Corrosion resistant alumina-­‐forming alloys for lead-­‐cooled fast reactors Jesper Ejenstam Doctoral Thesis in Chemistry
Royal Institute of Technology
Stockholm, Sweden 2015
Akademisk avhandling som med tillstånd av Kungliga Tekniska
Högskolan framläggs till offentlig granskning för avläggande av
teknologie doktorsexamen den 12 januari 2015 kl 10.00 i hörsal F3,
KTH, Lindstedtsvägen 26, 114 28 Stockholm.
Avhandlingen presenteras på engelska.
Corrosion resistant alumina-forming alloys for lead-cooled
fast reactors
Doctoral Thesis in Chemistry
KTH Royal Institute of Technology
School of Chemical Science and Engineering
Department of Surface and Corrosion Science
Drottning Kristinas väg 51
SE-100 44 Stockholm
Sweden
Denna avhandling är skyddad enligt upphovsrättslagen. Alla
rättigheter förbehålls.
Copyright © Jesper Ejenstam. All rights reserved. No parts of this
thesis may be reproduced without permission from the author.
TRITA-CHE REPORT 2015:1
ISSN 1654-1081
ISBN 978-91-7595-345-8
Printed at US-AB, Stockholm, Sweden 2014
The following parts are printed with permission:
PAPER I: © 2013 Elsevier B.V
PAPER II: © 2014 Elsevier B.V
Forever trust in who we are Nothing else matters Abstract Generation IV nuclear power technologies provide attractive
solutions to the common issues related to conventional nuclear
power plants currently in operation worldwide. Through a
significant reduction of the long-term radiotoxicity of nuclear
waste, a more efficient use of nuclear fuel resources, and
implementation of inherent safety features, Generation IV will
make nuclear power sustainable and thus increase the public
acceptance of nuclear power. Due to its attractive safety features,
the lead-cooled fast reactor (LFR) is one of the most studied
Generation IV reactor concepts currently. It is well known that
liquid lead is corrosive to steels at elevated temperatures, thus
limiting the operation temperature of the LFR. The use of aluminaforming FeCrAl alloys has been proposed to mitigate oxidation and
corrosion issues. Commercial FeCrAl alloys have Crconcentrations typically about 20 wt. % and are thus prone to α-α’
phase separation and embrittlement at temperatures up to about 500
°C. Reducing the Cr-concentration to levels around 10 wt. % would
theoretically resolve the said issue. However, the oxidation and
corrosion resistance may be impaired. In the scientific literature,
compositional limits indicating the formation of protective alumina
layers at various temperatures have been presented. Long-term
corrosion studies are however scarce. Moreover, in-depth studies
on the compositional limits regarding α-α’ phase separation are
lacking. In this thesis, the long-term (up to 10,000 h) corrosion
resistance and phase stability of alumina-forming alloys are studied
at temperatures up to 550 °C. In addition, the influence of reactive
elements (RE), e.g. Ti, Zr, and Y, on the liquid lead corrosion
resistance of Fe10CrAl alloys is evaluated. By balancing the
reactive element and the carbon content, with respect to carbide
v
formation, it is demonstrated in this thesis that it is possible to form
protective alumina layers on Fe10Cr4Al alloys from 450 °C,
despite the low Al and Cr concentrations. It was found that the
RE/carbon ratio needed to form protective alumina layers on
Fe10Cr4Al alloys must be larger than unity to mitigate the
detrimental effect of Cr-carbide formation. The underlying
phenomena are discussed, and a mechanism is suggested based on
the outcome of the long-term oxidation studies. The phase stability
of Fe10CrAl alloys was studied through thermal aging experiments
in the temperature interval of 450 to 550 °C. In addition, the results
were well reproducible using a developed Kinetic Monte Carlo
(KMC) model of the FeCrAl system. Furthermore, the model
indicated that the Cr-concentration should be limited to about 11
wt. % in a FeCr4Al alloy to mitigate α-α’ phase separation at all
temperatures of interest for an LFR. The liquid lead corrosion
resistance of alumina-forming austenitic stainless steels was shown
to be superior compared to regular stainless steels, albeit the effect
of ferrite stabilizing elements needs to be further addressed. The
results included in this thesis provide a valuable input on the key
issues related to the development of corrosion resistant aluminaforming alloys of interest for liquid lead applications. Moreover,
the superior oxidation properties of the studied alumina-forming
alloys make them of interest for use in other energy applications,
where corrosion issues limits the operation temperature and thus
the efficiency.
vi
Sammanfattning Generation IV-kärnkraftverk skulle kunna lösa många av de
problem som idag är relaterade till konventionella kärnkraftverk
världen över. Genom att kraftigt minska förvaringstiden för det
radiotoxiska kärnavfallet, mer effektivt utnyttja bränsleresurser,
och erbjuda systeminneboende säkerhetsmekanismer kommer
Generation IV-kärnkraftverk göra kärnkraften hållbar samt ge ökat
förtroende från allmänheten. Blykylda kärnkraftverk är särskilt
attraktiva just på grund av de möjliga säkerhetslösningar tekniken
erbjuder. Smält bly är dock ett korrosivt medium, och detta faktum
begränsar
maxtemperaturen
i
blykylda
kärnkraftverk.
Aluminiumoxid-bildande
FeCrAl-legeringar,
med
välkänt
fördelaktiga oxidations- och korrosionsmotstånd, har föreslagits
kunna lösa korrosionsproblemen. Kommersiella FeCrAl-legeringar
innehåller normalt runt 20 vikts-% Cr och drabbas därför av α-α’
fasseparation, vilket är kopplat till försprödning, i temperaturer upp
till ca 500 °C. En sänkning av Cr-halten till runt 10 vikts-% kan
teoretiskt
lösa
problemet,
men
oxidationsoch
korrosionsmotståndet
kan
dock
påverkas
negativt.
Sammansättningsgränser för bildning av skyddande aluminiumoxid
vid olika temperaturer har presenterats i vetenskapliga tidskrifter,
men få långtidsstudier har genomförts. Dessutom saknas ingående
undersökningar som indikerar sammansättningsgränsen för
fasseparation. I den här avhandlingen har långtidsstudier, med
avseende på korrosionsmotstånd och fasstabilitet, genomförts vid
temperaturer upp till 550 °C. Därutöver har effekten av reaktiva
element (RE), såsom Ti, Zr och Y, på legeringarnas
korrosionsmotstånd i smält bly undersökts. Studierna har visat att
det är möjligt för en Fe10Cr4Al-legering, trots lågt Al- och Crinnehåll, att bilda skyddande aluminiumoxid i temperaturer ner till
vii
450 °C. Detta genom att balansera RE- och kol-innehållet med
avseende på karbidbildning. Kromkarbider visade sig påverka
oxidbildningen negativt, och genom att välja en RE/kol-kvot större
än ett kan detta undvikas. I studierna diskuteras de bakomliggande
fenomenen kopplade till denna effekt, och en mekanism har
föreslagits. Fe10CrAl-legeringars fasstabilitet undersöktes genom
termisk åldring i temperaturintervallet 450-550 °C. Parallellt med
dessa undersökningar utvecklades en Kinetic Monte Carlo-kod
baserad på FeCrAl-systemet. Med denna kod erhölls
simuleringsresultat i linje med experimenten. Simuleringarna
indikerade även att en Fe11Cr4Al-legering skulle vara immun mot
fasseparation i hela det temperaturintervall som är av intresse för en
blykyld kärnreaktor. Aluminiumoxid-bildande austenitiska stål
undersöktes också i smält bly och visade sig vara överlägsna
konventionella rostfria stål, dock måste fasstabiliteten och de
mekaniska egenskaperna undersökas vidare. Studierna i den här
avhandlingen bidrar med viktig data för att lösa nyckelproblem
kopplade till aluminiumoxid-bildande legeringar för applikationer
som inkluderar smält bly. De överlägsna oxidationsegenskaperna
hos de aluminiumoxid-bildande legeringarna gör att de även har
stor potential att lösa korrosionsproblem, och därmed öka
prestandan, även inom befintlig energiproduktion där höga
temperaturer krävs.
viii
Acknowledgements This thesis presents the work I have carried out over the last couple
of years. However, it would not have been possible without the
unlimited support and guidance from a large number of people in
my vicinity.
First and foremost, I would like to express my sincere gratitude to
my supervisors Dr. Peter Szakalos (KTH), Adj. Prof. Bo Jönsson
(Sandvik Heating Technology AB) and Prof. Inger Odnevall
Wallinder (KTH) for guiding me through my Ph.D. studies.
Especially to Dr. Peter Szakalos for your endless commitment to
my project. During these years, we have proven to be a small yet
effective team, and I am looking forward to our future endeavors.
I would also like to thank all my co-authors, Dr. Mattias Thuvander
(Chalmers University of Technology), Assoc. Prof. Pär Olsson
(KTH), Prof. Mats Halvarsson (Chalmers University of
Technology), Dr. Jonathan Weidow (Chalmers University of
Technology), and Fernando Rave (Sandvik Heating Technology
AB). I have really enjoyed working with all of you, and I am proud
of the papers we have compiled together.
During this journey, I have received an incredible amount support
and enthusiasm from all my colleagues, too many to mention, at the
division of Surface and Corrosion Science (KTH), division of
Reactor Physics (KTH) and Sandvik Heating Technology AB. I
would like to thank you all for contributing to such a warm and
motivating working environment
ix
I have also received invaluable help and guidance from many
others. However, I would especially like to acknowledge Prof.
Janne Wallenius (KTH), Fernando Rave (Sandvik Heating
Technology AB), Dr. Alfons Weisenburger (Karlsruhe Institute of
Technology), Dr. Gunilla Herting (KTH), Dr. Merja Pukari
(KTH/Studsvik AB), Oskar Karlsson (Swerea KIMAB), and Peter
Byhlin (Sandvik Heating Technology AB). Without your help, I
would not have reached this far in my scientific career, and for that
I am truly grateful.
No research can be conducted without funding. Therefore, I
humbly thank the VR (Swedish Research Council) funded
GENIUS project, and the Elforsk funded KME 508 project and the
Euratom FP7 funded MatISSE project.
To friends and family near and far, especially to my wife Lina and
my son Henry, thank you for your unconditional love and support.
x
List of Publications Included publications The doctoral thesis is based on the following four publications:
Paper I:
Oxidation studies of Fe10CrAl–RE alloys exposed to
Pb at 550 °C for 10,000 h
Jesper Ejenstam, Mats Halvarsson, Jonathan Weidow,
Bo Jönsson and Peter Szakalos
Journal of Nuclear Materials, 443 (2013), 161-170
© 2013 Elsevier B.V
Paper II:
Microstructural stability of Fe-Cr-Al alloys at 450550 °C
Jesper Ejenstam, Mattias Thuvander, Pär Olsson,
Fernando Rave and Peter Szakalos
Journal of Nuclear Materials (in press)
© 2014 Elsevier B.V
Paper III: Long term corrosion resistance of alumina forming
austenitic stainless steels in liquid lead
Jesper Ejenstam and Peter Szakalos
Submitted to the Journal of Nuclear Materials
Paper IV: Optimizing the oxidation properties of FeCrAl
alloys at low temperatures
Jesper Ejenstam, Bo Jönsson and Peter Szakalos
Letter manuscript
xi
The author’s contribution to the included papers:
Paper I:
The author performed the major part of the planning
and the experimental work, the SEM-EDS evaluation,
and the interpretation of the results. The paper was
mainly compiled by the author.
Paper II:
The author performed the major part of the planning
and the experimental work, the evaluation, and the
interpretation of the results. The Atom Probe
Tomography evaluation and the theoretical calculations
were carried out by the co-authors. The paper was
mainly compiled by the author.
Paper III: The author performed the major part of the planning
and the experimental work, the SEM-EDS evaluation,
the thermodynamic calculations, and the interpretation
of the results. The paper was mainly compiled by the
author.
Paper IV: The author performed the major part of the planning
and the experimental work, the SEM-EDS evaluation,
the thermodynamic calculations, and the interpretation
of the results. The paper was mainly compiled by the
author.
xii
Publications not included in the thesis Paper V:
Finding the optimal nitride fuel and cladding
combination - A strategy for fuelling ELECTRA
Merja Pukari, Jesper Ejenstam, Peter Szakalos and
Janne Wallenius
FR13, Paris, France, 2013, IAEA-CN-199-226
Paper VI: ELECTRA-FCC: A Swedish R&D centre for
generation IV systems
Janne Wallenius et al.
FR13, Paris, France, 2013, IAEA-CN-199-149
Conference contributions Experimental and commercial FeCrAl alloys exposed for 2,500
h in liquid lead at 550 °C
Jesper Ejenstam, Peter Szakalos and Bo Jönsson
Oral presentation at the Nordic Generation IV (NOMAGE 4)
seminar, Halden, Norway (2011)
Oxidation properties of silicon-alloyed ferritic and austenitic
steels exposed to liquid lead at 550 °C for 2,500 h
Jesper Ejenstam and Peter Szakalos
Poster presentation at the 8th International Symposium on HighTemperature Corrosion and Protection of Materials (HTCPM), Les
Embiez, France (2012)
Evaluation of the oxidation properties of 10 wt. % Cr FeCrAl
alloys exposed to liquid lead at 550 °C for 10,000 h
Jesper Ejenstam, Mats Halvarsson, Jonathan Weidow, Bo Jönsson
and Peter Szakalos
Oral presentation at the Nuclear Materials Conference (NuMat),
Osaka, Japan (2012)
xiii
Selection of cladding material for ELECTRA-FCC
Jesper Ejenstam, Merja Pukari, Peter Szakalos and Janne Wallenius
Oral presentation at the European Nuclear Young Generation
Forum (ENYGF), Stockholm, Sweden (2013)
Corrosion resistance and microstructural stability of
Fe10CrAl-RE alloys
Jesper Ejenstam and Peter Szakalos
Oral presentation at the 4th Conference on Heavy Liquid Metal
Coolants in Nuclear Technologies (HLMC), Obninsk, Russia
(2013)
Effect of reactive elements in Fe10Cr4Al alloys exposed to
liquid lead up to one year
Jesper Ejenstam and Peter Szakalos
Oral presentation at the Nuclear Materials Conference (NuMat),
Clearwater Beach, USA (2014)
xiv
Preface Materializing the lead-cooled nuclear reactor requires
multidisciplinary efforts. All systems, physical or chemical, are
intimately connected, and complete understanding of the
interactions between these complex systems are imperative in order
to get a reactor licensed. The interaction of the coolant, nuclear fuel
and structural components under neutron irradiation is just an
example of such a complex system.
The development of corrosion resistant materials for the leadcooled reactor is indeed a single piece of a very large puzzle, albeit
an indisputable key issue to solve to ensure safe, economic and
sustainable operation.
Alumina-forming alloys, in particular FeCrAl alloys, have in recent
years surfaced as one the most promising materials to mitigate the
corrosion issue that liquid lead invokes. It has been shown that
protective alumina layers form even at temperatures as low as 450
°C, which is at least 300 °C lower than what was generally believed
to be the temperature limit some 5-10 years ago.
Although the thesis is written from a materials scientist’s point of
view, with the goal to bridge gaps between reactor physics and
materials science. In addition, the technical level has been adapted
so that it may be attractive for anyone interested in the field of
materials development, e.g. industry R&D personnel, regulators or
legislators.
Jesper Ejenstam
Stockholm, 2014
xv
Table of Contents Introduction ..................................................................................... 1
Aim of the thesis ................................................................................... 1
Climate change mitigation ................................................................... 1
Conventional nuclear power ................................................................ 3
Generation IV reactor systems ............................................................ 4
The lead-cooled fast reactor ................................................................ 6
Properties of lead and LBE ................................................................. 7
Alumina-forming alloys ..................................................................... 10
FeCrAl .............................................................................................. 10
FeNiCrAl .......................................................................................... 12
Theory ............................................................................................ 15
High-temperature oxidation of metals.............................................. 15
Metal oxide formation ...................................................................... 15
Kinetics of oxide growth .................................................................. 16
Transport properties in metal oxides ................................................ 18
The reactive element effect .............................................................. 19
The third element effect ................................................................... 20
Oxygen control in liquid lead ............................................................ 20
α-α’ phase separation in FeCr alloys ................................................ 24
Experimental ................................................................................. 27
Production of experimental alloys .................................................... 27
Thermodynamic modeling ............................................................... 27
Alloy production, analysis and forging ............................................ 28
Materials and sample preparation .................................................... 28
Experimental setup ............................................................................. 30
Main analytical techniques ................................................................ 31
Summary of results ....................................................................... 33
Oxidation properties of FeCrAl alloys (Papers I & IV) .................. 33
Influence of Al ................................................................................. 33
Influence of reactive elements ......................................................... 35
Microstructural stability of FeCrAl alloys (Paper II) ..................... 38
Oxidation properties of alumina-forming austenitic stainless steels
(Paper III) ........................................................................................... 40
xvii
Conclusions .................................................................................... 43
Future work ................................................................................... 45
References ...................................................................................... 46 xviii
Introduction Aim of the thesis The main aim of this work is to investigate the long-term corrosion
resistance of FeCrAl alloys, and in particular for alloy
compositions resistant to α-α’ phase separation (loss of ductility).
Moreover, the corrosion resistance of the studied FeCrAl alloys is
mainly evaluated at environmental conditions known to cause
severe corrosion damage to reference steels, i.e. at temperatures
above 500 °C and at dissolved oxygen (O) concentrations lower
than 10-6 wt. %. The ultimate goal is to obtain a novel mechanistic
understanding of prevailing oxidation properties of FeCrAl alloys
in liquid lead (Pb), knowledge that forms a scientific platform in
the development of alloys that are applicable in lead-cooled fast
reactors.
An important part of this work is to evaluate effects of reactive
element additions in low-Cr FeCrAl alloys in the liquid lead
environment.
Climate change mitigation The imminent threat of global warming requires radical changes in
human lifestyle. In order reach the 2 °C goal, i.e. the maximum
global average temperature increase that is tolerable to avoid
serious environmental effects, mankind needs to reduce its
dependency of fossil fuels. A large amount of the emission of
greenhouse gasses, e.g. CO2, originates from electricity production
through burning of fossil fuels. Coal, gas, and oil must hence be
phased out in order to meet the set goals. In the 2014 climate
1
change mitigation report [1] presented by the Intergovernmental
Panel on Climate Change (IPCC), it is stated that in order to meet
the 2 °C goal by 2050, it is imperative that the share of low-CO2
energy sources in the energy mix grow at least threefold, from a
current level of about 30 % (Fig. 1 [2]). A low-CO2 energy source
is defined by the source of energy that, on a life-cycle basis, emits
≤5 % compared to fossil fuel sources. Alongside renewable energy
sources, nuclear power meets the ≤5 % threshold and is thus a
cornerstone in the mitigation of climate change [1]. In order to
transition to low-CO2 energy sources, the IPCC foresees the need
to raise annual investments by 147 billion dollars globally. Today,
fossil fuels are the fastest growing sources of energy. Governments
worldwide are hence strongly encouraged to change their energy
policies in order to faster introduce low-CO2 energy sources in the
energy mix.
Historical worldwide electricity generation by source
25000
Terawatt hours (TWh)
20000
Coal and Peat
Oil
Natural gas
Nuclear
Hydro
Renewables and other
15000
10000
5000
0
1975
1980
1985
1990
1995
2000
2005
2010
(1971−2011)
Fig. 1. Compiled OECD data on the worldwide electricity generation by source between
1971 and 2011.
2
Conventional nuclear power Nuclear power, as we now it today, is to many the black sheep in
energy production. The accidents at Three Mile Island (1979),
Chernobyl (1986), and Fukushima (2011) have rightfully put the
safety of nuclear power plants under severe scrutiny. Despite the
severity of these accidents, relatively few human beings have been
lethally harmed due to radiation exposure. The United Nations
(UN) World Health Organization (WHO) claims in a report from
2005 that a total of 4000 tragic deaths can be traced back to the
Chernobyl accident [3]. In the case of Fukushima, WHO
anticipates no measurable increase in the cancer rate in Japan [4].
Despite the nuclear accidents, nuclear power has among the lowest
mortality of all available energy sources [5]. Having said that, an
accident at a nuclear power plant has a tremendous impact on the
society in terms of economic damage due to the massive
decontamination work needed.
The aging fleet of nuclear power plants in operation on a global
basis does not have passive emergency cooling systems. If a station
blackout occurs, i.e. all power is lost at the plant, auxiliary power
has to be transported to the plant. This if often done using mobile
diesel generators, which can be deployed to the plant on a short
notice. The need of active emergency cooling is the main Achilles
heel of most nuclear power plants in operation today. This is also
what caused the Fukushima accident. In modern designs,
generation III+, this issue is solved by integrating a passive cooling
system that will make the power plants independent of auxiliary
power supplies in case of a station blackout [6].
3
Another major problem with conventional nuclear power is the
waste issue. The spent nuclear fuel from a nuclear power plant
contains several radiotoxic elements with a very long half-life,
meaning that they need to be safely stored for up to about 100,000
years [7]. The main radionuclides responsible for the need of such a
long storage time are the transuranic elements: Np, Pu, Am and
Cm. Through recycling of said nuclear waste, the storage time
would decrease to about 500-1000 years. In the generation IV
reactor, the transuranic elements can be utilized as nuclear fuel,
hence solving the major issues with the nuclear waste.
Generation IV reactor systems Next generation nuclear reactor systems, denoted generation IV,
have several advantages that make them superior to preceding
reactor technologies. However, the technology leap needed to make
the technology viable on a commercial market requires significant
research efforts to be overcome. Nonetheless, the ideas behind the
generation IV concept are nothing new. In fact, they have been
around since the dawn of nuclear power reactors [8]. Several
research and demonstration reactors have been built and operated
over the years, yet none have truly fulfilled the criteria of a
generation IV reactor.
The members of the Generation IV International Forum (GIF) have
together decided on a set of criteria that a nuclear power system has
to fulfill in order to be entitled generation IV. These criteria invoke
sustainability, economics, safety and proliferation resistance. There
are in total eight criteria, which are stated as follows [9]:
4
ü Produce sustainable energy by enabling a more efficient use
of nuclear fuel
ü Minimize the amount of nuclear waste and its long-term
radiotoxicity
ü The cost of the produced energy should be considerably
lower than alternative energy sources
ü Offer a low financial risk for investors
ü Provide undisputable reliability and safety
ü Have a very low probability of suffering from core damage
ü Eliminate the need of active emergency cooling
ü Have improved protection against terrorism acts, and
operate on fuel that is undesirable for weapon production
Fulfilling each and every one of these goals is necessary to qualify
as a generation IV nuclear system, and it shines a light on why a
great technology leap is needed to materialize such a reactor
system. The benefit to the society when the generation IV reactors
start operating will be immense, thus motivating researchers all
around the world to continue their work.
There are in total six reactor concepts that all have been selected as
most likely to fulfill the above-described criteria
[9]:
-­‐
-­‐
-­‐
-­‐
-­‐
-­‐
Lead-cooled fast reactor (LFR)
Sodium-cooled fast reactor (SFR)
Gas-cooled fast reactor (GFR)
Molten salt reactor (MSR)
Super-critical water reactor (SCWR)
Very high temperature reactor (VHTR) 5
The lead-­‐cooled fast reactor The use of a lead-based coolant in a nuclear reactor has been
investigated for more than 60 years. Several countries have pursued
the technology, however, the central part of the research has been
carried out in Russia (and former Soviet). In 1952, the work on
using lead as a nuclear coolant was initiated at the Institute of
Physics and Power Engineering (IPPE) in Obninsk, Russia [10-12].
The work was supervised by A.I. Leipunsky, today considered the
founding father of the lead-cooled fast reactor (LFR). The efforts
carried out at IPPE resulted in the 70 MWth 27/VT reactor, which
reached first criticality in 1959 [11]. Following the success of the
program, Russia decided to use the LFR technology as a propulsion
system for submarines. A total of eight submarines were built and
operated. One submarine went under the designation ”project
645”, and had two reactors on board. The seven other submarines,
under the designation ”project 705”, all utilized a single reactor.
The project 705-submarines are also known by the NATO
designation ”Alpha class”. In addition, the hull of the Alpha class
submarine was completely made out of Ti, making the submarine
extremely light. As a result of the use of Ti and the LFR
technology, these submarines were (and are still) the fastest
submarines ever constructed1. The Alpha class submarines were in
operation for roughly 20 years, after which they were
decommissioned and scrapped.
As of today, no LFR is in operation in the world. However, there
are several projects ongoing worldwide that aims at making the
technology viable, e.g. MYRRHA, ALFRED, SVBR-100 and
1
Guinness world records
6
BREST-300. MYRRHA (Multi-purpose Hybrid Research Reactor
for High-Tech Applications) is a 50-100 MWth nuclear reactor
cooled by means of a lead-bismuth eutectic (LBE) [13]. This
reactor is developed by SCK-CEN in Belgium and is expected to
be operational around 2023. The ALFRED (Advanced Lead-cooled
Fast Reactor European Demonstrator) reactor will be cooled by
pure lead and have an effect of 120 MWe (300 MWth) [14]. The
concept is being developed by a European consortium, coordinated
by Ansaldo Nucleare, and is planned to be built in Romania,
starting around 2025. In Russia, two lead-cooled reactor concepts
are planned, the LBE-cooled SVBR-100 (Svintsovo-Vismutovyi
Bystryi Reaktor) [12, 15] and the lead-cooled BREST-300 (Bystry
Reaktor so Svintsovym Teplonositelem). First out is the SVBR-100
(100 MWe) reactor, designed by AKME Engineering, which is
planned to be operational around 2017. The larger Russian LFR,
BREST-300 (300 MWe), is aimed to be started around early 2020’s
[16], and like ALFRED, it will be cooled by pure lead [17]. The
BREST-300 is developed by the Moscow-based NIKIET institute
and is expected to be built near Seversk in the Tomsk region in
Russia.
Properties of lead and LBE Liquid lead and LBE are highly suitable nuclear coolants for
several reasons. Unlike water, lead and LBE do not moderate
neutrons. An LFR will hence utilize fast neutrons, i.e. a fast
spectrum. In a fast spectrum, not only 235U can be used as fuel. In
fact, the transuranic elements present in the current nuclear waste
are the cause for the long storage time of light-water reactor (LWR)
nuclear waste. The transuranic elements all have a high fission
probability in a fast spectrum and can thus be burned. From this
7
follows a significant reduction of both the amount of the nuclear
waste and its storage time. In addition, in a fast spectrum, depleted
U (238U) can be breed into fissile Pu (239Pu) [8]. The abundance of
238
U in natural U is about 99.3 %. The use of fast reactors can
hence significantly increase fuel resources. In comparison to
sodium, another fast reactor coolant, lead and LBE do not react
violently with water, an effect that eliminates the need of a
secondary heat exchanger circuit. A single steam generation circuit
can instead be inserted directly into the coolant. The high boiling
point of lead (1750 °C) and LBE (1670 °C) minimizes the risk of
losing the coolant during an accident, when the fuel may reach very
high temperatures. Moreover, the relatively large change in density
with temperature of lead and LBE makes it possible to cool a small
reactor with natural circulation of the coolant [18, 19]. In larger
reactors, the natural circulation can be utilized to cool the reactor
core after shutdown, i.e. to remove decay heat. This is a paramount
feature of the lead coolant as it makes the reactor inherently safe
through the naturally integrated passive cooling system. Chemical
bonding of radiotoxic fission products, such as I and Cs, to the
coolant is yet another important safety feature of lead and LBE
[20]. In the case of a fuel cladding tube rupture, the release of
hazardous gasses to the environment can thus be limited.
As presented, there are several evident advantages of using lead or
LBE as nuclear coolants. However, there are also some
disadvantages that need to be considered. The predominant issue is
the compatibility between the coolant and structural steels. At
elevated temperatures liquid metal corrosion (LMC) occurs in steel
components in contact with the liquid metal [21-23]. This
phenomenon is governed by the solubility of the steel constituents
in the molten metal. The corrosion rate increases with increasing
8
solubility of a particular element. Stainless steels show fast
dissolution rates in liquid lead at temperatures above 500 °C due to
the high dissolution rate of Ni [22]. Ferritic steels, on the other
hand, perform better, albeit are still attacked by the liquid metal at
the same temperatures. The formation of protective oxide layers on
the steel surfaces is required to reduce the corrosion rate. This can
be achieved by dissolving oxygen into the liquid lead. The amount
of dissolved oxygen needs to be high enough to ensure formation
of the metal oxide, often chromia (Cr2O3) on the stainless steel
surface, but less than the amount needed to oxidize the liquid metal
itself. Chromia-layers have been shown to provide adequate
corrosion resistance up to about 500 °C, after which dissolution
attack quickly occurs [24]. As the lead has a melting point of about
328 °C, a lead-cooled reactor needs to operate at a high enough
temperature to mitigate the risk of coolant freezing during
operation. The operation temperature interval often ranges from
400 °C to 600 °C, which challenges structural components. Here,
LBE has an advantage through its lower melting point of about 124
°C. An LBE-cooled reactor may hence operate at lower
temperatures, often in the interval of 200 to 500 °C, a temperature
interval where many conventional steels are able to form slow
growing adherent protective oxide layers [25]. On the other hand,
Bi is a rare metal, which makes LBE a much more expensive
coolant than pure lead. In addition, significant amounts of 210Po are
formed when Bi is irradiated in a fast neutron spectrum [10, 11].
This is a major issue as 210Po is a strong α-emitter and thus highly
radiotoxic.
Due to the availability of Bi and the issue with 210Po formation
under irradiation, a reactor cooled with pure lead is the most viable
option for a commercial LFR technology. However, material issues
9
at higher temperatures (>500 °C) must be solved in order to realize
the LFR concept. In recent years, alumina-forming (Al2O3) alloys
have gained a lot of interest as potential candidate materials for
LFRs due to their superior oxidation properties, compared to
chromia-forming steels. Several studies have shown that FeCrAl
alloys have excellent oxidation properties in the entire LFR
operation temperature range, thus realizing the lead-cooled
technology may be possible sooner than anticipated a few years ago
[26-32].
Alumina-­‐forming alloys Alloys that can form a protective alumina layer on their metal
surfaces at high temperatures are among the most oxidation
resistant alloys currently available. Alumina has high chemical
resistance, and the oxide layer becomes thin due to the low
diffusivity of oxygen [33]. Most often the ferritic FeCrAl alloys are
considered for high-temperature applications, however, lately also
austenitic FeNiCrAl alloys have gained a lot of interest in the
scientific literature [34-40].
FeCrAl Hans von Kantzow developed the FeCrAl alloy in the 1920’s [41].
He found that by adding Cr to a FeAl alloy, it was possible to
produce an alloy with excellent oxidation properties and
workability. In addition, this new type of alloy revealed high
electrical resistance, significantly higher than what was available at
that time. The reason the new FeCrAl alloy was superior to the
FeAl alloy was that the addition of Cr reduced the critical Alcontent needed for a protective alumina-layer to form. In FeAl
alloys, the Al-limit was said to be 16 wt. % for manufacturability
10
reasons. By adding Cr, the critical Al-content was significantly
decreased. In many cases the addition of 3 wt. % Al was enough to
ensure proper oxidation properties. At these Al-concentrations, the
final product, e.g. steel wires, were considerably easier to produce.
Many years later, the theory on the positive combinatory effect of
Cr and Al was presented as the “third element effect” [42-44]. von
Kantzow presented his findings in a patent, entitled ”Fire resistant
alloy with high electric resistance” [41]. In 1931, he created the
company AB Kanthal in Hallstahammar, Sweden. The main
applications of the newly developed FeCrAl alloy, nowadays
commonly known as the Kanthal alloy, were resistance wires and
heating elements operating at a temperature interval between 900
and 1400 °C [45]. The ability to form thin protective aluminalayers at these temperatures is crucial for the service life of the
components and the prerequisite for AB Kanthal to maintain a
world-leading position for decades. The conventional FeCrAl
alloys typically contain Fe (balance), 15-25 wt. % Cr and 3-6 wt. %
Al. In addition to the main constituents, small amounts (<1 wt. %)
of reactive elements (RE) are added to improve oxide formation
and its surface adherence [43, 46-48]. Reactive elements refer to
elements with high affinity to e.g. O, C, and N, and include rare
earth metals such as Y and Ce. Other common reactive elements
are Zr, Ti, and Hf.
The Cr-concentrations of conventional FeCrAl alloys make them
prone to α-α’ phase separation, and thus embrittlement at
temperatures up to about 500 °C [49-51]. FeCrAl alloys with lower
Cr-content have to be utilized for use below 500 °C. Lately, the
interest in low-Cr FeCrAl alloys has increased dramatically. Within
the framework of the development of accident tolerant fuel
cladding materials, initiated after the Fukushima accident, α-α’
11
phase separation resistant FeCrAl alloys been proposed as
successors to the conventional Zircaloy cladding tube material
currently used [52-54].
FeNiCrAl More recently, austenitic FeNiCrAl alloys were presented in the
scientific literature. These alloys are also known as aluminaforming austenitic stainless steels and often denoted AFA. The
austenitic crystal structure, face-centered cubic (FCC), provides the
alloy higher mechanical strength compared with the ferrite that has
a body-centered cubic (BCC) structure [34]. However, the
diffusivity of Al and Cr is slower in FCC compared to BCC
materials. Several research groups have tried to produce FeNiCrAl
alloys with Al-contents of 4 wt. % or more to compensate for the
lower diffusivity of Al [55-58]. However, even at this apparently
modest Al level, issues arise due to its strong ferrite stabilizing
effects. Until 2007, the production of a single-phase austenitic
FeNiCrAl alloy seemed impossible. Up to recently, all attempts
resulted in the formation of ferrite in the austenitic microstructure,
with a reduction in mechanical strength at temperatures above 500
°C as the evident consequence. In April 2007, Yamamoto et al.
presented the work by Oak Ridge National Laboratory (ORNL) in
the USA [34]. They found that it was possible to produce an AFA
alloy with only 2.5 wt. % Al that was able to form protective
alumina on its surfaces when exposed to an oxidizing environment
at high temperatures. Moreover, the alloy remained fully austenitic
in a wide temperature interval. It was found that Nb had a positive
effect on the oxidation resistance by increasing the solubility of Cr
and Al in the austenitic phase, hence enabling alumina to form even
at low Al-concentrations in the bulk [38]. Another positive effect of
12
the addition of Nb was that nano-sized Nb-carbides (NbC) were
evenly distributed in the matrix, which increased the high
temperature creep strength of the alloys.
13
14
Theory High-­‐temperature oxidation of metals Metal oxide formation Oxidation of a metal species occurs when it comes in contact with
an environment containing oxygen given that the oxygen activity
(partial pressure in the case of gases) in the environment is higher
than the equilibrium activity for the metal oxide [59-62]. The
standard Gibbs free energy of formation (∆𝐺 ° ) of a metal oxide is
written as
∆𝐺 ° = 𝑅𝑇𝑙𝑛(𝑝𝑂! )
(1)
where R is the Boltzmann constant, T is the temperature (in Kelvin)
and 𝑝𝑂! is the oxygen partial pressure (in bar). The relation
between ∆𝐺 ° and the temperature is given in an Ellingham diagram
[63], Fig. 2 [61]. Information can be obtained from the Ellingham
diagram about the standard Gibbs free energy for a metal oxide
equilibrium at a desired oxygen partial pressure and temperature.
The stronger the oxide former, the lower in the Ellingham diagram
it appears. The Ellingham diagram is a very helpful tool, which
quickly provides a good overview of thermodynamically stable
metal oxides in a certain oxidizing environment and vice-versa.
However, the diagram does not provide any information on the
kinetics of oxide growth.
15
H /H O
2 2
10-8
10-6
pO
2
10-4
10-2
1
0
O
=
+O 2
O4
e
4F 3
-200
O3
e
6F 2
2
+
Ni
1
O2
NiO
=2
2H 2
+ O2
10-4
O
= 2H 2
102
o
ΔG = RT ln pO (kJ/mole O )
2
2
-400
H
+
Cr
4/3
-600
O2
Si +
=
O2
104
SiO 2
A
4/3
10-10
10-12
l O3
/3A 2
-800
10-6
10-8
O
Cr 3
2/3 2
=
10-2
l
=2
+O2
106
10-14
10-16
-1000
108
10-18
10-20
-1200
0
400
800
1200
1600
1800
2000
Temperature (°C)
1010
10-22
H /H O
2 2
0 Kelvin
10-200
10-100
10-70 10-60
10-50
10-42
10-38
10-34
10-30 10-28 10-26
10-24
pO (bar)
2
Fig. 2. The Ellingham diagram shows the thermodynamic stability of metal oxides as a
function of temperature and oxygen partial pressure.
Kinetics of oxide growth The growth kinetics of a metal oxide is normally described by one
of the following three rate laws: logarithmic, parabolic or linear
law [59-62].
16
Logarithmic rate law The logarithmic rate law is governed by an electrical field across
the metal oxide and is commonly applied to metals oxidized at low
temperatures. The growth is characterized by an initially rapid
oxidation, after which the oxide growth is seemingly stopped. The
logarithmic rate law is written as
𝑥 = 𝑘! log (𝑎𝑡 + 𝑏)
(2)
where 𝑥 is the oxide thickness (in meters), 𝑘! is the rate constant, 𝑡
is the time (in seconds), and 𝑎 and 𝑏 are constants.
Parabolic rate law Parabolic growth occurs when the oxidation rate is governed by
diffusion of ions through the metal oxide layer. As the layer grows
in thickness, the oxidation rate decreases as the diffusion length
increases. Hence, the oxide thickness will be inversely proportional
to the growth rate. The parabolic rate law is written as
𝑥 ! = 𝑘! t + c
(3)
where 𝑥 is the oxide thickness (in meters), 𝑘! is the rate constant, 𝑡
is the time (in seconds), and 𝑐 is a constant.
Linear rate law For a typical non-protective oxide that cracks, spalls or by other
means fails to protect the underlying metal from further oxidation,
the oxidation follows a linear growth. The linear rate law is
normally written as
17
𝑥 = 𝑘! 𝑡
(4)
where 𝑥 is the oxide thickness (in meters), 𝑘! is the rate constant
and 𝑡 is the time (in seconds).
Transport properties in metal oxides Oxidation is a chemical process, which belongs to the group of
reactive diffusion processes [59-62]. The growth of the oxide
depends on the transportation process of charged particles, i.e.
anions and cations. Two diffusion processes, solid-state diffusion
and short-circuit diffusion, mainly govern the transport of ions,
given that the formed oxide is gas-tight. Solid-state diffusion is
controlled by microscopic defects, such as vacancies and
interstitials, in the oxide layer. Short-circuit diffusion, on the other
hand, is controlled by macroscopic defects in the oxide layer.
Examples of such defects are pores, grain boundaries and cracks.
The diffusion rate is faster at macroscopic defects due to the lower
activation energy needed for the ion to move. At lower
temperatures, short-circuit diffusion dominates due to the lower
activation energy, whereas solid-state diffusion dominates at higher
temperatures. The total area of grains is obviously much larger than
the grain boundary area, solid-state diffusion will thus dominate
when the temperature is high enough to assist ion transportation in
microscopic defects.
18
The reactive element effect The reactive element (RE) effect on the oxidation behavior of hightemperature alloys is central in high-temperature oxidation, and has
been known since Pheil patented it in 1935 [64]. By adding small
amounts, typically less than 1 wt. %, of elements with high affinity
to O, N or C, such as Y, Zr, Ti, Ce, Hf, Sc, La and Th, the
oxidation properties of alloys at high temperatures can be
improved. Several theories have been proposed, but the underlying
reasons are to a large extent still open and remain to be answered.
Despite the lack of a cohesive understanding of the effect, some
major influences of RE-additions are well known [43, 46-48]. The
oxide layer adhesion is improved by the addition of RE to a hightemperature alloy. In FeCrAl alloys, Y-additions have shown to
form Y-Al mixed oxides, denoted Y-Al garnets (YAG) that peg the
alumina layer to the metal substrate. During thermal cycling, the
pegging effect reduces buckling and thus increases the overall
adhesion of the oxide layer. Moreover, RE-additions have been
shown to alter the diffusivity of metal ions in the matrix. A
balanced oxide growth is achieved by balancing the outward
diffusion of metal ions and the inward diffusion of oxygen. It leads
to a reduced presence of interfacial pores and oxide layer stresses,
important to avoid spallation.
Moreover, Zr has been shown to thermodynamically stabilize finegrained microstructures in alloys despite long exposure times at
elevated temperatures [65, 66]. As the oxidation properties of a
high-temperature alloy is influenced by the grain size [67], the
effect of RE on the microstructure may be particularly important at
lower temperatures where grain boundary diffusion is the
dominating path for ion transportation.
19
The third element effect FeAl alloys form protective alumina layers at high temperatures at
a minimum Al-content around 9 wt. % [68]. The addition of Cr to
the FeAl system lowers the critical Al-content to about 3 wt. %
[33]. This effect is known as the “third element effect”. In 1965,
Wagner presented the classic theory that has gained wide
acceptance [42]. Wagner proposed that binary alloys that contain
elements with significantly different affinity to oxygen, e.g. Al and
Fe, would suffer from internal oxidation if the concentration of the
oxygen active element were too low. Upon the addition of an
element with an intermediate oxygen affinity, Cr in this case, the
oxygen activity at the metal-oxide interface is reduced through
oxidation of the same element. It has been suggested that a thin
layer of chromia forms during the initial stages of oxidation of
FeCrAl alloys, which is later to be incorporated in the alumina
layer [69, 70]. Moreover, theoretical calculations have shown that
the addition of Cr increases the Al driving force to diffuse to the
metal surface [71].
Oxygen control in liquid lead Controlling the oxygen concentration in liquid lead is crucial for
two main reasons [22, 72-75]:
-­‐
-­‐
to avoid precipitation of lead oxide
to avoid metal dissolution in structural steels through
formation of protective surface oxides
Typically when discussing oxidation of metals, the process takes
place in an oxygen-containing gaseous environment, where the
20
stability of a metal oxide depends on the oxygen partial pressure
( 𝑝𝑂! ) and the temperature. In a liquid metal, the oxygen
concentration depends on the oxygen activity. Whereas, the oxygen
partial pressure is equal to the oxygen activity in an ideal gas, the
activity of oxygen in a solution is given by
𝑎! = 𝛾! 𝑐!
(5)
where, 𝑎! is the oxygen activity, 𝛾! is the activity constant, and 𝑐!
is the oxygen concentration. For dilute solutions, Henry’s law may
be applied. Thus, the activity constant is independent of the activity
and the concentration [74, 76]. Liquid lead with dissolved oxygen
can be considered a dilute solution, and experiments have shown
that the activity constant is independent of the change in oxygen
concentration up to about 700 °C [76]. The activity constant is
preferably calculated at the oxygen saturation level where the
activity is unity. Now equation (5) can be re-written as follows
𝑎! =
𝑐!
𝑐!,!"#
(6)
The oxygen concentration in a lead melt that interacts with an
oxygen-containing gas can now be determined provided that
equilibrium conditions have been reached. At equilibrium, the
oxygen activity in the liquid lead must be the same as in the gas.
Sieverts’ law states that gaseous oxygen dissolves into liquid metal
as atomic oxygen, and that the square root of 𝑝𝑂! is proportional to
the dissolved oxygen concentration. As Henry’s law has been
shown to be valid up to saturation concentrations, it follows that
21
𝑎! =
𝑐!
𝑐!,!!"
=
𝑝𝑂!
𝑝𝑂!,!"#
(7)
hence, the oxygen concentration can be controlled by a gas phase.
To avoid lead oxide formation at e.g. 550 °C, the 𝑝𝑂! in the gas
phase has to be lower than 10-17 bar [73]. At these levels, it is not
possible to use e.g. Ar and 𝑂! mixtures. Instead, 𝐻! and 𝐻! 𝑂 gas
mixtures can be utilized, and the 𝑝𝑂! is given by
𝑝𝐻! 𝑂
𝑝𝑂! =
𝑝𝐻!
!
𝑒𝑥𝑝
2∆𝐺 𝐻! 𝑂
𝑅𝑇
(8)
where ∆𝐺 𝐻! 𝑂 is the Gibbs free energy of formation of water
vapor, 𝑅 is the gas constant and 𝑇 is the temperature (in Kelvin).
Combining equation (7) and (8) thus yields
𝑝𝐻!
𝑐!,!"#
∆𝐺 𝐻! 𝑂
=
𝑒𝑥𝑝
𝑝𝐻! 𝑂 𝑐! 𝑝𝑂!,!"#
𝑅𝑇
(9)
In 1998, Risold et al. [77] summarized the reported literature
findings on oxygen saturation concentrations. A good agreement
between experiments was found in the temperature interval of 350
°C (623 K) to 1050 °C (1323 K). The temperature dependence is
given as
𝑙𝑜𝑔𝑐!,!"# = 1.64 −
22
3503
, 623 𝐾 < 𝑇 < 1323 𝐾
𝑇
(10)
where 𝑐!,!"# is given in wt. %. By combining equations (9) and
(10), dissolved oxygen concentration levels may be introduced into
the Ellingham diagram, illustrated in Fig. 3 [73].
H2/H2O pO2 (bar)
-200
10
Oxygen concentration (wt. %) in Pb
-5
10-4
-250
10-14
PbO
10-3
ΔGo = RT ln pO2 (kJ/mole)
-300
10
10-16
-4
10
-400
10-18
Melting point of Pb
-350
-2
10-1
10 -6
10-20
1
10 -8
/FeO
Fe 3O 4
10
10
1
10
2
10
3
-450
1 0 -10
-500
200
300
400
500
Temperature (°C)
600
-22
10-24
10-26
1028
Fig. 3. The dissolved oxygen concentration in liquid lead depends on the oxygen activity
in the surrounding gas phase. Here, oxygen concentration lines are presented in a classical
Ellingham diagram.
23
α-­‐α’ phase separation in FeCr alloys FeCr alloys belong to the group of alloys that suffers from phase
separation. This means that the microstructure decomposes into a
Fe-rich phase (α) and a Cr-rich phase (α’) in a particular
temperature range [78-81], Fig. 4. This phenomenon leads to a
change in the local composition albeit the crystal structure is
maintained, which in turn leads to a loss of ductility. For the FeCrsystem, the critical temperature, at which the kinetics of the phase
transformation is rapidest, is around 475 °C. The effect is hence
often stated as 475 °C-embrittlement and occurs due to the
presence of a miscibility gap in the binary system [80]. In the
miscibility gap, two types of phase separation processes can occur,
nucleation and growth, or spinodal decomposition [82]. The energy
barrier between α and α’ phases governs nucleation and growth.
The nucleation occurs at sites where the thermodynamic barrier is
lower, e.g. at phase boundaries. When the α’-phase has nucleated, it
grows until it has reached a size where it is thermodynamically
stable in a given environment. Since nucleation and growth
typically occur for alloy compositions that just fall within the phase
transformation regime, the onset depends on the composition. In
contrast to nucleation and growth, spinodal decomposition has no
thermodynamic energy barriers to overcome. Within the spinodal
zone, the α’-phase is favored at all compositions, and hence phase
separation takes place spontaneously.
24
1500
Temperature (°C)
1200
α-FeCr
γ-Fe
900
σ
600
475
α-FeCr + σ
α-FeCr + σ
α-Fe + α’-Cr
300
0
20
40
60
Chromium (wt. %)
80
100
Fig. 4. The classical FeCr phase diagram, indicating stable phases as a function of
temperature and Cr-content.
In 1938, Becket first presented the loss of ductility in a FeCr alloy
at a temperature around 475 °C [83]. The presence of the
miscibility gap, which cause precipitation of the Cr-rich α’-phase
and subsequently embrittlement, was established by Williams and
Paxton in 1957 [80]. Around the same time, work on developing a
theoretical explanation of spinodal decomposition was carried out
by Hillert [84], and Cahn and Hilliard [85].
For the neutron irradiated FeCr alloys, α-α’ phase separation has
been shown to occur down to Cr-concentrations at around 9 wt. %
[86-88]. As the thermodynamic barrier governs the onset for
nucleation and growth, it can be overcome by adding energy to the
system. Heat and especially irradiation are examples of such energy
injections, where phase separation is accelerated due to an
increased diffusion rate in the system. In 1987, Andersson and
Sundman presented an extensive work on the FeCr system in which
the miscibility gap was thoroughly mapped [78]. The work was
25
based on experiments carried out at temperatures exceeding 427 °C
(700 K) and extrapolated to lower temperatures. Since then, several
studies, based on theoretical calculations in the low-temperature
range, have shown that the influence of Cr on the miscibility gap
has been overestimated [79, 89-91]. Rather than a reduction to
nearly 0 wt. % at 27 °C (300 K), the solubility of Cr in the α-phase
has been shown to be almost constant around 8 wt. % at
temperatures below 427 °C (700 K). The discrepancy in presented
results at low temperatures is important to resolve as many nuclear
candidate materials may operate within this temperature range.
Conventional FeCrAl alloys, with a Cr-content of about 20 wt. %
also suffer from α-α’ phase separation and are thus embrittled in
the temperature regime up to about 500 °C [49-51]. However, since
Al has been shown to stabilize the α-phase, the critical Cr-content
of FeCr alloys is increased when Al is added [51, 92].
26
Experimental Designing, producing and evaluating innovative alloy compositions
requires a rigorous work effort. Results communicated in the
scientific literature are often condensed for obvious reasons. The
immense work leading up to the conclusions are often taken for
granted, which is an issue for those not familiar with the field of
research. The goal of this study is to investigate the mechanisms
related to long-term corrosion resistance of low-Cr FeCrAl alloys.
In order to be able to understand the underlying effects the author
has been involved in each and every step of the alloy development,
an iterative process simplified in Fig. 5.
Thermodynamic
modeling
Casting
Chemical
analysis
Forging
Sample
preparation
Experiments
Analytical
evaluation
Conclusions
Fig. 5. Schematic overview of the iterative process of the alloy development.
Production of experimental alloys Thermodynamic modeling Thermodynamic modeling softwares are powerful tools when
designing a new alloy. By calculating the expected phases as a
function of temperature, it is possible to map what phases are
expected at the alloy melting temperatures and in the intended
27
operation temperature regime. A drawback of standard
thermodynamic calculations is that the kinetics is not taken into
account. Hence, calculated results provide steady-state conditions,
which may take a long time to reach. The models will thus not
provide a full answer, yet are useful starting points in the work of
creating a new alloy. In the work covered by this thesis, the
Thermo-Calc software [93, 94] was used. The main alloy databases
utilized included TCFE5, TCFE6, TCFE7, SSOL4 and SSOL5.
Alloy production, analysis and forging The alloys studied in this thesis were cast in a vacuum furnace
located at the R&D center at Sandvik Heating Technology AB
(former Kanthal AB) in Hallstahammar, Sweden. In this particular
melting furnace, 1 kg batches are typically made, a batch size
suitable for laboratory testing of experimental alloys. Upon
finishing the melting of the batch, the alloy’s chemical composition
is analyzed using inductively coupled plasma optical emission
spectroscopy (ICP-OES) and combustion analysis. Batches
showing significant deviations from the desired composition are
scrapped and recast. Approved batches are subsequently processed
into the desired shape. The central part of the work in this thesis is
based on hot-rolled material.
Materials and sample preparation Chemical compositions of all alloys included in this work are
presented in Tables 1-2. The main part of the work is related to the
development of ferritic FeCrAl alloys (Gen 1 and Gen 2), compiled
in Table 1. Table 2 presents the compositions of the studied
austenitic alloys.
28
Table 1. Analyzed (ICP-OES) chemical bulk compositions (wt. %) of the 1st and 2nd
generation FeCrAl alloys
Alloy
Fe
Cr
Al
Si
Mn
C
Ti
Zr
Y
10Cr-4Al
Bal.
10.00
4.03
0.10
N.A.
0.03
0.08
0.07
-
10Cr-6Al
Bal.
10.06
6.08
0.03
N.A.
0.03
0.07
0.08
-
10Cr-8Al
Bal.
10.09
7.65
0.04
N.A.
0.03
0.07
0.08
-
10Cr-6Al-(No RE)
Bal.
10.03
5.98
0.06
N.A.
0.03
-
-
-
Ref 21Cr-5Al
Bal.
21.05
5.04
0.07
N.A.
0.03
0.08
0.08
-
Zr 0.1
Bal.
10.12
3.98
0.12
0.11
0.04
0.08
0.11
-
Zr-0.2
Bal.
10.15
3.95
0.13
0.11
0.03
0.09
0.21
-
Zr-0.4
Bal.
10.20
4.06
0.12
0.12
0.03
0.07
0.39
-
Y-0.02
Bal.
10.26
4.24
0.07
0.12
0.03
0.07
-
0.02
Y-0.1
Bal.
10.21
4.14
0.12
0.13
0.03
0.07
-
0.09
Y-0.2
Bal.
10.12
4.05
0.12
0.11
0.03
0.08
-
0.19
Gen 1
Gen 2
N.A. - Not Analyzed, O<0.0035, N<0.003, S<0.005
Table 2. Analyzed (ICP-OES) chemical bulk compositions (wt. %) of chromia-forming
and alumina-forming austenitic alloys.
Alloy
Fe
Cr
Ni
Al
Mn
Mo
Si
Nb
Other
AISI 316L
Bal.
16.9
10.1
-
0.95
2.04
0.52
-
0.02 C, 0.03 P, 0.006 B
15-15 Ti
Bal.
15.1
15.0
0.01
1.80
1.22
0.44
-
0.09 C, 0.52 Ti, 0.01 P
20Ni AFA
Bal.
14.2
19.7
2.47
1.64
2.43
0.18
0.87
0.08 C
14Ni AFA
Bal.
14.4
13.9
2.49
1.60
2.50
0.19
0.89
0.08 C
The corrosion tests were carried out using small coupons, typically
measuring 1 × 8 × 30 mm in size. An extensive sample preparation
work was carried out prior to each corrosion study, which included
a final hot-rolling step, heat treatment, marking, polishing of
surfaces and sonication.
29
Experimental setup The central part of the work included in this thesis is related to
corrosion studies of alumina-forming alloys in liquid lead. In
parallel to the development of the first generation of FeCrAl alloys,
an extensive effort was put on the construction of the liquid metal
corrosion lab at the department of Surface and Corrosion Science at
KTH. Apart from installing the corrosion test facility, a gas
handling and climate control systems were installed and tested.
Handling of harmful heavy metals, like lead, requires a special
permission, thus necessary measures were taken to ensure a safe
working environment for everyone involved.
Corrosion studies in liquid lead were carried using a tube furnace.
In order to control the 𝑝𝑂! , and thus the dissolved oxygen
concentration in the liquid metal, hermetically sealed quartz tubes
were inserted into the tube furnace. The oxygen partial pressure
was controlled by means of an 𝐴𝑟-𝐻! -𝐻! 𝑂 gas mixture. Water
vapor was added to the mixture by bubbling the 𝐴𝑟-𝐻! gas through
a humidifier, and the temperature of the water bath controlled the
𝐻! 𝑂 concentration in the gas. The gas mixture was then flushed
through the furnace system, and the 𝑝𝑂! was measured at the
furnace gas outlets using a Zirox SGM5 oxygen analyzer. The
samples being studied were placed in lead-filled alumina crucibles,
and put on Ni trays that were inserted into the quartz tubes. The
corrosion test facility, developed by Karlsruhe Institute of
Technology (KIT), goes under the acronym COSTA, which stands
for “corrosion test facility for stagnant liquid alloys” [72, 73]. A
schematic overview of the COSTA facility is presented in Fig. 6.
30
Gas outlet
Oxygen
sensor
Tube furnace
Ar
Hot zone
Ar / H2
Humidifier
Quartz tube
Crucibles with samples
immersed in liquid lead
Fig. 6. Schematic overview of the corrosion test facility also referred to as COSTA.
Main analytical techniques The analytical results presented are widespread in terms of the size
of studied features and type of features studied. Therefore, several
complementary analytical techniques were employed. The postexposure evaluation was to a large extent related to corrosion
effects on the samples surfaces due to the exposure in liquid lead.
Cross sections for evaluation of corrosion effects are commonly
prepared through molding and polishing, which initially was
carried out for all samples. Thin oxides (≤100 nm) are difficult to
evaluate using this method, hence focused ion beam (FIB) milling
was utilized. In addition, FIB was used to prepare transmission
electron microscopy (TEM) samples by means of the standard liftout technique [95].
Scanning electron microscopy (SEM) was mainly used to
characterize the studied samples whereas TEM was used to the
characteristics of sub-µm alumina layers. Elemental analysis was
carried by means of energy dispersive spectroscopy (EDS), and
31
microstructural characterization using electron back-scatter
diffraction (EBSD). In order to evaluate phase separation on the
atomic level, atom probe tomography (APT) was utilized.
While SEM, TEM and EDS are considered standard techniques in
materials characterization, EBSD and APT may still be unknown to
many. EBSD [96] is mainly used to study crystallographic phases,
orientation and grain boundaries in materials by means of Kikuchi
pattern diffraction [97]. Mapping the material of interest, EBSD
may elucidate homogeneity, favored orientation, impurities and
internal strain. The technique is complementary to electron
microscopes and utilizes the backscattered information, which is
created by the incident electron beam. Conductive samples and fine
polished surfaces are required.
APT [98] or 3D atom probe (3DAP), is a destructive method that
analyzes small volumes of material, typically containing some 10
million atoms. Samples, with tip diameters up to 100 nm, are
prepared by means of electropolishing or FIB milling. The atoms in
the volume are excited and vaporized in an electrical field, and
detected by a position sensitive time-of-flight mass spectrometer.
When the material volume of interest has been consumed, the
volume can be digitally reconstructed for in-depth analysis on the
atomic scale. The technique is mainly used to acquire chemical
information, including isotope differences, of a material of interest,
e.g. homogeneity and impurities.
More information on the techniques and the particular instruments
used is given in Papers I-IV.
32
Summary of results Oxidation properties of FeCrAl alloys (Papers I & IV) Influence of Al The primary scientific goal was to study the long-term oxidation
properties of α-α’ phase separation resistant FeCrAl alloys in liquid
lead at temperatures up to 550 °C. It was anticipated that a two-fold
reduction of the Cr-content compared with a conventional FeCrAl
alloy would affect the oxidation properties negatively. Thus,
assessing the importance of the Al-content was the main goal of
Paper I. Given the compositions of the Gen 1-alloys, a critical Alcomposition in the 4-6 wt. % Al interval was needed to form
protective alumina layers. The influence of the long exposure time
(10,000 h) did not seem to affect the oxide thickness of the 6-8 wt.
% Al alloys. It is thus likely that the alumina layer forms early in
the oxidation process and then reach a near constant thickness, a
behavior that fits well to sub-parabolic growth, or even logarithmic
growth. This type of oxidation behavior is often observed at
relatively low oxidation temperatures. The impact of exposure time
was pronounced for the 4 wt. % Al alloy, which was not able to
form a thin protective oxide. A several µm thick, non-protective
oxide was observed after the shorter exposure (2,500 h). The oxide
was typically comprised of an inward growing mixed metal oxide.
In addition, Al-rich precipitates were found at the metal-oxide
interface. The 10,000 h exposure did not result in a significant
increase in the oxide thickness. This is likely due to the continuous
Al-rich oxide layer that formed at the metal-oxide interface. It was
found that the composition of the 4 wt. % Al alloy did not support
nucleation of alumina early on in the oxidation test. However, as
the non-protective oxide developed, it is likely that the effective
33
𝑝𝑂! present at the metal-oxide interface decreased to a level where
only alumina is thermodynamically stable. Examples of
representative non-protective and protective oxide layers are
presented in Fig. 7.
Pb
FeCrAl
Non-protective mixed oxide
Pt coating
5 µm
a.
Pt coating
Protective alumina layer
Protective alumina layer
FeCrAl
FeCrAl
1.5 µm
b.
c.
50 nm
Fig. 7. a. Backscattered electron SEM micrograph, showing an inward growing nonprotective mixed metal oxide formed on a Fe10Cr4Al alloy exposed to lead for 10,000 h
at 550 °C. Note that lead penetrated into the porous structure of the oxide. b. Typical
characteristics of a thin protective alumina layer when observed in the SEM. c. TEM
bright field micrograph, showing the morphology of a thin alumina layer formed during
long-term exposure to liquid lead at 550 °C.
The powerful influence of reactive elements on the oxidation
properties was also discussed in Paper I. The 6 wt. % Al alloy
without any RE-addition yielded similar oxidation properties as the
4 wt. % Al alloy with RE-additions. These results implied that the
formation of a protective oxide layer at a low Cr-content of 10 wt.
% is more sensitive to optimized RE-additions than findings related
to the critical Al-concentration. The optimization of RE-additions
at low Al-concentrations is proposed due to the loss of workability
and mechanical properties at a high Al-content.
34
Influence of reactive elements (RE) Based on the outcome of the oxidation study, carried out in Paper I,
the second generation of alloys (Gen 2) was produced. Ti, Y and Zr
additions, of varying concentrations, were studied in Fe10Cr4Albased alloys. Their long-term oxidation properties were studied in
the same environment as the Gen 1 alloys, apart from the exposure
time that was reduced to one year (8,760 h). In addition, a 1,000 h
oxidation study was carried out in liquid lead at 450 °C.
The results revealed that the oxidation properties of the alloys were
strongly dependent of the atomic ratio between the RE-content and
the C-content. Balanced RE and C concentrations, with respect to
carbide formation, resulted in the formation of thin protective
alumina-layers at 450 °C and 550 °C, albeit the low concentrations
of Al (4 wt. %) and Cr (10 wt. %).
Alloys containing an excess of C, in relation to carbide forming
RE-additions, all displayed non-protective Fe-based oxide layer
formation, similar to the behavior of the low-preforming alloys
discussed in Paper I. In addition, the formation of µm-sized Crcarbides was consequently observed in the outermost surface layer.
It is likely that the Cr-carbides form during the oxide layer
formation, thus nucleating at the metal-oxide interface. The
simultaneous depletion of Al and presence of free C during the
oxidation process are likely the driving forces for this process. The
proposed mechanism is schematically described in Fig. 8. It is well
known that the oxidation properties of the FeCrAl alloys are
dictated by the total amount of Al and Cr in the alloy. It is thus
reasonable to assume that at low Al and Cr concentrations, i.e. a
limited reservoir of protective oxide forming elements, the long-
35
term ability of the alloy to maintain a protective oxide is
compromised by the formation of competing phases (like Crcarbides) at or adjacent to the metal-oxide interface. In this
situation, rapid oxidation of the Cr-carbides is likely to occur, a
hypothesis supported by similarities in size and shape of Crcarbides and the non-protective oxides found at the metal-oxide
interface.
RE-deficit FeCrAl alloy
mixed
oxide
Cr-carbide
Alumina with RE
C
RE-carbides
C
Al, RE
C
Phase fraction
of precipitates
Activity
Al
RE
C
Cr-carbides
RE-carbides
Depth under oxide layer
Fig. 8. Illustration of the proposed mechanism, causing Cr-carbide precipitation under the
alumina-layer in alloys having an excess of C compared to RE. Al, which suppresses
carbide formation, is depleted at the metal-oxide interface, and thus Cr-carbide formation
is thermodynamically predicted.
36
The RE/C ratios of the investigated alloys are presented in Fig. 9.
Noticeable is that all alloys that had an RE/C ratio close to one still
performed poorly in the studied environment. A likely explanation
is that the solubility of the reactive elements in the ferritic phase, as
well as the consumption of RE during oxidation, must be taken into
account. Therefore, in order to balance the RE and C to suppress
the formation of detrimental Cr-carbides, the effective RE/C ratio
in the alloy should be larger than one.
The decrease in grain size, with increasing RE addition, was
evident. A difference of about 50 µm (average grain size) between
the alloy with largest grains (Zr-0.1) and the alloy with smallest
grains (Y-0.2) was observed. Smaller grains may provide improved
oxidation properties. However, no obvious beneficial effect could
be attributed to a reduced grain size in this study.
Variations in Zr and Y on the oxidation behavior of the
investigated alloys were mainly studied within the context of this
thesis. However, Zr was shown to be more beneficial compared to
Y for the protective layer formation. A plausible explanation is
differences in their respective reactions with C. While Zr forms
MC-type carbides, Y mainly forms intermetallics together with Fe
that may include C. Thus, the concept of FeCrAl-MC is proposed
for alloys at low temperatures.
37
2.5
0.4 Zr
RE/C ratio (at. %)
2
Non−protective oxide layer
Minor oxide defects
Protective oxide layer
0.2 Zr
1.5
0.2 Y
1
0.5
0
3
4
5
6
7
Al concentration (wt. %)
8
9
Fig. 9. The RE/C ratio as a function of Al-concentration is presented for all Fe10CrAlalloys investigated in this thesis.
Microstructural stability of FeCrAl alloys (Paper II) In Paper II, the phase stability of Fe10rAl alloys was assessed by
means of a thermal aging experiment conducted up to 10,000 h in
the temperature interval of 450-550 °C. Exposed samples were
evaluated using micro-hardness measurements, electrical resistance
measurements and atom probe tomography (APT).
The results showed that while the reference Fe21Cr5Al alloy
showed clear signs of α-α’ phase separation already after 48 h at
temperatures up to 500 °C, none of the Fe10CrAl alloys were
affected. The APT evaluation of the Fe10CrAl alloys displayed
homogeneous microstructures, down to the atomic level, whereas
the reference alloy was clearly phase-separated already after 48 h,
Fig. 10.
38
a.
b.
c.
Fig. 10. a. Example of a homogeneous Cr-distribution, here shown for the Fe10Cr8Al
alloy aged at 500 °C for 10,000 h. b. Clustering of Cr-atoms is seen already after 48 h at
475 °C for the reference Fe21Cr5Al alloy. c. The said clustering is more evident when
only displaying areas having a Cr-concentration of 30% or higher (iso-concentration
surfaces). All volumes measure 50×50×75 nm3 and display the distribution of 52Cr atoms.
In addition to the experiments, thermal aging was simulated using
Kinetic Monte Carlo (KMC) modeling. A model of the FeCr
system was extended with Al-Al, Al-Fe and Al-Cr interactions,
which yielded results in good agreement with the experiments. The
kinetics of the α-α’ phase separation is dependent on the Crconcentration and the temperature. It is thus possible that the time
frame of the experiment was too short. To compensate for this, the
model was used to carry out simulations up to 10,000,000 h,
however, no signs of α-α’ phase separation were observed. Based
on the model, the Cr-limit to avoid α-α’ phase separation was
calculated to approximately 13 wt. % in a 4 wt. % Al-alloy. At
even lower temperatures, around 300 °C, neutron irradiation injects
energy to the system that is known to accelerate α-α’ phase
separation in susceptible alloys. The model was hence used to
calculate the Cr-limit in the same system, yielding 11 wt. % Cr.
These predictions indicate that the studied Fe10Cr4Al alloy system
is suitable for use in LFR applications. In the aftermath of the
39
Fukushima accident, low-Cr FeCrAl alloys have also been
proposed as a possible replacement of the conventional Zircaloy
cladding material in light-water reactors. This application demands
a phase stable composition, thus widening the market for
Fe10Cr4Al alloys beyond the LFR application.
Oxidation properties of alumina-­‐forming austenitic stainless steels (Paper III) Based on the successful work on developing an alumina-forming
stainless steel (AFA) carried out at Oak Ridge National Laboratory
in the USA, two AFA alloys were studied within the framework of
this thesis. The goal was to investigate the corrosion resistance of
these AFA alloys at a temperature (550 °C) where standard
stainless steels, like 316L, typically suffer from dissolution attack
due to its Ni-content. Above 500 °C, chromia-layers are not
sufficient to protect the steel, however, based on the work on
FeCrAl alloys, alumina-layers would provide this capacity.
Improved mechanical properties of austenitic steels compared with
ferritic steels, make the AFA alloys interesting candidate materials
for use in an LFR application. Since it was anticipated that the high
Ni-content (~20 wt. %) of the conventional AFA alloy would suffer
from Ni dissolution in liquid lead, a low Ni-AFA (~14 wt. % Ni)
was produced and investigated.
In the one-year oxidation study in liquid lead at 550 °C, both AFA
alloys displayed improved corrosion resistance compared to the
two reference steels (15-15 Ti and 316L). Whereas localized
dissolution attacks occurred on the 20Ni AFA, no attacks were
observed on the 14Ni AFA. The protective oxide layer formed on
the 14Ni-AFA was mainly enriched in Al. No further investigation
40
of the oxide layer has been carried out so far. However, it is
reasonable to believe that the oxide layer is similar to layers formed
on FeCrAl alloys in the same environment.
The main issue with AFA alloys is to maintain the austenite
structure of the alloys while operating at temperatures below about
600 °C. Cr and Al in particular stabilize the ferrite phase. Large
amounts of these elements thus increase the driving force for ferrite
formation. The high creep strength of AFA alloys is strongly
related to the austenite phase, and the formation of ferrite has
shown to reduce the creep strength. The microstructural
investigation of the studied AFA alloys showed that only the 20Ni
AFA alloy maintained a fully austenitic structure, whereas about 17
% ferrite was formed in the 14Ni AFA alloy as a result of the oneyear exposure at 550 °C, Fig. 11. Further work has to be carried out
to address the impact of the ferrite on the mechanical properties of
the alloys.
Ferrite
20 µm
Fig. 11. The microstructure of the 14Ni AFA alloy exposed at 550 °C for one year,
showing the presence of ferrite (red) in the austenite (blue) matrix.
41
42
Conclusions The work presented in this thesis highlights the superior long-term
oxidation properties of FeCrAl alloys in liquid lead at temperatures
between 450 °C and 550 °C due to the formation of thin aluminalayers. The importance of RE-additions is highlighted since the
upper compositional levels of Cr is limited by α-α’ phase
separation and for Al by workability. Moreover, it was found that
low-Cr FeCrAl alloys are more sensitive regarding variations in
reactive element additions, likely due to the minimal reservoir of
Al and Cr, compared to conventional FeCrAl alloys.
The influence of reactive elements was clearly related to the Ccontent of the alloys. Balancing the RE-content and the C-content
resulted in the formation of a protective oxide layer. The alloys
with Zr additions displayed improved oxidation properties
compared to the alloys containing Y. The fact that Zr forms MCtype carbides, whereas Y does not, is a plausible explanation for
the difference in oxidation properties. It is proposed that a low-Cr
FeCrAl alloy is alloyed to allow MC-type carbides to form, thus
minimizing the formation of Cr-carbides. Moreover, the solubility
of RE as well as consumption of RE during the protective oxide
layer formation should be taken into account. In order to counteract
the increased driving force for Cr-carbide formation under the
alumina layer, it is proposed that an alloys RE/C ratio should be
larger than unity.
The thermal aging experiments up to 10,000 h did not show any
signs of α-α’ phase separation in the 10 wt. % alloys that were
studied. The APT evaluation displayed homogeneous
microstructures. A KMC model of the Fe-Cr system was extended
43
with Al-interactions, yielding results in good agreement with the
experimental results. In addition, the model predicted the Cr-limit
to avoid α-α’ phase separation to about 13 wt. % at 475 °C and 11
wt. % Cr at 328 °C, given a 4 wt. % Al alloy. These predictions
indicate that the Fe10Cr4Al alloy system mainly studied in this
thesis is suitable for use in LFR applications.
The oxidation properties of alumina-forming austenitic stainless
steels in liquid lead at 550 °C were evaluated based of their Nicontent. It was shown that a Ni-reduction, from 20 wt. % to 14 wt.
%, resulted in the formation of protective Al-rich oxide layers, at
least up to one year of exposure. Simultaneously, the austenite
structure was partly transformed into ferrite due to the higher
concentration of ferrite stabilizing elements (Al and Cr) compared
to Ni. Ferrite formation is related to a loss of mechanical
properties, thus further studies have to be carried out on the 14Ni
AFA system. Nevertheless, due to the beneficial oxidation
properties of AFA alloys, they as well are proposed as candidate
materials for liquid lead applications.
44
Future work The aim of the work carried out in this thesis was to investigate
whether protective alumina layers are formed on low-Cr FeCrAl
alloys in liquid lead. The target temperature was 550 °C, a
temperature where conventional chromia-forming steels do not
show protective behavior. In addition, a challenging intermediate
low dissolved oxygen concentration of 10-7 wt. % was selected for
the liquid lead exposure. Set goals within the framework of the
thesis were met, however, a large body of work remains in order to
license an alumina-forming alloy for use in a nuclear application.
The author foresees the following topics as being most urgent:
Ø Investigations of the effects of other MC-carbide formers,
in particular Hf, Ta, V, Nb, and combinations thereof, on
the oxidation properties of Fe10Cr4Al alloys.
Ø Screening of the alloys in flowing conditions as well as in a
wider temperature range (400-950 °C) and dissolved
oxygen concentration range (𝑐!,!"# -10-10 wt. %).
Ø Investigations of mechanical properties in general and the
influence of liquid lead in particular.
Ø Investigations of the weldability of the Fe10Cr4Al-MC
alloys.
Ø Verification of the simulated long-term aging results, by
means of irradiation experiments at critical temperatures.
Ø Further optimization of alumina-forming austenitic stainless
steels for use in liquid lead applications. The influence of
the ferrite phase on the mechanical properties is of great
interest. 45
References [1]
[2]
[3]
[4]
[5]
[6]
[7]
[8]
[9]
[10]
[11]
[12]
[13]
[14]
46
IPCC, Climate Change 2014: Mitigation of Climate Change
(2014)
OECD, OECD Factbook 2014 (2014)
WHO, Chernobyl: the true scale of the accident (2005)
WHO, Health risk assessment from the nuclear accident after the
2011 Great East Japan Earthquake and Tsunami (2013)
P.A. Kharecha, J.E. Hansen, Environ Sci Technol, 47 (2013)
4889-4895.
B. Barre, Cr Phys, 13 (2012) 352-358.
G. Choppin, J.-O. Liljenzin, J. Rydberg, C. Ekberg, Chapter 21 The Nuclear Fuel Cycle, in: Radiochemistry and Nuclear
Chemistry (Fourth Edition), Academic Press, Oxford (2013), pp.
685-751.
R.L. Murray, K.E. Holbert, Chapter 25 - Breeder reactors, in:
Nuclear Energy (Seventh Edition), Butterworth-Heinemann,
Boston, 2015, pp. 459-475.
Generation IV Forum (GIF), A Technology Roadmap for
Generation IV Nuclear Energy Systems (2002)
B.F. Gromov, Y.S. Belomitcev, E.I. Yefimov, M.P. Leonchuk,
P.N. Martinov, Y.I. Orlov, D.V. Pankratov, Y.G. Pashkin, G.I.
Toshinsky, V.V. Chekunov, B.A. Shmatko, Nuclear Engineering
and Design, 173 (1997) 207-217.
L. Cinotti, C. Smith, C. Artioli, G. Grasso, G. Corsini, LeadCooled Fast Reactor (LFR) Design: Safety, Neutronics, Thermal
Hydraulics, Structural Mechanics, Fuel, Core, and Plant Design,
in: Handbook of Nuclear Engineering, Springer US (2010), pp.
2749-2840
IAEA, Liquid Metal Coolants for Fast Reactors Cooled by
Sodium, Lead and Lead-Bismuth Eutectic (2012)
H.A. Abderrahim, P. Baeten, D. De Bruyn, R. Fernandez, Energ
Convers Manage, 63 (2012) 4-10.
A. Alemberti, M. Frogheri, L. Mansani, The lead fast reactor:
Demonstrator (ALFRED) and ELFR design, Proceedings of the
International Conference on fast reactors and related fuel cycles FR13, Paris, France (2013)
[15]
[33]
A.V. Zrodnikov, G.I. Toshinsky, O.G. Komlev, V.S. Stepanov,
N.N. Klimov, J. Nucl. Mater., 415 (2011) 237-244.
A. Alemberti, Ansaldo Nucleare, Private communication (2014)
A. Alemberti, V. Smirnov, C.F. Smith, M. Takahashi, Prog Nucl
Energ, 77 (2014) 300-307.
E. Adamov, V. Orlov, A. Filin, V. Leonov, A. SilaNovitski, V.
Smirnov, V. Tsikunov, Nuclear Engineering and Design, 173
(1997) 143-150.
J. Wallenius, E. Suvdantsetseg, A. Fokau, Nucl Technol, 177
(2012) 303-313.
J. Neuhausen, B. Eichler, Radiochim Acta, 94 (2006) 239-242.
J.S. Zhang, Corros. Sci., 51 (2009) 1207-1227.
OECD, Handbook on lead-bismuth eutectic alloy and lead
properties, materials compatibility, thermal-hydraulics and
technologies, OECD Nuclear Energy Agency (2007)
R.C. Asher, D. Davies, S.A. Beetham, Corros. Sci., 17 (1977)
545-557.
G. Muller, A. Heinzel, J. Konys, G. Schumacher, A.
Weisenburger, F. Zimmermann, V. Engelko, A. Rusanov, V.
Markov, J. Nucl. Mater., 335 (2004) 163-168.
A. Heinzel, A. Weisenburger, G. Muller, J. Nucl. Mater., 448
(2014) 163-171.
M. Kondo, M. Takahashi, J. Nucl. Mater., 356 (2006) 203-212.
P. Hosemann, H.T. Thau, A.L. Johnson, S.A. Maloy, N. Li, J.
Nucl. Mater., 373 (2008) 246-253.
J. Lim, H.O. Nam, I.S. Hwang, J.H. Kim, J. Nucl. Mater., 407
(2010) 205-210.
S. Takaya, T. Furukawa, M. Inoue, T. Fujisawa, T. Okuda, F.
Abe, S. Ohnuki, A. Kimura, J. Nucl. Mater., 398 (2010) 132-138.
M. Del Giacco, A. Weisenburger, A. Jianu, F. Lang, G. Mueller,
J. Nucl. Mater., 421 (2012) 39-46.
J. Lim, I.S. Hwang, J.H. Kim, J. Nucl. Mater., 441 (2013) 650660.
A. Weisenburger, A. Jianu, S. Doyle, M. Bruns, R. Fetzer, A.
Heinzel, M. DelGiacco, W. An, G. Muller, J. Nucl. Mater., 437
(2013) 282-292.
R. Prescott, M.J. Graham, Oxid. Met., 38 (1992) 233-254.
47
[16]
[17]
[18]
[19]
[20]
[21]
[22]
[23]
[24]
[25]
[26]
[27]
[28]
[29]
[30]
[31]
[32]
[34]
[35]
[36]
[37]
[38]
[39]
[40]
[41]
[42]
[43]
[44]
[45]
[46]
[47]
[48]
[49]
[50]
[51]
[52]
[53]
[54]
48
Y. Yamamoto, M.P. Brady, Z.P. Lu, P.J. Maziasz, C.T. Liu, B.A.
Pint, K.L. More, H.M. Meyer, E.A. Payzant, Science, 316 (2007)
433-436.
M.P. Brady, Y. Yamamoto, M.L. Santella, B.A. Pint, Scripta
Mater, 57 (2007) 1117-1120.
M.P. Brady, Y. Yamamoto, M.L. Santella, P.J. Maziasz, B.A.
Pint, C.T. Liu, Z.P. Lu, H. Bei, Jom-Us, 60 (2008) 12-18.
Y. Yamamoto, A. Takeyama, Z.P. Lu, C.T. Liu, N.D. Evans, P.J.
Maziasz, M.P. Brady, Intermetallics, 16 (2008) 453-462.
M.P. Brady, Y. Yamamoto, M.L. Santella, L.R. Walker, Oxid.
Met., 72 (2009) 311-333.
Y. Yamamoto, M.L. Santella, C.T. Liu, N.D. Evans, P.J.
Maziasz, M.P. Brady, Mat Sci Eng a-Struct, 524 (2009) 176-185.
M.P. Brady, J. Magee, Y. Yamamoto, D. Helmick, L. Wang, Mat
Sci Eng a-Struct, 590 (2014) 101-115.
H. von Kantzow, Fire-resistant alloy with high electric resistance,
US patent no. 1717284 (1929)
C. Wagner, Corros. Sci., 5 (1965) 751-764.
F.H. Stott, G.C. Wood, J. Stringer, Oxid. Met., 44 (1995) 113145.
Z.G. Zhang, F. Gesmundo, P.Y. Hou, Y. Niu, Corros. Sci., 48
(2006) 741-765.
http://www.kanthal.com, 2014-11-10.
J. Stringer, Mat Sci Eng a-Struct, 120 (1989) 129-137.
J. Jedlinski, Corros. Sci., 35 (1993) 863-869.
B.A. Pint, Oxid. Met., 45 (1996) 1-37.
J.L. Bartos, The nature of aging of Fe-Cr-Al-Y alloys at 450 C,
GEMP 708 (1969).
C. Capdevila, M.K. Miller, K.F. Russell, J. Chao, J.L. GonzalezCarrasco, Mat Sci Eng a-Struct, 490 (2008) 277-288.
S. Kobayashi, T. Takasugi, Scripta Mater, 63 (2010) 1104-1107.
T. Cheng, J.R. Keiser, M.P. Brady, K.A. Terrani, B.A. Pint, J.
Nucl. Mater., 427 (2012) 396-400.
B.A. Pint, K.A. Terrani, M.P. Brady, T. Cheng, J.R. Keiser, J.
Nucl. Mater., 440 (2013) 420-427.
K.A. Terrani, S.J. Zinkle, L.L. Snead, J. Nucl. Mater., 448 (2014)
420-435.
[55]
[56]
[57]
[58]
[59]
[60]
[61]
[62]
[63]
[64]
[65]
[66]
[67]
[68]
[69]
[70]
[71]
J.C. Pivin, D. Delaunay, C. Roquescarmes, A.M. Huntz, P.
Lacombe, Corros. Sci., 20 (1980) 351-373.
J.A. Mcgurty, R. Nekkanti, J. Moteff, Jom-J Min Met Mat S, 38
(1986) 22-25.
V. Ramakrishnan, J.A. Mcgurty, N. Jayaraman, Oxid. Met., 30
(1988) 185-200.
D.V.V. Satyanarayana, G. Malakondaiah, D.S. Sarma, Mat Sci
Eng a-Struct, 323 (2002) 119-128.
P. Kofstad, High temperature corrosion, ed., Elsevier Applied
Science, 1988.
A.S. Khanna, Chapter 6 - High temperature oxidation, in:
Handbook of Environmental Degradation of Materials, William
Andrew Publishing, Norwich, NY (2005) pp. 105-152.
N. Birks, G.H. Meier, F.S. Pettit, Introduction to the hightemperature oxidation of metals, ed., Cambridge University
Press, 2006.
G.Y. Lai, High-Temperature Corrosion and Materials
Applications, ed., ASM International (2007).
H.J.T. Ellingham, Journal of the Society of Chemical Industry, 5
(1944)
L.B. Pheil, W.T. Griffiths, Heat resistant alloy, US patent no.
2304353 (1942).
K.A. Darling, R.N. Chan, P.Z. Wong, J.E. Semones, R.O.
Scattergood, C.C. Koch, Scripta Mater, 59 (2008) 530-533.
K.A. Darling, B.K. VanLeeuwen, J.E. Semones, C.C. Koch, R.O.
Scattergood, L.J. Kecskes, S.N. Mathaudhu, Mat Sci Eng aStruct, 528 (2011) 4365-4371.
H.Y. Lou, F.H. Wang, B.J. Xia, L.X. Zhang, Oxid. Met., 38
(1992) 299-307.
R. Prescott, M.J. Graham, Oxid. Met., 38 (1992) 73-87.
F. Liu, H. Josefsson, J.E. Svensson, L.G. Johansson, M.
Halvarsson, Mater High Temp, 22 (2005) 521-526.
H. Gotlind, F. Liu, J.E. Svensson, M. Halvarsson, L.G.
Johansson, Oxid. Met., 67 (2007) 251-266.
E. Airiskallio, E. Nurmi, M.H. Heinonen, I.J. Vayrynen, K.
Kokko, M. Ropo, M.P.J. Punkkinen, H. Pitkanen, M. Alatalo, J.
Kollar, B. Johansson, L. Vitos, Corros. Sci., 52 (2010) 33943404.
49
[72]
[73]
[74]
[75]
[76]
[77]
[78]
[79]
[80]
[81]
[82]
[83]
[84]
[85]
[86]
[87]
[88]
[89]
[90]
50
G. Muller, G. Schumacher, F. Zimmerman, J. Nucl. Mater., 278
(2000) 85-95.
G. Muller, A. Heinzel, G. Schumacher, A. Weisenburger, J. Nucl.
Mater., 321 (2003) 256-262.
C. Schroer, J. Konys, Physical Chemistry of Corrosion and
Oxygen Control in Liquid Lead and Lead-Bismuth Eutectic
(2007).
J.S. Zhang, J Appl Electrochem, 43 (2013) 755-771.
C.B. Alcock, T.N. Belford, Transactions of the Faraday Society,
60 (1964) 822-835.
D. Risold, J.I. Nagata, R.O. Suzuki, J Phase Equilib, 19 (1998)
213-233.
J.O. Andersson, B. Sundman, Calphad, 11 (1987) 83-92.
G. Bonny, D. Terentyev, L. Malerba, Scripta Mater, 59 (2008)
1193-1196.
R.O. Williams, H.W. Paxton, J Iron Steel Inst, 185 (1957) 358374.
S. Hertzman, B. Sundman, Calphad, 6 (1982) 67-80.
D.A. Porter, K.E. Easterling, Phase transformations in metals and
alloy, 2nd ed., Nelson Thornes Ltd, 1992.
F.M. Becket, Transaction of the American Institute of Mining
and Metallurgical Engineers, 131 (1938)
M. Hillert, A theory of nucleation for solid metallic solutions,
Ph.D. thesis, Dept. of Metallurgy, Massachusetts Institute of
Technology (1956).
J.W. Cahn, J.E. Hilliard, The Journal of Chemical Physics, 28
(1958) 258-267.
M.H. Mathon, Y. de Carlan, G. Geoffroy, X. Averty, A. Alamo,
C.H. de Novion, J. Nucl. Mater., 312 (2003) 236-248.
V. Kuksenko, C. Pareige, P. Pareige, J. Nucl. Mater., 432 (2013)
160-165.
M. Bachhav, G.R. Odette, E.A. Marquis, Scripta Mater, 74
(2014) 48-51.
P. Olsson, J. Wallenius, C. Domain, K. Nordlund, L. Malerba,
Phys Rev B, 72 (2005)
W. Xiong, M. Selleby, Q. Chen, J. Odqvist, Y. Du, Crit Rev
Solid State, 35 (2010) 125-152.
[91]
[95]
[96]
[97]
[98]
W. Xiong, K.A. Gronhagen, J. Agren, M. Selleby, J. Odqvist, Q.
Chen, Solid-Solid Phase Transformations in Inorganic Materials,
Pts 1-2, 172-174 (2011) 1060-1065.
W. Li, S. Lu, Q.-M. Hu, H. Mao, B. Johansson, L. Vitos, Comp
Mater Sci, 74 (2013) 101-106.
B. Sundman, B. Jansson, J.O. Andersson, Calphad, 9 (1985) 153190.
J.O. Andersson, T. Helander, L.H. Hoglund, P.F. Shi, B.
Sundman, Calphad, 26 (2002) 273-312.
L.A. Giannuzzi, F.A. Stevie, Micron, 30 (1999) 197-204.
A.J. Wilkinson, P.B. Hirsch, Micron, 28 (1997) 279-308.
S. Nishikawa, S. Kikuchi, Nature, 121 (1928) 1019-1020.
M.K. Miller, R.G. Forbes, Mater Charact, 60 (2009) 461-469.
51
[92]
[93]
[94]