even alloys such as 304 and
316 can have very poor hot ductility if they contain much sulfur and oxygen or if they are reheated for long times or at temperatures above
1250 °C (2280 °F), which facilitates diffusion of
sulfur and oxygen to the grain boundaries and
also encourages very large grains. This poor hot
ductility manifests itself as “slivers,” which can
require grinding of the entire hot band surface.
These tendencies are fought by low oxygen and
sulfur levels and minimal slab-reheating temperatures and times, as well as slab surface working
in the caster pinch rolls. Sometimes, very poor
hot working alloys are given a single hot reduction pass on a hot mill to produce a full recrystallization that disperses grain boundary-weakening
elements on subsequent reheat.
REFERENCE
1. D. Peckner and I.M. Bernstein, Handbook of
Stainless Steels, McGraw-Hill, 1977, p 3–13
Stainless Steels for Design Engineers
Michael F. McGuire, p 161-171
DOI: 10.1361/ssde2008p161
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
CHAPTER 13
Thermal Processing
Summary
THE THERMAL PROCESSING of stainless
steel is a topic the end user should approach
with great respect. It is not simple in concept or
in practice. Before attempting to carry out any
thermal processing on stainless steel, the practitioner must understand the alloy design, composition, and processing history of the material in
question. The thermal processing then must be
designed and executed in a planned, controlled
manner. The consequences of failure in thermal
processing can become catastrophic to mechanical properties and corrosion resistance.
Introduction
The thermal processing of stainless steels can
have many purposes. Normally, the objectives
are simple: heating for hot working, annealing to
soften after cold working, solution annealing
to homogenize, heating to temper martensite, or
to stress relieve. However, even if the objective
is simple, the processes that occur are anything
but simple. Variations in temperature, times at
temperature, heating and cooling rates, and atmosphere can have complex and easily unintended consequences. There is no substitute for
understanding the processes that are occurring
when stainless steels are heated for successful
heat treating to be achieved.
Stainless steels have many alloying elements
in large amounts. Many of these elements are
highly reactive thermodynamically. The practical consequence of this is that many phases are
thermodynamically possible at different temperatures. Stainless also reacts with its environment
at high temperatures, causing changes in surface
alloy content. Some of the resulting phases are
desirable, and some are potentially very detrimental. Readers are encouraged to review the
earlier chapters on phases in stainless steel
(Chap. 6–10) to familiarize themselves with
these phases.
Each of the alloy groups of stainless steels has
radically different thermal processing objectives
and requirements; therefore, each is discussed
separately.
Austenitic Stainless Steels
Thermal processes applied to austenitic stainless steels include:
• Soaking for homogenization and preparation
for hot working
• Annealing to remove the effects of cold
work and to put alloying elements into solid
solution (solution annealing)
• Stress relieving
The temperatures at which these processes
are carried out are shown in Table 1 for typical
austenite compositions.
Soaking
Because virtually all stainless steel is continuously cast, the older soaking process of holding
ingots in soaking pits for many hours is rarely
used. The soaking had two functions. The obvious one was to equilibrate at the right temperature for hot working. The second, less-obvious,
one was to achieve greater chemical homogeneity. The lack of homogeneity comes from the
solute segregation that occurs as solute elements are rejected from the material that was
first to freeze. Solute segregation was exaggerated by the slow solidification of ingots, and
continuous casting helped make the stainless
162 / Stainless Steels for Design Engineers
Table 1
Recommended thermal processing temperatures for austenitic alloys
Alloy
Standard alloys
201, 202, 201LN
301, 301LN, all versions
304, 304L, 305, all versions
316, 316L, 316N, 317, 317L
308, 309, 309S, 310, 310S
Stabilized alloys
321
347, 348
20Cb-3
Moderately alloyed, Creq<26
S31725, N08028, JS700
Highly alloyed, lower sigma
alloys, Creq>30, high N
AL6-XN, 4565, 654SMO,
254SMO
Highly alloyed, sigma-prone
alloys, Creq>30, low N
AL6-X
Annealing
temperature, °C
Annealing
temperature, °F
ASTM A480(a)
2006, °F
Stress
relieving, °F
Stress
relieving, oC
...
...
...
1010–1120
1010–1120
1010–1120
1040–1175
1040–1175
...
955–1065
980–1025
925–955
1120–1175
1850–2050
1850–2050
1850–2050
1900–2150
1900–2150
...
1750–1950
1800–1950
1700–1750
2050–2150
1900 min
1900 min
1900 min
1900 min
1900 min
...
1900 min
1900 min
...
Various
1500–1600 non-L,
1000–1600 L grades
...
...
...
...
...
1000–1600
...
...
925–1010
1500–1600
815–870 non-L,
540–870 L grades
...
...
...
...
...
540–870
...
...
...
815–870
1120–1175
2050–2250
Various
Not recommended
...
1205–1230
2200–2250
Various
Not recommended
...
(a) Standard specification for general requirements for flat-rolled stainless and heat-resisting steel plate, sheet, and strip.
steel more homogeneous. Nevertheless, cast
slabs and blooms must be soaked to eliminate
as-cast segregations. This process, to the extent
it is done, occurs as they are reheated to the appropriate temperature for hot working. Soaking
dissolves the few percent of residual delta ferrite that are present on slab solidification. It is
important to soak at the highest temperature at
which delta ferrite is not a stable phase so that it
will dissolve, about 1250 °C (2280 °F) for most
austenitic stainless steels.
Soaking at higher temperatures causes ferrite
levels to increase, negating the homogenization
and causing very poor hot workability. Longer
times at temperature than the minimum required
for thermal uniformity also cause problems as
any sulfur and oxygen impurities are rejected
from austenite and can diffuse to grain boundaries, where they form weak, plastic films that
also degrade hot workability. Grain growth, by
reducing grain boundary area, exacerbates this
effect. Thus, soaking times are best minimized
and closely controlled. Alloys are therefore designed to have only a slight amount of delta ferrite to be redissolved during soaking. Ferrite is
useful because it has a high solubility for oxygen
and sulfur. Having none would result in impurity
rejection of these elements to grain boundaries
during solidification, which results in the worstpossible hot working characteristics. The oxygen
and sulfur trapped in the ferrite during solidification precipitates in the solid state as inclusions,
which also must be equilibrated with the sur-
rounding matrix by sufficient soaking. Welds that
are unannealed have such precipitated inclusions
in an unequilibrated state, and the result is diminished chromium concentration and poorer pitting
resistance.
Annealing
Annealing serves two main functions in
stainless steel: It removes the effect of cold
work by replacing strained microstructure with
new strain-free grains, that is, recrystallization.
New grains nucleate and grow. If stored strain
energy is insufficient, as happens often with
ferritic stainless steels, true recrystallization is
difficult to achieve, and the annealing process
may only produce recovery without recrystallization, leaving the same grains relieved of
stress. This leaves the surviving grains with the
same crystallographic orientation that deformation produced and may or may not be the desired outcome. Second, annealing returns into
solution solute that has been precipitated as unwanted phases, principally carbides, but also
intermetallic phases. Annealing also may help
to reduce solute segregation remaining from
the casting process, making the composition
more homogeneous. The homogenization
process is accelerated by the reduction in dimensions from hot and cold working. A reduction in dimension by a factor of two reduces the
time to achieve a given degree of homogenization by a factor of four.
Chapter 13: Thermal Processing / 163
Annealing to recrystallize is fairly rapid. To a
first approximation, it is instantaneous, and the
results are merely a function of the maximum
temperature attained. This may not be the case
for continuous annealing lines, in which transit
time can be short enough, less than a minute at
temperature, to limit the grain size attained. The
driving force for recrystallization is the strain
energy stored in the lattice from deformation.
The strain energy in a given material is proportional to the square of the material’s flow stress.
As the material is heated, recovery occurs first.
This is the change in physical and mechanical
properties associated with dislocation annihilation and polygonalization that occurs before the
nucleation and growth of new grains.
The nucleation of new grains occurs at highangle grain boundaries and proceeds by the
movement of roughly hemispherical growth
fronts into strained areas. The percentage recrystallized, once a sufficient temperature is
reached, grows sigmoidally. It is normal for
the time to fully recrystallize to be rather less
than the time to attain that temperature. Even
at the lower range of annealing temperatures,
times are generally less than 1 min. Recrystallization will not occur if the lattice contains insufficient strain energy. Thirty percent cold
work should be used as a rough threshold for
the required amount. Annealing after lower
amounts of cold work is characterized by
scarce nucleation sites and can result in very
large and nonuniform grain size. Hot-worked
material often has a composite structure that
may have already had some recrystallization
depending on the final reduction temperature.
Annealing may not produce a clear recrystallized structure in this case.
The relative rapidity of recrystallization annealing is due to the fact that it is rate controlled by short-range diffusion. Solution annealing requires longer-range diffusion and
thus can require much longer times. Some studies have shown that welds, for instance, do not
recover completely from their loss of corrosion
properties that arise from local alloy depletion
unless they have been annealed for times on the
order of 1 h (Ref 1). Others have seen homogenization in as little as 10 min (Ref 2). Wrought
materials can require much shorter times because reductions during hot working have reduced diffusion distances. It should be noted
that precipitates can be redissolved and not apparent in the annealed microstructure without
full homogeneity being achieved. For example,
carbides can be redissolved and carbon diffused away from the carbide, but this does not
mean that all composition gradients have been
reduced to zero. This means that precipitates
may re-form more rapidly in such a material
than they would in a completely homogeneous
alloy.
The annealing temperature for a given alloy is
chosen based on the temperature required to put
all alloying elements into solution. Higher carbon levels, for instance, require higher temperatures to dissolve all the carbon. This is one of
the principle values of accurate phase diagrams.
If it were simply a consideration of recrystallization, all alloys could be annealed at similar
temperatures at the low end of the recommended range. Within the recommended range,
the temperature selected should be determined
by the desired grain size. End use determines
whether a fine or coarse grain size is preferable.
Table 1 lists recommended annealing temperatures for austenitic stainless steels.
The overall interplay between prior cold work
and annealing temperature on mechanical properties of annealed material can be summarized
as (Ref 3):
• Grain size of a given alloy is the most important parameter in characterizing mechanical properties.
• Yield and tensile strength are essentially
constant for a given grain size regardless of
the amount of prior cold work; however, the
elongation depends on the prior reduction.
• Yield strength, tensile strength, and hardness
are essentially linear functions of grain size.
• Elongation decreases with finer grain size
and at an increasing rate as grain size becomes finer as long as cross-section size of
the test specimen is not extremely small.
This is not true for very coarse-grained material.
• Maximizing elongation comes from maximum prior cold work and medium annealing
temperatures
• Anisotropy coefficients, or plastic strain ratios r are constant up to about 40% reduction
after, which r45 and rn increase sharply, while
rt decreases. This leads to earing during
drawing.
• The increase in properties for a one ASTM
grain size increment is:
a. 13 MPa (2 ksi) for tensile strength
b. 20 MPa (3 ksi) for yield strength
c. 2 HRB for hardness
164 / Stainless Steels for Design Engineers
Atmospheres for annealing are important.
Austenitic stainless steels heated in air, of
course, form oxide scales. Beneath this oxide,
the metal matrix becomes significantly depleted
of chromium (Ref 4), often more than 5% lower
in chromium and to a depth of as much as 10 µ
(395 µin.). So, not only must any oxide be removed, so must the chromium-depleted layer.
This requires aggressive pickling, which while
done commonly, may not be practical for many
stainless users. The chromium-depleted zone,
however, does pickle rapidly precisely because
it does have less chromium. To avoid oxide
scale formation, vacuum, hydrogen, or inert gas
atmospheres may be used.
If vacuum is used, it should be less than
2 × 10–3 torr (0.3 Pa). If an inert gas or hydrogen
is used, the key consideration is moisture content. The dew point must be –40 °C (–40 °F) or
lower. More stringent levels may be required if
mirror finishes are desired. Cool down must be
rapid as oxidation potential increases as temperature decreases. Vacuum or inert gas is preferable
to hydrogen for alloys containing stable oxide
formers such as aluminum or titanium or for alloys containing boron.
Austenitic alloys that are subject to sensitization must be cooled rapidly enough from annealing temperatures to avoid carbide precipitation
during cooling. If forced air or water quenching
are impractical or if section size prohibits rapid
cooling, then using stabilized or low-carbon
grades is indicated.
Superaustenitic stainless steels, and even alloys like 317, present a special problem because
these alloys have significant sigma-forming tendencies. Sigma forms initially because solidification segregation causes local enrichment of
sigma-promoting elements, such as molybdenum. It can also form from slow cooling of slabs
or hot bands. This latter sigma forms at grain
boundaries and will cause embrittlement and reduced corrosion resistance, so it must not only
be redissolved, but also the alloy must be homogenized to remove the residual concentration
gradients from the sigma. If this is not done,
chromium- and molybdenum-depleted regions
will still exist, and sigma will re-form much
more rapidly during subsequent exposure to
high temperatures. For this reason, the higher
ends of the annealing ranges are recommended,
and annealing times should be generous. Newer
alloys have higher nitrogen contents to suppress
formation of sigma and other deleterious inter-
metallic phases. Use of high-chromium andmolybdenum alloys without enhanced nitrogen
is no longer recommended, and the use of lowernitrogen alloys should be reexamined and questioned if specified.
Last, stainless surfaces should be scrupulously clean before annealing. Even hard water
deposits can cause differential oxide growth,
which can cause etched spots on the surface,
where the postanneal pickling attacks the different oxide more strongly. Carbonaceous materials left on the surface are even more objectionable because they can cause carburization and
subsequent loss of corrosion resistance.
Stabilizing anneals are sometimes conducted
on stabilized alloys such as 321 and 347. This is
useful when carbon levels are sufficiently high
that significant dissociation of carbides occurs
at annealing temperatures. A second anneal at
lower temperature, about 900 °C (1650 °F),
then is done to permit the carbon to combine
with the stabilizing element rather than leaving
it available to form chromium carbides. Current
preferred practice for these alloys is to maintain
carbon and nitrogen below 0.03% for corrosionresistant service, which renders this stabilizing
unnecessary. Alloys used for high-temperature
service benefit from the creep-resisting contributions of higher carbon levels.
Stress Relieving
Austenitic stainless steel weldments often contain residual stresses, which can cause distortion
or lead to stress corrosion cracking in service.
They are commonly stress relieved at temperatures slightly below the annealing temperature,
so that residual stresses may be relieved by
creep. One hour at 900 °C (1650 °F) reduces
residual stress by about 85%. Lower temperatures require exponentially longer times for the
same stress relief, with times doubling for each
100 °C (180 °F) decrement as decreasing diffusion rates, which govern creep, are encountered.
Cold-worked austenitic stainless steels have a
markedly diminished proportional limit, particularly in compression. This Bauschinger effect,
which arises from the easy mobility of dislocations, can be eliminated by stress relieving at
around 350 °C (660 °F) for 2 h, which provides
the thermal energy for dislocation interactions
to lock into place. This produces a sharp yield
point without premature nonproportional elastic
deformation.
Chapter 13: Thermal Processing / 165
Ferritic Stainless Steels
Ferritic stainless steels, from an annealing
point of view, must be discussed in two categories. First are the modern, stabilized alloys,
which are ferritic at all temperatures. These alloys behave as interstitial-free (IF) alloys because the interstitial carbon and nitrogen are removed from solution as a stable precipitate. In
the second category are the older ferritic steels,
which have enough austenitizing elements, usually carbon, in solution to cause them to form
austenite at what would otherwise be a good annealing temperature. This makes them truly
quasi-martensitic alloys, and they must be
treated accordingly. Table 2, which lists heattreating temperatures for ferritic stainless alloys,
also shows which grades fit into which category.
Soaking
Heating of ferritic stainless for hot working is
straightforward. Whether stabilized or not,
these alloys are heated to the 1000 to 1100 °C
(1830 to 2010 °F) range for hot working. The
superferritics can be heated to up to 1300 °C
(2370 °F). At this temperature, no debilitating
phases occur, and ductility is good. The high
diffusion rate inherent to the ferritic structure
makes homogenization easy. As long as hot
working is completed at temperatures above
that at which austenite forms, good hot ductility
is expected. This is not a concern with IF alloys,
which do not form austenite.
Annealing
The IF ferritics do not undergo any phase
change with temperature during the course of
properly executed heat treatment. The objective
of annealing is generally simply to remove the
effects of cold work. This is because they do not
need to have carbon put into solution and, except in rare cases, do not have intermetallic
Table 2 Recommended annealing temperatures
for ferritic alloys
Alloy
Stabilized, Cr+Mo<20
409, 439,18 SR
Unstabilized, Cr+Mo<20
405, 430, 434, 436
Stabilized, Cr+Mo>20
29-4C, Monit, Seacure, 444
Unstabilized, Cr+Mo>20,
446
Annealing
temperature, oC
Annealing
temperature, oF
870–925
1600–1700
705–790
1300–1450
1010–1065
1850–1950
760–830
1400–1525
phases that require dissolution. Alloys with high
chromium and molybdenum contents can form
σ and/or α', the brittle, ordered body-centered
cubic (bcc) phase, at temperatures below annealing temperatures, so rapid cooling is prudent when chromium plus molybdenum exceeds 20%.
The driving force for recrystallization in these
alloys is limited by the lower stored energy
from deformation inherent in the bcc structure.
In addition, the pronounced deformation texture
leads to annealing responses that are more accurately characterized as recovery and grain
growth with diminished recrystallization. These
alloys retain this texture after annealing, and
this characteristic anisotropy is exploited for
good drawability. The major concern is to avoid
excessive annealed grain size, which greatly reduces toughness. Anneal at the higher end of the
range only if the loss of toughness associated
with large grain size is not a concern. Stabilizing anneals are normally unnecessary for stabilized ferritics as their high diffusion rates ensure
freedom from knife-edge attack due to sensitization from free unbonded carbon combining
with chromium at grain boundaries. The stabilizing additions of titanium and/or niobium tie
up the carbon as stable TiC or NbC, which does
not redissolve during annealing.
The interstitial-bearing ferritic stainless
steels must be annealed subcritically, or the formation of austenite at higher temperatures
would make martensite formation on cooling
virtually unavoidable. Thus, a typical primary
anneal cycle for a typical alloy such as 430
would be nearly 24 h at 750 °C (1380 °F), the
majority of which is thermal equilibration of
the large coil mass. The actual time at temperature required is less than 1 h. Continuous annealing is not practical because the diffusion of
carbon is too slow to occur in the dwell time at
temperature typical in continuous annealing
lines. This cycle also precipitates essentially all
the carbon and nitrogen as mixed Cr/Fe carbides and nitrides and homogenizes chromium
content. This necessarily slow process permits
subsequent subcritical annealing for mechanical properties (to alleviate the effects of cold
work) to be done in a few minutes since carbon
has been eliminated from solution by the formation of fairly stable carbides. Since the material is generally purchased in the annealed
condition, the user need never be concerned
with such lengthy anneals.
166 / Stainless Steels for Design Engineers
Stress relieving is rarely a concern for any
type of ferritic stainless. Unstabilized grades
should not be welded, and if they are, full subcritical annealing is required. Stabilized grades
have no need for postweld heat treatment. Lowtemperature heat treatment runs the risk of α'
formation and is best avoided.
Martensitic Stainless Steels
The martensitic stainless steels resemble the
unstabilized ferritic stainless steels described.
The martensitic stainless steels form essentially
100% austenite on heating and have very high
hardenability, so their ability to be softened by
annealing is limited. The traditional martensitic
stainless steels are iron/chromium/carbon alloys, sometimes with a small amount of nickel
and/or molybdenum. More recently, alloys have
been developed for petroleum applications that
contain high copper, nickel, and/or molybdenum and low carbon. The principles of heat
treatment of the two alloy categories are the
same. The more highly alloyed newer alloys
are, in fact, simpler to heat treat because their
low carbon and nitrogen levels alleviate the
need to temper.
martensite cannot be avoided by furnace cooling from austenitic temperatures, then only subcritical annealing is feasible. But, even for
nickel-free alloys the hardenability is so great
that annealing by slow cooling is quite difficult.
Martensitic alloys are put into the annealed condition for processing before they are quenched
and tempered for their final use. Thus, the more
economic subcritical anneal is the predominant
annealing heat treatment.
The nickel-bearing alloys have such high
hardenability that annealing in the critical range
cannot produce softening by any practical cooling rate, so subcritical annealing is always recommended for these alloys. Nickel reduces the
temperature at which austenite is stable as
shown in Chapter 9, Fig. 9. Other additions like
vanadium, molybdenum, and tungsten promote
secondary hardening and tempering resistance,
and subcritical annealing of these alloys becomes a slow, difficult process. This is a characteristic of the so-called super 12Cr alloys.
Martensitic alloys have lower corrosion resistance in the annealed condition than in the
hardened condition because in this state they
have the maximum amount of chromium tied up
as chromium carbide.
Austenitizing
Soaking
Hot working should be carried out in the
austenitic range. Temperatures for this are
listed in Table 3. Forging and hot working
should always be followed by annealing to
avoid stress cracking due to the deep hardening
of these alloys.
Annealing
Martensitic stainless steels can be annealed
by subcritical anneal and sometimes by full anneal depending on alloy level. If the alloy level
is such, as in the nickel-containing grades, that
Table 3
Table 3 lists the austenitizing and tempering
ranges for martensitic stainless steels. Full
austenitizing is crucial to producing a microstructure that is fully martensitic. Only austenite
transforms to martensite. If other constituents,
such as δ ferrite or carbides, exist during the
austenitizing heat treatment before quenching,
they will not transform to martensite. Some
alloys, such as the 440 group, have enough carbon that the austenitizing temperature determines how much carbon is put into solution.
The carbon in solution in the austenite will become the carbon level in the martensite, which
directly determines strength and corrosion
Recommended annealing, austenitizing, and tempering temperatures for martensitic alloys
Alloy
Subcritical anneal,
o
C (oF)
Full anneal,
o
C (oF)
Austenitizing,
o
C (oF)
Tempering, low
range, oC (oF)
Tempering, high
range, oC (oF)
Straight
Cr, C<0.20, 410,
416,403
Ni/Mo, C<0.20,
414, 431, 415,
425, C>0.20
440A,B,C F, 420
650–760
(1200–1400)
830–885
(1525–1625)
925–1010
(1700–1850)
205–370
(565–605)
400–700
(1050–1125)
620–705
(1150–1300)
Not recommended
980–1065
(1800–1950)
205–370
(565–605)
400–700
(1050–1125)
675–760
(1245–1400)
845–900
(1555–1650)
1010–1065
(1850–1950)
150–300
(300–700)
Not recommended
Chapter 13: Thermal Processing / 167
resistance because undissolved carbides contain
chromium, which diminishes that available for
corrosion resistance. Austenitizing temperature
and holding time become most significant when
carbon exceeds 0.20%, where its solubility is a
steep function of temperature.
The ␦ ferrite is a generally undesirable phase
that can be produced by temperature excursions
or composition variations. Excessive austenitizing temperatures can cause its formation, as can
low levels of carbon, which may be originally
present in the alloy or arise from decarburization. It will cause lower hardness and toughness
if present.
Heating rates should be such that a uniform
temperature is attained before the allotropic
transformation from bcc to face-centered cubic
(Fcc), which involves a more than 1% linear dimension change and can cause distortion or
cracking. Oxidation during austenitizing can
cause serious carbon loss on the surface, which
will result in serious loss of surface hardness.
Heating 410 in air for 10 min at 1100 °C (2010
°F) can cause surface carbon to decrease by
one-half, lowering hardness from HRC 45 to
under 20.
Quenching rate is not a significant issue for
the martensitic stainless steels since they have
such high hardenability, but some, especially
those with higher carbon levels, may have retained austenite, which can lower hardness and
cause problems with dimensional stability. The
quenching rate must be sufficient, however, to
avoid precipitation of carbides in the austenite
during quenching since the sensitization would
persist in the final microstructure. If this occurs, a subzero treatment at below –75 °C
(–100 °F) should be undertaken immediately
after quenching.
Tempering
Untempered martensite has insufficient toughness to be a useful engineering material. During
tempering, carbon is precipitated from the supersaturated state it is in when it is quenched into
the bcc martensite structure during the diffusionless transformation. The strain energy associated
with the lattice strain of the bcc martensite
caused by the poor fit of the carbon in the tetragonal interstices is very large. Heating to even
low temperatures allows carbon enough mobility to diffuse and precipitate as carbide. Since
carbon diffuses 106 times as fast as iron,
chromium, or other carbide formers, it tends to
precipitate with the abundant iron atoms first. At
higher temperatures and longer times, more thermodynamically stable carbides, such as Cr23C6,
form. Carbide formation is a complex function
of temperature, time, and composition. The
growth of carbides reduces strain and hardness.
There are exceptions, such as the precipitation of
Mo2C, whose morphology produces a precipitation hardening (PH), called secondary hardening. Niobium and vanadium also form carbides
that result in higher hardness at all tempering
temperatures. Had this been understood early
on, these steels could have been correctly included in the PH stainless group. The PH steels,
AM-350 and AM-355, both derive their PH
from the precipitation of Mo2C and Mo2N. In all
other cases, higher tempering temperatures lead
to lower hardness.
The nickel-bearing alloys have a restricted
upper tempering temperature because of the
danger of re-forming austenite, which would
then transform to untempered martensite during
cooling, requiring a second tempering operation.
Intermediate temperatures can lead to the
phenomenon of temper embrittlement, which is
caused by the precipitation of phosphorus and
other species, such as, but not limited to, carbides, at prior austenitic grain boundaries. This
phenomenon is distinct from the precipitation of
α', which causes the so-called 475 °C embrittlement, which occurs more severely in alloys
with higher chromium levels. Because the formation of martensite is diffusionless, the
austenite boundaries maintain the microcompositions they have at high temperatures. Austenite has low solubility for impurities such as
phosphorous, so phosphorous is more highly
concentrated in the grain boundary regions.
This compositional inhomogeneity can be eliminated by higher tempering temperatures or by
the addition of molybdenum, which combines
with the phosphorus and prevents the embrittlement. The existence of temper embrittlement is
the reason that Table 3 recommends avoiding
certain temperature ranges for tempering.
Lower tempering temperatures and higher
austenitizing temperatures are best for corrosion
resistance because both minimize the amount of
chromium tied up as carbide. Quenching from
higher temperatures also enhances toughness.
Stress Relieving
If quenched martensitic steels are not immediately tempered, then they should be promptly
168 / Stainless Steels for Design Engineers
stress relieved. Otherwise, the residual stress
from quenching could result in stress corrosion
cracking even in seemingly benign environments. Stress relieving is simply a low-temperature tempering operation, but some elimination
of residual stress does occur. Higher temperatures and longer times produce greater stress
relief and maximize elastic properties, but optimal toughness is obtained at intermediate times
and temperatures.
Precipitation Hardening Stainless Steels
There are three classes of PH steels, and each
requires totally different heat treatment. The
classes are martensitic, semiaustenitic, and
austenitic. The most straightforward alloys are
the martensitic PH grades. Like the plain martensitic alloys, the martensitic PH alloys are carefully designed to produce a nearly fully austenitic
structure at high temperature that quenches to a
nearly fully martensitic structure on cooling. The
martensite is low in carbon, so it is relatively soft
and not prone to brittleness. So, the hardness and
strength of these alloys is derived from a subsequent tempering-type heat treatment during
which various constituent elements form extremely fine coherent precipitates that greatly
strain and therefore strengthen the matrix. There
are numerous precipitates that can provoke this
effect, and they are described in detail in the PH
chapter (see Chapter 10). All require the shortrange diffusion of substitution elements to form
these optically invisible precipitates.
Martensitic PH Grades
Solution treatment of these alloys is conducted to achieve a fully austenitic structure.
The constituent elements are easily dissolved,
so excessive temperature or time is unnecessary
and could be counterproductive if it were to result in ferrite formation or surface oxidation,
which would be detrimental to final mechanical
properties. The presence of retained ferrite is
mainly a function of alloy and composition
within the allowed range. Earlier grades such as
17–4 and the obsolete stainless W intrinsically
contained some ferrite. The subsequent alloys
are substantially ferrite free. Most alloys may
retain some austenite after quenching to room
temperature, in which case subzero treatment
should be done within 24 h to avoid further stabilizing the austenite. Subzero treatment, by
eliminating retained austenite, enhances dimensional stability but diminishes toughness. Table 4
lists the solution treatments for all PH grades.
The as-quenched state is called condition A.
This is the normal condition in which the material is supplied from the mill and is intended to
be soft enough for machining and some forming.
If softer material is required, the H-1150M condition exists in which the material is first highly
overaged at 760 °C (1400 °F), allowing some
austenite to re-form. The subsequent aging then
overages that martensite while retaining some
stable austenite. The result is a very tough
microstructure.
Aging. The time and temperatures required
to produce this precipitation are also given in
Table 4. The condition code itself tells the temperature at which the aging is conducted in
that the code numbers are based on thermal
processing temperatures expressed in degrees
Fahrenheit (oF), for example, TH 900 means
transformed to martensite (T) and aged at 900
degrees Fahrenheit. The final properties are a
function of both aging time and temperature.
Lower temperatures result in higher possible
hardness but lower toughness. The precipitates, as mentioned, are optically invisible and
cause very little dimensional change. Contraction on the order of 0.0005 in./in. from aging is
typical, often permitting machining to final dimensions in condition A. All aging treatments
are above the temper embrittlement range to
which these alloys are susceptible. Service
temperatures in this range would result in embrittlement, so use above 350 °C (660 °F)
should be avoided. Molybdenum-bearing
grades should be selected to minimize this
phenomenon if high-temperature use is contemplated.
Solution Heat Treatment and Conditioning.
The semiaustenitic grades are more complicated
than the martensitic PH alloys. These alloys are
designed so that they are austenitic when
quenched from the solution heat treatment temperature. This is also called condition A, and it
permits them to be highly formable. This stabilization of the austenite comes mainly from
higher chromium and carbon levels. These alloys essentially resemble a lean 301 austenitic
alloy, many with some molybdenum substituting for part of the chromium. The key to these
grades is making them behave as a martensitic
alloy. This is done by precipitating some of the
carbon as chromium carbide at a temperature at
the high end of what would normally be consid-
Chapter 13: Thermal Processing / 169
Table 4
Recommended annealing and stress-relieving temperatures for martensitic PH grades
Alloy
Martensitics
13-8
15-5
17-4
450
455
465, 275
475
Semiaustenitics
17-7, 15-7
AM-350
Condition code
Solution anneal
A
RHxxx
Hxxx
A
Hxxx
A
Hxxx
A
Hxxx
A
Hxxx
A
Hxxx
A
Hxxx
925 °C 15 min, oil or air cool below 15 °C
925 °C 15 min, oil or air cool below 15 °C
925 °C 15 min, oil or air cool below 15 °C
1035 °C 30 min, oil or air cool below 15 °C
1035 °C 30 min, oil or air cool below 15 °C
1035 °C 30 min, oil or air cool below 15 °C
1035 °C 30 min, oil or air cool below 15 °C
1035 °C for 1 h, water quench
1035 °C for 1 h, water quench
830 °C for 1 h, water quench
830 °C for 1 h, water quench
980 °C for 1 h, cool rapidly
980 °C for 1 h, cool rapidly
925 °C 1 h, air cool
925 °C 1 h, air cool
A
T
1065 °C for 30 min, air cool
1065 °C for 30 min, air cool
C
R
1065 °C for 30 min, air cool
1065 °C for 30 min, air cool
THxxx
1065 °C for 30 min, air cool
CHxxx
RHxxx
1065 °C for 30 min, air cool
1065 °C for 30 min, air cool
1010–1065 °C
DA (double aged)
...
930 °C for 90 min, air cool,
730–760 °C 180 min
...
930 °C for 90 min, air cool
930 °C for 90 min, air cool.
180 min at –75 °C
930 °C for 90 min, air cool,
180 min at –75 °C
930 °C for 90 min, air cool.
730–760 °C 180 min
930 °C for 90 min,
air cool, 730–760 °C 180 min
A
1025–1040 °C
L (equivalent to T)
SC (equivalent to R)
SCTxxx 850 °F or
1000 °F
DA (double aged)
Equalized and
overtempered
Austenitic
A-286
–75 °C 8 h
–75 °C 8 h
...
...
...
...
...
...
...
...
...
–75 °C 8 h
–75 °C 8 h
–75 °C 8 h
–75 °C 8 h
...
760 °C (1400 °F), 90 min,
air cool to RT for 30 min.
Cold reduce
955 °C (1750 °F) 10 min, air
cool, chill to –75 °C for 8 h
760 °C (1400 °F)
90 min, air cool to RT for 30 min
Cold reduce
955 °C (1750 °F) 10 min, air
cool, chill to –75 °C for 8 h
...
930 °C for 90 min, air cool
930 °C for 90 min, air cool,
180 min at –75 °C
930 °C for 90 min, air cool,
180 min at –75 °C
A
1010–1065 °C
L (equivalent to T) 1010–1065 °C
SC (equivalent to R) 1010–1065 °C
SCTxxx 850 °F or
1000 °F
Am-355
Conditioning
ST1650
ST1650A
ST1650DA
900 °C 120 min, oil/water quench
900 °C 120 min, oil/water quench
900 °C 120 min, oil/water quench
ST1800
ST1800A
980 °C 120 min, oil/water quench
980 °C 120 min, oil/water quench
...
Aging
xxx °F for 4 h
xxx °F for 4 h
xxx °F for 4 h
xxx °F for 4 h
xxx °F for 4 h
xxx °F for 4 h
xxx °F for 4 h
xxx °F for 4 h
...
...
...
...
xxx °F for 4 h
xxx °F for 4 h
xxx °F for 4 h
...
...
xxx °F for 180 min
450–540 °C 180 min
...
...
...
xxx °F for 180 min
440–470 °C
for 180 min
540–590 °C for
180 min
...
730 °C 16 h
730 °C 16 h,
650 °C 8 h
730 °C 16 h
PH, precipitation-hardenable; RT, room temperature.
ered sensitization, 760 °C (1400 °F). The carbide precipitation occurs at the interfaces of the
small amount of residual ferrite these alloys
have and also at grain boundaries. The deletion
of carbon and chromium from the matrix
changes the matrix composition sufficiently that
the temperature for the start of the martensitic
transformation (Ms) of the depleted austenite increases from below zero to about 65 to 100 °C
(18 to 212 °F). The martensitic transformation
finishes (Mf) near room temperature. This
process is called austenite conditioning. The
heat treatment scheme just described would be
condition A-1400. This material, after quenching to room temperature, would be said to be
condition T.
The higher the temperature of the conditioning, the less carbon is precipitated and the lower
the resulting Ms. The highest conditioning temperatures of 955 °C (1750 °F) cause a sufficiently
170 / Stainless Steels for Design Engineers
low Ms that subzero treatment is required to
obtain the fully martensitic structure required for
age hardening. This would be called A-1750.
After subzero treatment at –73 °C (–100 °F), it
would be called condition R-100.
It is even possible to obtain the fully martensitic structure by cold work from full condition
A, which is quenched from 1065 °C (1950 °F).
This is called condition C and requires heavy
cold rolling to accomplish. Since there is no heat
treatment to precipitate carbon from the austenite matrix, the resulting martensite is the hardest.
The aging treatments of the semiaustenitic
alloys are identical to those for the martensitic
alloys because the treatments are standardized.
The resulting mechanical properties vary in a
complex fashion with the alloy composition and
thermomechanical treatment history before
aging. The principles at work that determine the
mechanical properties, besides the strain induced by the precipitate phase, are:
• Carbon level of the quenched martensite
• Amount of cold work of austenite prior to
aging
These factors do not come into play with the
martensitic PH grades because all of the carbon
is in solution, and they are not significantly cold
worked prior to aging. The semiaustenitic PH
alloys, because of the variable amount of carbide precipitation, have an equally variable
chromium content, so corrosion resistance will
vary accordingly. The highest-temperature solution and carbide precipitation treatments provide the best corrosion resistance, as do the
lowest aging temperatures.
Austenitic PH Alloys
In contrast to the martensitic and semiaustenitic PH alloys, the austenitic PH alloys
are austenitic under all conditions. Their
strengthening reaction comes from precipitation
of titanium- and aluminum-nickel intermetallic
phases within the austenite matrix as occurs in
superalloys. Solution treatment is very typical
Table 5
for a normal solution anneal of non-PH austenite. The subsequent precipitation aging requires
higher temperatures and longer times because
diffusion is much slower in austenite.
Duplex Stainless Steels
Duplex stainless steels are both ferritic and
austenitic, so their heat treatment combines the
same elements and principles as their principle
phases.
Soaking
Duplex alloys are multiphase at all useful
working temperatures, making their hot workability quite poor. It is extremely important to
drive sulfur to the lowest possible levels, less
than 0.001%, to achieve satisfactory hot ductility. Otherwise, soaking should proceed the same
as for ferritic stainless alloys since ferrite constitutes the continuous phase to be worked.
Because these alloys always contain high
chromium and generally high molybdenum,
they should be cooled as rapidly as possible
from high temperatures to avoid formation of
sigma or other intermetallic phases.
Annealing
The function of annealing in the duplex alloys
is generally to:
• Remove the effects of cold work
• Restore the balance between the volume
fraction of ferrite and austenite
• Achieve equilibrium composition within
both the austenite and ferrite
• Dissolve unwanted intermetallic phases
The annealing range of duplex alloys is somewhat restricted, approximating the overlap between what each of the two constituent phases
would be annealed at separately. Table 5 lists
the normal annealing temperatures for these alloys. The use of nitrogen as a key alloying
Recommended annealing and stress-relieving temperatures for duplex alloys
Alloy
Lean duplex, Cr+Mo<23,
2003-2101, 19-D-2304
Medium alloy, Cr+Mo<26,
2205
Cr+Mo>26, 2507, 52N+,
Zeron 100, 255
Annealing
temperature, °C
Annealing
temperature, °F
ASTM A480 2006
Stress relieving, °F
1010–1100
1850–2010
Various
Not recommended
1040–1100
1900–2010
1040 min
Not recommended
1050–1150
1925–2100
Various
Not recommended
Stress relieving, °C
Chapter 13: Thermal Processing / 171
element has improved the annealing behavior of
these alloys since its diffusion is quite rapid,
causing ferrite-austenite equilibrium to be attained very rapidly. Nitrogen also hinders (formation and facilitates the dissolution of secondary austenite, which can form after quenching
from welding temperatures and cause regions of
poor corrosion resistance.
These alloys are not very susceptible to carbide sensitization and normally have very low
carbon content. Thus, the guiding principle in
annealing is simply to achieve phase balance
and avoid cooling so slowly that intermetallic
phases may form.
The strengthening of duplex is normally
achieved by the strong grain refinement and
solid solution hardening. No strengthening heat
treatments are used. Stress relief would have to
occur at temperatures at which embrittling reac-
tions, either from ␣' beginning at 350 °C (660
°F) or from ␣/␹, which takes over at 600 °C
(1110 °F), would occur and so is not indicated
for these alloys.
REFERENCES
1. A. Garner, The Effects of Autogenous
Welding on Pitting Corrosion in Austenitic
Stainless Steel, Corrosion, Vol 35 (No. 3),
1979, p. 108–114
2. J.F. Grubb, personal communication, June
4, 2006
3. Data courtesy J&L Specialty Steel, October
2, 1998
4. J.F. Grubb, Proceedings of the International
Conference on Stainless Steels, 1991,
Chiba, ISIJ
Stainless Steels for Design Engineers
Michael F. McGuire, p 173-180
DOI: 10.1361/ssde2008p173
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
CHAPTER 14
Forming
Summary
STAINLESS STEELS RANGE in formability
from the extremely formable austenitic alloys to
the hard-to-form martensitic alloys. For steels
with equivalent corrosion resistances, formability increases with the level of austenitizing alloying elements. The ferritic alloys are the least
alloyed, least expensive, and least formable;
duplex steels are next, and austenitic steels are
the most formable but most expensive. However, if the specific structural anisotropy of
ferritic alloys, which gives particularly good
deep-drawing characteristics for a given level of
ductility, can be exploited, the best forming
economies can be gained.
Introduction
The technology for forming stainless steel is
quite similar to that for forming carbon steels.
The primary difference is the higher strength of
all types of stainless steels compared to drawingtype carbon steels. This higher strength requires
greater sophistication in tooling and lubricants,
and it requires more powerful forming equipment. The higher initial strength is also
accompanied by a higher work hardening rate in
austenitic stainless steels, which further distinguishes them from carbon steel. Galling also
must be recognized as a danger and prevented.
Stainless steel has lower thermal conductivity
than carbon steels, which can cause it to retain
heat from deformation and friction, thereby
decreasing lubricity. Last, stabilized stainless alloys contain abrasive carbide microconstituents.
The various types of stainless steel have very
different deformation characteristics in terms of
strain hardening and anisotropy. It will be
shown that it is important to understand and
exploit these characteristics to optimize forming
of stainless steels.
Flat, Rolled Stainless Steel
The vast majority of carbon steel and, especially, stainless products are flat products. These
steels are formed by bending, roll forming, spinning, hydroforming, and deep drawing. Roll
forming is most commonly used to produce
welded pipe and tubing and is simply bending
done in a continuous manner. Bending is a simple operation, and there is no meaningful change
in thickness of the sheet during the operation.
The higher yield strength and work-hardening
rates of most stainless steels will result in greater
springback than would be experienced in carbon
steel. Tooling must be adjusted to compensate
for this. In neither of these forming methods is
there large motion of the formed material across
the die surface, so lubricant is not normally
used. The reader is referred to the Forming and
Forging in Volume 14 of the ASM Handbook for
detailed charts on bend radii and springback related to bending.
Other forming techniques employ more complex deformation processes. Deep drawing is
the foremost of these. Figure 1 (Ref 1) shows
schematically what occurs during drawing. A
round blank is held between dies over a cavity,
and a punch pushes the material into the cavity
to produce the part. If the dies pinch the blank
to only a small degree, the process is normal, or
ordinary, drawing. If the dies significantly
restrain the periphery of the blank so it cannot
move, stretch forming occurs. Material properties determine whether a material is most successfully formed by stretch forming or drawing.
174 / Stainless Steels for Design Engineers
Fig. 1 Deep drawing schematic. Source: Ref 1
Stretch Forming. In stretch forming, the
material is constrained from moving wholly
into the die. Thus, the section that enters the die
is stretched more and must become thinner.
Whether the material becomes so thin locally
that it fails is governed by its work-hardening
rate. If it work hardens faster than it becomes
thinner, the strain is distributed, and local failure is prevented. Austenitic materials have the
face-centered cubic (fcc) crystal structure with
many slip systems and low stacking fault energies. This means that they can generate many
complex arrays of tangled dislocations, which
cause strain hardening. They can also transform
during stretching to the much harder martensite,
for which the deformation is greatest, again
redistributing the deformation away from the
potentially thinning area. This makes austenitic
stainless steels particularly suitable for stretch
forming.
Deep Drawing. Deep drawing, referred to
simply as drawing, without stretching requires a
different material characteristic. For drawing, a
low work-hardening rate is desirable so that the
material can compress in the circumferential
direction while elongating in the radial direction.
Obviously, a high ability to elongate is always
useful regardless of any other characteristic.
But, ferritic material has one other advantage:
Body-centered cubic (bcc) alloys have more
slip systems than fcc alloys. When bcc alloys
are rolled to become flat stock, they may retain
a preferred crystallographic orientation, called
texture, as a result of the deformation. This nonrandom crystal structure can cause the material
to have higher strength in the through-thickness
direction. This directional variation in properties is called anisotropy.
When a material with desirable texture is
stretched, it flows in the stretching direction and
contracts laterally at lower stresses than are
required to initiate plastic flow in the throughthickness direction. As long as the work-hardening rate keeps the flow stress below the
through-thickness yield strength, there will be
no thinning. The geometry of deep drawing
with constraint fits such materials’ capabilities.
This is why carbon steels and ferritic stainless
steels deep draw well. If the material were constrained from contracting while being stretched,
the tensile strength would be exceeded before
through-thickness flow occurred and the material would fracture with little deformation. So,
these materials cannot be stretch formed. Even
without stretching, some hold-down pressure is
required to prevent wrinkling of the blank
before it is pulled into the die. This is more pronounced for thinner blanks and for material
with higher work-hardening rates. In many
forming situations, adjustments to the drawing
process (i.e., hold-down force adjustment, draw
Chapter 14: Forming / 175
bead contour, blank size, lubrication, die radii,
etc.) may be more important than material properties in determining whether the desired part
can be made successfully.
The material properties that are important to
formability are ductility, as measured by tensile
test elongation; the work hardening rate, which
is the instantaneous slope of the true stress, true
strain curve and is called n; and the anisotropy.
The measure of anisotropy is the Lankford ratio
(Ref 2), expressed as:
r + 2r45 + r90
R= 0
4
(Eq 1)
where R is the average strain ratio, r0 is the strain
ratio in the longitudinal direction, r45 is the strain
ratio measured at 45° to the rolling direction (of
the sheet metal-forming operation), and r90 is the
strain ratio in the transverse direction. R determines the average depth (that is, the wall height)
of the deepest draw possible. When this expression equals 1, then a material may be considered
isotropic, that is, the material properties are the
same for all crystallographic orientations. As a
first approximation, the Lankford ratio equals
the ratio of the lateral strain to the through-thickness strain during the tensile test of a sheet specimen. As the value increases from 1, the drawability increases because the material tends to
maintain a constant thickness while changing
shape from a flat blank to a cup shape. The ability to be deep drawn is measured by the limiting
drawing ratio (LDR), the ratio of the diameter of
a disc to that of the deepest cylinder into which
it can be drawn.
The ferritic stainless steels in sheet form have
LDRs of around 2.2 compared to 2.0 for 304. For
austenitic steel, the ratio is about 1.0, while for
flat-rolled carbon steel and ferritic stainless steel,
it can be greater than 2.0, but values between 1.5
Table 1
Steel
YUS 190
YUS 436S
(B-added)
YUS 436S
(B-free)
YUS 4O9D
SUS 430
SUS304
Source: Ref 2
and 2.0 are more common. Figure 2 shows the
LDR as a function of the Lankford ratio. It
should noted that the very ductile 304 fares no
better than carbon steel because of the advantage
of a beneficial anisotropy, which the ferrous bcc
structure has. Indeed, the best-performing deepdrawing stainless steels are low interstitial ferritic
steels with boron added. A number of stainless
steels are compared in Table 1 (Ref 4).
It shows that when comparing different types
of materials, some tests are not good predictors
of deep-drawing performance. Optimizing the
material/drawn component combination is far
from simple, especially when other considerations, such as cost and material performance in
service must be factored in. End users are
encouraged to deal directly with the producing
mill early in the design stage of any new highproduction, deep-drawn component. The producing mills, while not necessarily exhaustive
sources of information, are certainly reservoirs
of knowledge of current practice.
The most widely used summary of a material’s formability is contained in its forming limit
diagram (FLD), developed by Keeler and Backofen (Ref 5). This diagram shows the locus of
failure under varying strain states. Figure 3
shows a comparison of the FLDs for austenitic
stainless steel and carbon steel. The FLD tells
Fig. 2 Limiting drawing ratio variations with Lankford ratio.
Source: Ref 3
Deep-drawing materials comparison
0.2% proof
stress,
N/mm1
Tensile
strength,
N/mm2
Elongation,
%
Hardness
HV
n value
ro
r45
r90
r̄
Erichsen
value,
mm
Conical cup
value,
mm
343
275
497
459
33.8
34.8
173
135
0.20
0.21
1.60
1.67
1.47
1.63
2.10
2.12
1.66
1.76
9.5
9.8
26.7
26.9
284
483
34.5
137
0.22
1.49
1.90
2.01
1.83
9.8
26.8
239
308
281
424
472
705
37.2
31.8
64.0
116
159
172
0.24
0.21
0.44
1.51
0.94
0.91
1.77
0.92
1.19
2.11
l.50
0.83
1.79
1.07
1.03
9.7
8.9
12.5
26.7
28.5
27.0
Lankford value
176 / Stainless Steels for Design Engineers
Fig. 4 Forming limit diagrams for categories of stainless
Fig. 3 Forming limit diagram of carbon steel compared to
austenitic stainless steel. Source: Ref 3
the point of failure for a given sheet material
with a given thermomechanical history over a
full range of combined strain states. These diagrams are generated by examining circle grids
printed on material that is deformed to failure in
various modes. The single most important value
on the curve is the intersection of the curve with
the major strain axis at zero minor strain. This
can be used as an index of formability and is the
elongation possible for plane strain conditions.
This value within a given class of materials is
proportional to the strain-hardening exponent
because a higher work-hardening rate causes
higher localized resistance to thinning, which is
the precursor of failure.
There is much more variety within the families of stainless steel than within carbon steel.
Figure 4 shows generalized FLDs for austenitic,
ferritic, high-strength ferritic, high-strength
austenitic, and duplex stainless steels.
Deep drawing of components is seen as a way
to obtain near-net shape. Since tooling is costly, it
is necessarily a high-volume application. Quite
often, designers push component design to the
limit of a material’s ability to be formed. There
are various drivers that cause this. One is to eliminate extra operations or components by consolidating them into one more complex deep-drawn
part. Another is to use the least-expensive alloy.
In some industries, such as the household appliance industry, as many components as possible
are deep drawn from ferritic stainless alloys and
the more costly austenitic alloys are used only
when the part cannot be made from a ferritic.
There has been much research to develop ferritic stainless alloys with improved formability.
This has been accomplished by reducing the
steels. A, austenitic stainless steel; F, ferritic stainless
steel; HAS, high-strength austenitic stainless steel; HSF, highstrength ferritic stainless steel; FA(50), ferritic-austenitic stainless
steel. Source: Ref 3
total interstitial content (i.e. carbon plus nitrogen) and by thermomechanical working to give
a fine-grained, fully recrystallized, yet beneficially anisotropic, microstructure. Figure 5
shows how the FLD of an enhanced 409 ferritic
stainless steel, 409 Ultra Form, compares to the
already highly evolved 409.
For austenitic stainless steel, maximum
drawability is obtained by low work-hardening
rates coupled with maximum elongation, as
exemplified by 305 and high-nickel 304. This
comes with a cost penalty as the easiest way to
improve formability is to increase the nickel
level, although replacing the expensive nickel
with copper or manganese, as in 204Cu, has
been shown to be effective. Many austenitic
components are made by stretched deep drawing. In this case, the preferable alloys are the
leaner austenitics, 201 and 301. These alloys
form martensite more rapidly than do 304 and
305 during stretching. Martensite has a 4%
greater volume than the austenite from which it
forms and a much greater strength. This gives
201 and 301 the ability to redistribute deformation from thinning areas elsewhere and stretch
extensively, making them an optimal material
for objects such as sinks.
The specific alloy composition is often finely
tuned for a given part and tooling design, and
small deviati ons can dramatically increase
breakage rates. Even such minor process
changes as blank temperature variations due to
ambient temperature can alter work-hardening
rates enough to cause breakage problems. Certainly, this can occur when designs push the
Chapter 14: Forming / 177
Fig. 5 Optimized 409 for forming versus normal 409. Source: Ref 5
envelope of a material’s capability, but cost
pressures generally drive designers to this
extreme.
On very severe forming, intermediate
annealing may be required to either enhance
ductility or reduce required pressure. Stainless
steel in the as-drawn condition will have residual stress and may have sufficient hardness to
be susceptible to delayed failure if placed in a
corrosive environment. Bright-annealed alloys
with high martensite levels from forming can
fail by hydrogen embrittlement with just the
residual hydrogen from annealing. Therefore,
the use of bright-annealed lean alloys such as
301 is not recommended.
A last key variable in which practice and material interact is strain rate. Ferritic steels flow
more easily at lower strain rates and are thus
more formable. Austenitic steels experience the
opposite effect if they are susceptible to martensite formation. Adiabatic heating can retard the
martensitic transformation and reduce the workhardening rate, changing their forming characteristics. When drawing, this is good, but for
stretching it may not be.
Hydroforming, a variation on deep drawing
in which hydrostatic pressure forces a blank
into the die cavity, can improve the degree to
which stainless steels can be deep drawn. The
hydroforming process avoids friction between
178 / Stainless Steels for Design Engineers
the blank and the tool. Deformation is spread
more evenly across the blank, and the material
forms close to its theoretical best. Productivity
using this technique is relatively low, so its use
is justified mainly when it is not otherwise possible to make a certain design in one drawn
component. An example of this may be the production of a complex exhaust manifold that requires a higher-alloyed ferritic stainless with
relatively low formability.
Besides failure by breakage, there are other
less-severe defects found on deeply drawn
parts. Austenitic steels can develop a surface
condition known as orange peel, the result of
slip planes within a grain disrupting the surface. Orange peel is prevented by keeping
grain size fine so that the surface relief is too
small to be seen. Austenitic stainless steels can
also develop anisotropy, which while less
severe than ferritic steels, can cause “earing,”
in which round blanks deform nonuniformly in
a four-, six-, or eight-fold symmetry, causing
excess ear-shaped material to extend beyond
the intended dimensions of the component.
Material is wasted because larger blanks have
to be used. The steel producer can minimize
the phenomenon by keeping cold roll reductions above about 60%. One measure of the
earing tendency is derived from the Lankford r
measurements:
r + r + 2 r45
Δr = 0 90
2
(Eq 2)
The left side of Eq 2, ⌬r, is a measure of the
variation of plastic strain ratio r with direction
in the plane of a sheet. Values of ⌬r near zero
generally indicate minimal tendency toward
earing, while ⌬ values significantly above or
below zero indicate increased tendency toward
earing. A combination of a high R value from
Eq 1 and a low ⌬r value provides optimal
drawability.
It should be noted that deformation always
produces some surface relief, so highly reflective surfaces become spectrally diffuse, or
cloudy, after plastic deformation. This has been
an issue for items such as automotive bright
trim. Mechanical buffing can restore the luster,
but the time and expense of buffing increase
dramatically if orange peel or roping (a similar
surface defect) is excessive.
Duplex stainless steel flat products exhibit
significant in-plane anisotropy that can cause
forming difficulties and lack of compliance with
mechanical property requirements. They also
have significant rolling anisotropy, which causes
the yield strength transverse to the rolling direction to be consistently higher than it is in the
rolling direction. This behavior is contrary to the
general behavior of single-phase alloys. The difference in yield strength is sufficient, reportedly
up to 15% lower in the longitudinal direction
than in the transverse direction required for tensile tests (Ref 7), that it is both serious design
and forming considerations.
Ferritics undergo a more specialized surface
relief because their anisotropic grain structure
can yield in a more concerted fashion and give
even greater surface relief, called ridging and
roping. This is minimized by refining grain
size, achieving full recrystallization versus just
recovery during annealing, and temper passing
(i.e., elongation of about 1% on a cold-rolling
mill) to suppress the yield point phenomenon
that is characteristic of ferrous bcc materials.
End users should always make their use of the
material known to the producing mill so that
the correct thermal processing path can be
employed for the manufacturing process the
material will undergo.
One of the most important material considerations for deep drawing is surface finish. Flatrolled stainless should be fully annealed and
pickled so that the surface holds lubricant well
and yields as readily as possible. Temper passing will drastically reduce the drawability of
stainless by smoothing the surface (rather than
increasing the yield strength). Temper passing
with roughened rolls does not significantly
harm drawability. The surface finishes that are
produced by temper rolls with special finishes,
such as Koolline, retain lubricant well and can
be drawn with minimal distortion.
Tooling for stainless must be strong and wear
resistant. Traditional tooling materials are D2
tool steel and high-strength aluminum bronze.
D2 tool steel must be hardened to HRC 60 to
62 and must have smooth surfaces. The use of
powder metal techniques to produce tool steels
for dies has permitted much higher volume fractions of ultrahard microconstituents such as
vanadium carbide to be introduced, thereby
vastly improving wear resistance without harming toughness or even raising overall hardness.
The benefits of cast aluminum bronzes are low
friction, high thermal conductivity, and low tendency to gall. They are preferred when finished
part surface appearance is more important than
Chapter 14: Forming / 179
Table 2
Suitability of various lubricants for use in forming of stainless steel
Lubricant
Fatty oils and
blends(a)
Soap-fat
pastes(b)
Wax-base
pastes(b)
Heavy-duty
emulsions(c)
Dry film (wax
or soap
plus borax)
Pigmented
pastes(b)(d)
Sulfurized or
sulfochlorinated
oils(e)
Chlorinated oils or
waxes(f) highviscosity types(g)
Chlorinated oils or
waxes(f) lowviscosity types(j)
Graphite or
molybdenum
disulfide(k)
Blanking
and
piercing
Pressbrake
forming
Press
forming
Multiple-slide
forming
Deep
drawing
Spinning
Drop
hammer
forming
Contour
roll
forming
Embossing
C
B
C
A
C
A
C
B
B
NR
NR
C
A
B
B
C
B
C
B
B
B
A
B
B
C
B
A
B
NR
B
A
B
B
NR
A
B
B
B
B
NR
B
A
B
NR
A
B
NR
A
B
A
C
NR
NR
NR
A
A
B+
A
C
NR
A
B
A
A(h)
NR
A
NR
A
NR
A(i)
A
NR
B+
A
A
A
B
NR
A(i)
A
A
NR
(l)
(l)
NR
(l)
NR
(l)
NR
NR
A, excellent; B, good; C, acceptable; NR, not recommended; SUS, Say-bolt universal second. Ratings consider effectiveness, cleanliness, ease of removal, and other
suitability factors. (a) Vegetable or animal types; mineral oil is used for blending. (b) May be diluted with water. (c) Water emulsions of soluble oils; contain a high concentration of extreme pressure (EP) sulfur or chlorine compounds. (d) Chalk (whiting) is most common pigment; others sometimes used. (e) EP types; may contain
some mineral or fatty oil. (f) EP chlorinated mineral oils or waxes; may contain emulsifiers for ease of removal in water-base cleaners. (g) Viscosity of 4,000 to 20,000
SUS. (h) For heavy plate. (i) For cold forming only. (j) Viscosity (200 to 1000 SUS) is influenced by base oil or wax, degree of chlorination, and additions of mineral oil.
(k) Solid lubricant applied from dispersions in oil, solvent, or water. (l) For hot forming applications only.
absolute die life and forming pressures are
moderate.
Lubricants for stainless steel forming must be
able to prevent metal-to-metal contact under
higher pressures than those seen with carbon
steel. The ASM Metals Handbook, Desk Edition, lists common lubricants as shown in Table
2 (Ref 1). Not listed in the table are the newer
thermoplastic acrylic polymers that, when applied to the surface at a density of around 1 g/m2
(0.004 oz/ft2), provide a dry film with lubricating properties surpassing any of those listed in
Table 2.
Stainless Long Products
Cold heading, one of the most important
forming operations conducted on stainless long
products, is a forming process that increases the
cross-sectional area of a room temperature
blank at one or more points along its length.
Cold heading is typically a high-speed process
in which the blank is progressively moved
through a multistation machine. The process is
widely used to produce a variety of small- and
medium-sized hardware items, such as screws,
bolts, nuts, rivets, and specialized fasteners.
As with flat-rolled forming operations, the
primary difference between carbon/alloy steels
and stainless steels comes from the higher yield
strength and higher work-hardening rates of
stainless. Anisotropy is not a significant consideration for long products. The ferritic stainless
steels are the most easily cold headed. The use
of the most formable stainless alloys for flat
products, the stabilized ferritic alloys, is limited
for long products because of the severe lack of
toughness these alloys show for cross sections
greater than about 2 mm (0.08 in.). But, the
nonstabilized ferritics, the martensitic, precipitation-hardenable (PH), austenitic, and duplex
grades are all cold formable.
In cold-heading terminology, the maximum
possible deformation an alloy can tolerate is expressed in terms of the length of long product
exposed beyond the die that can be successfully
forged into the upset. This is measured in the
number of diameters of initial stock. So, an optimal ferritic such as 430 can tolerate upsets up
to about 2.25 diameters, while a very low workhardening austenitic, such as 384, can tolerate
3.0. The martensitic, PH, and richer duplex
180 / Stainless Steels for Design Engineers
their work hardening is not instantly removed
by annealing. For example, the initial pressure
requirements for a given deformation range for
304 can be three times what is required to deform carbon steel and up to five times for higher
molybdenum alloys. For greater deformation,
this multiple increases (see Fig. 6a and b).
Austenitic stainless steel loses ductility when
heated above 1280 °C (2335 (F) because of low
melting phases in the grain boundaries. As
much as possible, all alloys should be hot
worked in a single phase field of the phase diagram to avoid mixtures of ferrite and austenite
since the great difference in their great strengths
can cause failures. Duplex stainless steels and
other alloys (e.g., 17-7 PH) that have high levels of ferrite in austenite or austenite in ferrite at
the hot-working temperature exhibit reduced
hot ductility compared to either fully austenitic
or fully ferritic stainless steels and are more difficult to hot work.
REFERENCES
Fig. 6 Forces required for hot working. Source: Ref 1
alloys are in the 1.5 to 2.0 range of formability.
The lean duplex, when they find their way into
wider use, should resemble 430 with 2.25 diameters maximum.
Hot forming of stainless steel is done as an integral part of their production; therefore, all
stainless alloys can be forged. The main issue is
that the high hot strength of stainless requires
much more force than would be required for
carbon steels. Martensitic stainless steels require 10 to 100% more force than 4340 alloy
steel, while the austenitics require much more
because of their high hot hardness and because
1. Forming of Sheet, Strip, and Plate: Deep
Drawing, in Metals Handbook, desk ed.,
J.R. Davis, Ed., ASM International, 1998, p
782–829
2. U.F. Kocks, C. Tomé, H.-R. Wenk, Ed., Texture and Anisotropy, Cambridge University
Press, Cambridge, UK, 1998
3. E. Schedin, “Forming Stainless Steel,”
ACOM Technical Paper, www.outokumpu
.com
4. H. Sumitomo and T. Tanoue, Nippon Steel
Technical Report 71, October 1996
5. S.P. Keeler and W.A.A. Backofen, ASM
Trans. Q, Vol. 56 (No. 163), 1963, p 25–48
6. “409 Ultra Form Stainless Steel,” Product
Data Bulletin, www.AKSteel.com
7. R. Cordewener et al., “Duplex Stainless
Steels,” Paper 109, TWI conference, Glasgow, 1994
SELECTED REFERENCE
• ASM Handbook, Vol. 14, Forming and
Forging, ASM International, 1988
Stainless Steels for Design Engineers
Michael F. McGuire, p 181-191
DOI: 10.1361/ssde2008p181
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
CHAPTER 15
Machining
Summary
MACHINING STAINLESS STEELS is a
complex operation. Not only does a shop need
the correct supporting equipment and supplies,
a better understanding of the metal itself is advantageous. Technology in the production of a
more machinable stainless steel is advancing.
The incorporation of complex oxides has led to
the development of materials that allow higher
machining speeds and increased productivities,
both of which are reducing machining costs and
keeping shops competitive.
Introduction
Stainless steel forgings, castings, plate, and
long products all are frequently machined. This
fundamentally involves the removal of a layer
of material from the workpiece with a cutting
tool one or multiple times until a finished or
semifinished part is produced. Machining, in itself, is a complex topic with many variables.
Rather than attempt to understand all aspects of
machining, it is helpful to consider a material’s
machinability, that is, its ability to be machined
and the factors that affect its ability to be machined. In Fig. 1, a macroview shows how the
machinability of a material is influenced by the
interaction of humans, machine, methods, material, and management. Some of the variables
can affect the appearance of the material, while
others affect the performance of the piece, making machining an art as well as a science. Optimum machinability is obtained when each of
these sectors come together, providing the best
possible conditions for efficient machining. Any
change in one of these sectors can change the
behavior or efficiency of a machining job.
From a more focused viewpoint, the machinability of a material is further described by:
1. Consistency: Does the material machinability stay the same when bundles are changed?
2. Tool life/wear: How long does the tool last
in the machining operation? This could be
minutes, hours, shifts, or days.
3. Productivity: How many parts were made
in an hour, shift, or day?
4. Cost per part: What is the cost of the final
geometry?
5. Cycle time: How fast can a part be completed?
6. Surface finish: How smooth or shiny is the
part?
7. Chip control: Are the chips manageable?
8. Maximum cutting speed: How fast can the
part be cut without affecting tool life?
9. Maintaining tolerances: How long can the
machining operation continue before adjustments are made?
10. Minimal operator intervention: Does the
operator need to constantly adjust setup?
Fig. 1 The 5 M’s of machinability
182 / Stainless Steels for Design Engineers
This list is somewhat empirical or job related, but it provides guidelines for defining a
job since cutting conditions can be very different for each material and part. For example, if
the surface finish of a part is very important, it
may be necessary for chip control and tool life
to be sacrificed. Clearly, machining involves
much more than simply cutting a piece of
metal.
Machining is a very empirically mature subject. The recommended feed rate, depth of cut,
tool material, and cutting fluid for a given material/material condition (thermomechanical history) can be found in readily available published tables. Books such as the ASM
Handbooks; Machinery’s Handbook, published
by Industrial Press; Marks’ Standard Handbook for Mechanical Engineers, published by
McGraw-Hill Book Company; or the Machining Data Handbook, 3rd edition, by the
Machinability Data center at the Institute of Advanced Manufacturing Services (IAMS; formerly known as Metcut Research Associates
Inc.) in Cincinnati, OH; include much of the
data used in industry today. Material manufacturers are also a source of valuable machining
data. A typical guide from ASM is shown as
Table 1 (Ref 1).
Rather than simply reproducing data, the
focus of this chapter is the metallurgical factors
governing the machinability of stainless steels.
Most of the information regards machining
stainless bar products; however, many of the
concepts could be applied to forgings as well as
castings.
Physical and Mechanical Properties
The machinability of stainless steels is very
difficult to characterize in definitive terms because of the broad nature of these materials. A
ferritic stainless steel, such as type 430, will
machine very differently from the martensitic.
In some sense, this is like comparing brass to
carbon steel. Both type 410 and type 430 are
stainless steels, but the chemistry and structural
differences create diversity in machining characteristics.
The machinability of stainless steels can be
thought of as a function of the steel’s chemistry,
cleanliness, structure, processing history, and
the cross-section size of the stock, with no one
factor more important than another:
Machinability of
stainless steels =
ƒ(chemistry, cleanlinesss,
structure, processing,
cross section)
Each variable contributes uniquely to machinability. Machine shops and users generally have
very little influence on these material variables.
Because no two mills are exactly identical, there
will be differences in machinability of a steel
grade provided by different mill suppliers. However, having an understanding of how these variables contribute to machinability is invaluable.
Armed with an understanding of the material and
how it is made, one can determine the tooling,
coolant, and setup of the machining job.
Let us take a closer look at these variables.
Chemistry
The role of chemistry is to define not only
the different grades of stainless steel (ferritic,
martensitic, etc.), but also how the grade
is chemically balanced within the specific
grade; for example, the amount of carbon in a
martensitic stainless can change tool wear characteristics, or a change in nickel content within
specification limits can alter the stringiness of a
chip. Combined, both will be the basis the material’s machinability.
Each of the elements used to produce stainless steels will contribute some general machining attributes. The effects of the elements as described next are general, and slight deviations
may be encountered depending on the stainless
grade. However, for the more common stainless
grades used today, these effects of these alloying elements are fairly accurate.
Iron is the base element in a stainless steel. It
is a soft, gummy material that has high workhardening characteristics. Iron is characterized
by surface finishes that are difficult to obtain
and chips that are stringy, and it has a high tendency toward tool built-up edge (BUE).
Chromium strengthens and reduces ductility
of stainless steel. Machine and tool setup require more rigidity. Chromium allows chips to
begin breaking.
Carbon content increases strengthen stainless
steels and promote carbide formation. Low carbon levels, typical in ferritic stainless steels, do
not help machinability much. Increasing
amounts of carbon to greater than 0.08% will aid
in chip breakability and reduced BUE in these
grades. However, as carbon content increases,
Chapter 15: Machining / 183
Table 1
Machining setup recommendations for turning wrought stainless steels
Uncoated
Coated
High-speed steel tool
Condition
Depth of
cut(a), in.
Speed,
fpm
Feed,
ipr
Tool
Tool
Tool material
Feed, material Speed, Feed, material
Brazed Indexable ipr
AISI
fpm
ipr
grade
grade
Ferritic steels
405, 409, 429, 135-185 Annealed
430, 434,
436, 442,
446(c)
0.040
0.150
0.300
0.625
150(235)
120(190)
95(150)
75(115)
0.007
0.015
0.020
0.030
M2, M3
M2, M3
M2, M3
M2, M3
0.040
0.150
0.300
0.625
0.040
0.150
0.300
0.625
95
75
60
45
80
65
50
40
0.040
0.150
0.300
0.625
0.040
0.150
0.300
0.625
0.040
0.150
0.300
0.040
0.150
0.300
155
125
100
80
145
115
90
70
95
75
60
65
50
40
Material
Hardness,
HB
Austenitic and duplex steels
201, 202, 301, 135-185 Annealed
302, 302B,
304, 304L,
305, 308,
309, 309S, 225-275 Cold drawn
or
310, 310S,
duplex
314, 316,
annealed
316L, 317,
321, 330,
347, 348,
384, 385(c)
2205, 2507
295-310 Annealed
Martensitic and PH steels
403, 410, 420, 135-175 Annealed
422, 501,
502(c)
175-225 Annealed
275-325 Quenched
and
tempered
375-425 Quenched
and
tempered
Speed, fpm
575
450
350
275
650
500
400
310
0.007
0.015
0.030
0.040
C-7
C-6
C-6
C-6
850
650
525
...
0.007
0.015
0.020
...
CC-7
CC-6
CC-6
...
0.007 M2, M3
0.015 M2, M3
0.020 M2, M3
0.030 M2, M3
0.007 T15, M42(b)
0.015 T15, M42(b)
0.020 T15, M42(b)
0.030 T15, M42(b)
325
300
225
175
300
250
290
140
375
325
250
200
325
275
215
165
0.007
0.015
0.020
0.030
0.007
0.015
0.020
0.030
C-3
C-3
C-2
C-2
C-3
C-3
C-2
C-2
500
425
325
...
425
350
275
...
0.007
0.015
0.020
...
0.007
0.015
0.020
...
CC-3
CC-3
CC-2
...
CC-3
CC-3
CC-2
...
0.007 M2, M3
0.015 M2, M3
0.020 M2, M3
0.030 M2, M3
0.007 M2, M3
0.015 M2, M3
0.020 M2, M3
0.030 M2, M3
0.007 T15, M42(b)
0.015 T15, M42(b)
0.020 T15, M42(b)
0.007 T15, M42(b)
0.015 T15, M42(b)
0.020 T15, M42(b)
475
400
320
240
460
385
300
235
360
280
225
290
225
180
620
480
380
300
570
450
350
265
465
360
280
320
250
200
0.007
0.015
0.030
0.040
0.007
0.015
0.030
0.040
0.007
0.015
0.020
0.007
0.015
0.020
C-7
C-6
C-6
C-6
C-7
C-6
C-6
C-6
C-7
C-6
C-6
C-7
C-6
C-6
800
625
500
...
850
550
450
...
700
450
375
475
300
250
0.007
0.015
0.020
...
0.007
0.015
0.020
...
0.007
0.015
0.020
0.007
0.015
0.020
CC-7
CC-6
CC-6
...
CC-7
CC-6
CC-6
...
CC-7
CC-6
CC-6
CC-7
CC-6
CC-6
PH, precipitation-hardenable. Source: Ref 1
(a) Caution: check horsepower requirements on heavier depths of cut. (b) Any premium high-speed steel (T15, M33, M41–M47). (c) Free machining versions.
the amount of carbide increases, the structure
changes to martensitic, and the wear on tools
increases.
Nickel increases the toughness and ductility
of stainless and reduces the work hardening
rate. Nickel also increases elevated temperature
mechanical properties. This causes chips to be
more difficult to break. Nickel will have a tendency toward increased BUE; however, better
tool life will generally result.
Sulfur reduces mechanical and corrosion
properties and can be a cause of hot cracking in
the resulfurized grades. It is best known as a
free-machining contributor that promotes better
tool life and greater machining speeds.
Manganese is generally added to combine
with sulfur to form manganese sulfide (MnS),
which acts as a self-lubricant and improves
machinability. In high-manganese grades, such
as duplex and 200 series alloys, manganese has
the same relative effect as nickel when used in
greater amounts, as, for instance, in the 200 series stainless steels.
Molybdenum increases the strength and elevated temperature mechanical properties. This
increase in hot hardness and strength means
more energy will be needed to cut the material,
thus creating hotter cutting conditions. While
the molybdenum helps in chip breakability, it
will require more rigid setups and will reduce
tool life.
Copper improves ductility and reduces the
strain-hardening or work-hardening rate (with
the exception of participation-hardening alloys,
for which copper is used as the precipitant).
Chips can be more difficult to break, which increases the tendency of BUE and promotes better tool life.
184 / Stainless Steels for Design Engineers
Nitrogen strengthens stainless steels. It aids
in chip breakability and reduces BUE but increases tool wear.
Titanium promotes carbide formation and increases tool wear.
Niobium promotes carbide formation and increases tool wear.
The production of stainless steels is identified by industry specifications, such as AISI,
UNS, EN, JIS, etc. These specifications are all
defined with fairly broad elemental chemical
compositions. For example, an AISI 304 has a
2 wt% window for the nickel content; that is,
this grade can have a nickel level of 8 to 10%.
A type 304 with 8% nickel can have different
machinability characteristics from a type 304
with 10% nickel. This 2% difference alters chip
morphology and surface finish capability. Since
today’s mill technology can meet very tight elemental targets within the grade specification,
how the mill balances the grade’s chemistry
will provide the foundation of its machinability
characteristics.
Cleanliness
The cleanliness of steel is determined by the
amount and type of inclusions it contains. Vacuum and argon oxygen decarburization (AOD)
melting and refining along with proper steelmaking techniques can reduce the inclusions to
negligible levels. It is beneficial to machinability to avoid hard inclusions. However, certain
inclusions are plastic and act as solid-state lubricants and chip breakers and prevent adhesion
of the material to the tool. The beneficial effect
of controlled inclusions is discussed in this
chapter.
Structure
Material structure consists of both the phases
that are present and the microstructure of those
phases. Each type of stainless steel belongs to a
larger family, which is characterized by a single
predominant phase or a combination of two.
These are ferritic, austenitic, martensitic, precipitation hardening, and duplex (see the chapters on stainless steels, Chapters 6 to 10, in this
Volume). Their machining characteristics are
described in the next section. The microstructure of a given alloy is independent of the grade
type and composition and is mainly influenced
by grain size. Grain size is not normally specified or reported on certifications; however, mills
measure and control it to varying degrees. The
material’s grain size results from the thermal
and mechanical history during manufacturing
and from the mill’s equipment capability and
practices.
The grain size of a particular product can dramatically change its machinability. It is entirely
possible for the grain size difference between
two lots of material to be large enough to prevent both lots from being effectively machined
with the same setup, requiring adjustments in
the machining setup to remedy the situation.
Finer grain sizes strengthen the stainless steel,
cause hotter cutting conditions, and have a
higher tendency of BUE. On the brighter side,
finer grain sizes yield better surface finishes and
smoother roll thread crests.
Process
The type of equipment used by the stainless
manufacturer, the manufacturing sequence, and
the practices employed by the mill can affect
machinability as well as mechanical properties,
but more important, processing affects how
consistently the material can be machined. The
melt type, hot rolling parameters, thermal treatments, cold-finishing parameters, and sequence
of these operations can affect how consistently
a material machines. Many times, the culprit is
equipment operational procedures or practices
that can vary one day to the next. Equipment
types can also play a role in machinability. For
example, machinability can vary when the
same-size material is drawn across two different
draw benches using different pulling mechanisms and two different straightening mechanisms. Whether the material is continuously annealed or batch annealed can cause different
strain distributions across the material cross
section as well as material strength differences.
Various annealing lines vary in time/temperature profiles and therefore result in different
grain size and mechanical properties.
With all this in mind, manufacturing consistency can be a great asset in machinability. A
machine shop can adjust when a material is consistently bad, but it is very difficult when one lot
is easy to machine followed by a lot that is
tough to machine, while a third bundle performs
differently from the first two. Mills that promote
machining consistency pride themselves by
practicing manufacturing consistency. Tolerance variation will be tighter and machining
costs will be lower with their products.
Chapter 15: Machining / 185
Cross-Section Size
Mill processing equipment dictates a manufacturing route based on size. Smaller diameters
are cold drawn, while larger diameters are
straightened/cut/turned, yielding a softer product. This can have an impact on machining performance despite all other factors being the
same. Cold finishing of stainless steels can be
accomplished via a couple of general manufacturing routes.
The first is by cold drawing to bar, and the second is simply a straightening-turning operation.
The mechanical properties of the straightenedturned bars will be softer than the bars made by
cold drawing. Typically, sizes greater than 1 in.
(25 mm) are annealed/turned and straightened,
with virtually no strain in the product.
Machinability of the Stainless
Steel Families
Comparing the machinability of stainless
steels with other materials such as carbon steels,
brass, or aluminum, there are some striking differences. In general, stainless steels have:
1. Low thermal conductivity
2. High work-hardening rates
3. High tensile strengths
4. High toughness
5. High ductility
6. Large spreads between the yield and tensile
strengths
will tend be stringy but can be broken through
aggressive chip breaking, and surface finishes
will be somewhat of a challenge. Ferritics are
the easiest of the stainless steels to machine, do
not require much horsepower, have a low workhardening rate and better tool wear, and will
generally have higher speed and feed capabilities than other stainless families.
Martensitic
Martensitic stainless are also very basic
straight chromium stainless steels, 400 series
stainless grades, and are similar to the ferritic
grades. The difference is that the martensitic
grades have much higher carbon levels, which
further strengthen the materials and allow these
materials to be hardenable by heat treatment.
These grades will have higher carbide levels,
which will lead to higher tool wear. This is especially true if the material is being machined in
the hardened condition. The higher strengths
will require more horsepower to cut and will
need more rigid setup than ferritic steels. The
work-hardening rate of martensitic stainless is
lower than for ferritic stainless. Martensitic
stainless also has a small yield-to-tensile ratio,
making chips easier to break.
Austenitic
Each stainless steel family (ferritic, martensitic, etc.) brings its own general set of machining rules. This is mainly due to the chemistry of
these families and its resultant effect on the
physical and mechanical properties. A general
description of the machining behavior is provided next. One must keep in mind that these are
general characteristics. Further alloying of these
families, such as with a sulfur addition, can result in a radical change in machining behavior.
The austenitic grades, the 300 series stainless
grades, are more difficult to machine than the
ferritic and martensitic families. Austenitic
stainless steels are more highly alloyed and are
more prone to higher work-hardening rates. This
leads to the need for higher horsepower and
more rigid setups. These grades are very prone
to BUE and hence are prone to poorer surface
finishes and tend to tear. The yield-to-tensile ratios of austenitic stainless steel is very large,
making chips hard to break. Chips in this family
of alloys tend to be long and stringy. The higher
strength and higher ductility of these grades also
tend to increase cutting temperatures, necessitating tooling with higher heat resistance.
Ferritic
Precipitation Hardening
Ferritic stainless steels are the most basic
stainless steels and are part of the 400 series
grades. Their basic chemical composition is
iron and chromium. These grades generally exhibit lower strengths, more ductility and softness, and a close yield-to-tensile ratio. These
grades will have a high tendency to BUE, chips
Precipitation hardening stainless steels are
characterized by higher strength and toughness.
The solution-annealed hardness of AISI 630, for
instance, is HRC 36 versus HRC 23 for a 304.
Higher horsepower requirements, high tendency
to BUE, higher tool wear, and difficulty breaking chips are familiar scenarios for this class of
186 / Stainless Steels for Design Engineers
stainless. Except for alloy A-286, the precipitation-hardenable (PH) grades are all martensitic
alloys and can be treated as such for machining
purposes.
Duplex
The duplex is a unique class of stainless characterized by a dual-phase structure. Duplex alloys have a structure that is roughly a 50% mix
of austenite and ferrite; thus, two hardness materials with different hardnesses coexist side by
side. The tool will alternate cutting between soft
and hard grains of the duplex structure, leading
to an automatic tendency to initiate chatter in
the cutting system. Strength levels of duplex alloys are quite a bit higher than austenitic grades.
Between the duplex structure and high-strength
levels, high horsepower is necessary, and highly
rigid setups are required. Some work at Ugitech
found that to effectively machine these grades,
highly alloyed carbide tooling with high hardness and high heat resistance, such as the
C7/C8-type carbides, should be used.
Super Stainless Steels
Super stainless steels are today’s highly specialized stainless grades. These grades, like the
duplex alloys, are being developed to increase
corrosion performance parameters to meet some
of today’s increasing performance requirements. These alloys are more highly alloyed
than the duplex materials. Strength levels are
higher and toughness is greater, driving machinability downward.
Role of Inclusions
Metallurgists have long known that the presence of a soft second phase dispersed in the matrix of a parent metal can improve its machinability. These particles provide a solid-state
lubricant between the chip and tool or a discontinuity in the material to aid in chip breaking.
The challenge to the alloy designer has been to
develop second phases that produce these beneficial effects with a minimal of degradation to
the material itself.
Lead, Selenium, Tellurium
The range of additions possible to stainless
steel is the same as for carbon steel. Lead additions probably are the best source for improving
machinability of a material. It was once a leading addition to carbon steels, but its reported environmental toxicity has diminished its role.
Many carbon mills are looking at other machining agents to replace lead. Lead, however, has
not been a large factor for stainless steels because of the extremely negative effect it has on
hot workability, always a serious consideration
in stainless steel design.
Selenium and tellurium have similar characteristics to lead as additives but also are noncompetitive due to cost, toxicity, and incompatibility with stainless.
Sulfur
It became obvious very early to metallurgists
that higher sulfur levels correlated to better
machinability, and sulfur remains the popular
additive choice. Sulfur is cost-effective as a
machinability additive and can be easily removed with modern refining methods. Sulfur is
a natural impurity and has negative effects on
mechanical and corrosion properties, discussed
separately here. The role of sulfur as a machinability agent in stainless steels is very complex
and not necessarily straightforward, but in general sulfur has been extremely beneficial in increasing the machinability of stainless steels.
Generally, as the molten metal cools to solid
form, sulfur combines with manganese to form
manganese sulfide inclusions. Manganese becomes a very important variable during this reaction. Two basic sulfide forms are found in
stainless steels—manganese sulfides and
chromium sulfides—and the form the sulfide
will take depends on the manganese content.
When manganese levels are less than 0.4%,
chromium sulfides and chromium-rich sulfides
will be present. As manganese levels reach 0.4
to 1.8%, chromium-rich manganese sulfides are
present. For manganese levels beyond 1.8%,
comparatively pure manganese sulfide will be
found. The manganese-to-sulfur ratio is also important. For highest machinability levels, a high
manganese-to-sulfur ratio is desired. However,
if corrosion resistance is desired, a low manganese level is preferred to encourage the formation of chromium sulfides or chromium-rich
manganese sulfides since these sulfide forms
have superior corrosion resistance.
Sulfides form initially as spherical inclusions
within the cast structure. Hot working, as well as
cold working, elongates these inclusions as
shown in Fig. 2. This elongation increases the
Chapter 15: Machining / 187
Fig. 2 Typical AISI 303. Source: Ref 2
Fig. 3 AISI 303 chip breaking at the sulfides. Courtesy of
Ugitech
surface area of the sulfides and will form weak
planes that mechanically weaken the steel, permitting chip breakage. The deformation caused
by the severe bending of the chip during cutting
causes the chip to break at the sulfide striations,
shown in Fig. 3. An additional benefit of these
manganese sulfide stringers is that when these
stringers touch and pass the cutting tool, a small
amount of manganese sulfide is deposited on the
tool surface, providing a lubricating layer between the chip and tool. This reduces friction,
thus reducing heat to the tool. Consequently,
machining speeds can be increased, improving
machining productivity.
Much has been written about the benefits of
sulfides in a machining operation. A metallurgical perspective shows the more complex nature
of sulfides. And, there is more to this than just
adding sulfur. The discussion of the combination of manganese and sulfur revealed that there
is a particular balance of manganese to sulfur to
achieve desired needs.
In addition to chemistry (manganese-sulfur
balance), the size and shape (relative elongation
or globular nature) of sulfides contributes to the
machinability of stainless steels. Sulfides are
defined into four categories based on morphology: type I to IV sulfides. Type I sulfides form
from the melt as large globular sulfides and are
assumed to be best for machinability. Type II
sulfides generally form as a eutectic-like distribution of finer rod-like sulfides in interdendritic
regions. Type III sulfides form as angularshaped particles. Type IV sulfides form from the
melt as plate-like sulfides in a ribbon-shaped
pattern. It has been shown that the larger, more
globular (less-elongated) type I, homogeneously dispersed sulfides are better than the
other types for enhancing machinability in both
carbon and stainless steels. However, it seems
that sulfides that are too large or too small can
be detrimental to machinability. The coarsest
type I sulfides, once thought to be best for
machinability, are difficult to attain with today’s
stainless manufacturing equipment. Ingot casting, with the slow cooling and solidification
rates, is beneficial to the creation of the coarser
type I sulfides. With the transition to continuously cast blooms, solidification rates are much
faster, effectively creating a finer type I sulfide.
The effect of sulfur on the machinability of
stainless steels is more effective in smaller
amounts than in larger amounts. Figure 4 shows
a graph of drill penetration results in varying
amounts of sulfur in an 18-9 stainless. It is easily seen that small additions of sulfur have the
greatest effect at sulfur levels <0.10%. This also
can be proven with a 304 or 316 stainless steel,
for which sulfur levels are inherently lower than
0.030%. Machinists readily can see differences
in tool wear between a 316 with 0.023% and
one with 0.028% sulfur. That small addition of
sulfur has a dramatic effect. On the opposite
side of the scale, with sulfur levels greater than
0.20%, the curve flattens out.
From a machining perspective, stainless
steels can be classified into three groups based
188 / Stainless Steels for Design Engineers
Fig. 4 Effect of sulfur on stainless machinability. Source: Ref 2
on the sulfur additions. In the first group, sulfur
content is limited to 0.010%. For these stainless
steels, corrosion resistance, weldability, or mechanical properties are more important than
machinability. Grades in this group are
316LVM and other remelted stainless grades,
many duplex and super stainless grades. The
second group has sulfur contents ranging from
0.010 to 0.030%. This group of stainless steels
still holds corrosion resistance, welding, and
mechanical properties as critical but has the advantage of increased machinability. Grades in
this group include 316/316L, 304/304L, 321,
347, 410, and 430, among others. The third
group with sulfur levels ⭓0.15% are considered
resulfurized grades. Sulfur levels are generally
0.25 to 0.35%, with some grades reaching
0.45%. The grades in the third group include
303, 420F, 430F, 430FR, 1.4570, and others. A
fourth group of grades can be included in this
classification, but the amount of sulfur is not the
grading criterion. This is where mills will rebalance chemistries and processing to enhance
machinability beyond what the first three groups
offer to create enhanced machining grades.
Product offerings include Ugitech’s Ugima and
Ugima XL (Ugima 2) materials, Sandvik’s Sanmac materials, Outokumpu’s Prodec materials,
and Carpenter’s Project 70+ materials.
Sulfides in stainless steels have a dark side,
especially when sulfur levels exceed 0.15%. It
is well documented that sulfides negatively affect corrosion resistance and mechanical properties of stainless steels. Sulfur tends to form
segregated films with low melting points during
solidification. And, since these films are low
strength, they may induce the formation of microcracks brought on by solidification shrinkage
stresses. Further processing of these materials
can induce cracking even further, leading to
poor processing yields.
Sulfur hurts corrosion resistance by locally depleting chromium from the matrix to precipitate
as a sulfide. These manganese sulfide inclusions
become exposed on the surface of the bar and become initiation sites for pitting corrosion. Passivation of the components will help, but the hole
left behind by the sulfide becomes a collection
site for contaminants. Other sulfide inclusions
that form on the grain boundaries contribute to
intergranular corrosion. Hot workability is also
hurt by sulfides. Sulfur increases the hot shortness of materials during hot-forming operations.
Manganese sulfides form stress risers within
the material, which lead to reduced mechanical
properties in notch-sensitive alloys, especially in
the transverse direction. As the amount of sulfides increases in stainless steels, a susceptibility
to longitudinal cracking can become an issue.
With these stress risers in place, any cold deformation can lead to cracking along the sulfide stringers. As a general rule, the smaller the
bar, the higher the sulfur, and the higher the
strain produce a high probability for a crack to
initiate.
Despite these deficiencies, the sulfides found
in stainless steels are very effective in improving
machinability, especially in austenitics, where
contact forces are very high. The benefits of sulfur to improve machinability outweigh the
losses due to defective parts, at least from the
machinist’s viewpoint. The question of performance of the finished component is another issue,
which concerns the end user; for the end user,
sulfur is a major negative factor for corrosion resistance. This has led to the development of alternative methods of improving machinability.
Oxides
The basic machinability-enhancing agent discussed thus far has been sulfur. The beneficial
effects of sulfur are undeniable, but the detrimental effect is equally evident. This has led
steel producers to look at other inclusion systems for a viscoplastic inclusion without the
negative effects of sulfur. Since the 1900s,
steelmakers have known that injecting calcium
into the melt converts refractory inclusions into
soft, malleable, complex oxides that act as freemachining agents with high-temperature lubricating capabilities. The oxide inclusion chemistry is based on the CaO-Al2O3-SiO2 system.
Figure 5 shows that small amounts of calcium
can greatly increase tool life. These calciumbased oxide formulations have been commer-
Chapter 15: Machining / 189
Fig. 6 Complex Ugima oxides populating the 303 matrix.
Courtesy of Ugitech
Fig. 5 Effect of calcium on machinability of 303. Source:
Ref 2
Fig. 7 X-ray examination showing manganese sulfides and Ugima oxides coating the tool surface. EDAX, energy dispersive analysis
by x-ray. Courtesy of Ugitech
cialized for stainless steels, but the machinability agent is, at this point, not standardized as
this process is difficult to reproduce consistently from heat to heat. However, Ugitech SA
(formerly Ugine Savoie and now part of
Schmolz and Bittenbach) developed a proprietary and patented process sold under the trade
name Ugima and Ugima XL (Ugima and Ugima
2 in Europe and Asia). Figure 6 shows the
Ugima oxide coexisting with sulfur in AISI.
The Ugima oxide works similarly to sulfur by
coating the cutting tool and acting as a lubricant.
Figure 7 shows EDAX (energy dispersive analysis by x-ray) spectra of the surface of a carbide
tool, proving the existence of coatings of manganese sulfide and Ugima oxide. The Ugima
oxide performs synergistically with sulfur. Since
the oxide alone has limited lubricating abilities
at slow machining speeds, the manganese sulfides in the stainless are the dominant machin-
190 / Stainless Steels for Design Engineers
ing agent. As machining speeds increase, manganese sulfide eventually loses its lubricity, and
the oxide acts as a high-temperature lubricant
that will allow faster machining speeds without
the need to use specialized tooling. Data generated at Ugitech as well as data generated by
other producers have shown that once sulfur
content reaches 0.30%, the contribution toward
machinability flattens and possibly decreases,
which seems to contradict the experience of
many machine shops. However, the combination of sulfur and the Ugima oxide extends the
machinability range for sulfur levels beyond
0.30%, as seen in Fig. 8.
Tool lubrication is only one function the complex oxide performs. Like sulfur, the Ugima
oxide is a discontinuity and will aid in chip
breakage. Figure 9 exhibits cross-sectional photomicrographs of 304L chips with a 0.025% sulfur level. As shown, the chip is shearing along
complex oxide stringers, helping the chip break.
Ugitech has seen synergistic effects with
copper additions as well. Machining tests at
Ugitech and many field experiences have shown
increases in machining performance of 30%
with the addition of 1.4 to 1.8% copper to a
303 stainless steel with the Ugima oxide. The
grade chemistry meets EN 1.4570 and is sold
under the trade name 303 Ugima UX (4570
Ugima in Europe). Further additions of copper
up to 4% have exhibited even better machining
performance.
Fig. 8 Comparison of machinability of AISI 303 at different
sulfur levels with and without the Ugima oxide.
The vertical axis, VB30/0.3, represents 0.3 mm of tool wear in
30 min.
Tooling and Coolants
The machining of stainless steels can be
much more complicated than for other materials. Machine and tool rigidity, machine power
requirements, sharp cutting tools, and optimum
lubrication type and amounts are very important. Feed rates need to be sufficiently more aggressive than with carbon steels to prevent work
hardening and glazing of the material as well as
to avoid reducing tool life.
High-Speed Tool Steels. The cutting tools
themselves are the main variable other than the
Fig. 9 Comparison of 304L chips with and without the Ugima oxide. Courtesy of Ugitech
Chapter 15: Machining / 191
material in determining the quality of the machined part. The earliest cutting tools were
those made of the high-speed steels (HSSs). The
wrought HSS tooling was very versatile, allowed resharpening of the tools many times,
and is still a preferred choice in many shops.
With time, mills perfected their ability to produce more highly alloyed tool steels to meet the
increasing demands of the machining industry.
Today, tools are being made with a powder metallurgy process, by which ingots of compacted
high-speed tool steel powders have more structural homogeneity and thus better wear and heat
resistance. Powder metallurgy techniques produce even more highly alloyed high-speed tool
steels with properties approaching those of carbide tooling, allowing better machining performance. Although high-speed tool technology
has improved, its limiting factors are heat and
wear resistance. Use of HSS tooling still limits
machining speeds and performance.
Carbides. Carbide tooling is the next generation of tool materials after HSS, although it
tends to be hard and brittle. In the past, unless it
was possible to feed hard, maintain fast speeds,
and have uninterrupted cuts, carbide tooling
was a good choice. However, carbide tool technology has come a long way in grades and
technology. Micrograin and nanograin carbides
are providing carbide tooling with increased
toughness. Single-point tooling is now able to
withstand the punishment of an interrupted cut,
and drills are able to withstand some flexing.
Generally, the carbide grades to be used when
machining stainless steels are the C5- to C8type carbides. These classes of carbides are
harder and have more heat resistance, properties
that are needed when cutting stainless. The discussion in this chapter has stated that the cutting
of stainless is more difficult and generates more
heat. Cutting materials that withstand these circumstances are needed. The C2- to C4-type carbides are not well suited for stainless steels because they do not have the heat resistance
needed. There are always exceptions. For example, it is possible to use the C2- to C4-type carbides when using the older cam-operated multispindle and Swiss machines, for which machine
speeds are limited. Tooling manufacturers such
as Kennametal, Sandvik, Iscar, etc. have their
own proprietary grade designations, but many
correspond to the C5 to C8 types.
Coatings. Tool coatings have contributed to
machining improvements. Coatings add a very
hard layer on the tool surface that will provide
additional lubricity between the tool and chip as
well as potentially providing heat resistance for
the tool. Like the carbide, there are certain coatings that are more beneficial than others. Titanium nitride (TiN) was a great coating when it
was first introduced, but further development in
this area has created other coatings that work
even better. Grades like titanium-aluminumnitride (TiAlN) and aluminum-titanium-nitride
(AlTiN) are great coatings for turning, form cutting, cutoff, grooving, drills, reamers, and
milling. These tools are usually subjected to
higher temperatures from high speeds, deep
cuts, and limited coolant. The aluminum in the
coating breaks down and combines with the
surrounding oxygen to form aluminum oxide
(Al2O3), a material with higher hardness and
better thermal resistance. When tooling will not
be subjected to much heat, such as with taps,
high wear resistance coatings such as titanium
carbonitride (TiCN) coatings are good.
Coolants. Stainless cutting is hot, has high
frictional forces, and has tendencies of the
metal to stick to the tools. Coolants need to remove this heat and provide lubrication to reduce
friction and minimize BUE. Coolants available
are petroleum based, semisynthetic, synthetic,
water soluble, and the new type: vegetable
based. All of these coolants need to be highly
fortified for use in stainless machining. For example, petroleum-based oils need high sulfur,
chlorine, and fat levels for lubricity at higher
temperatures. The other coolants need highpressure additives, again to help lubricity. It is
always a good idea to discuss the options with
the coolant supplier.
REFERENCES
1. Machining Data Recommendations, in Metals Handbook, desk ed., 2nd ed., J.R. Davis,
Ed., ASM International, 1998, p 917–950
2. T. Kosa and R.P. Ney, Sr., Machining of
Stainless Steels, in ASM Handbook, Vol
16, Machining, ASM International, 1989,
p 681–707
SELECTED REFERENCE
• www.ugitech.com
Stainless Steels for Design Engineers
Michael F. McGuire, p 193-199
DOI: 10.1361/ssde2008p193
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
CHAPTER 16
Surface Finishing
Summary
SURFACE TREATMENTS are extremely
important to the end user, and they are totally
within the end user’s control and specification.
They include the proper cleaning of stainless,
the various means of descaling after thermal
treatment, and the choice and application of surface finishes. The cost ramifications of improper
or suboptimal surface treatments are immense
because of the possibility of them compromising corrosion performance, which is a characteristic for which the end user pays dearly.
Beyond such considerations, one can also
create surfaces on the stainless that enhance its
beauty and performance. The surface can be
made reflective or matte, ground or mechanically
patterned, coated, painted, plated, or oxidized.
It can be treated chemically or electrolytically.
The surface can be altered on an atomic basis,
sometimes producing profoundly different mechanical and corrosion-resistant properties. All
of these are discussed in this chapter.
Function of Surface Treatments
Removal of Oxide Scale
Introduction
Surface finishing is usually very important for
stainless steel. The underlying economic justification for using stainless steel is that it does not
corrode if properly specified for the environment
it faces. Thus, its surface appearance remains
normally intact throughout its life. This appearance should therefore be aesthetically pleasing,
even in an industrial setting, and the surface finish should not detract from its performance.
Raw stainless surfaces resulting from rolling
and annealing operations are not considered attractive and are used only for functions in which
aesthetics are a negligible consideration. Even
so, there are surface treatments required for
stainless steel intended for such uses. The stainless surface must be freed of:
• Oxides resulting from annealing, joining, or
hot forming
• Accumulated ambient foreign material
• Applied process materials, such as forming
lubricants, fluxes, etc
• Contamination from other materials, especially iron.
Oxide scales form on stainless steel during
annealing, hot-forming, and joining operations.
Removal of this scale is important to proper corrosion resistance. This is because the chromium
in the steel oxidizes much more readily than
other elements, so the surface of the steel under
the oxide is chromium depleted and therefore
has lost possibly a significant amount of corrosion resistance. An oxide scale is quite different
from a passive film. Scales deplete chromium;
passive films enrich the surface in chromium by
selective loss of iron.
Oxide scales are arguably best removed by
pickling. Pickling is the chemical dissolution of
the oxide scale. The acids commonly used are nitric (HNO3), hydrofluoric (HF), and sulfuric
(H2SO4). HCL, which is commonly used to
pickle carbon steel, is not recommended for stainless because it locally attacks (i.e., pits) the surface. HNO3 is an oxidizing acid that by itself does
not pickle stainless. It is used in combination with
HF to modulate the attack by the strongly reducing action of the HF. This combination allows
good control of pickling rates by varying the ratio
of the two acids. H2SO4 is used alone. It is often
194 / Stainless Steels for Design Engineers
used as a preliminary pickling because of its aggressive, and less-controllable, action.
Pickling involves both scale dissolution and
metal dissolution. Thus, pickling can be considered a deliberately imposed corrosion process.
The scale formed during hot rolling may be an
embedded scale with a minimal chromiumdepleted layer beneath the scale. Annealing and
shotblasting are required before pickling so that
the acid can penetrate to the chromium-depleted
region under the scale and selectively dissolve
this layer. Nearly all the action of pickling derives from undercutting of the oxide scale and
dissolution of the chromium-depleted layer. It is
also vital that this layer be removed as it diminishes the corrosion resistance of the surface.
Simply removing the scale mechanically may
be insufficient if only the oxide is removed and
the compromised (depleted) metal surface is
left. Metal removal rates are proportional to the
chromium level. Thus, sensitized material will
undergo grain boundary attack if pickled. All
stainless has slightly greater oxidation in the
grain boundaries when it is oxidized, and this
leads to “ditching” of the grain boundaries during pickling, resulting in the matte appearance
of a pickled surface, as seen on two-dimensional
flat-rolled material.
Typical concentrations for H2SO4 pickling are
8 to 15% by weight. Bath temperatures are generally 150 °F (65 °C). The rate of pickling decreases rapidly as iron builds up in the bath, and
the bath must be replaced for efficient pickling.
The attack of H2SO4 on the base metal can be
severe, and undissolved constituents can remain
on the surface as “smut.” This smut must be
physically removed or dissolved by subsequent
HF/HNO3 pickling.
The HF/HNO3 pickling is carried out between
120 and 140 °F (50 and 60 °C). Higher temperatures cause excessive HF evaporation and can
also lead to visible emission of nitrogen oxides.
The NO2 (nitrogen dioxide) produced during
the reduction of the HNO3 (oxidation of the
metal) is visibly brown-red. The rate of NO2
formation increases at higher temperatures.
HNO3 concentrations are normally from 10 to
25%. HF concentrations can vary from 1% for
light scales to 8% for maximum aggressiveness
and difficult-to-pickle alloys.
The HF provides:
• a complexing agent for iron and chromium
• destabilization of the passive film
• stabilization of the redox potential
The HNO3 provides:
•
•
•
•
•
a source of H+ ions
stabilization of the passive film
elevation of the redox potential
an oxidizing agent of the base metal
a dissolving agent for the scale
The nitrates carried off from this pickling
process are an environmental problem. It has
been shown that the HNO3 can be eliminated
and replaced by hydrogen peroxide (H2O2).
This is the proprietary, patented UG3P process
developed by Ugine (Ref 1) or Henckel’s
“CleanOx” process.
Electrolytic pickling is commonly used in the
production of cold-rolled stainless. This process
uses alternating positive and negative polarization in baths using sulfates or nitrates. The H+
ions are produced by the polarization instead of
by an acid, so neutral solutions can be used.
The low solubility of trivalent chromium
(Cr+3), typically found in oxide scale, makes
such scale hard to remove. Use of strong acids
and complexing agents is required to get pickling to occur at an economically acceptable rate.
Hexavalent chromium (Cr+6) is much more
readily soluble in aqueous solutions, for example, oxidizing treatments such as Kolene,* a
molten mixture of nitrate salt and hydroxide, or
electrolytic conversion treatments such as electrolytic sodium sulfate. Both treatments convert
the trivalent chromium to hexavalent form, allowing for easy scale dissolution. Their use
may, however, lead to environmental problems
as hexavalent chromium compounds are toxic
and heavily regulated.
Pickling may be assisted by prior treatment of
the scale in molten 85% sodium hydroxide
(NaOH), 14% sodium nitrate, and 1% sodium
chloride bath. The strong oxidizing action of
such a bath chemically alters the chromium
oxide in the scale to a more soluble oxide without attacking the metal, making it more easily
removed by subsequent pickling.
Cleaning
Stainless steel is very resistant to chemicals
of many kinds, permitting it to be cleaned by
many aggressive agents. Because contamination
is “on” the surface of stainless rather than “in”
the surface as is the case with many materials,
simply using the cleaning agent appropriate for
* Kolene is a registered trademark of Kolene Corporation.
Chapter 16: Surface Finishing / 195
the contaminant will permit very thorough
cleaning of the stainless surface. The stainless
surface will not be harmed by cleaning as long
as strong halides, iron utensils, and abrasives
that alter the surface finish are avoided. Table 1
summarizes some cleaning recommendations
based on the contaminant to be cleaned.
The greatest controversy regarding cleaning
stainless steels involves the use of sodium
hypochlorite. Many highly qualified experts
flatly disapprove of its use because it can easily
leave a residue of chloride on the stainless surface, leading to localized corrosion. But, in reality, health considerations make this position untenable. Sanitation concerns in the food service
industry take precedence over this prohibition.
In fact, every stainless steel, as explained in the
“Corrosion and Oxidation” section of this Volume, is resistant to some level of chloride concentration at a given temperature and pH. Thus,
short-term, room temperature use of hypochlorite bleaches will generally not harm most types
of stainless steel if well rinsed after application.
Field experience has shown that stainless steels
with less than 16% chromium can be harmed by
hypochlorite bleaches. Damage occurs mainly
on abrasively polished surfaces in alloys with
16% chromium, such as 430, but does not occur
on roll-finished or bright-annealed surfaces.
Stainless steels with slightly higher chromium
and nitrogen, such as 201, 301, and 304, are not
attacked unless concentrated chloride solutions
are permitted to stay on the surface, particularly
in crevices.
The recommendation not to use abrasives
does not include soft abrasives such as calcium
Table 1
carbonate or brush material that is sufficiently
soft not to mar the stainless surface. The user is
encouraged to test whether a product meets this
requirement by testing it on a small, preferably
unexposed, area. Blasting the surface with carbon dioxide pellets is a rapidly growing process
for removing paint and other adherent, soft
coatings and deposits without damaging the surface of stainless steel.
These cleaning recommendations apply to industrial, architectural, and domestic uses of
stainless. For special levels of cleanliness required for medical, pharmaceutical, or semiconductor applications, refer to the chapters dealing
with those applications.
Passivation is a very commonly used surface
treatment to remove surface contamination, notably iron, and to form a passive film. The film
forms of its own accord when a clean surface
encounters moist air, but film formation can be
accelerated by controlling the environment.
The surface to be passivated should first be
cleaned by one of the methods discussed. This
allows uniform passivation and avoids contamination of the passivating solution, especially by
chlorides, which can cause a rapid attack. Those
stainless steels with more than about 17%
chromium can be pickled in 20% HNO3 at 50 to
60 °C (140 °F). Precipitation-hardenable (PH),
martensitic, and straight chromium grades with
less than 17% chromium should have 2.2 g/L
sodium dichromate added to that solution to
avoid attack. The free-machining grades are
the most easily attacked of all types. They
should be immersed in 5% NaOH for 30 min at
75 °C (165 °F) followed by a rinse before the
Recommended cleaning methods
Contaminant
Cleaning recommendation
Exterior soiling
Soap, detergent, or dilute ammonia
Fingerprints
Detergent and warm water or a hydrocarbon
solvent
Hydrocarbon solvent
Grease, oil
More severe stains, discolorations,
and rust stains
Nonscratching creams or polishes
Hard water scale, mortar
10–15% phosphoric acid, sulfamic acidcontaining, or oxalic-containing cleansers
Oxides, heat tint
If severe, treat by pickling; Scotchbrite©,
stainless scouring pad, or nonscratching
cream or polish
Paint
Alkaline, trisodium phosphate, or hydrocarbon
solvent
Comments
Use a soft cloth or sponge, clean water; dry with
forced air or a dry cloth.
Wax and oil polishes minimize fingerprinting. Glass
cleaner is appropriate for mirror finishes.
Alkaline cleaners may also be used in severe cases
but may require cleaning the entire surface to
maintain visual uniformity.
Do not use HCl-containing products.
Hypochlorite bleaches must be well rinsed to
avoid pitting.
Neutralize with ammonia, rinse, and dry.
Do not use HCl-containing products.
If abrasives must be used, they should blend
with existing surface finish in size and direction.
Follow manufacturer’s directions
196 / Stainless Steels for Design Engineers
nitric/dichromate passivation is performed. This
should be followed by a rinse, another 30-min
NaOH treatment, and a final rinse.
In recent years, citric acid has become a popular replacement for HNO3 because it avoids
the problem of toxic nitrates and hexavalent
chromium. When citric acid is used, a 10%
solution is applied for the same time and temperature as with HNO3. The use of NaOH before and after is still recommended for lowchromium and free-machining alloys. Refer to
ASTM documents A 967-01, A 380-99, B 91200, as well as Federal Specification QQ-P-35.
Brightening
Stainless steels can be brightened by chemical or electrolytic action, which selectively dissolves the surface in such a way that it becomes
microscopically smoother. Electropolishing is
the most effective means of accomplishing this.
It causes the surface roughness to decrease by
approximately one-half in Ra (the arithmetical
average of surface peaks and valleys as measured over a straight line). Electropolishing, like
pickling, also selectively removes any exposed
chromium-depleted regions, leaving only the
bulk alloy with the intended corrosion resistance on the surface. Apparent pitting resistance
is thus increased. The resulting surface is bright
and cleanable and provides the optimal corrosion resistance that a given alloy can achieve.
The reflectivity of the surface is a function of
the preelectropolished surface.
The most commonly used electropolishing
solution is 40% H2SO4 and 45% phosphoric
acid (H3PO4), balance water, used at 90 °C
(194 °F) with a current density of between 1.0
and 3.0 amp/m2. Other baths using perchloric
acid are technically good but carry the risk of
explosion.
Coloring
Stainless steel passive films are so thin that
they are quite invisible. Oxide films are thicker
and through optical interference can cause different colors. The films are formed thermally in a
process known as heat tint. The color depends on
oxide film thickness, which is a function of the
time at a given temperature, well metal composition, and oxygen partial pressure. The colors typically range from light yellow formed at 300 °C
(570 °F) through violet formed at 420 °C (790
°F) and dark blue formed at 600 °C (1110 °F).
Reproducibility is sometimes a problem.
Electrochemical methods permit more uniform and reproducible color and tougher films.
The most prominent is the International Nickel
Company (INCO) process. In this two-step
process, the steel is first immersed in a mixture
of one part chromic acid, two parts H2SO4, and
one part water at 30 °C (85 °F). This is followed
by an electrochemical treatment with the same
bath or with H3PO4 substituted for H2SO4.
The colors produced are correlated to thickness and treatment parameters (Ref 2) as shown
in Table 2. Rocha-Fila et al. confirmed that
these coloring treatments did not degrade the
pitting corrosion resistance since their formation mechanism more closely resembles passive
film formation than oxidation.
Aesthetic Surface Finishes
When steel is used in other than the asannealed and pickled state, it is often for aesthetic rather than functional reasons. The major
exception to this is temper-rolled strip, which
has a bright surface from cold rolling, but
whose normal use is strictly functional. The hotrolled or cold-rolled annealed and pickled surface finish is nonuniformly dull and unattractive
to most observers. When stainless steel first appeared on the market, a highly polished surface
was the paragon, as showcased on the Chrysler
Building in New York City (Fig. 1).
This surface finish was very expensive to produce since it had to be done by polishing and
buffing. Soon, the more economical abrasive
polishing became the standard, and it had the
benefit of removing the many cosmetic defects
that then were common to the manufacture of
stainless. Much higher levels of surface quality
were made possible with the development of the
Sendzimir mill to the bright anneal process permitted much higher levels of surface quality.
Table 2 Parameters for oxide film coloring of
stainless steel
Heat treatment time
10
15
20
25
30
35
40
50
Color
Oxide film
thickness, nm
Gold brown
Brown/red/blue
Brown/blue
Green/blue
Green/gold/blue
Gold/green
Gold/green/brown
Red/brown
70
100
120
140
165
185
210
250
Chapter 16: Surface Finishing / 197
Stainless steel surface finishing is governed by:
• High strength and high work-hardening rate
require more power for metal removal.
• High surface hardness influences which
abrasive will be effective.
• Low thermal conductivity can cause excessive temperatures during processing and
distortion.
• Oxidation (heat tinting) that occurs over
about 250 °C (480 °F)
• Residual stress due to surface working, especially in austenitic alloys
Fig. 1 View of the Chrysler Building in New York.
©iStockphoto.com/stevenallen
Mirror finishes could be achieved with cold
rolling followed by bright annealing. This has
not entirely replaced sheet buffing to obtain
mirror finishes, but sheet buffing has been relegated to those few applications where a high
degree of perfection is demanded. The main
methods of producing aesthetic surface finishes
on stainless steel are abrasive finishes and rolled
finishes. The latter are superior in uniformity
and corrosion resistance, but the polished finishes are still more common.
Polished Finishes
Polishing is carried out with coarse- to
medium-sized abrasives that are bonded to a
flexible backing. This is distinct from grinding,
in which abrasives are bonded rigidly to each
other on a rigid backing. Buffing is done by
very fine abrasives, which are loose and do their
work by being forced along the surface by a soft
material. The standard rule is that materials
used in finishing must not permit iron or iron
oxide particles to come in contact with the
stainless surface or passivation layer.
Grinding is a relatively coarse procedure with
usefulness that is largely confined to defect removal. Polishing for aesthetic purposes is generally done with abrasive coated belts and is
done to both coils and sheets. There are no official roughness values for the various surface
finishes. ASTM merely describes the finish by a
grit with which it is typically made. Table 3
shows the varying surface roughnesses that are
typical of polished stainless.
Neither the producer nor the customer is protected by a clear specification as of this writing.
Thus, the producer is encouraged to publish
standards, and the customer should specify minimum and maximum Ra values for his purpose.
The typical number 4 polish varies in roughness
from the beginning of the coil to the end because of the wear of the abrasive particles on
the belts. The difference can be more than 10
μin. (0.25 μm) in Ra, is quite visually apparent,
and will make adjacent pieces of stainless from
the same coil look different. This may be objectionable for certain products, especially appliances and architectural panels.
Polish appearance varies also with the presence or absence of lubricant. Lubrication during
polishing (i.e., wet polishing) gives sharper cuts
and less heat tint from frictional heating, resulting in more surface brightness. Polish appearance also depends on the length of the grit line
caused by an individual particle of abrasive. The
length of the grit line varies with the speed of the
material as it passes the rotating abrasive belt,
Table 3 Polished finish designations based on
grit sizes to achieve target surface roughnesses
Finish number
Grit number
Ra max, µm
Ra max, µin.
3
4
7
8
60–120
120–180
240–320
500
1.0
0.75
0.30
0.15
40
30
8
4
198 / Stainless Steels for Design Engineers
and the rotational speed and diameter of the
backup roll for the abrasive belt. This combination of variables, while seemingly controllable,
usually varies enough so that no two polished
sheets look identical to the trained eye. This
variability also causes problems in field repair.
Repair of a polished surface damaged, for
example, by welding or scratching is quite
challenging. A perfect match and blend to the
surrounding original surface is an art, and no
practitioner of that art can accomplish it perfectly. If the repaired area is small, then it can
be acceptable, if not undetectable, but for large
areas, repolishing of the entire surface is often
the best remedy.
A solution to the problem of matching polished surfaces is to use very long grit lines for
finishes of the same roughness. By eliminating
the variable of grit length, only correct pressure
and grit size are required to achieve good visual
matching. This can be done with a belt sander.
These finishes are called “hairline” in Japan and
“grainline” and other names in the United
States. Lack of uniformity, difficulty of repair,
and a decrease in corrosion resistance are the
chief drawbacks of polished finishes.
Bright Annealing
Aesthetic finishes that do not depend on abrasion are derived from bright annealing. Bright
annealing is annealing in a very low oxygen
atmosphere, either dissociated ammonia* or
hydrogen/argon. This process was originally
developed as a means of producing bright ferritic trim for automobiles. That use has largely
passed out of favor, but the process itself is intrinsically superior to annealing in air because
no oxide is formed only to be later removed,
usually at significant expense. Because the original product was strip, many older bright-annealing facilities are narrower than normal sheet
width. The wider, more modern lines are high
speed and wide enough to produce bright sheet.
Bright-annealed sheet will only be as bright and
flawless as the cold-rolled sheet that is annealed. Therefore, a quality product, generally
as mirror-like as possible, must be produced
with great care. The better producers have proprietary methods of prior pickling, annealing,
* At one time, “dissociated ammonia” referred to a 3:1
(mole ratio) mixture of hydrogen and nitrogen produced by
the catalytic decomposition of anhydrous ammonia. Today,
this composition is prepared by mixing gases. For many
stainless steels, the 3:1 H2:N2 can be varied without difficulty.
and rolling by which the surface is made flawless and very smooth. Final brightening occurs
with temper passing after bright annealing.
Rolled Finishes
The bright-annealed and temper-passed surface may be used in its mirror-like condition,
or it can be used as a basis for rolled finishes,
which can take on any appearance and can be
engraved onto a temper mill roll. All polished
finishes can be duplicated as rolled finishes.
But, in addition, many other designs such as
replicating fabric or leather, geometric designs, or matte finishes can also be made. The
pattern on the roll is impressed into the stainless surface, and elastic flattening of the roll
and the stainless cause about 50% less Ra on
the stainless than on the roll. This effect increases with the yield strength of the stainless
being processed. The benefits of rolled finishes
are:
• They can be made identically from coil to
coil since roll engraving is quite precise.
• They retain the enhanced corrosion resistance of bright-annealed material.
• They are less expensive to produce, so the
final product is generally priced lower than
polished material.
In theory, such finishes can be made from airannealed material, but finishes applied by roll to
two-dimensional surfaces that have been dulled
by pickling are somewhat gray and indistinct. If
the pickling is kept mild enough to retain surface
brightness, then it is possible insufficient prickling has occurred to remove the chromiumdepleted layer caused by the air anneal. Some of
the finishes produced by rolling are shown in the
architectural chapter (see Chapter 18). These finishes can be applied to all the normal sheet
alloys, including the lean duplex alloys. This has
permitted rolled finishes to be used on UNS
32003 for building exteriors in Doha, Qatar,
where corrosive conditions require a PREN
(percentage chromium equivalent) of 25.
The most critical applications from a surface
perfection point of view use rolled finishes
rather than abrasively finished number 4 polishes. These finishes also permit type 430 stainless steel to be used successfully in restaurant
applications without corrosion issues where
polished type 430 material previously had
corroded. There are also instances for which
polished UNS S31600 was corroding in coastal
Chapter 16: Surface Finishing / 199
architectural applications in the United States.
When the abrasively polished surface was
replaced with rolled finish UNS S30400
(Koolline), there was no further corrosion. Because alloying is such a high component of cost
in stainless, it makes sense to employ rolled finishes whenever possible. However, it must be
noted that when manual polishing is used to repair scratches or other damage in rolled finish
material, its corrosion resistance may be reduced.
Surface Alteration
In carbon steels, surface chemistry can be
changed to affect certain properties. Carburizing and nitriding are examples of such
processes. Simply using these processes on
stainless cannot be done because these elements combine too strongly with chromium as
carbides or nitrides, dramatically reducing the
corrosion resistance. There have been modifications to carburizing and nitriding that permit
austenitic stainless steels to have very high carbon levels implanted to a thin surface layer. By
exposing an activated stainless surface to a
high carbon fugacity at 470 °C (880 °F) for
around 200 h a 50 nm thick layer with 12 at.%
carbon in supersaturation can be achieved without carbide formation. The properties of this
layer are phenomenal; the hardness is 1000
Vickers 25 compared to 200 for the base alloy
(Ref 3). In addition, the corrosion resistance increases significantly. Only by such supersaturation with carbon could it be determined that
carbon, like nitrogen, is a powerful antipitting
alloying element when kept in solid solution.
These processes are at the very initial stages of
commercial use.
REFERENCES
1. Stainless Steels, Les Editions de Physiques,
1989
2. C. Rocha-Filo et al., J. Braz. Chem. Soc.,
Vol. 15 (No. 4), 2004, p 472–480
3. Y. Cao, F. Ernst, and G. Michal, Acta
Mater., Vol 51, 2003, p 4171–4181
Stainless Steels for Design Engineers
Michael F. McGuire, p 201-212
DOI: 10.1361/ssde2008p201
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
CHAPTER 17
Welding
NEARLY ALL WELDING of stainless steel
is done by end users or processors. Like thermal
processing, it is complex in theory and practice.
This chapter gives a basis for understanding the
influence of alloy composition and metallurgy
on the welding process, which must be respected as a process that combines melting, refining, and thermal processing. Knowledge of
each aspect is required for the process to be designed and executed properly.
The welding and joining of stainless steels requires knowledge of both the technology of the
welding or joining process and the response of
the steel to the thermal and mechanical effects
of the process. The welding process must, of
course, produce a sound joint, but it must also
result in the weld and its surrounding affected
metal having correct strength, toughness, corrosion resistance, etc. for the intended service
conditions. This chapter does not attempt to
teach welding. The main objective is to show
how standard welding technology is correctly
applied to stainless steels.
The foremost special consideration of welding stainless steel as opposed to carbon steel is
that the chromium in stainless steel, which is
what makes it stainless, must be protected from
oxidation, so that:
1. It stays in solution as a corrosion-resisting
element.
2. It does not form refractory oxides that would
diminish weld soundness.
Welding Characteristics of
Stainless Steels
Austenitic stainless steels are readily
welded by nearly all welding techniques. The
characteristics of austenitic stainless steels that
distinguish them from ordinary carbon steels in
welding are:
• Austenitic stainless steels have lower thermal
conductivity and higher thermal expansion
than carbon steels or ferritic stainless steels,
which can localize the heating, thus increasing the potential for residual stress and
therefore hot cracking.
• Stainless steels contain readily oxidized
chromium, which must be protected.
• Surface oxidation during welding depletes
chromium in all types of stainless steel from
the underlying surface, resulting in reduced
corrosion resistance unless this layer is removed.
• The possible formation of chromium carbides in the heat-affected zone (HAZ) can
cause susceptibility to grain boundary corrosion (sensitization).
• The possible precipitation of intermetallic
phases in the HAZ can lower toughness and
corrosion resistance.
• There is increased microsegregation in the
fusion zone with increasing alloy content.
• There are thermodynamically metastable
conditions due to the low diffusion rates in
the face-centered cubic (fcc) matrix.
The influence of carbon has been well addressed using low-carbon versions of all grades
whenever welding involves significant time between 600 and 900 °C (1110 and 1650 °F). This
prevents rapid precipitation by reducing the supersaturation of carbon. The older method of
preventing sensitization is to stabilize the alloys
with titanium, as in type 321, or with niobium,
as in 347. This is foolproof only if carbon levels
are low, less than 0.04%, since TiC can dissociate at elevated temperatures and not be able to
202 / Stainless Steels for Design Engineers
recombine successfully with titanium during
cooling, permitting a thin zone of sensitization
called knife-line attack. Fortunately, most 321
and 347 are produced with carbon levels below
0.03%. The higher carbon-stabilized alloys and
the high-carbon (>0.03%) unstabilized alloys
must be annealed after welding to redissolve
chromium carbides if the cooling was sufficiently slow for the carbides to have formed.
This is avoided only in thin-gauge (>1.5 mm,
0.06 in.) material or when the HAZ is drastically reduced, as in laser welding.
The high thermal expansion of austenitic
stainless steel can cause high residual stress
around welds, which may require annealing to
eliminate. Another serious threat posed by thermal stresses is hot cracking. This can occur to
material that has just solidified when geometric
constraints to contraction imposed by the surrounding material imposed act on weak grain
boundaries. This weakness occurs when the
steel solidifies in an austenitic mode. When
austenite freezes, it strongly rejects sulfur to the
intergranular areas, where it forms weak films.
This is solved by balancing the composition so
that alloys solidify first as ferrite, which does
not reject the sulfur, forcing it to precipitate as
sulfide inclusions within the grains. This approach is highly effective but cannot be used for
Fig. 1 The Schaeffler diagram. Source: Ref 1
some highly alloyed grades with compositions
that do not permit a ferritic solidification mode.
In such alloys, sulfur and other contaminants,
such as phosphorus, oxygen, zinc, and copper,
must be excluded from the weld zone. Welds of
less highly alloyed austenitics, generally those
with less than 20% chromium, which are balanced to freeze in a ferritic mode, retain some
ferrite at room temperature, normally between
3 and 10%. This is not harmful since the ferrite
is richer in chromium and in molybdenum, if
present.
The amount of ferrite expected can be measured by magnetic devices and estimated from
the Schaeffler diagram, a useful empirical mapping of weld metal phase composition shown in
Fig. 1. This diagram has an arbitrary cooling
rate resembling that of tungsten inert gas (TIG;
described in a separate section of this chapter)
welds. Faster or slower cooling will change the
relative amounts of ferrite and austenite because
of the need for diffusion to achieve the most stable phase balance. Very rapid cooling, as with
laser welding, tends to make austenitic welds
less ferritic and has the opposite effect in duplex
alloys.
The Schaeffler diagram has been improved
by the Welding Research Council’s adoption of
the modification shown in Fig. 2, which super-
Chapter 17: Welding / 203
Fig. 2 Welding Research Council’s (WRC’s) 1992 constitution diagram
imposes the solidification mode as a function of
the composition. The crucial line on this diagram is dotted-dashed line AF, which delineates
those compositions that solidify in a primary
ferrite mode, precluding the problem of intergranular solidification cracking.
Another problem particular to the more
highly alloyed grades is the formation of intermetallic phases from long cumulative exposure
to temperatures in the 600 and 900 °C (1110 and
1650 °F) range (coincidentally, the same as for
carbide precipitation). The slow diffusion of alloying elements in austenitics makes this a
lesser problem than in ferritics or duplex. This
adverse precipitation is largely prevented in the
modern, nitrogen-alloyed grades, so these alloys are recommended if extensive welding is
planned. The more highly alloyed grades also
suffer from greater microsegregation during solidification. This causes austenitic dendritic
cores to have lower chromium and molybdenum content and consequently lower corrosion
resistance. Thus, the welds have lower resistance to localized corrosion. This is addressed by
using more highly alloyed filler metal or by solution annealing the welds.
Restricting heat input to under 16 kJ/mm
(400 kJ/in.) and interpass temperature to under
150 °C (300 °F) helps to minimize each of
these risk factors inherent to the more highly
alloyed austenitic grades. Note that the influence of microsegregation of alloying elements
is separate from and in addition to the negative
influence of sulfur on the corrosion resistance
of welds. Austenitics at the alloy level of 316
and above should not have sulfur above
0.001% for these alloys to deliver the expected
corrosion resistance.
The austenitic stainless steel weld metal composition can be altered by the gases to which the
molten base metal is exposed. Lack of shielding
can lead to oxygen combining with chromium
and other elements, creating slag and depleting
the alloy of needed elements. Thus, oxygen-free
gas mixtures are used to exclude the ambient atmosphere from the molten pool during electric
arc welding. Inert gases provide the barrier,
while the addition of 3 to 5% by volume of nitrogen gives the necessary partial pressure to
ensure that welds will not be depleted of vital
nitrogen content. Figure 3 (Ref 1) shows the influence of nitrogen content of the shielding gas
204 / Stainless Steels for Design Engineers
Fig. 3 Effect of weld shielding gas composition on crevice
corrosion resistance of autogenous welds in AL-6XN
alloy tested per American Society for Testing and Materials
(ASTM) G-48B at 35 °C (95 °F)
on corrosion resistance of a highly alloyed
austenitic grade. Excess nitrogen in the shielding gas (e.g., more than 10%) can cause porosity in the weld, and greater than 5% is detrimental to the life of the tungsten electrode.
The heat from welding can produce a surface
oxide composed mainly of iron and chromium.
The underlying surface can be significantly depleted of chromium because of the loss of
chromium to this scale and therefore significantly lower in corrosion resistance. Pits can
start in this thin layer and propagate into sound
metal beneath. For heat-tinted surfaces, the
darker the tint, the stronger will be the effect. To
fully restore corrosion resistance, the area must
be ground to remove the oxide and any depleted
base metal. This should be followed by acid
pickling, which completes the removal of the
oxide and depleted zone.
Duplex stainless steels differ from austenitic
stainless steels in their metallurgical response to
welding mainly because their approximately
50% ferrite causes greater thermal conductivity
at lower temperatures, and ferrite has greater
diffusion rates. These alloys solidify in a completely ferritic mode, and since ferrite rejects
little sulfur on solidification, hot shortness is not
a problem. So, compared to austenitic stainless
steels, duplex stainless steels have the following
distinguishing factors:
• The ferritic solidification mode of duplex
stainless steels provides very good hot
cracking resistance. The rapid cooling of
welds produces welds and HAZ with more
ferrite than the parent metal by quenching in
the high-temperature ferrite.
• Duplex alloys are more sensitive to problems in the HAZ because their generally
high chromium and molybdenum content
plus their ferritic content make the precipitation of embrittling intermetallic phases more
rapid than in austenitics, so minimizing the
total time at high temperature is the overriding concern.
• While carbide sensitization is not an issue
with the duplex alloys, the formation of intermetallic phases can cause loss of corrosion resistance.
• Duplex, like all stainless types, must be protected from oxidation by shielding gas, and
since nitrogen is a crucial alloying element,
especially in duplex alloys, it must be a
component of the gas mixture.
• Cleaning before and after welding is equally
important in duplex as in austenitics.
Modern duplex alloys derive their impressive
strength, toughness, and corrosion resistance
from their nearly equal percentage of ferrite and
austenite. The nitrogen content of the austenite
brings its corrosion resistance up to that of the
ferrite phase, which is richer in chromium and
molybdenum. Nitrogen additions partition to
the austenite and thus both strengthens it and increases its corrosion resistance to close to that
of the ferrite. The early duplex alloys had a tendency to form excessive ferrite when welded
and formed embrittling intermetallic phases
rather rapidly. The additions of larger amounts
of nitrogen stabilized the austenite to higher
temperatures, so welds did not become so ferritic. The nitrogen also decreased the speed at
which intermetallic phases form, enlarging the
time window for welding without their precipitation. And, by promoting greater austenite formation at high temperature, the addition of high
(>0.12%) nitrogen actually reduces the tendency for chromium nitride precipitation. Despite these advances, the key precaution in
welding duplex alloys is to prevent the formation of embrittling phases while preserving as
close to a 50/50 austenite/ferrite structure as
possible. Minimizing time at red heat temperatures (500 to 900 °C, 930 to 1650 °F) is the objective. But, sufficient time must be spent above
about 1000 °C (1830 °F) to promote the formation of sufficient austenite. If the weld cannot be
annealed, increased nickel filler metal (e.g.,
2209 with 2205 base metal) should be used.
Thus, joint preparation must be done correctly
and not left to the welder to correct using timeconsuming remedial procedures.
Chapter 17: Welding / 205
Duplex stainless steels, because of their moderate thermal expansion and higher thermal
conductivity, can tolerate relatively high heat
inputs since these factors determine the stress
intensity that will be generated by thermal gradients. However, excessively low heat inputs
can result in fusion zones that are predominantly ferritic, with a resultant loss of toughness
and corrosion resistance. At the other extreme,
heat inputs that are too high lead to the formation of embrittling intermetallic phases. This
issue concerns the HAZ, which must dwell in
σ-forming temperatures for some period of
time. The key is to limit the time at those temperatures by not permitting interpass temperatures to exceed 150 °C (300 °F) because workpiece temperature has the greatest influence on
time at σ-forming temperatures. It is prudent to
impose this limitation when qualifying the weld
procedure and then monitoring the production
welding interpass temperature electronically to
ensure qualifying procedures are not more lenient than are those of production.
Postweld stress relief is not needed for duplex
weldments and indeed could be harmful because of the danger of embrittlement. Full annealing can be done and can restore the original
phase balance and composition that gives the
optimal toughness and corrosion resistance
found in wrought material.
Ferritic stainless steels can be split into two
groups for purposes of welding: the older semiferritic group and the more prevalent stabilized
ferritic group. The first group, in which
chromium is between 16 and 18% with carbon
up to 0.08%, is exemplified by the alloy 430.
These alloys form appreciable amounts of
austenite when heated above 800 °C (1470 °F).
Unless they are cooled extremely slowly, more
slowly than can be done in welds, the austenite
transforms to martensite, which is very brittle.
The stabilized grades commonly use titanium
or niobium to combine with the carbon and nitrogen, which otherwise would cause the formation of the high-temperature austenite, rendering the alloys ferritic at all temperatures.
The salient metallurgical characteristics for
welding of the two groups are:
• Both groups offer good thermal conductivity
and low thermal expansion.
• Both groups require protection from oxidation by shielding gases. The stabilized group
should not be exposed to nitrogen.
• The semiferritic group will form martensite,
which requires annealing to eliminate.
• The stabilized group can lose toughness via
excessive grain growth.
• The grades more highly alloyed with
chromium and molybdenum can form α' and
σ, leading to embrittlement.
The semiferritic alloys such as 430, 434, and
436 are seldom welded and often called unweldable. The reason is that the welds are invariably partially martensitic and thus normally
brittle. Only very specially controlled compositions of 430 can be welded successfully, and
these are not generally available commercially.
While the technical remedy for this is simply
annealing, it is seldom economically viable. It is
rare to see any welding more extensive than
spot welding of unexposed surfaces with these
alloys. If for some reason they must be used and
welded, then the techniques for welding
martensitic stainless steels should be employed.
The stabilized ferritic stainless steels are
commonly welded. The levels of stabilizing elements required to prevent austenite formation
and sensitization are well known and are reflected in the alloys’ chemistry specifications.
For 409, the required titanium level is Ti > 0.08
+ 8(C + N), while the requirement for the higher
chromium 439 is 0.20 + 4(C + N). These are
empirical relationships that take into account
that some titanium oxidizes before it can stabilize carbon and nitrogen. Niobium can replace
some titanium. This is discussed in detail in
Chapter 8 on ferritic stainless steels. Because of
the low toughness these alloys have in large
cross sections, these alloys are only rarely seen
with minimum section size of more than 3 mm
(0.11 in.) and normally have sections less than 2
mm (0.08 in.). Thus, successful welding is simplified to making a sound, well-shielded weld
without producing excessive grain growth in the
HAZ. In practice, this can be achieved by limiting heat input to less than 6 kJ/cm. An empirical
relationship between grain diameter D and heat
input E (kJ/cm) has been reported (Ref 2). In
the fusion zone, the relationship is:
D = 206 × E – 585.6
(Eq 1)
In the HAZ, it is:
D = 29.6 × E – 50.6 for up to 6.6 kJ/cm
(Eq 2)
and
D = 75 × E – 350 above 6.6 kJ/cm
(Eq 3)
The light gauges ensure sufficiently short
times at high temperature that precipitation of
206 / Stainless Steels for Design Engineers
intermetallic phases should not be a concern,
even though they can form, especially in superferritic alloys.
The impact properties of ferritic stainless
steels are always a concern because their transition temperature can become elevated to ambient levels. It has been determined that there exists an optimum level of titanium around 0.10%,
which ensures this minimum transition temperature (Ref 3). Because it is difficult to have low
enough carbon plus nitrogen to stabilize at this
titanium level, dual stabilization with titanium
and niobium as well as not having excessive
heat input are the best way to ensure weld
toughness.
Especially in the superferritics, maintaining
the benefits of having the fairly precise balance
of carbon plus nitrogen to the stabilizing elements titanium and niobium requires that neither carbon nor nitrogen come into contact with
the weld pool. Likewise, oxygen must be rigorously avoided because it will quickly deplete
the essential titanium, which is even more readily oxidized than chromium. Extraordinary surface cleaning at and near the weld will pay dividends in final quality.
Martensitic stainless steels vary little in alloy
content, ranging from 11 to 18% chromium with
small amounts of nickel and molybdenum.
Their carbon content ranges from 0.10 to over
0.30%. Thus, the major challenge they present
is avoiding the potential cracking, which is
most likely to occur in the HAZ from stresses
caused by the austenite-to-martensite transformation on cooling. Since this transformation
cannot be avoided, the desired approach is to
start with a well-tempered or annealed material
and then preheat and maintain high interpass
temperatures. For low carbon levels, below
0.10%, preheat can be omitted, but between
0.10 and 0.20% carbon, preheating to 250 °C
(480 °F) is advised and for higher carbon levels,
300 °C (570 °F). The problem becomes more
severe with increasing carbon level because the
transformation takes place at lower temperatures in more brittle material. Even with preheating, distortion may be encountered. For all
normal uses of martensitic stainless steels, a
final heat treatment is required to achieve the
quenched and tempered properties for which
these alloys are designed.
Aside from the cracking consideration,
martensitic welding considerations are similar
to, but less stringent than, those of low-alloy
stabilized ferritic stainless steels with regard to
cleanliness and shielding. If mechanical requirements permit, the use of austenitic (309L)
weld filler metal should be considered. The soft
joint may deform to accommodate thermal
strains and thus minimize weld cracking.
Precipitation-Hardening Stainless Steels.
Last, precipitation-hardening (PH) stainless
steels, while very complex metallurgically, are
straightforward from a welding perspective.
Obviously, any heat treatment to achieve the
properties of which these alloys are capable
must be a final step. The considerations in welding them are:
• Shielding must be sufficient to prevent loss
of oxidizable alloying elements such as titanium, aluminum, and, of course, chromium.
• Filler metal must match the base metal if
like properties are required.
• Postweld heat treatment solution annealing
must be adequate to homogenize weld solidification segregation.
• Austenitic PH grades are fully austenitic and
subject to hot short cracking.
• The high aluminum or titanium contents of
many PH alloys cause their welds to be
“slaggy,” and these slaggy welds have are irregular with objectionable recesses,
crevices, or prominences.
These alloys are easily welded and not prone
to cracking or developing embrittling phases.
But, because these alloys are designed for extreme mechanical performance, it is essential to
preserve their correct chemistry by shielding
with a fully inert gas mixture. If mechanical
properties equal to that of the base metal are not
required in the weld, then austenitic filler, such
as 309L, can be used.
Table 1 summarizes the major metallurgically
important parameters for the various types of
stainless alloys. It is prudent to consult with the
manufacturer’s data sheets for specific recommendations on alloys that they produce as they
are often privy to test data and user experience
that cannot be found elsewhere in the literature.
Material Selection and Performance
Stainless alloys that are prone to precipitation
of intermetallic phases require special prewelding consideration. Such alloys include duplex,
superferritic, and superaustenitic alloys. Any
amount of time for which these alloys have
been exposed to temperatures at which inter-
Chapter 17: Welding / 207
Table 1
Welding parameters for various stainless steels
Alloy group
Filler
Heat input
kJ/cm(max)
Shielding
gas
Preheat
Interpass
max
Postweld
heat treat
Austenitic
...
20–40
Ar+2% O2,
Ar/3% CO2/2% H2
He+7.5%Ar+2.5 CO2
Same
Same
Same
Same
Same
Same
Argon/helium or
argon + 3–5%
N2:no O2
Argon/helium
150 oC
150 oC
None or full anneal
301, 302, 304
304L
309
310
316L, 316Ti
321, 347
Superaustenitic
308, 308L
308L
309, 310
310
316L, 317L
347, 308L
22, 675, 276
Same
Same
Same
Same
Same
Same
16
...
...
...
...
...
...
50 oC
...
...
...
...
...
...
100 oC
...
...
...
...
...
...
None or full anneal
PH grades
Same as base alloy
20–40
no
...
Full solution anneal
Martensitic
410
410, 308, 309L
20–40
Ar+2% O2,
He+7.5%Ar+2.5 CO2
250oC
250 oC min
Slow cool
420
420, 308, 309L, 310
20–40
Ar+2% O2,
He+7.5%Ar+2.5 CO2
250 oC
250 oC min
Anneal
440
440, 308, 309L, 310
20–40
Ar+2% O2,
He+7.5%Ar+2.5 CO2
250 oC
...
...
Supermartensiic
Same as base metal
20–40
Argon/helium
no
...
Full solution anneal
430
430, 309L
20–40
Ar+2% O2,
He+7.5%Ar+2.5 CO2
no
...
Subcritical anneal
434
309 Mo L
20–40
Ar+2% O2,
He+7.5%Ar+2.5 CO2
no
...
Subcritical anneal
409
410L, 308, 309L
6.0
Ar+2% O2,
He+7.5%Ar+2.5 CO2
no
n.a.
none
439
439L, 309L, 316L
6.0
Ar+2% O2,
He+7.5%Ar+2.5 CO2
no
n.a.
none
Superferritic
2003, 2101, 2304,
19-D
2205
29-4C
2209
6.0
5–25
Argon/helium
Argon + 3% N2
no
no
n.a.
150 oC
None or full anneal
None or full anneal
2209
5–25
Argon + 3% N2
no
150 oC
None or full anneal
25 Cr duplex
2507
superduplex
25Cr-10Ni-4Mo-N
25Cr-10Ni-4Mo-N
5–25
5-25
Argon + 3% N2
Argon + 3% N2
no
no
150 oC
150 oC
None or full anneal
None or full anneal
Ferritic
PH, precipitation hardenable
metallic phases form without full subsequent
homogenization anneal is time that the welder
cannot use to complete a satisfactory weld before precipitation occurs. Thus, accurate knowledge of material history is vital. Likewise, variations within specification of nitrogen content
influence the time it takes intermetallic phases
to form. Once a welding procedure is qualified
for an alloy with given nitrogen content, use of
lower nitrogen alloys would not be prudent engineering practice.
Austenitic stainless steels that are intended
for autogenous welding are often specified with
elevated sulfur levels, on the order of 0.005 to
0.015%. This is done to improve weld penetration through the so-called Marangoni effect.
This effect exploits the temperature-dependent
surface concentration of sulfur in the weld pool,
which causes a decreased surface tension toward the hotter center of the pool, causing the
molten pool to flow toward the center on the
surface and then flow downward, shooting the
hottest metal to the bottom of the weld pool, as
shown in Fig. 4. This speeds welding and minimizes weld and HAZ width, which is a good
thing. The effect on corrosion resistance is less
desirable since the abundant MnS inclusions
that result from the higher sulfur levels decrease
pitting resistance. This decrease in corrosion resistance can only be eliminated by a long anneal. Unfortunately, the pipe purchaser cannot
know if the pipe has had a sufficient anneal.
without (right) sulfur. Source: Adapted from Ref 4
In-line induction annealing is insufficient for
this purpose. Furnace anneals of about an hour
are required. For alloys like 304L and 316L, the
user should always require material chemistry
certifications and assume that any sulfur levels
above 0.003% are going to result in decreased
pitting resistance of 1 to 5 PREN (pitting resistance equivalent number), which means up to 10
°C (18 °F) decrease in critical pitting temperature, roughly the difference between 304 and
316 in performance. This also applies to girth
welds done by the pipe user.
Welds are essentially a casting in the midst of
wrought material. In addition to inclusions decreasing weld corrosion resistance as mentioned, solidification segregation can also cause
microscopic regions to be poorer in corrosionresisting alloying elements chromium, molybdenum, and nitrogen. This effect is minimal for
low-alloy material, but for highly alloyed
austenitic grades, it is a major effect, as shown
in Fig. 5. Eliminating this effect requires a thorough homogenization anneal.
The use of filler metal with higher corrosion
resistance does not totally offset the influence of
welding on corrosion resistance because some
of the base metal melts and is not altered in
composition by the filler metal. This is called
the unmixed zone. It is essentially a zone with
properties equal to that which would occur in an
autogenous weld, that is, the corrosion resistance is lower depending on total alloy level and
sulfur content.
Welding Processes
All stainless steels should be very clean prior
to welding. The chemistries of both base metals
and filler metals are carefully formulated to produce the mechanical and corrosion properties
that these alloys have been designed to produce.
Virtually any contaminant can either interfere
with the welding procedure or detrimentally
90
85
80
75
70
65
60
55
50
45
40
35
30
25
20
15
10
5
0
−5
194
185
176
167
158
149
140
131
122
113
104
95
86
77
68
59
50
41
32
23
Unwelded
Welded
1
2
3
4
5
6
Molybdenum, wt%
Critical pitting temperature in 6% FeCI3, °F
Fig. 4 Metal flow directions in a weld pool with (left) and
Critical pitting temperature in 6% FeCI3, °C
208 / Stainless Steels for Design Engineers
7
Fig. 5 The influence of molybdenum on critical pitting temperature. Source: Adapted from Ref 5
alter the composition of the welded joint, which
in turn can alter corrosion and mechanical properties and compromise the entire structure.
Moisture, paint, dirt or grease, oil, and oxides
all can negate good material, good welding
technique, and good procedural qualification.
Cutting fluids, especially sulfurized oils, are especially detrimental and should be removed
completely prior to welding. Preheating is never
strictly forbidden since it is required to eliminate moisture.
Joint design does not differ in principle from
that of other steel weldments. There is, however
an increased need for dimensional uniformity
for the alloys susceptible to intermetallic precipitation since minimizing time at temperature is a
priority, and variations in joint geometry impede
the swift completion of the weld. This is also
true for alloys that are susceptible to excessive
grain growth, such as the stabilized ferritics, or
to sensitization.
Figure 6 shows some joint designs appropriate to stainless steels, including the more sensitive alloys. These, like all joint designs, aim to
ensure full penetration without burn through.
Gas tungsten arc welding (GTAW)/tungsten inert gas (TIG) is commonly used for the
automated production of stainless steel pipe and
tube, as well as manual short runs. It is versatile
and generally used when thicknesses are less
than 6 mm (0.2 in.). It can produce very highquality welds. A constant-current power supply is
preferred. It is best performed with the DCSP (direct current straight polarity) electrode negative
Chapter 17: Welding / 209
Groove
t
Process
Thickness
th, mm (in.)
Gap d,
mm (in.)
Root K,
mm (in.)
Bevel α(°)
GTAW
3–5
1–3
...
...
GMAW
3–5
1–3
...
...
SMAW
3–4
1–3
...
...
SMAW
4–15
1–3
1–2
55–65
d
a
GTAW
3–8
1–3
1–2
60–70
GMAW
5–12
1–3
1–2
60–70
SAW
9–12
0
5
60
SMAW
>10
1.5–3
1–3
55–65
GMAW
>10
1.5–3
1–3
60–70
SAW
>10
0
3–5
80
SMAW
>25
1–3
1–3
10–15
GMAW
>25
1–3
1–3
10–15
SAW
>25
0
3–5
10–15
GTAW
>3
0–2
...
...
GMAW
>3
0–2
...
...
SMAW
>3
0–2
...
...
SMAW
3–15
2–3
1–2
60–70
GTAW
25–8
2–3
1–2
60–70
GMAW
3–12
2–3
1–2
60–70
SAW
4–12
2–3
1–2
70–80
SMAW
12–50
1–2
2–3
10–15
GTAW
>8
1–2
1–2
10–15
GMAW
>12
1–2
2–3
10–15
SAW
>10
1–2
1–2
10–15
d
a
k
d
a
k
d
r = 6-8mm
t
d
a
d
a
d
r = 6-8mm
GMAW, gas metal arc welding; GTAW, gas tungsten arc welding; SAW, submerged arc welding; SMAW, shielded metal arc welding
Fig. 6 Joint designs. Courtesy Ugine S.A.
technique. It is helpful to incorporate a highfrequency circuit to aid in establishing the arc.
Thoriated electrodes containing 1.7 to 2.2% thoria are recommended because they have better
emissive properties and provide better arc stability at higher currents. If consumable electrodes
are used, the shielding gas precludes the need for
coatings. The weld metal alloys are not necessarily the same as the parent alloys but are chosen
based on their ability as weld metals to provide
the most acceptable corrosion and mechanical
properties. This sometimes means using
austenitic filler with a ferritic base or higher
nickel content in an austenitic or duplex base to
compensate for the solidification rate or inherently lower corrosion resistance of the weld.
The shielding gas must replicate the controlled gas mixtures used to refine stainless steel
and establish the original composition. The
weld pool exposes a great deal of surface area to
the atmosphere in a very turbulent manner. Gas
flows, usually 12 to 18 L/min, must be adequate
to prevent air infiltration by aspiration or turbulence before arc contact, ideally until temperatures cool to below oxidation temperatures.
For manual GTAW using a filler wire, the
wire should be fed continuously into the weld
pool. Intermittent wire addition can lead to creation of zones of essentially autogenous weld,
negating many of the benefits of filler metal addition. Moving the tip of the wire in and out of
the protection of the gas shield is especially
210 / Stainless Steels for Design Engineers
bad. The hot tip can carry oxides and nitrides
into the weld, defeating the action of the shield
gas and impairing weld quality.
Gas metal arc welding (GMAW) is arc
welding in which a consumable electrode provides larger amounts of filler weld metal than
practical in GTAW. There are three GMAW
techniques:
• Pulsed arc transfer
• Spray transfer
• Short-circuiting transfer
Pulsed arc transfer employs a power source
that is switched rapidly to provide transfer of
weld metal droplets at regular intervals. Spray
transfer uses a high current to form a stream of
fine drops from the end of the electrode. This is
done with high power, resulting in a large fluid
weld pool, and therefore limits the technique to
horizontal orientations and thick material.
Short-circuiting transfer uses arc contact with
the workpiece at low power to melt the electrode, after which the short circuit is broken,
and material transfer ceases. The technique creates a minimal weld pool and is viable in many
orientations. It is a low-heat process suitable for
thin material but may cause lack of penetration
defects if used for thick-section welding.
For all GMAW processes, excessive protrusion of the wire should be avoided; otherwise,
the full benefit of the inert gas shielding may be
lost.
Submerged arc welding (SAW) employs a
consumable electrode immersed in a conductive
flux that acts as a protective shield from the atmosphere. The arc is struck through the flux,
and gravity deposits the molten metal to the
workpiece. The large weld pool has high heat
input and can deposit large amounts of metal
relatively quickly. Thus, SAW may be preferable to multipass techniques for alloys such as
duplex for which time at temperature is limited.
It is restricted to horizontal orientations and requires postweld slag (flux) removal.
Shielded metal arc welding (SMAW) is
done manually with short lengths (“sticks”) of
coated electrodes. This method has great versatility with some trade-off in cost and quality.
This last aspect is arguable, but the lack of
shielding gas may introduce oxygen to the weld
metal, which can be detrimental to toughness.
Flux cored wire (FCW) welding is a method
that is able to accommodate a large range of
thickness and orientations while providing high
deposition rates. The equipment is the same as
for GMAW, but the consumable electrode, the
FCW filler metal, has a flux core that supplements the shielding gas. Because of the flux, the
shielding requirements are reduced; gases can
be argon/25% carbon dioxide for horizontal
welding with current and voltages from 150 to
200 amp and 22 to 38 V, respectively. Vertical
welds can use 100% carbon dioxide with amperage of 60 to 110 amp and voltage of 20 to 24
V. Flow rates of gas are 20 to 25 L/min. It is
possible to get high-carbon welds, which may
not resist corrosion as well as desired, so as always, weld qualification, including corrosion
evaluation, is critical.
Oxyfuel gas welding (OFW), “torch” welding, uses oxygen to accelerate fuel (typically
acetylene) combustion to produce temperatures
that can melt steels. By controlling the fuel-air
mixture, the flame can be made nonoxidizing
for low-alloy steels. However, these “neutral”
flames can simultaneously oxidize and carburize stainless steels. Thus, the OFW process is
not suitable for use with stainless steels.
Laser welding has become a major production method when it can be automated, as for
pipe and tube or high-production manufactured
items, such as air-bag canisters. Metallurgically,
it resembles resistance welding in that both
have minimal HAZ and very high quenching
rates, both of which can have a pronounced effect on some types of stainless steel. The effect
is to undercool the molten metal and suppress
the transformation that would normally occur.
So, an austenitic alloy that normally solidifies in
a ferritic mode before transforming to austenite
freezes directly as austenite. The freezing is so
rapid that the normal hot shortness of austenitic
solidification is avoided, so quality is not compromised. In fact, laser welds quench the material so rapidly that corrosion resistance is enhanced since inclusions cannot nucleate and
grow. Duplex alloys, on the other hand, freeze
in their high-temperature ferrite structure because the fast quench prevents the nucleation
and growth of austenite. Unless this ferrite is
heated to permit austenite to form, lower-toughness welds will result. Ferritic, martensitic, and
PH alloys are not harmed by the rapid quench.
Resistance welding is readily done on most
types of stainless steel. Allowance must be made
for the lower thermal and electrical conductivity
of stainless steels compared to other common
materials. Most resistance welds, including both
seam and spot welds, have deep, tight crevices
adjacent to the welds. The possibility of crevice
Chapter 17: Welding / 211
corrosion in these regions should be considered
when contemplating the use of spot welds in
stainless materials. The possibility of entrapment of foreign material and the difficulty of removing it from such crevices should also be
considered, especially in equipment for food
handling, pharmaceutical production, etc.
High-frequency induction welding of
stainless steel is more difficult than for lowalloy steel because of the refractory nature of
chromium oxide, which has a higher melting
temperature than does the stainless base metal.
This is opposite from the situation in low-alloy
steels, for which the iron oxide melts at a lower
temperature than does the iron base metal. The
presence of this refractory oxide on the surfaces
to be joined makes it more difficult to obtain a
defect-free weld.
Thermal cutting of stainless steels is routinely practiced, but the processes and parameters used are determined by the refractory nature of the chromium oxides that form on
stainless steels. The high temperatures attainable with lasers or plasma arc torches provide
good cutting action, and these processes are frequently used. To expand the range of thicknesses that can be cut or to increase cutting
speed, supplemental oxygen or nitrogen blast
jets may be used. Stainless steels may also be
cut using oxyfuel equipment if supplemental
iron powder is used. Combustion of the iron increases the temperature, while the iron oxide
helps flux the refractory chromium oxide. Thermally cut edges of stainless steel usually require
subsequent cleaning, typically by grinding or
milling. Chemical cleaning of all surfaces of cut
pieces to remove heat tint, fume deposits, and
other contaminants is advisable.
Soldering and brazing are possible with all
stainless steels. Soldering is done below 450 °C
(840 °F), while brazing is done above 450 °C
(840 °F). Solders are generally alloys of tin and
bismuth, lead, silver, or antimony or combinations of several of these. Brazes are normally either silver based or nickel based. The
chromium-rich oxide coating must be removed
by a suitable flux for bonding to occur. Fluxes
are typically acid type with chlorides. Thus,
after the soldering or brazing, the flux must be
thoroughly removed to prevent subsequent pitting corrosion. Brazing temperatures must be
chosen to avoid ranges at which unfavorable
phases form. The best range can be determined
from examining temperature ranges to be
avoided in the thermal processing chapter
(Chapter 13) of this book. Brazes and solders
rarely match the corrosion resistance of stainless steels, and careful attention should be given
to the potential for galvanic and other forms of
corrosion when considering the use of soldered
or brazed joints with stainless steels.
Welding Practices
Safety must always be considered when
welding. In addition to the normal hazards
(which are not discussed here) associated with
welding, welding of stainless steels presents a
special hazard: hexavalent chromium. The fume
created by welding stainless steel contains significant concentrations of chromium trioxide
and other forms of hexavalent (Cr+6) chromium.
Hexavalent chromium is a carcinogen and regulated by the Occupational Safety and Health
Administration (OSHA). Exposure to and inhalation of stainless steel welding fumes must
be avoided. The product exposure limit for
hexavalent chromium is 5 μg/m3 as of December 31, 2008. Refer to OSHA for further updates
on this limit. Use of fume extraction equipment
is generally the preferred method of minimizing
hexavalent chromium exposures. Positioning
and operation of the fume extraction device
must be done precisely to ensure effective fume
removal while avoiding excess turbulence,
which can cause loss of effective inert gas
shielding of the weld pool. Thermal cutting of
stainless steels also generates hexavalent
chromium, and similar procedures are required
to minimize exposure during such operations.
Nondestructive Evaluation (NDE) is used
almost universally to ensure weld quality. All of
the standard NDE techniques used with other
materials are applicable to stainless steel weldments. Allowance must be made for the differing physical properties of stainless steels, and
appropriate reference defect standards must be
provided. However, one technique—magnetic
particle inspection—is problematic. The presence of bands of persistent austenite in martensitic or PH stainless steels can lead to spurious
defect indications. For this reason, magnetic
particle examination of stainless steel welds is
best avoided.
Recent developments in stainless steel have
been made with weldability as a major consideration. Highly alloyed, low-carbon martensitic
alloys for line pipe have been developed with
212 / Stainless Steels for Design Engineers
the express purpose of use in the as-welded condition. The low carbon makes welds of this material that are tough and do not require tempering, so girth welds in the field are possible.
Likewise, the lean duplex alloys have very
delayed precipitation of intermetallic phases because of their higher nitrogen and lower
chromium and molybdenum contents. This
makes welding of these alloys much more foolproof than with the early duplex alloys, such as
S31803. The dual-stabilized ferritic alloys have
tougher welds than those stabilized with only titanium or niobium.
New developments in welding also have an
impact on stainless steels. The friction stir welding (FSW) process offers the promise of reliable
solid-state joining. By avoiding melting and
resolidification, issues associated with solute
redistribution are eliminated. The relatively low
temperatures involved essentially eliminate
generation of weld fume (see the discussion of
safety). Other new welding processes, such as
multiple (GTA or GMA) torch welding, laserassisted GMA or GTA welding, etc. promise
greater productivity.
REFERENCES
1. D.J. Kotecki, Welding of Stainless Steels,
Welding, Brazing, and Soldering, Vol 6,
ASM Handbook, ASM International, 1993,
p 677–707
2. B. Aziez and R. Feen, Sheet Metal Ind., 1,
1983, p 28–34
3. S.D. Washko and J.F. Grubb, Proc. Int’l
Conf on Stainless Steel, 1991, Chiba, ISIJ
4. Stainless Steels, Les Editions de Physiques,
1992, p 786
5. A. Garner, Corrosion, 37, 1981, p 178
Stainless Steels for Design Engineers
Michael F. McGuire, p 213-223
DOI: 10.1361/ssde2008p213
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
CHAPTER 18
Architecture and Construction
Summary
STAINLESS STEEL IS OFTEN EMPLOYED
as an architectural material; the material can itself be viewed as a metaphor for architecture, a
discipline that must balance aesthetics, economy,
and structural integrity. Stainless steel’s unique
combination of beauty, strength, and economy
makes it a remarkably appropriate material for
uses as diverse as sculpture and concrete reinforcing bar. However, stainless steels are complex; they come in many different grades (chemical analyses), and these grades have varying
strengths, appearance, resistance to corrosion,
availability, and costs. The success of a building
project involves careful planning for the
appropriate use of materials. This chapter deals
with the technology of stainless steel as it pertains to its proper use in architecture, art, and
construction.
can choose from a variety of stainless alloys
with sufficient corrosion resistance to withstand
any environment.
The ability of stainless steel to resist corrosion resides in its chromium-rich superficial
passive layer. Stainless steel by definition must
contain slightly more than 10% Cr. The passive
layer forms spontaneously in air or water, and
if it is removed, say by abrasion, it re-forms by
itself. This is explained in greater technical
Corrosion Resistance
Corrosion is the life-limiting factor for architectural metals. Steel, copper, aluminum, lead,
bronze, and other alloys react with the environment and degrade over time, as does wood,
stone, plastic, paint, and even glass. With stainless steel, it is possible to choose a material that
can withstand attack from the environment indefinitely. One of the first major uses of stainless steel in architecture was in New York’s
Chrysler building, completed in 1930. Despite
the rather crude methods of early production,
limited alloy options, and lack of application
experience involved in its construction, the
domed top of the Chrysler building still shines
undiminished by the harsh coastal and urban
climate (Fig. 1). With today’s technology, one
Fig. 1 The Chrysler building with its famous bright stainless
details. Copyright © iStockphoto.com/Steven Allen.
Used with permission
214 / Stainless Steels for Design Engineers
detail in this book in the chapters on corrosion,
but the key aspects are that the chromium
atoms on the surface of the metal react with
oxygen in air and water to form with neighboring iron atoms into a tight, ionically nonconductive layer that prevents any further oxygen
penetration. This layer is mere atoms thick and
completely invisible.
The strength of the layer in resisting corrosive
attack is proportional to several key alloying
elements: chromium, molybdenum, and nitrogen. They contribute according to the following
formula:
PREN = %Cr + 3.3(%Mo) + 16(%N)
(Eq 1)
PREN stands for pitting resistance equivalent
number. This number can be related to resistance to the mildest form of corrosive attack that
stainless steel undergoes, pitting corrosion.
Pitting is a “weakest link” phenomenon in
which corrosion begins in small, micron-size
parts of the surface and then grows by virtue of
the more aggressive media that form within
them because of the corrosion reaction’s products. Pitting occurs in environments that contain
chlorides. Chloride ions compete with oxygen
and disrupt the integrity of the protective passive layer. As alloys become richer in the alloying elements mentioned, their ability to maintain
the passive layer can overcome the chlorides’
ability to destroy it. The key is to choose an
alloy rich enough in chromium, molybdenum,
and nitrogen to withstand any environment the
structure will experience. Pitting corrosion, and
another similar form of corrosion called crevice
corrosion, can be prevented by proper choice of
alloy, finish, and design. Crevice corrosion occurs when recessed spaces are small enough to
act like a corrosion pit. The acidity within the
crevice increases because of restricted diffusion
in and out of the crevice, just as happens within
a pit. The buildup of iron and chloride ions
makes a very corrosive medium that disables
the passive film formation. A design that avoids
crevices is the best defense.
The decision criterion for material selection
with stainless steel should then be, Which grade
and finish will exclude the possibility of pitting
corrosion at the lowest cost? Then, proper design should be used to exclude the possibility of
crevice corrosion. Other factors involving
strength and fabrication should also be considered. Table 1 ranks a number of stainless alloys
by pitting resistance.
Balancing Corrosion Resistance,
Processing Characteristics, and
Economy
The rule of thumb for grade selection is usually somewhat oversimplified to recommend the
use of types 430 and 304 on interior applications, type 304 on exteriors where salt is not a
problem, and 316 where road salts or seacoast
effects make a more corrosion-resistant grade
necessary. A leading architectural metals company (Ref 1) makes the following recommendations, which mirror these traditional views:
• Type 304 should be used for most exterior
applications.
• Type 316 should be used within ten miles of
saltwater bodies. However, if the building is
subject to saltwater spray, a nobler grade of
stainless steel, such as 2101 or 2003, should
be specified.
• In close proximity to deicing salt use, even
on nearby roadways where vehicle traffic
can create airborne particles, type 316
should be used. If periodic rinsing will not
occur on all exterior surfaces, these areas
must be washed each spring. If dependable
maintenance is not predicted, a nobler grade
of stainless steel, such as 2101 or 2003,
should be used.
• Specify types with low carbon, less than
0.030%, if welding will be employed.
• Any grade, including type 430, may be used
in interior applications.
• In the most severe environments—high heat
and humidity, low rainfall, and high salinity,
such as are found in Middle Eastern countries—a grade with a PREN of 25 or above
is recommended.
These guidelines are based on the admittedly
easy availability of these alloys and a lack of
concern for cost during times of peak raw material prices. For projects where quick availability
is not more important than cost, 439 and 201
should be considered as viable replacements for
304. Stainless steel 2003 (UNS S32003) or an
equivalent lean-duplex grade can replace 316 at
a cost advantage during times of high alloy cost
Table 1 Ranking of common stainless steels by
pitting resistance equivalent number (PREN)
Alloy 430
PREN 15
439
17
201
17
304
19
316
24
2101 2003 2205 2507
26
28
35
38
Higher PREN values indicate greater pitting resistance.
Chapter 18: Architecture and Construction / 215
as have been experienced on occasion, such as
during the period of 2004 to 2007. It should be
noted that the leaner alloys suggested (439, 201,
and lean-duplex alloys such as 2003) can be
somewhat more difficult to form. If panel designs call for 90° bends, this is not an issue.
However, for applications requiring severe
forming, as in the case of a double-lock seam on
a standing seam roof, these grades can provide a
challenge to the fabricator/installer. Further,
these leaner grades can pose challenges with
certain finishing methods, such as abrasive polishing and embossing. To avoid unwanted complications related to grade selection, the specifier should consult a competent architectural
metals supplier. This effort will ensure a viable
specification is written that will balance cost
with the necessary performance attributes to
make the part as well as resist corrosion once
installed.
It is valuable to know, in times of high nickel
prices, that both the low-nickel 201 and the nonickel 439 can be used in place of 304, while
2101 (UNS S32101) and 2003 (lean duplexes)
can replace 316. To obtain these grades usually
involves working with a producing mill since
they are not typically stocked in service center
inventories. However, any competent architectural metals supplier will not shy away from the
use of specialty grades where appropriate.
Surface Finish and Corrosion Resistance
Surface finish is usually an aesthetic choice,
but it has a significant influence on corrosion resistance and must be factored into grade selection. Mill finishes such as 2B and 2D are inconsistent because they are annealed and pickled to
remove oxides. These unattractive surfaces,
however, have correct corrosion resistance for
their alloy content. Welding or abrading the surface degrades the corrosion resistance by a significant amount. An un-heat-treated weld has
lower resistance to corrosion in proportion to
the alloy content of the grade. Type 316 welds
have the corrosion resistance of wrought 304.
Abrasion has a similar effect. Type 316 with a
No. 4 polish behaves like 304 with a 2B mill
finish. Because welds are abraded, this compounds the effect.
Very smooth abrasively polished finishes mitigate this reduced corrosion resistance, as
shown in Fig. 2. The effect of reduced corrosion
resistance in abraded stainless steel surfaces is
not seen on finishes that are produced by patterns imprinted by hard-rolling mill rolls that
have been engraved with the desired pattern in
reverse; for this reason alone, this method of
surface finishing is recommended. These architecturally useful surface finishes are produced
by the preferred rolled-on, or embossed,
method. In addition to their advantage in corrosion resistance, they are extremely uniform
from batch to batch, unlike finishes produced by
abrasive-coated belts, which change in grit
coarseness with use.
Balancing Service Environment,
Design Requirements, and
Maintenance Considerations
An expert system has been developed that enables designers and specifiers to analyze the
trade-offs of climate, design requirements, and
maintenance on grade selection (Ref 3). Answering the questions in Fig. 3 for a particular
application yields a score that can be used to
identify an appropriate alloy according to the
scale shown Fig. 4. The Nickel Institute, formerly the Nickel Development Institute, also
offers excellent publications on topics related to
alloy selection for specific service environments
and design requirements (Ref 4).
Reviewing a map of the salinity of rainwater
in the United States is instructive of the degree
to which geography influences corrosion severity. The average atmospheric chloride levels
collected in rainwater are shown in Fig. 5 (Ref
5). The highest levels occur along the coastlines
of the Atlantic and Pacific Oceans and the Gulf
Fig. 2 The decrease in corrosion resistance with increasing
surface roughness by abrasion. Source: Ref 2
216 / Stainless Steels for Design Engineers
Fig. 3 Stainless steel selection expert system. Source: International Molybdenum Association (Ref 3)
Chapter 18: Architecture and Construction / 217
Fig. 4 Grades recommended based on the expert system. Source: International Molybdenum Association (Ref 3)
Fig. 5 Average chloride concentration (mg/L) in rainwater in the United States. Source: Ref 5
of Mexico. The maximum corrosion rate is related to the maximum chloride in the atmosphere. This will be related to the distance inland, the height above sea level, and the
prevailing winds (Ref 6).
Aesthetic Considerations
A correctly chosen grade of stainless steel
will have no degradation over time and, if prop-
erly maintained, will stay new looking indefinitely.
Surface finish aesthetics are arguably more
important architecturally than the influence surface finishes have on corrosion. Numerous finishes have been developed to try to meet various objectives. Finishes vary in reflectivity,
directionality, and subtlety. Figure 6 shows
some of the finishes that go beyond the familiar
brushed look (Ref 7), while Fig. 7 shows special
finishes created by one manufacturer.
218 / Stainless Steels for Design Engineers
Fig. 6 Various rolled-on stainless steel finishes. Source: Ref 7. Courtesy of Outokumpu
Fig. 7 Special finishes for 304/304L and 316/316L stainless steels available from one manufacturer. (a) Rolled-in low-glare finish
(InvariMatte). (b) Rolled-in no. 4 finish (InvariBlend). (c) Rolled-in moderate-glare finish (InvariLux). Source: Contrarian Metal
Resources (Ref 8)
The classic abrasively produced finishes are
No. 3, 4, and 8. These are American Society for
Testing and Materials (ASTM) designations for
abrasively produced finishes, which are traditionally produced by abrading the surface with
different grit size abrasives. Finish No. 3 calls
for 80 to 100 grit abrasive; No. 4 calls for 120 to
150 grit abrasive. Finish No. 8 is a mirror finish
obtained by final polishing with 800 grit abrasive. Finishes No. 3 and 4 are directional, with
grit lines typically 1 cm (0.4 in.) in length. Finish No. 3 has a surface roughness average (Ra)
of 0.4 to 0.8 µm (15 to 30 µin.). Specular gloss
at 85° is typically 40 to 60 (per ASTM D 523,
“Standard Test Method for Specular Gloss”).
The standards for appearance do not exist
within specifications, only the method of producing them. There is considerable difference
in appearance from sheet to sheet, coil to coil,
and manufacturer to manufacturer. The greatest
consistency of appearance comes from specifying a brand of rolled-on finish from a given
manufacturer. Any of the traditional finishes
can be replicated by a rolled-on finish with
greater uniformity, with the possible exception
of bright annealed having a difficult time matching the mirror quality of a No. 8 finish. This is
crucial in architecture, where the discovery of
Chapter 18: Architecture and Construction / 219
unacceptable visual nonuniformity on large
areas can be disastrous, especially when this appears late in the construction process, as is normally the case with exterior components.
Reflectivity or gloss can be a major consideration in the choice of a surface finish. Mirror
finishes are often used for high impact, but
more diffusely reflecting surfaces are more
common. Patterned surfaces provide consistent
reflectivity from a moderately reflective 40 to
60 specular gloss at 85° to a dull matte of less
than 20, the latter having been developed for
airport roofing, such as at Reagan Airport in
Washington, D.C., or the Pittsburgh Convention
Center (Fig. 8).
Flatness is a special consideration for panels
where lack of flatness, such as by “oil-canning,”
can cause a very shoddy appearance. Flatness is
measured in I units.
Flatness (Iunits) = 2 (πH / 2L) × 105
(Eq 2)
where H is the height of the deviation from flatness, and L is the distance between peaks of deviations, assuming a sinusoidal wave.
Because stainless in sheet form is usually reflective, small deviations from flatness can be
very visible. A good standard for flatness that
precludes visible distortion is five I units. Steel
producers have various means to produce this
level of flatness, the most extreme of which is
actually stretching the steel sheet or coil until
all distortions are eliminated. Sometimes,
rather than aiming for high flatness a controlled
deviation from flatness is used, such as slightly
concave panels or panels with a die-pressed
design. Another option to ensure flatness is to
back light-gauge stainless steel with a stiff
material.
Deviations of sheets from squareness and
straightness (camber) are also objectionable because such deviations can cause gaps between
panels. The degree to which this is objectionable is a function of design, and tolerances can
be held tightly at a cost. Width tolerance is
normally +1/16 in./–0 in 48 in., while length is
held to +1/8 in./–0 in 10 ft or less. Maximum
camber is 3/32 in. in 8 ft. Closer tolerances can
be negotiated.
Fig. 8 The Pittsburgh Convention Center with low-gloss finish stainless steel roof
220 / Stainless Steels for Design Engineers
Maintenance and Repair
Maintenance is a significant cost of any structure. One of the great values of stainless steel is
its low cost of ownership. Stainless steel can be
abused, however, and it does benefit from
proper maintenance. The main objective of the
maintenance of stainless is keeping it clean.
There are two reasons for this. The obvious first
reason is that whatever is soiling the surface is
probably not attractive. The second reason is
that it may harm the surface by allowing corrosion agents to concentrate. Table 2 provides recommended practices for removing various substances from stainless steel surfaces (Ref 9).
One of the most common complaints about
maintenance of stainless is the work involved in
removing fingerprints. The oil from fingerprints
makes an easily visible interference film on the
reflective stainless surface. The typical remedy
is to clean stainless steel with a solution containing light oil and a detergent. If the oil from a
hand contacts the uniformly thin film of cleaning oil, no visible mark is left. Alternatively,
Table 2
• Designs that can collect dirt, such as horizontal surfaces and recesses, should be
avoided.
• Designs that create uneven flow or drainage
patterns producing uncleansed areas should
be avoided.
• Sheltered areas and areas subject to splatter,
especially roadside spatter, should be designed so that they are easily cleanable.
Cleaning methods for uncoated stainless steel
Requirement
Routine cleaning of light soiling
Fingerprints
Oil and grease marks
Stubborn spots, stains, and light
discoloration; water marking;
light rust staining
Localized rust stains caused by carbon
steel contamination
Adherent hard water scales and mortar/
cement splashes
Heat tinting or heavy discoloration
Badly neglected surfaces with hardened
accumulated grime deposits
Paint, graffiti
polymer coatings are applied to stainless steel at
some producing mills to permanently provide a
film to which additional fingerprint oil cannot
add a noticeable discoloration.
The greatest ally of stainless on building exteriors is the cleansing action of rain. Rain does
not completely clean the surface, but it does dilute any harmful contaminants and forestall corrosion from accumulated chlorides. Without the
benefit of cleansing by rain, stainless exteriors
should be washed during routine window washing operations.
Given the importance of cleaning, the following design considerations are recommended:
Suggested method(a)
Comments
Soap, detergent, or dilute (1%) ammonia
Satisfactory on most surfaces
solution in warm clean water. Apply with a
clean sponge, soft cloth, or soft-fiber brush,
then rinse in clean water and dry.
Detergent and warm water; alternatively,
Proprietary spray-applied polishes available to
hydrocarbon solvent
clean and minimize re-marking
Hydrocarbon solvent
Alkaline formulations are also available with
surfactant additions.
Mild, nonscratching creams and polishes.
Avoid cleaning pastes with abrasive additions.
Apply with soft cloth or soft sponge; rinse
Cream cleaners are available with soft
off residues with clean water and dry.
calcium carbonate additions. Avoid chloridecontaining solutions.
Proprietary gels or 10% phosphoric acid
Small areas may be treated with a rubbing
solution (followed by ammonia and water
block comprising fine abrasive in a hard rubrinses) or oxalic acid solution (followed by
ber or plastic filler. Carbon steel wool and
water rinses)
pads that have previously been used on
carbon steel should not be used. A test
should be carried out to ensure that the
original surface finish is not damaged.
10–15 vol% solution of phosphoric acid.
Proprietary formulations available with
Use warm, neutralize with diluted ammonia
surfactant additions. Avoid the use of
solution, rinse with clean water and dry
hydrochloric acid-based mortar removers.
Nonscratching cream or polish. Apply
Suitable for most finishes.
with soft cloth or soft sponge. Rinse off
residues with clear water and dry.
Nylon-type pad
Use on brushed and polished finishes along the
grain.
A fine abrasive paste as used for car body
May brighten dull finishes. To avoid a patchy
refinishing. Rinse clean to remove all paste
appearance, the whole surface may need to
material and dry.
be treated.
Proprietary solutions or solvent paint stripper Apply as directed by manufacturer
depending on paint type. Use soft, nylon or
bristle brush on pretreated material.
(a) Cleaning agents should be approved for use under the relevant national environmental regulations and should be prepared and used in accordance with the company’s or supplier’s health and safety instructions. Source: Adapted from Ref 9
Chapter 18: Architecture and Construction / 221
• Contamination by rust from carbon steel is
corrosive and must be removed.
• Designs should facilitate easy access for
cleaning.
While any structure would benefit from these
guidelines, they are especially valuable in maximizing the benefits of stainless steel.
Repair of more severe damage done to surfaces, such as scratching, is difficult. If a decorative surface pattern has been damaged, the
challenge is in trying to replicate it in the field.
Very few surface finishes are wholly repairable
in the sense that they can be repaired in a spot
so that the repair is invisible. The reason for
that is mainly that abrasive finishes are applied
by rotating belts, and it is virtually impossible
to match the pressure, grit size, and arc of contact that created the original surface. It can be
more nearly done to a surface with a rolled-on
finish, which has a very consistent grit length
and depth. The exception to this rule are abrasively applied or rolled-on long-grain finishes.
These have very long grit lines, so grit length
is easily duplicated with a belt sander, the
usual tool available for field repairs. Even
welds can be removed and reblended to be indistinguishable from the surrounding original
surface. The ability to be repaired should be a
top criterion in the choice of a surface finish
whose appearance is critical and that may be
subject to damage.
Fabrication Considerations
Fabrication and joining of stainless steel employ the same techniques as for carbon steel and
other metals. The specifics of cutting, forming,
joining, soldering, and welding are described in
the processing section of this book and are not
repeated here. The main distinction in the use of
stainless steel in this regard is that its higher
strength and corrosion resistance permit the use
of lighter gauges. This in turn permits designs in
hollow or rolled-formed sections, which have
higher stiffness and low weight and potentially
lower overall cost than using less-expensive
metals. A second aspect of higher strength and
lighter gauge is greater spring back in forming
operations, such as press braking.
Certain processing principles related to architectural and building applications of stainless
steels should be emphasized:
• Separation of tools and work areas between
those used for stainless steel and carbon
steel is prudent. Contamination of stainless
surfaces with carbon steel from welding,
grinding, and cutting can stain the surface of
stainless steel and result in corrosion. This
can be remedied by passivation, but it is
much better to avoid it in the first place.
• Welding is better done in the shop than in
the field. Correct filler metals must be used,
and proper weld finishing is essential. The
ability of contactors to produce sound, attractive welds is an indicator of their overall
competence with stainless steel.
• Fasteners used with stainless steel should
also be stainless steel. Galvanized steel, carbon steel, and aluminum will corrode more
readily than the stainless, and this corrosion
is aggravated by galvanic contact. The resulting corrosion products are also harmful
as well as unsightly. Fasteners should not be
permitted to cause distortion of flat panels.
Additional Service Considerations
Fire resistance is an important consideration
in buildings. Stainless steel is the only common building material that remains strong and
tough at temperatures encountered in fires. Ordinary carbon steel undergoes a phase change
at about 760 °C (1400 °F). This change in
atomic structure results in a sudden shrinkage
of more than 1 linear percent. This can literally
pull a building apart. When this occurs to a
structure already weakened by heat, as carbon
steel is, catastrophic failure ensues. Austenitic
stainless steel keeps the same atomic structure
and remains much stronger than carbon steel at
elevated temperatures. Thus, austenitic stainless steel has great value as a material for
structures that must retain structural integrity
in a fire.
Tests have been conducted on glass-reinforced plastic, aluminum, galvanized steel, and
austenitic stainless steel ladders under load and
exposed to flame temperatures of more than
1000 °C (1830 °F). The plastic and aluminum
ladders failed in less than a minute. The galvanized carbon steel lasted 5 min, while the stainless steel remained intact (Ref 10). If the need
for fire resistance is serious, stainless steel
becomes the material of choice. It is used on
offshore oil platforms for stairways, ladders,
walkways, handrails, gratings, floor systems,
222 / Stainless Steels for Design Engineers
Fig. 9 Graphic depicting low release of metal ions from two grades of stainless steel (304 and 316) to the environment, based on a
4-yr multidisciplinary research project involving both field research and laboratory studies. Source: Ref 11
firewalls, blast walls, living modules, and so
forth.
Ecological considerations are never trivial
when considering a construction material.
Many materials used in buildings degrade environmentally, usually by corrosion, and enter the
general environment. Asbestos, lead-based
paints, lead coatings, and others are once-accepted materials whose long-term effects have
been dangerous and costly.
Stainless steel, because it does not corrode
when properly used, does not enter the environment. While this seems obvious, it has been the
object of interdisciplinary studies that have
demonstrated its innocuousness even under
conditions of heavy acid rain on freshly abraded
surfaces (Fig. 9) (Ref 11).
Stainless is a material that will never come
back to haunt an architect years later. Its intrinsic raw material content value ensures that even
with the end of the life of a structure, the stainless in the structure will be recycled.
Concrete Reinforcing Bar
Concrete reinforcing bar is one of the least
glamorous uses of stainless steel in structures.
In bridges, parking garages, and other concrete
structures, saltwater can penetrate the cement
over time. If the internal rebar corrodes, the
expansion of the corrosion products spalls
the concrete, leading to the failure of the
structure.
This can be delayed by treating the concrete
to repel the incursion of water or by coating
the carbon steel rebar with epoxy. These are
less than 100% effective. A more certain approach is to use stainless steel rebar. The stainless steel for this duty need not resist pitting
corrosion, which affects a tiny percentage of
the steel volume. Therefore, an inexpensive,
low-nickel grade, such as 409, 430, or 201, can
be used. Most of the work to date has been
with more expensive grades, such as 316 and
2205. The lean duplexes are ideal for this application because of their high strength, resistance to corrosion and SCC, and moderate cost.
The use of even these alloys reduces the longterm cost of these structures, so the future
adoption of less-expensive stainless steels
holds great promise. The more enlightened
transportation departments in the United Kingdom; Ontario, Canada; and Michigan, New
Jersey, and Oregon in the United States have
led this development.
Chapter 18: Architecture and Construction / 223
REFERENCES
1. Stainless Steel Selection Criteria, Contrarian Metal Resources, www.metalresources.
net, accessed June 2008.
2. Bulletin of the National Dairy Federation
189, 1985, p 3–12
3. C. Houska, “Which Stainless Steel Should
Be Specified for Exterior Applications?” International Molybdenum Association, www.
imoa.info, accessed June 2008
4. The Nickel Institute, www.nickelinstitute.
org, accessed June 2008
5. H. Guttman, Atmospheric and Weathering
Factors in Corrosion Testing, Atmospheric
Corrosion, W.H. Ailor, Ed., John Wiley and
Sons, 1982, p 51
6. R.B. Griffin, Corrosion in Marine Atmospheres, Corrosion: Environments and In-
dustries, Vol 13C, ASM Handbook, ASM
International, 2006, p 42–60
7. Guide to Stainless Steel Finishes, 3rd ed.,
Euro Inox, 2005, www.euro-inox.org, accessed June 2008
8. Contrarian Metal Resources, www.metalresources.net, accessed June 2008
9. Care and Maintenance of Stainless Steels,
Leda-Vannaclip, www.l-v.com.au, accessed
June 2008
10. The Nickel Institute, www.nickelinstitute.
org, accessed June 2008
11. D. Berggren et al., Release of Chromium,
Nickel and Iron from Stainless Steel Exposed Under Atmospheric Conditions and
the Environmental Interaction of These
Metals, European Confederation of Iron and
Steel Industries, Oct 2004, www.eurofer.
org, accessed June 2008
Stainless Steels for Design Engineers
Michael F. McGuire, p 225-232
DOI: 10.1361/ssde2008p225
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
CHAPTER 19
Automotive and Transportation
Applications
Summary
THE ADEQUATE DURABILITY and life
span of cars, trucks, or any transport system
requires freedom from corrosion. This has required subsystems, such as those for exhaust
and fuel, to resist more corrosive environments
for longer periods of time. The main result has
been a strong growth in the use of the leaner
ferritic stainless steels in many components.
As more exotic propulsion systems and fuel,
such as fuel cells and ethanol, emerge, stainless steels may be required to endure the corrosive environments.
the auto manufacturers have become the
largest users of stainless steel.
Exhaust systems constitute the largest use of
stainless steel in the automotive market, but
there are other important applications that cannot be ignored: valves and gaskets, hose
clamps, seat belt and air bag components, tubing, hardware, and filters. And, there will be
new applications that respond to new socioeconomic needs, such as for greater crash worthiness, lighter weight, or resistance to the corrosion of new fuels. But, since exhaust systems
currently predominate, they are covered first.
Exhaust Systems
Introduction
The use of stainless steel in automobiles
used to be mainly a story of decorative applications: wheel covers and trim with a minor
amount used for valves and hose clamps. However, as automobiles became more sophisticated technically and as durability and environmental demands grew, the role of stainless
became increasingly functional and less ornamental. Stainless alloys in common automotive use now are generally highly engineered
for their specific application and represent
some of the most highly evolved applications
engineering in any use of stainless. An examination of the preferred practices in materials
selection in automotive systems is an excellent
example of the rule of using the simplest and
lowest alloy content grade that can do the job.
Because automotive and steel-producing engineers have collaborated so well, both parties
have benefited greatly, as have consumers, and
Laws enacted in the United States in the
1970s mandated automotive emission standards that could be met only with catalytic
converters. The only practical materials that
could withstand the temperatures of the hot
end of an exhaust system using a catalytic converter were stainless steels. From AlleghenyLudlum’s MF-1 evolved a succession of ferritic alloys that grew in sophistication to meet
the increasing needs of corrosion resistance,
oxidation resistance, creep, thermal fatigue resistance, and formability. Soon, entire exhaust
systems were made of stainless; often, they
would last the life of the vehicle, rendering
obsolete an entire muffler replacement industry. The compelling need of the automotive industry for economy drove the widespread
adoption of argon oxygen decarburization
(AOD), the continuous caster, and other highvolume methods of the carbon steel industry.
So, even while the traditional automotive uses
226 / Stainless Steels for Design Engineers
of stainless steels—wheel covers and other ornamental trim—faded to nearly nothing, the
use of stainless in automobiles grew to about
30 kg (65 lb) per vehicle by the turn of the century. The stainless steel industry was changed
from a boutique industry to a mass production
industry by its embracing the needs of the automotive market.
An exhaust system normally consists of a
manifold to collect exhaust gases, a catalytic
converter to reduce NOx and CO emissions, and
a muffler; each of these are connected by piping. Each component of the system has different
requirements for formability, resistance to oxidation, resistance to external corrosion, resistance to internal corrosion, and mechanical
properties. At the outset, it should be noted that
ferritic stainless steels, as opposed to austenitic,
are optimal for oxidation resistance, especially
cyclic oxidation. It is not a difference in the
oxide scale.
The reason is that the thermal expansion of
ferritic stainless more closely matches that of
the oxide scale than does that of austenite. This
prevents the fracturing and spalling of the scale.
The intact scale of ferritic stainless is thus protective up to the temperature at which oxygen
diffusion through the scale becomes great
enough so that “breakaway” oxidation occurs.
At the breakaway temperature level, scale
growth is no longer parabolic with time but becomes linear and therefore no longer protective.
The temperature of this breakaway increases as
chromium content increases. We will see that
other alloying elements can also improve this
performance.
Not only are the steels in exhaust systems almost exclusively ferritic, they are also stabilized by titanium or niobium. This prevents sensitization and makes all the chromium content
useful as alloy. Titanium stabilization greatly
improves corrosion resistance by removing not
just the carbon and nitrogen, but also the oxygen and sulfur from solution. This sharply improves resistance to pitting corrosion. Niobium
is used to costabilize and fight creep. These alloys are therefore essentially interstitial free and
have excellent formability, which the designs of
exhaust system components require.
Their formability is further enhanced, when
necessary, by low additions of tramp substitutional alloying elements such as manganese,
nickel, and copper. Special thermomechanical
processing is also used to optimize texture and
grain size.
The metallurgy of ferritic stainless is discussed in depth in Chapter 8. Rather than reexplain these concepts here, we revisit only the
main points that are relevant to alloy selection
for exhaust systems.
• As chromium level increases, so does resistance to oxidation and corrosion, but
yield strength also increases, and ductility
decreases.
• Alloying with silicon, aluminum, and
molybdenum also increases oxidation resistance, but these elements have the same
detrimental effect on ductility while increasing hot strength.
• Niobium above that needed for stabilization
is a powerful solid solution hardener and is
effective at high temperatures.
• Ferritic stainless steels have very anisotropic
forming properties. They resemble highformability carbon steels in that they tend
not to thin when stretched, which greatly assists in formability.
• Ferritic alloys can form a hard, brittle phase
called α’ in a process commonly called 885
°F (or 475 °C) embrittlement. This is only a
factor in alloys with chromium of 18% or
more, especially those containing molybdenum and aluminum. Cold work accelerates
the formation of this phase.
• The σ phase does not readily occur in alloys
containing less than 20% Cr, so it is not a
consideration for exhaust systems unless silicon or molybdenum are also elevated.
• Coating ferritic steel with aluminum is effective in preventing oxidation and corrosion.
All these factors come into play in the design
of exhaust systems. Because the alloys have
evolved so well to fit the individual requirements for each component, we discuss them segment by segment through the exhaust system.
The exhaust manifold collects the hot, combusted gases from the engine and delivers
them to the front pipe. The exhaust manifold
must possess good high-temperature strength
and resistance to thermal fatigue. It must also
be able to resist oxidation at the exhaust temperature, which can reach 950 (C (1740 (F).
Exhaust manifolds had previously been heavy
castings but are generally now formed from
stamped sheet stainless steel or formed from
welded tubing that may have a double-wall
structure to insulate the gases from heat loss,
which could preclude successful catalytic conversion downstream.
Chapter 19: Automotive and Transportation Applications / 227
As the highest-temperature component of the
exhaust system, the exhaust manifold must possess the greatest resistance to high-temperature
oxidation damage. Risk of such damage is due
to the intermittent use of vehicles, which causes
cyclic oxidation and the ensuing spalling of the
oxide scale. There are numerous alloying approaches for optimizing the ferritic stainless alTable 1
loys for spalling resistance. All approaches involve raising chromium content but use different techniques to enhance the effect of
chromium. Table 1 lists a number of the grades
of stainless steel commonly used in exhaust systems and where they are used. The alloys are
listed in order of increasing severity of the requirements for each major system component.
Alloys normally used for the major elements of automotive exhaust systems
Service temperature
Component
Requirements
Alloys currently used, common name
(related designation)
High-temperature
•
strength, thermal
•
fatigue strength,
•
oxidation resistance, •
formability
•
•
•
•
•
•
•
T439HP (UNS S43035, dual-stabilized 439)
18CrCb (DIN 1.4509, 18CrCb)
441 (DIN 1.4509)
304/304L/304H (UNS S30400, S30403,
S30409)
321 (UNS S32100)
309S (UNS S30908)
310S (UNS S31008)
332Mo (S35125)
600 (N06600)
601 (N06601)
625 (N06625)
High-temperature
•
strength, thermal
•
fatigue strength,
•
oxidation resistance, •
formability
•
•
•
High-temperature
•
strength, thermal
•
fatigue strength,
•
oxidation resistance, •
formability, salt
•
attack resistance
•
High-temperature
•
strength, salt
•
attack resistance,
•
formability
•
•
•
409 ALMZ (aluminized 409)
T439HP (UNS S43035, dual-stabilized 439)
18CrCb (DIN 1.4509, 18CrCb)
441 (DIN 1.4509)
436S (type 436S)
444 (UNS S44400, T441)
433 (T443)
304/304L (UNS S30400, S30403)
T321 (S32100)
316/316L (S31600/S31603)
316Ti (S31635)
332Mo (S35125)
625 (N06625)
409HP (UNS S40930, dual-stabilized 409)
T439HP (UNS S43035, dual-stabilized 439)
441 (DIN 1.4509)
18CrCb (DIN 1.4509, 18CrCb)
444 (UNS S44400, T441)
433 (T443)
Exhaust manifold
ºC
750–950
ºF
1380–1740
Front pipe
600–800
1110–1470
Flexible pipe
600–800
1110–1470
Catalytic converter shell
600–800
1110–1470
Catalytic converter
substrate
1000–1200
1830–2190
Oxidation
resistance,
thermal shock
resistance
• ALFA-IV (FeCrAl)
Center pipe
400–600
750–1110
Salt damage
resistance
Muffler
100–400
210–750
Corrosion
resistance,
from inner and
outer surface
•
•
•
•
•
•
•
•
•
•
•
•
•
•
Tailpipe
100–400
210–750
Corrosion
resistance,
from inner
and outer surface
Source: Adapted from Allegheny Technologies Inc.
409HP (UNS S40930, dual-stabilized T409)
409 ALMZ (aluminized 409)
T439HP (UNS S43035, dual-stabilized T439)
441 (DIN 1.4509)
18CrCb (DIN 1.4509, 18CrCb)
444 (UNS S44400, T441)
433 (T443)
409HP (UNS S40930, dual-stabilized T409)
409 ALMZ (aluminized 409)
T439HP (UNS S43035, dual-stabilized T439)
436S (T436S)
441 (DIN 1.4509)
18CrCb (DIN 1.4509, 18CrCb)
Type 304/304L (UNS S30400, S30403)
• 409HP (UNS S40930, dual-stabilized T409)
• 409 ALMZ (aluminized 409)
228 / Stainless Steels for Design Engineers
The best choice for a given design is not obvious. We attempt to simplify the choices.
Thus, the basic alternatives for exhaust system alloys are:
• Straight chromium alloying at 11 to 12%
with stabilization by titanium or niobium,
the basic type 409 (UNS S40920)
• Straight chromium alloying at 17 to 18%
with stabilization by titanium or niobium,
the basic type 439 (UNS S43036)
Either of these basic alloys can enjoy enhanced oxidation resistance by additional alloying with molybdenum, aluminum, or silicon. In
addition, they can be coated with hot-dipped
aluminum-silicon alloy to increase oxidation resistance.
Chromium or molybdenum alloy additions
increase corrosion resistance, whereas aluminum or silicon additions do not improve that
trait. Aluminum coating is a powerful corrosion
fighter, and it has the aesthetic benefit of not
showing red rust.
Use of molybdenum or niobium enhances
high-temperature strength. Alloys with these
additions are thus useful for manifolds with a
design that constrains expansion and contraction, making thermal fatigue a problem.
All alloying additions detract from formability and toughness, as well adding to basic material costs. Thus, the objective must be to use
only those alloying elements that are indispensable to performance.
The front pipe connects the exhaust manifold
to the flexible joint and experiences nearly the
same temperatures as the exhaust manifold, but
not the same risk of thermal fatigue. To reduce
exhaust noise, a double-wall pipe is sometimes
used for this component.
The flexible joint is the one segment of the
exhaust system for which austenitic stainless
steels are preferred. The function of the flexible joint is to prevent vibration from the engine from being transmitted to the rest of the
exhaust system. It consists of a double-wall
pipe in a bellows configuration with an outer
covering of braided stainless steel wire. It must
have very good high-temperature fatigue
strength to withstand the cyclic stress of the vibration it absorbs. The material used must
have exceptional formability to be formed into
a bellows. The greater hot strength and formability of austenitic steels thus prevails. The
flexible joint is also exposed to road salt in
some regions, so it must resist hot salt corrosion. This may force the use of 316L versus the
normal choice of 304L.
The catalytic converter is the next component of the exhaust system. It exposes the exhaust gases to noble metal catalysts, which
complete the combustion of the gases to form
less-noxious compounds. This is an exothermic
reaction at temperatures equal to those in the
exhaust manifold. Thus, the housing, while not
requiring great hot strength, must resist oxidation. The catalyst itself is supported by a ceramic and ferritic stainless steel carrier that
must resist thermal shock and possess low heat
capacity for rapid heating. Exotic alloys of
20% Cr with 5% Al are used for the carriers.
The housing is generally made of a 17% Cr ferritic stainless. The converter is usually directly
beneath the passenger compartment, so a heat
shield of type 409 is used to separate it from
the floor.
The center pipe conveys the converted gases
to the muffler. The cooling exhaust gases no
longer present a major oxidation threat, but the
condensing water vapor creates an internal corrosion risk, and road salt presents an external
one. However, a simple grade such as 409
should generally provide sufficient resistance to
this environment.
The muffler, next in line, presents only a corrosion issue. The muffler must withstand corrosion from the outside, the worst of which comes
from road salt or coastal salt sources. Internal
corrosion is also a major consideration because
condensing exhaust gases create a hostile,
acidic environment. After startup, the heating of
the muffler to temperatures above 100 °C (212
°C) evaporates these condensates, and internal
corrosion ebbs. On short runs, this may not
occur. This represents a worst case for internal
corrosion. The dual internal and external corrosive attacks require the use of aluminized stainless for best performance.
The tailpipe is exposed to view in most vehicles, and its appearance is therefore important.
For this reason, an austenitic such as 304 can be
used, as can chromium plating or aluminizing.
The object here is to avoid visible corrosion.
Truck exhaust systems are beginning to require similar technical sophistication as their
emissions come under increased regulation.
However, they do not present any challenges
not already confronted and solved for passenger
vehicles.
Chapter 19: Automotive and Transportation Applications / 229
Structural Components
The driving forces of durability, safety, and
weight reduction have spawned other, more varied applications for stainless in automotive engineering. Across the board, the main distinguishing trait of stainless that qualifies it as the
optimal material is its corrosion resistance, but
this characteristic would be insufficient in many
cases without considering mechanical properties. Indeed, even if stainless steel were not corrosion resistant, its superior strength and toughness would qualify it for many automotive
applications.
Austenitic stainless steels are the toughest and
stiffest practical materials available to the automotive engineers. Common 301 can be cold
worked to yield strengths anywhere from its annealed level of about 300 MPa (44 ksi) up to
2000 MPa (290 ksi). In this higher-strength condition, it has become the standard material for
seat belt anchors and hose clamps. Type 301 in
the annealed condition is actually the original
transformation induced plasticity (TRIP) steel as
it can be tailored to have a controlled level of
austenite stability. This allows it to transform at
a known rate to martensite during deformation,
giving not only a very high work-hardening rate
but extraordinary resistance to localized thinning, necking, and therefore fracture. When
crash worthiness becomes a prime consideration, then this characteristic makes 301, or its
low-nickel counterpart 201, an ideal material for
structural, energy-absorbing components since
austenitic stainless can be rivaled for such applications only by heat-treated alloy steel, titanium,
or aircraft aluminum alloys, all of which are
more expensive, less durable, or less formable.
Tables 2 and 3 show the properties of specific
variations on basic 301 developed by Outokumpu and how they stack up against the
most competitive carbon steels, dual-phase
steels, and TRIP steels (Ref 1). The value of a
material as an energy-absorbing structure (i.e.,
one that enhances crash worthiness) is measured by the energy it can absorb per unit of
mass. The kinetic energy of a collision that a
structure can absorb in deformation is proportional to its strength multiplied by the amount it
can deform before fracturing. The superiority
of metastable stainless steels (i.e., 201 and 301,
those that most easily transform to martensite
during deformation) is shown in Fig. 1. Even
with its lower density, aluminum falls far short
of austenitic stainless in energy absorption per
unit weight.
Table 2 Comparison of tensile properties of carbon steels and stainless steels for automobile
structural components
Type
Thickness, mm
0.2% proof
strength, MPa
Ultimate tensile
strength, MPa
True stress at
ultimate tensile
strength, MPa
Uniform
elongation, %
Total elongation, %
Carbon steels
TRIP 700
DP 750
DP 800
1.58
1.48
1.44
473
513
573
703
811
896
818
920
976
16.4
13.4
8.9
17
18.8
9.9
1.16
1.55
306
639
937
1068
1429
1377
52.5
28.9
59.3
38.6
Austenitic stainless steels
HyTens X
HT 1000
Source: Ref 1
Table 3 Comparison of resilience and toughness of carbon steels and stainless steels for
automobile structural components
Type
Resilience, J/m3
Toughness, j/m3
0.996
1.131
1.32
105
101
74
0.536
1.726
364
269
Carbon steels
TRIP 700
DP 750
DP 800
Austenitic stainless steels
HyTens X
HyTens 1000
Source: Ref 1
230 / Stainless Steels for Design Engineers
Fig. 1
True stress-true strain curves for 301 variants (HyTens X and HyTens 1000) versus two duplex carbon steels (DP750 and
DP800) and a transformation induced plasticity (TRIP) steel (TRIP700). Source: Ref 1
These exceptional strength-to-weight and energy absorption-to-weight characteristics permit
automotive engineers to reduce weight and increase crash worthiness while designing vehicles with greater life span—because corrosion
resistance “comes along for the ride,” as it were.
Some components in which these virtues are
most readily exploited are bumper systems (Ref
2). Porsche uses austenitic stainless steel for
front and rear side members, internal push rods
on front and rear axles, and lower rear wishbones in its Carrera GT.
Another manufacturer, Audi, engineered various components of austenitic stainless steel into
its otherwise aluminum-intensive A6 series. The
use of stainless steel in strategic components
enables greater weight reduction than that
which the vehicle would have in all aluminum.
Volvo and Saab have designed austenitic stainless steel bumper systems that also serve to reduce overall vehicle weight.
While it is probably apparent to the reader that
essentially any body component can be made in
stainless and be made better in stainless, the
question of when doing so is a better engineering decision involves economic considerations.
Large automotive companies generally have
large fixed investments in painting and coating
systems to protect entire bodies from corrosion.
The incremental savings of eliminating coatings
on individual body components is thus essentially nil. However, if the entire system is stainless and the investment is avoided, then the initial cost of a stainless body actually can be lower
than one in coated carbon steel. This is the experience of Italian bus manufacturers, who began
in the 1980s using 304 stainless steel in buses.
Now, buses are 80% stainless.
Designers began the conversion to gain the
normal advantages the stronger stainless gives:
over 10% lighter weight and over 10% improvement in crash worthiness of the passenger
compartment, the accompanying savings in fuel
consumption, and the virtual elimination of
body maintenance. With essentially the entire
body now in stainless, coating and painting
could be eliminated. A stainless bus body is
shown in Fig. 2.
This swung the economic pendulum to stainless in a major way. Now, not only was the
long-term cost of operating the bus lower, but
the initial cost of the bus was lower. The economic analysis is shown in Table 4 (Ref 2).
The design key was to use rectangular 304
stainless structural tubing, which allowed strong,
stiff sections to be welded into space frames. It is
only a matter of time until this is improved on by
the use of 201 (with 3 to 4% Ni instead of 8 to
9% Ni) to lower cost and cold working of the
tubing to achieve higher strength levels.
Chapter 19: Automotive and Transportation Applications / 231
Microcars are now a familiar sight in Europe.
These vehicles are prized for their ability to be
driven and parked in very small or congested locations. Their economy of operation is also a
major attraction. These considerations combine
to make stainless the best material for many of
their components. Figure 3 shows a stainless
steel microcar frame.
The design by the famous design house Pininfarina employs a stainless frame to give maximum torsional stiffness and crashworthiness
while eliminating painting entirely.
Other Automotive Components
Stringent emissions controls regulations, led
in the United States by the state of California,
have made manufacturers reexamine the suitability of polymeric fuel tanks. These tanks contribute more to the required maximum 2 g/day
of hydrocarbon emissions than is tolerable, so a
Fig. 2
Stainless steel bus bodies. Source: Ref 2
few manufacturers, such as Volkswagen, have
installed stainless steel fuel tanks in their vehicles (Ref 2).
Trucks
Over-the-road trailers are an excellent example of stainless steel being used for utilitarian
purposes. Trailers used for hauling foodstuffs or
corrosive materials are now constructed almost
entirely of stainless steel; lined carbon steel
tanks are now largely obsolete. The engineering
basis for this is the same as for buses: high
strength, no coating costs, and a product with
long life and low maintenance costs. Structural
members in trailers are typically 304, while
tanks may also be 316L for corrosion resistance
when the transported material requires it. Tank
wrappers are often made of bright annealed and
buffed 304. Manufacturers of trailers would be
well advised to consider upgrading to duplex
grades such as 2003 or 2205 or to cold-worked
austenitic stainless, which would permit major
weight reduction. This weight reduction would
directly translate into greater load-carrying capacity because the payload of liquid-carrying
trailers is limited by total gross weight. The
ability to add a few thousand more pounds of
payload would quickly pay back a small premium in material cost.
Normal cargo-carrying trailers also use some
stainless where corrosion is problematic, such
as in doors and door frames. Weight reduction is
less important in these trailers, which reach
maximum load at a volume limit rather than a
weight maximum.
Table 4 Life-cycle cost calculation (LCC) for stainless steel versus carbon steel
for a bus application
Cost of capital
Inflation rate
Real interest rate
Desired LCC duration
Downtime per maintenance/replacement event
Monetary unit
Value of lost production
Material costs
Fabrication costs
Other installation costs
Total initial costs
Maintenance costs
Replacement costs
Lost Production
Material-related costs
Total operating cost
Total LCC cost
Source: Ref 2
Stainless steel
3.331
25.322
2.185
30.838
0
0
0
0
0
30.838
10.00%
5.00%
4.76%
20.0 years
1.0 day
U.S. $
101 U.S. $/day
Carbon steel
1.391
28.582
4.050
32.023
1.448
2.897
57
0
4.402
36.425
232 / Stainless Steels for Design Engineers
Fig. 3
Microcar frame fabricated from stainless steel. Source: Ref 2
Rail Transport
Passenger trains have exploited the high
strength-to-weight and toughness qualities of
the 301 family of stainless steels for many
years. The corrosion resistance of these alloys
makes them corrosion free in long use, obviating the need for painting and lowering maintenance costs. As with any other major use of type
301, a 5 to 10% increase in economy could be
achieved if type 201 were used instead of 301.
No loss in performance would occur. The transition to 201 has not occurred simply because of
inertia and resistance to change on the part of
designers and producers.
Hopper cars made of 12% Cr martensitic
stainless steels, typically 409Ni and 3Crl2, have
excellent abrasion and corrosion resistance as
well as high strength and therefore greater loadcarrying capacity. Curiously, the use of coldworked austenitic stainless in railcars, which has
been successful for decades, has not been carried
over into trucks and buses even though it is technically feasible to economically produce structural sections in very high-strength stainless.
REFERENCES
1. R. Andersson, E. Schedin, C. Magnusson,
J. Ocklund, and A. Persson, The Applicability of Stainless Steels for Crash Absorbing
Components, ACOM, No. 3–4, AvestaPolarit AB, 2002
2. F. Capelli, V. Boneschi, and P. Viganò,
“Stainless Steel: A New Structural Automotive Material, Vehicle Architectures: Evolution Towards Improved Safety, Low-Weight,
Ergonomics, and Flexibility,” paper
presented at Florence ATA 2005, 9th International Conference (Florence), May 2005,
www.centroinox.it, accessed June 2008
Stainless Steels for Design Engineers
Michael F. McGuire, p 233-242
DOI: 10.1361/ssde2008p233
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
CHAPTER 20
Commercial and Residential
Applications
Summary
STAINLESS STEEL HAS BECOME the essential material for products related to food,
health care, and laundry because it combines
strength and durability with an unexcelled ability to be cleaned, disinfected, and sterilized.
These qualities have long been apparent to commercial food, laundry, and health care professionals and have increasingly carried over into
equivalent domestic areas as consumers have
become more aware of the benefits of stainless.
Introduction
The last 20 years have seen the long-standing
pervasive commercial use of stainless steel for
food preparation and serving; laundry; heating,
ventilation, and air conditioning (HVAC); and
other appliances penetrate the domestic market
for the same types of goods. Whether this is a
fad of an increasingly affluent consumer or a reflection of more design engineers and consumers being more interested in lasting value
than they were in the “throwaway” society that
preceded that period remains to be seen. Stainless has been increasingly identified with highquality, high-end products. But, the case for
value rather than fad seems to be stronger if the
lessons of the harshly pragmatic automotive industry, in which decorative use of stainless has
virtually disappeared while utilitarian uses have
mushroomed, are any indication.
The case for using stainless in appliances of
all types, whether they are commercial or residential, relates to stainless being able to provide
the best value over the intended service life.
Stainless is without rival for ruggedness and
durability. Steel and aluminum corrode. Glass,
stone, and ceramics break. Plastic is weak. The
second, even more important, reason is that
stainless steel is essentially benign from a hygienic viewpoint. Stainless steel itself is inert,
both chemically and biologically, with respect
to food. Further, it provides minimal harbor for
unwanted biologic growth as do more porous
materials. Stainless also competes quite well esthetically with other materials, offering the designer numerous surface finishes. Last, stainless
is very amenable to nearly all manufacturing
techniques. Its lack of need of coatings often
makes components made from stainless less expensive to produce than equivalent designs that
must be coated with paint, porcelain, or metal.
Food Contact Qualifications
Setting aside cost, esthetics, and manufacturing considerations, a food contact material must
first meet three criteria: It must be chemically
inert, biologically inert, and cleanable and able
to be disinfected.
Chemical neutrality is achieved by a material when the material does not enter into the
food with which it comes in contact. This has
become an increasing concern as the effects of
ions or chemicals released from food preparation materials have been viewed as potential
toxic or disease agents. Medical knowledge is
not sufficiently advanced to convince consumers of the harmlessness of such contaminants, so it is preferable to demonstrate the absence of contamination if one is to win the
public confidence in a food contact material.
234 / Stainless Steels for Design Engineers
Stainless steel contains many constituent elements. Were they to enter the food with which
the stainless came in contact, then stainless
would be a poor food contact material. The distinguishing characteristic of stainless, however,
is the spontaneous passive film, which is so stable chemically. This film acts as a barrier to corrosion, which would result in metal release.
Stainless therefore is effectively inert. Tests
have been made of the rates at which metal ions
can enter foodstuffs (Ref 1). Table 1 shows the
vastly lower rates of metal ion release from
stainless than from aluminum and carbon steel,
both of which are permissible, if not optimal,
food contact materials. Aluminum releases aluminum ions into solution of both cooking oil
and 3% acetic acid at nearly equal rates of
15 mg/cm3 in 30 days. Carbon steel releases
iron at over 100 mg/cm3 in the same period.
Stainless, however, releases less than 0.010
mg/cm3 of iron. This is complemented by other
studies showing that the transfer of ions from
food contact vessel to food is diminishingly and
negligibly small. Tests have been conducted on
stainless steels, types 304, 439, and 444, that
had both industrial finishes (2B and BA) as well
as freshly abraded and air-aged finishes. These
sample steels were subjected to boiling solutions of oils, alcohols, water, and 3% acetic
acid. None caused the transfer of either
chromium or nickel to exceed the statutory 0.1
ppm level (Ref 2). Nickel levels of various
foods before and after cooking have been scientifically measured to assess the possibility of
leaching of that ion from 304 stainless steel
(Ref 2). No increase was noted from the natu-
rally occurring level of up to 0.3 μg/g in cereals
to 1.1 μg/g in meat and fish.
These negligible levels of leaching simply indicate that foodstuffs are a benign chemical to
stainless steel. Nevertheless, it is necessary to
apply the correct assessment of the corrosivity
of the foodstuff in question. In food production,
as opposed to preparation for serving, more extreme levels of acidity and salinity can be encountered. Nippon Steel reported (Ref 3) that
materials used in the manufacture of soy sauce,
which can have 15% salt, must withstand prolonged contact at 45 °C (115 °F). Under such
conditions, 316 stainless pits in about 1 day,
while higher alloy grades, the 6Mo alloys, of
which their YUS 270 is one (equivalent to UNS
S31254), are projected to last 20 years before
pitting.
This is significant because only pitting corrosion is likely to release metals ions into a food
substance. So, while guidelines exist for the
minimum alloy content permissible for normal
food contact, such as those promulgated by the
National Science Foundation (NSF), one must
still verify the corrosion due to a particularly
aggressive food ingredient. Choosing the proper
grade of stainless, based on pH, salinity, and
temperature, is the responsibility of the design
engineer. Referred to chapters in this book on
both corrosion and individual alloy families for
guidance in choosing an alloy based on resisting
pitting in a given environment. This having
been said, no alloy greater in pitting resistance
than 304 is required in residential or commercial cooking food contact. The higher alloy
requirements come from the more aggressive
Table 1 Net metal migration into acetic solution (3%)
Metal migration during indicated time, µg/cm3
Material
Austenitic stainless
Ferritic stainless
Aluminum
Carbon steel
Source: Ref 1
Time
Iron
Chromium
Aluminum
30 min
10 days
20 days
30 days
30 min
10 days
20 days
30 days
30 min
10 days
20 days
30 days
30 min
10 days
20 days
30 days
2.4
4.2
2.7
2.3
3.0
7.3
8.6
6.6
4.9
18.2
17.9
31.3
8,430
57,700
62,900
112,000
0.12
0.22
0.22
0.28
0.43
0.40
0.71
0.87
0.93
3.42
5.58
12.40
0.62
7.40
6.82
14.00
≤ 0.19
0.22
0.19
≤ 0.19
≤ 0.19
0.19
≤ 0.19
≤ 0.19
930
5,300
7,160
15,350
2.7
26.7
24.0
36.9
Nickel
≤ 0.12
≤ 0.12
≤ 0.12
0.31
≤ 0.12
≤ 0.12
≤ 0.12
≤ 0.12
≤ 0.12
≤ 0.12
≤ 0.12
0.22
≤ 0.12
≤ 0.12
≤ 0.12
≤ 0.12
Chapter 20: Commercial and Residential Applications / 235
concentrations and exposure periods that can be
found in food-processing plants.
Certain abuses can even damage stainless
cookware. Very high temperatures, such as can
occur when unattended pans have their liquids
boiled away, can damage stainless but are more
harmful to less-rugged alloys such as copper
and aluminum.
Stainless steel is primarily composed of iron,
chromium, and nickel along with small amounts
of manganese, silicon, and molybdenum. It contains trace amounts of copper, aluminum phosphorus, and sulfur. Each of these elements is
naturally occurring in food. Each can be found
in a typical multivitamin/multimineral supplement. Stainless is essentially devoid of heavy
metals, such as lead and mercury, which are vaporized at the temperatures at which stainless is
refined. Even if toxic metals were somehow to
be made to contaminate stainless, the passive
film would prevent their release. All these factors combine to make stainless the most chemically neutral metal found in food contact. Alternative alloys, such as copper and aluminum,
actively leach into foods. Copper and aluminum
have been linked to but not demonstrated to
cause Alzheimer’s disease.
Biological Neutrality. Microorganisms adhere to solid surfaces. When a clean surface
comes in contact with food, a surface deposit is
formed from the food. The film may also contain molecules left from previous cleaning and
disinfecting. The formation of this film is presumably influenced by material characteristics
such as roughness, although there are no specific studies on this. However, microorganisms
adhere to this film and, as colonies of them
grow, form a biofilm. This film consists of layers of microorganisms that can produce an exocellular polymeric matrix, which protects the
colony from cleaning and disinfecting. Geometric factors also can protect these colonies.
Rough surfaces are intuitively more difficult to
clean. The ability to maintain a microscopically
smooth surface is an asset in stainless that polymeric, enamel, and mineral surfaces lack. Stainless steel is much less roughened by abrasion,
keeping the surface smooth (Ref 4).This will be
seen to influence its ability to be cleaned and
disinfected.
There is some technology to go beyond biological neutrality in the use of coatings that actively discourage or eliminate growth of microorganisms. Polymeric coatings impregnated
with silver ions have been developed and com-
mercialized (Ref 5). Silver ions, like copper
ions, are powerful antimicrobial agents. The
combination of such a coating with stainless as
a corrosion-proof substrate may represent the
maximum in hygienic and chemical protection
and is already being used in medical applications where such concerns exceed those in ordinary food contact situations.
Cleanliness. A necessary quality in any material considered for food contact is the ability to
be cleaned. This includes the removal of both
organic and inorganic substances. The most important objective of cleaning is to remove the
visible and invisible materials that can provide a
growth medium for microorganisms. This
process is distinguished from disinfection,
which is the reduction of the microbial population to a satisfactory level. What this level is depends on the standards of hygiene in force. And,
although cleaning can and does reduce the population of microorganisms, true bacteriological
cleanliness is obtained only after disinfection.
The combination of cleaning and disinfecting is
important. Studies have shown that the efficacy
of disinfectants is weaker on bacteria that have
been established in a surface biofilm than on
bacteria in suspension. The most complete form
of disinfecting is sterilization, whose objective
is the complete removal of all microbial life and
viruses.
The purpose of cleaning stainless steel is to
rid it of contamination. Various stainless manufacturers and associations have identified a
number of effective of cleaning products (Ref 2,
6, 7):
• Alkalines, which dissolve fats and oils
• Chelating or sequestering agents, which agglomerate contaminants. These are often organic acids such as citric acid or oxalic acid
and amine acids such as sulfamic acid and
ethylene diamine tetraacetic acid (EDTA) or
salts of these compounds.
• Hydrocarbon solvents
• Water with soap, detergent, trisodium phosphate, or other surface active agents, which
emulsify
• Dilute oxidizing acids like nitric acid
• Mild acids such as phosphoric acid
The effectiveness of a cleaner relates mainly
to the contaminant to be removed. Some trial
and error may be required for a given contaminant. Some precautions are worth mentioning.
Abrasive cleaners should be used with caution.
The abrasive size and hardness must be chosen
236 / Stainless Steels for Design Engineers
so that the stainless surface finish is not affected
in an unwanted manner. If the abrasive is harder
than the stainless or coarser than the stainless
surface roughness, the underlying finish can be
disturbed. Care should also be taken to clean
with the polish grain if a polished surface is
being cleaned. Also, cleaners containing chlorides are common. Their use is not recommended on stainless. Use of hydrochloric
(muriatic) acid is especially detrimental. If
chloride-containing cleaners are used, then thorough rinsing should be conducted to avoid
chloride concentration through evaporation, especially in crevices. Steel wool or steel brushes
should not be used on stainless under any circumstances as iron residue interferes with the
integrity of the passive film.
The ability of stainless to be cleaned is best
measured by the actual removal of bacteria
colonies. This has been done to compare
unabraded and abraded (to simulate new vs.
used) stainless steel, enameled steel, mineral
resin, and polycarbonate materials, which can
be used for sinks, counters, food prep tables,
etc. (Ref 4). Figure 1 shows that the reduction in
bacteria count by the same cleaning technique is
ten times more effective on stainless than on the
other material types. Abrasion did not degrade
the ability of stainless to be cleaned as it did
softer materials. The surface of stainless, even
with the seemingly protected recesses due to
abrasive polishing, permits bacterial colonies to
be removed. The greater roughness of the other
materials may serve to protect the bacterial
colonies from shear forces and provide greater
specific surface area on which the colonies can
bond.
Disinfection. The ability of a surface to be
disinfected is measured by the concentration of
a given disinfectant required for a specific reduction in bacterial population. Numerous
studies have been published (e.g., Ref 8–10)
showing that glass and stainless steel have equal
aptitude for disinfection, and that polyesters,
polyurethanes, rubber, and aluminum all required about one to two orders of magnitude
greater concentrations of disinfectant for the
same result. These results indicate why stainless
is so essential to the food industry. Stainless can
be disinfected quite readily, which allows the
great invisible liability of food-borne diseases
to be minimized.
The effectiveness of sodium hypochlorite as a
disinfectant is inarguable, also. So, despite its
potential corrosivity, it will be commonly used.
Taking this into account requires that commercial and residential food equipment be able to
withstand some chloride level greater than otherwise projected. Industry practice in the United
States has shown that corrosion problems occur
at an unsatisfactory level with mechanically
polished 430 but not with bright-annealed 430.
Fig. 1 Bacterial retention as a function of material and cleaning time. Source: Ref 4
Chapter 20: Commercial and Residential Applications / 237
Thus, alloys with less than 16% Cr should not
be used unless corrosion can be accepted. Alloys containing 16% Cr can be used with optimal surface finish. Alloys as low in carbon as
12% are used for cutlery applications where
slight corrosive attack can be accepted. This is a
necessary trade-off required to achieve high
hardness for good cutting edge retention.
Higher chromium grades such as 304 can be
used even with mechanically polished surface
finishes. From a cost-effectiveness point of
view, there is no reason to use more expensive
alloys than 430 or 201 in the vast majority of
commercial and residential kitchen and laundry
applications from a corrosion standpoint as long
as surface finishes that have not been produced
by abrasive polishing are specified. Many such
finishes are widely used. In North America, the
rolled-on replicas of No. 4 finish, Koolline,
Lustrite, etc., are quite common, while in Europe the bright-annealed finish has been preferred. Both of these are preferable to mechanically polished finishes.
The food industry, an immense consumer of
stainless steel, could do more than any other industry to help conserve nickel by specifying alloys such as 430, 439, and 201 as their standard
alloys as well as by specifying nonabrasive finishes. This can be done with no loss of functionality or change of appearance and could save
23% to 50% in material cost.
Applications
Cookware. Any interaction between a food
contact material and the food is most likely to
occur during the cooking process when temperatures are greatest. Only glass and stainless are
excellent food contact materials. And, since
cookware must be flexible enough to handle any
potential food, the choice of material for cookware must be the most conservative. For this
reason and because of the brittleness of glass,
stainless is the material of choice. The qualities
discussed make aluminum and copper less desirable. Both leach into food. Copper can be
tinned to combat this. The tin also corrodes over
time but has very low toxicity. A larger drawback is the expense of retinning copper utensils.
Aluminum is known as a toxic metal, with its
toxicity causing symptoms similar to those of
Alzheimer’s and osteoporosis (Ref 11). These
two metals do have one advantage over stainless, however: their thermal conductivity. High
thermal conductivity in a cooking utensil minimizes differences in temperature across the surface in contact with the food, permitting better
control of the cooking process. The solution to
the problem of thermal conductivity is to make
composite materials. Stainless can be bonded to
copper and aluminum, which allows the stainless to be on both the food contact surface as
well as the exterior, with an inner layer of copper or aluminum effectively spreading the heat.
Aluminum and copper are nearly equally effective as inner conductive layers. Premium cookware features them both. The “sandwich” is the
optimal design because it optimizes heating uniformity even more than using aluminum or copper alone would since the high conductivity
inner core functions as an isotherm. The uniformity is the more important consideration than
the absolute thermal conductivity or even the
thermal diffusivity.
In a triple layer, the choice of the non-foodcontacting stainless is less stringent. Sometimes, the exterior is made of a ferritic stainless
steel. The ferromagnetism of ferritic stainless
steel makes it ideal for induction heating. Alloys such as 436 have been used for this application, while 304 is the pervasive choice for the
food contact surface. This is despite the fact that
201 or 301 are quite adequate for this application. It is also possible to produce a magnetic
carbon steel core with stainless bonded to both
sheet surfaces. The exposed edges are rolled to
shield them from corrosion.
Nonstick coatings, such as polytetrafluorethylene (PTFE), are very popular because of their
nonstick qualities. Above 350 °C (660 °F) these
coatings give off toxic fumes. This is a danger
for certain types of cooking, such as wok cooking or blackening, but more likely to be encountered by accidentally high temperatures above
those intended. Since they can be scratched and
are not impermeable, their use does not alter the
choice of the material to which they are applied.
Kitchen Appliances. Every type of commercial kitchen appliance can be, and usually is,
made of stainless steel, as are premium domestic kitchen appliances. This choice is based on
durability and ease of cleaning and disinfecting.
And, because many commercial appliances are
visible to the customer, aesthetics are also a
driving force. Choice of alloy for a given appliance is a crucial cost factor. As was noted that
430 is marginal for kitchen use, because of the
prevalence of chloride-containing cleaners, unless it has been bright annealed. All austenitics
238 / Stainless Steels for Design Engineers
are satisfactory under normal use. Designers
seem to generally neglect the possibility that
their equipment may be used in coastal climates. In the high ambient salinity of coastal
climates, corrosion will occur unless 304 with a
bright-annealed finish or a brushed finish rolled
onto a bright-annealed 304 is used. Mechanically polished 304 stainless is inadequate for
coastal environments. These are the same
guidelines used for architectural applications.
Figure 2 shows how different alloys withstand
coastal conditions. The corrosion on 430 would
be considered excessive, while that on the 201
and 304 is acceptable given that some routine
cleaning would have prevented the corrosion
that is present on these samples, which were exposed to coastal salt and humidity for 10 years
in North Carolina (Ref 12).
In the vast majority of ambient conditions,
coastal salinity is not a problem. This applies to
inland conditions or coastal conditions where
interior environments are protected by air conditioning or adequate cleaning of the stainless is
practiced. This is normally the case for commercial equipment. Under these conditions, 201
is quite adequate, and the use of 304 represents
wasteful overengineering. This choice is supported by decades of use by the major manufacturers of commercial appliances. Many who are
large enough to specify their desired grade on
bills of materials rather than simply buying
from service center inventories have routinely
used 201 and realized an approximately 8%
lower cost before surcharges. Use of 201 versus
304 reduces surcharges by almost 50%, which
can be a much larger savings than the base price
savings. Smaller manufacturers are often precluded from these savings because of the general, if inexplicable, practice of service centers
not stocking 201 despite its being the most costeffective general-purpose stainless grade. The
extended nickel price elevation from 2004 onward has a good chance of changing that situation as end users rebel against surcharges,
which cannot be passed on to their customers. It
has been pointed out that there is an array of
201-type grades, and that this is a drawback to
their wider adoption.
I recommend following American Society for
Testing and Materials (ASTM) A240 and specifying UNS S20100 when substituting for 304 as
this has very similar performance in forming,
welding, and appearance to 304 and can be
most easily interchanged without complications
in manufacturing and field performance. For
Fig. 2
Stainless steel samples exposed on a North Carolina
beach for 10 yr. Source: Ref 12
Chapter 20: Commercial and Residential Applications / 239
parts made by deep drawing, substitution is still
very possible, but deep-drawn grades are more
finely tuned to specific process paths and must
be more tightly specified than general-purpose
grades. The more commonly used alloys for appliances are listed in Table 2.
The greatest savings comes, of course, from
using ferritic grades, and they should be used
whenever forming requirements permit, which
is the majority of the time, since most appliance
components experience little more than cutting,
bending, and welding. There are important precautions, however. Mechanical polishing results
in unacceptable corrosion resistance, and the
low work hardening rate of ferritics causes the
mechanical polish to take on a different color
shade. This subtle difference can be magnified
to objectionable levels when a mechanically
polished ferritic stainless, such as 430, is put
side by side with an austenitic such as 201 or
304. This can be solved by specifying rolled-on
finishes, which look the same on ferritics and
austenitics. These finishes also supply the added
corrosion resistance that makes alloys such as
430 acceptable. It is still preferable to use a
dual-stabilized grade such as 468, which can be
welded without adverse corrosion effects and
has high formability and corrosion resistance at
as little as half the cost of 304 when alloy surcharges are factored in. Use of dual stabilization
permits keeping titanium levels to a minimum,
making it possible to avoid TiN-caused surface
defects, which occur if significant TiN precipitation occurs before solidification in the original
steel production. This occurrence is strictly a
thermodynamic phenomenon related primarily
to the titanium and nitrogen levels, which
should be minimized so that the product of titanium times nitrogen is less than 0.0025 when
Table 2
concentrations are in weight percent. This is difficult to achieve for 17% Cr alloys if stabilization is by titanium alone.
Interior or working parts of appliances, to the
degree they require high cleanability or contact
food, are also often made of stainless. This is
especially true of dispensing machines, such as
for beverages, ice cream, and ice. Stainless interiors are often found in refrigerators and dishwashers. In the case of dishwashers, forming requirements are often severe enough to require
the use of austenitic stainless. Rolled-on finishes are generally preferred. Not a small reason
for this is that this finish requires only a single
temper pass to both flatten and provide the finish. This yields very consistent forming characteristics, meaning much lower breakage during
press-forming operations. Rolled-on finishes
also have very high visual consistency, which is
usually a very important quality criterion for appliance manufacturers.
Canisters, chafing dishes, serving pans, etc.
are generally made from austenitic stainless
steel, which lends itself to the typical deepforming operations used in their manufacture.
Coatings are rarely used. If antimicrobial coating were to be used in food contact, this would
be an ideal application since already cooked
food is most often in the intermediate temperature danger zone at which bacteria can multiply. Food preparation tables also fit into this
category.
Appliance facades are increasingly using
stainless. These include refrigerators, stoves,
microwaves, drawers, etc. Shelves and exhaust
hoods also benefit from being made of stainless.
The drivers here are cleanability, durability,
and esthetics. There are important visual considerations in these applications. Consistent,
Stainless steels commonly used for appliances
Composition, %
Alloy
UNS No.
C
N
Cr
Ni
Mn
Si
Mo
201
301
304
316
430
439
S20100
S30100
S30400
S31600
S43000
S43035
0.15
0.15
0.08
0.08
0.12
0.07
0.25
...
0.10
0.10
...
0.04
16.0–18.0
16.0–18.0
18.0–20.0
16.0–18.0
16.0–18.0
17.0–19.9
3.5–5.5
6.0–8.0
8.0–10.5
10.0–14.0
0.75
0.50
5.5–7.5
2.00
2.00
2.00
1.00
1.0
1.00
1.00
1.00
1.00
1.00
1.0
...
...
...
2.0-3.0
...
...
468
S46800
0.030
0.030
18.0–20.0
0.50
1.00
1.00
436
444,
YUS190
29-4C
S43600
S44400
0.12
0.025
...
0.035
16.0–18.0
17.5–19.5
...
1.0
1.00
1.0
1.00
1.0
S44735
0.025
0.025
28.0–30.0
0.5
1.00
0.75
Ti/Nb
...
...
...
...
...
0.20 + 4 × (C + N),
to 1.10
...
Ti + Nb: 0.20 + 4 ×
(C + N), to 0.80
...
Nb + Ta: 5 × C, to 0.70
1.75–2.50 Ti + Nb: 0.20 + 4 ×
(C + N), to 0.80
3.5–4.5
Ti + Nb: 0.20 + 4 ×
(C + N), to 0.80
240 / Stainless Steels for Design Engineers
defect-free surface finishes are paramount. This
again can really only be achieved by rolled-on
finishes since abrasively polished finishes vary
excessively in roughness, reflectivity, and color.
Panel flatness is often very important and another benefit from rolled finishes. If visible
welds are required, as is often the case for products such as hoods and counters, then special
finishes with very long polish grains have a
major advantage in that the weld can be ground
and polished with a belt sander of the appropriate grit size so that the weld blends imperceptibly with the adjoining original surface. This is a
practical impossibility with abrasively polished
finishes and very difficult with rolled finishes.
Freedom from fingerprinting can be another
valuable attribute for faÁade applications. This
can be obtained on stainless by the mill application of a thin, bonded polymer film. All bare
stainless finishes show fingerprints. With uncoated stainless, it is best avoided by using mineral oil-based cleaners.
Although very high alloy stainless steels are
used for high-temperature kitchen applications,
such as heating element sheathing (American
Iron and Steel Institute [AISI] type 334), it is
seldom used for oven interiors because it does
take on a heat tint when exposed to temperatures above 300 °C (570 °F). Range tops, which
see lower temperatures, are normally stainless.
Outdoor cooking grills, because they must endure exterior environments without corrosion,
are almost always 304 or a similar grade. Heat
tint does not occur with these to a problematic
degree. Gas burner manifolds are also stainless.
In this case, ferritics are required because of the
need for extreme high-temperature oxidation resistance and the desirability of a low coefficient
of thermal expansion. The preferred alloys are
those developed for automotive exhaust systems, variations on 409 and 439.
Flatware and cutlery were among the original uses of stainless. Stainless filled the gap between carbon steel, which was hard but whose
rusting was an obvious problem, and silver,
which was soft and whose cost prohibited its
use to all but the wealthy few. Cutlery is the domain of martensitic stainless steel. The corrosion resistance of martensitic grades cannot be
improved above modest levels, never reaching
that of 304, but this criterion is secondary to
hardness because of the need to keep a sharp
cutting edge. Maximum corrosion resistance is
achieved in the as-quenched condition. But,
some toughness is a valuable but not crucial
characteristic, so most cutlery is tempered at
low temperatures. The vast majority of requirements for high-quality cutlery are satisfied by
420 stainless. If greater cutting edge retention is
desired, then more or harder carbides are engineered into the martensitic matrix. This is done
by adding more carbon and chromium, as is
found in 440A and to a greater extent in 440C.
The wear resistance added by carbides is proportional to their hardness and amount. The
chromium carbides of these straight-chromium
martensitic stainless steels are very hard, 1800
HV, versus the 1100 HV hardness of iron carbides. The addition of higher levels of carbon
ties up chromium so that it cannot add to corrosion resistance, however, so that it can become
barely rust resistant. Furthermore, at high carbon levels, carbides precipitate in the liquid and
are much coarser. These large carbides can pull
out during edge honing, making a ragged rather
than a fine, smooth cutting edge.
However, vanadium and tungsten have even
harder carbides, 2800 and 2100 HV, respectively. Through conventional casting and hot
working, only a small amount of these carbides
can be introduced into the matrix. The problem
is that if primary carbides form during solidification, they tend to be coarse and to embrittle the
alloy. Hard particles are much more useful for
wear resistance if they are small and widely dispersed. To a degree, this refinement of the primary carbides can be achieved by raising nitrogen levels. These problems can be circumvented
by the use of powder metallurgy, which permits
the solidification step on a macroscale to be
skipped. Larger volume fractions of hard carbides such as vanadium carbide and tungsten
carbide can be added and dispersed. Table 3 lists
the martensitic alloys used for cutlery. It is reasonable to say that most of these alloys far exceed the requirements of food preparation.
Improved corrosion resistance of these alloys
is achieved by adding molybdenum at the expense of chromium, which would cause excessive δ-ferrite retention if it were raised. This can
be seen in alloys above the basic 420.
Flatware has no hardness requirement, so
grade selection is based on the need for perceived quality. At the high end is 304, which has
all the corrosion resistance that could be needed
for flatware. However, type 301 is commonly
used also, as are ferritic steels, such as 430, for
low-cost flatware. Depending on the shape of
the final utensil, material is stamped or forged
and then finished.
Chapter 20: Commercial and Residential Applications / 241
Table 3 Stainless steels commonly used for cutlery
Composition, %
Alloy
Designation
420
UNS S42000
4116
DIN 1.4116, nominal
440A
UNS S44002
440C
UNS S44004
BG-42
Nominal composition
ATS-34
Nominal composition
14-4 CrMo Nominal composition
154 CM
Nominal composition
CPM S30V Nominal composition
CPM S60V Nominal composition
CPM S90V Nominal composition
Form
C
Mn
S
Si
Cr
Mo
Ni
Other
Wrought
Wrought
Wrought
Wrought
Wrought
Wrought
Wrought
Wrought
PM
PM
PM
0.15 min
0.50
0.60–0.75
0.95–1.20
1.15
1.05
1.05
1.05
1.45
2.15
2.20
1.00
...
1.00
1.00
...
0.4
0.5
0.45
...
0.40
...
0.030
...
0.030
0.030
...
...
...
...
...
...
...
1.00
...
1.00
1.00
0.3
0.35
0.3
0.3
...
...
...
12.0–14.0
14.5
16.0–18.0
16.0–18.0
14.5
14.0
14.0
14.0
14.0
17.0
13.0
...
0.65
0.75
0.75
4.0
4.0
4.0
4.0
2.0
0.40
1.0
...
...
...
...
...
...
...
...
...
...
...
...
0.15 V
...
...
1.2 V
...
...
...
4.0 V
5.5 V
9.0 V
Many kitchen utensils are also made entirely
or in part with stainless. Type 304 is the alloy
most commonly used, but again any of the
stainless steels with at least 16% Cr are adequate, and grade selection depends on forming
and joining requirements.
Laundry appliances have converted significantly to stainless. This trend began in Europe
with the development of high spin speed, horizontal axis washing machines. These washers
use far less water and energy to achieve higher
levels of cleaning with less damage to clothing.
These features have eroded the share enjoyed by
vertical axis, agitator-type washers, whose low
speed allows them to be constructed of lowstrength materials such as plastic or porcelaincoated steel. The stresses induced by the high
spin speeds, which are necessary in horizontal
axis machines to take water removal from 80%
to 95%, require the strength of stainless steel.
Porcelain-coated carbon steel obviously can be
strong, but the coating is cracked by strains that
the steel itself easily tolerates. An additional benefit to stainless over porcelain is that stainless
starts smooth and becomes even smoother with
use, while porcelain becomes quite abrasive over
time as wear opens voids with edges that can be
quite sharp and cause significant damage to
clothing. Washer tubs and drums are made of
both ferritic and austenitic stainless. The selection is based on forming requirements rather than
corrosion or strength. If components can be made
by bending rather than stretching, then the lowercost ferritics can be used. Ferritics should be a
17% stabilized grade, such as 439 or 468, and
austenitics can be 201, 301, or 304. Unstabilized
ferritic alloys, such as 430, should never be used
in welded applications. Dryers are less challenging, and it is difficult to make a strong case for
the functional value of stainless. Those designs
that use stainless will last longer and be gentler to
clothing. Those, along with the implied quality of
stainless, are the main drivers for its use.
Heating and Water Heating. With the development of high-efficiency, natural-gas-fired,
forced-air furnaces, stainless has come into domestic use as a heat exchanger material. These
furnaces gain their extra efficiency by condensing water from combustion gas exhaust. This
condensate can, depending on incoming air,
contain corrosion elements, which has led to the
use of very highly alloyed ferritic stainless steel
in their construction. Alloy 29-4C (UNS
S44735) was the original alloy used nearly universally in the United States. The worst consequence of perforation by pitting could be the release into the home of toxic gas, so pitting
corrosion must not be allowed.
The intermediate efficiency furnaces (80 to
90%) require the use of corrosion-resistant vent
pipe to prevent corrosion from condensation in
the flue. High-temperature plastics were tried,
but failed joints in them caused their recall after
several fatalities were reported. High-efficiency
(90% or higher) furnaces can use low-temperature plastic pipe, but these units require the use
of a corrosion-resistant secondary heat exchanger to recover the latent heat of vaporization of the water from combustion. Alloy 29-4C
was the original choice for most secondary heat
exchangers, but at least one used alloy 6XN
(UNS N08367) alloy for formability. Some
manufacturers have always used lower-alloyed
stainless steels.
The proper handling of combustion products
is an interesting problem in materials selection.
The variability of the use environments leads to
a huge spread in corrosion conditions and materials performance. In the end, one has to balance
materials selection between cost (fortunately,
29-4C alloy is nickel free) and probability of
failure. Given the number of units produced and
242 / Stainless Steels for Design Engineers
the potentially serious consequences of failure,
failure rates must be less than 10 – 4, while failure rates much less than 10 – 6 are impossible to
verify and hard to justify. In any case, the competition is always between various stainless
steels. Galvanized steel will not work. The issue
is difficult enough for natural-gas-fired furnaces. Oil fired is a developing situation for
which there is no good consensus. Wood burners and other unconventional furnaces (such as
corn burners) present additional challenges, and
answers are even less obvious.
Water heaters are sometimes made of stainless steel. It is not uncommon for water to have
a sufficient level of chlorides to lead to stress
corrosion cracking if an austenitic stainless is
used. Therefore, the recommended alloy for this
application is UNS S44400. More recently, lean
duplex alloys have been developed, such as
2101 and 2003, which can perform quite well
without corrosion or stress corrosion cracking.
More highly alloyed duplex alloys such as 2205
are more expensive but would work well.
REFERENCES
1. M.J. Julio, M.L. Martin, and J.M. Baena,
Cation Migration Tests in Metal Containers,
Innovation Stainless Steel (Florence), Oct
1993, p 1.221–1.226
2. “Stainless Steel in Contact with Food,”
Document Ugine, June 1996
3. “The Application of High Corrosion
Resistance Stainless Steel YUS270 in
Food Processing Facilities and Equipment,”
Nippon Steel Technical Report 87, Jan
2003
4. J.T. Holah and R.H. Thorpe, Bacteria Retention on Cleaned Surfaces, J. Appl. Bacteriol., Vol 69, 1990, p 599–608
5. Agion Technologies, www.agion-tech.com,
accessed June 2008
6. Removal of Stains and Discolourations, Outokumpu, www.outokumpu.com, accessed
June 2008
7. “The Care and Cleaning of Stainless Steel,”
Specialty Steel Industry of North America,
www.ssina.com, accessed June 2008
8. E.P. Kysinski et al., J. Food Processing, Vol
55, 1992, p 246–251
9. A.A. Mafu et al., J. Dairy Sci., Vol 73,
1990, p 3428–4332
10. P. Gelinas and J. Goulet, Can. J. Microbiol.,
Vol 29, 1983, p 1715–1730
11. R.A. Goyer, Toxicity of Metals, Properties
and Selection: Nonferrous Alloys and Special-Purpose Materials, Vol 2, ASM Handbook, ASM International, 1990, p 1233–
1269
12. Allegheny Ludlum research, as presented in
D.S. Bergstrom and C.A. Botti, AL
201HPTM (UNS S20100) Alloy: A HighPerformance, Lower-Nickel Alternative to
300 Series Alloys, Stainless Steel World,
KCI Publishing, 2005
Stainless Steels for Design Engineers
Michael F. McGuire, p 243-246
DOI: 10.1361/ssde2008p243
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
CHAPTER 21
Marine Systems Applications
Summary
AS RECENTLY AS THE 1960s AND 1970s,
handbooks on stainless steel were stating that
“stainless steels are not stainless in seawater,”
and “successful prolonged corrosion-free service of stainless steel in seawater requires sophisticated corrosion engineering, or enormous good
fortune” (Ref 1). The advances in stainless steel
made since then have thankfully made these
statements obsolete. Not only have basic corrosion problems been solved, stress corrosion
cracking also can be avoided. More impressively, this can be done with alloys with
strengths much higher than those of the alloys,
such as 316, that they replace and that have been
only marginally successful in marine environments. The inertia in changing from the weaker,
less-corrosion-resistant, more expensive austenitic stainless steels is large because of less
availability of the newer, better alloys, and lack
of familiarity with their benefits. Those who
understand and use these newer duplex alloys
will be rewarded. This chapter reviews the major
marine applications of stainless steels, including
desalination equipment, shipping containers,
and heat exchangers that handle seawater.
Desalination
At one time not long ago stainless steel was
thought to be an inadequate to marginal material
for use in seawater. Its use in heated seawater
was therefore all the more suspect. This was first
changed with the development of superferritic
and superaustenitic alloys. The superferritic alloys such as Seacure (UNS S44660) and 29-4C
(UNS S44735) are quite resistant to seawater,
even at high temperature. Their low toughness
restricts their use to items of rather thin gauge,
less than about 1.0 to 2.0 mm, depending on
alloy. Thus, their use is limited to tubing. Superaustenitic alloys can be used at any thickness,
although they are a costly material. The success
story for stainless steel and seawater and therefore desalination is that of duplex stainless steel.
With the same corrosion resistance as any super
austenitic or superferritic alloy, it has nearly
double the strength plus resistance to stress corrosion cracking. And while duplex stainless steel
is not a cheap material, it does contain much less
nickel than an equivalently corrosion resistant
austenitic stainless steels, which is a major cost
saving factor.
Desalination technology is relatively new if
one ignores the fact that distillation has been
around for a very long time. Desalination in commercially viable quantities began with multistage flash technology in the 1950s. The underlying principle of this process is the evaporation of
water vapor from salt water with the subsequent
condensation of the salt-free water vapor. In the
multi-stage flash (MSF) approach feedwater is
heated and the pressure is lowered so that the
water “flashes” into steam. A variation on this
technology is multiple-effect distillation (MED),
another low-temperature distillation process. The
differences in all distillation-based systems
reduce to the efficiency of the design in minimizing energy consumed per unit of pure water output. All distillation processes require heating of
the input water and some process power.
The other basic engineering approach to desalination is reverse osmosis (RO). The invention of polymer membranes that could separate
the salt ions from the water made this technology possible. No thermal energy is required. The
water is pumped at high pressure through these
permeable membranes physically separating the
244 / Stainless Steels for Design Engineers
salt from the water. The change in salt concentration across the membrane is a function of the
pressure and the membrane itself. A second
treatment may be used to improve water quality.
The distillation methods require about 5
kWh/metric ton of water output, while the RO
methods require twice that. The distillation
methods require another 20 kWh of thermal energy from some source for feedwater heating,
while the RO method requires none. Thus the
ability to find energy from cogeneration or a
source such as solar, etc may determine which
process is preferred.
Materials Selection for Desalination
Materials used for distillation processes have
evolved from use of type 316 (UNS S31600)
stainless steel, first as lining and then as
cladding. The superaustenitic alloys, the 6Mo
variations, came next because they truly solved
the corrosion problem, but at a price. Then, separately the duplex alloys were developed, with
the first market the petroleum industry, whose
demands and research made these alloys possible. It was not a stretch to see that high-strength
alloys that could withstand seawater in offshore
applications could do well on land as well. To
give full credit, the pulp-and-paper industry was
also beginning to employ duplex stainless steels
for their processes. Type 316 stainless steel has
passed from consideration as a material for handling brackish water or seawater.
In distillation systems, the rule of thumb is
that 2205 alloy (UNS S32205), with its pitting
resistance equivalent number (PREN) of 35, is
sufficient for seawater up to 20 °C (70 °F); alloys 2507 (UNS S32750) or Zeron 100 (UNS
S32760) should be used for seawater at elevated
temperatures or high salinity. For the output of
fresh water, lesser alloying is required. Stainless
steel types 304 (UNS S30400), 316 (UNS
S31600), 2101 (UNS S32101), 2003 (UNS
S32003), or even 439 (UNS S43035) may be
used depending on the combination of salinity
and temperature of the output water.
Besides their high corrosion resistance for
lower total alloy cost, the duplex stainless steels
have higher strength, which is a significant factor since distillation plants are large. The use of
duplex allows wall thickness reductions that
bring about larger savings than those based
solely on their cost per unit weight. Figure 1
shows the difference among the candidate stainless steels in corrosion resistance (Ref 2). The
viable materials for seawater are those that can
withstand roughly 20,000 ppm Cl– level at the
appropriate temperature.
The strengths of the various candidate materials are given in Table 1. These are typical values.
Fig. 1 Corrosion resistance (pitting) as a function of salinity and temperature. 1. 304L (UNS S30403); 2. 316L (UNS S31603);
3. 2205 (UNS S32205); 4. 904L (UNS N08904); 5. 254SMO (UNS S31254). Source: Ref 2
Chapter 21: Marine Systems Applications / 245
Table 1
Typical analyses and properties of major marine alloys
Composition, %
Alloy
2101
2003
2205
2507
304L
316L
317L
6XN
254SMO
Zeron 100
Yield strength
UNS
Cr
Mo
N
Ni
PREN(a)
S32101
S32003
S32205
S32750
S30403
S31603
S31703
N08367
S31254
S32760
21.5
20.5
22
25
18
16
18
21
20
25
0.3
1.5
3
4
0
2
3
6
6
3.5
(+0.75 W)
0.22
0.18
0.17
0.27
0.05
0.05
0.05
0.22
0.20
0.27
1.5
3
5
7
8
10
14
24
18
7
26
29
35
42
18
24
29
45
43
42
Tensile strength
MPa
ksi
MPa
ksi
Elongation,
%
515
515
515
550
220
220
230
380
380
550
75
75
75
80
32
32
33
55
55
80
650
725
760
800
520
520
540
760
750
750
94
105
110
116
75
75
78
110
109
109
40
40
35
35
50
50
45
45
45
35
(a) PREN, pitting resistance equivalent number.
Refer to the appropriate design code for your
particular application to find minimum properties. The reader is cautioned that duplex longitudinal properties are slightly lower than the
transverse properties that testing requires.
Pumps for seawater follow the same guidelines as piping, tanks, and all other components.
Cast or wrought duplex are the alloys of choice.
Shipping
The major uses of stainless steel in shipping
are in bulk storage containment. Cargos range
from food and beverages to chemicals and liquid
natural gas (LNG). Practice in the past has been
to use austenitic grades of stainless with cathodic
protection when necessary to address inadequate
corrosion resistance. However, since 2000 marine chemical tankers have become the largest
consumer of duplex stainless steel. The reason
for this is that cargo tanks ideally have the widest
potential range of cargos possible. This range is
defined by corrosion resistance. This factor alone
is reason to choose duplex over austenitic alloys
such as 316L (UNS S31603) or 317L (UNS
S31703). An equally decisive factor is strength.
With codes permitting the tank’s design to be
based on yield strength, the use of duplex alloys—with strengths about double those of
austenitic steels—permits significant weight reduction. This is a major economic factor for ship
owners in that dead weight can be replaced by
fee-paying cargo at the same operating cost.
These incremental revenues, over the life of the
vessel, are many times the original cost of the
material. Based on the high value for strength in
ship economics, it would seem that the higheststrength alloys, such as 2507 (UNS S32750),
may be justified based on strength alone; their
exceptional corrosion resistance would be simply
an excellent side benefit. Corrugated stainless
bulkheads are positioned within the carbon steel
hull. The stiff corrugated bulkheads are themselves structural strengtheners for the entire ship.
The vertical corrugations also facilitate tank
cleaning as internal stiffeners are eliminated.
Cryogenic containers are still the bastion of
austenitic stainless steels. As leaner austenitic
alloys have become less expensive than 9% Ni
alloy steel, a martensitic grade, they have become the material of choice. In this case, the
201 types are preferred to 304 because 201 has
greater strength at the cryogenic operation temperature and is, of course, less expensive. The
expanding market for LNG has made ocean
transport increasingly important because large
disparities in prices often are due to the difficulty in transporting it. The two best materials
are UNS S20153 and S20400, which perform
equally well. If higher strength is valuable to a
design for cryogenic uses, then UNS S21904
(21-6-9 or Nitronic 40) could be used. This
alloy has yield strengths of 460 MPa (65 ksi) at
room temperature and 1200 MPa (175 ksi) at
–196 °C (–320 °F). It is completely resistant to
martensite formation.
Other shipboard systems benefit equally from
the use of duplex stainless steel. This extends to
piping, hardware, propellers, shafts, etc.
Heat Exchangers
Coolers for captive water systems such as for
power plants often need to resist corrosion by
brackish water or seawater. To the extent that
these are thin-wall tubing, ferritic alloys such as
246 / Stainless Steels for Design Engineers
Seacure (UNS S44660) or 29-4C (UNS S44735)
have been used quite successfully. If thicker
tubes are required, then the equivalent duplex or
austenitic alloys can be used. This would include
types 2003 (UNS S32003), 2205 (UNS S32205),
or 2507 (UNS S32750) duplex stainless steels,
depending on salinity and temperature; for
austenitics, the 6Mo alloys such as 254SMO
(UNS S31254) and AL6XN (UNS N08367) may
be used. The duplex alloys have the advantage
of lower cost. Both are resistant to stress corro-
sion cracking to very high temperature and
salinity.
REFERENCES
1. Peckner and I. Bernstein, Stainless Steel
Handbook, McGraw-Hill, 1966, p 37-1
2. Stainless Steel for Desalination Processes,
Feb 2006, Outokumpu, www.outokumpu.
com, accessed June 2008
Stainless Steels for Design Engineers
Michael F. McGuire, p 247-255
DOI: 10.1361/ssde2008p247
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
CHAPTER 22
Petroleum Industry Applications
Summary
The petroleum industry has driven large segments of the steel industry since both their be-
ginnings. Demand for steel for drill pipe, casing, and tubing has led to many developments,
such as the technology for producing highquality seamless and welded pipe and tubing.
Pipeline needs have fueled the market for highstrength, low-alloy plate. Offshore production
in often-hostile environments has presented severe material challenges. And, as the light sweet
crude that was easily found and produced on
land is exhausted, future supplies of hydrocarbons are increasingly likely to contain sulfides,
carbon dioxide, and saltwater in sufficient
amounts to make corrosion a top priority in selecting materials. For reference in this chapter,
Tables 1 through 5 list the relevant alloys for
petroleum industry applications. Many, but not
all, of these alloys are listed in the National Association of Corrosion Engineers (NACE)
MR0175, Sulfide Stress Corrosion Cracking Resistant Metallic Materials for Oil Field Equipment; the tables in this chapter also include
some newer alloys not in the NACE document.
Table 1 Ferritic stainless steels for petroleum
industry applications
Table 2 Martensitic stainless steels for
petroleum industry applications
THE PETROLEUM INDUSTRY has had to
deal with increasingly hostile environments in
its search for new supplies of oil. And that petroleum, when found, often contains harmful ingredients. The result is increasing demand for
steels with greater strength and corrosion resistance. Martensitic and duplex stainless steels
have provided the corrosion resistance and
strength to deal with higher levels of hydrogen
sulfide, carbon dioxide, chlorides, and acidity.
This chapter reviews the selection of stainless
steels for petroleum applications, including oil
country tubular goods (OCTGs), line pipe, offshore platforms, and refinery equipment.
Introduction
UNS
Common name
S40500
S40900
S43000
S43035
S43400
S43600
S44200
S44400
S44500
S44600
S44626
S44627
S44635
S44660
S44700
S44735
S44800
S46800
405
409
430
439
434
436
442
444 (18-2)
…
446
26-1 Ti, E-Brite
26-1 Cb
26-4-4, Monit
Seacure, SC-1
29-4
29-4C
29-4-2
468
Note: See Appendix 1 for alloy compositions. Source: Adapted from NACE
MR0175, “Sulfide Stress Corrosion Cracking Resistant Metallic Materials for
Oil Field Equipment”
UNS
Common name
J91150
J91151
J91540
K90941
S14125
S41000
S41426
S41427
S42000
S42400
S42500
CA15
CA15M
CA6 NM
9Cr 1Mo
S/W 13Cr
410
13CRS
…
420
F6NM
15Cr
JFE KL-12G
JFE KNHP12Cr
Nippon NT-CRS
Nippon NT-CRSS
420M
L80 13 Cr
Hardness,
HRC, max(a)
...
...
...
...
28
22
...
29
22
23
22
...
...
...
...
...
...
Note: See Appendix 1 for alloy compositions. (a) As specified in NACE
MR0175. Source: Adapted from NACE MR0175, “Sulfide Stress Corrosion
Cracking Resistant Metallic Materials for Oil Field Equipment”
248 / Stainless Steels for Design Engineers
Table 3 Precipitation hardening stainless steels
for petroleum industry applications
UNS
Common name
S13800
S15500
S15700
S17400
S17700
S35000
S35500
S45000
S45500
S46500
S66286
13-8 PH
15-5 PH
15-7 PH
17-4 PH
17-7 PH
AM-350
AM-355
Custom 450
Custom 455
Custom 465
A-286
Custom 465 (275)
Custom 475
Hardness,
HRC, max(a)
...
33
32
33
...
...
...
31
...
...
35
...
...
Note: See Appendix 1 for alloy compositions. (a) As specified in NACE
MR0175. Source: Adapted from NACE MR0175, “Sulfide Stress Corrosion
Cracking Resistant Metallic Materials for Oil Field Equipment”
The modern dilemma that makes stainless necessary is the addition presence of wet carbon
dioxide, which is extremely corrosive to carbon
and alloy steel. As if this is not a sufficient material problem, sometimes the wetness is from
saltwater, which further aggravates corrosivity.
This corrosion problem is compounded by the
accelerating influence of high temperature in
deeper formations. What is the answer to the
corrosion problem? Inhibitors, coatings, cathodic protection, or more corrosion-resistant
materials are the main responses. The first three
responses are not always practical. They also
represent an ongoing cost rather than a one-time
cost. Each situation must be evaluated regarding which is the optimal solution.
Combating Corrosion in Alloys for
Petroleum Applications
Alloying steel with chromium, copper,
molybdenum, and nickel can lower the corrosion rate of steel by a factor of 10,000. Figure 1
shows the influence of chromium alone, which
produces a 100-fold reduction in corrosion of
steel in seawater and carbon dioxide (Ref 1).
Molybdenum is the most powerful alloying
addition to magnify the benefit of chromium.
The effects of copper and nickel are also very
significant, as Fig. 2 (Ref 2) indicates. These additions must be made in a very balanced way if
a tough, fully martensitic structure is to be
maintained. Carbon must be kept low to avoid
the formation of chromium carbides during
tempering, which would counteract the benefit
of the chromium. Nickel is necessary to prevent
σ-ferrite formation, which reduces toughness.
Table 4 Duplex stainless steels for petroleum
industry applications
UNS
Common name
PREN(a)
J93345
J93370
J93380
J93404
S31200
S31260
S31500
S31803
S32001
S32003
S32101
S32205
S32304
S32404
S32520
S32550
S32750
S32760
S32803
S32900
S32906
S32950
S32977
S39274
Escoloy
CD4MCu
Z100
958
44LN
DP3
3RE60
2205 (old)
19D
2003
2101
2205 (new)
2304
U50
52N+
255
2507
Zeron 100
2803Mo
329
2906
7-Mo Plus
AF 918
DP3W
31-47
30-34
38-46
39-47
30-36
34-43
27-31
30-36
20-24
27-31
25-29
34-38
23-27
27-32
37-48
32-44
38-44
40-46
33-41
26-35
36-45
32-43
39-46
39-47
Note: See Appendix 1 for alloy compositions. (a) PREN, pitting resistance
equivalent number. Source: Adapted from NACE MR0175, “Sulfide Stress Corrosion Cracking Resistant Metallic Materials for Oil Field Equipment”
The chapters on martensitic and precipitation
hardening stainless steels discuss this in detail.
The martensitic stainless steels used for these
applications are resistant to carbon dioxideenhanced corrosion up to partial pressures of
100 atm, after which further alloying is necessary. This cannot be achieved with a martensitic
structure, but the duplex alloys have the corrosion resistance and strength to work in this
regime. They have high annealed strength and
can also be cold worked to higher strength levels.
If hydrogen sulfide is present, the selection
process can become more difficult. Highstrength martensitic steels are susceptible to
brittle delayed failure in the presence of hydrogen sulfide. Being stainless does not by itself
provide immunity. If localized corrosion occurs,
hydrogen uptake ensues, and delayed failure
follows. Only keeping hardness below wellestablished levels can render a martensitic alloy
immune. If the localized corrosion can be prevented, however, then the stress corrosion
cracking (SCC) cannot be initiated. Molybdenum alloying expands the pH and chloride
range from which an alloy can be free of the pitting corrosion that initiates SCC, as shown in
Fig. 3 (Ref 2). Martensitic steels of all types
have a maximum in susceptibility to SCC via
hydrogen embrittlement near room temperature.
Duplex alloys and austenitic alloys become
susceptible at higher temperatures and do not
Chapter 22: Petroleum Industry Applications / 249
Table 5
Austenitic stainless steels for petroleum industry applications
UNS
Common name
J92500
J92600
J92602
J92701
J92710
J92800
CF-3
CF-8
CF-20
CF-16F
CF-8C
CF-3M
CF-12M
CG-8M
CK3MCuN
CH-20
CN-3MN
AL 22
CN-7M
20Cb-3
AL 20
20Mo-4
20Mo-6
Sanicro 28
Nicrofer 3127 hMo
20Mod
AL-6X
AL-6XN
JS-700
332
25-6Mo
Cronifer 1925 hMo
URSB-8
904L
201
201LN
202
Nitronic 30 (204L)
204
205
Nitronic 50
Nitronic 60
Nitronic 40(219)
21-6-9 LC
Nitronic 33
J93000
J93254
J93402
J94652
N06022
N08007
N08020
N08020
N08024
N08026
N08028
N08031
N08320
N08366
N08367
N08700
N08800
N08925
N08926
N08932
N80904
S20100
S20153
S20200
S20400
S20430
S20500
S20910
S21800
S21900
S21904
S24000
PREN(a)
...
...
...
...
...
...
...
...
...
...
...
...
...
28
28
38
...
39
54
38
...
49
36
...
46
47
49
39
...
...
...
...
...
...
...
...
...
...
...
UNS
S30100
S30153
S30200
S30215
S30300
S30400
S30403
S30409
S30415
S30453
S30500
S30800
S30815
S30900
S31000
S31008
S31254
S31266
S31600
S31603
S31609
S31635
S31700
S31703
S31725
S31726
S31753
S32100
S32109
S32200
S32654
S33000
S33400
S34565
S34700
S34709
S35125
S35315
Common name
301
301LN
302
302B
303
304
304L
304H
153MA
304LN
305
308
253MA
309
310
310S
254SMO
B66
316
316L
316H
316Ti
317
317L
317LM
317LMN
317LN
321
321H
NIC 25
654SMO
330
334
4565
347
347H
332Mo
353MA
Cronifer 2328
PREN(a)
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
46
59
...
25
...
...
...
...
...
...
...
...
...
...
64
...
...
54
...
...
...
...
...
Note: See Appendix 1 for alloy compositions. (a) PREN, pitting resistance equivalent number. Source: Adapted from NACE MR0175, “Sulfide Stress Corrosion
Cracking Resistant Metallic Materials for Oil Field Equipment”
Fig. 1 Influence of chromium on the corrosion rate of steel in
environments experienced by oil country tubular
goods. Test conditions: synthetic sea water; CO2 partial pressure, 0.1 MPa; test temperature, 60 (C °140 °F); test duration,
150 h; flow velocity, 2.5 m/s; specific volume, 800 mL/cm2.
SSC, stress corrosion cracking. Source: Ref 1
Fig. 2 Influence of copper and nickel on the corrosion rate of
martensitic stainless alloys used for oil country tubular
goods. Source: Ref 2
250 / Stainless Steels for Design Engineers
The NACE recommendations of suitable materials are defined by MR0175. Table 6 summarizes these recommendations. The reader is encouraged to refer to the latest version of this
document for further details.
Oil Country Tubular Goods
Fig. 3 Influence of molybdenum on susceptibility to stress
corrosion cracking in solutions containing (a) 3.5%
NaCl and (b) 0% NaCl. Source: Ref 2
exhibit the same increasing susceptibility with
strength. So, when hydrogen sulfide, which enhances hydrogen uptake, levels exceed about
10–2 atm, the martensitic alloys should no longer
be used, and the duplex alloys are preferred. As
temperatures and hydrogen sulfide partial pressures increase, alloying must also, until at 1 atm
of hydrogen sulfide nickel base alloys are required. Figure 4 shows this progression with the
alloy recommendations of Sumitomo. The requirements behind this diagram are generic.
Any producer’s alloys must comply with this
diagram’s regions, which have been defined by
NACE. Stainless steels are required above a
certain carbon dioxide level for all levels of hydrogen sulfide. Martensitic alloys, commonly
called “13Cr,” are the first step up from alloy
steels. At higher levels of carbon dioxide and
hydrogen sulfide, duplex alloys are required,
with the 22CR alloys such as UNS S32205 used
at temperatures up to 200 °C (390 °F) and the
25CR alloys such as UNS S32507 at temperatures up to 250 °C (480 °F). Nickel-base alloys
are required at hydrogen sulfide levels above 1
atm partial pressure.
Oil country tubular goods (OCTG) include
the drill pipe, casing, and tubing and associated hardware used to construct oil and gas
wells. Drill pipe is used to twist the drill bit
and convey drilling fluids to the point of contact and flush away debris. Casing is put in
place to stabilize the well walls, while tubing is
placed within the casing to carry oil and gas to
the surface. Each of these components sees significant stresses, and high strength-to-weight
materials are needed. Drill pipe is in tension,
torsion, and compression alternately throughout
its life. Casing hangs from the wellhead under
its own weight for distances from hundreds of
meters to 7000 or 8000 m and must withstand
very high collapse as well as burst pressures.
Well strings, the exact sequence of size and
strength pipe for each level of the well, are optimized for the conditions of each well. The variety of strengths and sizes are standardized by
the American Petroleum Institute. The use of
the highest strengths has always been limited by
hydrogen embrittlement accelerated by hydrogen sulfide, so that the maximum hardness for a
given material must be strictly adhered to when
hydrogen sulfide is present.
The terms 13Cr, 22Cr, and 25Cr are commonly used in the industry even though this
greatly oversimplifies the alloying, and therefore performance, options that exist. The 13Cr
alloys are a family of martensitic stainless
steels. The 22Cr and 25Cr alloys are duplex
grades. The former are used in the quenched
and tempered condition, while the duplex alloys
are used as annealed or cold worked.
The 13Cr grades began as simply variations
on 420, which is a straight-chromium martensitic often used for cutlery. This alloy, while far
better (about 100 times) than alloy steel in corrosion resistance, has nearly the least corrosion
resistance of all stainless steels. To achieve
higher corrosion resistance molybdenum is
added. Molybdenum at 1% increases resistance
to general corrosion in a sodium chloride/
hydrogen sulfide/carbon dioxide environment by
about tenfold. Another 1% increases it another
Chapter 22: Petroleum Industry Applications / 251
Fig. 4 Alloy suitability as a function of H2S and CO2 partial pressure. Source: Ref 1
tenfold. The 2% level of molybdenum also
greatly reduces pitting, which in turn eliminates
the initiation point of SCC. Simply adding
molybdenum would cause the alloy to have excessive δ-ferrite, which cannot transform to
martensite and would therefore reduce mechanical properties. Thus, nickel must be added to
counter the ferrite stabilizing effect, unfortunately, but necessarily increasing the cost. The
nickel does help lower the general corrosion
rate. Carbon and nitrogen in these alloys are
kept at low concentrations. These alloys are otherwise almost identical to precipitation hardening martensitic stainless steels without the precipitating phase.
Martensitic alloys are susceptible to SCC by a
hydrogen embrittlement mechanism. This susceptibility is strongly temperature dependent. It
decreases with temperature from a maximum at
ambient to none at around 100 °C (210 °F). If the
hydrogen sulfide level exceeds 0.03 atm, then
22Cr alloys should be used rather than 13Cr
because of this risk. Hydrogen sulfide may be
contained in the petroleum, or it may come from
sulfate-reducing bacteria, introduced by flooding,
for example. This can cause a sulfide-free system
to become sulfide rich after the fact and make initial materials choice wrong after the fact.
The 22Cr and 25Cr alloys have significantly
higher resistance to chlorides and wet hydrogen
252 / Stainless Steels for Design Engineers
Table 6 Restrictions in use recommended by NACE MR0175 for selected stainless steels used for
petroleum industry applications
UNS
J91150
J91151
J91540
J93254
J95370
N08926
S15500
S15700
S17400
S20910
S41000
S41425
S41426
S41427
S41429
S41500
S42000
S42400
S42500
S45000
S66286
Common
name
CA15
CA15M
CA6NM
...
...
...
15-5
15-7
17-4
...
410
...
...
...
...
F6NM
420
...
...
450
A286
Austenitic A-2
Duplex
Superaustenitic,
type 3a
Superaustenitic,
type 3b
PREN(a)
Temperature,
ºC (ºF)
pH, min
H2S, kPa
Hardness,
HRC
Cl–,
mg/L
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
<40
>40 <45
(Ni+2Mo) >30
...
...
...
...
150 (300)
121 (250)
...
...
...
66 (150)
...
...
...
...
...
...
...
...
...
...
65 (150)
...
232 (450)
232 (450)
...
3.5
3.5
3.5
...
...
3.5
...
...
...
...
3.5
...
3.5
3.5
4.5
3.5
3.5
3.5
3.5
...
...
...
...
...
...
10
10
10
...
700
700
3.4
...
3.4
100
10
10
10
10
10
10
10
10
10
1
100
...
10
20
...
22
22
23
100 HRB
94 HRB
...
33
32
33
...
22
28
27
29
27
23
22
23
22
31
35
22
...
...
...
...
...
...
...
90,000
60,700
...
...
...
35
...
...
...
6,000
...
...
...
...
...
...
...
...
...
...
...
121 (250)
149 (300)
171 (340)
...
...
...
700
310
100
...
...
...
5,000
5,000
5,000
>40
Notes: See NACE MR0175 for further use and processing restrictions. See Appendix 1 for alloy compositions. (a) PREN, pitting resistance equivalent number. Source:
Adapted from NACE MR0175, “Sulfide Stress Corrosion Cracking Resistant Metallic Materials for Oil Field Equipment”
sulfide and can resist SCC at ten times higher
concentrations than the 13Cr alloys. Besides the
inherently greater corrosion resistance that derives from the chromium, molybdenum, and nitrogen levels of the 22Cr and 25 Cr alloys (see
the chapters on corrosion in this book), the duplex alloys have very fine grain size and a
roughly 50/50 mixture of ferrite and austenite.
This acts as a crack arrestor should one phase be
susceptible to cracking while the other is not.
There have been no reported downhole failures of annealed or cold-worked duplex alloys.
There was one instance of very high-strength
tubing cracking after cathodic contact with carbon steel casing. This was after removal from
the well and after handling damage had occurred. The affected microstructure was found
to be high (70%) in ferrite (Ref 3).
For corrosion resistance above that furnished
by superduplex materials such as the 25Cr alloys, super austenitic alloys fill a gap before
nickel base alloys are needed. These alloys
achieve a tenfold increase in hydrogen sulfide
resistance and very elevated SCC resistance.
These are the so-called 6Mo grades. The more
advanced of them contain high levels of nitro-
gen. The more common alloys are UNS S32654
and N08367.
The recent development of lean duplex alloys
has not yet made its way into OCTGs. These alloys offer an inherent alloy savings over the
13Cr grades in nickel and molybdenum content
while offering better corrosion and SCC resistance. Their strength levels in the annealed condition, 450 MPa (65 ksi), are lower than those
of the martensitic alloys, 600 MPa (87 ksi), so
for most downhole applications they will require cold working. It is likely, however, that
these alloys will see their first service as line
pipe, where they will not need to be cold worked
to higher strength levels to be widely used.
Line Pipe and Flow Lines
With the awesome cost of corrosion, the
case for stainless line pipe is easily made.
Whether to use stainless depends on vulnerability of carbon steel. This evaluation is made
based on the carbon dioxide, hydrogen sulfide,
water, salinity, temperatures, pressures, flow
conditions, and so forth. The normal basis for
Chapter 22: Petroleum Industry Applications / 253
these calculations follows that published by C.
de Waard of Shell (Ref 3). The competing technology when corrosion dangers arise is the use
of corrosion inhibitors, cathodic protection, or
to line carbon steel with a protective coating.
The use of inhibitors is subject to the risk of velocity limitations, temperature limitations, and
simply of the inhibitor working appropriately,
not to mention cost. Cathodic protection is
costly and complex. Coatings can be damaged
by numerous occurrences in acidity, mechanical damage, or fluctuations in temperature or
pressure. Use of a stainless corrosion-resistant
alloy can have well-defined and controlled
costs and performance over the life of an installation. Line pipe differs from downhole in having strength requirements more in line with that
of annealed duplex alloys. These requirements
gave birth to modern duplex alloys, starting
with UNS S31803 and evolving to UNS
S32205 as the value of higher nitrogen became
understood.
Nitrogen not only enhances corrosion resistance, but also suppresses the formation of undesirable and embrittling intermetallic phases that
might otherwise form at welding temperatures.
It also keeps the desirable austenite/ferrite ratios
in weld metal.
Since the development of the first widely accepted duplex alloys, more alloys have
emerged. Superduplexes, such as UNS S32750,
have become accepted alloys. Then, the need to
improve costs led in the 1990s to the use of
martensitic alloys with high levels of nickel and
molybdenum, which at the time were lower
cost. The emergence in the early 2000s of lean
duplex alloys provided strength and more corrosion resistance with lower nickel levels, giving
them a cost advantage during periods of high
nickel cost.
The main attribute required by line pipe that
is not as important in OCTGs is weldability.
This is not an overwhelming challenge for duplex alloys, but for martensitic alloys, it requires a very low interstitial level so that the
martensite is self-tempering and ductile in the
as-welded condition. This can be achieved by
stabilizing the alloy with small amounts of titanium. It would appear that under current conditions that alloy 2101 (UNS 32101) has a
cost/performance edge over the martensitic
competition and should for the long term. The
main ingredients required in a duplex for
strength and corrosion resistance are chromium
and nitrogen, both relatively inexpensive alloying elements. Alloys S32001 and S32101 are
well formulated for medium and high levels of
corrosion resistance required for wet carbon
dioxide, hydrogen sulfide, and trace chlorides.
The main precaution for duplex alloys is maintaining a nominally 50/50 mixture of ferrite and
austenite with no embrittling intermetallic
phases. The modern alloys have high (greater
than 0.14%) nitrogen, which helps to preserve
austenite levels after welding and suppress intermetallic formation. Nevertheless, minimization of time above 350 °C (660 °F) is important.
This tendency increases with chromium and
molybdenum content, which is another reason
why the lean duplex alloys are so attractive.
For subsea use, 22Cr duplex generally requires cathodic protection because of the risk of
crevice corrosion. 25Cr duplexes are used without cathodic protection. Duplex pipelines have
been in service in the North Sea since the 1970s.
Umbilical Tubing and Risers
Increasingly, wells are located undersea. It is
standard practice to control and monitor these
wells via bundled umbilical tubing. The tubing
can provide hydraulic and electrical power, control and adjust pressure, carry communications,
and even introduce chemical to the well. The
depth of wellheads can increase collapse pressures to levels beyond the capability of thermoplastics, which has led to the use of duplex
stainless steel because of its strength and resistance to corrosion and SCC. When resistance to
seawater is the main concern, the rule of thumb
is that a pitting resistance equivalent number
(PREN) of 35 or greater is required, whereas resistance to crevice corrosion requires a PREN
of at least 40. This has made the superduplex
UNS S32750 the standard. Such a critical item
as an umbilical may seem like a poor application on which to economize, but again the lean
duplexes offer possibilities to do so. By zinc
coating lean duplexes such as alloy 19D (UNS
S32001) and 2101 (UNS S32101), very long
service lives can be safely extrapolated. These
alloys are being promoted on their lower susceptibility to σ formation during welding, and if
welding thermal cycles cannot be controlled
that may be an issue, but superduplex seems to
have become a pervasive choice because it is
superbly reliable.
254 / Stainless Steels for Design Engineers
Risers are now produced in coiled tubing of
over 100 mm (4 in.) diameter, so that very economical long lengths are feasible.
Platforms
Platforms present a special case in which the
costs of maintenance are high, the corrosion environment is severe, and the penalty for excess
weight is also high. A savings of 1 ton in weight
topside can save over $100,000 in steel in the
subsea jacket. This leads to a rapid payback for
the use of materials that are sufficiently resistant
to corrosion such that corrosion loss allowance
can be eliminated. Both titanium and stainless
alloy UNS S32750 are equal candidates for this
service, depending on availability and current
alloy prices. Except in rare cases, stainless steel
wins the cost battle between these alloy systems.
Almost any structure is a candidate for stainless topside processing: piping, pumps, flanges,
fittings, etc. Hardware of any type and construction materials benefit from being stainless. Seawater systems often employ 22Cr duplex with
cathodic protection or unprotected 25Cr duplex.
A wise preventive action is to paint stainless
that is covered by insulation or similar material,
which otherwise can result in concentration
cells and consequent pitting.
Table 7
Liquefied Natural Gas Vessels
Liquefied natural gas (LNG) is becoming an
increasingly important commodity as the value
of stranded gas makes it economically desirable
to convert it to a transportable state. Converting
natural gas to a cryogenic liquid presents a material problem. Vessels to contain it must have
strength and toughness at temperatures below
–150 °C (–240 °F). The traditional material, 9%
Ni martensitic steel, has become expensive
compared to the lower-nickel austenitic stainless steels, such as 201LN (UNS S20153),
which have no transition temperature and
strengthen with decreasing temperature. Alloy
201LN is cheaper, easier to weld and fabricate,
and of course is stainless, which 9% Ni steel is
not. The extreme ductility of 201LN compared
to martensitic steel gives it a decided advantage
in terms of rupture resistance, which is a major
design and political concern with this potentially explosive commodity.
Refinery Equipment
Corrosion resistance is a major factor in the
choice of materials in refinery operations. As we
discussed, crude oil itself is sometimes a very
corrosive fluid, but in refining the by-products,
Stainless steels used in various refinery processes
Process
Corrosive agents
Applications
Preheaters, distillation
tower
Towers
405, 409, 410
…
Vacuum fractionalization
Sulfur-containing acids
(SCAs)
SCA, chlorides
405, 410, 316
Coker
SCA, H2S
Condensers
Coke drums
S44735, 2205
409
Gas plants
H2S, water, Cl–, ammonia
AL-6XN, 2205, 2507
Amine plant
Ammonia, MEA, DEA
410S, 316L
304L, 316L
…
…
Sulfuric acid alkylation
Sulfuric acid
Dilute sulfuric acid
H2S, ammonia, PTA(a)
Compressor coolers,
reboiler tubes
Trays
Reboilers, trays, filters,
condenser tubing
Contactor, mixer
Effluent piping
Hot sections
Depending on crude
corrosivity
Depending on chloride level
Depending on crude corrosivity
…
20Cb3
316L
321, 347
General
Reactor internals
410S, 304
304
Low pH excursions possible
…
Long exposure at high temperature
…
HCl catalyst regeneration
Heat exchangers
Trays
Cyclones, vapor lines
Tubing, heat exchangers
Heat exchangers
2205
410
304
304
409, 321, 347, 2205,
6Mo
304, 20Cb3, 2205
Crude distillation
Hydrotreating
Catalytic re-forming
Fluid catalytic cracking
High-temperature strength
needed
HCl residue
High temperature
Hydrogen plant
Hydrocracking
Sulfides, chlorides
Sour water stripping
Sulfuric acid, ammonium
bisulfide, chlorides
(a) PTA, polythionic acid. Source: Ref 4
Stripper
Alloys
Notes
…
Condensers may need 6Mo
…
…
Depending on temperature,
risk of chlorides
Severity depends on presence of sulfuric acid
Chapter 22: Petroleum Industry Applications / 255
chemicals used in refining and the temperatures
used may further aggravate that corrosivity. The
aggressive chemical agents that refinery materials must withstand include wet hydrogen sulfide
and carbon dioxide, napthenic acids, polythionic
acids, chlorides, sulfuric acid, and alkalines as
well as simple oxidation. Sometimes, temperatures of use are such that embrittling or sensitizing phase transformations may occur. Table 7
lists some major refinery processes and the materials used in them (Ref 4).
Most of these situations are discussed elsewhere in this book in detail. One that is quite
specific to refinery applications is polythionic
acid (PTA) attack. These acids usually form accidentally when sulfide corrosion products react
with moisture and air. The attack is intergranular, and materials respond to it much as they do
to the Strauss test. The remedies are to prevent
the inadvertent formation of PTA and to avoid
using austenitics, which are prone to grain
boundary chromium depletion by sensitization,
and instead use low-carbon grades and stabilized grades.
REFERENCES
1. Sumitomo Products for the Oil and Gas Industries, www.sumitomometals.co.jp, accessed June 2008
2. H. Asahi et al., “Development of High
Chromium Stainless Line Pipe,” Nippon
Steel Technical Report 72, January 1997
3. C. de Waard and U. Lotz, “Prediction of
CO2 Corrosion of Carbon Steel,” Paper 69,
presented at Corrosion/93, National Association of Corrosion Engineers, 1993
4. C.P. Dillon, Corrosion Resistance of Stainless Steels, Marcel Dekker, 1995
Stainless Steels for Design Engineers
Michael F. McGuire, p 257-263
DOI: 10.1361/ssde2008p257
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
CHAPTER 23
Chemical and Process Industry
Applications
Summary
Single- and Dual-Environment Systems
ENGINEERS IN THE PROCESS industries
must have materials that can contain a huge variety of chemical species at many temperatures,
pressures, and flow rates. This is applied corrosion engineering combined with physics and
structural design. It is obvious that this task depends on the availability of corrosion data, more
than can be presented here. This chapter covers
what data are necessary and how they can be
found.
Under ideal conditions, a material may need
to resist one single major corrosion threat. If the
most potentially damaging species can be
clearly identified, then candidate materials can
be found by searching published data. These
data are available freely online from Web sites
(such as Ref 1 and 2) or for a charge from
sources such as the National Association of
Corrosion Engineers (NACE; Ref 3) and ASM
International (Ref 4). It is difficult for any published data to keep up with the latest developments. The testing alone of new materials can
take a long time, and then it must wait for publication. All materials are not covered, especially
when a manufacturer publishes data on proprietary alloys and excludes competitive materials.
That having been said, any improvements over
standard alloys will first be reported by the developers of the alloy, and they will logically
tout its strongest points. For this reason, dialogue with the primary steel producers is encouraged. No one has more exposure to the latest trends in applications.
A single-environment system is typically one
in which the aggressive chemical species is the
only consideration. This is normally the case for
piping, tanks, or reaction vessels holding the
species or materials immersed in the aggressive
species.
A dual-environment system is typically encountered in heat exchangers, but it must also
be extended to single-environment systems in
which the exposure of the nonreactant side of
the material to the ambient environment cannot
be neglected, as in the case of marine ambient
environments.
Introduction
The need to work with hostile chemicals begins with the manufacture of those chemicals. It
was in the production of nitric acid that stainless
had its first industrial application. These are industries with purely need-driven material challenges. New processes are constantly in development, and they present new environments in
which materials must perform. The choices are
highly pragmatic. In an industrial environment,
the costs of a poorly performing material can be
well known by its effect on downtime, maintenance, liability, etc. The essential knowledge is
which materials will work.
The selection of materials for the chemical
and power industries is first a study of corrosion
resistance, including resistance to stress corrosion cracking (SCC). Strength plays a secondary
role but can be an important cost factor. These
considerations may occur at very high or very
low temperatures, in which case corrosion resistance may become oxidation resistance and
strength may mean creep strength.
258 / Stainless Steels for Design Engineers
The challenges that must be met are primarily
ensuring adequate corrosion resistance and secondarily having acceptable mechanical properties. The corrosion issues run the full gamut of
potential forms of corrosion:
• General corrosion
• Pitting corrosion and crevice corrosion
• Intergranular corrosion
• Stress corrosion cracking
• Erosion corrosion
In addition to these forms of corrosion associated with liquids, there are considerations of gas
phase attack, which may be oxidation, sulfidation, or attack by other gases.
Mechanical design considerations are normally limited to static stress allowances. Previously, handbooks dealt very lightly with this
topic because all the normally recommended
steels had similar strength. The proliferation of
duplex stainless steels has changed that. Now,
high-strength alloys of high corrosion resistance
and SCC resistance are available and are making traditionally chosen stainless steels less than
optimal.
Corrosion Types
A designer wants to deal with general corrosion. Its rate can be predicted, and thickness
can be chosen to allow for its occurrence. Corrosion data for general corrosion are normally
presented in isocorrosion charts. These present
the temperatures and concentrations for a
given environment at which various materials
will exhibit the same corrosion rate. This rate
is most often 0.1 mm/yr, an amount that can be
thought of as a tolerable level for many uses.
Figure 1 shows an isocorrosion chart for stainless steels in sulfuric acid (Ref 1). The data are
clear when presented in this fashion. It can further be appreciated that in general reducing
alloy performance to a mathematical formula,
such as the pitting resistance equivalent number (PREN) equation, would not be reasonable
since the relative performance of alloys
changes considerably with concentration.
Thus, the design engineer must rely on experimentally developed data. Since these data are
available both online and in print, no attempt
will be made to reproduce them fully here. Examples are given in Tables 1 and 2 (Ref 1).
Such tables are very useful, although the presentation is not visually compact. The inclusion
Fig. 1 Isocorrosion chart for sulfuric acid. Source: Ref 1
of carbon steel and titanium gives a valuable
frame of reference for the engineer.
If the forms of localized corrosion discussed
next can be avoided, the corrosion tables are
sufficient to guide the designer to a reasonable
selection of candidate materials for any process
in which the chemical species involved have
been identified. If the data have not been developed for a certain environment, then the tables
give a first approximation of which materials
may be resistant from examination of similar
environments, and a final decision can only be
reasonably made though direct corrosion testing
of candidate materials. Refer to the chapters on
corrosion for a more thorough discussion of
uniform corrosion.
Pitting and Crevice Corrosion
Stainless steel is unique among metals and alloys in that it derives its corrosion resistance
from constituent alloying elements working together to form a thin passive layer that, when
intact, is highly resistant to corrosion. The
strength of the passive layer in resisting attack
by halide ions, which are the most disruptive
ions to the layer, is proportional principally to
the chromium, nitrogen, and molybdenum contents of the alloy. This relationship follows the
formula:
PREN = %Cr + 3.3%Mo + 30%N
(Eq 1)
This formula is one of the commonly used
versions, none of which is universally correct.
Both tungsten and carbon can increase pitting
resistance, while sulfur diminishes it. This is
discussed in the corrosion section of this book.
Chapter 23: Chemical and Process Industry Applications / 259
The important consideration is that this formula
assumes that the key alloying elements are homogeneously distributed in solution. This will
only be true if correct thermomechanical processing occurs because, thermodynamically,
these alloys are not used in an equilibrium condition. Were they to attain equilibrium, say by
overheating, alloy segregation by precipitation
could occur, causing localized loss of corrosion
Table 1
resistance, which is what causes pitting.
Chromium is very reactive: Its affinity for oxygen makes the passive film strong. Pitting has
nearly always been associated with manganese
sulfide inclusions, and although there is still debate over the precise mechanism, it appears that
chromium depletion at the metal-inclusion interface is to blame. Eliminating inclusions by
eliminating either manganese or sulfur improves
Corrosion table for sulfuric acid (H2SO4)
Concentration, %
Temperature, °C
Carbon steel
13% Cr steel
18-2 (UNS S44400)
3R12 (UNS S30400)
3R60 (UNS S31600)
18-13-3
17-14-4
2RK65
(UNS N08904)
Sanicro 28
(UNS N08028)
254SMO
(UNS S31254)
654 SMO
(UNS S32654)
SAF 2304
(UNS S32304)
SAF 2205
(UNS S31803)
SAF 2507
(UNS S32750)
Titanium
0.1
100 = BP
0.5
20
0.5
50
0.5
100 = BP
1
20
1
50
1
70
1
85
1
100 = BP
2
20
2
50
2
60
3
20
3
35
3
50
2
2
2
2
1
1
1
0
2
2
0
0
0
0
0
0
2
2
2
1
0
0
0
0
2
2
2
2
1
1
1
1
2
2
0
0
0
0
0
0
2
2
2
1
0
0
0
0
2
2
2
1
0
0
0
0
2
2
2
2
1
1
0
0
2
2
2
2
1
1
1
1
2
2
0
0
0
0
0
0
2
2
2
1
0
0
0
0
2
2
2
1
0
0
0
0
2
2
0
0
0
0
0
0
2
2
2
1
0
0
0
0
2
2
2
1
0
0
0
0
...
0
0
0
0
0
0
0
0
0
0
0
0
0
0
...
0
0
...
0
0
0
0
1
0
0
0
0
0
0
...
0
0
...
0
0
0
0
0
0
0
0
0
0
0
1
0
0
...
0
0
0
0
1
0
0
0
0
0
0
...
0
0
1
0
0
0
0
...
0
0
0
0
0
0
...
0
0
...
0
0
0
0
0
0
0
0
0
0
0
0
0
0
1
Concentration, %
Temperature, °C
3
85
1
3
100 = BP
5
20
5
35
1
5
60
0
5
75
0
5
85
1
5
101 = BP
1
10
20
1
10
50
0
10
60
0
10
80
1
10
102 = BP
0
20
20
20
40
Carbon steel
13% Cr steel
18-2 (UNS S44400)
3R12 (UNS S30400)
3R60 (UNS S31600)
18-13-3
17-14-4
2RK65 (UNS
N08904)
Sanicro 28
(UNS N08028)
254SMO
(UNS S31254)
654 SMO
(UNS S32654)
SAF 2304
(UNS S32304)
SAF 2205
(UNS S31803)
SAF 2507
(UNS S32750)
Titanium
2
2
2
2
1
1
1
0
2
2
2
2
2
2
2
1
2
2
2
1
0
0
0
0
2
2
2
1
0
0
0
0
2
2
2
2
1
0
0
0
2
2
2
2
1
1
1
0
2
2
2
2
2
2
2
1
2
2
2
2
2
2
2
2
2
2
2
2
0
0
0
0
2
2
2
2
1
1
0
0
2
2
2
2
1
1
1
0
2
2
2
2
2
2
2
1
2
2
2
2
2
2
2
2
2
2
2
2
0
0
0
0
2
2
2
2
1
1
1
0
...
1
0
0
0
0
0
2
0
0
0
0
2
0
0
...
1
0
0
0
0
1
2
0
0
0
0
2
0
0
0
0
0
0
0
0
0
2
...
...
0
0
...
0
0
...
1
0
0
0
0
0
2
0
0
0
2
2
1
2
...
1
0
0
0
0
0
2
0
0
0
1
2
0
0
...
1
0
0
0
0
...
...
0
0
0
0
2
0
0
1
2
0
1
1
2
2
2
1
2
2
2
2
2
2
(continued)
Notes: 0, corrosion rate of less than 0.1 mm/yr. The material is corrosion proof. 1, corrosion rate of 0.1–1.0 mm/yr. The material is not corrosion proof but useful in certain cases. 2, corrosion rate of more than 1.0 mm/yr. Serious corrosion. The material is not usable. BP, boiling solution. Source: Adapted from Ref 1
260 / Stainless Steels for Design Engineers
Table 1
(continued)
Concentration, %
Temperature, °C
20
50
20
60
20
80
20
100
30
20
30
40
30
60
30
80
40
20
40
40
40
60
40
90
50
20
50
40
50
70
Carbon steel
13% Cr steel
18-2 (UNS S44400)
3R12 (UNS S30400)
3R60 (UNS S31600)
18-13-3
17-14-4
2RK65
(UNS N08904)
Sanicro 28
(UNS N08028)
254SMO
(UNS S31254)
654 SMO
(UNS S32654)
SAF 2304
(UNS S32304)
SAF 2205
(UNS S31803)
SAF 2507
(UNS S32750)
Titanium
2
2
2
2
1
1
1
0
2
2
2
2
2
1
1
0
...
...
...
...
...
...
...
1
2
2
2
2
2
2
2
2
2
2
2
2
1
1
1
0
2
2
2
2
2
1
1
0
2
2
2
2
2
2
2
1
...
...
...
...
...
...
...
...
2
2
2
2
2
2
2
0
2
2
2
2
2
2
2
0
2
2
2
2
2
2
2
1
2
2
2
2
2
2
2
2
2
2
2
2
2
2
2
0
2
2
2
2
2
2
2
0
2
2
2
2
2
2
2
2
0
0
...
2
0
0
1
...
0
0
1
2
0
0
1
0
0
...
2
0
0
1
2
...
1
...
...
0
1
...
0
0
0
2
...
...
...
...
0
0
0
...
0
0
...
2
2
2
2
2
2
2
2
2
2
2
2
2
2
2
0
1
2
2
0
1
2
2
2
2
2
2
2
2
2
0
0
1
2
...
0
1
2
0
1
2
2
1
1
2
2
2
2
2
2
2
2
2
2
2
2
2
2
2
2
Concentration, %
Temperature, °C
60
20
60
40
60
70
70
20
70
40
70
70
80
20
80
40
80
60
85
20
85
30
85
40
85
50
90
20
90
30
Carbon steel
13% Cr steel
18-2 (UNS S44400)
3R12 (UNS S30400)
3R60 (UNS S31600)
18-13-3
17-14-4
2RK65
(UNS N08904)
Sanicro 28
(UNS N08028)
254SMO
(UNS S31254)
654 SMO
(UNS S32654)
SAF 2304
(UNS S32304)
SAF 2205
(UNS S31803)
SAF 2507
(UNS S32750)
Titanium
2
2
2
2
2
2
2
0
2
2
2
2
2
2
2
1
2
2
2
2
2
2
2
1
2
2
2
2
2
2
2
0
2
2
2
2
2
2
2
1
2
2
2
2
2
2
2
1
2
2
2
2
1
1
1
0
2
2
2
2
2
2
2
1
2
2
2
2
2
2
2
2
0
1
1
1
1
1
1
0
1
1
1
1
1
1
1
0
2
2
1
1
1
1
1
1
2
2
2
2
2
2
2
1
0
0
0
0
0
0
0
0
1
1
1
0
0
1
1
0
0
0
1
0
0
1
...
1
1
0
0
0
0
0
0
0
1
...
0
1
...
0
1
2
0
...
...
...
1
...
0
1
...
0
1
...
...
...
...
...
...
...
...
1
...
2
...
...
...
...
...
...
...
...
1
1
...
...
...
1
2
2
2
1
...
...
2
2
2
1
...
...
...
1
1
...
...
...
...
2
2
...
2
2
1
1
...
...
0
0
2
2
2
2
2
2
2
2
Concentration, %
Temperature, °C
90
40
90
70
94
20
94
30
94
40
2
94
50
2
96
20
2
96
30
2
96
40
2
96
50
2
98
30
2
98
40
98
50
98
80
Carbon steel
13% Cr steel
18-2 (UNS S44400)
3R12 (UNS S30400)
3R60 (UNS S31600)
18-13-3
17-14-4
2RK65
(UNS N08904)
Sanicro 28
(UNS N08028)
2
2
2
2
1
1
1
1
2
2
2
2
2
2
2
2
0
0
0
0
0
0
0
0
2
1
0
0
0
0
0
0
2
2
2
1
0
1
1
1
2
2
2
1
1
1
1
1
0
0
0
0
0
0
0
0
1
1
0
0
0
0
0
0
2
2
1
0
0
1
1
1
2
2
2
1
1
1
1
1
1
1
0
0
0
0
0
0
1
1
1
0
0
0
0
1
2
2
2
2
0
1
1
1
2
2
2
2
2
2
2
2
0
1
0
0
0
0
0
0
0
1
0
0
0
1
(continued)
Notes: 0, corrosion rate of less than 0.1 mm/yr. The material is corrosion proof. 1, corrosion rate of 0.1–1.0 mm/yr. The material is not corrosion proof but useful in certain cases. 2, corrosion rate of more than 1.0 mm/yr. Serious corrosion. The material is not usable. BP, boiling solution. Source: Adapted from Ref 1
Chapter 23: Chemical and Process Industry Applications / 261
Table 1
(continued)
Concentration, %
Temperature, °C
90
40
90
70
94
20
94
30
94
40
94
50
96
20
96
30
96
40
96
50
98
30
98
40
98
50
98
80
254SMO
(UNS S31254)
654 SMO
(UNS S32654)
SAF 2304
(UNS S32304)
SAF 2205
(UNS S31803)
SAF 2507
(UNS S32750)
Titanium
1
...
...
...
...
...
1
...
...
...
...
...
0
2
2
2
...
...
...
2
0
1
...
2
...
...
1
1
1
...
...
...
...
...
1
...
...
0
...
...
0
1
1
...
0
...
...
...
0
0
1
...
0
0
1
1
0
...
0
0
0
1
0
0
0
1
0
0
0
1
2
2
2
2
2
2
2
2
2
2
2
2
2
2
Notes: 0, corrosion rate of less than 0.1 mm/yr. The material is corrosion proof. 1, corrosion rate of 0.1–1.0 mm/yr. The material is not corrosion proof but useful in certain cases. 2, corrosion rate of more than 1.0 mm/yr. Serious corrosion. The material is not usable. BP, boiling solution. Source: Adapted from Ref 1
Table 2 Corrosion table for fuming sulfuric
acid (oleum), H2SO4 + SO3
When a crevice is permitted to exist, it mimics the pH-altering action found within pits in
which transport restriction leads to a buildup of
metal and hydrogen ions and oxygen depletion.
All alloys undergo crevice corrosion under
less-aggressive conditions than those required
to induce pitting, so care must be taken to avoid
crevices.
Conc. H2SO4, %
Conc. SO3, %
Temperature, °C
100
7
60
100
11
60
100
11
100
100
60
20
100
60
70
100
60
80
Carbon steel
13% Cr steel
18-2
(UNS S44400)
3R12
(UNS S30400)
3R60
(UNS S31600)
18-13-3
17-14-4
2RK65
(UNS N08904)
Sanicro 28 (UNS
N08028)
254SMO
(UNS S31254)
654 SMO
(UNS S32654)
SAF 2304
(UNS S32304)
SAF 2205
(UNS S31803)
SAF 2507
(UNS S32750)
Titanium
0
0
0
0
0
0
2
2
...
...
...
...
...
...
...
...
2
...
0
0
1
0
0
0
0
0
0
0
0
0
Intergranular Corrosion
0
0
0
0
0
0
...
...
...
...
0
0
0
0
0
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
2
2
2
2
2
2
Intergranular corrosion is a problem that can
be avoided entirely by correct alloy selection
and proper thermal processing. The principle
cause of grain boundary attack is alloy depletion at the grain boundaries. The most familiar
form of this problem occurs when austenitic
alloys having carbon levels above 0.03% are
welded. The region near the weld where temperatures reach 600 to 900 °C (1100 to 1650 °F)
may have carbon migrate to and along grain
boundaries, the fast diffusion paths, where it
combines with less-mobile chromium atoms
and precipitates as chromium carbide. This
lowers the chromium level in solution, resulting in poor corrosion resistance only at the
grain boundaries. This is easily prevented by selecting alloys with low carbon levels. Duplex
alloys, curiously, undergo chromium carbide
precipitation under the same conditions but do
not undergo significant chromium depletion
because the neighboring ferrite grains, in which
chromium diffuses more rapidly, contribute
chromium, mitigating the depletion. Precipitation segregation of all types, not just by carbides,
must be guarded against. Sigma phase, nitrides,
secondary austenite, and others can cause local
breakdown of corrosion resistance if alloys are
heated to a dangerous temperature for sufficient
time. It is important to learn these potential
Notes: 0, corrosion rate of less than 0.1 mm/yr. The material is corrosion proof. 1,
corrosion rate of 0.1–1.0 mm/yr. The material is not corrosion proof but useful in
certain cases. 2, corrosion rate of more than 1.0 mm/yr. Serious corrosion. The
material is not usable. Conc., concentration. Source: Adapted from Ref 1
the potential at which passive film breakdown
occurs. This is especially important for welds,
which, if not annealed, can have maximum
deleterious segregation by both inclusions and
solidification segregation. All austenitic and
duplex stainless alloys have best corrosion resistance when quenched from the solution annealing temperature. The precipitation hardening, martensitic, and ferritic alloys are more
complicated but are less relevant to this topic. If
information on them is needed, they are discussed in detail in their respective chapters.
262 / Stainless Steels for Design Engineers
vulnerabilities by reviewing the metallurgy of
any alloy selected for service.
Stress Corrosion Cracking
The theory of SCC is still under debate. The
reader will find the arguments confusing as the
debate generates more heat than light. We will
skip the theory; it can be found in the corrosion
chapters. SCC, like excessive general corrosion
or pitting, is avoided by referring to published
test data from the corrosion tables. If a material
must be used where a risk of SCC occurs, then
stress levels must be managed to stay below the
threshold stress for SCC. Figure 2 shows how
various alloys resist SCC as a function of chloride concentration and temperature, the two
most important aggravating factors. Material
comparisons are made difficult because tests are
normally run at a given fraction of a material’s
yield strength. Thus, the data in Fig. 2 (Ref 1)
must be interpreted. Higher-strength duplex alloys, while having better SCC performance than
austenitics of equal corrosion resistance (e.g.,
316 vs. 2304), have much better SCC resistance. Furthermore, the stress at which failure
will occur is much higher since the yield
strength at which the testing takes place is about
twice as high for duplex alloys. SCC also exhibits a threshold stress below which failure
does not occur. This is about 60% of tensile
strength for duplex and about 30% for
Fig. 2 Stress corrosion cracking (SCC) resistance in neutral
chloride solutions containing 8 ppm oxygen. Testing
time, 1000 h. Applied stress equal to proof strength at testing
temperature
austenitics. Designing within this limit is sensible practice. And, if alloy selection uses a rule
of avoiding situations in which pitting can
occur, SCC will also be avoided even if stress
excursions occur since in general pitting is a
necessary precondition for SCC.
Erosion
Flow velocities can reach levels at which erosion becomes problematic, especially if hard
particles are suspended in a fluid. Assuming that
the material has sufficient corrosion resistance
to survive well in the static environment, the
best performance under erosive conditions is
obtained by materials with higher surface hardness. Accordingly, the duplex perform better
than austenitic alloys of the same corrosion resistance level.
Specific Environments
The list of specific environments against
which stainless steels are sufficiently resistant
to select for use in the chemical process industries is too long to provide here. Some of the
most important specific corrosives, such as
nitric, sulfuric, phosphoric, hydrochloric, and
organic acids and others, are covered in the
chapter on corrosion. The main caution to the
designer is to make sure that the source material from which design guidance is sought is
current. Many otherwise excellent handbooks
are somewhat obsolete in that they do not
include the very importance duplex stainless
steel family or only include the oldest alloys in
the group, such as 2205 (UNS S32205). Many
new alloys now exist that range in corrosion
performance from that of 316 to that of the
6Mo-plus-N austenitics. These alloys are usable in all gauges, have high strength and
toughness, resist SCC, and can achieve the
corrosion resistance levels of any ferritic or
austenitic alloy. They can also provide significant savings in alloy cost at the same corrosion
level because they have lower nickel levels.
REFERENCES
1. Sandvik Materials Technology, www.smt.
sandvik.com, accessed June 2008
Chapter 23: Chemical and Process Industry Applications / 263
2. Outokumpu Corrosion Handbook for Stainless Steels, www.outokumpu.com, accessed
June 2008
3. P.A. Schweitzer, Corrosion Resistance
Tables, 5th ed., National Association of Corrosion Engineers, NACE 37755, 2004 \aq2\
4. D.B. Anderson and B.D. Craig, Handbook
of Corrosion Data, 2nd ed., ASM International, 1995
Stainless Steels for Design Engineers
Michael F. McGuire, p 265-267
DOI: 10.1361/ssde2008p265
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
CHAPTER 24
Pulp-and-Paper Industry Applications
Summary
THE PULP-AND-PAPER INDUSTRY has
seen more benefits from developments in stainless steel than any other industry. The harsh
chemicals used in this industry called for better
materials than the normal austenitic stainless
steels without the expense of the 6Mo grades.
This need has been met through the use of the
duplex alloys, which have become the new
standard.
Introduction
The proximity of the Scandinavian paper industry to that region’s specialty steel industry
has been symbiotic. As a result of the strong
interaction between engineers having wellspecified needs for improved materials and metallurgists capable of providing them, the advances in materials in the pulp-and-paper
industry have been a model of rapid technology
transfer and innovation. Beginning in 1988, duplex stainless steels first went into production in
kraft digesters, and there has been no turning
back in the replacement of austenitic stainless
steels by duplex. So, a discussion of the materials selection for the pulp-and-paper industry has
changed from a fairly complicated analysis of
which austenitic steel to use while guarding
against stress corrosion cracking and pitting
corrosion and when to use clad materials for
cost savings, to a fairly simple discussion of
which duplex stainless steel is most economical
for a given piece of equipment. Since this revolution occurred in the 1990s before the latest
surge in nickel prices, it is safe to say that future
pulp-and-paper projects will be essentially the
exclusive domain of duplex stainless steels because of their lower cost per unit of corrosion
resistance, high strength, and near immunity to
stress corrosion cracking. Pricing changes
mainly with alloying element costs, principally
those of nickel and molybdenum. At prices between the highs and lows of the first decade of
the 2000s, duplex costs have been roughly onethird less than that of an equivalent corrosionresisting austenitic stainless. This factored in
with strength nearly double that of the equivalent austenitic make them an overwhelmingly
superior choice for pulp-and-paper equivalent
except if very special corrosion requirements
differ from the norm, such as in bleaching.
Paper-Making Processes
The kraft (German for “strong”) process was
introduced in 1937, replacing the sulfite process.
In the kraft process, the lignin-connecting wood
fibers are dissolved under conditions of elevated
temperature and pressure in acidic conditions of
pH 2.0 to 4.0. This leaves a long fiber, which enables paper of high strength, hence the name
kraft. Over the years the materials used for the
vessels, called digesters, in which this process is
carried out have been sequentially carbon steel,
stainless steel, and stainless steel clad onto carbon steel. In the previous sulfite process, acidresistant brick vessels were used. Now, the digesters, essentially large vertical tanks, are
constructed of 2205 (UNS S32205) as a rule (see
Fig. 1).
The digestion is typically carried out at 150
and 180 °C (300 and 360 °C) and 10 to 12 bar.
The pH of the sulfate is around 2.0 to 4.0. In
this environment, 316L can survive, but it
266 / Stainless Steels for Design Engineers
Fig. 1 The first kraft digester fabricated from alloy 2205.
Courtesy of Outokumpu
requires maintenance and has a finite life. The
2205 is twice as resistant to corrosion, 0.005
mm/yr versus 0.011 mm/yr (Ref 1). In the
nonchloride environment, molybdenum is not
an essential alloying element, so the introduction of the use of 2101 (UNS S32101) or 2304
(UNS S32304) is a logical cost-saving move
without strength or corrosion compromises. The
reduction in wall thickness allowed by the
higher-strength duplex depends on the engineering code required. The American Society of
Mechanical Engineers (ASME) code requirement is based on tensile strength and permits
only a 24% reduction in wall thickness, while
the total kjeldahl nitrogen (TKN) code, based
on yield strength, would allow a 46% reduction.
This large a difference in strength levels required by codes is unfortunate and reflects an
orientation to materials in which the yield/
tensile ratio is closer to unity, unlike either duplex or austenitic stainless steel. In the more unusual case of digesters using the sulfite process,
the materials selected would be the same.
As one proceeds downstream in the process,
environments change greatly, but the optimal
materials remain duplex for various reasons.
The subsequent stage is blow tanks in which the
pulp suspension is injected at high velocity. The
environment is a mixture of alkaline liquid,
while the vapor phase can contain organic acids.
The hardness of the duplex helps mitigate erosion, while the alloy level is beneficial against
corrosion. 2205 is the alloy of choice here, but
2003 (UNS S32003) would suffice.
The next step, washing and screening, has
seen increasingly severe environments as closed
systems required for pollution control have become more common. This has rendered the previous choice of carbon steel untenable. This
stage also sees erosion potential from hard particles, such as sand, in the pulp. The optimum
solution is a lean duplex such as 2101, 2304, or
2003.
The delignification of the pulp comes next.
This oxygen process dates from the 1970s. At
first, highly alloyed austenitic alloys were
used. Subsequently, it was found again that duplex performed better in that they were sufficiently corrosion resistant, but also offered
freedom from stress corrosion cracking as well
as materials savings because of their higher
strength.
The bleaching of the pulp is important for
many types of paper, and this can be done by
chlorine bleaches or ozone/peroxide bleaches.
The chlorine bleaching now must generally be
done in closed systems, which results in a
buildup of chloride levels to a point at which
corrosion levels are unacceptable unless very
highly alloyed materials are used. The 6Mo
grades have been successful, but now they can
be replaced by duplex alloys such as 2507
(UNS S32750), which again save cost by virtue
of their higher strength.
Bleaching can be accomplished without chlorine in the so-called TCF, totally chlorine free,
process. This reduces the corrosivity of the environment as the ozone and hydrogen peroxide
used in the process are relatively harmless to
stainless steel. Alloys such as 316 are adequate
for this environment, but lean duplex, 2101 or
2304, offer cost reductions through their greater
strength.
In plants that use recycled paper and mechanical wood chip processing, the materials selection criteria remain the same. Duplex stainless
has become the clear choice.
Further downstream, containers and process
equipment benefit equally from duplex down to
the handrails and walkways. This wholesale use
of duplex can make plants nearly maintenance
Chapter 24: Pulp-and-Paper Industry Applications / 267
free from a corrosion point of view, a dramatic
change in an industry in which the thousand-fold
greater corrosion rates of carbon steel presented
operators with endless equipment downtime
problems.
Additional detail about corrosion challenges
and the use of stainless steels in the pulp-andpaper industry can be found in Ref 2.
REFERENCES
1. A. Tuomi et al., Duplex America 2000 Conference, Houston, KCI Publishing, 2000
2. H. Dykstra et al, Corrosion in the Pulp and
Paper Industry, Corrosion: Environments
and Industries, Vol 13C, ASM Handbook,
ASM International, 2006, p 762–802
Stainless Steels for Design Engineers
Michael F. McGuire, p 269-278
DOI: 10.1361/ssde2008p269
APPENDIX 1
Compositions
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
0.15
0.03
0.03
0.15
0.15
0.08
0.12–0.25
0.0–3
0.08
0.08
0.08
0.04
0.06
0.10
0.12
0.12
0.15
0.08
0.03
0.15
0.03
0.15
0.08
0.15
0.15
0.15
0.15
0.08
0.03
0.04–0.10
0.08
0.08
0.03
0.08
0.04–0.06
0.12
0.018
0.16–0.24
0.08
0.05–0.10
0.20
0.08
S21500
S21600
S21603
S30100
S30153
S30200
S30430
S30215
S30300
S30223
S30310
S30400
S30403
S30409
S30451
S30452
S30453
S30424
S30415
S30500
S30600
S30615
S30800
S30815
S30900
S30908
C
S20100
S20103
S20153
S20161
S20200
S20300
S20500
S20400
S24100
S24300
S21900
S21904
S20910
S21800
S21400
S21460
Designation(a)
0.10–0.20
...
...
...
...
...
...
0.10
0.10
...
0.10–0.16
0.16–0.30
0.10–0.16
0.10
0.12–0.18
...
...
...
...
0.14–0.20
...
...
...
0.25–0.50
0.25–0.50
0.25
0.25
0.25
0.08–0.20
0.25
...
0.32–0.40
0.15–0.30
0.20–0.40
0.20–0.40
0.15–0.40
0.15–0.40
0.20–0.40
0.08–0.18
0.35
0.25–0.50
N
Cr
14.0–16.0
17.5–22.0
17.5–22.0
16.0–18.0
16.5–18.0
17.0–19.0
17.0–19.0
17.0–19.0
17.0–19.0
17.0–19.0
17.0–19.0
18.0–20.0
18.0–20.0
18.0–20.0
18.0–20.0
18.0–20.0
18.0–20.0
18.0–20.0
18.0–19.0
17.0–19.0
17.0–18.5
17.0–19.5
19.9–21.0
20.0–22.0
22.0–24.0
22.0–24.0
16.0–18.0
16.0–18.0
16.0–18.0
15.0–18.0
17.0–19.0
16.0–18.0
16.5–18.0
15.0–17.0
16.5–19.0
17.0–19.0
19.0–21.5
19.0–21.5
20.5–23.5
16.0–18.0
17.0–18.5
17.0–19.0
Composition of austenitic stainless steels
5.5–7.5
5.5–7.5
5.5–7.5
4.0–6.0
7.5–10.0
5.0–6.5
14.0–15.5
7.0–9.0
11.0–14.0
11.5–14.5
8.0–10.0
8.0–10.0
4.0–6.0
7.0–9.0
14.5–16.0
14.0–16.0
5.5–7.0
7.5–9.0
7.5–9.0
2.0
...
2.0
2.0
2.0
2.0
2.0
2.5–4.5
2.0
2.0
2.0
2.0
2.0
2.0
...
0.8
2.0
2.0
2.0
2.0
0.8
2.0
2.0
(continued)
9.0–11.0
5.0–7.0
7.5–9.0
6.0–8.0
6.0–8.0
8.0–10.0
8.0–10.0
8.0–10.0
8.0–10.0
8.0–10.0
7.0–10.0
8.0–10.5
8.0–10.5
8.0–10.5
8.0–10.5
8.0–10.5
8.0–10.5
12.0–15.0
9.0–10.0
10.5–13.0
14.0–15.5
13.5–16.0
10.0–12.0
10.0–12.0
12.0–15.0
12.0–15.0
Mn
3.5–5.5
3.5–5.5
3.5–5.5
4.0–6.0
4.0–6.0
5.0–6.5
1.0–1.75
1.5–3.0
0.5–2.5
2.25–3.75
5.5–7.5
5.5–7.5
11.5–13.5
8.0–9.0
0.75
5.0–6.0
Ni
...
2.0–3.0
2.0–3.0
...
...
...
...
...
0.60 optional
0.60 optional
0.60 optional
...
...
...
...
...
...
2.0
...
...
0.2
...
...
...
...
...
...
...
...
...
...
0.5
...
...
...
...
...
...
1.5–3.0
...
...
...
Mo
Composition, %
Si
1.00
1.00
2.00–3.00
1.00
1.00
1.00
1.00
1.00
1.00
1.00
1.00
1.00
0.75
1.00–2.00
1.00
3.75–4.25
3.2–4.0
1.00
1.4–2.0
0.75
0.75
1.20
1.00
1.00
1.00
1.00
0.75
0.75
3.00–4.00
1.00
1.00
1.00
1.0
1.00
1.00
1.00
1.00
1.00
3.50–4.50
0.30–1.00
1.00
Notes: All compositions include Fe as balance. Single values are maximum values unless otherwise noted. (a) Unified Number System, UNS numbers are S or N followed by 5 digits.
201
201L
201LN
Gall-Tough
202
230 EZ
205
Nitronic 30
Nitronic 32
Nitronic 33
Nitronic 40 (219)
21-6-9 LC
Nitronic 50
Nitronic 60
Tenelon
Cryogenic
Tenelon
Esshete1250
216
216L
301
301LN
302
302Cu
302B
303
303Se
303 Plus X
304
304L
304H
304N
304HN
304LN
304BI
153MA
305
Cronifer 1815
RA 85 H
308
253MA
309
309S
Name
Table A1.1
P
0.040
0.045
0.045
0.045
0.045
0.045
0.045
0.045
0.045
0.045
0.045
0.045
0.045
0.045
0.045
0.045
0.045
...
...
...
...
0.045
0.045
...
0.045
0.045
0.060
0.045
0.045
0.040
0.060
0.040
0.060
0.040
0.060
0.060
0.060
0.060
0.040
0.040
0.060
0.060
S
0.030
0.030
0.030
0.030
0.030
0.030
0.030
0.030
0.15 min
0.06 min
0.25 min
0.030
0.030
0.030
0.030
0.030
0.030
...
...
...
...
0.030
0.030
...
0.030
0.030
0.030
0.015
0.015
0.040
0.030
0.18–0.35
0.030
0.030
0.030
0.030
0.030
0.030
0.030
0.030
0.030
0.030
Nb 0.75–1.25
...
...
...
...
...
Cu 3.0–4.0
...
...
Se 0.15 min
...
...
...
...
...
...
...
B 1.00–1.20
Ce 0.04
...
Cu 0.50
Al 0.8–1.5
...
1.0 Al
...
...
...
...
...
...
...
Cu 1.75–2.25
...
...
...
...
...
...
Nb 0.1–0.3
...
...
...
Other
V 0.15–0.40
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
0.03–0.08 Ce
...
...
...
...
...
...
...
...
...
...
...
...
...
...
V 0.1–0.3
...
...
...
Other
270 / Stainless Steels for Design Engineers
S30909
S30940
S30941
DIN 1.4828
S31000
S31008
S31009
S31040
S31041
S31042
S31050
DIN 1.4841
S31400
S31600
S31609
S31620
S31603
S31653
S31651
S31635
S31700
S31703
S31753
S31725
S31726
S32100
S32109
S33000
N08800
S35125
S33400
S34700
S34709
S34800
S34809
S37000
S38400
S35315
309H
309Cb
309HCb
309Si
310
310S
310H
310Cb
310HCb
310HCbN
310MoLN
310Si
314
316
316H
316F
316L
316LN
316N
316Ti
317
317L
317LN
317LM
317LMN
321
321H
330
332
332Mo*
334
347
347H
348
348H
370
384
353MA
24.0–26.9
0.12–0.18
0.08
12.5–14.5
15.0–17.0
0.005
...
17.0–19.0
17.0–19.0
0.03–0.05
0.08
...
0.08
17.0–19.0
...
...
0.08
17.0–20.0
19.0–23.0
20.0–23.0
18.0–20.0
17.0–19.0
17.0–19.0
22.0–24.0
22.0–24.0
22.0–24.0
19.0–21.0
24.0–26.0
24.0–26.0
24.0–26.0
24.0–26.0
24.0–26.0
24.0–26.0
24.0–26.6
24.0–26.0
23.0–26.0
16.0–18.0
16.0–18.0
16.0–18.0
16.0–18.0
16.0–18.0
16.0–18.0
16.0–18.0
18.0–20.0
18.0–20.0
18.0–21.0
18.0–20.0
17.0–20.0
17.0–19.0
Cr
0.04–0.10
...
...
0.10
...
...
0.10
0.40–0.10
0.08
0.03
0.10
0.08
0.08
...
...
...
0.11
...
...
...
...
...
0.15–0.35
0.09–0.15
0.11
...
0.10
...
...
0.10
0.10–0.16
0.10–0.16
0.10
0.10
0.10
0.10–0.22
0.10
0.10–0.20
0.10
N
0.04–0.10
0.08
0.04–0.10
0.04–0.10
0.25
0.08
0.04–0.10
0.08
0.04–0.10
0.04–0.10
0.02
0.20
0.25
0.08
0.04–0.10
0.08
0.03
0.03
0.08
0.08
0.08
0.03
0.03
0.03
0.03
0.08
C
34.0–36.0
14.5–16.5
17.0–19.0
9.0–13.0
9.0–13.0
9.0–13.0
34.0–37.0
30.0–35.0
31.0–35.0
19.0–21.0
9.0–13.0
9.0–12.0
12.0–15.0
12.0–16.0
12.0–16.0
11.0–13.0
19.0–22.0
19.0–22.0
19.0–22.0
19.0–22.0
19.0–22.0
19.0–22.0
20.5–23.5
19.0–22.0
19.0–22.0
10.0–14.0
10.0–14.0
10.0–14.0
10.0–14.0
10.0–14.0
10.0–14.0
10.0–14.0
11.0–15.0
11.0–15.0
11.0–15.0
13.5–17.5
13.5–17.5
9.0–12.0
Ni
1.0
(continued)
1.65–2.35
2.0
2.0
2.0
2.0
2.0
1.5
1.0–1.5
1.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
3.0–4.0
3.0–4.0
3.0–4.0
4.0–5.0
4.0–5.0
2.0
Mn
...
1.5–2.5
...
...
...
...
...
...
2.0–3.0
...
...
...
...
...
...
...
...
...
...
...
...
...
1.6–2.6
...
...
2.0–3.0
2.0–3.0
1.75–2.5
2.0–3.0
2.0–3.0
2.0–3.0
2.0–3.0
2.0
2.0
2.0
2.0
2.0
...
Mo
Composition, %
0.6–1.0
0.5–1.0
1.00
1.00
1.00
1.00
0.75–1.50
1.00
0.75
1.00
1.00
1.00
...
...
0.045
0.045
0.045
0.045
0.030
...
...
...
0.045
0.045
P
0.045
0.040
0.040
0.040
0.045
0.045
0.045
0.045
0.045
0.045
0.020
0.045
0.045
0.045
0.045
0.20
0.045
0.045
0.045
0.045
0.045
0.045
0.045
0.045
0.045
0.045
Si
0.75
1.00
1.00
1.50–2.50
1.00
1.00
1.00
1.00
1.00
1.00
0.5
1.50–2.50
1.50–3.00
1.00
1.00
1.00
1.00
1.00
1.00
1.00
1.00
1.00
1.00
1.00
0.75
1.00
Notes: All compositions include Fe as balance. Single values are maximum values unless otherwise noted. (a) Unified Number System, UNS numbers are S or N followed by 5 digits.
Designation(a)
Name
Table A1.1 (continued)
S
...
...
0.030
0.030
0.030
0.030
0.030
...
...
...
0.030
0.030
0.030
0.030
0.030
0.015
0.030
0.030
0.030
0.030
0.030
0.030
0.010
0.015
0.030
0.030
0.030
0.10 min
0.030
0.030
0.030
0.030
0.030
0.030
0.030
0.030
0.030
0.030
Other
Other
...
...
Nb 10xC to 1.10 . . .
Nb 10xC to 1.10 . . .
...
...
...
...
...
...
...
...
Nb 10xC to 1.10 . . .
Nb 10xC to 1.10. . .
Nb 10xC to 1.10 . . .
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
5xC to 0.70 Ti . . .
...
...
...
...
0.030 max P
...
...
...
...
...
Ti 5x(C+N)
...
to 0.70
Ti 4x(C+N)
...
to 0.70
...
...
Ti 0.15–0.60
Al 0.15–0.60
Nb 0.25–0.60 . . .
Ti 0.15–0.60
Al 0.15–0.60
Nb+Ta 10xC
...
to 1.10
Nb+Ta 10xC
...
to 1.10
Nb+Ta 0xC
Co 0.2
to 1.10
Nb+Ta 8xC
...
to 1.0
Ti 0.10–0.40
Co 0.05
Nb+Ta 10xC
Co 0.2
to 1.10
Ce 0.03–0.10 . . .
Appendix 1: Compositions / 271
...
0.15–0.25
0.15–0.25
0.35–0.60
...
...
...
0.04
0.015
0.02
0.03
0.03
0.07
0.05
...
N08031
N08932
S31266
S32200
N08007
20Mod
24.0–26.0
23.0–26.0
20.0–23.0
19.0–22.0
21.0–23.0
22.0–24.0
26.0–28.0
20.0–22.5
19.0–21.0
26.0–28.0
25.0–29.0
20.0–25.0
20.0–22.0
19.25–21.50
20.0–22.0
22.0–24.0
18.0–21.0
19.0–21.0
22.5–25.0
22.0–26.0
26.0–28.0
20.0–22.0
20.0–22.0
19.0–23.0
19.0–23.0
19.50–20.50
23.0–25.0
24.0–25.0
19.0–21.0
Cr
24.0–26.0
21.0–24.0
23.0–27.0
27.5–30.5
25.0–27.0
26.0–28.0
24.0–26.0
balance
24.0–26.0
31.0–33.0
32.0–37.0
30.0–38.0
3.25–4.50
1.50–2.75
10.5–12.5
7.9–9.0
8.0–11.0
32.0–38.0
35.0–40.0
33.0–37.0
29.9–32.5
23.5–25.5
23.5–25.5
24.0–26.0
23.0–28.0
17.50–18.50
16.0–18.0
21.0–23.0
32.0–38.0
Ni
2.0
2.0
1.0
1.5
2.5
0.75
2.0
...
2.0
1.0
1.5
1.0
8.0–10.0
7.0–9.50
1.0–1.5
1.5–3.5
0.75–1.50
0.75–1.50
1.0
1.0
2.5
2.0
2.0
2.0
2.0
1.0
5.0–7.0
2.0–3.0
2.0
Mn
4.5–6.5
5.0–7.0
2.5–3.5
2.0–3.0
4.0–6.0
2.5–3.0
6.0–7.0
12.5–14.5
6.0–7.0
...
...
4.0–4.8
...
...
...
...
1.0–1.75
2.0–3.0
3.5–5.0
5.0–6.7
3.0–4.0
6.0–7.0
6.0–7.0
4.3–5.0
4.0–5.0
6.0–6.5
4.0–5.0
7.0–8.0
2.0–3.0
Mo
Composition, %
Si
0.40
1.0
0.5
1.5
1.0
0.75
0.30
0.08
0.50
0.03
1.0
0.6–1.0
0.25
0.25
0.70–1.25
0.69–0.90
0.03–0.80
1.00
0.50
0.50
1.00
1.00
1.00
1.00
1.00
0.80
1.0
0.5
1.0
Notes: All compositions include Fe as balance. Single values are maximum values unless otherwise noted. (a) Unified Number System, UNS numbers are S or N followed by 5 digits.
0.15–0.25
0.015
0.02
N06022
N08926
AL 22
Cronifer
1925 hMo
Cronifer 2328
Nicrofer
3127HMo
URSB-8
B66
NIC 25
CN-7M
N08320
...
...
...
0.28–0.50
0.20–0.40
0.15–0.25
0.28–0.38
...
...
...
...
...
...
0.18–0.25
...
...
0.18–0.22
0.40–0.60
0.45–0.55
...
0.04–0.08
0.06–0.10
0.08
0.48–0.58
0.50–0.60
0.15–0.25
0.28–0.38
0.28–0.36
0.07
0.03
0.03
0.02
0.035
0.030
0.04
0.02
0.02
0.03
0.02
0.07
S33228
S25045
S35135
S63008
S63012
S63017
S63018
S63198
N08020
N08024
N08026
N08028
N08366
N08367
N08700
N80904
S31254
S34565
S32654
N08020
AC66
Incoloy 803
Incoloy 864
21-4N
21-2N
21-12N
23-8N
19-9DL
20Cb-3
20Mo-4
20Mo-6
Sanicro 28
AL-6X
AL-6XN
JS-700
904L
254SMO
4565
654SMO
AL 20
N
C
Designation(a)
Name
Table A1.1 (continued)
P
...
...
...
...
...
0.045
...
...
...
...
...
...
0.045
0.050
0.045
0.045
...
0.045
0.0035
0.030
0.030
0.030
0.030
0.040
0.045
0.030
...
0.030
...
S
...
...
...
...
...
0.035
...
...
...
...
0.015
0.015
0.015
0.030
0.030
0.030
...
0.035
0.035
0.030
0.030
0.030
0.030
0.030
0.035
0.010
...
0.005
...
Other
Other
3.0–4.0 Cu
...
1.0–2.0 Cu
0.5–3.0 Cu
2.5–3.5 Cu
1.0–1.4 Cu
...
W 1.0-3.0
...
...
...
Ti 0.4-0.7
...
0.05–0.10 Ce Al 0.025
Ti 0.15–0.60
0.15–0.60 Al
Ti 0.4–1.0
...
...
...
...
...
...
...
...
...
0.1–0.35 Ti
0.25–0.60 Nb
8xC to 1.00 Nb Cu 3.0–4.0
Nb 0.15-0.35 Cu 0.5–1.5
8xC Nb
2.0–4.0 Cu
...
0.6–1.4 Cu
0.030 P
...
0.040 P
...
8xC to 0.5 Nb 0.5 Cu
1.0–2.0 Cu
...
0.5–1.0 Cu
...
0.10 Nb
...
0.3–0.6 Cu
...
3.0–4.0 Cu
8xC to 1.00
Nb+Ta
2.5 Co,0.35 V W 2.5-3.5
0.5–1.5 Cu
...
272 / Stainless Steels for Design Engineers
S40910
S40920
409
409
...
0.01
0.030
0.020
S43035
S43932
439
439LT
ATI, AK
439 HP
alloys
439 ultraform
S46800
468
18 Cr–Cb
13.0
...
0.07
0.12
0.10
...
0.03
0.04
18.0
18.0–20.0
17.5
17.0–19.0
17.0–19.0
16.0–18.0
16.0–19.5
...
...
...
...
0.04
14.0–16.0
16.0–18.0
16.0–18.0
...
0.08
0.12
0.12
0.12
13.5
...
13.0
12.0
0.015
10.5–11.7
11.35
0.025
0.025
0.020
0.03
0.015
10.5–11.7
10.5–11.7
10.5–11.75
10.5–11.7
10.5–11.75
Cr
11.5–14.5
12.0–13.0
10.5–11.75
...
0.50
0.2
0.50
0.50
...
1.00
0.75
0.75
...
...
...
...
...
0.5–1.0
0.20
0.50
0.50
0.05
0.50
0.50
Ni
0.60
...
0.50
Notes: All compositions include Fe as balance. Single values are maximum values unless otherwise noted
AK alloy
typical
0.01
0.03
S43023
S43036
430Se
430Ti
429
430
430F
4724
Alfa II
Alfa I
12 SR
0.03
0.010
S40975
AK alloy
typical
AK alloy
typical
ATI alloy
typical
ATI alloy
typical
Outukumpu
typical
S42900
S43000
S43020
409Ni
11 Cr–Cb
...
0.06
S40940
409Cb
0.020
0.020
0.030
0.030
N
...
...
...
0.02
0.02
0.03
0.03
C
0.08
0.05
0.08
409 ultraform AK alloy
S40930
466
405
400
409
designation
S40500
AK alloy
S40900
UNS
Composition of ferrite stainless steels
Name
Table A1.2
0.30
1.00
0.35
1.00
1.00
1.25
1.00
1.00
1.00
1.25
0.70
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
0.45
1.00
0.45
1.0
1.00
1.00
1.00
1.00
1.00
1.00
1.0
0.03
0.03
...
1.0
1.30
...
0.040
...
0.040
0.040
0.060
0.040
0.040
0.040
0.060
...
...
...
...
0.040
...
...
...
0.040
1.00
1.00
1.0
0.040
0.040
P
0.040
...
0.045
1.0
1.00
Si
1.00
1.00
1.00
Composition, %
Mo
...
...
...
(continued)
0.035
0.035
...
1.00
0.25
1.00
0.75
1.00
1.00
1.00
Mn
1.00
1.00
1.00
...
0.030
...
0.030
0.030
0.030
0.030
0.150
min
0.060
0.030
...
...
...
...
0.030
...
...
...
0.030
0.030
0.030
S
0.030
...
0.045
Ti+Nb:0.20+
4x(C+N)
to 1.10
0.25
...
0.20+4x(C+N)
to 1.10
0.20+4x(C+N)
to 1.10
0.20+4x(C+N)
to 0.75
Ti+Nb
0.35
...
...
...
...
0.40
0.40
0.30
Ti
...
...
6x(C+N)
to 0.75
6x(C+N)to 0.5
8x(C+N)
to 0.15–0.50
8x(C+N)
0.8+ 8x(C+N)
Ti+Nb
10xC
to 0.75 Nb
...
...
0.55
...
...
...
...
...
...
...
...
...
...
...
...
0.60
...
0.35
...
...
...
0.17
...
Nb
...
...
...
...
...
...
Al 0.15
...
SE 0.15
AL 0.15
...
...
...
Al 1.0
AL 4.0
AL 3.0
AL 1.2
...
...
...
...
...
...
...
Other
Al 0.10–0.30
Al 0.25 Al
...
Appendix 1: Compositions / 273
S44735
S44600
29-4C
446
0.25
...
0.035
23.0–27.0
28.0–30.0
25.0–27.0
24.5–26.0
25.0–27.5
22.0
24.0
20.0
17.5–19.5
18.0–23.0
17.3
17.5–18.5
16.0–18.0
16.0–18.8
18.0
17.30
Cr
0.6
0.5
1.5–3.5
3.5–4.5
0.50
0.3
...
0.25
1.0
0.6
0.3
...
...
...
...
Ni
0.25
Notes: All compositions include Fe as balance. Single values are maximum values unless otherwise noted.
0.20
0.025
0.025
S44660
Sea–cure
0.015
0.01
0.035
...
0.03
0.025
...
0.08
S44635
...
0.035
...
0.015
0.01
Monit
433
0.025
0.20
0.01
S44200
ATI alloy
typical
S44400
...
...
...
0.12
0.12
0.030
...
N
. . ..
0.08
C
0.015
S44100
designation
AK alloy
typical
Outukumpu
typical
S43400
S43600
UNS
(continued) of ferrite stainless steels
Composition
ATI alloy
typical
Outukumpu
4762
typical
453
ATI alloy
typical
E-Brite, 26-1 S44627
444, YUS
190-EM
441,
4509,
430J1L
442
436S
434
436
4742
18SR
Name
Table A1.2
1.50
1.00
1.00
1.00
0.40
0.3
0.7
0.30
1.0
1.0
0.20
1.00
1.0
1.0
0.7
Mn
0.30
...
3.5–4.5
2.5–3.5
3.5–4.5
0.75–1.25
...
...
...
0.75–1.25
...
1.2
...
0.75–1.25
0.75–1.25
...
1.00
0.75
1.00
0.75
0.40
0.3
1.4
0.4
1.0
1.0
0.4
1.0
1.0
1.0
1.3
Si
. . ..
Composition, %
Mo
. . ..
0.040
0.040
0.040
0.040
0.020
...
...
...
0.040
0.040
...
0.040
0.040
0.040
...
P
. . ..
0.030
0.030
0.030
0.030
0.020
...
...
...
0.030
0.030
...
0.030
0.030
0.030
...
S
...
...
...
...
Ti
...
Ti+Nb:0.20+
4x(C+N)
to 0.80
Ti+Nb:0.20+
4x(C+N)
to 0.80
Ti+Nb:0.20+
4x(C+N)
to 0.80
...
0.02
...
Ti+Nb:0.20+
4x(C+N)
to 0.80
...
...
8x(C+N) min
0.1–0.6
0.25
...
...
...
...
0.5–0.20
...
...
10x(C+N)
...
...
...
...
Nb+Ta
5xC:0.70
9xC
0.3–1.0
...
Nb
...
...
...
...
0.60 Al
0.10 REM
0.2 Cu
0.5 Cu+Ni
...
Al 1.5
...
...
...
...
...
...
...
Al 1.0
Other
Al 1.7
274 / Stainless Steels for Design Engineers
designation
S40300
S41000
S41003
S41008
S41040
S41003
S41400
S41425
S41500
S41600
S41623
S41800
S42000
DIN 1.4116
nominal
S42020
S42023
S42200
S42400
S42500
...
...
JFE nominal
JFE nominal
Nippon nominal
Nippon nominal
JFE nominal
JFE nominal
S43100
S44002
S44003
S44004
S44020
S44023
Nominal PM
Nominal PM
Nominal PM
Nominal PM
Nominal PM
Nominal PM
Nominal PM
UNS
0.15 min
0.15 min
0.20–0.25
0.06 max
0.08–0.20
0.50–0.55
0.30 max
0.025
0.025
0.03
0.02
0.01
0.01
0.20 max
0.60–0.75
0.75–0.95
0.95–1.20
0.95–1.20
0.95–1.20
1.15
1.05
1.05
1.05
1.45
2.15
2.20
C
0.15 max
0.15 max
0.03
0.08
0.18 max
0.030 max
0.15 max
0.05
0.05 max
0.15 max
0.15 max
0.15–0.20
0.15 min
0.50
...
...
...
...
...
...
...
...
...
0.040
0.015
0.010
0.010
...
...
...
...
...
...
...
...
...
...
...
...
...
N
...
...
...
...
...
...
...
0.06 0.12
...
...
...
...
...
...
12.0–14.0
12.0–14.0
11.0–13.5
12.0–14.0
14.0–16.0
13.0–14.0
12.0–14.0
13.0
13.0
12.7
12.3
11.0
12.0
15.0–17.0
16.0–18.0
16.0–18.0
16.0–18.0
16.0–18.0
16.0–18.0
14.5
14.0
14.0
14.0
14.0
17.0
13.0
Cr
11.5–13.5
11.5–13.5
10.5–12.5
11.5–13.5
11.5–13.5
10.5–12.5
11.5–13.5
12.0–15.0
11.5–14.0
12.0–14.0
12.0–14.0
12.0–14.0
12.0–14.0
14.5
Composition of martensitic stainless steels
Notes: All compositions include Fe as balance. Single values are maximum values unless otherwise noted
420F
420FSe
422
424
425
425 mod
Trinamet
HP13Cr-1
HP13Cr-2
NT-CRS
NT-CRSS
KL-12Cr
KL-HP 12Cr
431
440A
440B
440C
440F
440FSe
BG-42
ATS-34
14-4 CrMo
154 CM
CPM S30V
CPM S60V
CPM S90V
403
410
410S
410
410Cb
412
414
414 mod
415
416
416Se
418
420
4116
Name
Table A1.3
...
...
0.50–1.00
3.50–4.50
1.00–2.00
0.50
...
4.0
5.0
4.5
5.8
2.4
5.5
1.25–2.50
...
...
...
0.75
0.75
...
...
...
...
...
...
...
Ni
...
...
1.5
...
...
1.5
1.25–2.50
4.0–7.0
3.50–5.50
...
...
1.80–2.20
...
...
1.25
1.25
1.00
0.50–1.0
1.00
1.00
1.00
0.45
0.45
1.45
2.0
...
...
1.00
1.00
1.00
1.00
1.25
1.25
...
0.4
0.5
0.45
...
0.40
...
1.00
1.00
0.75
0.30–0.60
1.00
1.00
1.00
...
...
...
...
...
...
1.00
1.00
1.00
1.00
1.00
1.00
0.3
0.35
0.3
0.3
...
...
...
Mo
0.50
1.00
...
1.00
1.00
1.00
1.00
0.60
0.60
1.00
1.00
0.50
1.00
...
Composition, %
Mn
1.00
1.00
1.00
1.50
1.00
1.50
1.00
0.5–1.0
0.50–1.0
1.25
1.25
0.50
1.00
...
1.00
1.00
0.75
0.30–0.6
1.00
1.00
1.00
...
...
...
...
...
...
1.00
1.00
1.00
1.00
1.00
1.00
0.3
0.35
0.3
0.3
...
...
...
Si
0.50
1.00
...
1.00
1.00
1.00
1.00
0.60
0.60
1.00
1.00
0.50
1.00
...
0.040
0.040
0.040
0.040
0.040
0.040
0.040
...
...
...
...
...
...
0.040
0.040
0.040
0.040
0.040
0.060
...
...
...
...
...
...
...
P
0.040
0.040
0.040
0.040
0.040
0.040
0.040
0.040
0.040
0.060
0.060
0.040
0.040
0.040
0.15
0.06
0.030
0.030
0.010
0.030
0.030
...
...
...
...
...
...
0.030
0.030
0.030
0.030
0.10–0.35
0.060
...
...
...
...
...
...
...
S
0.030
0.030
0.030
0.030
0.030
0.030
0.030
0.005
0.030
0.15 0.30
0.060
0.030
0.030
...
...
Se 0.15 min
0.75–1.25 W
...
...
...
CU 2.0–3.0
...
...
1.5 Cu
1.5 Cu
0.5 Cu
...
...
...
...
...
...
Se 0.15 min
1.2 V
...
...
...
4.0 V
5.5 V
9.0 V
Other
...
...
...
...
Nb 0.05–0.30
...
...
Cu 0.30
...
...
Se 0.15 min
W 2.50–3.50
...
...
Appendix 1: Compositions / 275
0.21
0.02
0.01
0.09
0.09
0.07–0.11
0.10–0.15
0.08
0.08
0.07
0.07
0.05
0.05
0.05
...
...
...
...
...
0.07–0.13
0.07–0.13
...
...
...
...
0.010
...
...
10.0
11.0–12.50
10.5–11.50
16.0–18.0
14.0–16.0
16.0–17.0
15.0–16.0
13.5–16.0
15.0–17.0
115.5–17.5
14.0–15.5
12.25–13.25
14.0–16.0
11.0–12.50
Cr
5.5
10.75–11.25
7.5–8.5
6.5–7.75
6.5–7.75
4.0–5.0
4.0–5.0
14.0–27.0
6.0–7.5
3.0–5..0
3.5–5.5
7.5–8.5
5.0–7.0
7.5–9.5
Ni
0.10
0.25
0.50
1.00
1.00
0.50–1.25
0.50–1.25
2.00
1.0
1.0
1.0
0.2
1.0
0.50
Mn
S32900
S31200
S31260
S31500
S31830
S32001
S32003
S32101
S32205
S32304
S32520
S32550
S32750
S32760
S32906
S32950
S39274
S39277
Designation
0.08
0.03
0.03
0.30
0.03
0.03
0.03
0.04
0.03
0.03
0.03
0.04
0.03
0.03
0.03
0.03
0.03
0.025
C
...
0.14–0.20
0.10–0.30
0.05–0.10
0.08–0.20
0.05–0.17
0.14–0.20
0.20–0.25
0.14–0.20
0.05–0.20
0.20–0.35
0.10–0.25
0.20–0.30
0.20–0.30
0.30–0.40
0.15–0.35
0.24–0.32
0.23–0.33
N
23.0–28.0
24.0–26.0
24.0–26.0
18.0–19.0
21.0–23.0
19.5–21.5
19.5–21.–
21.0–22.0
22.0–23.0
21.5–23.5
24.0–26.0
24.0–27.0
24.0–26.0
24.0–26.0
28.0–30.0
26.0–29.0
24.0–26.0
24.0–26.0
Cr
Composition of selected duplex stainless steels
2.5–5.0
5.5–6.0
5.5–7.5
4.25–5.25
2.5–3.5
1.0–3.0
3.0–4.0
1.35–1.70
4.5–6.5
3.0–5.0
5.5–8.0
6.0–8.0
6.0–8.0
6.0–8.0
5.8–7.5
3.5–5.2
6.0–8.0
6.5–8.0
Ni
Notes: All compositions include Fe as balance. Single values are maximum values unless otherwise noted
329
44LN
DP3
3RE60
2205 (old)
19 D
2003
2101
2205
2304
Uranus 52N+
255
2507
Zeron 100
2906
7-Mo Plus
DP3W
AF 918
Name
Table A1.5
2.0
0.75–1.25
4.5–5.5
...
2.0–3.0
2.5–3.25
2.5–3.25
1.0–1.5
...
...
...
2.0–2.5
0.5–1.0
0.5
Mo
Composition, %
1.0
2.0
1.0
1.2–2.0
2.0
4.0–6.0
2.0
4.0–6.0
1.0
2.5
1.5
1.5
1.2
1.0
0.8–1.5
2.0
1.0
0.8
Mn
1.0–2.0
1.2–2.0
2.5–3.5
2.5–3.0
2.5–3.5
...
1.5–2.0
0.1–0.8
3.0–3.5
...
3.0–5.0
2.9–3.9
3.0–5.0
3.0–5.0
1.5–2.6
1.0–2.5
2.5–3.5
3.0–4.0
Mo
Composition, %
Notes: All compositions include Fe as balance. Single values are maximum values unless otherwise noted. (a) Nominal value
...
S46500
S17600
S17700
S15700
S35000
S35500
S66286
Custom 465
Custom 475
17-7 PH
15-7 PH
AM-350
AM-355
A-286
Ferrium S53(a)
S17600
S17400
S15500
S13800
S45000
S45500
Stainless W
17-4 PH
15-5 PH
13-8 PH
Custom 450
Custom 455
N
Designation
Name
C
Composition of selected precipitation–hardenable stainless steels
Table A1.4
Si
Si
0.75
1.0
0.75
1.4–2.0
1.0
1.0
1.0
1.0
2.0
1.0
0.8
1.0
0.8
1.0
0.5
0.6
0.8
0.8
0.10
0.25
0.50
1.00
1.00
0.50
0.50
1.00
1.00
1.00
1.00
0.10
1.00
0.50
P
...
0.015
0.015
0.040
0.040
0.040
0.040
0.040
0.040
0.040
0.040
0.010
0.030
0.040
P
0.040
0.045
0.030
0.030
0.030
0.040
0.040
0.040
0.030
0.040
0.035
0.040
0.035
0.030
0.030
0.035
0.030
0.030
S
...
0.010
0.010
0.030
0.030
0.030
0.030
0.030
0.030
0.030
0.030
0.008
0.030
0.030
Other
0.030
0.030
0.020
0.030
0.020
0.030
0.030
0.030
0.020
0.040
0.020
0.030
0.020
0.010
0.030
0.010
0.020
0.020
S
Cu
Other
...
...
0.1–0.5
...
...
...
...
...
...
...
...
...
...
0.5–1.0
...
...
1.5–2.5
0.8–1.2
W
...
Al 1.0–1.5
Al 0.75–1.5
Al 0.75–1.5
...
...
Al 0.35
B 0.001–0.010
W 1, V 0.3
Al 0.4
Nb .015–0.45
Nb 0.15–0.45
Al 0.90–1.35
Nb 8XC
Nb +Ta 0.1–0.5
0.75
1.0
0.75
1.4–2.0
1.0
1.0
1.0
1.0
2.0
1.0
0.8
1.0
0.8
1.0
0.5
0.6
0.8
0.8
Ti 0.4–1.2
Cu 3.0–5.0
Cu 2.5–4.5
...
Cu 1.25–1.75
Cu 1.5–2.5,
Ti 0.08–1.4
Ti 1.50–1.80
Co 8.0–9.0
...
...
...
...
Ti 1.9–2.35
V 0.10–.050
Co 14
276 / Stainless Steels for Design Engineers
Wrought
equivalent(a)
410
...
420
420F
431,442
446
...
S41500
422
17–4PH
15–5PH
2205 (S32205)
255 (S32550)
(S32760)
...
...
...
2507 (S32750)
...
312
304L
316L
316LN
304
347
316
304H
316H
316H
Nitronic™60
316
303
302
Nitronic™50
317
308
309S
309H
309
254SMO™
310
904L
AL–6XN®
J91150
J91151
J91153
J91154
J91803
J92613
J91650
J91540
J91422
J92180
J92110
J92205
J93373
J93380
J93370
J93372
J93371
J93404
J93345
J93423
J92500
J92800
J92700
J92600
J92710
J92900
J92590
J92901
J92971
J92972
...
J92701
J92602
J93790
J93000
J93001
J93400
J93401
J93402
J94653
J94202
J94652
J94651
0.15
0.15
0.40
0.2–0.4
0.30
0.30
0.06
0.06
0.20–0.28
0.07
0.07
0.03
0.03
0.03
0.04
0.04
0.06
0.03
0.08
0.30
0.03
0.03
0.03
0.08
0.08
0.08
0.04–0.10
0.04–0.10
0.10
0.10
0.12
0.16
0.20
0.06
0.08
0.12
0.08
0.04–0.10
0.20
0.025
0.20
0.03
0.03
C
...
...
...
...
...
...
...
...
...
0.05
0.05
0.10–0.30
0.22–0.33
0.20–0.30
...
0.10–0.25
0.15–0.25
0.10–0.30
0.10–0.30
...
...
...
0.10–0.20
...
...
...
...
...
...
0.08–0.18
...
...
...
0.20–0.40
...
...
...
...
...
0.18–0.24
...
...
0.18–0.24
N
11.5–14.0
11.5–14.0
11.5–14.0
11.5–14.0
18.0–22.0
26.0–30.0
10.5–12.5
11.5–14.0
11.0–12.5
15.5–17.7
14.0–15.5
21.0–23.5
24.0–26.7
24.0–26.0
24.5–26.5
24.5–26.5
24.0–27.0
24.0–26.0
22.5–25.5
26.0–30.0
17.0–21.0
17.0–21.0
17.0–21.0
18.0–21.0
18.0–21.0
18.0–21.0
18.0–21.0
18.0–21.0
15.0–18.0
16.0–18.0
18.0–21.0
18.0–21.0
18.0–21.0
20.5–23.5
18.0–21.0
20.0–23.0
22.0–26.0
22.0–26.0
22.0–26.0
19.5–20.5
23.0–27.0
20.0–22.0
20.0–22.0
Cr
1.0
1.0
1.0
1.0
2.0
4.0
6.0–8.0
3.5–4.5
0.5–1.0
3.6–4.6
4.5–5.5
4.5–6.5
5.6–6.7
6.5–8.5
4.75–6.0
4.7–6.0
4.0–6.0
6.0–8.0
8.0–11.0
8.0–11.0
8.0–12.0
8.0–12.0
9.0–13.0
8.0–11.0
9.0–12.0
9.0–12.0
8.0–11.0
9.0–12.0
13.0–16.0
8.0–9.0
9.0–12.0
9.0–12.0
8.0–11.0
11.5–13.5
9.0–13.0
10.0–13.0
12.0–15.0
12.0–15.0
12.0–15.0
17.5–19.5
19.0–22.0
23.0–27.0
23.0–27.0
(continued)
Ni
0.50(b)
0.15–1.00
0.50(b)
...
...
...
...
0.4–1.0
0.9–1.25
...
...
2.5–3.5
2.9–3.8
3.0–4.0
1.75–2.25
1.75–2.25
1.75–2.25
4.0–5.0
3.0–4.5
...
...
2.0–3.0
2.0–3.0
...
...
2.0–3.0
...
2.0–3.0
1.75–2.25
...
2.0–3.0
1.5
...
1.5–3.0
...
...
...
...
...
6.0–7.0
...
4.5–5.5
6.0–7.0
Mo
Composition, %
1.00
1.00
1.00
1.00
1.00
1.00
0.50
1.00
0.5–1.0
0.70
0.70
1.50
1.20
1.00
1.00
1.00
1.00
1.50
1.00
1.50
1.50
1.50
1.50
1.50
1.50
1.50
1.50
1.50
1.50
7.0–9.0
1.50
1.50
1.50
4.0–6.0
1.50
1.50
1.50
1.50
1.50
1.20
2.00
2.00
2.00
Mn
1.50
0.65
1.50
1.50
1.50
1.50
1.00
1.00
1.00
1.00
1.00
1.00
1.10
1.00
1.00
1.00
1.00
1.00
1.50
2.00
2.00
2.00
1.50
2.00
2.00
2.00
2.00
1.50
1.50
3.5–4.5
2.00
2.00
2.00
1.00
1.50
2.00
1.50
2.00
2.00
1.00
2.00
1.00
1.00
Si
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.17
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
P
Notes: All compositions include Fe as balance. Single values are maximum values unless otherwise noted. (a) The wrought equivalent composition is not the same as the cast. (b) Mo is not an intentional addition.
CA-15
CA-15M
CA-40
CA-40F
CB-30
CC-50
CA-6N
CA-6NM
CA-28MWV
CB-7Cu-1
CB-7Cu-2
CD-3MN
CD-3MCuN
CD-3MWCuN
CD-4MCu
CD-4MCuN
CD-6MN
CE-3MN
CE-8MN
CE-30
CF-3
CF-3M
CF-3MN
CF-8
CF-8C
CF-8M
CF-10
CF-10M
CF-10MC
CF-10SMnN
CF-12M
CF-16F
CF-20
CG-6MMN
CG-8M
CG-12
CH-8
CH-10
CH-20
CK-3MCuN
CK-20
CN-3M
CN-3MN
UNS
designation
Composition of Alloy Casting Institute (ACI) heat– and corrosion–resisting casting alloys
(continued)
Corrosion–resisting alloys
Name
Table A1.6
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
S
...
...
...
0.20–0.40 S
...
...
...
...
0.9–1.25 W, 0.2–0.3V
2.5–3.2 Cu, 0.2–0.35Nb
2.5–3.2 Cu, 0.2–0.35Nb
Cu 1.0
Cu 1.4–1.9
Cu 0.5–1.0, W 0.5–1.0
Cu 2.75–3.25
Cu 2.75–3.25
Cu 1.75–2.5
...
...
...
...
...
...
...
Nb 8XC min
...
...
...
(10xC)–1.2 Nb
...
...
Se 0.2–0.35
...
0.1–0.3 Nb, 0.1–0.3 V
...
...
...
...
...
Cu 0.5–1.0
...
...
...
Other
Appendix 1: Compositions / 277
0.20
0.50
0.50
0.20–0.50
0.20–0.40
0.20–0.50
0.20–0.50
0.20–0.60
0.25–0.35
0.35–0.45
0.20–0.60
0.20–0.50
0.35–0.75
0.45–0.55
0.35–0.75
0.25–0.35
0.35–0.75
0.35–0.75
0.35–0.75
J82090
J92605
J93005
J93403
J92603
J93505
J94003
J94224
J94203
J94204
N08604
J94213
N08705
...
N08605
N08603
N08005
N08006
N06050
504
446
327
312
302B
309
...
310
...
...
...
...
...
...
330
...
...
...
...
...
...
...
...
...
0.2
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
...
N
8–10
26–30
26–30
26–30
19–23
24–28
26–30
24–38
23.0–27.0
23.0–27.0
28–32
19–23
24–28
24–28
13–17
13.0–17.0
17–21
10–14
15–19
19.0–22.0
18.0–20.0
19.0–21.0
Cr
...
4 max
4–7
8–11
9–12
11–14
14–18
18–22
19.0–22.0
19.0–22.0
18–22
23–27
33–37
33–37
33–37
33.0–37.0
37–41
58–62
64–68
27.5–30.0
22.0–25.0
31.0–34.0
Ni
1.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.5
2.5
2.5
2.5
2.5
2.5
0.9–1.2
0.5(b)
0.5(b)
0.5(b)
0.5(b)
0.5(b)
0.5(b)
0.5(b)
0.5(b)
0.5(b)
0.5(b)
0.5(b)
0.5(b)
0.5(b)
0.5(b)
0.5(b)
0.5(b)
0.5(b)
0.5(b)
2.0–3.0
2.5–3.0
...
Mo
Composition, %
1.50
1.50
0.15–1.5
Mn
Si
0.35–0.65
1.0
1.5
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
2.0
1.50
3.50
0.50–1.5
P
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.035
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
Notes: All compositions include Fe as balance. Single values are maximum values unless otherwise noted. (a) The wrought equivalent composition is not the same as the cast. (b) Mo is not an intentional addition.
0.07
0.07
0.05–0.15
C
N08007
J94650
N08151
UNS
designation
320
...
...
Wrought
equivalent(a)
(continued)
CN-7M
CN-7MS
CT-15C
Heat resisting alloys
HA
HC
HD
HE
HF
HH
HI
HK
HK-30
HK-40
HL
HN
HP
HP-50WZ
HT
HT-30
HU
HW
HX
Name
Table A1.6
S
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.035
0.04
0.04
0.04
0.04
0.04
0.04
0.04
0.04
Other
...
...
...
...
...
...
...
...
...
...
...
...
...
W 4.0–6.0, Zr 0.1–1.0
...
...
...
...
...
Cu 3.0–4.0
Cu 1.5–2.0
Nb 0.5–1.5
278 / Stainless Steels for Design Engineers
Stainless Steels for Design Engineers
Michael F. McGuire, p 279-280
DOI: 10.1361/ssde2008p279
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
APPENDIX 2
Physical and Mechanical Properties of
Select Alloys
Table A2.1
201
301
304
304L
305
316L
321
904L
AL6-XN©
409
430
439
468
410
2101
2003
2205
2507
Physical properties of major stainless steel engineering alloys
UNS
Density, kg/dm3
Modulus of
elasticity, GPa
S20100
S30100
S30400
S30403
S30400
S31603
S32100
N08904
N08367
S40920
S43000
S43035
S46800
S41000
S32101
S32003
S32205
S32750
7.86
8.03
7.90
7.90
7.90
8.00
7.92
7.95
8.06
7.76
7.70
7.70
7.76
7.65
7.8
7.72
7.8
7.8
207
193
200
200
200
200
193
190
200
200
200
200
200
200
200
210
200
200
Coefficient of
thermal exp.,
10–6 × K–1
Thermal
conductivity,
W/M·°K
Specific heat,
J/kg·°K
Electrical
resistivity,
Ω·mm2/m
16.6
16.6
16.6
16.6
16.6
16.5
16.6
15.3
15.3
10.5
10.3
10.2
10.5
10.5
13.5
13.5
14.6
12.5
16.3
16.3
16.3
16.3
16.3
14.6
16.3
13.2
11.8
25.0
23.9
24.2
25.0
24.9
17.0
17.0
16.5
13.5
502
500
500
500
500
480
500
460
474
477
460
460
477
460
500
510
500
500
0.67
0.73
0.72
0.72
0.72
0.74
0.72
0.95
0.89
0.60
0.60
0.63
0.60
0.56
0.80
0.80
0.80
0.80
280 / Stainless Steels for Design Engineers
Table A2.2
Typical minimum mechanical properties of representative stainless steel engineering alloys
Name
Condition
UNS
201
Annealed
S20100
201F
2B
301
Annealed
301 tensile
Tensile strength,
MPa
Elongation, %
Hardness
260 min
550 min
40 min
100 Rb max
S20100
330
700
51
89 Rb
S30100
205 min
515 min
40 min
95 Rb max
2D
1 hard
/4
S30100
320
850
49
88 Rb
S30100
580
900
32
25 Rb
301
1
/2 hard
S30100
815
1150
23
35 Rc
301
3
S30100
1000
1270
17
40 Rc
301
/4 hard
Full hard
S30100
1160
1380
12
42 Rc
301 sink
2D
S30100
270
690
57
82 Rb
304
Annealed
S30400
205 min
515 min
40 min
92 Rb max
304
Hot rolled
S30400
335
640
51
86 Rb
304
2D
S30400
265
625
55
81 Rb
304
2B
S30400
305
635
52
85 Rb
304
#4 polish
S30400
325
650
51
85 Rb
304
2BA
S30400
315
640
53
85 Rb
304
1
S30400
705
890
23
29 Rc
304L
/4 hard
Annealed
S30403
170 min
485 min
40 min
92 Rb max
304L
2D
S30403
255
590
53
80 Rb
304LT
2D
S30403
255
600
51
81 Rb
304DD
2D
S30400
270
610
55
82 Rb
304EDD
2D
S30400
260
600
56
78 Rb
305
Annealed
S30500
170 min
485 min
40 min
88 Rb max
305
2D
S30500
245
560
52
73 Rb
316L
Annealed
S31603
170 min
485 min
40 min
95 Rb max
316L
2B
S31603
310
595
51
82 Rb
321
Annealed
S32100
205 min
515 min
40 min
95 Rb max
321
2B
S32100
285
570
49
78 Rb
904L
Annealed
N08904
220 min
490 min
35 min
90 Rb max
904L
2B
N08904
270
605
50
79 Rb
AL6-XN©
Annealed
N08367
310 min
690 min
30 min
100 Rb max
AL6-XN©
2B
N08367
365
745
47
88 Rb
409
Annealed
S40920
170 min
380 min
20 min
88 Rb max
409
2D
S40920
260
440
31
60 Rb
430
Annealed
S43000
205 min
450 min
20 min
89 Rb max
430
2B
S43000
345
515
27
67 Rb
439
Annealed
S43035
205 min
415 min
22 min
89 Rb max
439
2D
S43035
315
455
32
76 Rb
468
Annealed
S46800
205 min
415 min
22 min
90 Rb max
468
2D
S46800
205
415
32
76 Rb
29-4C
Annealed
S44735
415 min
550 min
18 min
25 Rc max
29-4C
2D
S44735
550
650
20
20 Rc
410
Annealed
S41000
205 min
450 min
20 min
96 Rb max
410
2B
S41000
320
515
28
81 Rb
2101
Annealed
S32101
530 min
700 min
30 min
301
Yield strength, MPa
...
(a) Finish conditions: 2D is cold rolled, annealed, and pickled; 2B is 2D with an added temper mill pass (approximately 0.5% reduction); 2BA is cold rolled, bright annealed, and temper passed.
Stainless Steels for Design Engineers
Michael F. McGuire, p 281-283
DOI: 10.1361/ssde2008p281
Copyright © 2008 ASM International®
All rights reserved.
www.asminternational.org
APPENDIX 3
Introduction to Thermo-Calc and
Instructions for Accessing Free
Demonstration Version
WITHIN THE MAIN BODY of this textbook, a number of diagrams have been plotted
and attributed to a software package called
Thermo-Calc. The purpose of this appendix is
to give a brief introduction to Thermo-Calc, explain what it is, and what are its uses. Also provided are instructions for accessing a demonstration version of the software.
What Is Thermo-Calc?
Thermo-Calc (Ref 1) is a powerful, flexible
software package available from Thermo-Calc
Software AB for performing various kinds of
thermodynamic and phase diagram calculations
for multicomponent systems.
The software is based on the so-called CALPHAD (CALculation of PHAse Diagrams)
method (Ref 2), which describes mathematically the thermodynamics of a system through a
representation of the Gibbs energies of the different crystalline phases relevant to that system
and defined by the chemical composition of the
system. Thermo-Calc minimizes the total Gibbs
energy of the system with respect to various
constraints such as temperature, pressure, and
chemical composition and thus predicts the
most stable energy state (or equilibrium state)
that can form. By suspending certain phases
(i.e., manually removing certain selected phases
from the system and thus restricting the formation of such phases), Thermo-Calc can also be
used to investigate meta-stable equilibria-type
problems.
Thermo-Calc is used in conjunction with
thermodynamic databases containing polynomial functions that describe the Gibbs energies
of the different phases according to certain
models that take into consideration nonideal
chemical interactions in solution phases. These
databases are based on the critical evaluation of
thermodynamic and phase equilibria data for binary, ternary, and some higher-order systems,
which are then assembled into self-consistent
databases. Different databases are available for
different broad classifications of materials, systems, or applications. For example, there are
databases for steels and iron-based alloys; ironbased slags; nickel superalloys; aluminum,
magnesium, titanium, and zirconium alloys; cemented carbides; nuclear materials; and more.
Further information on the different databases
available can be found at the Thermo-Calc Web
site: www.thermocalc.com
The thermodynamic database for steels (Ref
3), as developed by Thermo-Calc Software AB,
was used in conjunction with Thermo-Calc for
all the calculations made during the preparation
of this book. The version of the database used
for these calculations contains data for 20 elements and 85 phases.
Although the databases are based primarily
on the critical assessments of binary, ternary,
and some quaternary systems, the CALPHAD
methodology provides a theoretical framework
on which extrapolations can be made to predict
the phase equilibria for higher-order, multicomponent systems (the higher the order of the
system, the weaker the nonideal interaction
282 / Stainless Steels for Design Engineers
parameters become). Thermo-Calc can therefore be used in conjunction with such databases
to make predictions for multicomponent systems and alloys of industrial relevance as illustrated by some of the examples given in the
main body of this book. These calculations can
be validated against real alloy data if this information is available but is not based on (or adjusted to fit) such higher-order alloy data.
Higher order in this sense means more than four
elements (i.e., larger than a quaternary system).
Four specific types of calculation can be performed using Thermo-Calc, although the range
of problems to which these can be applied is
broader:
1. Single-Point Equilibria: The temperature,
pressure, composition/activity of a component (or the amount of a phase) are fixed and
the stable or meta-stable equilibrium for the
specified conditions is calculated.
2. Step: The amount of one state variable parameter (or condition) can be changed, while
the other conditions remain fixed. For example, to see how the different phases and their
amounts and compositions would vary with
temperature for a given alloy, one would
“step in” temperature. Alternatively, one can
vary the composition of one of the components/elements and calculate how the phase
amounts change for a fixed temperature or
predict how the solidus or liquidus would
change with varying alloy composition.
3. Map: Two axis variables (such as temperature, pressure, composition, or activity of
the components) are changed at the same
time. Isoplethal sections are generated by
varying temperature and composition of one
of the components. Isothermal sections are
the result of varying the amounts of two of
the components for a fixed temperature. Examples of each of these kinds of diagrams
are given in the main body of the text.
4. Scheil: Thermo-Calc includes a ScheilGulliver model for nonequilibrium solidification and a modified Scheil model that considers partial equilibrium for components
that are selected by the user.
Applications of Thermo-Calc
Thermo-Calc is used around the world within
academia, in government research laboratories,
and by commercial industry. The software can
be used to perform calculations for most applications involving phase equilibria, meta-stable
equilibria, phase transformations, phase diagrams, and various thermodynamic properties,
as well as critical assessments and data evaluations for multicomponent systems.
While many types of calculations can be
made using Thermo-Calc, the software typically
is used to predict:
• Stable and meta-stable phase equilibria for
binary, ternary, and higher-order systems
(calculations for alloy compositions with 6,
10, 15 elements are not uncommon, as illustrated by some of the examples in the main
body of this book).
• Amounts of phases (mass, volume and mole
fractions) formed (phase balance) as a function of temperature, pressure, and composition and also the chemical compositions of
the phases formed
• Phase transformation temperatures such as
liquidus, solidus, and solvus temperatures.
Phase transformation temperatures can be
predicted based on the actual chemistry (not
nominal chemistry).
• Thermochemical data such as enthalpies,
heat capacity, and activities
• Driving forces for precipitation
• Phase diagrams (isothermal and isoplethal
sections for multicomponent, multiphase
systems as illustrated in this book)
• Molar volume, density, and thermal expansion
• Scheil-Gulliver (nonequilibrium) solidification simulations
Thermo-Calc is not restricted just to modeling the alloy. Complex systems representing
processing, for example, can also be considered. For example, another application is to calculate the carbon potential of multicomponent
gas phase systems as a function of composition,
temperature, and pressure and then predict what
phases an alloy might form at a given temperature when exposed to such a carbon potential.
Thermo-Calc can thus be applied to a number
of practical problems related to metallurgy, processing, in-service performance, etc. as summarized by:
• Alloy Design: Modification of alloy chemistries to improve properties or reduce costs
using calculations to guide which compositions may be most suitable before preparing
them for testing
Appendix 3: Introduction to Thermo-Calc / 283
• Heat Treatment: Prediction of formation
of problematic phases prior to thermal processing
• Casting: Calculation of liquidus and solidus
temperatures; calculation of thermodynamic
properties of the alloy for input into casting
modeling codes
• Welding and Joining: Prediction of the
phases formed at the joining of two dissimilar
materials or the interaction with filler material
• Quality Control: Investigation of properties
and phase balance within designated compositional tolerances
More examples are available in the literature
(search on key terms Thermo-Calc or CALPHAD). A list of published articles citing
Thermo-Calc is available at www.thermocalc.
com.
How to Obtain a Free Demonstration
Version of Thermo-Calc
Thermo-Calc is available in two formats:
Thermo-Calc Classic, which has a command
line interface and can be run under a number of
different operating systems (including Microsoft Windows and Linux/Unix), and
Thermo-Calc for Windows, which has an easyto-learn graphical user interface but only operates in the Microsoft Windows environment.
Demo versions are available for both of these
versions of the software.
The demo versions are free to use, subject to
the terms outlined in the Thermo-Calc Software
End User License Agreement. It should be
noted that the demo versions are limited to
using just three elements (whereas in the full
product the current upper limit is 40 elements)
and are supplied with only certain small databases that are for demonstration purposes and
the evaluation of the software only.
A link to register and download the demonstration version of the software can be accessed via a link on the Thermo-Calc web site
at www.thermocalc.com. All fields in the registration form should be completed before continuing to the download page, where further instructions regarding installation of the software
will be provided.
On installation of the software, additional
documentation, including a Users Guide/Examples manual in the form of PDF files will also be
installed. Technical support for the demo versions of the software is limited, but problems
related to installation or general inquiries can be
addressed by visiting www.thermocalc.com and
linking to their support.
The demo version will run for approximately
1 month on a single computer, and installation
on a network system is not supported. If you
wish to run the software after the demo license
has expired, it can be downloaded again (i.e.,
obtaining a new demo license).
REFERENCES
1. J.O. Andersson, T. Helander, L. Höglund,
P.F. Shi, and B. Sundman, Thermo-Calc and
DICTRA, Computational Tools for Materials
Science, Calphad, Vol 26, 2002, p 273–312,
2002
2. N. Saunders and A.P. Miodownik, Pergamon
Materials Series, CALPHAD (Calculation of
Phase Diagrams): A Comprehensive Guide,
1, Elsevier, 1998
3. TCFE5—TCS Steel/Fe-Alloys Database,
Version 5.0, 2007, Thermo-Calc Software
AB, www.thermocalc.com
Copyright © 2008 ASM International®. All rights reserved.
Stainless Steels for Design Engineers (#05231G)
www.asminternational.org
Index
A
acetic acid
corrosion rates for various alloys of, plus formic acid,
103(F)
duplex alloys, 102
isocorrosion curves in, 36(F)
isocorrosion performances of various alloys, 102(F)
acids, corrosion in, and bases, 31–36
adsorption-induced brittleness, 51, 53
adsorption-induced plasticity, 51, 52–53
aeration cells, differential, 38–39
aerobic bacteria, influencing corrosion, 55
aesthetic finishes
Chrysler Building in New York City, 196, 197(F), 213(F)
considerations, 217–219
flatness, 219
surface, 217–219
aging treatments, precipitation hardening stainless
steels, 168, 170
aggressive chemical agents, refinery applications, 255
Alloy Casting Institute (ACI)
composition of, heat- and corrosion-resisting casting
alloys, 277(T), 278(T)
naming system, 147
alloy design
avoiding unwanted phases, 94
duplex alloys, 92–94
Thermo-Calc software, 93, 282
alloying, sensitization, 47–48
alloying elements
alloys, 1
influence on alloy families, 147, 149
influence on corrosion rate in contaminated sulfuric acid,
33 (F)
influence on thermodynamic activity of C, N, S and O,
157(T)
influence on uniform corrosion, 29–30
martensitic precipitation-hardening stainless steels,
141(T)
martensitic stainless steels, 130, 131(T)
alloy oxidation, behaviors, 66
alloys
influence of, content on corrosion rate in hydrochloric
acid, 35(F)
isocorrosion curves for, in sulfuric acid, 33(F)
isocorrosion curves for, sulfuric acid with chlorides, 34(F)
alloy systems
alloying elements, 1
austenitic alloys, 69, 72
austenitic stainless family, 71(F)
body-centered cubic (bcc) phase, 1
families in perspective, 69–72
most widely used, 1
Schaeffler–Delong stainless steels constitution diagram,
5(F), 70(F)
alpha prime
formation kinetics for duplex alloys, 94, 96(F)
influence of, formation on hardness, 116(F)
iron-chromium phase, 8
martensite, 7–8, 73
transition temperature change with, formation with
aging, 98(F)
alteration, surface, 199
aluminum
inclusions and pitting, 40, 41
influence on thermodynamic activity of C, N, S and O,
157(T)
metal migration into acetic solution from, 234(T)
migration into acetic solution from stainless, aluminum
or carbon steel, 234(T)
oxidation resistance, 79, 226, 228
protective layer formation, 64–65
aluminum alloys, 1
aluminum oxide, 191
aluminum-titanium-nitride (AlTiN), 191
aluminum/titanium precipitates, possible, 138(F)
American Iron and Steel Institute (AISI), 240
American Petroleum Institute (API), 135
American Society for Testing and Materials (ASTM),
238–239
American Society of Mechanical Engineers (ASME), 266
anisotropy
deep drawing, 174
ferritic stainless steels, 120–121
Lankford ratio, 120, 175
stainless long products, 179
annealing
austenitic stainless steels, 162–164
bright, 198
286 / Index
annealing (continued)
deep drawing, 177
duplex stainless steels, 170–171
ferritic stainless steels, 165–166
long-term, of welds, 43
martensitic stainless steels, 166
precipitation hardening steels, 139
anode
electrochemical reactions, 12
polarization, 20–21, 23
appliances
facades, 239–240
kitchen, 237–240
laundry, 241
stainless steels commonly used for, 239(T)
architecture and construction
aesthetic considerations, 217–219
average chloride concentration in rainwater in United
States, 217(F)
balancing corrosion resistance, processing and economy,
214–215
balancing service environment, design and maintenance,
215, 217
cleaning methods for uncoated stainless steel, 220(T)
concrete reinforcing bar, 222
corrosion resistance, 213–214
design considerations, 216(F)
ecological considerations, 222
environment, 216(F)
fabrication and joining, 221
fabrication considerations, 221
fire resistance, 221–222
flatness, 219
grades recommended by expert system, 217(F)
graphic depicting low release of metal ions from 304
and 316 stainless steels, 222(F)
local weather pattern, 216(F)
maintenance, 220–221
maintenance schedule, 216(F)
roof, 219(F)
ranking common stainless steels by pitting resistance
equivalent number (PREN), 214(T)
repair, 221
rolled-on stainless steel finishes, 218(F)
salt exposure, 216(F)
special finishes, 218(F)
stainless steel selection expert system, 216(F)
surface finish aesthetics, 217–219
surface finish and corrosion resistance, 215
argon oxygen decarburization (AOD)
adoption, 70
alloy adjustment, 157
automotive industry, 225
cleanliness, 184
control of nitrogen in refining by, 92
ferrite, 4
first commercial use, 109
foundry practice, 154
inclusions in steel, 40
production process, 155
atmospheres, oxidation, 66–67
atmospheric corrosion
uniform corrosion, 36–37
atomic rearrangements, 2
attraction, interatomic, 2
austenite
alloying elements, 5
carbide precipitation, 9
carbon and nitrogen, 6, 9
diffusion rates, 6
face-centered cubic (fcc), 5
γ-austenite in precipitation hardening alloys, 138–139
interstitial elements, 6
lattice expansions, 6(F)
lean alloy of martensite and, 73–74
mechanical properties, 7
metastable state, 6
phase in duplex alloy at room temperature, 91
Schaeffler–Delong constitution diagram, 5(F), 70(F)
Schaeffler diagram, 202(F)
secondary type, 7
semiaustenitic precipitation-hardenable stainless steel,
143
sulfur and oxygen, 6
austenite conditioning, 169
austenitic-ferritic “C” alloys, 151–152
austenitic “H” alloys
high temperature HE–HP, 152–154
precipitation hardening stainless steel, 170
austenitic precipitation-hardenable stainless steels
cold work and aging, 146(F)
composition, 145(T)
corrosion resistance, 145–146
mechanical properties, 145
austenitic stainless steels
alloy families, 69–72
annealing, 162–164
automotive structural components, 229
carburization, 82
composition of, 270(T), 271(T)
composition of high-temperature, 82(T)
compositions of commonly used lean, 72(T)
compositions of corrosion-resistant, 86(T)
corrosion resistance ratings, 87(T)
corrosion-resistant alloys, 84–89
corrosive environments, 88–89
critical pitting temperature (CPT), 43, 44(F)
drawability, 176
ductility, 180
face-centered cubic (fcc), 174
family, 71(F)
forming limit diagram of carbon steel and, 176(F)
halogens, 82
heat exchangers, 246
high-temperature alloys, 79–83
high-temperature mechanical properties, 82–83
impact strength variation with temperature, 75(F)
intermetallic phases, 82, 203
isocorrosion curves for, in hydrochloric acid, 34(F)
kitchen appliances, 239
Index / 287
lean alloys, 72–78
machinability, 185
machining setup recommendations, 183(T)
martensite and austenite, 73–74
mechanical properties, 74–76, 82–83
mechanical properties after cold work and annealing, 163
metal migration into acetic solution from, 234(T)
nitriding, 82
oxidation resistance, 79–81
petroleum industry applications, 249(T)
pitting resistance equivalent number (PREN), 43, 78, 85
precipitation of carbides and nitrides, 76–78
recommended thermal processing temperatures, 162(T)
resilience and toughness of carbon steel vs. for automotive components, 229(T)
SCC (stress corrosion cracking), 49–50
secondary phases in, 82(T)
sensitization, 46–47
soaking, 161–162
stabilization, 78
stainless steel in shipping, 245
SCC performance, 262
stress-strain curve for single crystals of stable, 53(F)
surface finish, 89
tensile properties of carbon steel vs. for automotive
components, 229(T)
thermal processing, 161–164
water vapor, 81–82
weaknesses, 69
welding characteristics, 201–204
welding parameters, 207(T)
austenitizing, martensitic stainless steels, 131, 132(F),
166–167
automotive and transportation
alloy selection for exhaust systems, 226
alloys for major elements of automotive exhaust systems, 227(T)
automotive emission standards, 225
bus bodies, 231(F)
car manufacturers, 230
catalytic converter, 227(T), 228
center pipe, 227(T), 228
decorative to highly engineered applications in automobiles, 225
exhaust manifold and high-temperature, 227
exhaust systems, 225–228
ferritic stainless, 226
flexible pipe, 227(T), 228
front pipe, 227(T), 228
fuel tanks, 231
life-cycle cost calculation for stainless vs. carbon steel
for bus, 231(T)
microcar frame, 231, 232(F)
muffler, 227(T), 228
rail transport, 232
resilience and toughness of carbon and stainless steels
for automobiles, 229(T)
stress-strain curves for 301 variants vs. duplex steels and
transformation steel, 230(F)
structural components, 229–231
tailpipe, 227(T), 228
tensile properties of carbon and stainless steels for automobiles, 229(T)
trucks, 231
B
bacteria influencing corrosion, 55
bacterial retention, food contact materials, 236(F)
bases, corrosion in acids and, 31–36
basic oxygen furnace (BOF), 156
Bauschinger effect, 164
biocorrosion, 55–56
biological neutrality, food contact, 235
bleaching, pulp, 266
body-centered cubic (bcc) phase
carbon and alloy steels, 1
change to face-centered cubic (fcc), 127(F)
ferrite, 4, 110
ferritic material, 174
metals, 2
boron
additions to ferritic stainless steels, 121
ferrite, 4
brazing, 211
bright annealing, 198
brightening stainless steels, 196
buffing, 197
built-up edge (BUE)
austenitic stainless grades, 185
carbon and, 182
coolants minimizing, 191
copper and, 183
ferritic stainless steels, 185
grain sizes, 184
iron and tool, 182
nickel and, 183
nitrogen and, 184
precipitation hardening stainless steels, 185
bus
life-cycle cost calculation for stainless vs. carbon steel,
231(T)
stainless steel body, 230, 231(F)
Butler–Volmer equation, 20–21
C
calcium
effect on machinability of 303, 188–189, 189(F)
inclusions and pitting, 41
calcium-fluoride based slag
electroslag remelting (ESR), 158
“C” alloys
austenitic–ferritic alloys, 151–152
corrosion resisting, 147
duplex alloys, 151
mechanical properties of corrosion resisting cast, 150(T)
metallurgy of, 149, 151–152
precipitation hardening, 151
carbide, precipitation kinetics, 7(F)
288 / Index
carbides
duplex alloys, 94
flatware, 240
precipitation, 76–78
stainless steel, 9
tooling, 191
carbon
alloying element, 1
austenite, 5
ferrite, 3–4
influence of alloying elements on thermodynamic activity of, 157(T)
influence on thermodynamic activity of C, N, S and O,
157(T)
influence on uniform corrosion, 29
interstitial atoms of, in austenite, 6
machinability of stainless steels, 182–183
precipitation of carbides, 76–78
precipitation rates by, content, 76(F)
steel content, 156
variation of martensite hardness with, 128(F)
welding of austenitic stainless steels, 201
carbon diffusion, 115
carbon dioxide, wet, 248
carbon solubility, austenitic stainless, 76(F)
carbon steel
activities and activity coefficients of elements in, 40(T)
body-centered cubic (bcc) phase, 1
corrosion rates of stainless vs., 135(F)
metal migration into acetic solution from, 234(T)
resilience and toughness of, vs. stainless steel for automotive components, 229(T)
tensile properties of, vs. stainless steel for automotive
components, 229(T)
carburization, 82
carburizing, 199
casting
stainless steel processing, 158–159
Thermo-Calc software, 283
casting alloys
Alloy Casting Institute (ACI), 147
austenitic-ferritic alloys, 151–152
austenitic HE–HP alloys, 152–154
chromium alloys, 148(T)
chromium-nickel alloys, 148(T)
composition of cast heat-resistant stainless and nickel
base alloys, 149(T)
compositions of cast stainless corrosion resisting alloys,
148(T)
duplex alloys, 151
ferritic HA, HC, HD alloys, 152
foundry practice, 154
high-temperature mechanical properties of “H” alloys,
153(T)
influence of alloying elements, 147, 149
mechanical properties of heat-resistant stainless, at room
temperature, 152(T)
metallurgy of “C” alloys, 149, 151–152
metallurgy of “H” alloys, 152–154
molten metal transfer, 153
naming system, 147
precipitation hardening, 151
room temperature mechanical properties of corrosion
resisting stainless, 150(T)
welding, 154
catalytic converter, 227(T), 228
catastrophic oxidation, 65
cathode
effect of, polarization, 24(F)
electrochemical reactions, 12
mass transfer limitations, 24
polarization, 20–21, 23
caustic solutions, 50
center pipe of exhaust systems, 227(T), 228
cerium
inclusions and pitting, 41
oxidation resistance, 79–80
protective layer formation, 65
Charpy V toughness
high-temperature austenitic alloys, 83(F)
niobium-stabilized alloy, 119(F)
titanium-stabilized alloy, 119(F)
chemical agents, 255
chemical and process industry. See also corrosion types
corrosion table for fuming sulfuric acid, 261(T)
corrosion table for sulfuric acid, 259(T), 260(T), 261(T)
corrosion types, 258–262
erosion, 262
forms of corrosion, 258
intergranular corrosion, 261–262
isocorrosion chart for sulfuric acid, 258(F)
pitting and crevice corrosion, 258–259, 261
single- and dual-environment systems, 257–258
specific corrosives, 262
stress corrosion cracking (SCC), 257, 262
chemical neutrality, food contact, 233–235
chemistry, machinability of stainless steels, 182–184
chi
intermetallic phase, 9
precipitation kinetics, 7(F)
chip breaking, 187(F)
chloride concentration in rainwater, 217(F)
chloride-containing solutions, 50–51
chloride ion, aggressive against stainless steel, 85
chlorinated oils or waxes, 179(T)
chromium
alloying element, 1
austenitic alloys, 72
austenitic stainless steels, 85–86
chromium-oxygen system volatility vs. temperature and
oxygen pressure, 64(F)
corrosion resistance, 228
depletion from austenite near grain boundaries, 77(F)
ferrite, 3
ferritic stainless alloys with low, medium and high, 110
inclusions and pitting, 42, 43
influence on thermodynamic activity of C, N, S and O,
157(T)
influence on uniform corrosion, 29, 30(F)
ion release from stainless steel grades, 222(F)
Index / 289
machinability of stainless steels, 182
migration into acetic solution from stainless, aluminum
or carbon steel, 234(T)
oxidation resistance, 71, 79, 80(F), 228
paralinear oxidation from evaporation of chromium
superoxide, 64(F)
Pourbaix diagram, 17(F)
sulfide formation, 186
thermodynamics of oxidation, 57–59
volatile nature of Cr2O3, 63, 64(F)
chromium alloys, cast stainless, 148(T)
chromium-nickel alloys, 148(T)
Chrysler Building
architecture using stainless steel, 213
highly polished surface, 196, 197(F), 213(F)
cleaning
passivation, 195–196
recommended methods, 195(T)
stainless steel, 194–196
cleaning methods
stainless steels, 235–236
uncoated stainless steel, 220(T)
cleanliness
food contact materials, 235–236
machinability of stainless steels, 184
coastal climates, 237–238
coatings
cookware, 237
tooling, 191
cold heading, 179–180
cold work, 75
coloring of stainless steels, 196
commercial use
applications, 237–242
cookware, 237
flatware and cutlery, 240–241
food contact, 233–237
heating and water heating, 241–242
kitchen appliances, 237–240
laundry appliances, 241
stainless steel, 233
composition
Alloy Casting Institute (ACI) heat- and corrosionresisting casting alloys, 277(T), 278(T)
austenitic precipitation-hardenable (PH) stainless steel,
145(T)
austenitic stainless steels, 270(T), 271(T)
duplex alloys commercially available, 97 (T)
duplex stainless steels, 276(T)
ferrite stainless steels, 111(T), 112(T), 273(T), 274(T)
martensitic PH stainless steels, 140(T)
martensitic stainless steels, 124(T), 125(T), 275(T)
PH stainless steels, 276(T)
semiaustenitic PH stainless steel, 143(T)
tool and cutlery martensitic stainless steels, 134(T)
concrete reinforcing bar, 222
constitution diagram
Schaeffler–Delong stainless steels, 5(F), 70(F)
Schaeffler diagram, 202(F)
Welding Research Council’s 1992, 203(F)
construction. See architecture and construction
contamination, 235
continuous slab casting, 158
cookware, 237
coolants, 191
copper
acid resistance, 71
machinability of stainless steels, 183
copper sulfate. See also sulfuric acid plus copper sulfate
corrosion of stainless steel and titanium in, plus sulfuric
acid, 31(F), 32(F)
corrosion
combating, in alloys for petroleum industry, 248, 250
definition, 11
erosion, 262
intergranular, 261–262
isocorrosion chart for sulfuric acid, 258(F)
pitting and crevice, 258–259, 261
pitting resistance equivalent number (PREN), 258
single- and dual-environment systems, 257–258
stress corrosion cracking (SCC), 257, 262
table for sulfuric acid, 259(T), 260(T), 261(T)
tendency, 15–16
types, 258–262
corrosion cost, 252–253
corrosion kinetics
Butler–Volmer equation, 20–21
introduction, 19–20
mass transfer control, 21
migration and ionic diffusion, 21–22
mixed potential theory and polarization diagrams,
22–23
passivation, 23–25
Tafel regime: electrode-kinetics control, 21
corrosion rate
alloys in simulated evaporator liquid, 37(F)
influence of alloying element on, in contaminated
sulfuric acid, 33(F)
stainless oil country tubular goods, 135(F)
stainless vs. carbon steel, 135(F)
Tafel slope, 23(F)
vs. surface roughness, 215(F)
corrosion resistance
architecture, 213–214
austenitic precipitation-hardenable (PH) stainless steel,
145–146
balancing, processing and economy, 214–215
duplex alloys, 99–106, 204
ferritic stainless steels, 109, 110, 121–122
function of salinity and temperature, 244(F)
martensitic PH stainless steels, 141–142
material selection for desalination, 244–245
pulp-and-paper industry, 265–267
rail transport applications, 232
ratings of austenitic stainless steels, 87(T)
semiaustenitic PH stainless steel, 144
stainless steel for refinery equipment, 254–255
stainless steel in shipping, 245
sulfur hurting, 188
surface finish and, 215
290 / Index
corrosion resisting alloys
austenitic stainless steels, 84–89
“C” alloys, 149, 151–152
composition of, austenitic stainless steels, 86(T)
composition of Alloy Casting Institute (ACI), 277(T),
278(T)
compositions of cast stainless, 148(T)
duplex alloys, 91
mechanical properties of stainless, 150(T)
corrosion theory
corrosion tendency, 15–16
electrochemical reactions, 11–12
Faraday’s law, 12
galvanic vs. electrochemical cells, 14
Nernst equation, 12–14
Pourbaix diagrams, 16–17
standard half-cell reduction potentials vs. normal hydrogen electrode, 14(T)
corrosion types
atmospheric, 36–37
biocorrosion and microbiologically induced, 55–56
chromium influence, 29
corrosion fatigue, 55
corrosion in acids and bases, 31–36
corrosion with fatigue or fraction, 48–55
crack initiation, 48–49
crack propagation, 49, 52(F)
crevice, 38–39, 45–46
critical current density, 29, 30(F)
dissimilar metals and differential aeration cells, 38–39
environmental variables, 50–51
environmental variables influencing uniform corrosion,
28–29
grain boundary, 46–48
hydrochloric acid, 33–34, 35(F)
hydrogen embrittlement, 54–55
influence of alloying elements, 29, 30(F)
localized, 37–38
material variables, 29–31, 49–50
molybdenum role, 29–30
nickel, 30
nitric acid, 34, 35(F)
nitrogen, 30
organic acids, 35, 36(F)
phosphoric acid, 34–35, 36(F)
pitting, 39–45
pitting resistance, 43–45
preventing crevice, 45–46
SCC (stress corrosion cracking), 48–54
SCC mechanisms, 51–54
sensitization, 46–48
sodium chloride/carbon dioxide environment, 30
strong bases, 35
sulfuric acid, 31–33
sulfuric acid plus copper sulfate, 31(F), 32(F)
uniform, 27–37
corrosive environments
austenitic stainless steels, 88–89
platforms, 254
refinery equipment, 254–255
creep rupture strength, 83, 84(F)
creep strength, 83(F)
crevice corrosion
austenitic stainless steels, 85, 87
corrosion type, 214, 258–259, 261
critical, temperature with alloy content, 45(F)
critical crevice temperature (CCT) and critical pitting
temperature (CPT), 105(F)
dissimilar metals and differential aeration cells,
38–39
duplex alloys, 103–104
geometry, 45
preventing, 45–46
critical crevice temperature (CCT), 105(F)
critical current density, 29, 30(F)
critical pitting temperature (CPT)
austenitic steels, 43, 44(F)
critical crevice temperature (CCT) and CPT, 105(F)
duplex alloys, 103, 104(F), 105(F)
stainless steels for unwelded and welded material,
44(F)
vs. pitting resistance equivalent number (PREN), 85(F),
104(F)
cryogenic containers, 245
current density, 23(F)
cutlery
flatware and, 240–241
martensitic stainless steels, 133–134, 240
stainless steels commonly used for, 241(T)
cutting tools, 133–134
D
deep drawing
anisotropy, 174
forming stainless steel, 173, 174–179
geometry, 174–175
hydroforming, 177–178
intermediate annealing, 177
materials composition, 175(T)
schematic, 174(F)
strain rate, 177
texture, 174
tooling, 176, 178–179
defects, 160
delignification of pulp, 266
demand for steel, 247
desalination
materials selection for, 244–245
multi-stage flash (MSF), 243
reverse osmosis (RO), 243–244
technology, 243–244
design, balance, 215, 217
designers
car manufacturers, 230
pitting corrosion, 39
development
precipitation–hardening stainless steels,
137–138
welding, 211–212
Index / 291
differential aeration cells
active alloys, 38–39
microfouling, 56
schematic, 20(F)
differential aeration corrosion cell, 12
diffusion
atomic rearrangements, 2
ionic transport, 21–22
diffusion rates, austenite vs. ferrite, 6
digesters
first kraft from alloy 2205, 266(F)
pulp-and-paper industry, 265
disinfection, 236–237
dissimilar metals, 38–39
dissociated ammonia, 198–199
dissolution
equation, 37
term, 27
dryers, laundry appliances, 241
dry film, forming stainless, 179(T)
dual-environment system, 257
duplex alloys
acetic acid, 102
annealing, 170–171
composition of selected, stainless steels, 276(T)
compositions, 97(T)
concept, 91–92
corrosion resistance, 99–106
corrosion-resistant “C” alloys, 151
crevice corrosion, 103–104
deep drawing, 178
fastest-growing stainless steel family, 91–92
fatigue, 98
Fe-Cr-Ni phase diagrams, 92(F)
formation kinetics, 96(F)
formic acid, 103(F)
forming and machining, 99
heat exchangers, 246
hot forming, 180
hydrochloric acid, 100–101
impact strength, 97–98
impact strength variation with temperature, 75(F)
iron-nickel diagrams, 93(F)
machinability, 186
machining setup recommendations, 183(T)
mechanical properties, 94–98
nitric acid, 101
organic acids, 102, 103(F)
partitioning of elements, 93–94
petroleum industry applications, 248(T)
phosphoric acid, 101–102
photomicrographs, 95(F)
pitting corrosion, 102–103
pitting resistance equivalent number (PREN), 43
PREN influencing fatigue, 98, 99(F)
pulp-and-paper industry, 265, 266–267
recommended annealing and stress–relieving temperatures, 170(T)
SCC (stress corrosion cracking), 49, 104–106
sensitization, 47
soaking, 170
sodium hydroxide, 101
stainless steel for line pipe, 253
stainless steel in shipping, 245
strength, 96
stress corrosion cracking (SCC) performance, 262
stress-strain curves for 301 variants vs., 230(F)
structure and alloy design, 92–94
sulfuric acid, 100
thermal processing, 170–171
umbilical tubing and risers, 253–254
variations of ferrite, austenite, and duplex with temperature, 98(F)
welding characteristics, 204–205
welding parameters, 207(T)
wrought 2205 duplex microstructure, 91(F)
E
earing
deep drawing, 178
measuring tendency, 178
ecological considerations, 222
electrochemical cell
closed circuit, 11–12
potential, 38
electrochemical corrosion, 19
electrochemical reactions, 11–12
electrode-kinetics control Tafel regime, 21
electrodes, polarization, 20–21
electrolysis cell, 14(F)
electrolyte resistance, 22
electrolytic cells, galvanic vs., 14
electrolytic pickling, cold-rolled stainless, 194
electromotive force, 13
electropolishing, 196
electroslag remelting (ESR), 157–158
embrittlement. See also hydrogen embrittlement (HE)
alpha prime, 8
high-temperature, 114
σ phase at higher temperatures, 151
engineering alloys
minimum mechanical properties of stainless steel,
280(T)
physical properties of major stainless steel, 279(T)
environment, stainless steel selection expert system,
216(F)
environmental variables
stress corrosion cracking, 50–51
uniform corrosion, 28–29
epsilon martensite, 7–8, 73–74
equilibrium, argon oxygen decarburization (AOD),
155–156
equivalent weight (EW), 19
erosion, corrosion, 258, 262
expert system
recommended stainless steel grades, 217(F)
stainless steel selection, 216(F)
exposure to salt, stainless steel selection expert system,
216(F)
292 / Index
F
fabrication, 221
facades of appliances, 239–240
face-centered cubic (fcc) phase
aluminum alloys, 1
austenite, 5
austenitic materials, 174
change to body-centered cubic (bcc), 127(F)
metals, 2
Faraday’s law, 12, 21
fatigue
corrosion, 55
duplex alloys, 98, 99(F)
fatty oils and blends, suitability in forming stainless
steel, 179(T)
ferrite
carbon and nitrogen, 3–4
carbon diffusion rate, 115
chromium, 3
diffusion rates in austenite vs., 6
δ-ferrite in precipitation hardening alloys, 138–139
hydrogen and boron, 4
mechanical properties, 4
molybdenum, 4
phase diagram of iron chromium, 3(F)
phase in duplex alloy at room temperature, 91
Schaeffler–Delong constitution diagram, 5(F), 70(F)
Schaeffler diagram, 202(F)
stabilization with titanium, 4
thermal conductivity and thermal expansion, 4–5
ferritic “H” alloys, 152
ferritic stainless steels
alpha prime formation, 116, 117(F)
annealing, 165–166
automotive exhaust systems, 226
body-centered cubic (bcc), 174
carbon diffusion rate in, 115
composition, 111(T), 112(T), 273 (T), 274 (T)
corrosion and oxidation resistance, 109, 121–122
deep drawing, 178
embrittling phenomenon, 116
forming limit diagrams, 176 (F)
groups of low, medium and high chromium, 110, 113
heat exchangers, 245–246
high-temperature properties, 121
hot rolling, 159
impact strength variation with temperature, 75(F)
intermetallic phases, 116
iron-chromium phase diagrams, 113(F), 114(F)
kitchen appliances, 239
lowest cost and simplest stainless, 109–110
machinability, 185
machining setup recommendations, 183 (T)
mechanical behavior, 116–117
metallurgy, 113–116
metal migration into acetic solution from, 234(T)
petroleum industry applications, 247(T)
pitting resistance equivalent number (PREN), 43
recommended annealing temperatures, 165(T)
sensitization, 47
soaking, 165
stabilization, 109, 115, 118–120
stress corrosion cracking (SCC), 49
stress relieving, 166
superferritics, 113
texture and anisotropy, 120–121
time-temperature-transformation (TTT) curve for 430,
115(F)
titanium and niobium, 118
titanium for carbide and nitride formation, 115
toughness, 116(F), 117(F), 118–119
welding characteristics, 205–206
welding parameters, 207(T)
ferromagnetism, 5
fingerprints, cleaning methods for uncoated stainless,
220(T)
fire resistance, stainless steel, 221–222
flatness, surface aesthetic, 219
flatware, 240–241
flexible pipe, 227(T), 228
flow lines, 252–253
flux cored wire (FCW) welding, 210
food contact
bacterial retention by material and cleaning time, 236(F)
biological neutrality, 235
chemical neutrality, 233–235
cookware, 237
flatware and cutlery, 240–241
heating and water heating, 241–242
kitchen appliances, 237–240
material cleanliness, 235–236
metal migration into acetic solution, 234(T)
qualifications, 233–237
stainless steels commonly used for appliances, 239(T)
stainless steels commonly used for cutlery, 241(T)
surface disinfection, 236–237
formability, ferritic stainless steel, 120
formic acid
austenitic stainless steels, 89
corrosion in, 35
corrosion rates for various alloys of acetic plus formic
acid, 103(F)
duplex alloys, 102
isocorrosion curves in, 36(F)
forming limit diagram (FLD), 175–176
forming technology
deep drawing, 173, 174–179
deep drawing materials composition, 175(T)
deep drawing schematic, 174(F)
duplex alloys, 99
duplex stainless steel, 178
ferritics, 120, 178
flat, rolled stainless steel, 173–179
forces for hot working, 180(F)
forming limit diagram of carbon steel vs. austenitic
stainless steel, 176(F)
forming limit diagrams for stainless steel categories,
176(F)
hot, of stainless steel, 180
Index / 293
hydroforming, 177–178
hydrogen embrittlement, 177
limiting drawing ratio (LDR) vs. Lankford ratio, 175(F)
optimized 409 for forming vs. normal 409, 177(F)
orange peel, 178
stainless long products, 179–180
stainless steel, 173
stretch forming, 174
suitability of lubricants for use in, 179(T)
surface finish, 178
tooling, 176, 178–179
foundary practice, casting alloys, 154
free energy, phases, 2
friction stir welding, 212
front pipe, alloys in automotive exhaust systems,
227(T), 228
fuel tanks, 231
fuming sulfuric acid, 261(T)
furnace, stainless steels in, 241–242
G
Gallionella, 55
galvanic cell
electrochemical reaction, 38
schematic, 14(F)
vs. electrolytic cells, 14
gas metal arc welding (GMAW)
joint design, 209(F)
process, 210
gas tungsten arc welding (GTAW)
joint design, 209(F)
process, 208–210
geometry
crevice corrosion, 45
deep drawing, 174–175
pitting corrosion, 39
Gibbs free energy
electrochemical reactions, 12–13
oxidation, 57, 58(F)
grade selection, corrosion resistance, processing and
economy, 214–215
graffiti, cleaning methods for uncoated stainless,
220(T)
grain boundaries
austenite, 6
austenite, of martensite, 8
boron additions to ferritics, 121
carbide precipitation, 9, 76–77, 77(F)
corrosion, 46–48
defects in stainless steel, 160
depletion of chromium from austenite near, 77(F)
ferrite-austenite, 8
grain size
austenitic stainless steel annealing, 163
martensitic stainless steels and toughness, 131,
132(F)
material structure, 184
graphite, suitability in forming stainless steel, 179(T)
grinding, coarse polishing, 197
grit sizes, 197(T)
Guinier–Preston (GP) zones, 138
H
half-cell reactions
reduction potentials, 13–14
vs. normal hydrogen electrode, 14(T)
“H” alloys
austenitic HE–HP alloys, 152–154
corrosion resisting, 147
ferritic HA, HC, HD, 152
high-temperature mechanical properties of, 153(T)
mechanical properties of heat-resistant, 152(T)
metallurgy, 152–154
halogens, 82
hardening
austenitic stainless steels, 75
ferritic stainless steels, 116–117
heat-affected zone (HAZ)
austenitic stainless steel, 207
chromium carbide formation in, 201
duplex stainless steels, 204
ferritic stainless steels, 109, 205
laser welding, 210
martensitic stainless steels, 206
secondary austenite, 7
heat exchangers, 245–246
heat-resistant alloys
composition of Alloy Casting Institute (ACI), 278(T)
compositions, 149(T)
“H” alloys, 152–154
mechanical properties of cast stainless, 152(T)
heat tint, coloring stainless steels, 196
heat tinting, cleaning method, 220(T)
heat treatment, Thermo-Calc software, 283
heat treatment and conditioning, 168–170
heavy-duty emulsions, 179(T)
heavy metals, elimination, 157
high-frequency induction welding, 211
high-speed tool steels, 190–191
high-temperature alloys
austenitic stainless steel, 82(T)
intermetallic phases of austenitic stainless steel, 82
martensitic stainless steels, 133, 134(F)
mechanical properties of austenitic, 82–83, 84(F)
oxidation resistance of austenitic, 79–81
water vapor, 81–82
high-temperature embrittlement, 114
high-temperature properties, ferritic stainless steels, 121
hopper cars, 232
hot ductility defects, 160
hot forming, 180
hot mill defects, 160
hot rolling, 159–160
hot Steckel mills, 159
hot strip tandem mills, 159
hydrochloric acid
austenitic stainless steels, 88, 89(F)
corrosion in, 33–34
294 / Index
hydrochloric acid (continued)
duplex alloys, 100–101
influence of alloy content on corrosion rate in, 35(F)
isocorrosion curves for austenitic stainless steels in, 34(F)
isocorrosion curves for stainless steels in, 34(F)
isocorrosion performance of duplex, 101(F)
hydrofluoric acid, 193–194
hydroforming, 177–178
hydrogen/argon atmosphere, bright annealing, 198–199
hydrogen embrittlement (HE)
corrosion fatigue, 55
crack growth, 49
ferritic stainless steels, 121–122
mechanisms, 54–55
stress corrosion cracking, 51, 52–54
hydrogen ion reduction, 15(F), 28(F)
hypochlorite bleaches, 195
I
impact strength
duplex alloys, 97–98, 204
variation with temperature for stainless steels, 75(F)
inclusion-related defects, 160
inclusions
chip breaking at sulfides, 187(F)
lead, selenium, tellurium, 186
oxides, 188–190
pitting corrosion, 40–43
role in machining stainless steels, 186–190
stainless steel, 10
sulfur, 186–188
induction welding, high-frequency, 211
ingot method, 158, 159
initiation
pitting, 39–40, 43
stress corrosion cracking, 48–49
interatomic attraction, thermodynamics, 2
intermetallic phases
austenitic stainless steel, 82, 203
ferritic stainless steels, 116
stainless steel, 8–9
International Nickel Company (INCO) process, 196
interstitial elements, 6
ionic current, 11–12
ionic diffusion, 21–22
iron
body- and face-centered cubic transformations, 2
electrochemical corrosion, 19
ferrite, 3
ion release from stainless steel grades, 222(F)
machinability of stainless steels, 182
migration into acetic solution from stainless, aluminum
or carbon steel, 234(T)
penetration rates, 19(T)
Pourbaix diagram, 16(F)
pseudo-binary-phase diagram for, and sulfur, 41(F)
iron-chromium
phase diagram, 3(F)
phase diagram from Thermocalc, 113(F)
phase diagrams with varying carbon, 114(F),
130(F)
phase diagrams with varying chromium, 130(F)
iron dissolution, 20(F)
iron reduction, 15(F), 28(F)
J
JFI Steel, oil country tubular goods and line pipe
alloys, 135(T)
joining
stainless steel, 221
Thermo-Calc software, 283
joint design, 208, 209(F)
K
kinetics, alpha prime formation, 94, 96(F)
kitchen appliances
austenitic stainless steel, 239
coastal conditions, 238
exposure of stainless samples to North Carolina beach,
238(F)
facades, 239–240
ferritic stainless grades, 239
food contact, 237–240
interior or working parts, 239
knife-line attack, 48, 202
kraft process
paper-making, 265–267
pulp-and-paper industry, 36, 265
L
Lankford r, earing tendency, 178
Lankford ratio
anisotropy measure, 120, 175
limiting drawing ratio as function of, 175
lanthanum, protective layer, 65
laser welding, 210
lattice expansions, 6(F)
laundry appliances, 241
laves, precipitation kinetics, 7(F)
laves phase, 9
leaching, elements from stainless to foods, 234
lead, stainless steel machinability, 186
lean alloys
austenitic, 72–78
compositions of austenitic, 72(T)
martensite and austenite, 73–74
lime content, electroslag remelting, 158
limiting current
electrode reaction kinetics, 21
increasing mass transfer, 23, 24(F)
limiting drawing ratio, 175
line pipe
martensitic stainless, 134–135
stainless steel application, 252–253
liquefied natural gas (LNG) vessels, 254
localized corrosion, 37–38
Index / 295
lubricants
oxides, 189–190
suitability in forming stainless steel, 179(T)
lubrication, 197–198
M
machinability, material’s, 181(F)
machining stainless steels
austenitic, 185
carbides, 191
carbon, 182–183
chromium, 182
cleanliness, 184
coatings, 191
coolants, 191
copper, 183
cross-section size, 185
duplex, 186
duplex alloys, 99
ferritic, 185
high-speed tool steels, 190–191
introduction, 181–182
iron, 182
lead inclusions, 186
machinability of stainless steel families,
185–186
manganese, 183
martensitic, 185
material’s machinability, 181(F)
molybdenum, 183
nickel, 183
niobium, 184
nitrogen, 184
oxide inclusions, 188–190
physical and mechanical properties, 182–185
precipitation hardening, 185–186
process, 184
role of chemistry, 182–184
role of inclusions, 186–190
selenium inclusions, 186
setup recommendations for turning wrought stainless
steels, 183(T)
structure, 184
sulfur, 183
sulfur inclusions, 186–188
super stainless steels, 186
tellurium inclusions, 186
titanium, 184
tooling and coolants, 190–191
maintenance
balancing service, design and, 215, 217
stainless steel, 220–221
stainless steel selection expert system, 216(F)
manganese
alloying element, 1
austenite, 5
inclusions and pitting, 41
influence on thermodynamic activity of C, N, S and O,
157(T)
machinability of stainless steels, 183
sulfide formation, 186
manganese sulfides
inclusions, 41
stress risers, 188
x-ray examination, 189(F)
Marangoni effect, 207
marine systems
corrosion resistance vs. salinity and temperature, 244(F)
desalination, 243–245
heat exchangers, 245–246
materials for desalination, 244–245
shipping, 245
typical analyses and properties of marine alloys, 245(T)
martensite
carbon and nitrogen, 9
composition range, 7
formation, 126–127
forms, 7–8, 73–74
lattice expansions, 6(F)
lean alloy of, and austenite, 73–74
platelets from surface, 126(F)
reversion of, formed by cold work, 75(F)
Schaeffler–Delong constitution diagram, 5(F), 70(F)
Schaeffler diagram, 202(F)
tempering, 7
varying hardness with carbon content, 128(F)
martensitic alloys
“C” alloys, 149, 151
composition of, precipitation hardening (PH) alloys,
140(T)
corrosion resistance of, PH alloys, 141–142
mechanical properties of, PH alloys, 139(T)
microstructures of, PH alloys, 140(F)
precipitation hardening stainless steels, 139–142
martensitic stainless steels
annealing, 166
applications, 133–135
austenitizing, 166–167
composition, 124(T), 125(T)
composition of, 275(T)
compositions of tool and cutlery, 134(T)
corrosion rates of stainless oil country tubular goods
(OCTG) alloys, 135(F)
corrosion resistance, 123
distinction from other alloys, 123, 126
expanding austenite stability range with nickel, 131(F)
flatware and cutlery, 240–241
hardness variation with carbon content, 127, 128(F)
high-temperature use, 133, 134(F)
hot rolling, 159
influence of alloying elements, 130, 131(T)
iron-chromium phase diagrams, 130(F)
machinability, 185
machining setup recommendations, 183(T)
martensite formation, 126–127
OCTG and line pipe, 134–135
passivation, 195
petroleum industry applications, 247(T)
phase structure, 127–128, 130–131
296 / Index
martensitic stainless steels (continued)
photomicrographs, 129(F)
recommended annealing, austenitizing, and tempering
temperatures, 166(T)
sensitization, 47
smallest stainless steel category, 123
soaking, 166
strain energy, 126(F), 127(F)
stress relieving, 167–168
tempering, 167
tempering and toughness, 132, 133(F)
thermal processing, 131–133, 166–168
tool and cutlery alloys, 133–134
toughness by austenite grain size and phosphorus,
132(F)
welding characteristics, 206
welding parameters, 207(T)
mass transfer, 23, 24(F)
mass transfer control, 21
mass transport, 24–25
material selection, welding, 206–208
material structure, 184
material variables
stress corrosion cracking (SCC), 49–50
uniform corrosion, 29–31
mechanical behavior, ferritic stainless steels, 116–117
mechanical properties
austenite, 7
austenitic precipitation-hardenable (PH) stainless steel,
145
corrosion resisting cast stainless alloys, 150(T)
deep-drawing stainless steels, 175(T)
duplex alloys, 94–98
ferrite, 4
heat-resistant cast stainless alloys, 152(T)
high-temperature, of austenitic stainless steels, 82–83,
84(F)
high-temperature, of “H” alloys, 153(T)
lean austenitic alloys, 74–76
machinability of stainless steels, 182–185
marine alloys, 245(T)
martensite, 8
martensitic PH stainless steels, 139(T)
minimum, of stainless steel engineering alloys, 280(T)
semiaustenitic PH stainless steel, 144(T)
stainless steels, 10
mechanisms
pitting corrosion, 39–40
precipitation-hardening, 138
stress corrosion cracking (SCC), 51–54
melting production process, 155–157
metal dusting, oxidation, 67
metal flow directions, 208(F)
metal ions release, 222(F)
metallurgy
“C” alloys, 149, 151–152
ferrite stainless steels, 113–116
ferritic stainless, 226
“H” alloys, 152–154
introduction, 1–2
metal oxides
parabolic rate constants for growth, 59(F)
standard Gibbs free energy of, formation vs. temperature, 57, 58(F)
metals
oxidation, 57
with oxide scale, 61(F)
metastable condition, 2
metastable pitting, 40
microbiologically induced corrosion, 55–56
microcar frame, 231, 232(F)
microorganisms, food contact, 235
microstructures, martensitic precipitation-hardening
stainless steels, 140(F)
migration, ionic transport, 21–22
mineral resin, bacterial retention, 236(F)
mischmetal, 41
mixed potential theory, 22–23
mold powder, continuous casting, 158
molybdenum
alloying element, 1
austenitic alloys, 72
carbide precipitation, 9
corrosion resistance, 228
disulfide, 179(T)
ferrite, 4
influence on resistance to stress corrosion cracking,
50(F)
influence on thermodynamic activity of C, N, S and O,
157(T)
influence on uniform corrosion, 29–30
influencing critical pitting temperature in welded vs.
unwelded austenitic grade, 208(F)
machinability of stainless steels, 183
oxidation resistance, 226, 228
protective layer formation, 64–65
muffler, automotive exhaust systems, 227(T), 228
multistage flash (MSF), 243
N
naming system, Alloy Casting Institute (ACI), 147
National Association of Corrosion Engineers (NACE)
alloy listing, 247
diagram showing alloy suitability, 250, 251(F)
regulating high-strength alloys, 142
restrictions in use recommendations for stainless steels,
252(T)
natural gas, vessels for liquefied, 254
Nernst equation
open circuit potential, 20
thermodynamics of electrochemical reactions, 12–14
neutrality
biological, in food contact, 235
chemical, in food contact, 233–235
New York City’s Chrysler Building, 196, 197(F),
213(F)
nickel
alloying element, 1
austenite, 5
Index / 297
austenitic alloys, 72
carbide precipitation, 9
corrosion rates for stainless steels and, base alloys,
80(F)
influence on thermodynamic activity of C, N, S and O,
157(T)
influence on uniform corrosion, 30
influencing oxidation of iron-chromium alloys, 79,
80(F)
ion release from stainless steel grades, 222(F)
machinability of stainless steels, 183
migration into acetic solution from stainless, aluminum
or carbon steel, 234(T)
resistance to stress corrosion cracking, 50(F)
nickel base alloys, 149(T)
niobium
carbide former, 78
creep resistance, 71–72
high-temperature martensitic stainless, 133, 134(F)
machinability of stainless steels, 184
replacing titanium, 205
role in sensitization, 47
stabilization, 226
stabilization of ferritic stainless steels, 118–119
Nippon Steel, oil country tubular goods and line pipe
alloys, 135(T)
nitric acid
austenitic stainless steels, 88
corrosion behavior of high-silicon alloys in concentrated, 35(F)
corrosion in, 34
duplex alloys, 101
isocorrosion curve for, 35(F)
pickling oxide scale, 193–194
nitrides
duplex alloys, 94
stainless steel, 9–10
nitriding
austenitic stainless steel and, 82
surface alteration, 199
nitrogen
austenite, 5
austenitic alloys, 72
austenitic stainless steels, 71, 86
delay in carbide precipitation by, 78(F)
ferrite, 3–4
influence of alloying elements on thermodynamic activity of, 157(T)
influence on thermodynamic activity of C, N, S and O,
157(T)
influence on uniform corrosion, 29, 30
interstitial atoms of, in austenite, 6
machinability of stainless steels, 184
solubility in austenite, 77–78
stainless steel for line pipe, 253
nondestructive evaluation (NDE), 211
normal hydrogen electrode (NHE)
half-cell reduction potential vs., 14(T)
standard hydrogen electrode, 14
North Carolina, exposure of stainless steel, 238(F)
O
Occupational Safety and Health Administration
(OSHA), 211
Ohm’s law, 22
oil and grease marks, cleaning methods, 220(T)
“oil-canning,” 219
oil country tubular goods (OCTG)
influence of chromium on corrosion rate of steel,
249(F)
influence of copper and nickel on corrosion rate of
martensitic alloys, 249(F)
martensitic stainless, 134–135
stainless steels in petroleum industry, 250–252
oleum, 33
open circuit potential, 13
orange peel, 178
organic acids
austenitic stainless steels, 89
corrosion in, 35
duplex alloys, 102
isocorrosion curves in, 36(F)
overpotentials, 20
oxidation
effect of chromium, 57–59, 60(T)
effect of rare earth additions, 65
effect of silicon, aluminum, and molybdenum,
64–65
electrochemical nature of, 60–61
influence of nickel on, of iron-chromium alloys, 79,
80(F)
iron-chromium-oxygen phase diagram, 59(F)
kinetics and, rates, 61–63
metal dusting, 67
metal with oxide scale, 61(F)
oxidation-resisting grades of stainless steel, 60(T)
parabolic rate constants for growth of oxides, 59(F)
paralinear, from evaporation of chromium superoxide,
64(F)
quasi-steady-state approximation of moving boundary
problem of internal, 66(F)
reaction at anode, 12
schematic predicting thermal stresses, 65(F)
spalling and cracking of scale, 63–65
standard Gibbs free energy of formation of metal oxides
vs. temperature, 58(F)
temperature dependence of metal dusting of iron,
67(F)
thermodynamics of, 57–60
transient, 60
under less-oxidizing atmospheres, 66–67
volatile nature of Cr2O3, 63, 64(F)
Wagner’s theory, 61–63
oxidation resistance
austenitic stainless steels, 79–81
ferritic stainless steels, 109, 121–122
isooxidation curves, 80(F)
oxide film
coloring stainless steels, 196
parameters for, coloring stainless steel, 196(T)
298 / Index
oxides
effect calcium on machinability of 303, 188–189, 189(F)
elongated, 190(F)
inclusions, 10
inclusions and pitting, 41
machinability of stainless steel, 188–190
metal with, scale, 61(F)
pickling to remove, layer, 25, 193–194
removal of oxide scale, 193–194
stabilization of austenitic alloy, 78
Ugima, in 303 matrix, 189(F)
un-deformed, 190(F)
x-ray examination showing Ugima, and manganese sulfides, 189(F)
oxyfuel gas welding (OFW), 210
oxygen
austenite impurity, 6
ferrite impurity, 3
impurity, 1
inclusions and pitting, 40, 41–42
influence of alloying elements on thermodynamic activity of, 157(T)
influence on thermodynamic activity of C, N, S and O,
157(T)
influence on uniform corrosion, 29
steel content, 156
oxygen gas reduction, 15(F), 28(F)
oxygen pressure, 64(F)
P
paint, cleaning method for uncoated stainless, 220(T)
paper-making processes
bleaching pulp, 266
digestion, 265–266
kraft process, 265
process equipment, 266–267
washing and screening, 266
partitioning elements, 93–94
passenger trains, 232
passivation
effect on polarization diagrams, 23–25
removing surface contamination, 195–196
stainless steel, 25
theory, 23
transpassive regime, 24
passive behavior, 27
passivity, 27
penetration equation, 19
penetration rates, 19(T)
petroleum industry
alloy suitability vs. H2S and CO2 partial pressure, 251(F)
austenitic stainless steels for, 249(T)
chromium influence on corrosion rate in environments
by oil country tubular goods (OCTG), 249(F)
combating corrosion in applications, 248, 250
copper and nickel influencing corrosion rate of martensitic stainless alloys for OCTG, 249(F)
demand for steel, 247
duplex stainless steels for, 248(T)
ferritic stainless steels for, 247(T)
line pipe and flow lines, 252–253
liquefied natural gas (LNG) vessels, 254
martensitic stainless steels for, 247(T)
molybdenum influence on SCC susceptibility, 248,
250(F)
National Association of Corrosion Engineers (NACE),
247, 248, 250, 252(T)
OCTG, 250–252
platforms, 254
precipitation-hardening stainless steels for, 248(T)
presence of wet carbon dioxide, 248
refinery equipment, 254–255
restrictions in stainless steel use recommended by
NACE, 252(T)
stainless steels for refinery processes, 254(T)
stress corrosion cracking (SCC), 248
umbilical tubing and risers, 253–254
pH, corrosion tendency, 15–16
phase diagrams
computer models, 2
Fe-Cr-Ni, 92(F)
iron-chromium, 3(F), 73(F), 113(F), 114(F), 130(F)
iron-chromium-oxygen, 58, 59 (F)
iron-nickel, 93(F)
pseudo-binary-, for iron and sulfur, 41 (F)
phases
alloy systems, 1
ferrite, 3–5
free energy, 2
intermetallic, of stainless steel, 8–9
structure of martensitic stainless steels, 127–128,
130–131
phosphoric acid
austenitic stainless steels, 88
corrosion in, 34–35
duplex alloys, 101–102
electropolishing solution, 196
isocorrosion curves in, 36(F)
minimum temperatures for wet, with duplex alloys,
102(F)
phosphorus
ferrite impurity, 3
impurity, 1, 156–157
martensitic stainless steels and toughness, 131, 132(F)
photomicrographs
duplex alloys, 94, 95(F)
martensitic stainless steels, 127, 129(F)
physical properties
major stainless steel engineering alloys, 279(T)
stainless steels, 10
pickling
oxide layer removal, 25, 193–194
uniform corrosion, 28
pigmented pastes, 179(T)
Pilling–Bedworth ratio (PBR), 63
pitting
activities and activity coefficients in liquid steels, 40(T)
austenitic stainless steels, 85, 88
corrosion type, 39–40, 258–259, 261
Index / 299
critical, temperatures, 44(F)
geometry, 39
inclusions, 40–43
influence of sulfur level on, resistance, 42(F)
metastable, 40
mischmetal, 41
passive anode polarization curve, 40(F)
pit initiation, 39–40
pseudo-binary-phase diagram for iron and sulfur, 41(F)
resistance, 43–45
“weakest link” phenomenon, 214
pitting corrosion
CPT (critical pitting temperature) vs. NaCl concentration, 103, 105(F)
CPT vs. pH, 103, 105(F)
CPT vs. pitting resistance equivalent number (PREN),
103, 104(F)
duplex alloys, 102–103
varying pitting potential with temperature, 103, 104(F)
pitting resistance equivalent number (PREN)
austenitic alloys, 43, 78, 85
corrosion, 37, 258
critical pitting temperature vs., 85(F)
duplex alloys, 43, 92, 102
ferritic alloys, 43
influence on duplex alloy fatigue strength, 98, 99(F)
pitting corrosion, 43–45, 214
ranking stainless steels by PREN, 214(T)
umbilical tubing and risers, 253
Pittsburgh Convention Center, 219(F)
platforms, stainless steel, 254
polarization
anode, 23
cathode, 23
influence on uniform corrosion, 29, 30(F)
overpotentials, 20
passivating alloys, 39(F)
passive anode, curve, 40(F)
stainless steel in chloride-containing solution, 40
polarization diagrams
effect of cathode polarization, 24(F)
effect of mass transport, 24–25, 25(F)
mixed potential theory and, 22–23
passivation, 23–25
schematic, 22(F)
schematic of passive anode polarization curve, 23, 24(F)
polishing
grit sizes for target surface roughness, 197(T)
surface finishing, 197–198
polycarbonate, bacterial retention, 236(F)
polythionate, 50
polythionic acid, 255
porosity, 52
Porsche, auto components, 230
postweld stress relief, 205
potassium hydroxide, 88–89
Pourbaix diagrams
chromium, 17(F)
construction of, 16–17
iron, 16(F)
powder metallurgy, 159
precipitated phases
Guinier–Preston zones, 138
stainless steels, 8(T)
precipitation
carbides, 9
carbides and nitrides, 76–78
possible aluminum/titanium, 138(F)
precipitation-hardening stainless steels
advantage over martensitic, 137
annealed condition, 139
austenitic, 144–146
austenitic alloys, 170
cast PH alloys, 151
cold work influence on aging of A–286, 145, 146(F)
composition, 276(T)
composition of austenitic, 145(T)
composition of martensitic, 140(T)
compositions of semiaustenitic, 143(T)
corrosion resistance of martensitic, 141–142
corrosion resistance of semiaustenitic, 144
development, 137–138
influence of alloying elements, 141(T)
machinability, 185–186
machining setup recommendations, 183(T)
martensitic, 139–142
martensitic grades, 168–170
mechanical properties of martensitic PH alloys, 139(T)
mechanical properties of semiaustenitic, 144(T)
mechanism of PH, 138
microstructures, 140(F)
passivation, 195
petroleum industry applications, 248(T)
phases in stainless steel, 10
possible aluminum/titanium precipitates, 138(F)
presence of δ-ferrite and γ-austenite, 138–139
processing routes for S15700, 142, 143(F)
properties of A-286 vs. test temperature, 145(F)
recommended annealing and stress-relieving temperatures for martensitic grades of, 169(T)
semiaustenitic, 142–144
specialized family, 137
stress corrosion cracking (SCC), 141, 142
thermal processing, 168–170
welding characteristics, 206
welding parameters, 207(T)
precipitation kinetics, 7(F)
prevention, crevice corrosion, 45–46
production processes
basic oxygen furnace (BOF), 156
casting, 158–159
defects, 160
electroslag remelting (ESR), 157–158
hot rolling, 159–160
hot Steckel mills, 159
hot strip tandem mills, 159
impurities, 156–157
influence of alloying elements on thermodynamics,
157(T)
melting and refining, 155–157
300 / Index
production processes (continued)
remelting, 157–158
semiaustenitic precipitation-hardenable stainless steel,
142, 143(F)
stainless steel, 155
thermodynamics, 156
vacuum arc remelting (VAR), 157–158
vacuum induction melting (VIM), 157
vacuum oxygen decarburization (VOC), 156
propagation
crack, rates of metals vs. current density, 52(F)
stress corrosion cracking (SCC), 49
pulp-and-paper industry
duplex stainless steels, 265, 266–267
kraft process, 36, 265
paper-making processes, 265–267
pulsed arc transfer, 210
Q
quality control, Thermo-Calc software, 283
R
rain, cleansing action of, 220
rainwater, average chloride concentration, 217(F)
rare earth metals
inclusions and pitting, 41
protective layer formation, 65
reduction, 12
reduction potential
iron and hydrogen ion reductions vs. pH, 15(F),
28(F)
iron and oxygen gas reductions vs. pH, 15(F),
28(F)
reference electrode, 13–14
refinery equipment, 254–255
refining production process, 155–157
repair, 221
residential applications
cookware, 237
domestic goods, 233
flatware and cutlery, 240–241
heating, 241–242
kitchen appliances, 237–240
laundry appliances, 241
water heaters, 241, 242
resistance
Ohm’s law, 22
pitting, 43–45
resistance welding, 210–211
resistivity, 22(T)
reverse osmosis, 243–244
ridging, ferritics, 178
rolled finishes
applications, 198–199
benefits, 198
roping, ferritics, 178
rust staining, cleaning method for uncoated stainless,
220(T)
S
Saab, auto components, 230
safety, welding, 211
salinity, corrosion resistance vs., 244(F)
salt exposure, stainless steel selection expert system,
216(F)
sanitation, cleaning stainless steels, 195
scale
metal with oxide, 61(F)
spalling and cracking of, 63–65
Schaeffler–Delong constitution diagram, 5(F), 70(F)
Schaeffler diagram, 202–203
seawater. See also marine systems
desalination, 243–245
secondary austenite, 7
secondary phases, 82(T)
selenium, 186
semiaustenitic precipitation-hardenable stainless steels
austenite-stabilizing elements, 142
compositions, 143(T)
corrosion resistance, 144
mechanical properties, 144(T)
processing by T route, 142, 143(F)
sensitization
austenitic, 46–47
duplex steels, 47
effect of alloying, 47–48
ferritic, 47
ferritic stainless steels, 115
heat treatment vs. time, 46(F)
intergranular corrosion, 9
knife-line attack, 48, 202
martensitic steels, 47
schematic of, due to chromium-rich precipitates, 46(F)
welding, 48
service, design, and maintenance, 215, 217
shielded metal arc welding (SMAW)
joint design, 209(F)
process, 210
shielding gas
welding austenitic stainless steel, 203–204
welding parameters for various stainless steels, 207(T)
shipping, 245
short-circuiting transfer, 210
sigma
intermetallic phase, 8–9
precipitation kinetics, 7(F)
silicon
alloying element, 1
content in cast alloys, 147
corrosion of high- austenitic steels in nitric acid, 34, 35(F)
inclusions and pitting, 41
influence on thermodynamic activity of C, N, S and O,
157(T)
oxidation resistance, 71, 79, 226, 228
protective layer formation, 64–65
single-environment system. See also chemical and
process industry
aggressive chemical species, 258
Index / 301
slabs, 158
slip dissolution, 51
soaking
austenitic stainless steels, 161–162
duplex stainless steels, 170
ferritic stainless steels, 165
martensitic stainless steels, 166
soap-fat pastes, 179(T)
sodium chloride/carbon dioxide environment, 30
sodium hydroxide
austenitic stainless steels, 88–89
corrosion in, 35–36
corrosion rates of duplex alloys, 101(F)
corrosion rates of duplex alloys with contaminated environment, 101(F)
duplex alloys, 101
sodium hypochlorite
cleaning stainless steels, 195
disinfecting stainless, 236–237
software package, Thermo-Calc, 2, 281–283
soldering, 211
solution treatment, precipitation-hardening stainless
steels, 168–170
specialization, precipitation-hardening stainless steels,
137
Sphaerotilus, 55
spots, cleaning method for uncoated stainless, 220(T)
spray transfer, 210
stability
expanding austenite, with nickel, 131(F)
lean alloy of martensite and austenite, 73–74
stabilization
ferrite, 4
ferritic stainless steel, 109, 115, 118–120
ferritic steels for exhaust systems, 226
lean austenitic alloys, 78
stacking fault, 73–74
stainless long products
cold heading, 179–180
hot forming, 180
stainless steel alloys
austenite, 5–7
ferrite, 3–5
stainless steels. See also casting alloys
activities and activity coefficients of elements in, 40(T)
bacterial retention by material and cleaning time,
236(F)
carbides, 9
casting, 158–159
casting alloys, 147, 149
categories for oxidation resistance, 59–60
classifications by sulfur content, 187–188
cleaning methods for uncoated, 220(T)
composition of austenitic, 270 (T), 271 (T)
composition of duplex, 276(T)
composition of ferrite, 273(T), 274(T)
composition of martensitic, 275(T)
composition of precipitation-hardenable (PH), 276(T)
concrete reinforcing bar, 222
corrosion rates of, vs. carbon steel, 135(F)
corrosion table for, in sulfuric acid plus copper sulfate,
31(F)
deep drawing, 173, 174–179
defects in, hot–rolled bands, 160
flat, rolled, 173–179
hot rolling, 159–160
inclusions, 10, 186–190
isocorrosion curves for, in sulfuric acid plus copper sulfate, 32(F)
lubricants for forming, 179 (T)
machinability, 185–186
machining setup recommendations, 183 (T)
melting and refining, 155–157
minimum mechanical properties of, engineering alloys,
280 (T)
nitrides, 9–10
oxidation-resisting grades, 60(T)
passivation, 25
penetration rates, 19(T)
physical properties of major, engineering alloys, 279 (T)
precipitated phases, 8 (T)
precipitation-hardening process, 10
precipitation kinetics in 316, 7 (F)
properties, 10
ranking by pitting resistance equivalent number (PREN),
214(T)
ranking common, by PREN, 214(T)
refinery processes, 254(T)
remelting, 157–158
resilience and toughness of carbon steel vs. for
automotive components, 229(T)
Schaeffler–Delong constitution diagram, 5 (F)
selection expert system, 216(F)
tensile properties of carbon steel vs. for automotive
components, 229(T)
thermodynamics, 2
welding parameters, 207(T)
stains, cleaning method for uncoated stainless, 220(T)
standard Gibbs free energy, 57, 58(F)
standard hydrogen electrode, 14
Steckel mills, hot, 159
Steel Founder’s Society of America, 147
strain energy, martensitic stainless steels, 126(F), 127(F)
strain rate, deep drawing, 177
stress corrosion cracking (SCC)
advantages of duplex alloys, 105–106
austenitic stainless steels, 87–88
corrosion form, 258
crack initiation, 48–49
crack propagation, 49
crack propagation rates of metals vs. current density,
52(F)
debating mechanisms, 105
dilation of austenite due to hydrogen in solution, 53, 54(F)
duplex alloys, 91, 104–106
environmental variables, 50–51
ferritic stainless steels, 121
influence of molybdenum on resistance, 50(F)
martensitic precipitation-hardening (PH) stainless steels,
141, 142
302 / Index
stress corrosion cracking (continued)
material variables, 49–50
mechanisms, 51–54
petroleum industry, 248
resistance, 257, 262(F)
stress-strain curve for single crystals of austenitic steel
with and without hydrogen, 53(F)
susceptibility of martensitic stainless steels, 123
susceptibility to, with oxygen and chloride content for
304 stainless, 51(F)
theory, 262
threshold stress for, for various alloys, 88(F)
varying resistance to, with nickel content, 50(F)
water heaters, 242
zones of susceptibility, 48(F)
stress relief annealing (SRA), 48
stress relieving
austenitic stainless steels, 164
ferritic stainless steel, 166
martensitic stainless steels, 167–168
stress risers, 188
stress sorption, stress corrosion cracking, 51, 53
stretch forming, 174
strip casters, 158–159
strip tandem mills, 159
strong bases
austenitic stainless steels, 88–89
corrosion in, 35–36
structure
duplex alloys, 92–94
machinability of stainless steels, 184
phase, of martensitic stainless steels, 127–128, 130–131
submerged arc welding (SAW)
joint design, 209(F)
process, 210
submerged entry nozzle, 158
sulfate process, pulp-and-paper, 36
sulfides
inclusions, 10
inclusions in stainless steel, 186–187
size and shape and machinability, 187
stabilization of austenitic alloy, 78
sulfite process, 265
sulfur
austenite impurity, 6
comparing machinability, 190(F)
effect on stainless machinability, 187, 188(F)
ferrite impurity, 3
impurity, 1, 156
inclusions and pitting, 40, 41
influence of alloying elements on thermodynamic activity of, 157(T)
influence on thermodynamic activity of C, N, S and O,
157(T)
machinability of stainless steels, 183
metal flow directions in weld pool with and without,
208(F)
pitting resistance of unannealed welds, 42(F)
pseudo-binary-phase diagram for iron and, 41(F)
stainless steel machinability, 186–188
sulfuric acid
austenitic stainless steels, 88
corrosion in, 31–33
corrosion table, 259(T), 260(T), 261(T)
corrosion table for fuming, 261(T)
duplex alloys, 100
electropolishing solution, 196
influence of alloying element on corrosion rate in contaminated, 33(F)
isocorrosion, 88(F)
isocorrosion chart for, 258(F)
isocorrosion curves for various alloys in, 33(F)
isocorrosion curves for various alloys in, with chlorides,
34(F)
isocorrosion curves of duplex grades, 100(F)
isocorrosion rates for various stainless steels, 32(F)
oleum, 33
pickling oxide scale, 193–194
sulfuric acid plus copper sulfate
corrosion table for stainless steels and titanium in, 31(F)
isocorrosion curves for stainless steel and titanium in,
32(F)
sulfurized or sulfochlorinated oils, 179(T)
superaustenitic stainless steels, 164
superferritics, 113
super stainless steels, 186
surface finishing
aesthetics, 217–219
aesthetic surface finishes, 196–199
austenitic stainless steels, 89
bright annealing, 198
brightening, 196
cleaning, 194–196
coloring, 196
and corrosion resistance, 215
deep drawing, 178–179
function of surface treatments, 193–196
introduction, 193
parameters for oxide film coloring of stainless, 196(T)
passivation, 195–196
pickling, 193–194
polished finishes, 197–198
recommended cleaning methods, 195(T)
removal of oxide scale, 193–194
rolled finishes, 198–199
rolled-on finishes, 218 (F)
surface alteration, 199
surface roughness, 215(F)
surface treatments
brightening, 196
cleaning, 194–196
coloring, 196
passivation, 195–196
removal of oxide scale, 193–194
T
Tafel slope, 23(F)
tailpipe, automotive exhaust systems, 227(T), 228
tellurium, 186
Index / 303
temperature
chromium-oxygen system volatility, 64(F)
corrosion resistance vs., 244(F)
critical crevice corrosion, with alloy content, 45(F)
critical pitting, (CPT), 43, 44(F)
impact strength variation with, for stainless steels,
75(F)
partitioning ratio varying with, 93(F)
standard Gibbs free energy of metal oxide formation vs.,
58(F)
variation of pitting potential with, for duplex alloys,
104(F)
tempering
influencing martensitic stainless hardness, 132,
133(F)
martensite, 7
martensitic stainless steels, 167
tensile properties
austenitic precipitation-hardenable stainless steel,
145(F)
austenitic stainless steels, 75
tensile strength equation, 74
texture
deep drawing, 174
ferritic stainless steels, 120–121
thermal conductivity
duplex alloys, 205
ferrite, 4–5
thermal cutting, 211
thermal expansion
austenitic stainless steels, 202
duplex alloys, 205
ferrite, 4–5
thermal processing
annealing, 162–164, 165–166, 170–171
austenitic stainless steels, 161–164
austenitizing, 166–167
duplex stainless steels, 170–171
ferritic stainless steels, 165–166
martensitic stainless steels, 131–133, 166–168
precipitation-hardening stainless steels, 168–170
soaking, 161–162, 165, 166, 170
stress relieving, 164, 166, 167–168
tempering, 167
thermal stresses, predicting, 64, 65(F)
Thermo-Calc
alloy design, 93, 282
applications, 282–283
casting, 283
free demonstration version, 283
heat treatment, 283
iron-chromium phase diagram, 113(F)
map, 282
phase determination program, 2
quality control, 283
Scheil–Gulliver model, 282
single-point equilibria, 282
software package, 281–282
step, 282
welding and joining, 283
thermodynamics
argon oxygen decarburization (AOD), 156
influence of alloying elements on, activity of C, N, S,
and O, 157(T)
oxidation, 57–60
stainless steel, 2
thiosulfate, 50
time-temperature-transformation (TTT) diagram
high-alloy stainless steel, 94, 96(F)
unstabilized 430-type alloy, 115(F)
titanium
carbide and nitride formation, 115
carbide former, 78
corrosion table for, in sulfuric acid plus copper sulfate,
31(F)
deoxidizer in chromium–iron alloys, 156
ferritic alloy stabilization, 4, 205
inclusions and pitting, 41
influence on thermodynamic activity of C, N, S and O,
157(T)
isocorrosion curves for, in sulfuric acid plus copper sulfate, 32(F)
isocorrosion curves in phosphoric acid, 36(F)
machinability of stainless steels, 184
possible aluminum/titanium precipitates, 138(F)
role in sensitization, 47
stabilization of ferritic stainless steels,
118–119
stabilization of ferritic steels, 226
titanium-aluminum-nitride (TiAlN), 191
titanium carbonitride (TiCN), 191
titanium nitride (TiN), 191
tooling
carbides, 191
coatings, 191
coolants, 191
costs in deep drawing, 176
high-speed tool steels, 190–191
lubricants, 189, 190
materials in deep drawing, 178–179
tools, martensitic stainless steels, 133–134
toughness
austenitic stainless steels, 75–76
duplex alloys, 97–98, 204
ferritic stainless steels, 117, 118(F), 118–119
high-temperature austenitic alloys, 83(F)
martensitic stainless steels, 131, 132(F)
trains, 232
transient oxidation, 60
transpassive dissolution, 27
transpassive regime, 24
transportation. See automotive and transportation
trucks, 231
tubular goods, 134–135
tungsten
carbides for flatware, 240
influence on thermodynamic activity of C, N, S and O,
157(T)
tungsten inert gas (TIG), 208–210
tuyeres, oxygen injection, 156
304 / Index
U
Ugima oxide
coating cutting tool and lubricant, 189–190
comparing 304L chips with and without, 190(F)
machinability by sulfur levels with and without, 190(F)
x-ray showing, 189(F)
umbilical tubing and risers, 253–254
uniform corrosion. See also corrosion types
environmental variables influencing, 28–29
material variables, 29–31
stainless steel, 27–28
United States, chloride concentration in rainwater,
217(F)
unmixed zone, 208
V
vacuum arc remelting (VAR), 157–158
vacuum induction melting (VIM), 157
vacuum oxygen decarburization (VOD)
cleanliness, 184
refining process, 156
vanadium
carbides for flatware, 240
high-temperature martensitic stainless, 133, 134(F)
Volvo, auto components, 230
W
Wagner’s theory, 61–63
washer tubs and drums, 241
water heaters, 241, 242
waterline corrosion, 38
water marking, cleaning method for uncoated stainless,
220(T)
water vapor, 81–82
wax-base pastes, 179(T)
wax or soap plus borax, 179(T)
weather pattern, stainless steel selection expert system,
216(F)
weldability, 253
welding
austenitic stainless steels, 201–204
cast stainless alloys, 154
characteristics of stainless steels, 201–206
duplex stainless steels, 204–205
ferritic stainless steels, 205–206
flux cored wire (FCW), 210
gas metal arc welding (GMAW), 210
gas tungsten arc welding (GTAW), 208–210
high-frequency induction, 211
joint design, 208, 209(F)
laser, 210
martenistic stainless steels, 206
material selection and performance, 206–208
metal flow directions in weld pool, 208(F)
new developments, 212
nondestructive evaluation (NDE), 211
oxyfuel gas welding (OFW), 210
parameters for various stainless steels, 207(T)
practices, 211–212
precipitation-hardening (PH) stainless steels, 206
processes, 208–211
recent developments, 211–212
resistance, 210–211
safety, 211
Schaeffler diagram, 202(F)
sensitization, 48
shielded metal arc welding (SMAW), 210
soldering and brazing, 211
submerged arc welding (SAW), 210
thermal cutting, 211
Thermo-Calc software, 283
tungsten inert gas (TIG), 208–210
Welding Research Council’s 1992 constitution diagram,
203(F)
weld shielding gas composition and crevice corrosion
resistance, 204(F)
Welding Research Council, constitution diagram,
203(F)
welds
influence of sulfur on pitting resistance of unannealed,
42(F)
long-term annealing, 43
wet carbon dioxide, 248
Y
yield strength
austenitic precipitation-hardenable stainless steel, 145(F)
equation, 74
high-temperature austenitic alloys, 83, 84(F)
yttrium, 65
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