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ISBN print: 978-3-0357-1402-9
ISBN eBook: 978-3-0357-3402-7
DOI book: 10.4028/b-V7RsD7
© 2020 Trans Tech Publications Ltd, All Rights Reserved

High Temperature Corrosion of Stainless Steels:
An Alain Galerie Festschrift

Edited by
Somrerk Chandra-ambhorn
Yves Wouters
Wanlop Harnnarongchai

High Temperature Corrosion of Stainless Steels: An Alain Galerie Festschrift

Selected, peer reviewed papers from the
Alain Galerie Symposium on
High Temperature Corrosion,
9 July, 2018, Bangkok, Thailand

Edited by

Somrerk Chandra-ambhorn, Yves Wouters and Wanlop Harnnarongchai

Copyright © 2020 Trans Tech Publications Ltd, Switzerland

All rights reserved. No part of the contents of this publication may be reproduced or transmitted in any form or by any means without the written permission of the publisher.

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Solid State Phenomena, ISSN: 1662-9779, Vol. 300, pp -2-0

© 2020 Trans Tech Publications Ltd, All Rights Reserved

Accepted:
2019-10-28

Online:
2020-02-12

Alain Galerie: A Biography

The original version of this paper is available on https://www.scientific.net/SSP.300.-2.pdf

Shigenari Hayashi

Hokkaido University, Japan

Foreword

The original version of this paper is available on https://www.scientific.net/SSP.300.-2.pdf

In industries, there are many processes operated at elevated temperature under severely corrosive atmospheres. High temperature corrosion (HTC) may degrade the performance and the life of metals and alloys which are involved in the processes. The protection of the materials from HTC is a key issue.

High temperature oxidation (HTO) is the most frequent phenomena in HTC. Metallic alloys are exposed to the gaseous oxidants ( O2,H2O,CO2,SO2,NO2 etc.) at elevated temperatures. Metallic constituents are oxidised and form the oxide scale (oxide layer) on the surface. In the case of the oxide scale being dense enough to separate physically the metal and the oxidant, the oxidation process is rate determined by diffusion of ions through the layer and the oxidation rate is slowed down dramatically. Prior to the scale formation, less noble constituents are oxidised by dissolved oxygen in the alloy and oxide precipitates form internally (internal oxidation). The increase in the number of precipitates and their coarsening result in the combination of precipitates and the formation of continuous oxide scale (external oxidation). During mass transport process of oxidation, complex microstructure developments such as oxide grain growth and void (pore) formation strongly affect the stress generation and exfoliation of oxide scales. Furthermore, the grain boundaries in the oxide scale may serve short circuits of ionic diffusion.

High temperature oxidation seems to be so complicated for engineers to understand and protect materials from HTO. The clarification of HTO requires more researches and basic knowledge of thermodynamics and kinetics.

Professor Alain Galerie is one of the most recognised figures in the French and global HTC community. He started his work on HTC from 1970 and he has continued in five decades. His style of research is to focus on elemental phenomena which govern the chemical and mechanical features in HTC by using microscopic analytical techniques, which helps largely for understanding the mechanism of HTO/HTC. He has been invited to give lectures from major scientific congresses such as Gordon Research Conference on HTC, International Symposium on HT Corrosion and Protection of Materials (the Embiez Congress) and so on. He is also author or co-author of several books. There are excellent professors and engineers in HTC community not only in France but also in the world, whom he educated as a teacher and a supervisor.

This book is published by the contribution of Professor Alain Galerie's former students, the colleagues in his research group and leading researchers in the field. This book offers briefly the basic knowledge and the recent understanding on HTC of stainless steels so that it is beneficial to graduate students and young researchers in HTC community.

Toshio Maruyama
Professor Emeritus,
Tokyo Institute of Technology, Japan

Preface

The original version of this paper is available on https://www.scientific.net/SSP.300.-2.pdf

This book provides the basic concepts and the recent research on the high temperature corrosion of stainless steels, starting from the introduction to stainless steels in the early part of Chapter 1. Thermodynamics and kinetics of high temperature corrosion phenomena, as well as mechanical behaviour of thermal oxide formed on the metal are treated in Chapters 1 and 2. Chapter 3 explains tools to characterise the thermal oxide scale. Applications on stainless steel oxidation are given in these three chapters. The effects of alloying elements and the atmosphere emphasising on water vapour on stainless steel oxidation are further addressed in Chapters 4 and 5. The last chapter devotes to the specific application of stainless steel as an interconnect in solid oxide fuel cells (SOFCs).

Because this festschrift is published in honour of Professor Alain Galerie, selection of the presented topics and recent research therefore reflect his style and the works conducted by his group and his alliance. For example, the treatment of the high temperature corrosion phenomena by using defect reactions is favoured here following his style. The experimental set-ups developed in his group are highlighted particularly the photoelectrochemical technique, the set-up to measure chromium species volatilisation and the mechanical adhesion energy measurement. His interest in the oxidation in water vapour and the development of stainless steel SOFC interconnect are also emphasised.

This book stems from the "Alain Galerie Symposium on High Temperature Corrosion" which was held at King Mongkut's University of Technology North Bangkok (KMUTNB), Bangkok, Thailand, on 9 July 2018. Fortunately, the book was finished in the period of the celebration of the 30-year anniversary of the KMUTNB Thai-French Innovation Institute established since 1990. The editors appreciate Professor Toshio Maruyama for his foreword. They sincerely thank the co-authors and anonymous reviewers for the contribution to this book. The acknowledgement is given to the assistants to the editors: Ms. Thammaporn Thublaor and Mr. Panya Wiman.

The editors hope that this book will be useful for the high temperature corrosion community. Further, they do hope that this festschrift will be an academic monument to show the contribution of Professor Alain Galerie to this field, and to express our love and respect to him.

The editors

Somrerk Chandra-ambhorn
Professor, KMUTNB, Thailand

Yves Wouters
Professor,
Université Grenoble Alpes, France
Wanlop Harnnarongchai
Assistant Professor,
KMUTNB, Thailand

Yves Wouters1,a*and Somrerk Chandra-ambhorn2,b

The original version of this paper is available on https://www.scientific.net/SSP.300.-2.pdf

1 Université Grenoble Alpes, SIMaP, 1130 rue de la Piscine, BP 75, 38402 St Martin d'Hères, France
2 High Temperature Corrosion Research Centre, Department of Materials and Production Technology Engineering, Faculty of Engineering, King Mongkut's University of Technology North Bangkok, 1518, Pracharat 1 Road, Bangsue, Bangkok, 10800, Thailand
a  yves.wouters@univ-grenoble-alpes.fr, b  somrerk.c@eng.kmutnb.ac.th

Professor Alain Galerie is, to this day, one of the most recognised figures in the French and global high temperature corrosion community. His entire career demonstrates his great attachment to his job as a teacher-researcher.

Since 1970, he has been studying the theoretical study of the kinetics of reactions between composite gases and metals, including sulphurisation, developed in his thesis in 1977. Thereafter, and until now, he has continued to tackle new scientific projects often characterised by innovative research and with multiple outcomes. During several postdoctoral research stays at the Harwell Nuclear Center (UK), he got passion with ion implantation and its effects on oxidation, then on the development of making solidified surface alloys under a laser beam. The research on ion implantation into metals led him win the "Prix Lucien Chatain" from the French South-East Scientific Foundation. In the 90s, when few scientists saw the interest, he began, with Marcel Caillet, DR CNRS, a vast program of study of the effects of water vapour on the oxidation at high temperature of pure metals. It is only a decade later that the community will become aware of the inevitable interest of this subject, water vapor being revealed, in many industrial environments, as an extremely aggressive agent. Thanks to this expertise, he was invited to give two plenary lectures to the Gordon Research Conferences in 2001 and 2009. He was also interested in the contribution of photoelectrochemistry, technique which has proven to be very successful in bringing innovative information thanks to the semiconducting signature of oxide scales. Later, he developed a new major axis around the adhesion of thermal oxides to refractory alloys by implementing tools for determining the interfacial rupture energy of the protective layers (blister test and SEM in-situ tensile testing). He also studied the monitoring of damage during high temperature oxidation by acoustic emission. In addition, he was involved in a perpetual action of newness, for example with the use of Raman imaging and FIB/SEM in several studies. All of this work was done with 30 PhD students. As the industrial development is growing up in Asia, he and his group have supervised 5 PhD students from Thailand who are now actively working in research in their region, as well as a number of interns from Thailand.

As a result of the research works, he produced more than 180 papers in scientific journals and proceedings ( 163 articles referenced "Web of Science", h-index =23 ) and more than 50 invited lectures. He is regularly invited as a member of thesis and HDR (habilitation thesis to supervise
research) juries. He involved, nationally as well as internationally, in the organisation of thematic schools and major scientific congresses - the most famous being probably the International Symposium on High-Temperature Corrosion and Protection of Materials (better known as name of the Embiez Congress). Although now retired, Professor Alain Galerie is still very active. He is regularly invited in the major congresses of the high-temperature oxidation community in which he gives very popular scientific historical lectures.

Professor Alain Galerie is also author or co-author of several books, the most notable of which is "Les mécanismes de la corrosion sèche" with his colleagues Pierre Sarrazin and Jacques Fouletier, published in French in 2000. This book won the "Roberval Prize" in 2001 and was translated into English version (Mechanisms of High Temperature Corrosion: A Kinetic Approach) in 2008. He also wrote the eight papers published in the "Techniques of the Engineer" with Laurent Antoni. He won the "Ugine-René Castro Prize" from the French Society for Metallurgy and Materials in 2005, and the "Grande Médaille" from the French Centre of Anticorrosion in 2010.

He spent most of his career at ENSEEG (National School of Electrochemistry and Electrometallurgy of the Polytechnic National Institute of Grenoble), where he was for many years the Director of Studies. He has also held, at the national level, the position of President of the "P" and "PSI" Concours Communs Polytechniques. His teaching has always been, for him, a priority and he constantly privileged the close relationship with students, advocating for the strengthening of practical work and laboratory projects. In addition to mineral chemistry, he has taught in particular heterogeneous kinetics, high temperature corrosion and surface treatments. The school PHELMA / Grenoble INP, today formed by the meeting of three Grenoble schools, owes a lot to him in his new organisation. Having trained many engineers and PhD students whom he likes to meet often in the research and development centres of the big French industrial groups, Professor Alain Galerie is known for his remarkable pedagogy and his great attachment to the training of the students, qualities which have undoubtedly helped to mark several generations of young researchers.

Research Group

Walairat Chandra-ambhorn and Patthranit Wongpromrat
King Mongkut's Institute of Technology Ladkrabang, Thailand

Somrerk Chandra-ambhorn, Wanlop Harnnarongchai, Thanasak Nilsonthi, Piyorose Promdirek, Anusara Srisrual and Thammaporn ThublaorKing Mongkut's University of Technology North Bangkok, ThailandSébastien Chevalier, Lionel Combemale and Ioana PopaUniversité de Bourgogne Franche-Comté, FranceMuriel Braccini, Laurence Latu-Romain, Valérie Parry, Céline Pascal and Yves Wouters Université Grenoble Alpes, France

Group Pictures at Alain Galerie Symposium on High Temperature Corrosion, King Mongkut's University of Technology North Bangkok, Bangkok, Thailand, 9 July 2018.

The original version of this paper is available on https://www.scientific.net/SSP.300.-2.pdf

From left to right: Somrerk Chandra-ambhorn, Thanapat Kaewmaneekul, Toshio Maruyama, Shigenari Hayashi, Yves Wouters, Suchart Siengchin, Alain Galerie, Vincent Drapeau, Laurance Latu-Romain, Sébastien Chevalier, Valérie Parry, Thammaporn Thublaor and Wattana Kaewmanee

From left to right: Supat Ieamsupapong, Somrerk Chandra-ambhorn, Thanasak Nilsonthi, Walairat Chandra-ambhorn, Piyorose Promdirek, Alain Galerie, Yves Wouters, Patthranit Wongpromrat, Valérie Parry, Laurence Latu-Romain and Anusara Srisrual

Table of Contents

Preface

Thermodynamics and Kinetics of the High Temperature Oxidation of Stainless Steels

Somrerk Chandra-ambhorn, Thammaporn Thublaor, Celine Pascal

CHAPTER 2
Mechanical Behaviour of Thermal Oxide Scales on Stainless Steels

Valérie Parry, Somrerk Chandra-ambhorn, Thanasak Nilsonthi, Muriel Braccini

Characterisation of Thermal Oxide Scales on Stainless Steels

Anusara Srisrual, Thammaporn Thublaor, Patthranit Wongpromrat

CHAPTER 4
High Temperature Oxidation of Stainless Steels

Somrerk Chandra-ambhorn, Shigenari Hayashi, Laurence Latu-Romain, Patthranit Wongpromrat

Effect of Water Vapour on the High Temperature Oxidation of Stainless Steels

Somrerk Chandra-ambhorn, Patthranit Wongpromrat, Thammaporn Thublaor, Walairat Chandra-Ambhorn

Sébastien Chevalier 1,a*, Lionel Combemale 1, b, Ioana Popa 1,c, Somrerk Chandra-ambhorn 2, d, Walairat Chandra-ambhorn 3,e, Piyorose Promdirek 2,f and Patthranit Wongpromrat 3, g

Sébastien Chevalier, Lionel Combemale, Ioana Popa, Somrerk Chandra-ambhorn, Walairat Chandra-Ambhorn, Piyorose Promdirek, Patthranit Wongpromrat

Keywords Index

Authors Index

The original version of this eBook is available on https://www.scientific.net/book/high-temperature-corrosion-of-stainless-steels-an-alain-galerie-festschrift/978-3-0357-3402-7

Solid State Phenomena, ISSN: 1662-9779, Vol. 300, pp 1-24

doi: 10.4028/www.scientific.net/SSP.300.1

© 2020 The Author(s). Published by Trans Tech Publications Ltd, Switzerland.

Submitted:
2019-02-15

Revised:
2019-07-24

Accepted:
2019-08-08

Online:
2020-02-12

Thermodynamics and Kinetics of the High Temperature Oxidation of Stainless Steels

The original version of this paper is available on https://www.scientific.net/SSP.300.1.pdf

Somrerk Chandra-ambhorn 1,a*, Thammaporn Thublaor 1, b and Céline Pascal 2,c

1 High Temperature Corrosion Research Centre, Department of Materials and Production Technology Engineering, Faculty of Engineering, King Mongkut's University of Technology North Bangkok, 1518, Pracharat 1 Road, Bangsue, Bangkok, 10800, Thailand

2 Université Grenoble Alpes, CNRS, Grenoble INP, SIMaP, 1130 rue de la Piscine, BP 75,38402

St Martin d'Hères, France

a  somrerk.c@eng.kmutnb.ac.th, b  thammaporn.tt@gmail.com, c  Celine.Pascal@simap.grenoble-inp.fr

Keywords: high temperature oxidation, stainless steels, alloying elements, thermodynamics, kinetics

Abstract

This chapter introduces stainless steels and their classification for the high temperature applications. The enabling theories for the high temperature corrosion i.e. thermodynamics and kinetics are further addressed. The basic concept of thermodynamics is given and the stability of the formation of thermal oxide on stainless steel is exemplified. Types of defect in the oxide and Fick's first law for the diffusion of defect though the oxide are introduced. Oxidation kinetics is explained with the emphasis on the derivation of the parabolic rate law.

CHAPTER 1

The original version of this paper is available on https://www.scientific.net/SSP.300.1.pdf

1.1 Introduction to Stainless Steels

As noted by A. Galerie [1], the first idea of producing stainless steel might trace back to the year 1882 when the French metallurgist P. Berthier recorded that the sufficient addition of chromium to iron could make the steel better resistant to the corrosion in water and acid. However, such steel with good enough mechanical stability had not been realised until H. Brearley from Sheffield produced the first stainless steel containing 13wt.% of chromium and 0.24wt.% of carbon in 1913 [1]. The commonly used AISI 304 stainless steel was produced by W.H. Hatfield in 1924 [1].

Nowadays stainless steel is defined as a ferrous alloy containing chromium with a minimum content of about 10.5−12wt.% [ 2,3 ], while it is suggested that the chromium content to combat the high temperature corrosion should be at least 18wt.% [4,5] if other elements beneficial to the use in this condition are not additionally alloyed. The classification of stainless steels is normally based on its structure which closely relates to the steel mechanical properties e.g. ferritic, austenitic, martensitic, austeno-ferritic (duplex) or precipitation-hardened stainless steels. Some elements are added to the steel and have the effect on stabilising the ferrite matrix e.g. chromium, molybdenum, silicon and niobium [4], while some elements stabilise the austenite matrix e.g. nickel, manganese, nitrogen and carbon [4, 6]. For example, the AISI 304 grade which has the austenite structure at room temperature is obtained by the addition of 9.5wt.% of nickel to the ferritic Fe-19Cr (unless mentioned otherwise, the chemical composition is in wt.%) [1, 7].

However, in the high temperature corrosion aspect, stainless steels are typically classified by the type of the thermal oxide scale formed at high temperatures i.e. chromia- or alumina-forming ones. Table 1.1 lists chemical compositions of some important chromia- and alumina-forming stainless steels. It is noted that, for titanium and niobium contents which are reported as a range, the lower limit is in fact a function of carbon and nitrogen contents in the steel [7], but in this table it is calculated by using the maximum content of carbon and neglecting the nitrogen content.

Table 1.1 Chemical compositions (wt.%) of stainless steels.

The maximum service temperature of the steel depends on many factors especially the composition of the hot gas and the loading condition at high temperatures. However, a primary guidance for the maximum use temperature of some stainless steels for isothermal oxidation is shown in Table 1.2. For each steel grade in this table, only important elements especially the ones that help increase the high temperature corrosion resistance are listed with the contents averaged from the values reported in Table 1.1.

Table 1.2 Maximum isothermal service temperatures ( Tmax ) of stainless steels [1, 7, 13].

Tmax (°C) | Stainless steel grade
Chromia former | Alumina former
| Ferritic grade | Austenitic grade |
750 | 430 (Fe–17Cr) |  |
800 | 409 (Fe–11Cr–0.42Ti) |  |
815 | 430Ti (Fe–17Cr–0.45Ti) |  |
850 |  | 304 (Fe–19Cr–9.5Ni) 316 (Fe–17Cr–12Ni–2.5Mo) |
950 | 441 (Fe–18Cr–0.35Ti–0.7Nb) 444 (Fe–18.5Cr–0.53Ti–2.15Mo) |  |
1000 |  | 309 (Fe–23Cr–13.5Ni–1Si) |
1150 | 446 (Fe–24.5Cr) | 310 (Fe–25Cr–20.5Ni–1.5Si) |

1200

Aluchrom YHf
(Fe–20.3Cr–6.1Al–
0.06Y–0.05Zr–
0.03Hf)

1300 |  |  | MA956 (Fe–19.2Cr–4.3Al– 0.36Ti–0.63 Y₂O₃)

For the chromia-forming stainless steels, we might consider that the basic grade for the high temperature application is an AISI 430 ferritic stainless steel which consists of about 16−18wt% of chromiun and has the maximum isothermal use temperature of 750∘C. The addition of some alloying elements could increase the high temperature corrosion resistance and therefore the maximum use temperature. Mechanisms of the alloying elements to help reduce the stainless steel oxidation rate are explained in Chapter 4. However, the empirical observation of this effect especially from Tables 1.1 and 1.2 may be noted here as follows.

The first group of the elements that can help increase the high temperature corrosion resistance consists of chromium, aluminium and silicon. These elements could react with oxygen and form the protective oxide layer [1, 4, 5] or internal precipitates [14] that make the steel resistant to the high temperature oxidation. For example, when comparing with AISI 430 (Fe-17Cr), the addition of chromium up to 23-26 wt.% in the steel gives AISI 446 which has the maximum isothermal use temperature up to 1150∘C. The increase in chromium content from 18-20 wt.% in AISI 304 to 22-24 wt.% in AISI 309 and the increase in silicon content from 0.75wt.% in AISI 304 to 1wt% in AISI 309 significantly increases in the maximum isothermal use temperature from 850∘C for AISI 304 to 1000∘C for AISI 309. The addition of aluminium by 1.2wt% to Fe−12Cr, giving 12SR stainless steel, could form the internal alumina precipitate at the interface between the chromia scale and steel substrate and therefore allows the maximum isothermal use temperature up to 900∘C [1, 7]. It is noted that 12SR and 18SR are chromia-forming stainless steels when they are oxidised
in air [1]. However, in a more reducing atmosphere chromia may become unstable and in this situation these alloys could turn to be the alumina-formers [1].

The second group of the elements beneficial to the high temperature corrosion is titanium and niobium. These elements are primarily added to the steel to avoid chromium carbide formation thus mitigating the sensitisation problem of the steel [4, 8]. However, as will be explained in Chapter 4, the addition of these elements could also promote the high temperature corrosion resistance. For example, the addition of 0.18−0.65wt.% of titanium to Fe−11.5Cr makes AISI 409 which has the maximum isothermal use temperature of 800∘C. This temperature is higher than the one of AISI 430 which contains higher chromium content (16-18 wt.%) but is not stabilised by titanium. The addition of 0.45wt% of titanium to AISI 430 gives the 430 Ti grade which has the higher maximum isothermal use temperature up to 815∘C[1,13]. The slight increase in chromium from 16-18 wt.% in AISI 430 to 17.5-18.5 wt.% in the 441 grade with the stabilisation by titanium and niobium increases the maximum isothermal use temperature of the steel up to 950∘C.

For the alumina-forming stainless steel, aluminium must be sufficiently added to the Fe−Cr so that the stable continuous alumina layer can be formed. Once the alumina forms the oxidation rates of the steels tend to be improved compared to those of the chromia-forming stainless steels. This is because alumina has very low deviation from stoichiometry, i.e. this deviation is about 10−3 for chromia at 600∘C while it is about 10−4 for alumina even at 1000∘C [13]. Furthermore, the cation self-diffusion coefficient at 1000∘C in alumina is about 10−15 cm2 s−1 while that of the chromia is higher i.e. 10−12 cm2 s−1 at the same temperature [13]. From Tables 1.1 and 1.2 it can be seen that by adding aluminium to form the alumina protective layer together with the addition of titanium and/or reactive elements e.g. yttrium, zirconium, hafnium and cerium, the alumina-forming stainless steels can be used at temperatures higher than 1000∘C. For example, the addition of 6.1wt% aluminium with small amounts of yttrium, zirconium and hafnium to Fe−20.3Cr gives Aluchrome YHf which has the maximum isothermal use temperature of 1200∘C. The addition of 4.3wt.% of aluminium, 0.36wt.% of titanium and 0.63wt.% of Y2O3 to the Fe−19.2Al gives the oxide-dispersion strengthened (ODS) MA956 stainless steel which can be used isothermally at temperatures up to 1300∘C.

From this introduction, it can be seen that we need to know that when stainless steels with different chemical compositions are used at high temperatures under atmospheres with different gas compositions, which oxide phases are stable and could be formed. Furthermore, with these oxides, what the reaction rate and mechanism are. To answer these questions, we need thermodynamics to understand stability of the phases formed and we need kinetics to identify the rate and the oxidation mechanisms. The basics of these subjects will be given in the following sections.

1.2 Thermodynamics of the High Temperature Corrosion

1.2.1 Basic Concepts

The prediction of oxide formation requires to determine the most stable states of the reacting system. At constant temperature and pressure, the stability of a system is measured by its Gibbs free energy. The total Gibbs free energy, G, of a system is defined as

G=H−TS=U+PV−TS(1.1)

where H is the enthalpy, S the entropy, U the internal energy, V the volume of the system, P the pressure and T the temperature.

The Gibbs free energy change, dG, for a system in which change of composition takes place through chemical reaction, is given by

dG=(∂G∂T)P,ni dT+(∂G∂P)T,ni dP+∑i(∂G∂ni)P, T,nj≠i dni=V dP−S dT+∑iμi dni(1.2)

where ni is the mole number of component i and μi=(∂G/∂ni)P,T,nj≠i the chemical potential.
The general expression of chemical potential of the component i is given by

μi=μi∘+RTlnai(1.3)

where μi∘ is the standard chemical potential and ai the activity of compound i .
Change in mole number of component i can be expressed from the extent of reaction (ξ) and the stoichiometric coefficient (vi), which are negative for reactants and positive for products of the reaction, i.e.

dni=vi· dξ(1.4)

The combination of Eqs. 1.2 and 1.4 yields

dG=V dP−S dT+∑iμivi dξ=(∂G∂T)P,ni dT+(∂G∂P)T,ni dP+(∂G∂ξ)P, T dξ(1.5)

and hence

(∂G∂ξ)P, T=∑iviμi=ΔrG(1.6)

ΔrG is the change in the Gibbs free energy of a system that occurs during a reaction. Equilibrium and spontaneity criteria are defined from the second law of thermodynamics as follows.

For the case I, equilibrium is achieved at constant pressure and temperature when G is minimised, i.e. (dG)P,T=0 or ΔrG=0 from Eqs. 1.5 and 1.6.

For the case II, a system proceeds spontaneously during a reaction at constant pressure and temperature when (dG)P,T<0 or ΔrG<0.

We consider an oxidation reaction of a metal M such as

2a bM+O2( g)↔2 bMaOb(1.7)

From Eqs. 1.3 and 1.6, the expression of Gibbs free energy of reaction is written as follows:

ΔrG=∑μivi=2 bμMaOb−2a bμM−μO2=2 bμMaOb∘−2a bμM∘−μO2∘+RTlnaMaOb2/baM2a/b·aO2ΔrG=ΔrG∘+RTlnaMaOb2/baM2a/b·aO2

(1.8)(1.9)

According to the second law, when equilibrium is achieved, Eq. 1.9 becomes

ΔrG∘=−RTln(aMaOb2/b)eq(aM2a/b)eq·(aO2)eq=−RTlnK(1.10)

where ΔrG∘ is the standard Gibbs free energy of reaction and K the equilibrium constant used to describe the equilibrium state of the reaction system.

1.2.2 Construction and Use of Thermodynamic Diagrams

Diagrams allow determination of the oxide likely to form during selective oxidation of alloys for given operating conditions. The Ellingham diagram, i.e. the plot of the standard Gibbs free energy of reaction ( ΔrG∘ ) versus temperature, is useful to compare the relative stabilities of each oxide.

From Eq. 1.1, the standard Gibbs free energy of reaction ΔrG∘ is written as follows:

ΔrG∘=ΔrH∘−TΔrS∘(1.11)

ΔrH∘ and ΔrS∘ are calculated in temperature range without change of state from standard molar heat of formation and molar entropies assuming that the heat capacity contribution is negligible. From Eq. 1.11, ΔrG∘ can be plotted as a linear function of T, giving the Ellingham diagram as shown in Fig. 1.1. The M/MaOb line in the diagram indicates the oxidation reaction in the form of Reaction 1.7 which is normalised to consume one mole of O2. Because the values of ΔrG∘ are expressed in kJ per one mole of the oxygen gas so the stabilities of various oxides can be compared directly, i.e. the lower the position of the line on the diagram, the more stable is the oxide. If the activities of M and MaOb are taken as unity (i.e. pure condensed phases), Eq. 1.10 can be used to express the activity or the partial pressure of oxygen at which the metal and the oxide coexist. Then the dissociation pressure of the oxide ( pO2)eq  can be calculated as follows:

(aO2)eq=(pO2)eqp∘=exp(ΔrG∘RT)(1.12)

where p∘ is the standard pressure equalling to 1 bar.
From Eqs. 1.9 and 1.12, expression of Gibbs free energy of reaction is established as follows:

ΔrG=ΔrG∘−RTlnpO2p∘=RTln(pO2)eqpO2(1.13)

Equilibrium and spontaneity criteria relying on the second law are deduced from Eq. 1.13. If the oxygen partial pressure pO2 is higher than the equilibrium value ( pO2 ) eq , the metal will be oxidised (ΔrG<0), and if it is lower than the equilibrium value then the oxide will be reduced (ΔrG>0). For a given temperature T, the value of ( pO2)eq  can be determined directly from the Ellingham diagram by drawing a straight line from the origin marked Ω(T=0 K and ΔrG∘=0) through the point of the M/MaOb line at the temperature of interest T and reading the oxygen partial pressure from its intersection with the logarithmic scale at the right side labelled log(pO2)eq  (Fig. 1.1). This
log(pO2)eq  scale is built from a network of straight lines radiating from the origin Ω and corresponding to RTln((pO2)eq /p∘).

[picture]

Fig. 1.1. Ellingham diagram for selected oxides constructed using thermodynamic data from Kubaschewski et al. [15].

Fig. 1. 1. Ellingham diagram for selected oxides constructed using thermodynamic data from Kubaschewski et al. [15].

The second type of diagram is the stability diagram of condensed phases. It relies on Eq. 1.14 and consists to plot log (pO2)eq  with (pO2)eq  in bar versus the reciprocal temperature, as presented in Fig. 1.2 plotted using data from Ref. [16].

log(pO2)eq=ΔrG∘(ln10)RT=(ΔrH∘104(ln10)R)·104T−ΔrS∘(ln10)R(1.14)

[picture]

Fig. 1.2. Stability diagram of oxides calculated using FactSage and FactPS database (FactSage 7.0) [16].

Fig. 1. 2. Stability diagram of oxides calculated using FactSage and FactPS database (FactSage 7.0) [16].

For instance, for a usual stainless steel, AISI 304L (Fe-17.5Cr-8Ni-1.72Mn-0.32Si-0.02C) [17], diagram in Fig. 1.2 indicates that MnO and SiO2 can be formed selectively under Cr2O3 as shown in the upper figure of Fig. 1.3. Depending on the activity and diffusion coefficient of these elements in the metal, but also depending on the solubility and diffusivity of oxygen, these oxides can precipitate internally or form a continuous layer.

According to the diagram in Fig. 1.2, the oxide MnO is more stable than Cr2O3 and might be expected to form a sublayer beneath a Cr2O3 scale. This does not occur for several reasons. Firstly, MnCr2O4 spinel-type oxide is stable with respect to the binary oxides ( MnO and Cr2O3 ) at high temperature. Secondly, manganese is soluble and diffuses rapidly in Cr2O3 but relatively slowly in the alloy [18]. Then, the duplex oxide scale MnCr2O4/Cr2O3 is commonly observed on manganesecontaining stainless steels as shown in the lower figure of Fig. 1.3. The cross-section micrograph of AISI 304L oxidised at 850∘C in O2 in Fig. 1.4 reveals this typical duplex microstructure. When duration or temperature increases, fast growing iron-rich nodules appear and the oxide scale is no longer protective [17].

Even if the diagrams in Figs. 1.1 and 1.2 allow to explain the successive oxide formed from the oxide/gas interface to the metal/oxide interface, a more realistic approach is required because activity of elements is not equal to unity in the alloy, and mixed oxide (e.g. MnCr2O4 ) and solid solution oxide (e.g. Cr2−xFexO3 ) must be taken into account. Thermo-Calc software package [19] in combination with the suitable database can be used for thermodynamic calculations, based on the CALculation of PHAse Diagram (CALPHAD) method. The calculation of phase equilibria in a multicomponent system requires the total Gibbs free energy minimisation of the system. The CALPHAD method employs various phenomenological models (e.g., substitutional solution model for most of the disordered solution phases, sub-lattice model for phases composed of several sublattices on which the various components can mix as for carbides…) to describe the temperature,
pressure, and concentration dependencies of the Gibbs free energy functions of each phase that takes part in the equilibrium and to assign them an analytic expression. The assessed parameters for thermodynamic models for the individual phases, evaluated that the model fits the available experimental data for a given phase, are stored in thermodynamic databases linked to software for minimisation process and the calculation of phase equilibria. In fact, accuracy of calculations is dependent on the databases behind the package. A complete description of the CALPHAD technique can be found in literature [20, 21].

[picture]

Fig. 1.3. Schematic sketch of the oxide scale grown on the AISI 304L stainless steel.

Fig. 1. 3. Schematic sketch of the oxide scale grown on the AISI 304L stainless steel.

[picture]

Fig. 1.4. SEM-BSE cross-section micrograph (upper picture) and EDX element mapping after oxidation (lower picture) of the AISI 304L stainless steel oxidised at 850∘C in oxygen. Reproduced and adapted from A. Col et al., Corros. Sci. 114 (2017) 17-27 [17].

Fig. 1. 4. SEM-BSE cross-section micrograph (upper picture) and EDX element mapping after oxidation (lower picture) of the AISI 304L stainless steel oxidised at 850 ∘ C in oxygen. Reproduced and adapted from A. Col et al., Corros. Sci. 114 (2017) 17-27 [17].

For instance, for the AISI 304L stainless steel oxidised at O2 in oxygen, the TCFE database devoted to steels and Fe-base alloys is used [17]. Useful diagrams consist to plot the fractions of phase and the composition of each phase versus M2O3 activity as shown in Fig. 1.5. According to Fig. 1.5(a), once the corundum-type oxide 10−26 is formed, the oxygen partial pressure at the metal/oxide interface decreases to Cr2O3 bar. Oxygen partial pressure is then controlled by the decomposition pressure of the corundum oxide, which is O2 according to Fig. 1.5(b). Figs. 1.5(b) and (c) show the evolution of spinel and corundum oxide composition depending on oxygen activity. The iron enrichment of both oxides corresponds to an increase of oxygen activity and then to a less protective oxide scale.

[picture]

Fig. 1.5. (a) Volume fraction of phases, (b) atomic fraction of elements in corundum, and (c) atomic fraction of elements in spinel at different oxygen activities for AISI 304L oxidised at (MaOb) in oxygen. Redrawn and adapted from A. Col et. al., Corros. Sci. 114 (2017) 17-27 [17].

Fig. 1. 5. (a) Volume fraction of phases, (b) atomic fraction of elements in corundum, and (c) atomic fraction of elements in spinel at different oxygen activities for AISI 304L oxidised at ( M a O b ) in oxygen. Redrawn and adapted from A. Col et. al., Corros. Sci. 114 (2017) 17-27 [17].

1.3 Point Defects in Metal Oxide

The former section assumes that the interested material is a kind of bulk without defects. However, the oxide always contains defects, and importantly the growth of thermal oxide on metal consists of elementary steps including the diffusion process relevant to these defects. This section introduces the formation of point defects in the oxide by the interaction with oxygen gas in the atmosphere which could lead to the scenarios of the possible mechanisms of the oxide growth.

At high temperatures, metal is oxidised by the hot gas giving thermal oxide scale on its surface. The global reaction for the oxidation of metal ( M ) with oxygen gas (

aM(s)+b2O2( g)↔MaOb( s)(1.15)

) resulting in the metal oxide O−s is as follows:

O2+2 s↔2O−s(1.16)

The oxidation mechanism may start from the adsorption of the gas molecule from the atmosphere on the surface. An example of the surface reaction is Reaction 1.16 which is the dissociative adsorption of oxygen on the oxide/gas interface [13]. The s notation denotes an unoccupied site on the oxide/gas interface and VO represents the adsorbed oxygen at that interface.

VO••

For the inward growth of a compact oxide, after the adsorption process the adsorbed oxygen is incorporated into the oxide possibly by jumping into the oxide lattice or the oxide grain boundary. It therefore diffuses via these paths from the external oxide/gas to the internal metal/oxide interface. For the transport through the oxide lattice, the diffusion may occur by the jumping of oxygen vacancy or oxygen interstitial.

The vacancy at the oxygen site is simply called an oxygen vacancy and denoted as VOz•. If two electrons are still localised at the oxygen vacancy site, the effective charge which is the charge of the present defect ( -2 ) relative to the charge of oxygen in a perfect crystal ( -2 ) will be zero. If the oxygen vacancy is formed with a loss of one electron, the actual charge of the defect will be -1 , giving the effective charge of +1 . In such case we may describe the defect as 0≤z≤2 where one dot superscript denotes the effective charge of +1 . This defect may also be described in general as

a<M>↔aMM+bVOz•+bze′(1.17)

where z is a degree of ionisation ( <M> for oxygen). As for the oxidation mechanism, this defect is formed at the internal interface by the following reaction:

MM

where e′ is the metal in the substrate,

O−s+VOz•+ze′↔OO+s(1.18)

is the metal on the normal metal lattice site in the oxide, and

b2O2+bVOz•+bze′↔bOO(1.19)

is the electron. After the formation of the oxygen vacancy at the internal interface, oxygen on the oxygen site of the oxide jumps into this oxygen vacancy inwardly, resulting in the outward diffusion of the oxygen vacancy to the external interface. At the external interface, the adsorbed oxygen according to Reaction 1.16 jumps into the oxygen vacancy according to Reaction 1.18:

Oiz′

If it is assumed that the rate constants of the interfacial forward Reactions 1.16, 1.17 and 1.18 are very rapid compared with the diffusion rate through the oxide, these reactions thus attain equilibrium [13]. As a result, Reactions 1.16 and 1.18 which take place at the external interface can lead to Reaction 1.19. Fig. 1.16 (case I) schematises the possible oxidation mechanism by the oxygen vacancy for the growth of the metal oxide MO. The mentioned assumption will also be applied to the other three following cases where the oxygen interstitial, metal vacancy or metal interstitial are dominant.

h•

As for the oxygen interstitial, the oxygen that occupies an interstitial site is called an oxygen interstitial and denoted as

b2O2↔bOiz′+bzh•(1.20)

where one prime notation represents one negative effective charge. At the external interface, oxygen gas jumps into the oxide and forms the oxygen interstitial with electron hole (

a<M>+bOiz′+bzh•↔aMM+bOO(1.21)

) according to Reaction 1.20:

VMz′

Then it diffuses inwards to form the oxide at the internal interface by the following reaction:

0≤z≤

Fig. 1.6 (case II) schematises the oxidation mechanism by this defect for the growth of MO.
For the outward growth of the compact scale, the metal from the substrate diffuses from the internal to the external interface in order to react with oxygen in the atmosphere giving the new construction unit of metal oxide at the external interface. For the diffusion through the oxide lattice, firstly through the metal vacancy

b2O2↔bOO+aVMz′+azh•(1.22)

where

a<M>+aVMz′+azh•↔aMM(1.23)

higher oxidation state of M , this defect is formed at the external interface by the following reaction:

Miz•

Thermal oxide is grown by the outward jumping of metal on the metal site of oxide to the metal vacancy, resulting in the inward diffusion of metal vacancy through the oxide to the internal interface. At the internal interface, the defect is annihilated by the following reaction:

a<M>↔aMiz•+aze′(1.24)

The oxidation mechanism by this defect for the growth of MO is sketched in Fig. 1.6 (case III).
For the last case, the metal may jump into the interstitial site giving metal interstitial

b2O2+aMiz•+aze′↔aMM+bOO(1.25)

, this defect is formed at the internal interface according to the following reaction:

(VOz•,Oiz′,VMz′

[picture]

Fig. 1.6. Possible oxidation mechanisms of MO by Miz•) and (VOz•,Oiz′) transport.

Fig. 1. 6. Possible oxidation mechanisms of MO by M i z • ) and ( V O z • , O i z ′ ) transport.

This defect diffuses from the internal interface through the metal interstitial sites of the oxide to the external interface. At the external interface, it reacts with oxygen in the atmosphere giving the new construction unit of the oxide according to Reaction 1.25. Fig. 1.16 (case IV) schematises the oxidation mechanism by metal interstitial for the growth of MO.

VMz′

These four defects Miz• and Oiz′ may be categorised as anionic defects VMz′ and cationic defects ( VOz• and Miz• ). It is noticed from Fig. 1.16 that if the predominant defect is anionic (cases I and II) it constructs the new unit of oxide at the internal metal/oxide interface, giving the inward growth of the oxide. However, if the dominant defect is cationic (cases III and IV) the new oxide unit is constructed at the external oxide/gas interface, giving the outward growth of the oxide. Furthermore, those defects may be categorised as p-type defects ( Δm/A and

ΔmA=klt(1.26)

) and n -type defect ( Δm=mt−mt=o and t(mt) ) where z is not zero. The p-type defect is created with the formation of electron hole (Eqs. 1.20 and 1.22), while the n-type defect is created with the generation of electron (Eqs. 1.19 and 1.25).

1.4 Phenomenological Oxidation Kinetics

Macroscopically, the oxidation resistance could be assessed by the change of mass per unit area of the sample, a mass gain ( mt=0 ), as a function of the oxidation period of time. The relationship between mass gain and oxidation time may be linear, parabolic, linear-parabolic, cubic, logarithmic etc. In the case of linear oxidation kinetics, the relationship between mass gain and the oxidation time is

A

where k1 is the change of mass measured at the oxidation time CO2 with respect to initial mass before oxidation ( kp ),

(ΔmA)2=kpt(1.27)

is the oxidised surface area of the metal and d(Δm/A)/dt is the linear rate constant. If the oxidation kinetics is linear, it indicates that the rate determining step of the oxidation is a surface reaction [13]. The example for this is the oxidation of AISI 441 stainless steel oxidised in MMaOb atmosphere as explained in Chapter 6.

In the case of parabolic oxidation kinetics, the relationship between mass gain and time can be expressed by the following relation where MO2 is the parabolic rate constant:

MaOb

In this case the scale is protective since the increase in mass gain with time ( O2,ρMaOb ) is reduced at longer periods of oxidation. The parabolic oxidation indicates that the oxide growth is controlled by the diffusion of defect through the oxide scale. Derivation of this rate law based on the diffusion of point defects will be described in Section 1.5.2.

It is noted that the relationships between mass gain and time in Eqs. 1.26 and 1.27 might also be presented as the relationships between the oxide thickness and time. The relation between the oxide thickness ( X ) and mass gain for the global oxide formation Reaction 1.15 is given in Eq. 1.28.

In this equation, A and ΔmO2 are respectively the molar masses of the oxide Δm and the oxygen gas ρMaOb=(ΔmMaOb/ΔV)=(ΔmMaOb/AX) is the density of the oxide, ΔmMaOb is the oxidised surface area of the sample, and

X=2 b(MMaObMO2)1ρMaOb(ΔmA)(1.28)

is the mass of oxygen gas consumed to form the oxide which is in fact equal to the increase of mass that is measured from the experiment ( k ). The relation in Eq. 1.28 was derived using the stoichiometric relations in the oxidation Reaction 1.15 and the relation

k=k∘exp(−EaRT)(1.29)

where k∘ is mass of the oxide formed during oxidation.

k,Ea

Furthermore, it is often found from the experiment that the rate constant ( T ) relates with temperature in the following Arrhenius form:

kp

Here

logk=(−Ea104(ln10)R)·104T+logk∘(1.30)

is the pre-exponential factor having the same unit as Ea the apparent activation energy, R the gas constant, and 700−900∘C the absolute temperature. Eq. 1.29 can also be written as shown in Eq. 1.30. This relation indicates that the Arrhenius relation between the logarithm of kp and the reciprocal temperature is linear.

R2

As explained by Kofstad [22], when the reaction proceeds, a transition state is assumed to exist at the top of the free energy barrier separating the initial and the final states. It is considered that the species in the transition and the initial states are in equilibrium and the activation energy 141 kJ mol−1 is the height of the free energy barrier [22]. The change in the activation energy may suggest the possible change in the rate determining step of the oxidation reaction [13, 22].

Fig. 1.7(a) exemplifies the oxidation results of the iron oxidised in oxygen at the 1000∘C [23]. The value of mass gain is squared and plotted as a function of time as shown in Fig. 1.7(b). The linear relationship of the plot between square of mass gain and time exhibits the parabolic oxidation kinetics, and the slop of this plot is the value of the parabolic rate constant. Fig. 1.7(c) shows the logarithm of 10−6 g2 cm−4 s−1 as a function of the reciprocal temperature. The linear relation in this plot was obtained with the 10−10.5 of 0.997 . This means that the oxidation of iron follows the Arrhenius relation with the apparent activation energy, which was extracted from slop of the line in Fig. 1.7(c), of 10−12 g2 cm−4 s−1.

Fig. 1.8 shows the Arrhenius relation of different oxides [23-25]. It is seen that the parabolic rate constants of alumina and chromia are significantly lower than the rate of the formation of iron oxides on iron. For example, at 2×10−11 g2 cm−4 s−1(10−10.70 g2 cm−4 s−1), the parabolic rate constant for the formation of iron oxides is about 10−12.2. This value is lower to be about 10−13 g2 cm−4 s−1 to 2.7×10−13 g2 cm−4 s−1(10−12.57 g2 cm−4 s−1) for the chromia formation, such as Fe−25Cr−5Al for the growth of chromia on Fe-20Cr [13]. It is even lower to about

j→δ=−Dδ∂Cδ∂xu→x(1.31)

to j→δ for the alumina formation, such as δ(molcm−2 s−1),Dδ for the growth of alumina on δ(cm2 s−1),Cδ [13]. These results indicate the effectiveness of the chromia and alumina scale to combat the high temperature corrosion of metals.

[picture]

Fig. 1.7. (a) Mass gain as a function time, (b) square of mass gain as a function of time, and (c) the Arrhenius relation of the iron oxidised in oxygen at δ(molcm−3). Redrawn and adapted from S. Chandra-ambhorn et al., Corros. Sci. 148 (2019) 355-365 [23].

Fig. 1. 7. (a) Mass gain as a function time, (b) square of mass gain as a function of time, and (c) the Arrhenius relation of the iron oxidised in oxygen at δ ( m o l c m − 3 ) . Redrawn and adapted from S. Chandra-ambhorn et al., Corros. Sci. 148 (2019) 355-365 [23].

[picture]

Fig. 1.8. Arrhenius plot for the formation of different oxides [23-25].

Fig. 1. 8. Arrhenius plot for the formation of different oxides [23-25].

1.5 Derivation of the Parabolic Rate Law

1.5.1 Fick's First Law for the Diffusion of Defect through the Oxide

To derive the parabolic rate law, we start from the uni-directional phenomenological Fick's first law as follows [26, 27]:

u→x

Here,

j→δ=−DδCδRT∂μ~δ∂xu→x(1.32)

is the molar flux of a species μ~δ the diffusivity of the species δ the concentration of the species δ(μδ), and

μ~δ=μδ±zδFϕ(1.33)

the unit vector in x direction. The generalised Fick's first law [13, 26] is

μδ=μδ∘+RTlnaδ(1.34)

where zδ is the electrochemical potential of the species δ,F which might be charged. It relates to the chemical potential of the species 96485Cmol−1 as follows [13]:

ϕ

where

x(V),μδ∘

Here, δ is a degree of ionisation of the spicies aδ the Faraday's constant ( δ ) [28], δ the electric potential at the coordinate δ the chemical potential of the species Cδ at
standard state (J), and molcm−3 the activity of the species 1 mol cm−3. The positive (negative) sign is used in the right-hand side of Eq. 1.33 if the electric charge of the species

μδ=μδ∘+RTlnCδ or dμδ=RT dlnCδ=RTCδdCδ

(1.35)(1.36)

is positive (negative).

By assuming that the activity of A,μ~δ=μδ equals its concentration dμ~δ=dμδ=(RT/Cδ)dCδ in the unit of Cδ relative to the concentration at standard state i.e. molcm−3, it can be written that

δ([δ])

It is noticed that for the electrically neutral species Ω. From that relation and Eq. 1.36 we can obtain that (cm3 mol−1). In this case, the generalised Fick's first law in Eq. 1.32 can be reduced to be the phenomenological one in Eq. 1.31.

It is also noted that the molar concentration (

Cδ=[δ]Ω(1.37)

) in the unit of MaOb relates to the mole fraction of the species rMaOb according to Eq. 1.37 where molcm−2 s−1 is the molar volume

rMaOb=1ΩdX dt(1.38)

of the interested solution:

δ

Furthermore, the rate of the construction of ∂μδ/∂x=0 per unit area ( v→ ), in the unit of E→, can be written as a function of the rate of increase in oxide thickness ( X ) as follows [13]:

E→

In the diffusion of point defects, the species NC−1 may be charged and as a result affected by the applied electric field. To understand the phenomena, let consider the case when ϕ(Vm−1) and a positively charged species of 1 Coulomb is transported with the velocity of

\stackrel⇀E=−∂ϕ∂x\stackrel⇀ux(1.39)

along the direction of the electric field −kδv→δ. Electric field ( kδ ) is defined as a force applied on a positive charge per its charge ( E→ ) and equals the negative gradient of the electric potential E→=kδv→δ in the following form [29]:

Uδ

If it is supposed that the movement of this charged species causes the friction force exerted on it to be proportional to the velocity of the charged species, such friction force is therefore equal to (1/kδ) where m2 V−1 s−1 is a constant of proportionality [26]. The stationary limiting velocity is attained when this friction force is balanced by δ which in this case is the force applied by the electric field, i.e.

v→δ=±UδE→(1.40)

[26]. This relation may be expressed in general as in Eq. 1.40 where σ equals

iδ→=zδF→jδ→(1.41)

and called the electrical mobility in the unit of j→δ. The positive (negative) sign is used in the right-hand side of Eq. 1.40 when the charge of the species A is positive (negative) [30].

dl

Electrical mobility of the charged species contributes to many properties of materials, including the electrical conductivity ( i→δ ) of material. The relationship between the electrical mobility and conductivity can be obtained, starting from the following Faraday's law for the migration of the charged species under the pure electric field [26]:

Iδ=Δϕ/Rδ

Consider the case where the positively charged species with the molar flux of Iδ passes perpendicularly to the area Δϕ along the length Rδ, causing the current density of Iδ=∫i→δ·dA→. The Ohm's law states that Δϕ=∫E→· dl→ where R=dl/(σ·dA) is the electric current, dA→, dl→ is the potential difference between two terminals of interest, and E→ is the electrical resistance of the interested part. By inserting the terms

iδ→=±σE→(1.42)

and ∂μδ/∂x=0 and δ into the Ohm's law and by taking into account the fact that the directions of

j→δ=Cδv→δ=±CδUδE→(1.43)

and

σδ=zδFCδUδ(1.44)

are the same, it can be obtained that [26, 29]

∂μδ/∂x=0

where the positive (negative) sign is used in the right-hand side of this equation when the charge of the interested species is positive (negative).

Furthermore, in the case that δ, the molar flux of the species

j→δ=±zδFDδCδRTE→(1.45)

can also be written as [26]

Uδ=zδFRTDδ(1.46)

By inserting Eqs. 1.42 and 1.43 into Eq. 1.41, it can be obtained that

j→δ=−1Ω(Dδ∂[δ]∂x±Uδ[δ]∂ϕ∂x)u→x(1.47)

This equation expresses the relationship between the electrical mobility and conductivity. It will be used to compare the electrical conductivities of the charged species during oxide growth in the derivation of the parabolic rate law in Section 1.5.2.

Furthermore, in the case that kp, from Eqs. 1.32, 1.33 and 1.39, the molar flux of the species MaOb can be written in the following form:

j→VMz′=−1Ω(DVMz′∂[VMz′]∂x−UVMz′[VMz′]∂ϕ∂x)u→xj→h•=−1Ω(Dh·∂[ h•]∂x+Uh·[ h]∂ϕ∂x)u→x

(1.48)(1.49)

From Eqs. 1.43 and 1.45, the following relation can be obtained:

z[ VMz′]=[h•](1.50)

From Eqs. 1.32-1.34, 1.37 and 1.46, we can obtain

∂ϕ/∂x

The Fick's first law in this form will be used to derive the parabolic rate law as the following.

1.5.2 Parabolic Rate Law

In this section, we will derive the relationship between RT/zF with the defect concentration and oxygen partial pressure, starting from the case that the oxide DvMz′/UvMz′ is p-type with a dominant metal vacancy and a dominant electronic conductivity [13]. In such case, the relationship of the point defect with electron hole and oxygen partial pressure is shown in Reaction 1.22. From Eq. 1.47, the molar fluxes of metal vacancy and electron hole can be written as follows:

j→VMz′UVMz′[VMz′]+j→h•Uh•[ h·]=−(z+1)DVMz′ΩUVMz′∂ln[VMz′]∂xu→x(1.51)

with the electroneutrality condition of

Uh·[ h•]≫UVMz′[VMz′](1.52)

By eliminating

j→VMz′=−1Ω(z+1)DVMz′∂[VMz′]∂xu→x(1.53)

using Eqs. 1.48 and 1.49, then replacing the electron hole concentration with metal vacancy one using Eq. 1.50, and changing the constant ( z=0 ) to (z+1) using Eq. 1.46, we obtain

x=0

In the case that the electrical conductivity of electronic defect i.e. electron hole is significantly higher than that of the ionic defect i.e. metal vacancy, from Eq. 1.44 we may write that

[VMz′]e

From the relationship in Eq. 1.52, Eq. 1.51 can be simplified to be the following expression:

[VMz′]i

From this equation it can be seen that if the diffusing species is not charged (

∫[VMz′]e[VMz′]i d[ VMz′]u→x=∫0x−Ω·j→VMz′(z+1)DVMz′dx(1.54)

), Eq. 1.53 is simply the Fick's first law. However, if the diffusing species is charged with the degree of ionisation of z , the molar flux will be increased by the factor of

j→VMz′=1Ω(z+1)DVMz′([VMz′]e−[VMz′]iX)u→x(1.55)

. This equation can be integrated as shown in Eq. 1.54 from the position at the external interface ( MaOb ) where the metal vacancy concentration is

X2=kp′t(1.56)

to the position at the internal interface, i.e. at the position X corresponding to the value of the oxide thickness, where the metal vacancy concentration is

kp′=2a(z+1)DVMz′([VMz′]e−[VMz′]i)(1.57)

.

kp′

If it is assumed that the flux and diffusivity in equation are constant throughout the oxide, it can be obtained that

(kp)

From Eqs. 1.38, 1.55 and the stoichiometric relation between M and

kp=b22a(MO2MMaOb)2ρMaOb2(z+1)DVMz′([VMz′]e−[VMz′]i)(1.58)

in the oxide formation, we can obtain that

[VMz′]=(K1.22zaz)1a(z+1)pO2b2a(z+1)(1.59)

where

(Cr2O3)

Here pO2,e is the parabolic rate constant from the relation between the oxide thickness and time. By inserting X from Eq. 1.28 into Eq. 1.56, we can obtain the relationship between mass gain and the oxidation time in the form of Eq. 1.27 with the parabolic rate constant pO2,i of

5×10−28

From the equilibrium constant of Eq. 1.22 and the electroneutrality condition in Eq. 1.50, it can be obtained that

800∘C

It is noted that in the case of chromia

kp=b22a(MO2MMaOb)2ρMaOb2(z+1)DVMz′[VMz′]e(1.60)

thermally grown on chromium in air, the oxygen partial pressure at the external interface (

kp=(K1.22zaz)1a(z+1)b22a(MO2MMaOb)2ρMaOb2(z+1)DVMz′pO2,eb2a(z+1)(1.61)

) is 0.21 bar while that pressure at the internal chromium/chromia interface ( MaOb ) is drastically lower i.e. about kp bar at kp [15]. If the oxygen partial pressure at the external interface is drastically higher than that at the internal one, the metal vacancy at the external interface is also tremendously higher than that at the internal interface. In such case Eq. 1.58 can be reduced to be

1/2(z+1)

From Eqs. 1.59 and 1.60, we can obtain that

kp=(K1.201/bzz)1z+1 b2(MO2MMaOb)2ρMaOb2(z+1)DOiz′pO2,e12(z+1)(1.62)

Eqs. 1.60 and 1.61 express that the parabolic rate constant depends on many factors particularly the oxygen partial pressure at the external interface which prescribes the metal vacancy
concentration there. It depends on the diffusivity of metal vacancy in the oxide and the degree of ionisation of the defect. Also, it also depends on the molar mass and density of the oxide as well as a and b which represent the stoichiometric proportion between metal and oxygen in the oxide.

In the case of another p-type MaOb with a dominant oxygen interstitial and a dominant electronic conductivity,

kp=b2(MO2MMaOb)2ρMaOb2(z+1)DVOz•([VOz•]i−[VOz•]e)(1.63)

can be obtained in the form shown in Eq. 1.62. It is still observed that the [e′]=z[VOz•] also depends on the oxygen partial pressure but with the exponent of

[VOz•]=(1 K1.191/bzz)1/(z+1)·1pO21/2(z+1)(1.64)

.

800∘C

For the n-type pO2 which has oxygen vacancy as a dominant defect and also has a dominant electronic conductivity, we can derive in the similar manner as the case of metal vacancy to obtain the parabolic relation between mass gain and time in the form of Eq. 1.27. However, the rate constant in this case is

5×10−28

From the equilibrium constant of Eq. 1.19 and the electroneutrality condition i.e. 1/pO2,i1/2(z+1), it can be obtained that

1/pO2,e1/2(z+1)

In the case of chromia thermally grown on chromium in air at

kp=b2(MO2MMaOb)2ρMaOb2(z+1)DVOz[ VOz•]i(1.65)

, the

kp=(1 K1.191/bzz)1/(z+1)b2(MO2MMaOb)2ρMaOb2(z+1)DVOz1pO2,i1/2(z+1)(1.66)

in equilibrium between chromium and chromia is about kp bar as previously mentioned. This value is drastically lower than the oxygen partial pressure at the external interface which is 0.21 bar. Thus,

kp=(1 K1.25zaz)1/a(z+1)b22a(MO2MMaOb)2ρMaOb2(z+1)DMiz1pO2,ib/2a(z+1)(1.67)

is drastically higher than kp, and as a consequence the oxygen vacancy concentration at the internal interface is also tremendously higher than that at the internal interface. In such case Eq. 1.63 can be reduced to be

kp

From Eqs. 1.64 and 1.65, we can obtain that

kp

In the case of the n-type oxide with a dominant metal interstitial and a dominant electronic conductivity, VMz′ can also be obtained as follows:

Oiz′

It can be seen that the VOz• of the n-type oxide growth according to Eqs. 1.66 and 1.67 are independent of the oxygen partial pressure at the external interface, while the Miz• of the p -type oxide growth according to Eqs. 1.61 and 1.62 depend on the oxygen partial pressure in the atmosphere. Thus the variation of the Oiz′ with the oxygen partial pressure in the atmosphere can be used to identify if the dominant defect responsible for the growth is p-type (i.e. VMz′ or H2 ) or n-type (i.e. Fe−15.7wt.%Cr−8.5wt.%Mn and 900−1,100∘C ). For the p-type defect, we can further identify the dominant defect as an anionic defect ( O2 ) if the oxide growth is inward and as a cationic defect ( O2+H2O ) if the oxide growth is outward. The same consideration can be done for the case of the n-type defects. By this way, the dominant defect responsible for the oxide growth can be identified.

1.6 Summary

This chapter introduces the classification of stainless steels for the high temperature applications. The guidance of the maximum isothermal service temperature of some important stainless steels is given. Thermodynamic concept is further explained in order to give the framework for predicting the stable phases in the studied condition. Example of the thermodynamic analysis of the formation of thermal oxide on stainless steel is given. Further, types of defect in the oxide and the Fick's first law for the diffusion of defect though the oxide are explained. The oxidation kinetics is reviewed starting from the phenomenological laws of linear and parabolic reaction rate as well as the Arrhenius relation. The derivation of the parabolic rate law is given especially for the case that the oxide is p-type with a dominant metal vacancy and a dominant electronic conductivity.

References

The original version of this paper is available on https://www.scientific.net/SSP.300.1.pdf

A. Galerie, High temperature corrosion of chromia-forming iron, nickel and cobalt-base alloys, in: R.A. Cottis, M.J. Graham, R. Lindsay, S.B. Lyon, J.A. Richardson, J.D. Scantlebury, F.H. Stott (Eds.), Shreir's Corrosion, fourth ed., Elsevier, The Netherlands, 2010, pp. 583-605.

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S. Chandra-ambhorn, P. Saranyachot, Effect of the content in shielding gas on the microstructure and oxidation resistance of steel GTA welds, J. Mater. Process. Technol. 268 (2019) 18-24.

L. Antoni, A. Galerie, La corrosion sèche des métaux: cas industriels (M4 229), Techniques de l'Ingénieur, France.

A.J. Sedriks, Corrosion of Stainless Steels, second ed., John Wiley & Sons, USA, 1996.

S. Chandra-ambhorn, F. Roussel-Dherbey, F. Toscan, Y. Wouters, A. Galerie, M. Dupeux, Determination of mechanical adhesion energy of thermal oxide scales on AISI 430Ti alloy using tensile test, Mater. Sci. Technol. 23 (2007) 497-501.

B.A. Pint, High temperature corrosion of alumina-forming iron, nickel and cobalt-based alloys, in: R.A. Cottis, M.J. Graham, R. Lindsay, S.B. Lyon, J.A. Richardson, J.D. Scantlebury, F.H. Stott (Eds.), Shreir's Corrosion, forth ed., Elsevier, The Netherlands, 2010, pp. 606-645.

I.G. Wright, B.A. Pint, P.F. Tortorelli, High-temperature oxidation behavior of ODS-Fe3Al, Oxid. Met. 55 (2001) 333-357.

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P. Sarrazin, A. Galerie, J. Fouletier, Mechanisms of High Temperature Corrosion: A Kinetic Approach, Trans Tech Publications, Switzerland, 2008.

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This ePub file was created using machine-generated methods and automated software to convert the source text into an ePub format and may contain formatting inconsistencies from the original version. Formatting may vary depending on the device used.

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Solid State Phenomena, ISSN: 1662-9779, Vol. 300, pp 25-46

doi: 10.4028/www.scientific.net/SSP.300.25

© 2020 The Author(s). Published by Trans Tech Publications Ltd, Switzerland.

Submitted:
2018-12-27

Revised:
2019-08-06

Accepted:
2019-08-08

Online:
2020-02-12

CHAPTER 2
Mechanical Behaviour of Thermal Oxide Scales on Stainless Steels

The original version of this paper is available on https://www.scientific.net/SSP.300.25.pdf

Valérie Parry 1,a*, Somrerk Chandra-ambhorn 2,b, Thanasak Nilsonthi 2,c and Muriel Braccini 1, d

1 Université Grenoble Alpes, CNRS, Grenoble INP, SIMaP, F-38000 Grenoble, France

2 High Temperature Corrosion Research Centre, Department of Materials and Production

Technology Engineering, Faculty of Engineering, King Mongkut's University of Technology North Bangkok, 1518, Pracharat 1 Road, Bangsue, Bangkok, 10800, Thailand

a valerie.parry@grenoble-inp.fr, b somrerk.c@eng.kmutnb.ac.th, c thanasak.n@eng.kmutnb.ac.th, d  muriel.braccini@grenoble-inp.fr

Keywords: stress, strain, plane stress analysis, mechanical failure, interface, adhesion energy, mode-mixity, tensile testing, inverted blister test, modified 4-point bending test

Abstract

The mechanical behaviour and adhesion properties of thermal oxide scales are key issues for steel processing and long-term durability. This chapter aims at taking up the various aspects to be considered for such studies. The first part is devoted to a description of the origin of stress and stress quantification. Then, description of mechanical failure and damaging patterns of thermal oxide scales will be given. Finally, definitions of adhesion energy as well as quantitative methods to measure adhesion energy will be proposed. An appendix describing the hypotheses and the constitutive equations for plane stress analysis, which suits to oxide scales, is also given The purpose is enriched by references in particular to Alain Galerie's co-workers' publications.

The durability at high temperature of stainless steels depends on the resistance of the protective chromia-rich scale to chemical and/or mechanical failure. The tricky point is the coupling between the mechanical behaviour of the metal-oxide system and its chemical and microstructural changes during the oxide growth. Mechanical stress accumulation in oxide layer and metallic substrate is the driving force for the mechanical failure. Whereas resistance to high-temperature oxidation assumes the growth and the maintenance of a protective oxide scale which separates the metal from its environment, steel processing requires several de-scaling operations. Since it is desirable to be able to predict oxide scale adhesion, stress generation and relief in oxide scales are of importance. Although stresses can result from an external loading applied during high temperature treatment, additional stresses are generated during the oxidation process: growth stresses and thermal stresses. For this section, readers are also invited to consult Refs. [1-4].

Growth stresses happen during the isothermal part of the oxidation process. It can result from volume difference or epitaxial relationship between the metal and the oxide or from compositional changes or gradient in point defects.

The specific volume of the oxide is different from that of the metal which is consumed. Oxidation causes volume changes which are adjusted by strain in the oxide. In the case of scale formation occurring by inward diffusion of oxygen, the molar volume ratio described by Pilling and Bedworth (Pilling-Bedworth Ratio, PBR), may be first used to indicate the type of growth stress.

PBR=ΩMaOba·ΩM(2.1)

where ΩMaOb is the molar volume of the metal oxide MaOb and ΩM is the molar volume of the metal M.

If the ratio is greater than one, the resulting expansion could put the oxide into compression. In the case of a ratio lower than one, the oxide scale is subjected to tensile stress and a discontinuous oxide could result. Examples of molar volume ratios are displayed in Table 2.1.

Table 2.1 Pilling-Bedworth Ratio (PBR) of some common metal/oxide couples [2,5].

Oxide/metal | Al₂O₃/Al | ZrO₂/Zr | Cu₂O/Cu | NiO/Ni | FeO/α-Fe | TiO₂/Ti
PBR | 1.28 | 1.56 | 1.64 | 1.65 | 1.68 | 1.73
Oxide/metal | SiO₂/Si | Cr₂O₃/Cr | Fe₃O₄/α-Fe | Fe₂O₃/α-Fe | Ta₂O₅/Ta | Nb₂O₅/Nb
PBR | 1.9 | 2.07 | 2.10 | 2.14 | 2.5 | 2.68

The volume change linked to new oxide formation has to be accommodated at the metal/oxide interface. Considering no stress relieving mechanism, the resulting strain in the oxide εox is

εox=(PBR)1/3−1(2.2)

Considering an elastic behaviour for the oxide, the biaxial stress state (see Appendix) in the oxide would be

σox=−Eox1−voxεox(2.3)

where Eox is the elastic modulus and vox is Poisson's ratio of the oxide.
Although the PBR description is qualitatively useful, its quantitative use is limited. First, it does not apply to oxide growth by outward diffusion where oxide formation takes place at the outer free surface. In this case, oxide is not constrained to occupy the volume of the consumed metal and the stress in the scale is relatively low. In the case of oxide growth by inward anionic diffusion, the stress level calculated from Eq. 2.3 is overestimated. This approach assumes that no additional stress relaxation mechanisms such as plastic deformation (creep, dislocation motion) or fracture, take place. Moreover, complex transport phenomena such as mixed transport, grain boundary diffusion are not taken into account.

At the early stages of the oxide growth, the first germs to form are considered to have an epitaxial relationship with the metal. This lattice parameter mismatch between the metal and the oxide results in stress development.

The selective oxidation results in the depletion of one or more elements in the underneath metal substrate and thus the change in the alloy lattice parameter. In the same manner, dissolution of oxygen into metal with high solubility, volume change associated with internal precipitation of oxide or oxidation of internal carbides can cause volume expansion and produce stresses.

Variation of lattice parameter across the scale due to a gradient in point defects can generate stresses in oxide allowing deviations from stoichiometry. Oxide growth by outward metal diffusion results in vacancy injection in the substrate. Relaxation around the vacancies induces a latticeparameter variation resulting in stresses in the substrate or stress relief by enhancing creep within the metal.

Since metals and oxides have different thermal expansion coefficients ( α ), thermoelastic stress can be generated during temperature change. α is generally defined as

α=1L dL dT(2.4)

where L is the length of the considered material and dL/dT is the change of length with respect to the temperature T.

Examples of thermal expansion coefficients are displayed in Table 2.2.

Table 2.2 Thermal expansion coefficients of metals and oxides [3, 6].

System | Metal coefficient (× 10⁻⁶ K⁻¹) | Oxide coefficient (× 10⁻⁶ K⁻¹) | Ratio | Temperature range (°C)
Fe/FeO | 15.3 | 12.2 | 1.25 | 100–900
Fe/Fe₂O₃ | 15.3 | 14.9 | 1.03 | 20–900
Ni/NiO | 17.6 | 17.1 | 1.03 | 20–1000
Co/CoO | 14.0 | 15.0 | 0.93 | 25–350
Cr/Cr₂O₃ | 9.5 | 7.3 | 1.30 | 100–1000
Cu/Cu₂O | 18.6 | 4.3 | 4.32 | 20–750
Cu/CuO | 18.6 | 9.3 | 2.00 | 20–600

An intact and adherent oxide scale is considered. This implies that any deformation of the metallic substrate is transferred to the scale. Considering an elastic behaviour for both metal and oxide and a homogeneous stress, the resulting strain in the oxide dεox is

dεox=αox dT+(1−voxEox)dσx(2.5)

The biaxial stress state in the oxide is

dσox=−Eox(αox−αm)dTEoxhoxEmhm(1−vm)+(1−vox)(2.6)

where hox and hm are the thicknesses of scale and metallic substrate respectively, Eox is the elastic modulus of the oxide, vox and vm are Poisson's ratios of the oxide and metal respectively.

For a thin oxide scale on thick metallic substrate, the biaxial stress state in the oxide can be approximated by

dσox=−Eox(αox−αm)dT(1−vox)(2.7)

The thermally induced stress depends on the magnitude of temperature change and the difference between the thermal expansion coefficient of the oxide and the metal. According to Table 2.2, the coefficients of thermal expansion of metals are typically higher than that of oxides. It means that during cooling, the oxide scale will be subjected to compressive stress. It should be noted that this approach assumes that no stress relaxation mechanisms are in effect.

The stress level in the oxide cannot be measured directly. Its determination requires assumptions. X-Ray diffraction or Raman spectroscopy are based on the measurements of strain of the oxide lattice [7-10]. Fig. 2.1 displays the evolution of compressive stress in the oxide scale thermally grown on stainless steels, deduced from the Cr2O3 Raman peak shift [10]. Deflection technique is sensitive and permits to obtain a stress distribution. However creep data for both oxide and metal phases are needed. For more details readers are invited to consult Ref. [11].

[picture]

Fig. 2.1. Evolution of residual compressive stress, deduced from Cr2O3 Raman peak shift, according to the thickness of oxide scale thermally grown on ferritic stainless steels. Redrawn and adapted from S. Chandra-ambhorn et al., J. Power Sources 171 (2007) 688-695 [10].

Fig. 2. 1. Evolution of residual compressive stress, deduced from C r 2 O 3 Raman peak shift, according to the thickness of oxide scale thermally grown on ferritic stainless steels. Redrawn and adapted from S. Chandra-ambhorn et al., J. Power Sources 171 (2007) 688-695 [10].

Stresses generated by oxide growth and thermal mismatch can be accommodated by elastic or plastic deformation of the oxide and/or of the metal. However, when the stored energy in the oxide scale exceeds the fracture resistance of the oxide or of the metal/oxide interface, mechanical failure happens. Fracture is the most effective relaxation mode but it has severe consequence. Since fresh metal is exposed to oxidising environment, the corrosion rate increases.

Damaging patterns depends on the stress state, the properties and the microstructure of both oxide and metal. It varies from micro-cracks to spallation. If the oxide scale is put into tension, the appearance of cracks happens as soon as oxide yield stress value is reached. This limit corresponds to less than 1% of deformation, even at high temperature. If the oxide scale is under compressive stresses, which is the most common case, the scale will fracture and spall.

Two conditions are necessary for spallation: transverse cracking and interface decohesion. Decohesion can happen at the metal/oxide interface (adhesive failure) or along oxide scale layer interfaces (cohesive failure). Two cracking and spallation routes caused by compressive stress in the oxide, illustrated in Fig. 2.2, can be considered.

[picture]

Fig. 2.2. Spallation routes of oxide scale under compressive stress depending on the relative fracture resistance of the oxide and of the metal/oxide interface. Adapted from H.E. Evans, Inter. Mater.

Fig. 2. 2. Spallation routes of oxide scale under compressive stress depending on the relative fracture resistance of the oxide and of the metal/oxide interface. Adapted from H.E. Evans, Inter. Mater.

Rev. 40 (1995) 1-40 [4].

The first, called wedging, is characteristic of mechanically weak oxides with a strong adhesion to the substrate. It supposes primarily the failure of the oxide scale by compressive shear crack due, for example, to temperature drop or mechanical loading. The origin of oxide scale failure is linked to defects present in the scale.

The second, called buckling, describes the behaviour of a strong oxide with a weak interface with the substrate. Spallation starts from the buckling of the oxide scale, i.e. the metal/oxide interface decohesion. The origin of buckling may be linked to voids accumulation or segregation of elements at the metal/oxide interface.

In the case of a thin oxide of thickness h under biaxial compressive stress (see Appendix), the critical stress σc necessary to form a buckle of radius a is

σc=1.22E(1−v2)(ha)2(2.8)

Since σc increases with the square of the scale thickness, buckling route is less likely for thick oxide.

Depending on the mechanical stress experienced by a crack, three linearly independent fracture opening modes can occur: mode I, or mode of normal opening of the crack, and modes II and III, opening modes in shear as shown in the Fig. 2.3.

[picture]

Fig. 2.3. The three fracture modes.

Fig. 2. 3. The three fracture modes.

In a homogeneous isotropic material, the propagation of a crack is usually done such that the stress field at the crack tip is purely tensile (mode I). However, in the case of an interfacial crack, the propagation is generally done in mixed mode (mode I and mode II combined). This mode-mixity can be defined by an angle ψ expressing the ratio between shear ( τ ) and normal ( σ ) stresses as follows:

ψ=tan−1(τσ)(2.9)

Adhesion of oxide scales on metals can be assessed by qualitative consideration on the metal/oxide interface or by direct measurement of spallation in service. Adhesion and thus fracture are key parameters to better understand spallation phenomena. Two main interrelated conditions for fracture are generally proposed: energy or stress intensity factor approaches.

The energy criterion for fracture arises from Griffith's work [12] in which the presence of a preexisting crack is considered (Fig. 2.4). The energy stored in the volume ( Etot  ) is the sum of the potential energy ( Epot  ) arising from the stored elastic energy ( Eel  ) and the work performed by external forces ( Wext  ) and of the surface energy ( Esurf  ):

Epot=Eel−WextEsurf=2γA

(2.10)(2.11)

where γ is the interfacial energy and A is the crack area.

[picture]

Fig. 2.4. Crack of surface A in a volume submitted to external forces.

Fig. 2. 4. Crack of surface A in a volume submitted to external forces.

The energy releases by the crack growth is the energy release rate G. The energy balance requires a net decrease in energy and G is defined as

G=−∂Etot∂A(2.12)

According to Griffith, fracture occurs when the energy release rate is sufficient to supply the energy requirements for the creation of new fracture surfaces.

∂Etot∂A+2γ=0(2.13)

This approach yields a measure of GC, the fracture energy, or toughness, of the interface, which is

Gc=2γ(2.14)

GC is the energy value required to extend a crack over unit area. When the stored elastic energy in the oxide scale exceeds the fracture resistance of the interface, crack propagation and spallation will occur.

The value of Gc cannot be directly equated to that needed solely to rupture molecular bonds. So Gc differs from the thermodynamic adhesion energy Wad derived from the metal, oxide and interfacial surface energies:

Wad=γm+γox−γi(2.15)

Gc is several orders of magnitude higher than Wad, since it encompasses the dissipative-energy losses occurring around the crack tip ( Wdis  ), that is, viscoelastic and plastic energy losses in the crack tip damage zone. In many cases Wad  lies in the order of 1 J m−2 while Gc is between 1 and 1000 J m−2.

Moreover, GC value depends on the propagation mode of the crack (i.e. the mode-mixity). According to Fig. 2.5, the lowest values of Gc are obtained for the lowest value of ψ and corresponds to a fracture opening normal to the plane of the crack (mode I). Increasing mode II contribution increases GC values.

Phenomenological laws have been designed to describe this dependence, for example the one proposed by Hutchinson and Suo [13] which includes a parameter η rendering the dependency of the fracture energy to mode-mixity in the following form:

Gc(ψ)=GIc(1+tan2(ηψ))(2.16)

[picture]

Fig. 2.5. Evolution of Gc according the mode-mixity angle ψ. Redrawn and adapted from M. Braccini, Chapter 4 Interface adherence, in: M. Braccini, M. Dupeux (Eds.), Mechanics of Solid Interfaces, John Wiley & Sons, USA, 2012 [14].

Fig. 2. 5. Evolution of G c according the mode-mixity angle ψ . Redrawn and adapted from M. Braccini, Chapter 4 Interface adherence, in: M. Braccini, M. Dupeux (Eds.), Mechanics of Solid Interfaces, John Wiley & Sons, USA, 2012 [14].

The stress-intensity factor approach of fracture energy is related to Irwin [15]. According to his work, the stress field around a sharp crack in a linear-elastic material can be defined by a single parameter, the stress-intensity factor K. Fracture occurs when the value of K exceeds some critical value Kc. Thus, K is a stress-field parameter independent of the material, whereas Kc often referred to as the fracture toughness, is a measure of a material property.

Griffith and Irwin approaches are linked. The energy criterion clears the defect size and shape whereas values of KI depend on crack features. In the case of an infinite plane with a crack perpendicular to the loading direction (propagating in mode I) submitted to an uniaxial tensile stress σ, the relationship between energy release rate and stress intensity factor in mode I is

G=KI2·1E(2.17)

where E is the elastic modulus and σ is the stress field.

Eqs. 2.18 and 2.19 are two examples of stress intensity factors in infinite volumes. The case of a straight crack of length 2a and of a circular crack of radius a perpendicular to the loading direction (mode I) are illustrated Fig. 2.6.

[picture]

Fig. 2.6. Examples of stress-intensity factors for a straight crack of length 2a (left) and a circular crack of radius a (right) perpendicular to the loading direction (mode I) in infinite volumes. Redrawn and adapted from T.L. Anderson, Fracture Mechanics: Fundamental and Applications, third ed., CRC Press, USA, 2005 [16].

Fig. 2. 6. Examples of stress-intensity factors for a straight crack of length 2 a (left) and a circular crack of radius a (right) perpendicular to the loading direction (mode I) in infinite volumes. Redrawn and adapted from T.L. Anderson, Fracture Mechanics: Fundamental and Applications, third ed., CRC Press, USA, 2005 [16].

KI=σπa for a straight crack of length 2aKI=2σaπ for a circular crack of radius a(2.19)

To assess the scale adhesion, several experiments have been designed to determine quantitative values of adhesion energy of metal-oxide system. Each configuration is characterised by loading conditions which act on the crack propagation. Indeed, in the case of a metal/oxide interface the difference between the mechanical behaviour of the two materials of both sides of the interface generates asymmetry even if the geometry and the loading are symmetrical. As long as the interface represents the easiest break path, the crack is trapped and will not bifurcate easily as it would in a homogeneous fragile material. Thus mixed-mode conditions of loading affect the propagation of the interfacial crack and thus the value of adhesion energy. As a consequence, comparison between values of adhesions energy obtained from different tests must be done cautiously.

In the following, three methods will be described: the inverted blister test, the modified 4-point bending test and in-situ tensile testing. The angle ψ expressing the ratio between shear (mode II) and normal (mode I) stresses has been evaluated for the three tests [17]. Results are presented in Table 2.3. The loading of the inverted blister testing is mainly in mode I, whereas tensile testing loading is mainly in mode II. The modified 4-points bending test corresponds to equivalent contributions of mode I and mode II.

Table 2.3 Mixed mode angles for different method of adhesion energy measurements [17].

Test | Inverted blister test | Modified 4-points bending test | Tensile testing
ψ angle (°) | 25 | 40 | close to 90

For a given test, the reproducibility and the energy released during crack initiation and crack propagation are different. Crack initiation is a stochastic mechanism which may require a high value of energy to happen whereas crack propagation is more reproducible. Thus it provides a better description of the adhesion properties of an interface. Adhesion test measurements have to be performed during crack propagation although it is not always possible.

The inverted blister test is a configuration of the blister test [18] adapted to the brittle nature of the oxide [19-20]. One side of an oxidised metallic foil is glued, on the oxide side, over the central hole of a sample holder as sketched in Fig. 2.7 (top). The relative thicknesses of the different layers are not scaled in the figure. The thickness of the stainless steel foil is around 150μ m while the thermal oxide is about 1−3μ m thick, see for examples Ref. [21]. Different glues can be used depending on the toughness of the metal/oxide interface. For example, Araldite® allows measuring adhesion energies up to ~200 J m−2. The oxide layer is locally removed by grinding through the sample holder hole in order to reveal the bare metal. The test, described on Fig. 2.7 (middle and bottom), consists in injecting in the hole an inert and incompressible fluid at a controlled pressure with a micro-syringe to create a blister. The altitude and the diameter of the blister can be measured by projecting interference fringes on the bulge surface, giving iso-altitude lines by contour analysis profilometry. The evolution of the fluid pressure (P) as a function of the bulge attitude (h) is shown in Fig. 2.8.

[picture]

Fig. 2.7. Schematics of the experimental set-up (top) and the two stages of the inverted blister test (middle and bottom). Adapted from J. Mougin et al., Mater. Sci. Eng., A 359 (2003) 44-51 [21] and from R.J. Hohlfelder et al., Mat. Res. Soc. Symp. Proc. 356 (1995) 585-590 [22].

Fig. 2. 7. Schematics of the experimental set-up (top) and the two stages of the inverted blister test (middle and bottom). Adapted from J. Mougin et al., Mater. Sci. Eng., A 359 (2003) 44-51 [21] and from R.J. Hohlfelder et al., Mat. Res. Soc. Symp. Proc. 356 (1995) 585-590 [22].

[picture]

Fig. 2.8. Evolution of the fluid pressure according to the bulge altitude. Redrawn and adapted from M. Braccini, Chapter 4 Interface adherence, in: M. Braccini, M. Dupeux (Eds.), Mechanics of Solid Interfaces, John Wiley & Sons, USA, 2012 [14].

Fig. 2. 8. Evolution of the fluid pressure according to the bulge altitude. Redrawn and adapted from M. Braccini, Chapter 4 Interface adherence, in: M. Braccini, M. Dupeux (Eds.), Mechanics of Solid Interfaces, John Wiley & Sons, USA, 2012 [14].

When increasing the amount of injected fluid, the metallic part of the sample blow up, forming a bulge of constant diameter. The bulge altitude increases with the fluid pressure. This "bulge state" allows measuring the Young's modulus and the residual stress of the metallic foil. For a critical pressure, damaging happens and a circular crack propagates along the metal/oxide interface. The diameter of the bulge increases with the amount of injected fluid. The bulge altitude increases as well whereas the fluid pressure decreases. It is assumed that no energy release occurs in the oxide layer during the crack propagation along the metal/oxide interface, in particular the long-range average biaxial stress in the oxide layer (see Appendix) do not contribute to crack propagation. Only the energy afforded by the injected fluid is taken into account. So, the pressure drop curve of the "blister stage" is described by an energy balance between the energy supplied to the system by the fluid under pressure and the energy consumed by the deformation of the metallic foil and by the propagation of the interfacial crack. In the case of an elastic behaviour of the metallic substrate, the pressure drop is described by a hyperbole [21,22].

Gc=CPh(2.20)

where P is the fluid pressure, h is the bulge altitude and C is a constant depending on the mechanical properties and the residual stress of the metallic foil.

Refs. [17, 21, 23, 24] from Alain Galerie's co-workers, provide applications of the inverted blister test for adhesion energy measurements.

The geometry of the modified 4-point bending test is presented in Fig. 2.9 [25-27]. The proportions between the different layers in Fig. 2.9 are not scaled. Stiffeners or counterplates are glued on the oxide scale. The notch in the counterplates corresponds to a pre-crack. The system is submitted to a vertical loading at a constant rate, leading first to elastic bending of the sample, then to interface decohesion and crack propagation. The force Fp corresponding to crack propagation is constant. The typical evolution of the applied force ( F ) as a function of the displacement ( Δl ) during the test is shown in Fig. 2.10. In the case of identical thicknesses and mechanical properties of the substrate and of the counterplates, the adhesion energy is derived from Fp as follows:

Gc=2116Fp2l2E¯h3b2(2.21)

Here h=h1+h2 and E¯ equals E or E/(1−v2) for the plane stress and plane strain assumption respectively. E and v are respectively the elastic modulus and the Poisson's ratio of the substrate which is equal to the counterplate under this assumption [14]. Illustrations of adhesion energy measurements with 3 - and 4 -point bending tests can be found in Refs. [17, 27].

[picture]

Fig. 2.9. Schematic of the 4-point bending test and typical force displacement curve. Adapted from P.G. Charalambides et al., Mech. Mater. 8 (1990) 269-283 [25], I. Hofinger et al., Int. J. Fract. 92 (1998) 213-220 [26] and M. Zhe et al., J. Adhes. Sci. Technol. 26 (20120 1-17 [27].

Fig. 2. 9. Schematic of the 4-point bending test and typical force displacement curve. Adapted from P.G. Charalambides et al., Mech. Mater. 8 (1990) 269-283 [25], I. Hofinger et al., Int. J. Fract. 92 (1998) 213-220 [26] and M. Zhe et al., J. Adhes. Sci. Technol. 26 (20120 1-17 [27].

[picture]

Fig. 2.10. Typical force-displacement curve during the 4-point bending test. Adapted from M. Zhe et al., J. Adhes. Sci. Technol. 26 (2012) 1-17 [27].

Fig. 2. 10. Typical force-displacement curve during the 4-point bending test. Adapted from M. Zhe et al., J. Adhes. Sci. Technol. 26 (2012) 1-17 [27].

For the tensile testing, oxidised specimens are loaded in a tensile machine designed to be placed under an optical microscope or in the scanning electron microscope chamber. The central region of the specimen is observed whereas a regular elongation rate is applied. The displacement motion is interrupted in order to acquire images of the sample surface. Sequential micrographs of the evolution of the damage in the oxide scale are collected according to the applied strain [9, 10, 17, 24, 28-38].

[picture]

Fig. 2.11. The possible scenarios for scale decohesion during tensile testing. σx,yx and σy,xx are normal stresses exerted on the oxide film in the loading direction and in the direction perpendicular to the loading direction respectively.

Fig. 2. 11. The possible scenarios for scale decohesion during tensile testing. σ x , y x and σ y , x x are normal stresses exerted on the oxide film in the loading direction and in the direction perpendicular to the loading direction respectively.

At the beginning of the test, the substrate and the oxide scale are elastically strained. Whereas the plastic deformation of the substrate occurs, cracks happen in the oxide scale. These cracks allow a stresses redistribution, in particular shear stresses appearance at the metal/oxide interface. The volume conservation in plasticity imposes a compressive strain of the scale in the direction perpendicular to the traction direction. Depending on the scale adhesion, different scenarios for decohesion, displayed in Fig. 2.11, are possible: delamination, wedging (cracking of the layer then detachment) and buckling (detachment then buckling).

The evolution of the spallation ratio as a function of the imposed strain for thermal oxide scales grown on AISI 444 at 800∘C in synthetic air for different oxidation times, giving different scale thicknesses, are displayed in Fig. 2.12. The spallation ratio is measured from electronic or optical images and corresponds to the spalled area normalised by the total area of the sample (metal and oxide) in the image. According to Fig. 2.12, the beginning of spallation as well as the saturation of the spallation ratio of the thickest oxide happen for the lowest imposed strain value. Thus, the related curves of the spallation ratio with respect to the imposed strain are characterised by high slope which indicates the higher sensibility of thick oxides to be spalled out with the increase of the imposed strain. This behaviour was observed by François Toscan [29] and many times later in the studies under Galerie's supervision [10, 30]. It is therefore called the "François effect" as named by A. Galerie. The determination of adhesion energy from tensile testing experiments at the strain initiating the first spallation is described in the following part.

[picture]

Fig. 2.12. Spallation ratio as a function of the imposed strain of AISI 444 oxidised at 800∘C in synthetic air to obtain different oxide thicknesses. Redrawn and adapted from S. Chandra-ambhorn et al., J. Power Sources 171 (2007) 688-695 [10].

Fig. 2. 12. Spallation ratio as a function of the imposed strain of AISI 444 oxidised at 800 ∘ C in synthetic air to obtain different oxide thicknesses. Redrawn and adapted from S. Chandra-ambhorn et al., J. Power Sources 171 (2007) 688-695 [10].

Adhesion energy can be determined from tensile testing. This approach relies on the calculation of the energy stored in the oxide until spallation takes place.

Oxides scales thermally grown on thick substrate can be considered as thin flat plates with an homogeneous microstructure that are acted upon only by load forces that are parallel to them. The scales are under plane stress state in the plane of the oxide ( x,y ), the stress vector along the direction z perpendicular to the oxide film is zero.

Considering an elastic behaviour of the oxide scale, the strain energy is

w=∑i,j12σijεij(2.22)

M. Dupeux and A. Galerie proposed that the evolution of the stress in the oxide during tensile testing can be described as a succession of 3 stages [9, 10, 17, 24, 29-37]. This approach considers an elastic-plastic behaviour of the metallic substrate.

Stage I describes the residual compressive stress state of the oxide before tensile testing. The residual elastic strain energy in the oxide is related to growth and thermal stresses occurring during oxidation and cooling down respectively. These stresses put the oxide scale into compression. The equibiaxial stress state in the plane of the oxide (x,y) is

σxx=σyyσij=0 for i≠jσzz=0

(2.23)(2.24)(2.25)

The resulting strain in the oxide in the direction i where i,j stand for x or y(εi) is

εi=σiEox−voxσjEox=σiEox−voxσiEox=(1−voxEox)σi=σiM(2.26)

with

M=Eox1−vox(2.27)

where Eox is the elastic modulus of the oxide scale and vox is the Poisson's ratio of the oxide scale.
Finally, the residual elastic strain energy in the direction i during stage I ( wres, i ) is

wres,i=12σiεi=12σi2M=12σres2M(2.28)

and the total residual elastic strain energy ( wres  ) is

wres=wres,x+wres,y=σres2M(2.29)

The residual elastic strain energy in each direction corresponds to the shaded area on the stressstrain representation displayed in Fig. 2.13.

[picture]

Fig. 2.13. Representation of the elastic stored energy in the oxide scale under compressive stress (stage I).

Fig. 2. 13. Representation of the elastic stored energy in the oxide scale under compressive stress (stage I).

Stage II describes the beginning of tensile testing, until the metallic substrate reaches its yield point. Assuming an intact and adherent oxide scale, the substrate deformations are transferred to the scale:

εx,ox=εx, mεy,ox=εy, m=−vmεx, m

(2.30)(2.31)

Here, εx,OX and εx,m are respectively the strain components in the tensile (x) direction in oxide and metal respectively, while εy,ox and εy,m are the strain components in the perpendicular (y) direction in the oxide and metal respectively.

According to Eq. 2.60 from Appendix describing the plane stress analysis, the planar stress in the oxide in the direction i where i stands for x or y(σi,ox) is

σi,ox=Eox1−vox2(εi,ox+voxεj,ox)(2.32)

According to Eq. 2.32 the stresses in the oxide can be written as a function of εx,m as follows:

σx,ox=Eox1−vox2(1−voxvm)εx, mσy,ox=Eox1−vox2(vox−vm)εx, m

(2.33)(2.34)

The components of the elastic strain energy due to tensile testing during Stage II ( wel ) in oxide in both x and y directions are

welx,ox=12(Eox1−vox2)(1−voxvm)εx, m2wely,ox=12(Eox1−vox2)(vm2−voxvm)εx, m2

(2.35)(2.36)

In the particular case of the same Poisson's ratios for metal and oxide, the stress along y axis is zero and there is no elastic stored energy in this direction. The elastic strain energy is simplified as follows:

wel,ox=12Eoxεx, m2(2.37)

Stage III corresponds to the phase for which plastic deformation of the substrate occurs without oxide scale spallation, the oxide is still intact and adherent and so the plastic deformations of the substrate are transferred to the scale.

εox=εm(2.38)

However, during plastic deformation, the volume of the metal remains constant.

Vm=Vm(1+εx, m)(1+εy, m)(1+εz, m)(2.39)

The plastic deformation of the metal is supposed to be isotropic in the plane perpendicular to the tensile direction. This implies that strains in the metal along y and z directions are identical. According to Eq. 2.39, strain in the metal in the direction y can be written as a function of the εx,m, the strain component in the metal in the tensile direction. The strains components are then

εy, m=11+εx, m−1(2.40)

According to Eq. 2.60 from Appendix, the planar stress in the oxide in the direction i is

σi,ox=Eox1−vox2(εi,ox+voxεj,ox)(2.41)

The components of the stress in the oxide due to plastic deformation of the substrate during tensile testing are

σx,ox=Eox1−vox2(εx, m+vox(11+εx, m−1))(2.42)

σy,ox=Eox1−vox2(11+εx, m−1+voxεx, m)(2.43)

The plastic deformation of the substrate remains small [35] and Taylor expansions of (1+εx,m)−1/2 in the neighbourhood of zero can be used.

(1+εx, m)−1/2=1−12εx, m(2.44)

The simplified components of the stress in the oxide due to plastic deformation of the substrate during tensile testing are

σx,ox=(Eox1−vox2)(1−12vox)εx, mσy,ox=(Eox1−vox2)(vox−12)εx, m

(2.45)(2.46)

Finally the components of the elastic strain energy during stage III of tensile testing ( wpl ) are

wplx,ox=12(Eox1−vox2)(1−12vox)εx, m2wply,ox=−12(Eox1−vox2)(vox−12)εx, m2

(2.47)(2.48)

It is noted that the strain energy in the oxide can also be written as a function of the strain of the metal in y direction, as shown in Eqs. 2.49 and 2.50 for the situations in stages II and III respectively.

wely,ox=12(Eox1−vox2)(1−voxvm)εy, m2wplyy,ox=12(Eox1−vox2)(1−2vox)εy, m2

(2.49)(2.50)

Fig. 2.14 shows evolution of stress in the oxide in x and y directions during straining. Stages I, II, III correspond to the situations from A to B,B to C and C to D respectively.

Adhesion energy ( GC ) can be calculated using Eq. 2.51 where wox is the total strain energy in the oxide and hox is the oxide thickness. As in Eq. 2.52, the total strain energy is the sum of the stain energies in stages I, II and III i.e. wres,ox , wel,ox  and wpl,ox respectively. The quantified values of the adhesion energies were obtained in the range of tens to hundreds Jm−2 for AISI 430Ti oxidised in Ar−20%O2 at 900∘C [31].

Gc=woxhox(2.51)

wox=wres,ox+wel,ox+wpl,ox(2.52)

[picture]

Fig. 2.14. Stress-strain curves representing the strain energies of an oxide scale thermally grown on AISI 430Ti along loading ( x ) (top) and transverse ( y ) (bottom) directions. Redrawn and adapted from S. Chandra-ambhorn et al., Mater. Sci. Technol. 23 (2007) 497-501 [31].

Fig. 2. 14. Stress-strain curves representing the strain energies of an oxide scale thermally grown on AISI 430Ti along loading ( x ) (top) and transverse ( y ) (bottom) directions. Redrawn and adapted from S. Chandra-ambhorn et al., Mater. Sci. Technol. 23 (2007) 497-501 [31].

This chapter reviews theoretical bases and experimental methods allowing appreciating mechanical behaviour of oxide scale thermally grown on stainless steel.

Mechanical stress accumulation in oxide layer and metallic substrate is the driving force for failure. Stress originates from the coupling between the mechanical behaviour of the metal-oxide system and its chemical and microstructural changes during the oxidation process. Growth stress due to the inward diffusion of oxygen is evaluated by the Pilling-Bedworth ratio (PBR), the ratio between the molar volumes of oxide and metal. PBR values higher than one indicate the compressive stress development in the oxide while PBR values lower than one indicate the tensile stress development in the oxide which can make the oxide susceptible to fail. Growth stress can also arise from epitaxial stress, compositional changes or gradient of point defects. Stress level in the oxide cannot be directly measured and its experimental determination requires assumptions. However, considering oxide layer grown on metallic substrate as thin flat plates with a homogeneous microstructure and free from stress relieving mechanism, linear elasticity equations allow calculating the strain and the equibiaxial stress state in the plane of the oxide generated by oxide growth and thermal mismatch.

Stresses can be accommodated by elastic or plastic deformation of the oxide and/or of the metal. However, when the store energy in the oxide scale exceeds the fracture resistance of the oxide or of the metal/oxide interface, mechanical failure happens. The resulting damaging pattern depends on the stress state, the properties and the microstructure of both oxide and metal. It varies from microcracks to spallation. Under compression, the oxide is failed by wedging (respectively buckling) if the cohesive strength in the oxide is weaker (respectively stronger) than the metal/oxide interfacial adhesive strength. Definitions of adhesion energy (or fracture toughness) are presented through two main interrelated conditions for fracture: energy criterion and stress intensity factor. Several experiments designed to determine quantitative values of adhesion energy are also described: the inverted blister test, the modified 4-point bending test and in-situ tensile testing. Each configuration is characterised by loading conditions. Mode-mixed conditions of loading affect the propagation of the interfacial crack and thus the value of adhesion energy. As a consequence, comparison between values of adhesions energy obtained with different tests must be done cautiously.

Appendix: Plane Stress Analysis

The original version of this paper is available on https://www.scientific.net/SSP.300.25.pdf

This appendix summarises the hypothesis and the formulation of stress and strain in the frame of plane stress analysis.

A cubic element is submitted to stresses in every direction (Fig. 2.15). For each face of the cube:

σi is the normal stress, parallel to i axis ( i stands for x,y or z ).

τi,j and τi,k are shear stresses lying to the plane perpendicular to i axis. τi,j applies along the j direction.

[picture]

Fig. 2.15. Stress element in Cartesian coordinate.

Fig. 2. 15. Stress element in Cartesian coordinate.

σi is a normal stress applying in the direction i on an isotropic material. Considering a linear relation between stress and strain, Hooke's law states that the resulting strain εi,σj in the direction i is

εi,σj=σiE(2.53)

where E is the elastic or Young's modulus. The strains induced by σi in transverse directions j and k tend to shrink the element. It is generally assumed that εj,σj is linked to the normal stress via Poisson's ratio.

εj,σj=−vεi,σj=−vσiE(2.54)

Now considering normal stresses in every direction, the strain εi along the i direction is given by the Young's relation.

εi=εi,σi+εi,σj+εi,σk(2.55)

Combining Eqs. 2.54 and 2.55, the strain εi along the i direction is

εi=εi,σi−vσiE−vσkE(2.56)

Conversely, the expression of σi, the stress along the i direction, according to strains is given by the Lamé's relation.

σi=E(1+v)(1−2v)[(1−v)εi+v(εj+εk)](2.57)

Oxides scales thermally grown on thick substrate can be considered as flat plates that are acted upon only by load forces that are parallel to them. In such case, the stress vector along the direction perpendicular to the film is zero. The material is said to be under plane stress.

Let consider an oxide scale in the (x,y) plane submitted to residual stresses and/or tensile testing. The oxide is under plane stress and σz is zero. In order to avoid shear stress components, τzx and τzy, the x axis of the frame is chosen to correspond to the tensile direction. The expression of εi and σi ( i stands for x or y ), the strain and the stress in the i direction, are

εi=σiE−vσjEεz=−vσxE−vσyEσi=E1−v2(εi+vεj)

(2.58)(2.59)(2.60)

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The original version of this paper is available on https://www.scientific.net/SSP.300.25.pdf

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Solid State Phenomena, ISSN: 1662-9779, Vol. 300, pp 47-80

doi: 10.4028/www.scientific.net/SSP.300.47

© 2020 The Author(s). Published by Trans Tech Publications Ltd, Switzerland.

Submitted:
2019-04-11

Revised:
2019-05-26

Accepted:
2019-08-08

Online:
2020-02-12

Characterisation of Thermal Oxide Scales on Stainless Steels

The original version of this paper is available on https://www.scientific.net/SSP.300.47.pdf

Anusara Srisrual 1,a*, Thammaporn Thublaor 2, b and Patthranit Wongpromrat 3,c

1 Department of Industrial Physics and Medical Instrumentation, Faculty of Applied Science, King Mongkut's University of Technology North Bangkok, 1518, Pracharat 1 Road, Bangsue, Bangkok, 10800, Thailand

2 High Temperature Corrosion Research Centre, Department of Materials and Production Technology Engineering, Faculty of Engineering, King Mongkut's University of Technology North Bangkok, 1518, Pracharat 1 Road, Bangsue, Bangkok, 10800, Thailand

3 Department of Chemical Engineering, Faculty of Engineering, King Mongkut's Institute of Technology Ladkrabang, Chalongkrung 1 Road, Lat Krabang, Bangkok, 10520, Thailand

a  anusara.s@sci.kmutnb.ac.th, b  thammaporn.tt@gmail.com, c  patthranit.wo@kmitl.ac.th

Keywords: thermogravimetric method, chromium volatilisation, scanning electron microscope, energy dispersive X-ray spectroscopy, focused ion beam, Raman spectroscopy, X-ray diffractometer, X-ray photoelectron spectroscopy, photoelectrochemistry

Abstract

This chapter aims at reviewing the characterisation techniques that are commonly used for high temperature oxidation study, especially on stainless steels. In addition, the experimental studies about the high temperature oxidation i.e. thermogravimetric method and chromium volatilisation measurement are explained. The various kinds of characterisation techniques for physico-chemical and electronic properties of thermal oxide scales are reviewed, starting from optical microscopy (OM), scanning electron microscopy (SEM), energy dispersive X-ray spectroscopy (EDS), focused ion beam coupled with scanning electron microscope (FIB/SEM), X-ray diffractometer (XRD), X-ray photoelectron spectroscopy (XPS), Raman spectroscopy (RS), and photoelectrochemical characterisation (PEC). The review focuses on the basic concepts and shows how the characterising tools can be applied to thermal oxide characterisation.

CHAPTER 3

The original version of this paper is available on https://www.scientific.net/SSP.300.47.pdf

3.1 Introduction

In high temperature corrosion study, thermal oxide scales are interested in many aspects such as interfaces, morphologies, species, semiconducting types, and growth mechanisms, because they affect lifetime of material. The various techniques are used to charaterise the materials.

The surface morphology of thermal oxide grown on alloy can be simply observed by optical microscope and scanning electron microscope. The literature [1-5] always reports the oxide morphology obtained from electron micrographs because it helps visualise the oxide morphology and topography. The focused ion beam coupled with scanning electron microscope is the superior technique because the focused ion beam is used for the cross-sectioning preparation inside the vacuum chamber when the electron microscope is operated, allowing to reveal the information of matter from surface to the desired depth such as morphology [6-7]. Chemical compositions of oxides are generally analysed by energy dispersive X-ray spectroscopy [8-10] and X-ray photoelectron spectroscopy [11-14]. The oxide phases are widely identified by X-ray diffractometer [12-16] and Raman spectroscopy [17-19]. Photoelectrochemical technique, which is generally used for finding electronic characteristics of matter, can be applied to identify thermal oxide phases with a high performance and a superior local analysis [20,21].

Apart from the characterisation techniques, the experimental methods for simulating oxidation conditions are also important. The set-up generally consists of a furnace and a gas system. A simple set-up uses a typical furnace such as a tubular furnace and the mass gain can be subsequently calculated by weighting outside the furnace. Thermobalance is a superior method because a balance
is equipped with a furnace, providing a real-time weighting during oxidation and permitting an oxidation kinetic study. In addition, the experiments can be set-up depending on the study aspects such as the study of chromium volatilisation. In some oxidising conditions, trivalent chromium compound ( Cr3+ ) is oxidised to the hexavalent chromium ( Cr6+ ) compound, which can be volatile such as the oxyhydroxyde CrO2(OH)2, leading to a low chromium/iron ratio in the scale. Both of the experimental methods and the characterisation techniques will be reviewed in this chapter.

3.2 Analysis of Gas Composition in the Furnace: a Case of Humidified Oxygen

The types of the oxide formed at high temperatures can be varied depending on many parameters particularly the atmospheres the metal is exposed to. The analysis of gas composition in the atmosphere is important to predict thermodynamic stability and oxidation kinetics of the metal. One of the gases that could significantly affect the oxide formation and will be treated here is water vapour. This gas is typically present in the ambient atmosphere of the industrial processes [22-27] or in the cathode side of the solid oxide fuel cells [28-31]. The experimental set-up to produce the gas mixture containing water vapour can be done by purging oxygen into a flask containing water, giving oxygen with water vapour at the outlet of the flask. The water vapour content is controlled by water temperature in the flask according to the Clausius-Clapeyron equation [32]:

ln(pH2O(g),TpH2O(g))=−ΔHl→vR(1T−1373.15)(3.1)

where pH2O(g),T is the partial pressure of water vapour at temperature T,pH2O(g) is the partial pressure of water vapour at 100∘C(1bar),ΔHl→v is the enthalpy of vaporisation (Jmol−1),R is the ideal gas constant ( 8.314 J mol−1 K−1 ) and T is the considered temperature in K .

At ambient temperatures of 15−20∘C, the water vapour pressure might be calculated by assuming that ΔHl→v is constant with the value of 40893 J mol−1 [33]. In this case, the water vapour pressure is in the range of 0.020−0.027 bar. In the condition where water vapour exists in the atmosphere with other gases i.e. H2 and O2, these three gases can be in equilibrium according to Reaction 3.2.

2H2O(g)↔2H2( g)+O2( g)(3.2)

As an example for the case when the oxidising atmosphere contains A% of H2O in argon inert gas, Table 3.1 shows the calculation method to determine a partial pressure of each gas in the atmosphere. From this table, A mol of water vapour is in the initial state. When the reaction occurs, 2x mol of A is consumed while 2x mol of hydrogen gas and x mol of oxygen gas are formed. This gives at equilibrium A−2x mol of water vapour, 2x mol of hydrogen gas, and x mol of oxygen, giving the total number of mole of 100+x. If the total pressure is 1 bar, a mole fraction equals the partial pressure. The partial pressure of each gas can then be expressed as follows:

pH2O=A−2x100+xpH2=2x100+x

(3.3)(3.4)

pO2=x100+x(3.5)

where x can be found out from the equilibrium constant of Reaction 3.2 which is

K=exp(−ΔrG3.2∘RT)=(pH22·pO2pH2O2)eq =(2x100+x)2(x100+x)(A−2x100+x)2(3.6)

By assuming that x≪A and x≪100, it can be obtained that

x=(25A2K)1/3(3.7)

Table 3.1 Determination of partial pressure of gases in water dissociation reaction.

| 2H₂O(g) = 2H2(g) + O2(g) | Ar | Total mole
Initial number of mole | A | 0 | 0 | 100−A | 100
Reacted number of mole | −2x | 2x | x |  |
Final number of mole at equilibrium | A−2x | 2x | x | 100−A | 100+x
Mole fraction | A−2x ──── 100+x | 2x ──── 100+x | x ──── 100+x |  |

For example, at 900∘C in argon with 20% of water vapour, by using [32]

ΔrG3.2∘=117800−26.2Tcal(3.8)

it can be obtained that x=8.45×10−5 and pO2=8.45×10−7 bar [34]. It is noted that, in pure water vapour, the oxygen dissociation from water vapour is essentially small. Thus the role of water vapour on steel oxidation is not only mainly from the oxygen dissociated but from other mechanisms which will be explained in Chapter 5.

Furthermore, after the oxidation test, mass gain or scale thickness can be measured. Such mass gain is in fact from the oxygen gas consumed by the oxidation reaction. Once the mass gain is measured, the oxide thickness can be calculated using Eq. 1.28. The calculated thickness can be used to compare with the measured oxide thickness. For example, when AISI 441 stainless steel is oxidised in air at 850∘C for 1 h , mass gain was measured and calculated to be 0.03mgcm−2 and then oxide thickness was calculated to be 0.18μ m by assuming that the oxide formed is chromia and has the density of 5.21 g cm−3 [35]. The measured oxide thickness gives the value of 0.16μ m. This case shows that the scale thickness calculated from the measured mass gain is similar to the thickness measured directly from the sample.

3.3 Chromium Species Volatilisation Measurement

In order to measure chromium species volatilisation from the alloys, the experimental set-up as shown in Fig. 3.1 is invented. The oxygen tank is connected to the flow meter which is continuously connected to the humidifier set. The role of humidifier set is to adjust the humidity in the gas. After this section, the humidified O2 is flowed to the quartz column installed inside the furnace. The specimen is hung in the quartz column at the position " S ". At the end of the column, the condenser is installed and immersed in the thermos cryostat where the temperature is set to be 7∘C in order to condense the chromium volatile species from the specimen. Before venting the outlet gas to the atmosphere, the gas is flowed through the water-containing Erlenmeyer flask which is the second bubbler of the setup for making sure that all chromium volatile species are captured into the solution. After the experiment, the temperature is cooled down to room temperature. The column is cleaned by 0.1 M HCl . The cleaning acid solution together with the solutions collected from the condenser and the second bubbler are then analysed by inductively coupled plasma (ICP) to measure the amount of chromium volatile from the specimen.

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Fig. 3.1. Schematic sketch of chromium volatilisation experimental set-up. Reproduced from W. Wongpromrat et al., Corros. Sci. 106 (2016) 172-178 [28].

Fig. 3. 1. Schematic sketch of chromium volatilisation experimental set-up. Reproduced from W. Wongpromrat et al., Corros. Sci. 106 (2016) 172-178 [28].

3.4 Physico-Chemical Characterisation

Physico-chemical properties of materials can be characterised by various techniques. The sources that are used in these techniques can be photon, electron, or ion. Microscopy is a basic technique for physical property observation that the image is usually performed by photon (visible light or X-ray or laser) or electron as optical microscopy or electron microscopy, respectively. Apart from benefits in image performing, when photons or electrons or ions interact with matter, the various signals can be produced.

In the case of photon, X-ray is broadly used to be the source for characterisation techniques. When X-ray is illuminated on matter, the photoelectron and secondary X-ray can be generated. For X-ray photoelectron spectroscopy (XPS) technique, the binding energy of photoelectron, which is the fingerprint of matter, can be evaluated to identify element and its chemical environment. X-ray diffraction (XRD) is one of the techniques that use X-ray as a source and analyse diffracted X-ray to generally identify a phase of matter by finding the lattice spacing and crystal structure. Laser source,
which is a monochromatic and intense light, is also typically used as in Raman spectroscopy where matter is characterised from the frequency change of light.

In the case of electron, the interaction of electron beam on matter might generate secondary electron, backscattered electron, Auger electron, and X-ray. These kinds of signals are useful for characterisation. Scanning electron microscopy (SEM) uses secondary and backscattered electrons to perform image that reveals surface topography. Auger electron, which is produced at a very near surface, is the important one of surface analysis techniques, named Auger spectroscopy (AES). The energy dispersive X-ray spectroscopy (EDS or EDX), which is typically equipped with SEM, utilises the characteristic X-ray for elemental analysis technique. Besides, some electrons that can pass through the thin matter are utilised to find many characteristics by transmission electron microscopy (TEM). Nowadays, many techniques with electron source are developed such as scanning transmission electron microscopy (STEM) which combines the principles of SEM and TEM. Moreover, some techniques are developed to support and enhance performance of other techniques such as focused ion-beam (FIB) instrument that is equipped with SEM.

In the case of ion, the energetic ion beam focused on matter can generate sputtered atoms or molecules. The charge of these species might be neutral, negative or positive, according to momentum transfer from the primary ion to the sample surface. In secondary ion mass spectroscopy (SIMS), the intensity of the obtained secondary ion is measured and reported.

Most of the techniques mentioned above will be presented in more detail concerning the principles, applications, advantages, disadvantages and limitations in the following parts.

3.4.1 Optical Microscopy

Microscopy is a technique for image performing. The technique is generally used to observe morphology of materials. The typical microscopes can be categorised by sources as optical microscope and electron microscope. Magnification and resolution of the image obtained from optical microscope is low, comparing with the image obtained from electron microscope. However, optical microscope is still useful and widely used because the operation is simple and fast. In addition, the optical microscope is inexpensive and affordable.

Optical microscopy or OM is generally referred to an optical instrument containing lenses that produce an enlarged image of an object placed in the focal plane of the lenses using light source. Typically, visible light with 400 to 700 nm in wavelength is used to study the microstructure of materials, leading to resolution limitation in submicron scale. Magnification of the image is also limited at around 1500 times. Apart from morphology, optical microscopes can be used to determine other characteristics of materials, depending on model or set-up, such as phase distribution, transparency or opacity, colour, refractive indices, dispersion of refractive indices, crystal system, degree of crystallinity, and polymorphism.

The conventional optical microscope could be operated in two modes as transmission and reflection mode, depending on the path of light. Transmission is suitable for the transparent sample, whereas reflection is suitable for the opaque sample. Hence, the use of reflective light is appropriate for metallurgy and the oxide grown on surface of alloys studies. The image colours of the reflected light microscope depend on amounts of collected light reflected from surface. If the area reflects more light to objective lenses, a lighter tone is presented on image.

Fig. 3.2 illustrates the example of optical micrographs obtained from the 2205 duplex stainless steel surface before and after oxidation at 850∘C under water vapour during 5 min . Since the duplex stainless steel consists of two phases of ferrite and austenite, the micrographs reveal the different phases with some defects in structure as twin. Moreover, in Fig. 3.2(b), micrograph illustrates the different oxide morphology between oxides grown on ferritic phase (dark area) and on austenitic phase (light area). However the technique still has a limit. It cannot provide the perfect oxide morphology image because one shot of image is obtained from only one focal plane, resulting in some blur area in the image when the high magnification image is performed as illustrated in Fig. 3.3. The light microscope previously mentioned is generally used in high temperature oxidation
and is sometimes specifically called a bright field microscope. However, other light microscopes for the other particular functions are evolved such as a dark field microscope, ultraviolet microscope, fluorescence microscope, and phase-contrast microscope.

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Fig. 3.2. Optical micrographs of the 2205 duplex stainless steel surface (a) before and (b) after oxidation at 850∘C under water vapour in nitrogen gas during 5 min [36].

Fig. 3. 2. Optical micrographs of the 2205 duplex stainless steel surface (a) before and (b) after oxidation at 850 ∘ C under water vapour in nitrogen gas during 5 min [36].

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Fig. 3.3. Optical micrograph of the AISI 304 austenitic stainless steel oxidised at 800∘C in laboratory air during 25 h [36].

Fig. 3. 3. Optical micrograph of the AISI 304 austenitic stainless steel oxidised at 800 ∘ C in laboratory air during 25 h [36].

3.4.2 Scanning Electron Microscopy

As already mentioned in the previous part, electron microscopy is one of microscopy techniques that the images are performed by electron source. The technique is also mainly categorised in scanning electron microscopy (SEM) and transmission electron microscopy (TEM). Fig. 3.4 shows signals obtained from materials interacted with incident electron beam. The signals are achieved from both quasi-elastic and inelastic events. For the quasi-elastic event that the trajectory of incident electron is affected while the kinetic energy of incident electron is not affected, some incident electrons can pass through materials either with or without diffraction that are called diffracted electrons or transmitted electrons respectively. The other incident electrons which cannot pass through material are retarded by the electromagnetic field of the nucleus and scattered back from the matter surface, giving backscattered electrons. For the inelastic event that the electron energy is transferred to the matter, the secondary electron, the characteristic X-ray and the Auger electron are generated with different phenomena.

Scanning electron microscopy mainly utilises signals of the secondary electron (SE) and the backscattered electron (BSE) that escape from the surface. The interaction volume of an incident electron beam on a sample in Fig. 3.5 illustrates that the secondary electron comes from the shallow depth and small volume compared with the backscattered electron because the energy of the secondary electron is very low compared with the energy of the backscattered electron. Therefore the secondary electron provides near-surface information with the small volume, leading to the higher image resolution compared with the backscattered electron image. Fig. 3.6 shows the high resolution secondary electron micrographs at different magnifications of 2205 duplex stainless steel oxidised at 850∘C under water vapour in nitrogen gas during 2 h . The different surface morphology is clearly revealed between the oxide grown on ferritic and austenitic phases. The electron micrographs obtained from secondary electrons and backscattered electrons of 2205 duplex stainless steel oxidised at 850∘C under hydrogen gas during 5 min is shown in Fig. 3.7. It is obviously seen that the secondary electron micrograph provides the better information about morphology.

However, the backscattered electron micrograph is still useful because the backscattered electron can provide the superior information about atomic number of matter. The backscattered electron yield obtained from the interaction between incident electron and material strongly depends on atomic number when the atomic number is less than 40 . In this range of atomic number, a higher amount of backscattered electron is emitted from a heavier element sample. In Fig. 3.7(b), the backscattered electron micrograph reveals phases of ferrite ( α ) and austenite ( γ ). The presence of nickel ( Ni ) element in austenite phase leads to a lighter zone at austenitic grain because the atomic number of nickel is higher than iron and chromium. Moreover, according to the ferrite phase, the darker zone on the top left area implies that the oxide composes of a higher chromium content, comparing to the lighter zone.

Therefore, scanning electron microscope provides not only topography and morphology of matter but also element distribution recognition. Secondary electron micrograph is more preferable for morphology observation, whereas backscattered electron micrograph is preferred when the phase contrast or phase distribution is needed.

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Fig. 3.4. Signals obtained from material interacted with electron beam. Redrawn and adapted from P. Sarrazin et al., Mechanisms of High Temperature Corrosion: A Kinetic Approach, Trans Tech Publications, Switzerland, 2008 [37].

Fig. 3. 4. Signals obtained from material interacted with electron beam. Redrawn and adapted from P. Sarrazin et al., Mechanisms of High Temperature Corrosion: A Kinetic Approach, Trans Tech Publications, Switzerland, 2008 [37].

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Fig. 3.5. Interaction volume of an incident electron beam on a sample. Redrawn and adapted from P. Sarrazin at al., Mechanisms of High Temperature Corrosion: A Kinetic Approach, Trans Tech Publications, Switzerland, 2008 [37].

Fig. 3. 5. Interaction volume of an incident electron beam on a sample. Redrawn and adapted from P. Sarrazin at al., Mechanisms of High Temperature Corrosion: A Kinetic Approach, Trans Tech Publications, Switzerland, 2008 [37].

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Fig. 3.6. Secondary electron micrographs of the 2205 duplex stainless steel oxidised at 850∘C under water vapour in nitrogen gas during 2 h (a) 250× and (b) 5000× magnification [36].

Fig. 3. 6. Secondary electron micrographs of the 2205 duplex stainless steel oxidised at 850 ∘ C under water vapour in nitrogen gas during 2 h (a) 250 × and (b) 5000 × magnification [36].

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Fig. 3.7. Electron micrographs obtained from (a) secondary electrons and (b) backscattered electrons of the 2205 duplex stainless steel oxidised at 850∘C under hydrogen gas during 5 min [36].

Fig. 3. 7. Electron micrographs obtained from (a) secondary electrons and (b) backscattered electrons of the 2205 duplex stainless steel oxidised at 850 ∘ C under hydrogen gas during 5 min [36].

3.4.3 Energy Dispersive X-ray Spectroscopy

Energy dispersive X-ray spectroscopy is generally known as EDX or EDS. This technique utilises the X-ray signal obtained from the interaction between incident electron and matter as presented in Fig. 3.4. Hence, the energy dispersive X-ray spectroscopy is often found as a supplementary technique combined with the scanning electron microscopy. The obtained X-ray radiation composes of continuous X-ray and characteristic X-ray as presented in Fig. 3.8(a). Only characteristic X-ray, which is emitted from electron transition between the different atomic energy levels as in Fig. 3.8(b), is used for the energy dispersive X-ray spectroscopy technique. Since energy of the characteristic X-ray is unique depending on element, the spectroscopy is used for elemental analysis and reported as the intensity versus the characteristic X-ray energy.

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Fig. 3.8. (a) X-ray radiation of copper target and (b) characteristic X-ray emission. Redrawn and adapted from B.D. Cullity, S.R. Stock, Elements of X-Ray Diffraction, third ed., Prentice Hall, USA, 2001 [38].

Fig. 3. 8. (a) X-ray radiation of copper target and (b) characteristic X-ray emission. Redrawn and adapted from B.D. Cullity, S.R. Stock, Elements of X-Ray Diffraction, third ed., Prentice Hall, USA, 2001 [38].

Fig. 3.9 shows EDS spectrum obtained from austenitic phase of 2205 duplex stainless steel oxidised at 850∘C in water vapour for 60 min . The pattern shows the presence of iron (Fe), chromium (Cr), nickel (Ni), manganese (Mn), molybdenum (Mo) and oxygen (O) elements. Since the analysis volume of EDS is not small, the analysis signal comes from not only surface oxide but also the bulk alloy, allowing the presence of the signals of nickel and molybdenum. In addition, the quantitative analysis of elements shown in Fig. 3.9 can be quantified in terms of weight percent (wt.%) and atomic percent (at.%) as reported in Table 3.2. The technique typically has the limitation to detect the elements with atomic number less than 11. Because oxygen is a constituent of the oxide and has the atomic number of 8, the quantitative analysis is then unfavorable for high temperature oxidation field. However, the technique provides the idea about the elemental components in the oxide phases. Moreover, X-ray mapping of individual element can be performed to observe the element distribution of sample.

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Fig. 3.9. EDS spectrum obtained from austenitic phase of the 2205 duplex stainless steel oxidised at 850∘C in water vapour for 60 min [36].

Fig. 3. 9. EDS spectrum obtained from austenitic phase of the 2205 duplex stainless steel oxidised at 850 ∘ C in water vapour for 60 min [36].

Table 3.2 Quantitative analysis of EDS spectrum in Fig. 3.9 [36].

Element | O K | Mo L | Cr K | Mn K | Fe K | Ni K | Total
wt.% | 24.87 | 2.22 | 19.19 | 2.26 | 46.77 | 4.69 | 100.00
at.% | 53.51 | 0.80 | 12.70 | 1.42 | 28.83 | 2.75 | 100.00

3.4.4 Focused Ion Beam Coupled with Scanning Electron Microscope

The focused ion beam coupled with scanning electron microscope or FIB-SEM consists of 2 columns that are FIB and SEM columns. They are installed with the intercolumn angle of 54∘ as shown in Fig. 3.10. FIB is the perpendicular column to the specimen surface. Gallium ion ( Ga+) source installed inside FIB column is used for cutting, technically called milling, the localised area while SEM records the image of such area. However, the specimen can be damaged due to the milling process. Thus, before milling, the interesting surface needs to be protected usually by coating with carbon.

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Fig. 3.10. Schematic sketch of FIB-SEM. Redrawn and adapted from H. Iwai et. al., J. Power Sources, 195 (2010) 955-961 [39].

Fig. 3. 10. Schematic sketch of FIB-SEM. Redrawn and adapted from H. Iwai et. al., J. Power Sources, 195 (2010) 955-961 [39].

Normally, FIB-SEM has been used for many purposes in the study of high temperature oxidation of alloys. The first one is to make a thin lamella. Before characterising the specimen by TEM, thin lamella with the thickness of 100−150 nm is prepared by FIB in order to observe the transmission of electron. Moreover, the characterisation of alloys by SEM and EDX has one big problem concerning the too deep penetration depth of electron beam. For example, if the penetration depth of electron beam is about 10μ m but the depth of interested oxide layer is 5μ m, the characterisation obtained from SEM and EDX would be disturbed from phases under the interested oxide layer. In order to get rid of this effect, thin lamella from FIB may be prepared to limit the area of characterisation. The preparation of thin lamella is shown in Fig. 3.11 [40]. First, carbon is deposited on the surface in order to protect the surface from the milling process as shown in Fig. 3.11(a). After that, the FIB-boxes are made at both sides of the lamella as shown in Fig. 3.11(b). The right and left sides as well as beneath the lamella are cut (the red dash lines refer to the cutting lines) while the small connection bridge at the left side is still remained in order to maintain the lamella as shown in Figs. 3.11 (c) and (d). Later, the lamella is connected with a pin by carbon-deposition as illustrated in Fig. 3.11(e) and then the small bridge can now be cut as illustrated in Fig. 3.11(f). The lamella is moved and connected to the copper grid as shown in Figs. 3.11 (g) and (h), and then the pin is disconnected. Finally, the lamella is slenderised to be 100−150 nm thick. After this preparation, the thin lamella can be characterised by TEM, SEM or other techniques. The example of the advantage of thin lamella is shown in Fig. 3.12 showing the cartography and EDX localised analysis of AISI 441 after oxidation at 800∘C in 5%H2O−O2 for 24 h [40].

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Fig. 3.11. Thin lamella preparation from FIB. Reproduced and adapted from W. Wongpromrat, PhD Thesis, Université Grenoble Alpes, France, and King Mongkut's Institute of Technology Ladkrabang, Thailand, 2015 [40].

Fig. 3. 11. Thin lamella preparation from FIB. Reproduced and adapted from W. Wongpromrat, PhD Thesis, Université Grenoble Alpes, France, and King Mongkut's Institute of Technology Ladkrabang, Thailand, 2015 [40].

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Fig. 3.12. (a) Cartography and (b) EDX of the localised area of the nodule grown during oxidation at 800∘C in 5%H2O−O2 for 24 h . Reproduced from W. Wongpromrat, PhD Thesis, Université Grenoble Alpes, France, and King Mongkut's Institute of Technology Ladkrabang, Thailand, 2015 [40].

Fig. 3. 12. (a) Cartography and (b) EDX of the localised area of the nodule grown during oxidation at 800 ∘ C in 5 % H 2 O − O 2 for 24 h . Reproduced from W. Wongpromrat, PhD Thesis, Université Grenoble Alpes, France, and King Mongkut's Institute of Technology Ladkrabang, Thailand, 2015 [40].

The FIB-SEM technique can also be applied to observe the cross section of the small localised area where other techniques are unable. For example, in the recent work of Parry et al. [29], FIB was used to observe the revolution of the nodule grown on AISI 441 ferritic steel after oxidation at 800∘C in 5%H2O−O2 for 4,14 and 24 h as shown in Fig. 3.13. Furthermore, the FIB-SEM technique can also be used to reconstruct the area of interest. In this case, the FIB and SEM columns would work iterately in milling and capturing the images of the scoped area. All images were then combined in the software. The results attained from this process may be the revolution of the area of interest as shown in Figs. 3.13 and 3.14 and the reconstruction of phase(s) on the specimens as shown in Fig. 3.15.

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Fig. 3.13. FIB cross sections of the nodule grown on AISI 441after oxidation at 800∘C in 5%H2O−O2 for (a) 4 h , (b) 14 h and (c) 24 h . Reproduced from V. Parry et al., Corros. Sci. 141 (2018) 255-263 [29].

Fig. 3. 13. FIB cross sections of the nodule grown on AISI 441after oxidation at 800 ∘ C in 5 % H 2 O − O 2 for (a) 4 h , (b) 14 h and (c) 24 h . Reproduced from V. Parry et al., Corros. Sci. 141 (2018) 255-263 [29].

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Fig. 3.14. The revolution of the nodule extracted from stack of 330 SEM-images where Fig. (a), (b) and (c) correspond to image stack numbers of 1, 165 and 330 respectively. Reproduced from W. Wongpromrat et al., Mater. High Temp. 32 (2015) 22-27 [30].

Fig. 3. 14. The revolution of the nodule extracted from stack of 330 SEM-images where Fig. (a), (b) and (c) correspond to image stack numbers of 1, 165 and 330 respectively. Reproduced from W. Wongpromrat et al., Mater. High Temp. 32 (2015) 22-27 [30].

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Fig. 3.15. Reconstruction of niobium phases, including NbC and intermetallic niobium, above the metallic grain boundary of AISI 441 oxidised at 800∘C in O2−5%H2O for 24 h . Reproduced from W. Wongpromrat, PhD Thesis, Université Grenoble Alpes, France, and King Mongkut's Institute of Technology Ladkrabang, Thailand, 2015 [40].

Fig. 3. 15. Reconstruction of niobium phases, including NbC and intermetallic niobium, above the metallic grain boundary of AISI 441 oxidised at 800 ∘ C in O 2 − 5 % H 2 O for 24 h . Reproduced from W. Wongpromrat, PhD Thesis, Université Grenoble Alpes, France, and King Mongkut's Institute of Technology Ladkrabang, Thailand, 2015 [40].

3.4.5 X-ray Diffractometer

X-ray diffractometer or XRD is used to observe the crystalline information of the solid samples that can be used to identify the phase of matter such as oxide species. The diffraction pattern produced by X-ray through the lattice in a crystal is recorded and interpreted to determine the nature of the lattice [41]. When X-ray is applied to the crystal, the diffraction occurs on the lattice plane. The incident X-ray primarily interacts with electron in the atom. Some photons of X-ray are deflected from the incident direction. This diffraction phenomenon is illustrated in Fig. 3.16 and can be described by Bragg's law in Eq. 3.9 [42] i.e.

2dsinθ=nλ(3.9)

where d is the lattice spacing, θ the incident angle to specimen surface, n an integer ( 1,2,3,… ) and λ the wavelength of X-ray.

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Fig. 3.16. Diffraction of X-ray in crystalline structure at different incident angles. Redrawn and adapted from R. Abbaschian et al., Physical Metallurgy Principles, fourth ed., Cengage Learning Press, UK, 2009 [42].

Fig. 3. 16. Diffraction of X-ray in crystalline structure at different incident angles. Redrawn and adapted from R. Abbaschian et al., Physical Metallurgy Principles, fourth ed., Cengage Learning Press, UK, 2009 [42].

Due to the characteristics of crystal structure and lattice spacing of each compound, X-ray diffraction pattern is unique and can be analysed by scanning X-ray with various incident angles and observing the intensity of the diffracted X-ray at each angle. The data attained from XRD are the crystal structure of the unit cell and the orientation of the lattice plane. In the crystalline solid, the atoms are arranged periodically.

The crystal types can be classified through the lattice parameters. Table 3.3 shows the crystal systems in the unit cell.

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Fig. 3.17. Unit cell with axial lengths (a, b, c) and interaxial angles ( α,β,γ ). Redrawn and adapted from W.D. Callister, D.G. Rethwisch, Materials Science and Engineering, eighth ed., John Wiley & Sons, USA, 2011 [43].

Fig. 3. 17. Unit cell with axial lengths (a, b, c) and interaxial angles ( α , β , γ ). Redrawn and adapted from W.D. Callister, D.G. Rethwisch, Materials Science and Engineering, eighth ed., John Wiley & Sons, USA, 2011 [43].

Table 3.3 Classification of crystal systems. Redrawn and adapted from W.D. Callister, D.G. Rethwisch, Materials Science and Engineering, eighth ed., John Wiley & Sons, USA, 2011 [43].

Crystal system | Axial relationships | Interaxial angle | Unit cell Geometry
Cubic | a = b = c | α = β = γ = 90° | Cubic unit cell with equal edges a
Tetragonal | a = b ≠ c | α = β = γ = 90° | Tetragonal unit cell with square base a and height c
Orthorhombic | a ≠ b ≠ c | α = β = γ = 90° | Orthorhombic unit cell with three unequal edges a, b, c
Rhombohedral | a = b = c | α = β = γ ≠ 90° | Rhombohedral unit cell with equal edges a and equal non-90° angles α
Hexagonal | a = b ≠ c | α = β = 90° γ = 120° | Hexagonal unit cell with hexagonal base and height c
Monoclinic | a ≠ b ≠ c | α = β = 90° ≠ γ | Monoclinic unit cell with three unequal edges and one non-90° angle β
Triclinic | a ≠ b ≠ c | α ≠ β ≠ γ ≠ 90° | Triclinic unit cell with three unequal edges and three unequal non-90° angles α, β, γ

For the lattice plane in the cubic structure, three integers called Miller indices ( h,k,l ) are used to represent the planes and the directions in the crystal lattices. The examples of the lattice planes in cubic structure and their Miller indices are shown in Fig. 3.18.

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Fig. 3.18. The lattice planes in cubic structure and their Miller indices. Redrawn and adapted from R. Abbaschian et al., Physical Metallurgy Principles, fourth ed., Cengage Learning Press, UK, 2009 [42].

Fig. 3. 18. The lattice planes in cubic structure and their Miller indices. Redrawn and adapted from R. Abbaschian et al., Physical Metallurgy Principles, fourth ed., Cengage Learning Press, UK, 2009 [42].

After analysis, the peaks can be observed at different incident angles. These results can be interpreted by comparing the patterns obtained from the experiment with the standard database called "Joint Committee of Power Diffraction Standards (JCPDS)", currently renamed to be "International Centre for Diffraction Data (ICDD)". The example of XRD patterns obtained from the oxidised AISI 441 specimen is presented in Fig. 3.19. This figures shows the effect of oxidising gas, time, and temperature on thermal oxide phase where x and y is chromium mole fraction in Fe2−xCrxO3 and Mn3−yCryO4 respectively.

[picture]

Fig. 3.19. XRD patterns of AISI 441 specimens oxidised in air: (a) 650∘C,5 min; (b) 850∘C, 5 min ; (c) 850∘C,60 min; in water vapour: (d) 650∘C,5 min; (e) 850∘C,5 min; (f) 850∘C,60 min. Redrawn from A. Srisrual et al., Corros. Sci. 51 (2009) 562-568 [44].

Fig. 3. 19. XRD patterns of AISI 441 specimens oxidised in air: (a) 650 ∘ C , 5 m i n ; (b) 850 ∘ C , 5 min ; (c) 850 ∘ C , 60 m i n ; in water vapour: (d) 650 ∘ C , 5 m i n ; (e) 850 ∘ C , 5 m i n ; (f) 850 ∘ C , 60 m i n . Redrawn from A. Srisrual et al., Corros. Sci. 51 (2009) 562-568 [44].

3.4.6 Raman Spectroscopy

Raman spectroscopy is a vibrational spectroscopic technique based on inelastic scattering or the frequency change of monochromatic light typically delivered by laser. The incident monochromatic light interacts with sample and then a spontaneous scattering of incident light is re-emitted. Frequency of the re-emitted photons is usually shifted up or down compared with original frequency, which is named the Raman scattering or Raman effect, following the name of the first observer, C.V. Raman, an Indian physicist. The intensity of Raman scattering for each frequency shift is recorded and plotted. The shift provides information about vibrational and rotational transitions in molecules that can indicate the structure, orientation, and chemical composition of the interested molecule.

The incident light or photon can excite molecule of matter from ground state in a rotational or vibrational energy state to virtual energy state, and then it emits or scatters a photon and returns to a rotational or vibrational state when the molecule relaxes. From the phenomena, the light scattering can be categorised in three types as presented in Fig. 3.20. Rayleigh scattering is an elastic scattering that the scattering light has the same frequency as the exciting light. Stokes scattering and anti-Stokes scattering are an inelastic scattering. In the case of the Stokes scattering, the final energy state is higher than the initial energy state. Hence, the scattering light shifts to the lower frequency comparing with the exciting light. In the case of the anti-Stokes scattering, the final energy state is lower than the initial energy state. Thus the scattered light shifts to the higher frequency comparing with the exciting light. In general, the total amount of inelastic Stokes and anti-Stokes radiation is small comparing with the elastic Rayleigh radiation. However, only inelastic radiation, typically Stokes radiation, is used in Raman spectroscopy.

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Fig. 3.20. Rayleigh, Stokes and anti-Stokes scattering phenomena. Redrawn and adapted from J.E. Pemberton, A.L. Guy, Raman spectroscopy, in: ASM Handbook Committee (Ed.), ASM Handbook Vol. 10, ninth ed., ASM International, USA, 1986, pp. 261-286 [45].

Fig. 3. 20. Rayleigh, Stokes and anti-Stokes scattering phenomena. Redrawn and adapted from J.E. Pemberton, A.L. Guy, Raman spectroscopy, in: ASM Handbook Committee (Ed.), ASM Handbook Vol. 10, ninth ed., ASM International, USA, 1986, pp. 261-286 [45].

Raman shift is generally presented in wavenumber or in the unit of reciprocal length ( cm−1 ). The correlation between spectral wavelength and Raman shift is as the following equation:

Δw=(1λ0−1λ1)(3.10)

where Δw is the Raman shift, λ0 and λ1 are the exciting wavelength and the Raman scattering wavelength respectively.

Raman spectroscopy is widely used for material identification because the technique is easy to use and sample preparation is not required. The related technique is infrared spectroscopy that is less sensitive to surface species and difficult to analyse an aqueous system, comparing to Raman spectroscopy. Fig. 3.21 presents Raman spectra recorded from the AISI 441 ferritic stainless steel oxidised in air and in water vapour. The spectra reveal the identical peaks of haematite ( Fe2O3 ), chromia (Cr2O3), corundum solid solution (Fe2−xCrxO3), and spinel solid solution (Mn3−yCryO4) [44]. Peaks of chromia grown under water vapour are always higher and sharper than those of films grown in air. Haematite decreases while solid solution and chromia increase when temperature increases both in air and in oxidising conditions containing water vapour.

In addition, Raman spectroscopy can be performed at microscopic scale as called micro-Raman spectroscopy. Fig. 3.22 shows Raman spectra recorded from austenite and ferrite phases of the oxidised 2205 duplex stainless steel. Both spectra show the characteristic peaks of haematite, chromia and corundum solid solution. The peak of corundum solid solution is shifted from 668 cm−1 (austenitic grains) to 674 cm−1 (ferritic grains) indicating that the solid solution oxide with a lower x-value is formed on austenite [46].

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Fig. 3.21. Raman spectra of AISI 441 ferritic stainless steel oxidised in air: (a) 650∘C,5 min; (b) 850∘C,5 min; (c) 850∘C,60 min; in water vapour: (d) 650∘C,5 min; (e) 850∘C,5 min; (f) 850∘C,60 min. Redrawn and adapted from A. Srisrual et al., Corros. Sci. 51 (2009) 562-568 [44].

Fig. 3. 21. Raman spectra of AISI 441 ferritic stainless steel oxidised in air: (a) 650 ∘ C , 5 m i n ; (b) 850 ∘ C , 5 m i n ; (c) 850 ∘ C , 60 m i n ; in water vapour: (d) 650 ∘ C , 5 m i n ; (e) 850 ∘ C , 5 m i n ; (f) 850 ∘ C , 60 m i n . Redrawn and adapted from A. Srisrual et al., Corros. Sci. 51 (2009) 562-568 [44].

[picture]

Fig. 3.22. Raman spectra recorded from individual austenitic (lower curve) and ferritic (upper curve) grains of the 2205 duplex stainless steel oxidised at 650∘C in water vapour atmosphere for 5 min . Redrawn and adapted from A. Srisrual et al., Mater. High Temp. 28 (2011) 349-354 [46].

Fig. 3. 22. Raman spectra recorded from individual austenitic (lower curve) and ferritic (upper curve) grains of the 2205 duplex stainless steel oxidised at 650 ∘ C in water vapour atmosphere for 5 min . Redrawn and adapted from A. Srisrual et al., Mater. High Temp. 28 (2011) 349-354 [46].

As the micro-Raman spectroscopy can be locally performed on individual area [47], Fig. 3.23(a) reports Raman spectra collected from different areas of the oxidised Inconel 690 alloy. Both of Cr2O3 and NiFe2O4 identical shifts are found on grain and grain boundary. Cr2O3 is a main oxide at grain boundary, whereas the presence of NiFe2O4 is more evident at grain area and nodules. Fig. 3.23(b) illustrates Raman mapping of Cr2O3 and NiFe2O4. The Cr2O3 covers the whole surface but the oxide intensely presents at grain boundary. NiFe2O4 clearly presents at the grain area. Moreover, the different oxide distribution over the grain between black and white areas should be affected from the different orientation, causing by twin defect in the alloy.

Even though the Raman spectroscopy is very useful in high temperature oxidation area, the characterisation should be carefully performed with a suitable laser power to avoid sample deterioration especially with a dark sample. Moreover, fluorescence effect can be occurred and can affect the interpretation of the results.

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Fig. 3.23. (a) Raman shift at white grain (WG), black grain (BG) and grain boundary (GB), and (b) Raman mapping obtained from selected area of Inconel 690 oxidised at 900∘C in air for 21 h . Redrawn and adapted from A. Srisrual, PhD Thesis, University of Grenoble, France, 2013 [47].

Fig. 3. 23. (a) Raman shift at white grain (WG), black grain (BG) and grain boundary (GB), and (b) Raman mapping obtained from selected area of Inconel 690 oxidised at 900 ∘ C in air for 21 h . Redrawn and adapted from A. Srisrual, PhD Thesis, University of Grenoble, France, 2013 [47].

3.4.7 X-ray Photoelectron Spectroscopy

X-ray photoelectron spectroscopy (XPS), which is sometimes known as Electron Spectroscopy for Chemical Analysis (ESCA), is used for observing the element composition, chemical stoichiometry, chemical state and electronic state of the elements existing in the materials [41]. The advantage of XPS is the characterisation of the near surface about 10 nm depth which is much thinner than other techniques. Fig. 3.24 shows the principle of XPS. A high energy photon (X-ray) strikes the inner-shell electron in the atom of material at the surface, and then the escaped innershell electron becomes photoelectron meanwhile the hole is created as shown in Figs. 3.24 (a) and (b) respectively. To stabilise the atom, electron from the other shell drops into the inner-shell hole. In the meantime, the excess energy is emitted as photon or characteristic X-ray. If the photon is absorbed by electron and electron has enough energy, the characteristic electron will escape from atom in the name of Auger electron as in Fig. 3.24(c).

Since the binding energy is characteristics of any element, it is used to identify the element. The binding energy (B.E.) can be evaluated from photon energy ( hν ) and kinetic energy of photoelectron (K.E.) following Eq. 3.11 when a work function of sample is ignored. Kinetic energy of photoelectron is generally measured by cylindrical mirror analyser or concentric hemispherical analyser. The result of XPS is reported as a spectrum line. The example of XPS spectrum is shown in Fig. 3.25. The quantitative analysis can be achieved by calculating the ratio of the curve area divided by the sensibility factor of that element as shown in Eq. 3.12 [40]. In this equation, SA and SB are respectively the curve areas of A and B , while αA and αB are the sensitivity factors of A and B respectively.
B.E. = hv - K.E.

 Ratio of A to B=(SA/αA)(SB/αB)(3.12)

[picture]

Fig. 3.24. Principle of XPS. Redrawn and adapted from J.B. Lumsden, X-ray photoelectron spectroscopy, in: ASM Handbook Committee (Ed.), ASM Handbook Vol. 10, ninth ed., ASM International, USA, 1986, pp. 1025-1056. [48].

Fig. 3. 24. Principle of XPS. Redrawn and adapted from J.B. Lumsden, X-ray photoelectron spectroscopy, in: ASM Handbook Committee (Ed.), ASM Handbook Vol. 10, ninth ed., ASM International, USA, 1986, pp. 1025-1056. [48].

[picture]

Fig. 3.25. XPS spectrum of AISI 441 oxidised in water vapour at 850∘C for 5 min [36].

Fig. 3. 25. XPS spectrum of AISI 441 oxidised in water vapour at 850 ∘ C for 5 min [36].

3.5 Photoelectrochemical Characterisation

Photoelectrochemistry or PEC is the technique for measuring and interpreting electrical currents generated by illuminating a semiconductor material in an electrolyte. In the last decades, photoelectrochemistry has been used in thin solid film applications such as photovoltaic cell and corrosion studies. This technique could provide the information on electronic structure and indirectly (through the optical band gap values) chemical composition of passive films in situ and under controlled potential in long-lasting experiments. It could also provide information on mechanisms of generation (recombination effect) and transport of photocarrier in amorphous material. These informations are utilised to find the physico-chemical property of oxide film. The PEC could also give information on kinetics of growth under illumination of photoconducting films and interference effect in absorbing material at constant growth rate.

The main components of typical photoelectrochemical set-up compose of lamp, monochromator, chopper, electrochemical cell, potentiostat, lock-in amplifier, and computer. The lampmonochromator system provides a variable wavelength in a wide range of photon energy and photocurrents will be recorded according to photon energies as called photocurrent energy spectra. However, the system may include a laser source at constant wavelength when the imaging study is preferred.

The Science et Ingénierie des Matériaux et Procédés (SIMaP) laboratory, Université Grenoble Alpes, France, has developed the photoelectrochemical set-up that has evolved from macroscopic, microscopic, and mesoscopic scale of characterisation, as named macroscopic PEC, microscopic PEC, and mesoscopic PEC, respectively. Both macroscopic and microscopic PEC are typical modes of photoelectrochemical technique that are performed to obtain photocurrent energy and potential spectra and the photocurrent image, respectively. Mesoscopic PEC is the unique mode that is firstly developed and proved for the local photocurrent energy spectra with 30μ m spatial resolution [46]. Generally, the technique is used to find the band energy gap and semiconducting type of thermal oxides. On the other hand, the thermal oxide phases can be identified when the band energy gap of thermal oxide is known. However, each mode of PEC is utilised for the different objectives as will be presented in the following.

In the three operating modes of the photoelectrochemical set-up, some components are shared such as the photoelectrochemical cell, the potentiostat, the lock-in amplifier, and the mechanical light modulator (chopper and optical path device). The electrochemical cell consists of three electrodes equipped with a flat quartz optical window, which allows light illuminated on sample surface. The oxidised sample should be used as a working electrode and the oxide at the back side must be removed to provide the electrical contact. A platinum foil is used as a counter electrode. An Hg/Hg2SO4 mercury sulphate electrode (MSE) is used as a reference electrode ( +0.650 V/SHE ). A sodium sulphate solution with 0.1 M at pH8.0±0.1 is typically used as the electrolyte in the cell. Potentiostat is used to control the electrochemical cell where the voltage is proportional to overall electrochemical current issued from the matter interface. The electrochemical current is collected by a lock-in amplifier. Chopper and lock-in amplifier are necessary components for extracting a weak photocurrent signal in dark current (modulus and phase shift) from the collected electrochemical current.

3.5.1 Macroscopic Photoelectrochemical Characterisation

Macroscopic photoelectrochemistry chacterisation is generally performed with a white light and monochromator system [20,21]. The technique provides information about semiconducting type and bandgap energy, evaluating from photovoltammograms and photocurrent energy spectra. Fig. 3.26 presents the schematic diagram of macroscopic PEC set-up at SIMaP laboratory [47]. A xenon arc bulb connected to a monochromator is used as a light source to provide variable wavelength photons with a minimum spectral width of ±1 nm. The characterised area is around 0.8 cm2 while the whole sample surface is illuminated.

[picture]

Fig. 3.26. Schematic diagram of the macroscopic photoelectrochemical set-up at SIMaP laboratory. Redrawn from A. Srisrual, PhD Thesis, University of Grenoble, France, 2013 [47].

Fig. 3. 26. Schematic diagram of the macroscopic photoelectrochemical set-up at SIMaP laboratory. Redrawn from A. Srisrual, PhD Thesis, University of Grenoble, France, 2013 [47].

Photovoltammogram is the curve between photocurrent and applied potential. It is usually plotted at the different selected constant photon energy provided by monochromator. Applied potential is varied and controlled by potentiostat. The as-measured photocurrent, IPH, is recorded according to the applied potential. Semiconducting type can be evaluated from the change of the curve. If the modulus of the photocurrent increases when the applied potential is increased, the oxide exhibits an n-type semiconducting behaviour, as demonstrated in Fig. 3.27(a). On the other hand, if the modulus of the photocurrent decreases when the applied potential is increased, the sample exhibits a p-type semiconducting behaviour, as demonstrated in Fig. 3.27(c). In addition, Fig. 3.27(b) indicates the presence of at least two components in the thermal scale that one being n-type as the curve is recorded at 2.41 eV whereas another one being more or less insulating as the V-shaped part is recorded at 3.53 eV .

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Fig. 3.27. Plots of the typically measured photocurrent amplitude as a function of the applied potential obtained at hν=2.41eV (dashed line) and 3.53 eV (solid line) on the whole surface of the 2205 duplex stainless steel oxidised in water vapour (a) at 650∘C for 5 min , (b) at 850∘C for 60 min , and in oxygen (c) at 850∘C for 60 min . Redrawn from A. Srisrual et al., Oxid. Met. 79 (2013) 337-347 [49].

Fig. 3. 27. Plots of the typically measured photocurrent amplitude as a function of the applied potential obtained at h ν = 2.41 e V (dashed line) and 3.53 eV (solid line) on the whole surface of the 2205 duplex stainless steel oxidised in water vapour (a) at 650 ∘ C for 5 min , (b) at 850 ∘ C for 60 min , and in oxygen (c) at 850 ∘ C for 60 min . Redrawn from A. Srisrual et al., Oxid. Met. 79 (2013) 337-347 [49].

Photocurrent energy spectrum or the curve plotting the photocurrent versus energy is performed at variable wavelength provided by monochromator. The spectrum is typically plotted at different applied potentials depending on the presence of oxides, as illustrated in Figs. 3.28 and 3.29. Since the photon flux from the light source is different and depends on the wavelength, the as-measured photocurrent must be corrected to the calculated photocurrent, JPH, by

JPH=IPHΦN( hν)(3.13)

where ΦN(hν) is the normalised photon flux at each wavelength. Fig. 3.29 illustrates the photocurrent energy spectra obtained from the same sample in Fig. 3.27. Each energy spectrum reveals the combination of the photocurrent contributions, implying the presence of different oxide types with different bandgap energies. Bandgap energy of semiconductor, Eg, can be determined by the simple equation of Gärtner-Butler model in Eq. 3.14, i.e. for each sample

(η×E)1/2=C×(E−Eg)(3.14)

where η is the quantum yield of the photocurrent that equals JPH,E is the photon energy, and C is a constant. Fig. 3.30 shows the linear transformed photocurrent energy spectra of the curves in Fig. 3.28 performed at −250mV/MSE and at −700mV/MSE. The transforms in Figs. 3.30 (a) and (b) indicate the bandgap energy of the first contribution at around 2.0 eV corresponding to the bandgap energy of haematite [50,51], and the bandgap energy of the second contribution at around 2.6 eV which corresponds to the bandgap energy of the rhombohedral solid solution [51].

[picture]

Fig. 3.28. Photocurrent energy spectra at two potentials for AISI 441 ferritic stainless steel oxidised at 650∘C in air for 5 min . Redrawn from A. Srisrual et al., Corros. Sci. 51 (2009) 562-568 [44].

Fig. 3. 28. Photocurrent energy spectra at two potentials for AISI 441 ferritic stainless steel oxidised at 650 ∘ C in air for 5 min . Redrawn from A. Srisrual et al., Corros. Sci. 51 (2009) 562-568 [44].

[picture]

Fig. 3.29. Photocurrent energy spectra at several applied potentials obtained from the whole surface of the 2205 duplex stainless steel oxidised (a) at 650∘C in water vapour for 5 min , (b) at 850∘C in water vapour for 60 min , and (c) at 850∘C in oxygen for 60 min . Redrawn from A. Srisrual et al., Oxid. Met. 79 (2013) 337-347 [49].

Fig. 3. 29. Photocurrent energy spectra at several applied potentials obtained from the whole surface of the 2205 duplex stainless steel oxidised (a) at 650 ∘ C in water vapour for 5 min , (b) at 850 ∘ C in water vapour for 60 min , and (c) at 850 ∘ C in oxygen for 60 min . Redrawn from A. Srisrual et al., Oxid. Met. 79 (2013) 337-347 [49].

[picture]

Fig. 3.30. Linear transforms of curves in Fig. 3.28; (a) −250mV/MSE, (b) −700mV/MSE. Redrawn from A. Srisrual et al., Corros. Sci. 51 (2009) 562-568 [44].

Fig. 3. 30. Linear transforms of curves in Fig. 3.28; (a) − 250 m V / M S E , (b) − 700 m V / M S E . Redrawn from A. Srisrual et al., Corros. Sci. 51 (2009) 562-568 [44].

3.5.2 Microscopic Photoelectrochemical Characterisation

Microscopic photoelectrochemical is generally performed with the laser source to obtain a photoelectrochemical image that provides the information about the spatial distribution of oxide. Fig. 3.31 presents the schematic diagram of the microscopic PEC set-up at SIMaP laboratory [47]. A monochromatic light from argon laser is used instead of a xenon lamp in macroscopic PEC. An inverted microscope is modified by the 50/50 beam splitter to allow the positioning of the focused beam on the area selected on the optical image provided by CCD camera. A positioning system with a high resolution in XY axes motorised microscope stage is mounted. The laser is used for raster scanning and photocurrent (modulus and phase shift) is collected to perform the image. Spatial resolution of the image can be achieved 1μ m by this set-up.

[picture]

Fig. 3.31. Schematic diagram of the microscopic photoelectrochemical set-up at SIMaP laboratory. Redrawn from A. Srisrual, PhD Thesis, University of Grenoble, France, 2013 [47].

Fig. 3. 31. Schematic diagram of the microscopic photoelectrochemical set-up at SIMaP laboratory. Redrawn from A. Srisrual, PhD Thesis, University of Grenoble, France, 2013 [47].

Fig. 3.32 shows a set of images obtained from the 2205 duplex stainless steel oxidised in water vapour and in oxygen at 850∘C for 60 min . The characterisation was performed at the same applied potential. The 2.60 eV photon energy contributes Fe2O3 and iron-rich Fe2−xCrxO3, whereas the 3.53 eV photon energy contributes Fe2O3,Fe2−xCrxO3, and Cr2O3.

[picture]

Fig. 3.32. Image set of the 2205 duplex stainless steel oxidised at 850∘C for 60 min in water vapour (top images) and in oxygen (bottom images): (a) and (e) optical image; (b) and (f) PEC image at hν=2.60eV; (c) and (g) PEC image at hν=3.53eV; and (d) and (h) processed image. Reproduced and adapted from A. Srisrual et al., Oxid. Met. 79 (2013) 337-347 [49].

Fig. 3. 32. Image set of the 2205 duplex stainless steel oxidised at 850 ∘ C for 60 min in water vapour (top images) and in oxygen (bottom images): (a) and (e) optical image; (b) and (f) PEC image at h ν = 2.60 e V ; (c) and (g) PEC image at h ν = 3.53 e V ; and (d) and (h) processed image. Reproduced and adapted from A. Srisrual et al., Oxid. Met. 79 (2013) 337-347 [49].

Both in Figs. 3.32 (b) and (f), the ferrite area is darker than the austenite area, indicating the higher Fe2O3 and Fe2−xCrxO3 on the austenite. To see Cr2O3 fraction in the scale, the processed images in Figs. 3.32 (d) and (h) are constructed by dividing, point by point, the photocurrent measured at 2.60 eV by that measured at 3.53 eV for each sample. In the case of water vapour oxidation giving the image shown in Fig. 3.32(d), the Cr2O3 fraction is higher on ferrite which
should be from the higher chromium content in ferritic phase and the faster chromium diffusion rate in ferritic phase [49]. In the austenite phase, iron content is high, inducing higher iron diffusion in the corundum oxide and therefore reducing the chromia fraction [49]. In the case of the oxidation in oxygen giving the image shown in Fig. 3.32(h), the higher oxygen partial pressure plays a negative role and opposite results are observed as the higher Cr2O3 fraction occurred on the austenitic phase.

3.5.3 Mesoscopic Photoelectrochemical Characterisation

Mesoscopic photoelectrochemical is developed for recording the photovoltammograms and photocurrent energy spectra on a specific microscopic area, thus promoting the superior performance of photoelectrochemical technique. Fig. 3.33 presents the schematic diagram of the mesoscopic photoelectrochemical set-up. It is similar to the microscopic photoelectrochemical setup in Fig. 3.31, except the part of the light source unit that xenon arc bulb and monochromator system as in macroscopic PEC are used instead of argon laser. Moreover, the optical path is specially designed to obtain a good alignment and to allow an intense micro-focusing light on the interesting area. Spatial resolution of the technique can achieve 30μ m according to the spot size of the focussed light.

[picture]

Fig. 3.33. Schematic diagram of the mesoscopic photoelectrochemical set-up at SIMaP laboratory. Redrawn from A. Srisrual, PhD Thesis, University of Grenoble, France, 2013 [47].

Fig. 3. 33. Schematic diagram of the mesoscopic photoelectrochemical set-up at SIMaP laboratory. Redrawn from A. Srisrual, PhD Thesis, University of Grenoble, France, 2013 [47].

Fig. 3.34 shows the typical photovoltammogramms obtained at 3.53 eV photon energy on the whole surface of the oxidised 2205 duplex stainless steel sample and grains of austenite and ferrite oxidised at 850∘C for 60 min in water vapour and in oxygen. Photon energy at 3.53 eV is higher than the bandgap energies of Fe2O3,Fe2−xCrxO3, and Cr2O3. As a result, the recorded photocurrents can be obtained from all of these oxides. The photovoltammograms in Figs. 3.34 (a) and (b) report the effect of oxygen activity on semiconducting type of thermal oxide. It is seen that the oxides act as insulator and p-type behaviour in water vapour and in oxygen atmosphere respectively. Moreover, Fig. 3.34(a) shows the flatband potential of the macroscopic spectra located at around 350 mV that is an intermediate value between the ones of individual ferritic and austenitic grains, indicating that the macroscopic response averages the ones of the different grains.

[picture]

Fig. 3.34. Typical photovoltammogramms obtained at hν=3.53eV on the whole surface of the 2205 duplex stainless steel and the individual grain of austenitic and ferritic phases oxidised at 850∘C for 60 min (a) in water vapour and (b) in oxygen. Redrawn from A. Srisrual et al., Oxid. Met. 79 (2013) 337-347 [49].

Fig. 3. 34. Typical photovoltammogramms obtained at h ν = 3.53 e V on the whole surface of the 2205 duplex stainless steel and the individual grain of austenitic and ferritic phases oxidised at 850 ∘ C for 60 min (a) in water vapour and (b) in oxygen. Redrawn from A. Srisrual et al., Oxid. Met. 79 (2013) 337-347 [49].

Fig. 3.35 presents the photocurrent energy spectra obtained from individual austenitic and ferritic grains of the 2205 duplex stainless steel oxidised at 850∘C in water vapour for 60 min . The characterisation was performed at the applied potential of -350 mV . Both amplitude and dephasing angle of photocurrent are reported. Dephasing angle that shifts might correlate with either semiconducting behaviour or oxide type. The amplitude and dephasing angle coincidently report the lower bandgap energy of Fe2−xCrxO3 solid solution obtained from austenite comparing to ferrite, indicating a lower x value in oxide grown on austenitic grain.

[picture]

Fig. 3.35. Photocurrent amplitude (white or gray circle) and dephasing angle (dashed line) versus energy spectra obtained on individual (a) austenitic and (b) ferritic grains at the applied potential of -350 mV of the 2205 duplex stainless steel oxidised at 850∘C in water vapour for 60 min . Redrawn from A. Srisrual et al., Oxid. Met. 79 (2013) 337-347 [49].

Fig. 3. 35. Photocurrent amplitude (white or gray circle) and dephasing angle (dashed line) versus energy spectra obtained on individual (a) austenitic and (b) ferritic grains at the applied potential of -350 mV of the 2205 duplex stainless steel oxidised at 850 ∘ C in water vapour for 60 min . Redrawn from A. Srisrual et al., Oxid. Met. 79 (2013) 337-347 [49].

Fig. 3.36 shows the photocurrent energy spectra of the sample from Fig. 3.35. In this case the characterisation was performed at the applied potentials of 350 and 0 mV for the sample oxidised in water vapour and in oxygen respectively. For the results of the oxidation in water vapour (low oxygen activity) shown in Fig. 3.36(a), the differences in photocurrent of ferritic and austenitic oxidised grains are mainly located in the low photon energy range with higher photocurrent issued from the iron rich oxide on the austenitic grain. On the contrary, for the results of the oxidation in oxygen (high oxygen activity) shown in Fig. 3.36(b), the difference in the photocurrent of the ferritic and austenitic oxidised grains is mainly located in the high photon energy range with higher photocurrent issued from the chromium rich oxide on the austenitic grain.

[picture]

Fig. 3.36. Photocurrent energy spectra obtained on individual austenitic (open circle) and ferritic (closed circle) grains of the 2205 duplex stainless steel oxidised at 850∘C for 60 min (a) in water vapour at 350 mV , (b) in oxygen at 0 mV . Redrawn from A. Srisrual et al., Oxid. Met. 79 (2013) 337-347 [49].

Fig. 3. 36. Photocurrent energy spectra obtained on individual austenitic (open circle) and ferritic (closed circle) grains of the 2205 duplex stainless steel oxidised at 850 ∘ C for 60 min (a) in water vapour at 350 mV , (b) in oxygen at 0 mV . Redrawn from A. Srisrual et al., Oxid. Met. 79 (2013) 337-347 [49].

3.6 Summary

This chapter firstly reviews the thermogravimetric method and the analysis of gas composition in the atmosphere emphasising the one with water vapour. The experimental set-up for chromium volatilisation measurement is further addressed. The characterisation techniques for finding physico-chemical properties and semiconducting characteristics of thermal oxide scale are later reviewed. OM and SEM can reveal surface morphology. In addition, the elemental composition can be analysed and mapped by EDS that is typically integrated with SEM. FIB-SEM allows the in-situ milling during SEM performing and can lead to study the oxide in cross-section, to make a thin lamella, and to reconstruct the phases of interest. XPS provides the element composition and chemical state of material which can help identify the oxide species. XRD, Raman spectroscopy, and PEC technique are generally performed to identify the oxide phases. Moreover, PEC is also used to find the semiconducting behaviour of the oxide.

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Solid State Phenomena, ISSN: 1662-9779, Vol. 300, pp 81-106

doi: 10.4028/www.scientific.net/SSP.300.81

© 2020 The Author(s). Published by Trans Tech Publications Ltd, Switzerland.

Submitted:
2019-04-12

Revised:
2019-06-13

Accepted:
2019-08-08

Online:
2020-02-12

CHAPTER 4
High Temperature Oxidation of Stainless Steels

The original version of this paper is available on https://www.scientific.net/SSP.300.81.pdf

a  somrerk.c@eng.kmutnb.ac.th, b  hayashi@eng.hokudai.ac.jp,

c  laurence.latu-romain@simap.grenoble-inp.fr, d  patthranit.wo@kmitl.ac.th

Keywords: high temperature oxidation, stainless steel, alloying element, chromia, alumina

Abstract

This chapter is dedicated to the description of high temperature oxidation of both chromia and alumina forming alloys. The defect structures of iron and chromium are firstly reviewed. The effects of elements on stainless steel oxidation behaviour are further addressed. For the chromiaforming stainless steel, the oxidation rate is reduced with the increased silicon content but not in a monotonic manner. Titanium and niobium can reduce breakaway oxidation of Fe−18Cr−10Ni austenitic stainless steel. Titanium can enhance the adhesion of scale to the Fe−18Cr by mechanical keying effect of TiO2 formed at the steel/scale interface. For the alumina-forming stainless steel, the formation of alumina and its transformation during oxidation are reviewed. Chromium can be added to reduce the critical aluminium content in the steels in order to form alumina at high temperatures. The addition of reactive elements with appropriate level can improve scale adhesion and reduce the steel oxidation rate. Refractory element like molybdenum can increase strength of material but also accelerate the oxidation rate of the steels containing reactive elements. The development of new alumina-forming austenitic alloy grades is finally described.

Somrerk Chandra-ambhorn1,a*, Shigenari Hayashi2,b, Laurence Latu-Romain3,cand Patthranit Wongpromrat4,d

The original version of this paper is available on https://www.scientific.net/SSP.300.81.pdf

1 High Temperature Corrosion Research Centre, Department of Materials and Production Technology Engineering, Faculty of Engineering, King Mongkut's University of Technology North Bangkok, 1518, Pracharat 1 Road, Bangsue, Bangkok, 10800, Thailand
2 Faculty of Engineering, Hokkaido University, N13W8 Kita-ku, Sapporo 060-8268, Japan
3 Université Grenoble Alpes, CNRS, Grenoble INP, SIMaP, 1130 rue de la Piscine, BP 75, 38402 St Martin d'Hères, France
4 Department of Chemical Engineering, Faculty of Engineering, King Mongkut's Institute of Technology Ladkrabang, Chalongkrung 1 Road, Lat Krabang, Bangkok, 10520, Thailand

4.1 Oxidation of Iron and Chromium

4.1.1 Oxidation Mechanism of Iron

The formation of thermal oxide on pure iron is well known, although complex, and is briefly described in this section. At temperatures higher than the eutectoid temperature of the Fe−O system i.e. 570∘C, the formed thermal oxide scale typically consists of the thickest wüstite, with magnetite sitting on the wüstite layer and the outermost haematite sitting on magnetite [1]. Païdassi [2, 3] reported that, for the iron oxides grown at 700−1250∘C in air, haematite was about 1% while magnetite was about 3.5−5% thick. It was reported that the predominant defects in wüstite and magnetite are iron vacancy, while the dominant defect in haematite is oxygen vacancy [4]. Based on this defect structure, the growth mechanism of iron oxides on iron may be described [4] as shown in Fig. 3.1.

For the growth of haematite, at the scale/gas interface oxygen adsorbs on the solid surface i.e.

O2( g)+2 s↔2O−s(4.1)

[picture]

Fig. 4.1. Typical thermal oxide scale on iron at 900∘C in air with the proportion of oxide phases according to Païdassi [2, 3] and with dominant defects in each oxide according to Sarrazin et al. [4].

Fig. 4. 1. Typical thermal oxide scale on iron at 900 ∘ C in air with the proportion of oxide phases according to Païdassi [2, 3] and with dominant defects in each oxide according to Sarrazin et al. [4].

where s is the available site on the surface and O−s is the adsorbed oxygen on the surface. At the interface between magnetite and haematite, the oxygen vacancy is created by the following reaction:

2(FeFe)m↔2(FeFe)h+3( VO)h+2( VFe)m(4.2)

Subscripts h, m, w indicate that the considered entity is in haematite, magnetite and wüstite phase respectively. The formed oxygen vacancy according to Reaction 4.2 diffuses outwards to the haematite/gas interface. At that interface, the adsorbed oxygen according to Reaction 4.1 jumps into the oxygen vacancy in the haematite according to the following reaction:

O−s+(VO)h↔(OO)h+s(4.3)

For the growth of magnetite, at the magnetite/haematite interface, oxygen on the oxygen site in haematite ( (OO)h ) is used to form the new magnetite unit. However, because the iron activity at this interface is lower than that at the wüstite/magnetite interface, the iron vacancy can then be formed as follows:

4(OO)h↔4(OO)m+3( VFe)m+4( VO)h(4.4)

The formed iron vacancy diffuses to the wüstite/magnetite interface. At that interface, iron on the normal iron site in wüstite ( (FeFe)w )) jumps into the metal vacancy in magnetite according to the following reaction:

3(FeFe)w+3( VFe)m↔3(FeFe)m+3( VFe)w(4.5)

For the growth of wüstite, at the wüstite/magnetite interface, one unit of magnetite is used to construct the new wüstite unit. Because the ratio of oxygen and iron in magnetite is higher than that in wüstite, the formation of wüstite from magnetite then gives iron vacancy as follows:

3(FeFe)m+4(OO)m↔3(FeFe)w+4(OO)w+(VFe)m(4.6)

The formed iron vacancy in wüstite ( (VFe)w ) diffuses towards the wüstite/iron interface. At that interface, iron in the matrix, <Fe>, jumps into the iron vacancy in wüstite according to the reaction:

<Fe>+(VFe)w↔(FeFe)w(4.7)

It is noted that at the temperature of 570∘C, there exists the eutectoid reaction among wüstite, iron and magnetite. Thus when wüstite is cooled down from high temperatures to 570∘C or lower temperatures, wüstite can decompose giving the eutectoid structure consisting of iron and magnetite, as typically observed in the thermal oxide on hot-rolled low carbon steel where the cooling rate of the strip is essentially slow [1,5-7].

4.1.2 Oxidation Mechanism of Chromium

Latu-Romain et al. [8-10] have reported that the chromia thermally grown on chromium consists of the columnar external part and the equiaxed internal one. It is proposed that the dominant defect in the external chromia is the p-type chromium vacancy, while the dominant defect in the internal chromia is the n-type oxygen vacancy as sketched in Fig. 4.2, when the external oxygen partial pressure is superior or equal to 10−12 atm.

[picture]

Fig. 4.2. Structure of chromia thermally grown on chromium. Redrawn and adapted from L. Latu-Romain et al., Oxid. Met. 86 (2016) 497-509 [8].

Fig. 4. 2. Structure of chromia thermally grown on chromium. Redrawn and adapted from L. Latu-Romain et al., Oxid. Met. 86 (2016) 497-509 [8].

As suggested by Latu-Romain et al. [8], the defect concentrations in chromia can be analysed as a function of the oxygen partial pressure, giving the Brouwer diagram, as follows.

The relevant defect reactions are

32O2↔3OO+2 VCr′′′+6 h•32O2+3 VO••+6e′↔3OO nil ↔2 VCr′′′+3 VO•• nil ↔e′+h•

(4.8)(4.9)(4.10)(4.11)

with the equilibrium constants of

K4.8=[VCr′′′]2[ h•]6pO23/2 K4.9=1pO23/2[ VO••]3[e′]6 KS=[VCr′′′]2[ VO••]3 Ke=[e′][h•]

(4.12)(4.13)(4.14)(4.15)

K4.8 and K4.9 are equilibrium constants of Reactions 4.8 and 4.9 respectively. KS and Ke are equilibrium constants of Schottky defect reaction i.e. Reaction 4.10, and Reaction 4.11 respectively. The electroneutrality condition for this problem is

2[ VO••]+[h•]=3[ VCr′′′]+[e′](4.16)

The dominant regimes of defects may be considered as follows.
Regime 𝐈 is the case at relatively low pO2 where oxygen vacancy and electron are dominant as can be seen from Reaction 4.9. In this case the electroneutrality condition described in Eq. 4.16 can be reduced to the following one:

2[ VO••]=[e′](4.17)

Regime II is the case at intermediate pO2. By considering that the electronic defect was dominant in this regime [8], the following electroneutrality condition can be drawn:

[e′]=[h•](4.18)

Regime III is the case at relatively high pO2 where metal vacancy and electron hole are dominant as suggested by Reaction 4.8. In this case the electroneutrality condition described in Eq. 4.16 can be reduced to the following one:

[h•]=3[ VCr′′′](4.19)

As for Regime I, from Eqs. 4.13 and 4.17, the oxygen vacancy concentration can be expressed in the following form:

[VO••]=2−2/3 K4.9−1/9pO2−1/6(4.20)

From Eq. 4.20 and the electroneutrality condition according to Eq. 4.17, it can be obtained that

[e′]=2[ VO••]=21/3 K4.9−1/9pO2−1/6(4.21)

From Eq. 4.21 and electronic defect equilibrium constant in Eq. 4.15, we obtain that

[h•]=2−1/3 Ke K4.91/9pO21/6(4.22)

From Eq. 4.20 and the Schottky defect equilibrium constant in Eq. 4.14, the chromium vacancy concentration can be obtained as follows:

[VCr′′′]=2 K4.91/6 KS1/2pO21/4(4.23)

For Regime II, from Eq. 4.15 and the electroneutrality condition in Eq. 4.18, we obtain that

[e′]=[h•]=Ke1/2= constant (4.24)

Inserting Eq. 4.24 into Eq. 4.12 gives

[VCr′′′]=K4.81/2 Ke−3/2pO23/4(4.25)

Inserting Eq. 4.24 into Eq. 4.13 gives

[VO••]=K4.9−1/3 Ke−1pO2−1/2(4.26)

For Regime III, by inserting the electroneutrality condition according to Eq. 4.19 into Eq. 4.12, the chromium vacancy concentration can be obtained in the following form:

[VCr′′′]=3−3/4 K4.81/8pO23/16(4.27)

From Eq. 4.27 and electroneutrality condition according to Eq. 4.19, we obtain that

[h•]=3[ VCr′′′]=31/4 K4.81/8pO23/16(4.28)

From Eqs. 4.28 and 4.15, it can be obtained that

[e′]=3−1/4 Ke K4.8−1/8pO2−3/16(4.29)

From Eq. 4.27 and the Schottky defect equilibrium constant in Eq. 4.14, the oxygen vacancy concentration can be expressed in the following form:

[VO••]=31/2 KS1/3 K4.8−1/12pO2−1/8(4.30)

By using Eqs. 4.20-4.30, the Brouwer diagram of the chromia thermally grown on chromium in oxygen can be drawn as shown in Fig. 4.3. It was reported [9] that the transition zone between the high pO2 region where metal vacancy is dominant and the intermediate pO2 region was reported to be in the pO2 range of 5×10−13 to 10−12 atm. The transition zone between the low pO2 region where oxygen vacancy is dominant and the intermediate one was reported to be in the pO2 range of 10−14 to 5×10−14 atm [9].

[picture]

Fig. 4.3. Brouwer diagram of chromia thermally grown on chromium in oxygen. Redrawn and adapted from L. Latu-Romain et al., Oxid. Met. 86 (2016) 497-509 [8], L. Latu-Romain et al., Corros. Sci., 141 (2018) 46-52 [9] and L. Latu-Romain et al., Corros. Sci. 126 (2017) 238-246

Fig. 4. 3. Brouwer diagram of chromia thermally grown on chromium in oxygen. Redrawn and adapted from L. Latu-Romain et al., Oxid. Met. 86 (2016) 497-509 [8], L. Latu-Romain et al., Corros. Sci., 141 (2018) 46-52 [9] and L. Latu-Romain et al., Corros. Sci. 126 (2017) 238-246

[10].

4.2 Chromia-Forming Stainless Steel

For the Fe−2.25Cr steel (unless stated otherwise, the alloy composition is in wt.% ), it was reported that after oxidation in dry air at 600∘C, the oxide formed on this steel was poorly protective and comprised haematite sitting on magnetite and Fe−Cr spinel [11]. When chromium increases up to 18wt.%, the chromia topped with Cr−Mn spinel can be formed as previously shown and discussed in Fig. 1.4. In the case that other alloying elements are added to the steel, they may form the oxide as layer or nodule and affect properties of the scale as will be reviewed in the following.

4.2.1 Effect of Silicon

In the case of carbon steel, Chandra-ambhorn et al. [12] oxidised iron with silicon contents of 0.23 and 1.03wt.% in Ar−20%H2O. It was found that the oxidation rate was linear-parabolic. The addition of silicon to the steel formed fayalite (Fe2SiO4) as an intermediate layer between substrate and iron oxides. It was reported that diffusivity of iron in fayalite is slower than that in wüstite [12, 13]. As a result, the formation of this layer could reduce the steel oxidation rate by inhibiting the outward diffusion of iron from the substrate to form iron oxides. However, the scale of steel with 1.03wt.% of silicon was found to be thicker than that of the steel with 0.23wt.% of silicon. To explain this result, even though the external iron oxides of the former steel was thinner than that of the later steel, the internal fayalite layer of the former steel was thicker in higher degree thus resulting in increasing the total scale thickness. In the case of stainless steel, Bamba et al. [14] studied the oxidation of Fe−15Cr without and with 0.5 and 1wt.% of silicon. It was found that increasing silicon content could reduce the steel oxidation rate but the oxidation rate of Fe−15Cr− 0.5 Si was slower than that of Fe−15Cr−1Si. The discontinuous silica precipitates was found at the metal/scale interface and beneath the oxide scale for the oxidised Fe−15Cr−0.5Si, while the
continuous silica layer was found at the interface between the chromia scale and the Fe−15Cr−1Si substrate. They proposed that the most reduced oxidation rate of the Fe−15Cr−0.5Si is not from the formation of silica as a barrier to prevent the incorporation of chromium into the oxide scale but probably from the formation of silica precipitates which inhibits the diffusion of chromium from steel to the metal/scale interface by providing sinks for the injected metal vacancies [14]. If chromium vacancy is inhibited to enter the metal, it can be accumulated at the scale/steel interface and as a result reduces the scale adhesion as observed in that work [14].

4.2.2 Effects of Titanium and Niobium

Wongpromrat et al. [15] studied the formation of titanium and niobium rich oxide on Fe−18Cr−0.13Ti−0.55Nb (AISI 441) oxidised at 800∘C in O2−5%H2O for 24 h . It was found that the oxide nodule could be formed above Fe2Nb which was precipitated at the metal grain boundary. This Fe−Nb intermetallic compound appears in the Fe-Nb system during tempering operations [15,16]. The nodule could also be formed at the heart of the grain as shown in Fig. 4.4. From this figure, we can see the thin continuous chromia layer outside the nodule. However, in fact this layer also lies at the middle of the nodule indicating the original formation of the chromia layer in that area. Titanium can diffuse to the chromia layer and dissolves there. It can also transport to the top surface to form titania. Niobium is found to accompany titanium and can dissolve in titania. Fig. 4.4 and the Raman spectroscopy results reported in literature [15] indicated that the nodule external part consists mainly of Ti(Nb)O2. Fig. 4.5 shows the three dimensional image of the oxide nodule above the grain of AISI 441 after oxidation at 800∘C in O2−5%H2O for 24 h . We can see that the nodule is connected to the surrounding scale and situates above the intragranular niobium-rich precipitates. The size of precipitate tends to be smaller when it is closer to the nodule, suggesting that the nodule growth is a driving force for the precipitate dissolution and the transport of titanium and niobium to the nodule. This precipitate could correspond to NbC and/or (Ti, Nb)(C, N) [15,16].

The addition of titanium and niobium can also affect properties of both metal and oxide scale formed at high temperatures. Toscan et al. [17] studied the oxidation of ferritic stainless steels (Fe−18Cr,Fe−18Cr−0.41Ti,Fe−18Cr−0.58Nb) and austenitic ones (Fe−18Cr−10Ni,Fe−18Cr−10Ni−0.32Ti,Fe−18Cr−10Ni−0.46Nb ) at 950∘C in laboratory air. The results are shown in Fig. 4.6.

[picture]

Fig. 4.4. Cross-sectional STEM/EDX maps of the oxide nodule above the grain of AISI 441 after oxidation at 800∘C in O2−5%H2O for 24 h . Reproduced from W. Wongpromrat et al., Mater. High Temp. 32 (2015) 22-27 [15].

Fig. 4. 4. Cross-sectional STEM/EDX maps of the oxide nodule above the grain of AISI 441 after oxidation at 800 ∘ C in O 2 − 5 % H 2 O for 24 h . Reproduced from W. Wongpromrat et al., Mater. High Temp. 32 (2015) 22-27 [15].

[picture]

Fig. 4.5. Three dimensional image of the oxide nodule above the grain of AISI 441 after oxidation at 800∘C in O2−5%H2O for 24 h (red: oxide layer with chromium and titanium; blue: silicon; yellow: niobium). Reproduced from W. Wongpromrat et al., Mater. High Temp. 32 (2015) 22-27 [15].

Fig. 4. 5. Three dimensional image of the oxide nodule above the grain of AISI 441 after oxidation at 800 ∘ C in O 2 − 5 % H 2 O for 24 h (red: oxide layer with chromium and titanium; blue: silicon; yellow: niobium). Reproduced from W. Wongpromrat et al., Mater. High Temp. 32 (2015) 22-27 [15].

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Fig. 4.6. The ratio of diffusivity of chromium in stainless steel to parabolic rate constant of oxidation at 950∘C in laboratory air [17].

Fig. 4. 6. The ratio of diffusivity of chromium in stainless steel to parabolic rate constant of oxidation at 950 ∘ C in laboratory air [17].

Fig. 4.6(a) shows the diffusivity of chromium in stainless steels. It can be seen that the diffusivities of chromium in ferritic steels are faster than those in austenitic ones. This is because the ferritic crystal structure was loosely packed compared with the austenitic one [17]. The addition of stabilisers (titanium and niobium) does not considerably change the chromium diffusivity in ferritic grades as it was observed that the difference in diffusivities between the stabilised grades and the non-stabilised one is only in the range of 8−15%. However, even though the diffusivities of chromium in austenitic grades are relatively low, the addition of stabilisers significantly increases the chromium diffusivity i.e. by about 170 and 264% for the addition of titanium and niobium respectively.

After the isothermal oxidation test at 950∘C in laboratory air up to 60 h , the oxidation kinetics of the studied ferritic stainless steels were parabolic with the accelerated oxidation rate for the stabilised steels as shown in Fig. 4.6(b). Toscan et al. [17] observed that the addition of titanium gave internal oxides which were possibly responsible to the faster oxidation rate. For the austenitic stainless steels, they found that the oxidation rate of Fe−18Cr−10Ni was also parabolic but only up to about 9 h before breakaway oxidation occurred. The addition of stabilisers in austenitic stainless steel could suppress the breakaway oxidation at that early period of oxidation [17]. Fig. 4.6(b) shows the rate constants of the austenitic stainless steels in the parabolic oxidation regime. For the cyclic oxidation test, the oxidation rates of the ferritic stainless steels were still parabolic [17]. On
the contrary, spallation of the thick and porous nodular iron oxides formed on austenitic stainless steels was observed with the greater mass loss for the non-stabilised austenitic stainless steel [17].

As a matter of discussion, the diffusivity of chromium in metal indicates the supply rate of chromium from the metal to the metal/scale interface, while the parabolic rate constant indicates the consumption rate of chromium for oxidation. The ratio of chromium diffusivity to parabolic rate constant could then be used to assess the tendency of chromium depletion at the metal/scale interface which could lead to the formation of the non-protective iron oxides. From the data of chromium diffusivities and parabolic rate constants shown in Figs. 4.6 (a) and (b), that ratio was calculated and plotted in terms of log(10DCr/kp) as shown in Fig. 4.6(c). It can be seen that the ratios of chromium diffusivity to parabolic rate constant for ferritic grades are about 58-325. These values are considerably higher than those for the austenitic ones which are in the range of 0.5−7.0. One of the major reasons for this difference is that chromium diffusivity in the ferritic grade is faster than that in the austenitic one. The high chromium supply rate could effectively maintain chromium content at the metal/scale interface to form the protective chromia. For the non-stabilised austenitic stainless steels (Fe−18Cr−10Ni), this ratio is about 0.53 . The very low ratio indicates that the chromium supply rate is not enough to maintain the chromium content at the metal/scale interface to form the protective chromia and thus the breakaway oxidation is possible. The addition of stabilisers could increase chromium diffusivity in the austenitic matrix as shown in Fig. 4.6(a), contributing to the increased ratio of chromium diffusivity to parabolic rate constant as shown in Fig. 4.5(c). As a result, the suppression of the breakaway oxidation is possible. The addition of stabilisers is then required for austenitic grades for this purpose.

Titanium and niobium do not only affect the chromium diffusion in the steel and the oxidation rate constant, they also affect the scale adhesion. Mougin et al. [18] investigated the effects of stabilisers on the adhesion of scale on stainless steels. In this work, Fe−18Cr,Fe−18Cr−0.14Ti−0.47Nb,Fe−18Cr−0.45Ti and Fe−18Cr−0.44Nb were oxidised at 900∘C in Ar−15%O2. For the grades containing titanium, the nodular titania was observed on top of the oxide scale as well as along the metal grain boundary including the one at the metal/scale interface. For the grades containing niobium, the Fe2Nb intermetallic was also found along the metal grain boundary. It was also observed that the spontaneous spallation of the oxide scale on Fe−18Cr−Ti was drastically reduced compared to the oxidised Fe−18Cr and Fe−18Cr−Nb respectively. Those samples were further subjected to the inverted blister test for the scale adhesion assessment. It was found that the adhesion energy of the scale on Fe−18Cr−Ti was drastically higher than that of the scale on Fe−18Cr−Nb, while the adhesion energy of the scale on Fe−18Cr−Ti−Nb was between those former two values and slightly lower than the adhesion energy of the scale on Fe−18Cr. In case of Fe- 18Cr−Ti, after the inverted blister test the scale and stainless steel substrate were debonded as schematised in Fig. 4.7. It is seen that the TiO2 which precipitates at the metal/scale interface remains in the metal side while the oxide scale including Cr−Mn spinel, chromia and silica (when existing) are removed from the substrate. This result exhibits the mechanical keying effect of TiO2 on improving the scale adhesion. On the contrary, in case of Fe−18Cr−Ti−Nb, even though TiO2 nodule exists at the metal/scale interface, the formation of Fe2Nb is found to be under the titania nodule. This Fe2Nb reduces the adhesion of the nodular titania to the metal substrate, and as a result after the inverted blister test the oxide scale together with the nodular titania is detached from the metal substrate. For the other stainless steels, the Fe-Mo intermetallic compound was also observed at the interface between scale and the AISI 444 stainless steels. Its effect on reducing the scale adhesion was also reported [19].

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Fig. 4.7. Schematic sketch showing the roles of titanium and niobium on the adhesion of scale to Fe-18Cr-Ti-Nb stainless steel. Redrawn and adapted from J. Mougin et al., Mater. Sci. Eng. A 359 (2003) 44-51 [18].

Fig. 4. 7. Schematic sketch showing the roles of titanium and niobium on the adhesion of scale to Fe-18Cr-Ti-Nb stainless steel. Redrawn and adapted from J. Mougin et al., Mater. Sci. Eng. A 359 (2003) 44-51 [18].

4.2.3 Effects of Reactive Elements

Reactive element (RE) is the element that is more active to oxygen than the elements that form the oxide scale i.e. chromium and aluminium [20, 21]. These elements are such as yttrium, lanthanum, zirconium, cerium and hafnium [21, 22]. The addition of the reactive elements could change the growth mechanism of the oxide, thus affecting behaviour and properties of the oxidesteel system e.g. the selective oxidation, the oxidation rate and the scale adhesion.

For the chromia-forming stainless steel, the thermal scale is typically chromia topped by the spinel of manganese and chromium [19]. Comparing to the undoped steel, when the reactive element is alloyed it was reported that this element could reduce the alloy grain size and therefore increase the cationic outward flux, thus promoting the formation of the protective chromia scale earlier [21]. Furthermore, when the oxidation proceeds, it was reported that the reactive element tends to diffuse to the scale/gas interface via the oxide grain boundary where it segregates. Because the ionic radius of the reactive element is typically bigger than that of chromium, the reactive element segregation thus blocks the outward chromium diffusion resulting in the reduced oxidation rate [21, 23].

The other beneficial effect of the reactive element is to improve the scale adhesion. The example of this effect exhibits in Fig. 6.4 in Chapter 6 that shows the oxide scale on Crofer 22 APU (Fe-22Cr-La-Ti) and AISI 441 stainless steels oxidised in air and water vapour before and after the tensile test for scale adhesion assessment. It can be seen that the scale on Crofer 22 APU which contains lanthanum well adheres to the substrate while the scale on AISI 441 is spalled under the imposed strain [24]. The mechanism proposed to explain the improved scale adhesion by the reactive element is the surface sulphur cleaning effect suggesting that the interaction between the reactive element and sulphur could reduce the sulphur segregation at the metal/scale interface [23, 25]. The other mechanism proposed is the reduction of void formation by the reactive element. This might be because the reduction of the cationic outward diffusion resulted in the lower vacancy injection to the metal/scale interface [23, 26]. It was also proposed that the reactive element might reduce the surface energy, thus reducing the driving force for void formation and therefore the formation of interfacial voids [22, 23].

4.3 Alumina-Forming Stainless Steels

4.3.1 High Temperature Oxidation of Alumina-Forming Stainless Steels

Most of commercial stainless steels form a protective Cr2O3 scale when exposed at higher temperatures in oxidising atmospheres. The Cr2O3 scale shows good oxidation resistance but it is much less protective compared to an Al2O3 scale. Al2O3 is one of the most thermodynamically stable oxides and protects the alloy substrate even at the temperatures higher than 1300∘C. Heat resistant steels which are protected by an Al2O3 scale are, therefore, expected to have excellent oxidation resistance and extended lifetime at higher temperatures. However, because there are several problems for Al2O3-forming alloys, commercial grade of Al2O3-forming stainless steels has not been widely used.

One of the typical problems of the Al2O3-forming alloys is formation of Al2O3 scale at the temperatures of most of heat resistant alloys are used. The temperatures in which most of heat resistant steels are used as the components of chemical plants, incinerators, boilers, etc. are not so high, ~800∘C. At such temperature region, Al2O3 scale formation is difficult. Fig. 4.8 shows the example of initial Al2O3 formation on a Fe−24Cr−10Al (at.%) stainless steel measured by in-situ high-temperature XRD using a synchrotron source [27]. In this test, the specimen was exposed to air and heated with the heating rate of 50∘Cmin−1 to 1000∘C. Then the sample was held isothermally at 1000∘C for 1 h in the same atmosphere. During the experiment, i.e. heating, isothermal oxidation and cooling, the XRD measurement was conducted on the sample surface every 10 s using 6 s for data collection in each measurement. Initial Al2O3 formation is always observed at the temperatures above 800∘C.

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Fig. 4.8. In-situ high-temperature XRD patterns of oxide scale formed on Fe-24Cr-10Al (at.%) heating to 1000∘C. Redrawn and adapted from S. Yoneda et al., Oxid. Met. 89 (2018) 81-97 [27].

Fig. 4. 8. In-situ high-temperature XRD patterns of oxide scale formed on Fe-24Cr-10Al (at.%) heating to 1000 ∘ C . Redrawn and adapted from S. Yoneda et al., Oxid. Met. 89 (2018) 81-97 [27].

The second problem is the preferential formation of a less protective metastable Al2O3. Although this problem will be discussed below, metastable Al2O3 scale formation is not avoided on the alloys with high aluminium contents and nickel-based alloys even though the alloys are oxidised at temperatures over 1100∘C.

In addition, high alloy aluminium contents are also not acceptable for heat resistant steels. As a matter of fact, despite the requirement of a sufficient aluminium content in the alloy for the formation of a protective Al2O3 scale, aluminium additions are usually very limited because of the adverse effect of aluminium on the ductility, weldability, and fabricability of the substrate. Thus, aluminium concentration in the substrate needs to be lowered as much as possible.

Currently, most of successful commercial Al2O3-forming stainless steels are ferritic stainless steels. They are used in metal honeycombs of automobile emission system i.e., catalytic converters and in heating wires, which are usually used at temperatures more than 1100∘C. In the Sections 4.3-4.5, the oxidation behaviour of Al2O3-forming stainless steels, which contain not only steels called as "stainless" but also iron-based alloys which form a protective Al2O3 scale, will be discussed.

4.3.2 Metastable and Stable Alumina Scale Formation

It is well-known that Al2O3 has several isomorphous structures, such as γ,θ, and α−Al2O3. Among those Al2O3 phases, γ,θ−Al2O3 are metastable phases and transition eventually to α−Al2O3 [28, 29]. When Al2O3-forming alloys are oxidised, a metastable Al2O3 scale is usually formed initially, and it later transforms to the stable Al2O3 scale [30-33]. This initial metastable Al2O3 formation is one of the issues for the use of Al2O3-forming alloys, because the transition to α−Al2O3 takes longer time at lower temperatures [31] and the growth rate of metastable Al2O3 is much higher than that of stable Al2O3 scale [34, 35]. In order to accelerate the transformation, much attention has been paid towards understanding the mechanism of the metastable-stable Al2O3 transformation [3639], even though the details of the mechanism, particularly for the thermally grown oxide scale of Al2O3, have not been fully understood yet.

Tanigushi and Andoh [31] studied the isothermal transformation of Al2O3 scales formed on Fe−Cr−Al alloys with/without aluminium-deposition. In the air oxidation between 1073−1373∘C, they found that transformation to α−Al2O3 occurs in shorter time at higher temperatures on both samples. Regarding aluminium contents, it is clear that transformation is delayed on the alloys with higher surface aluminium contents. Higher aluminium addition in alloys is also detrimental for rapid formation of α−Al2O3.

This delayed transformation on the alloy is strongly related to the initial transient oxidation behaviour of alloys. Transformation of metastable to stable α−Al2O3 scale is known to be accelerated by the presence of Fe2O3 and/or Cr2O3 scale. The beneficial effect of such oxides is referred as the "template" effect, since Fe2O3 and Cr2O3 have the same corundum crystal structure as α−Al2O3. Those oxides act as a "template" for the nucleation of α−Al2O3. From the previous studies, the template effect is interpreted as a precipitation of α−Al2O3 from the transient Fe2O3 with Al2O3 saturation [35, 40-43]. Hayashi et al. [44] studied the meta-stable phase transformation behaviour of Al2O3 scale on Fe−Ni−Al alloys. They found that the oxidation mass gain tended to be reduced with the increased iron content in the alloy. The increased iron content in the alloy reduced the time needed for the transition from the rapid oxidation stage in the early period to the slower oxidation stage in the later stage. Fig. 4.9 shows the phase transformation behaviour of Al2O3 scale which formed on Fe−41.5Al(at.%), Ni−41Al−41Fe(at.%) and Ni−41.5Al(at.%) [44]. It was found that high iron content is beneficial and promotes the rapid transformation to α−Al2O3. The samples on which α−Al2O3 developed rapidly always formed Fe2O3 prior the formation of Al2O3 scale in the transient oxidation period as indicated in Fig. 4.9(d). However, transformation to α−Al2O3 is delayed when alloys are oxidised in atmospheres where the oxygen partial pressure is lower than that of formation of Fe2O3[45]. Iron-based stainless steels might have advantage to form α−Al2O3 due to initial Fe2O3 or Cr2O3 formation. However, rapid establishment of α−Al2O3 results in smaller α−Al2O3 grain sizes [46], which might influence its growth rate of the α−Al2O3 scale.

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Fig. 4.9. In-situ high-temperature XRD patterns of oxide scale formed on (a) Fe-41.5Al (at.%), (b) Ni−41Al−41Fe (at.%), (c) Ni−41.5Al (at.%) and (d) Fe−41.5Al (at.%) which is the enlarged patterns of (a) during heating with the rate of 20∘Cmin−1 to 1000∘C followed by isothermal oxidation up to 1 h. Redrawn and adapted from S. Hayashi et al., Oxid. Met. 86 (2016) 151-164 [44].

Fig. 4. 9. In-situ high-temperature XRD patterns of oxide scale formed on (a) Fe-41.5Al (at.%), (b) N i − 41 A l − 41 F e (at.%), (c) N i − 41.5 A l (at.%) and (d) F e − 41.5 A l (at.%) which is the enlarged patterns of (a) during heating with the rate of 20 ∘ C min − 1 to 1000 ∘ C followed by isothermal oxidation up to 1 h. Redrawn and adapted from S. Hayashi et al., Oxid. Met. 86 (2016) 151-164 [44].

4.4 Alumina-Forming Ferritic Stainless Steel

4.4.1 Effect of Chromium

As mentioned above, higher aluminium additions to the alloy substrate are restricted. Chromium additions are known to reduce the critical aluminium content to form a protective Al2O3 scale. Therefore, most of the commercial Al2O3-forming stainless steels contain relatively high chromium contents, in the range 13~20%. The composition of Al2O3-forming steels is typically Fe−20Cr−10Al (at.%) with additions of small amounts of various reactive elements such as yttrium, cerium, lanthanum, zirconium and so on. Regarding the mechanical properties of alloy, higher chromium content is detrimental due to the possibility of formation of the brittle chromium-rich bcc phase ( α′ ) in the ferritic ( α ) matrix of Fe−Cr−Al system because the miscibility gap at lower temperatures. For the Fe−Cr−Al system in the region of ferrite substrate with an aluminium content about 10at.% at 475∘C [47], the maximum chromium content allowed in order to avoid α′ precipitation is about 12~15at.%. The Al2O3-forming stainless steels are usually used at higher temperatures more than 1000∘C. Therefore, α′ formation can be neglected during service time. However, if high-chromium stainless steels are exposed at lower temperatures for longer times, embrittlement due to precipitation of α′-phase should be considered.

The effect of chromium on promoting Al2O3 scale formation on the substrate with low aluminium contents has been widely investigated. It is believed that the initially formed Cr2O3 scale decreases the oxygen potential at the metal/scale interface, which decreases inward oxygen diffusion in the substrate, thus promoting a protective Al2O3 scale without internal aluminium oxidation [48-50]. Wagner suggested this effect of chromium as an "oxygen getter" [51]. Yoneda et al. [27, 52] investigated the initial oxidation behaviour of Fe−Cr−Al alloys with 4−24at.% of chromium to verify this mechanism by the in-situ high-temperature XRD technique using a synchrotron source. Fig. 4.10 shows XRD patterns of Fe−6Al (at.%) and Fe−4Cr−6Al (at.%) alloys during heating to 1000∘C in air and the cross-section of the oxide scale after isothermal oxidation at 1000∘C for 1 h [52]. The diffracted signals obtained from both alloys during heating are only Fe2O3. However, an exclusive Al2O3 scale has developed on Fe−4Cr−6Al (at.%) after 1 h of oxidation. They also found that the addition of chromium by 24 at. % to Fe−6Al (at.%) alloy does not prevent internal oxidation at the initial oxidation during heating. Apparently, an initial Cr2O3 scale formation cannot be the
main reason for promoting Al2O3 scale formation. The interdiffusion coefficients of aluminium in the Fe−Cr−Al system at 1000∘C was evaluated [53] and the results suggested that chromiumdepletion in the subsurface region in the initial oxidation period may enhance the aluminium outward diffusion flux due to the positive cross-term coefficient, which promotes the Al2O3 scale formation. Airiskallio et al. [54] also made the discussion similar to this "cross-term" effect. Additionally, it is remarked that the alumina morphology is the opposite of the chromia one, and composed of an inner columnar and an outer equiaxed subscales.

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Fig. 4.10. In-situ high-temperature XRD patterns of (a) Fe-6Al (at.%) and (b) Fe-4 Cr−6Al (at.%) alloys during heating, and (c) the TEM cross-section of Fe−4Cr−6Al (at.%) alloy after oxidation at 1000∘C in air for 1 h . Redrawn and adapted from S. Yoneda et al., Oxid. Met. 86 (2016) 357-370 [52].

Fig. 4. 10. In-situ high-temperature XRD patterns of (a) Fe-6Al (at.%) and (b) Fe-4 C r − 6 A l (at.%) alloys during heating, and (c) the TEM cross-section of F e − 4 C r − 6 A l (at.%) alloy after oxidation at 1000 ∘ C in air for 1 h . Redrawn and adapted from S. Yoneda et al., Oxid. Met. 86 (2016) 357-370 [52].

4.4.2 Effect of Refractory Elements

In order to satisfy the demand to reduce the greenhouse gas and fuel consumption rate, the operation temperature of recent automobile engines is increasing to improve the thermal efficiency. Therefore, even though the Al2O3-forming ferritic stainless steels are predominantly used for catalytic converters, increasing the strength at higher temperature of the steels is still required. In order to increase the strength of the steels, refractory metals such as molybdenum, tungsten and tantalum are used as solid-solution hardening elements. Among those elements, molybdenum is
widely used for Fe−Cr−Al steels which are used for not only the catalytic system but also heating elements. The addition of molybdenum is reported to reduce the oxidation performance of Fe−Cr−Al−La and Fe−Cr−Al−La−Zr alloys by increasing the growth rate of the Al2O3 scale [55, 56]. Mizutani et al. [56] investigated the detrimental effect of molybdenum on the oxidation of Fe−Cr−Al−La−Zr alloys and found that molybdenum reduces the beneficial effect of lanthanum and zirconium on Al2O3 scale growth as shown in Fig. 4.11.

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Fig. 4.11. Effect of molybdenum addition on the oxidation kinetics of (a) Fe−20Cr−6Al (at.%), (b) Fe−20Cr−6Al−0.08La (at.%) and (c) Fe−20Cr−6Al−0.08La−0.03Zr (at.%) during oxidation at 1100∘C in air. Redrawn and adapted from A. Mizutani et al., Oxid. Met. 87 (2017) 109-124 [56].

Fig. 4. 11. Effect of molybdenum addition on the oxidation kinetics of (a) F e − 20 C r − 6 A l (at.%), (b) F e − 20 C r − 6 A l − 0.08 L a (at.%) and (c) F e − 20 C r − 6 A l − 0.08 L a − 0.03 Z r (at.%) during oxidation at 1100 ∘ C in air. Redrawn and adapted from A. Mizutani et al., Oxid. Met. 87 (2017) 109-124 [56].

As it will be shown in following section, reactive elements greatly reduce the growth rate of Al2O3 scales by reducing grain boundary diffusivities of anion and oxide ions via their grain boundary segregation. The results by Mizutani et al. [56] indicated that molybdenum additions reduce the grain boundary segregation of reactive elements and eliminate the blocking effect of reactive elements for ion diffusion. They concluded that higher reactive element additions are
necessary for molybdenum-containing alloys in order to reduce the growth rate of Al2O3 scales to similar levels as the Al2O3 scale formed on the alloys without molybdenum. This detrimental effect of molybdenum was speculated to be caused by molybdenum reducing the thermodynamic activity of lanthanum and/or zirconium in the Fe−Cr−Al substrate.

4.4.3 Effect of Reactive Elements

Similar to the case of chromia-forming stainless steels, reactive elements have two important effects on the oxidation of alumina-forming steels. They increase the adhesion and reduce the growth rate of an Al2O3 scale [57]. Various mechanisms to increase adhesion of the scale have been proposed such as sulphur and/or carbon gettering, keying effect and void formation suppression at the metal/scale interface [57-59]. The reduced Al2O3 scale growth rate by the reactive element additions is explained by their segregation at grain boundaries [22]. The grain boundary diffusivity of anion and oxide ions in Al2O3 is greatly reduced by doping with reactive elements [60]. Among those beneficial effects, Pint et al. [61] suggested that improved adhesion of the Al2O3 scale is the most important effect for increasing the lifetime of steels. However, reduction of the growth rate by reactive element additions is also an important effect, particularly for foils and heating elements, since aluminium supply from the matrix in thin foils and wires is very limited.

Over-doping of reactive elements in alloys causes rapid oxidation, thus optimisation of reactive element concentrations is strongly required [61, 62]. Because the tolerance of concentration of reactive elements depends on reactive elements and alloy systems, it is necessary to establish the method to find the best concentration of reactive elements. Shizukawa et al. [63] investigated the oxidation behaviour of ODS Fe−Cr−Al−Y2O3 alloys with different excess oxygen, which is the total content of oxygen subtracted by the one used to form Y2O3 during the alloy production, to control the activity of yttrium, aY, in the substrate. Fig. 4.12 shows the effect of excess oxygen on the oxidation mass gain of an ODS alloy of composition Fe−7Al−15Cr−0.5Ti−0.4Zr−0.5Y2O3 at 1200∘C in air. Lower excess oxygen level causes higher oxygen mass gain due to over-doping of zirconium, but oxidation mass gain drastically decreases on the alloys with higher excess oxygen content. Based on the result shown in Fig. 4.12, they suggested that over-doping by zirconium is strongly related to the stability of ZrO2 and Al2O3 in the alloy substrate, i.e. activity of zirconium ( aZr ) and Al(aAl) which are given by Eq. 4.31.

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Fig. 4.12. Oxidation kinetics of Fe−7Al−15Cr−0.5Ti−0.4Zr−0.5Y2O3 with excess oxygen at 1200∘C in air. Redrawn from Y. Shizukawa et al., The effect of excess oxygen on the high temperature oxidation in zirconium added FeCrAl ODS steels, in: Y. Murata et al. (Eds.), Proceedings of the First International Conference on Advanced High-Temperature Materials Technology for Sustainable and Reliable Power Engineering, Japan, 2015, pp. 149-152 [63].

Fig. 4. 12. Oxidation kinetics of F e − 7 A l − 15 C r − 0.5 T i − 0.4 Z r − 0.5 Y 2 O 3 with excess oxygen at 1200 ∘ C in air. Redrawn from Y. Shizukawa et al., The effect of excess oxygen on the high temperature oxidation in zirconium added FeCrAl ODS steels, in: Y. Murata et al. (Eds.), Proceedings of the First International Conference on Advanced High-Temperature Materials Technology for Sustainable and Reliable Power Engineering, Japan, 2015, pp. 149-152 [63].

aZr<aAl4/3[exp(ΔGZrO∘RT)exp(ΔGAl2O3∘RT)](4.31)

where R is the gas constant, T is the absolute temperature, ΔGZrO2∘ and ΔGAl2O3∘ are the Gibbs free energies of formation of ZrO2 and Al2O3, respectively. Gheno et al. [64] also proposed a similar mechanism for hafnium-tolerance in γ+γ′Ni−Cr−Al alloys. They further suggested that the level of hafnium content varies due to structural changing of alloy subsurface region by the selective oxidation of aluminium because γ′ dissolution in the substrate [64]. Thus, the factors which may control the activities of aluminium and reactive elements are expected to affect the optimised concentration of reactive elements. Optimisation of reactive element concentration is one of the most important factors for producing higher oxidation resistant commercial stainless steels. Therefore, further studies are still necessary.

Because the Al2O3 scale grows mainly by grain boundary diffusion at the temperature range of our interest, the microstructure of the Al2O3 and its evolution during oxidation strongly affect the growth rate of the Al2O3 scale. Naumenko et al. [65] reported that the grain size of columnar grains of an inwardly growing Al2O3 scale formed on Fe−Cr−Al−Y alloy at 1200∘C in air linearly increases with the depth from the gas/scale interface. They supposed that coarsening of columnar grains, i.e. lateral growth of the grains, does not occur during oxidation. However, our recent study clearly indicated that the columnar grains of Al2O3 formed on Fe−Cr−Al alloys without reactive elements showed grain coarsening with oxidation time [66].

Fig. 4.13 shows the size of columnar grains of the Al2O3 scale formed on Fe−Cr−Al alloy without and with different zirconium contents as a function of the distance from the outer/inner interface at 1000∘C. Figs. 4.13 (a), (b) and (c) show the results for the Fe−20Cr−10Al (at.%), Fe−20Cr−10Al−0.01Zr (at.%) and Fe−20Cr−10Al−0.1Zr (at.%) respectively as a function of distance from the outer/inner interface at 1000∘C in air [66]. The solid and blank circles in the figures indicate the samples after oxidation for 100 and 400 h , respectively. It can be seen that grain coarsening is not observed for up to 400 h on the alloy with 0.1at.% of zirconium and for up to 100 h on the alloy with 0.01at.% of zirconium. However, grain coarsening is observed on the alloy without zirconium after 100 h and alloy with 0.01at.% of zirconium after 400 h . Apparently zirconium prohibits the grain coarsening of Al2O3 during oxidation, resulting in smaller columnar grain sizes. A similar effect of zirconium on the formation of a fine grained Al2O3 scale formed on Fe−Cr−Al−Y−Zr alloy at 1200∘C in air was also observed [67]. Zirconium obviously has an effect on the microstructure of Al2O3 scale. Moreover, not only the growth rate of the scale, but also different properties such as its adhesion and creep deformation should be controlled by the microstructure of Al2O3 scale. Thus, it would be interesting to investigate the effect of different reactive elements on the microstructure of the Al2O3 scale.

4.5 Alumina-Forming Austenitic Stainless Steels

As mentioned previously, Al2O3 scale formation is difficult at relatively lower temperatures. Therefore Al2O3-forming stainless steels are expected to be used at much higher temperatures compared to Cr2O3-forming steels. The temperature capability of most of the stainless steels with ferritic substrate is much lower than that of austenitic steels, and as a result ferritic FeCrAl alloys are not used at higher temperatures even they have excellent oxidation resistance. Structural components such as boiler tubes and turbine blades need to have sufficient high-temperature strength. These components should employ the austenitic matrix, which has much higher
mechanical strength. However, commercial " Al2O3-forming austenitic stainless steels" have not been well developed.

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Fig. 4.13. Grain size of outer equiaxed and inner columnar Al2O3 grains formed on (a) Fe−20Cr−10Al (at.%), (b) Fe−20Cr−10Al−0.01Zr (at.%) and (c) Fe−20Cr−10Al−0.1Zr (at.%) as a function of distance from the surface at 1000∘C in air after oxidation for 100 h (solid circles) and 400 h (blank circles) [66].

Fig. 4. 13. Grain size of outer equiaxed and inner columnar A l 2 O 3 grains formed on (a) F e − 20 C r − 10 A l (at.%), (b) F e − 20 C r − 10 A l − 0.01 Z r (at.%) and (c) F e − 20 C r − 10 A l − 0.1 Z r (at.%) as a function of distance from the surface at 1000 ∘ C in air after oxidation for 100 h (solid circles) and 400 h (blank circles) [66].

Compared to ferritic stainless steels, austenitic steels need higher aluminium contents to form an exclusive Al2O3 scale because the much lower diffusivity of aluminium in the fcc structure than that in the bcc is expected [68, 69]. Similar to the case of ferritic steels, maximum aluminium additions to the austenitic substrate are also very limited. Aluminium is very strong bcc stabiliser in this system, therefore, addition of aluminium in the austenitic Fe−Ni−Cr substrate results in A 2 -bcc and/or B2-bcc phase formation. From the Fe−25Ni−15Cr−Al (at.%) phase diagram calculated by FactSage software [70], although the accuracy of this diagram is low particularly at lower temperature region, it is clear that aluminium addition promotes formation of A2 and B2 phases. A2 phase formation decreases high-temperature mechanical properties. Precipitation of B2 phase in the fcc substrate could increase the mechanical strength. However, it also decreases the oxidation performance because aluminium concentration in the substrate is decreased due to aluminium consumption by the B2 phase formation. Fig. 4.14 shows the oxidation behaviour of Ni-20Cr-10Al (at.%) with different iron contents [71]. All alloys have an austenite matrix with B2 precipitates at 1000∘C. Oxidation mass gain increases with increasing iron content in the alloy. The aluminium content in the substrate was found to decrease with the increase of iron content in the alloy i.e. from 8at.% for 20Fe−50Ni−20Cr−10Al (at.%) to 6at.% for 40Fe−30Ni−20Cr−10Al (at.%). This decrease was due to the increase in the volume fraction of B 2 phase from 0.08 to 0.15%, which results in formation of a protective Al2O3 scale difficult. B2−(Ni,Fe)Al phase formation is clearly detrimental for an exclusive Al2O3 scale formation.

[picture]

Fig. 4.14. Oxidation kinetics of Fe−Ni−20Cr−10Al (at.%) with different Fe/Ni ratios at 1000∘C in air [71].

Fig. 4. 14. Oxidation kinetics of F e − N i − 20 C r − 10 A l (at.%) with different F e / N i ratios at 1000 ∘ C in air [71].

Brady et al. [72-74] in Oak Ridge National Laboratory proposed Al2O3-forming wrought and cast austenitic steels. Their proposed steels contain 20−30wt.% of nickel, 12−20wt% of chromium, 0.6−3.0wt.% of niobium, 0.1wt.% of carbon and 2.5−5.0wt.% of aluminium and are strengthened by MC carbide and Fe2Nb laves phases. These alloys show good high-temperature creep strength but aluminium level is very critical for an exclusive Al2O3 scale formation, suggesting that a higher aluminium level is necessary. Although they supposed that B2−NiAl may act as a reservoir compensating the aluminium consumption by Al2O3 scale formation, higher aluminium addition promoted the B2-NiAl formation. This result is different from our previous study as shown in Fig. 4.14.

We are now investigating new Al2O3-forming austenitic stainless steels [75]. The target of this steel development is to alternate current conventional Cr2O3-forming austenitic stainless steels such as grade 310, 308 steels, which are produced by conventional techniques such as hot-rolling. The strategy for our alloy design is to produce steels with B2-free austenite matrix, allowing production by conventional processing. Fig. 4.15 shows the effect of copper content on oxidation mass gain of copper-modified Fe−17Ni−17Cr−7Al−0.06Zr (at.%) alloy at 1000∘C in air [75]. Small addition of copper, ~4at.%, promotes a protective Al2O3 scale formation and was found to decrease the critical aluminium content for an Al2O3 scale formation, although the detailed mechanism of this beneficial effect of copper is not known yet. The oxidation resistance of these new austenitic steels can be increased by an exclusive Al2O3 scale formation.

[picture]

Fig. 4.15. Effect of copper on the oxidation mass gain of Fe−17Ni−17Cr−8Al−xCu−0.06Zr (at.%) alloys after 100 h of oxidation at 1000∘C in air. (a) Oxidation mass gain as a function of copper content, and cross-sections of (b) 0 at. % copper and (c) 5 at. % copper alloys after oxidation. Reproduced and adapted from S. Hayashi et al., High-temperature oxidation of Cu-modified Fe-Ni-Cr-Al alloys, in: Y. Murata et al. (Eds.), Proceedings of the First International Conference on Advanced High-Temperature Materials Technology for Sustainable and Reliable Power Engineering, Japan, 2015, pp. 371-374 [75].

Fig. 4. 15. Effect of copper on the oxidation mass gain of F e − 17 N i − 17 C r − 8 A l − x C u − 0.06 Z r (at.%) alloys after 100 h of oxidation at 1000 ∘ C in air. (a) Oxidation mass gain as a function of copper content, and cross-sections of (b) 0 at. % copper and (c) 5 at. % copper alloys after oxidation. Reproduced and adapted from S. Hayashi et al., High-temperature oxidation of Cu-modified Fe-Ni-Cr-Al alloys, in: Y. Murata et al. (Eds.), Proceedings of the First International Conference on Advanced High-Temperature Materials Technology for Sustainable and Reliable Power Engineering, Japan, 2015, pp. 371-374 [75].

4.6 Summary

The first part of this chapter introduces the defect structures of thermal oxides on iron and chromium. The typical thermal oxides formed on iron at temperatures higher than the eutectoid one consist of the oxygen-deficient haematite sitting on the metal-deficient magnetite which sits on the major metal-deficient wüstite. The thermal oxide on chromium is duplex chromia with an internal (resp. external) n-equiaxed (resp. p-columnar) subscales. The Brouwer diagram for chromia formation is also described.

The second part of this chapter reviews the effects of alloying elements on the high temperature oxidation of the chromia-forming stainless steels. Silicon tends to reduce the steel oxidation rate but not in a monotonic manner. Titanium and niobium can help suppress the breakaway oxidation for the austenitic stainless steel since it significantly increases the ratio between the diffusivity of chromium in steel which represents the chromium supply from metal to metal/scale interface and the oxidation parabolic rate constant which represents the chromium consumption rate for oxidation. Titanium can form TiO2 at the internal metal/scale interface and serves as a mechanical anchor to hold oxide and metal together resulting the improved scale adhesion. The addition of titanium and niobium together into the steel can form TiO2 but also with Fe2Nb particles which present between the titania and the metal substrate. This Fe−Nb intermetallic compound can reduce the scale adhesion by reducing the TiO2 mechanical keying effect. For the AISI 444, Fe-Mo intermetallic compound can also form at the metal/scale interface and consequently reduces the scale adhesion.

The third part of this chapter reviews the high temperature oxidation behaviour of aluminaforming stainless steels. The oxidation rate of this steel is significantly lower than that of the chromia-forming stainless steel. However, the application of this steel is limited especially by the formation of metastable alumina and the limited aluminium content that can be alloyed without degrading the steel mechanical properties. High aluminium content at the surface can delay the transformation of metastable to stable alumina. The presence of Fe2O3 on the metal surface can serve as a template for alumina nucleation, thus accelerating the transformation. For the aluminaforming ferritic stainless steel, the addition of chromium can help reduce the critical content of aluminium added to the steel in order to form alumina at high temperatures. However, care must be taken when chromium is added with relatively high content since it can form the brittle chromiumrich bcc phase at 475∘C. Reactive elements can also be added to improve the scale adhesion and reduce and the steel oxidation rate. Refractory metals e.g. molybdenum, tungsten and tantalum can be added to the steel to increase its strength. However, the addition of refractory metal like molybdenum can increase the oxidation rate of stainless steel containing reactive elements e.g. zirconium and lanthanum. The alumina-forming austenitic stainless steel has capacity to be used at high temperatures because of its superior oxidation resistance and mechanical properties at high temperatures. However, since the diffusivity of aluminium in austenite is relatively low, the high aluminium content is needed in order to form alumina at high temperatures. This high aluminium content can form B2 phase which can decrease the aluminium content in the matrix and therefore reduces the oxidation performance. The reduced iron content in Ni−20Cr−10Al (at.%) can reduce the volume fraction of B2 phase and therefore gives the lower oxidation rate. Some aluminaforming austenitic grades have been developed e.g. the Fe−17Ni−17Cr−7Al−0.06Zr−0.6Cu (at.%). Addition of copper in the mentioned alloy can reduce the critical content of aluminium to form alumina at high temperatures.

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Solid State Phenomena, ISSN: 1662-9779, Vol. 300, pp 107-134

doi: 10.4028/www.scientific.net/SSP.300.107

© 2020 The Author(s). Published by Trans Tech Publications Ltd, Switzerland.

Submitted:
2019-04-22

Revised:
2019-07-24

Accepted:
2019-08-08

Online:
2020-02-12

Effect of Water Vapour on the High Temperature Oxidation of Stainless Steels

The original version of this paper is available on https://www.scientific.net/SSP.300.107.pdf

Somrerk Chandra-ambhorn 1,a, Patthranit Wongpromrat 2, b, Thammaporn Thublaor 1,c and Walairat Chandra-ambhorn 2, d*

1 High Temperature Corrosion Research Centre, Department of Materials and Production Technology Engineering, Faculty of Engineering, King Mongkut's University of Technology North Bangkok, 1518, Pracharat 1 Road, Bangsue, Bangkok, 10800, Thailand

2 Department of Chemical Engineering, Faculty of Engineering, King Mongkut's Institute of Technology Ladkrabang, Chalongkrung 1 Road, Lat Krabang, Bangkok, 10520, Thailand

a  somrerk.c@eng.kmutnb.ac.th, b  patthranit.wo@kmitl.ac.th, c  thammaporn.tt@gmail.com, d walairat.ch@kmitl.ac.th

Keywords: high temperature oxidation, stainless steel, alloying element, chromia, alumina

Abstract

This chapter primarily reviews the nature of water vapour when it presents in bulk gas. The change in a ratio between water vapour and corresponding dissociated hydrogen, which determine the thermodynamic stability of the oxide formation, is analysed when the oxidation kinetics are linear and parabolic. When water vapour reaches the solid/gas interface, chromium species volatilisation and oxidation controlled by surface reaction can occur. The adsorbed water vapour can be further incorporated into the oxide possibly in the form of hydrogen defects. The role of these defects on altering the defect structure of the oxide is discussed. Finally, characteristics of the oxide scale on stainless steels formed in the atmosphere containing water vapour are reviewed.

CHAPTER 5

The original version of this paper is available on https://www.scientific.net/SSP.300.107.pdf

5.1 Water Vapour in the Bulk Gas

In practice, water vapour presents in various atmospheres such as in ambient air of the industrial processes [1-4] or in atmospheres on cathode or anode sides of solid oxide fuel cells [5-7]. The presence of water vapour could significantly change the oxidation rate and behaviour of the metals. In the classic 2001 paper titled "The kinetic behaviour of metals in water vapour at high temperature: Can general rules be proposed?", Galerie et al. summarised that "water vapour may react more rapidly with metals than oxygen provided that the formed oxide can grow by transport of hydroxyl ion. If such a transport is not possible, the rate of the reaction will depend on the surface acidity of the oxide: less acidic is the oxide, slower is the reaction." [8]. From the time of that publication, the extensive research on water vapour effect on stainless steel oxidation has been conducted under his guidance. A part of those works is reviewed here together with the other ones relevant to this topic. The chapter starts from the analysis of water vapour in the bulk gas. It is then followed by the situations when water vapour arrives at the scale/gas interface and when the adsorbed water is incorporated into the oxide scale in the form of hydrogen defect.

As for the water vapour in the bulk gas, when water vapour presents in the atmosphere, it can dissociate giving H2 and O2 according to the following global reaction:

H2O↔H2+12O2(5.1)

with

Kw=exp(−ΔGw∘RT)=pH2pO21/2pH2O(5.2)

Kw and ΔGw∘ are the equilibrium constant and standard Gibbs free energy of water dissociation reaction as Reaction 5.1, respectively. If the bulk gas contains water vapour mixed with an inert gas, the values of the equilibrium partial pressure of oxygen, ( pO2 ) eq , dissociated from Reaction 5.1 can be calculated as explained in Section 3.2. Fig. 5.1 shows the temperature-dependent equilibrium partial pressure of oxygen dissociated from water vapour and from the Ar−H2O gas mixture. It can be used to identify the equilibrium oxygen partial pressure in Ar−H2O when the water vapour content in the atmosphere is varied.

[picture]

Fig. 5.1. Equilibrium partial pressure of oxygen in Ar−20%H2O gas mixture constructed using thermodynamic data from Kubaschewski et al. [9].

Fig. 5. 1. Equilibrium partial pressure of oxygen in A r − 20 % H 2 O gas mixture constructed using thermodynamic data from Kubaschewski et al. [9].

To check whether the oxygen dissociated from water vapour significantly affects the steel oxidation, Chandra-ambhorn et al. [4] oxidised a low carbon steel at 900∘C in Ar−20%H2O. As can be seen in Fig. 5.1, the equilibrium oxygen partial pressure is 8.45×10−7 bar in that atmosphere at 900∘C. In their experiment, this gas mixture was flowed with the rate of 6 L h−1 to oxidise the steel in a thermobalance where the mass change was recorded every 5.3 s . The calculation showed that, within 5.3 s , the oxygen gas dissociated from water vapour is supplied to the reactor only with the amount of 0.0024μ g. However, H2O in this atmosphere has the partial pressure of 0.2 bar. With the same gas flow rate, H2O supplies oxygen associated with its molecule to the reactor with the amount of 290μ g in 5.3 s . The mass gain measured by thermobalance in each 5.3 s was found to be in the range of 1.4-27.0 μg [4]. These results indicate that water vapour plays a dominant role to provide oxygen for oxidation much more than the one from oxygen gas which is dissociated from water vapour.

When water vapour takes part in oxidation process, water is consumed and hydrogen is produced. These phenomena could change the ratio of water vapour to hydrogen ( pH2O/pH2 ) in the atmosphere, and therefore possibly affect stability of the oxide formation. For the oxidation of metal M in H2O giving MaOb, the global oxidation reaction can be written as follows:

aM+bH2O↔MaOb+bH2(5.3)

Let dnH2O, in  is the differential of the number of mole of water vapour that enters the reactor, dnH2O,consumed  is the differential of the number of mole of water vapour consumed by oxidation and dnH2, released  is the differential of the number of mole of hydrogen molecule produced as a result of the oxidation. If it is assumed that all hydrogen produced by Reaction 5.3 is released to the atmosphere, the pH2O/pH2 at the time t, denoted as (pH2O/pH2)t, can be expressed as follows:

(pH2OpH2)t=dnH2O, in −dnH2O, consumed dnH2, released (5.4)

We can find out dnH2O, in from the experimental parameters using the following relation:

dnH2O, in =n˙H2O, in dt=pH2O, in V˙RT dt(5.5)

where pH2O, in  is the partial pressure of water vapour entering the reactor, V˙ is the gas flow rate, R is the ideal gas constant, T is the absolute temperature and dt is the differential of time. From Reaction 5.3, we see that the number of mole of H2O consumed equals the number of mole of O used to form the oxide which is the one obtained from the experiment, thus giving

dnH2O, consumed =dnmeasured (5.6)

where dnmeasured  is the differential of the number of mole measured from the thermogravimetric experiment. From Reaction 5.3, we can also see that the number of mole of H2O consumed equals the number of mole of H2 gas produced by the oxidation reaction. If it is assumed that all H2 molecules produced from Reaction 5.3 release to the atmosphere, we can write that

dnH2, released =dnH2O, consumed (5.7)

Substituting Eqs. 5.5-5.7 into Eq. 5.4 gives

(pH2OpH2)t=n˙H2O,indt dnmeasured −1(5.8)

This equation shows that dnmeasured , which depends on the oxidation kinetic behaviour of the metal influences (pH2O/pH2)t.

In the case of linear oxidation kinetics, the relation between mass gain ( Δm/A ) and time ( t ) is expressed as

(ΔmA)=k1t(5.9)

where Δm is the difference between the mass measured at the time t and the mass before oxidation, A is the area of the metal and k1 is the linear rate constant. We can further derive that

dnmeasured =AklMO dt(5.10)

where MO is the atomic mass of oxygen. Inserting Eq. 5.10 into Eq. 5.8 gives

(pH2OpH2)t=MOn˙H2O,inAkl−1(5.11)

From this equation, we can see that at a fixed flow rate and k1, the ( pH2O/pH2)t gas ratio does not change with time. Furthermore, higher is the k1, lower is the (pH2O/pH2)t.

In the case of parabolic oxidation, the relation between mass gain and time can be described as

(ΔmA)2=kpt+C(5.12)

where kp is the parabolic rate constant and C is a constant which could be obtained when fitting the experimental results. From this relation, it can be obtained that

dnmeasured =kpA dt2MOkpt+C(5.13)

Inserting Eq. 5.13 into Eq. 5.8 gives

(pH2OpH2)t=2MOn˙H2O, in kpt+CAkp−1(5.14)

From this equation, we can see that at the fixed gas flow rate, the (pH2O/pH2)t increases parabolically with time.

Fig. 5.2 demonstrates the variation of (pH2O/pH2)t as a function of time for steels oxidised at 900∘C in Ar−20%H2O [4]. In this case, the oxidation kinetics was found to be linear in the early stage, thus giving a constant (pH2O/pH2)t as a function of time. After a linear oxidation stage, the kinetics turns into parabolic, thus giving an increased value of ( pH2O/pH2)t as function of time in accordance with the relation in Eq. 5.14. For ( pH2O/pH2 ) in the range of 101 to 102.5(10−320), magnetite is a stable oxide phase. According to thermodynamic data from Ref. [9], this oxide phase is stable for pH2O/pH2 in the range of 5.5−22422. However, haematite was found in experiment even though it is thermodynamically stable at the pH2O/pH2 higher than 22422 [4]. It is possible that hydrogen produced by Reaction 5.3 is not totally released to the atmosphere. This thus causes, as described in Eq. 5.4, an increase of the (pH2O/pH2)t up to a value corresponding to the stability domain of haematite. Furthermore, it should be noted that the variation of pH2O/pH2 as analysed in
this section can be applied only if the gas pressure at the metal surface equals the bulk gas pressure value [10].

[picture]

Fig. 5.2. Variation of (pH2O/pH2)t as a function of time in the oxidation of steels at 900∘C in Ar−20%H2O. Redrawn from S. Chandra-ambhorn et al., Corros. Sci. 87 (2014) 101-110 [4].

Fig. 5. 2. Variation of ( p H 2 O / p H 2 ) t as a function of time in the oxidation of steels at 900 ∘ C in A r − 20 % H 2 O . Redrawn from S. Chandra-ambhorn et al., Corros. Sci. 87 (2014) 101-110 [4].

5.2 Water Vapour Effect at the Solid/Gas Interface

5.2.1 Chromium Species Volatilisation

When water vapour from the bulk gas arrives at the metal surface it can oxidise the metal and also volatilise the oxide formed. Chromium species volatilisation, so-called chromium species evaporation or vaporisation, from chromia grown on chromium alloys in oxidising conditions at elevated temperatures has been observed by many researchers since 1960s. When chromium alloy was exposed to oxidising atmosphere at high temperatures, the CrO3 crystal has sometimes been found to deposit at cooler parts of equipment due to the reduction of CrO3 back to Cr2O3 [11, 12]. The research group of Caplan and Cohen [11] was one of the early groups who studied the chromium volatilisation by observing the weight loss of Cr2O3 pellets exposed to various dry and humidified atmospheres at 1000−1200∘C for 5 to 100 h . In their experiment, the gas flow rate of gas was varied from 10 to 200 mL min−1 at room temperature, corresponding to a linear velocity in the range of 0.025−0.5 cm s−1. They found for both dry and humidified oxygen that the rate of weight loss of Cr2O3 increased with increasing temperatures while no weight change occurred under dry and humidified argon atmospheres. They also found that the rate of weight loss of Cr2O3 exposed to humidified oxygen was about two times higher than in dry oxygen.

Tedmon [13] derived the model equation to describe the oxidation behaviour of chromium and iron-chromium alloys based on the assumption that not only chromium oxidation into Cr2O3 occurred when exposing to oxidising atmospheres at high temperatures but also the volatilisation of Cr2O3 formed into volatile species. By considering the two phenomena simultaneously, Tedmon found that a transition from a parabolic oxidation kinetics into linear one can be observed. This model is in good agreement with the experimental data.

Stearns et al. [14] measured the volatilisation rates of Cr2O3 at low pressure ( 1.51×10−4 atm ) in gas steam with different partial pressures of oxygen over the temperature range of 1000−1300∘C. Taking into account both diffusion and reaction effects, the reaction-controlled rates of chromium species volatilisation were calculated. They observed that the Cr2O3 volatilisation rate decreased
with decreasing oxygen partial pressure. Furthermore, they found a volatilisation rate dependence to the oxygen partial pressure to the power 0.75 . For temperature dependence study, the Arrhenius relation between the Cr2O3 volatilisation rate versus reciprocal temperature gave the experimental activation energy of 251±21 kJ mol−1.

Fryburg et al. [15] investigated the formation of volatile species vaporised from various types of metals and alloys including chromium at 1050∘C in oxidising atmospheres containing water vapour. They found that besides the volatile species of CrO3, under the oxidising atmosphere containing water vapours, CrO2(OH)2 was also formed resulting in increasing the rate of chromium species evaporation.

Asteman et al. [16, 17] investigated chromium volatilisation from 304L stainless steel exposed to dry O2 and O2 containing 10% water vapour at 873∘C from 1 to 672 h . The weight gain of the sample exposed to dry O2 increased with increasing exposure time. They also observed that the presence of water vapour in O2 did not significantly affect the weight gain for samples exposed for less than 168 h . However, for longer exposure time the weight gain of the samples exposed to wet O2 was considerably lower than that exposed to dry O2. They explained that after prolonged exposure to humidified oxygen, chromium species volatilisation and chromium-manganese spinel formation led to the decrease in the rate of weight gain. Chromium volatile species was suggested to be CrO2(OH)2.

Asteman et al. [18, 19] also investigated the effect of chromium species volatilisation on the oxidation behaviour of 304 L and 310 steels in water vapour and oxygen mixtures at 500−800∘C. In their experiment, the gas flow rate was varied. The results revealed that the oxidation strongly depended on the gas flow rate. In the oxidation of AISI 310 at 600∘C in humidified oxygen using the gas flow rate in the range of 0.3−10 cm s−1, they found that at relatively low gas flow rate the protective α−(Cr,Fe)2O3 oxide was formed with the outer part depleted in chromium. At high flow rates beyond the critical value the protective oxide could be failed [19]. In that case, the α−Fe2O3/(Cr,Fe)2O3 with the thickness higher than 5μ m was found at the heart of the alloy grain while the α−(Cr,Fe)2O3 close to the steel grain boundaries remained partially protective. They further suggested that durability of stainless steel in humidified oxygen is connected to its ability to compensate chromium loss by sublimation.

Yamauchi et al. [20] also investigated the effect of water vapour on the volatilisation of chromia. Their results were in good agreement with Asteman et al. [16, 17]. Furthermore, they compared the volatilisation behaviour of chromia under N2−O2−H2O atmosphere with those under N2−O2 and N2−H2O atmospheres within the temperature range of 1173−1473 K. In their work, samples were prepared by hot-pressing chromia powder in a graphite die using sintering temperature of 1673 K . Before the test, all samples were polished with a fine diamond abrasive and then ultrasonically cleaned in ethanol. Mass loss of each sample was measured for various exposure times. It was observed that the mass losses in N2−O2 and N2−H2O atmospheres were almost identical and very low in comparison with that in N2−O2−H2O atmosphere. From this observation, they suggested that the chromium species volatilisation based on the reaction that both oxygen and water vapour were involved. Considering the effect of oxygen partial pressure on the mass loss, they found that the rate of mass loss increased with increasing oxygen partial pressure.

Gindorf et al. [21] developed the transpiration apparatus used to measure the amount of vaporised chromium species. They measured the amount of chromium species evaporated from Cr2O3 powder at 1223 K at different flow rates of humid air. They observed that a chromium sublimation rate increases with increasing flow rate. However, the chromium volatilisation rate remains unchanged for flow rate higher than 400 mL min−1. Gindorf et al. also performed the thermodynamic computation using databases of IVTANTHERMO and Ebbinghaus and found that most abundant volatilisation reactions should be as follows:

Cr2O3( s)+32O2( g)→2CrO3( g)Cr2O3( s)+32O2( g)+2H2O(g)→2CrO2(OH)2( g)Cr2O3( s)+O2( g)+H2O(g)→2CrO2(OH)(g)

(5.15)(5.16)(5.17)

Furthermore, they compared their experimental results with the computational one and found that, at low partial pressure of water vapour, the evaporated chromium species is in the form of CrO3 while, at high partial pressure of water vapour, CrO2(OH)2 should be major chromium volatile species.

Opila et al. [22] studied the effect of oxygen and water vapour partial pressure on the amount of produced chromium volatile species. They found that partial pressure of the chromium volatile species depended on the partial pressure of water vapour to the power 0.96 and partial pressure of O2 to the power 0.77 . These values corresponded to the power law exponents of the rate law of Reaction 5.16. Kurokawa et al. [23] also investigated the effect of gas flow rate on the amount of chromium species volatilisation from chromium oxide pellets using experimental setup inspired by the transpiration apparatus of Gindorf's research group. Their results are in agreement with the ones by Gindorf et al. [21].

Stanislowski et al. [24] studied chromium species volatilisation on the surface of chromiaforming Ducrolloy exposed to humidified air at 800∘C by varying the flow rate from laminar to turbulent regime. The evolution of chromium volatilisation flux according gas flow rate can be divided into three zones called "equilibrium volatilisation", "non-equilibrium volatilisation" and "volatilisation limit". For the "equilibrium volatilisation" related to the lower flow rate part of laminar flow region, the reactions are at equilibrium. The rates of forward and backward reactions are equivalent. Hence, chromium species volatilisation rate should be equal to the one calculated by thermodynamics, and thus increases with increasing the gas flow rate. In the "non-equilibrium volatilisation" zone that covers higher flow rate part of laminar flow region and the beginning part of turbulent zone, a thin film of gas called "boundary layer" forms on the surface. A concentration gradient happens due to the diffusion limit of chemical species with boundary layer. In laminar regime, both inside and outside boundary layers have laminar flow pattern. The faster the gas flows, the thinner the boundary layer is formed resulting in less concentration gradient. At the beginning when the laminar flow changes into the turbulent flow due to the increase of fluid flow rate, the flow pattern at the top part of boundary layer becomes turbulent while the bottom part of boundary layer which is close to the surface still maintains laminar flow pattern. In the turbulent zone of boundary layer, a turbulent stream causes the fluid to well mix. Therefore, a concentration gradient in the turbulent zone becomes less as the degree of turbulent increases. This may be the reason why chromium species volatilisation rate increases with increasing rate at first and then the increasing rate becomes slower as the flow is more turbulent. For the last zone, volatilisation limit zone, the flow is fully turbulent and, instead of diffusion limit, chromium volatilisation is limited by kinetics.

In our research group, Wongpromrat et al. [6] developed the experimental set-up as described in Section 3.3 in Chapter 3 to measure the chromium species volatilisation from AISI 430 stainless steel subjected to O2−5%H2O at 800∘C. It was found that in the linear gas velocity range of 0.8−1.95 cm s−1, the chromium species volatilisation rate increases with the increased velocity as shown in Fig. 5.3. This indicates that the mass transfer rate of the chromium volatile specie is slower than the surface reaction rate. Thus the mass transfer of the volatile specie across the oxide/gas interface is a rate-determining step.

[picture]

Fig. 5.3. Mass flux of the Cr volatile species ( j ) measured from AISI 430 oxidised at 800∘C in O2−5%H2O for 24 h [6] compared with the mass fluxes calculated using thermodynamic data from Opila et al. [22] and Stanislowski et al. [24]. Redrawn and adapted from W. Wongpromrat et al., Oxid. Met. 79 (2013) 529-540.

Fig. 5. 3. Mass flux of the Cr volatile species ( j ) measured from AISI 430 oxidised at 800 ∘ C in O 2 − 5 % H 2 O for 24 h [6] compared with the mass fluxes calculated using thermodynamic data from Opila et al. [22] and Stanislowski et al. [24]. Redrawn and adapted from W. Wongpromrat et al., Oxid. Met. 79 (2013) 529-540.

To compare the results with the theoretical value, mass flux of the chromium volatile species ( j ) is calculated using the following equation [6,25]:

j=kmRT(psurface −pbulk )(5.18)

Here, km is the mass transfer coefficient of the chromium volatile species, R is the ideal gas constant, T is the absolute temperature. psurface  and pbulk  are the partial pressures of the volatile species at the solid surface and in the bulk gas respectively.

In the following calculation of the mass flux, it is assumed that psurface  is significantly higher than pbulk  and thus the pbulk  term is neglected. The values of psurface  can be obtained by using thermodynamic data reported in literature. As for km, its value can be obtained from Eq. 5.19 [6, 25] where v is the kinematic viscosity, vl is the linear velocity of the gas and l is the sample length. DAB is the binary gas (O2−CrO2(OH)2) diffusion coefficient which can be calculated using the ChapmanEnskog kinetic theory [26]. The calculation shows that the values of km are 0.5−1.5 cm s−1 for the linear velocity of the gas in the range of 0.85−1.95 cm s−1.

km=0.664(DABv)16(vll)12(5.19)

By determining the values of km,psurface  and pbulk  previously discussed as well as T, the mass flux of the volatile specie according to Eq. 5.18 can be calculated and plotted in Fig. 5.3. In that figure, there are also two calculated lines which were obtained by using different thermodynamic data i.e. from Opila et al. [22] and Stanislowski et al. [24] in order to find out the value of psurface . It can be seen that our measured mass fluxes are lower than the calculated ones. The measured mass fluxes were obtained from the chromium volatile species in the bulk gas which was dissolved in water and washing solution as described in Section 3.3 in Chapter 3, while the calculated mass flux is the flux that transfers from the solid surface to the bulk gas. The volatilisation in this case is controlled by
mass transfer of the volatile species, therefore the mass flux of the volatile species in the bulk gas which was the measured value was lower than the mass flux passing through the oxide/gas interface which is calculated using Eq. 5.18.

Wongpromrat et al. [7] also investigated the effect of gas velocity on the chromium species volatilisation rate of the AISI 441 stainless steel up to the linear gas velocity of 10 cm s−1. The result is shown in Fig. 5.4. As seen from the figure, the pattern of the graph is similar to the results of Gindorf et al. [21] and Kurokawa et al. [23] i.e. at the low linear velocity range of 1−4 cm s−1, volatilisation rate increases with increasing gas velocity. Meanwhile, at the gas velocity more than 4 cm s−1, volatilisation rate seems unchanged. They explained that, at low gas velocity, volatilisation rate is limited by the diffusion rate of chromium volatile species through the boundary layer, while at high gas velocity, diffusion rate becomes faster resulting in constant volatilisation rate as it is limited by surface reaction. Comparing with the result of Stanislowski et al. [24], it is possible that Gindorf et al. [21] and Kurokawa et al. [23] and Wongpromrat et al. [7] performed the experiments at low velocity which might lie on "laminar region" discussed in Stanislowski et al. [24] as they performed the experiments with low gas velocities.

[picture]

Fig. 5.4. Chromium species volatilisation rate of the AISI 441 oxidised at 800∘C in O2−5%H2O for 96 h as a function of linear gas velocity. Redrawn from W. Wongpromrat et al., Corros. Sci. 106 (2016) 172-178 [7].

Fig. 5. 4. Chromium species volatilisation rate of the AISI 441 oxidised at 800 ∘ C in O 2 − 5 % H 2 O for 96 h as a function of linear gas velocity. Redrawn from W. Wongpromrat et al., Corros. Sci. 106 (2016) 172-178 [7].

Expecting to reduce the chromium species volatilisation rate of AISI 441, Wongpromrat et al. coated their specimens with Mn−Co spinel and found that unlike the case of AISI 430,Mn−Co spinel coating could not suppress the chromium volatilisation [6]. Therefore, under the supervision of Professor Alain Galerie and the collaboration between KMITL, KMUTNB (Bangkok, Thailand) and SIMaP (Grenoble, France), instead of using conventional way like coating the steel surface with ceramics, Wongpromrat et al. [7] applied the pre-oxidation method to form thin pre-oxidation film on AISI 441. As seen in Figs. 5.5 and 5.6, superior suppression of chromium species volatilisation is obtained when flowing 5%H2O in O2 at low rates (0−4 cm s−1) for 96 h through the pre-oxidised specimens in argon at 250∘C and in CO2/CO at 250 and 850∘C. This means that the pre-oxidation could be a promising method to inhibit chromium species volatilisation from AISI 441 which will probably be able to use as interconnector of solid oxide electrochemical devices. The possible mechanism for this suppression is described in Section 6.4 in Chapter 6 concerning the chromium volatilisation of stainless steel SOFC interconnect.

[picture]

Fig. 5.5. Chromium species volatilisation rate of the AISI 441 without and with pre-oxidation in argon at 250∘C after exposure to O2−5%H2O at 800∘C for 96 h . Redrawn from W. Wongpromrat et al., Corros. Sci. 106 (2016) 172-178 [7].

Fig. 5. 5. Chromium species volatilisation rate of the AISI 441 without and with pre-oxidation in argon at 250 ∘ C after exposure to O 2 − 5 % H 2 O at 800 ∘ C for 96 h . Redrawn from W. Wongpromrat et al., Corros. Sci. 106 (2016) 172-178 [7].

[picture]

Fig. 5.6. Chromium species volatilisation rate of AISI 441 without and with pre-oxidation in CO/CO2 at 250 and 850∘C after exposure to O2−5%H2O at 800∘C for 96 h . Redrawn from W. Wongpromrat et al., Corros. Sci. 106 (2016) 172-178 [7].

Fig. 5. 6. Chromium species volatilisation rate of AISI 441 without and with pre-oxidation in C O / C O 2 at 250 and 850 ∘ C after exposure to O 2 − 5 % H 2 O at 800 ∘ C for 96 h . Redrawn from W. Wongpromrat et al., Corros. Sci. 106 (2016) 172-178 [7].

5.2.2 Oxidation Controlled by Surface Reaction

Water vapour could not only volatilise the oxide on the surface but also oxidise the metal. When water vapour does not impact the diffusion process or modify the microstructure of the oxide, the surface reaction can be a rate determining step for the oxidation process [27]. In pure water vapour, Galerie et al. [8] and Young [27] suggested that water vapour can be dissociated giving OH
adsorbed on the site s(OH−s) and H2 at the solid/gas interface according to Reaction 5.20. The adsorbed OH could further dissociate giving the adsorbed O and H2 according to Reaction 5.21. They suggested that the energy needed to break the first OH bond is significantly lower than the one needed to break the second OH bond. As a result, the covering of hydroxide on the surface is more probable even at room temperature, while the covering of oxygen on the surface is less likely [8]. The oxygen species need the surface that can tightly fix the OH particle by its oxygen electron pair so that, by the great induced polarisation, the second OH bond can be broken; and in such case the surface is acidic [8]. Galerie et al. [8] plotted the evolution of a linear rate constant of metals oxidised in water vapour with the cation hydration enthalpies, giving the linear relationship between these two variables as shown in Fig. 5.7. It is noted that the oxidation rate of chromium and niobium are in fact parabolic, but a linear mean rate is assumed in the first 1 h [8]. From this figure, they proposed to use the cation hydration enthalpy to assess the acidity of the oxide in the manner that more negative value of the cation hydration enthalpy indicates more acidity of the oxide, and as a consequence the faster oxidation rate.

H2O(g)+s→OH−s+12H2( g)OH−s→O−s+12H2( g)

(5.20)(5.21)

[picture]

Fig. 5.7. Relation between the linear rate constant for the oxidation of metals in water vapour at 800∘C and the cation hydration enthalpies. Redrawn from A. Galerie et al., Mater. Sci. Forum 369-372 (2001) 231-238 [8].

Fig. 5. 7. Relation between the linear rate constant for the oxidation of metals in water vapour at 800 ∘ C and the cation hydration enthalpies. Redrawn from A. Galerie et al., Mater. Sci. Forum 369-372 (2001) 231-238 [8].

However, the prediction of the oxidation rate using the hydration enthalpy is limited by many reasons such as the oxides formed in oxygen and in water vapour must be in the same form so that the comparison is relevant [8]. The prediction cannot be used for the case of vanadium and niobium since these metals form V2O5 and Nb2O5 in oxygen, while they form VO2 and NbO2 in water vapour [8]. In the case of chromia, the diffusion in this oxide is relatively slow and remains a rate
determining step for the oxidation in water vapour [27], even though in some other atmospheres like carbon-containing ones, the surface reaction can be a rate determining step for stainless steel oxidation [28-30] as described in Section 6.3.2.2 in Chapter 6. Furthermore, water vapour significantly changes the defect structure and therefore the diffusion process, as well as the microstructure of Cr2O3. These changes influence the oxidation mechanism as will be explained in the next section.

5.3 Water Vapour Effect on the Diffusion-Controlled Oxidation through Chromia

5.3.1 Diffusion-Controlled Oxidation Phenomena

Figs. 5.8-5.10 exemplifies the effect of water vapour on the oxidation of iron, chromium and iron-chromium alloys at high temperatures. Fig. 5.8 shows the mass gain of iron when it was oxidised in oxygen without and with water vapour [31]. It can be seen that, at relatively low temperature of 700∘C, the oxidation rates of iron in both atmospheres are nearly identical. However, at the higher temperatures of 800 and 900∘C, water vapour in the atmosphere accelerates the iron oxidation rate.

[picture]

Fig. 5.8. Mass gains of iron oxidised in dry and humidified oxygen at 700,800 and 900∘C. Redrawn and adapted from S. Chandra-ambhorn et al., Corros. Sci. 148 (2019) 355-365 [31].

Fig. 5. 8. Mass gains of iron oxidised in dry and humidified oxygen at 700,800 and 900 ∘ C . Redrawn and adapted from S. Chandra-ambhorn et al., Corros. Sci. 148 (2019) 355-365 [31].

For the chromia thermally grown on chromium, Hänsel et al. [32] conducted the oxidation test at 1000∘C in argon containing water vapour, oxygen and hydrogen. The oxidation kinetics was found to be parabolic. For the results at relatively high oxygen partial pressure i.e. 0.04 bar, the parabolic rate constant is independent on the water vapour content as shown in Fig. 5.9. However, at the relatively low oxygen partial pressure in the range of 10−15 to 10−14 bar, three important results were observed. First, the rate constant increases with the increased water vapour content at constant oxygen partial pressure as shown in Fig. 5.10(a). Second, the rate constant increases with the increased water vapour at constant hydrogen partial pressure as in Fig. 5.10(b). Third, the rate constant increases with the increased hydrogen partial pressure at constant oxygen partial pressure as in Fig. 5.10(c). It is noted that the oxygen partial pressure in Figs. 5.9 and 5.10 were calculated using thermodynamic data from Ref. [9].

[picture]

Fig. 5.9. Parabolic rate constant as a function of the partial pressure of water vapour at constant oxygen partial pressure for the chromium oxidation at 1000∘C in the relatively high oxygen partial pressure [32].

Fig. 5. 9. Parabolic rate constant as a function of the partial pressure of water vapour at constant oxygen partial pressure for the chromium oxidation at 1000 ∘ C in the relatively high oxygen partial pressure [32].

[picture]

Fig. 5.10. Parabolic rate constant as a function of (a) the partial pressure of water vapour at constant oxygen partial pressure (b) the partial pressure of water vapour content at constant hydrogen partial pressure, and (c) the hydrogen partial pressure at constant oxygen partial pressure for the chromium oxidation at 1000∘C in the relatively low oxygen partial pressure range [32].

Fig. 5. 10. Parabolic rate constant as a function of (a) the partial pressure of water vapour at constant oxygen partial pressure (b) the partial pressure of water vapour content at constant hydrogen partial pressure, and (c) the hydrogen partial pressure at constant oxygen partial pressure for the chromium oxidation at 1000 ∘ C in the relatively low oxygen partial pressure range [32].

Fig. 5.11 shows the oxidation behaviour of Fe−15.7Cr−8.5Mn stainless steel in dry and humidified oxygen at 700∘C [33]. The kinetics is also found to be parabolic with the enhanced rate when water vapour presents in the atmosphere. This result was also observed such as in the case of Fe- 17Cr−1Mn stainless steel [34]. According to Figs. 5.8-5.11, the oxidation rate is parabolic which indicates that the oxidation is controlled by the diffusion through the oxide.

[picture]

Fig. 5.11. Mass gains of Fe−15.7Cr−8.5Mn stainless steel oxidised in dry and humidified oxygen at 700∘C. Redrawn and adapted from S. Chandra-ambhorn, Corros. Sci. 148 (2019) 39-47 [33].

Fig. 5. 11. Mass gains of F e − 15.7 C r − 8.5 M n stainless steel oxidised in dry and humidified oxygen at 700 ∘ C . Redrawn and adapted from S. Chandra-ambhorn, Corros. Sci. 148 (2019) 39-47 [33].

To understand the role of water vapour on defect structure and therefore the oxidation mechanism, we start from the simple case of an oxide in the form of M2O3 where metal vacancy is a dominant defect at the outer part and oxygen vacancy is a major defect in the inner part of the oxide. This defect structure is in agreement with Cr2O3 formed in oxygen as described in Section 4.1.2 in Chapter 4. The relevant defect formation reactions are shown in Reactions 5.22 and 5.23. These two reactions give the Schottky defect reaction as shown in Reaction 5.24. The electronic defect reaction also exists in the oxide as shown in Reaction 5.25.

32O2↔3OO+2 VM′′′+6 h•32O2+3 VO••+6e′↔3OO nil ↔2 VM′′′+3 VO•• nil ↔e′+h•

(5.22)(5.23)(5.24)(5.25)

When water vapour presents in the atmosphere, water vapour might be incorporated in the oxide in many forms, at least in the following three cases: (i) an incorporation of water vapour in the form of the hydrogen interstitial with positively unit effective charge ( Hi• ) [35, 36], (ii) a formation of Hi• which is bounded with oxygen in the oxide giving the substitutional hydroxyl ion ( OHO• ) together with the hydrogen interstitial with negatively unit effective charge ( Hi′ ) [32], and (iii) a formation of the hydroxyl ion which plays a major role on enhancing the oxygen inward diffusion [8, 37]. Each possibility will be discussed as follows.

5.3.2 Defect Structural Analysis when 𝐇𝐢∘ Presents in 𝐌2𝐎3

Norby [35] proposed the beautiful analysis leading to the construction of the Brouwer diagram for M2O3 oxidised in humidified oxygen which is a framework of this section. It was argued that
water vapour in the atmosphere could dissociate giving hydrogen interstitial and liberate oxygen according to Reaction 5.26 [35]. From this reaction and Reaction 5.1, we may consider that water vapour could dissociate giving hydrogen gas according to Reaction 5.1 and the hydrogen gas dissolves into the oxide in the hydrogen interstitial form according to Reaction 5.27 [35].

12H2O↔Hi•+14O2+e′12H2↔Hi•+e′

(5.26)(5.27)

From Reactions 5.22 and 5.26, Reaction 5.28 showing the relation between metal vacancy and hydrogen interstitial can be obtained. Similarly, from Reactions 5.23 and 5.26 , we can have Reaction 5.29 which shows the relation between oxygen vacancy and hydrogen interstitial.

3H2O↔3OO+2 VM′′′+6Hi•3H2O+3 VO••↔3OO+6Hi•

(5.28)(5.29)

Equilibrium constants of some reactions used in the following analysis are given in Eqs. 5.30-5.35. K5.22, K5.23, K5.26, K5.28 are the equilibrium constants of Reactions 5.22, 5.23, 5.26 and 5.28, respectively. KS is the equilibrium of the Schottky defect reaction according to Reaction 5.24, while Ke is the equilibrium constant of the electronic defect reaction according to Reaction 5.25.

K5.22=[VM′′′]2[ h•]6pO23/2 K5.23=1[ VO••]3[e′]6pO23/2 K5.26=[Hi•][e′]pO21/4pH2O1/2 K5.28=[VM′′′]2[Hi•]6pH2O3 KS=[VM′′′]2[ VO••]3 Ke=[e′][h•]

(5.30)(5.31)(5.32)(5.33)(5.34)(5.35)

Norby [35] assumed that at relatively lower partial pressure of water, hydrogen interstitial is a minor defect (case I). Reaction 5.29 shows that, in this condition, the oxygen vacancy becomes dominant. Its charge is compensated by electron as can be seen from Reaction 5.23. When the partial pressure of water vapour increases, Norby proposed that hydrogen interstitial became dominant and was compensated by electron (case II) and eventually by metal vacancy (case III) as can be seen from Reaction 5.28. From these assumptions, the dependencies of defect concentrations with the gas partial pressures can be derived as follows.

For the case I corresponding to the condition that hydrogen interstitial is a minor defect, the electroneutrality condition is

[e′]=2[ VO••](5.36)

From Eqs. 5.31 and 5.36, it can be obtained that

[VO••]=122/3 K5.231/9·1pO21/6(5.37)

From Eqs. 5.34 and 5.37, we obtain that

[VM′′′]=(2 K5.231/6 KS1/2)·pO21/4(5.38)

From Eqs. 5.36 and 5.37, it can be obtained that

[e′]=21/3 K5.231/9·1pO21/6(5.39)

From Eqs. 5.35 and 5.39, it can be obtained that

[h•]=K5.231/9 Ke21/3·pO21/6(5.40)

From Eqs. 5.32 and 5.39, we can obtain that

[Hi•]=K5.231/9 K5.2621/3·pH2O1/2pO21/12(5.41)

For the case II corresponding to the condition that hydrogen interstitial is the major defect compensated by electron, the electroneutrality condition is

[e′]=[Hi•](5.42)

From Eqs. 5.32 and 5.42, the following relation is obtained:

[e′]=[Hi•]=K5.261/2·pH2O1/4pO21/8(5.43)

From Eqs. 5.35 and 5.43, it can be obtained that

[h•]=Ke K5.261/2·pO21/8pH2O1/4(5.44)

From Eqs. 5.31 and 5.43, we can obtain that

[VO••]=1 K5.231/3 K5.26·1pH2O1/2pO21/4(5.45)

From Eqs. 5.34 and 5.45, we obtain

[VM′′′]=(K5.231/2 K5.263/2 KS1/2)·pH2O3/4pO23/8(5.46)

For the case III corresponding to the condition that hydrogen interstitial is the major defect compensated by metal vacancy, the electroneutrality condition is

3[ VM′′′]=[Hi•](5.47)

From Eqs. 5.33 and 5.47, it can be obtained that

[VM′′′]=K5.281/833/4·pH2O3/8(5.48)

From Eqs. 5.47 and 5.48 we can obtain that

[Hi•]=(31/4 K5.281/8)·pH2O3/8(5.49)

From Eqs. 5.34 and 5.48 we obtain

[VO••]=31/2 KS1/3 K5.281/12·1pH2O1/4(5.50)

From Eqs. 5.32 and 5.49, it can be obtained that

[e′]=K5.2631/4 K5.281/8pH2O1/8pO21/4(5.51)

From Eqs. 5.35 and 5.51, we can obtain

[h•]=31/4 K5.281/8 Ke K5.26·pO21/4pH2O1/8(5.52)

From Eqs. 5.37-5.41, 5.43-5.46 and 5.48-5.52, the defect concentrations can also be plotted as a function of the oxygen partial pressure at a fixed water vapour pressure, giving the Brouwer diagram as shown in Fig. 5.12. This diagram is similar to the one presented by Norby [35] but with slight correction to make all lines follow Eqs. 5.36-5.52.

Based on this analysis, Chandra-ambhorn et al. [33] proposed the graphical representation of the role of water vapour on the defect concentrations and dominant regimes as illustrated in Fig. 5.13 [33]. When water vapour content in the atmosphere increases, the defect concentration lines shift from the solid to the dash ones. The shifts of these lines further change the dominant regime of each defect by shifting the vertical dotted lines which separate each dominant defect area.

[e′]=2[ VO••][e′]=[Hi•]3[ VM′′′]=[Hi•]

[picture]

Fig. 5.12. Brouwer diagram for M2O3 in humidified oxygen based on Norby's work [35]. Redrawn from S. Chandra-ambhorn et al., Corros. Sci. 148 (2019) 39-47 [33].

Fig. 5. 12. Brouwer diagram for M 2 O 3 in humidified oxygen based on Norby's work [35]. Redrawn from S. Chandra-ambhorn et al., Corros. Sci. 148 (2019) 39-47 [33].

[picture]

Fig. 5.13. The effect of water vapour on metal and oxygen vacancy concentrations in a Brouwer diagram. Redrawn from S. Chandra-ambhorn et al., Corros. Sci. 148 (2019) 39-47 [33].

Fig. 5. 13. The effect of water vapour on metal and oxygen vacancy concentrations in a Brouwer diagram. Redrawn from S. Chandra-ambhorn et al., Corros. Sci. 148 (2019) 39-47 [33].

It can be seen from Fig. 5.13 that in the relatively high water vapour content region (case III), when water vapour content increases the metal vacancy concentration increases and the dominant regime of this defect is expanded. If the metal vacancy is dominant for the oxide growth, the increase in its concentration enhances the oxidation rate of the metal.

5.3.3 Defect Structural Analysis when OH𝐎+ and 𝐇𝐢′ Presents in 𝐌2𝐎3

In this analysis, Hänsel et al. [32] proposed the defect situations in chromia, which are the framework of this section, to explain the results shown in Figs. 5.9 and 5.10. When VM′′′ is the dominant defect, the defect formation reaction is like Reaction 5.22. It is still considered that water vapour gives the hydrogen defect in the oxide. However, if hydrogen interstitial is bounded with the oxygen atom in the crystal, Reaction 5.27 could be re-written as Reaction 5.53 with the equilibrium constant shown in Eq. 5.54.

12H2+OO+h•↔OHO•K5.53=[OHO•][h•]pH21/2

(5.53)(5.54)

The charge balance condition in this case is written as

3[ VM′′′]=[OHO•]+[h•](5.55)

From Eqs. 5.30, 5.54 and 5.55, it can be obtained that

[VM′′′]=K5.221/833/4·pO23/16·(1+K5.53pH21/2)3/4(5.56)

From Eq. 5.56 and the equilibrium constant of water dissociation from Eq. 5.2, it can be obtained that

[VM′′′]=K5.221/8 Kw3/833/4·pH2O3/8pH23/8·(1+K5.53pH21/2)3/4(5.57)

When oxygen partial pressure is relatively high, it is presumed that the hydrogen partial pressure is low giving the following assumption [32]:

K5.53pH21/2=[OHO•][h•]≪1(5.58)

which further means

[OHO•]≪[h•](5.59)

By using the assumption in Eq. 5.58, Eq. 5.56 is reduced to be

[VM′′′]=K5.221/833/4·pO23/16(5.60)

This relation indicates that at relatively high oxygen partial pressure, the rate constant due to the metal vacancy is a function of the oxygen partial pressure only. Thus, varying water vapour content
in the atmosphere does not change the rate constant, which is congruent to the experimental results shown in Fig. 5.9.

However, in the case of relatively low oxygen partial pressure, the hydrogen partial pressure is presumed to be high giving the following assumption [32]:

K5.53pH21/2=[OHO•][h•]≫1(5.61)

which means

[OHO•]≫[h•](5.62)

By using the assumption in Eq. 5.61, Eq. 5.57 can be reduced to be

[VM′′′]=K5.221/8 K5.533/4 Kw3/833/4·pH2O3/8(5.63)

This relation indicates that at relatively low oxygen partial pressure, the defect concentration and therefore the parabolic rate constant increases with the increased water vapour content in the atmosphere. This analysis is in agreement with the experimental results in Figs. 5.10 (a) and (b). Nevertheless, this relation predicts that in this low oxygen partial pressure range the hydrogen partial pressure does not affect the rate constant, while the experimental results in Fig. 5.10(c) show that it does. In this case, the defect situation according to this model is probably oversimplified.

It is noted that when the assumptions according to Reaction 5.61 and 5.62 are applied, the electroneutrality condition according to Eq. 5.55 reduces to be 3[ VM′′′]=[OHO•]. Furthermore, the OHO• formation reaction i.e. Reaction 5.53 stems from the Hi• formation reaction i.e. Reaction 5.27 and therefore the dependency of these two defect concentrations on gas pressure are in the same form. As a result, the relationship of the metal vacancy and gas pressure according to Eq. 5.63 is equivalent to that derived from the case III where 3[ VM′′′]=[OHO•] and metal vacancy is dominant, i.e. Eq. 5.48 in Section 5.3.2.

Hänsel et al. [32] further considered the possibility that the metal interstitial ( Mi•• ) is a dominant defect. As described in Reaction 1.25 in Chapter 1, the defect reaction for M2O3 formation in the present case may be written as follows:

32O2+2Mi•••+6e′↔2MM+3OO(5.64)

with the equilibrium constant of

K5.64=1[Mi•••]2[e′]6pO23/2(5.65)

They also proposed that in the region near the metal/scale interface, the formation of negatively charged hydrogen interstitial might be possible by the following reaction:

12H2+e′↔Hi′(5.66)

with the equilibrium constant of

K5.66=[Hi′][e′]pH21/2(5.67)

The charge balance condition in this case is

3[Mi•••]=[e′]+[Hi′](5.68)

From Eqs. 5.65, 5.67 and 5.68, it can be obtained that

[Mi••]=(133/4 K5.641/8)·1pO23/16·(1+K5.66pH21/2)3/4(5.69)

From Eqs. 5.69 and 5.2, it can be obtained that

[Mi•••]=(133/4 K5.641/8 Kw3/8)·pH23/8pH2O3/8·(1+K5.66pH21/2)3/4(5.70)

In the case when the oxygen partial pressure is relatively low, we may consider that the hydrogen partial pressure is relatively high leading to the following assumption [32]:

K5.66pH21/2=[Hi′][e′]≫1(5.71)

which also means

[Hi′]≫[e′](5.72)

From the assumption in Eq. 5.71, Eqs. 5.69 and 5.70 are reduced to be

[Mi•••]=(K5.663/433/4 K5.641/8)·pH23/8pO23/16=(K5.663/433/4 K5.641/8 Kw3/8)·pH23/4pH2O3/8(5.73)

This relation shows that, in the low oxygen partial pressure range, the defect concentration and therefore the rate constant increase with the increased hydrogen partial pressure at constant oxygen partial pressure. The prediction according to Eq. 5.73 is in agreement with the experimental result in Fig. 5.10(c), while the one according to Eq. 5.63 is not. However, this relation fails to explain the experimental results in Fig. 5.10(b) which shows that the rate constant increases with water vapour at the fixed hydrogen partial pressure.

From the analysis in this section, it was summarised [27, 32] that no single defect model can be used to explain the oxidation mechanism in the atmosphere containing water vapour. However, the linear combination of the mass-transfer contributions by metal vacancy and metal interstitials according to Eqs. 5.56 and 5.69 together with the equilibrium of the hydrogen dissolution reaction i.e. Reaction 5.74, contains all relations between the rate constant and the gas partial pressure at the low oxygen partial pressure regime [27, 32].

H2+OO↔OHO•+Hi′(5.74)

5.3.4 Effect of Hydroxyl Ion on the Oxygen Transport in the Oxide

Sections 5.3.2 and 5.3.3 consider the possibility that hydrogen interstitial is formed and might be bounded with the oxygen in the oxide giving the substitutional hydroxyl ion, OHO•. Thereafter, this defect affects the defect structure of the oxide, such as increasing the metal vacancy concentration in
the condition that water vapour content is high, thus contributing to the increased oxidation rate. However, the hydroxyl ion could also directly affect the diffusion of oxygen in the oxide as suggested by Galerie [8, 37].

Again in the classic 2001 paper, Galerie et al. [8] suggested that hydrogen from water vapour could be incorporated into the oxide by jumping into the substitutional site usually occupied by O2− or into the oxide grain boundary. The species stemed from water vapour that can jump into the substitutional site are the adsorbed OH(OH−s) or H+which might form the substitional hydroxyl by the following reactions [8]:

OH−s+VO••+e′→OHO•H++OO→OHO•

(5.75)(5.76)

For the lattice transport, the oxygen vacancy is needed for the diffusion of OHO• [8]. Wouters et al. [38] discussed that the ionic radius of the hydroxyl ion ( OH−) is about 90 pm while that of oxygen ion ( O2− ) is about 140 pm , and the charge of OH−is also less than that of O2−. It is then expected that the diffusion coefficient of the hydroxyl ion through vacancies is higher than that of the oxygen ion [38]. For the transport along the grain boundary, the charge of the diffusing species must go along with the semiconducting property of the oxide, i.e. the p-type metal-deficient oxide allows the anion transport and the n-type oxygen-deficient oxide allows the cation transport [8].

For thermally grown chromia on chromium and stainless steels, the photoelectrochemistry measurement showed that the semiconducting type of the oxide turned from amphoteric p+n when oxidised in oxygen into n when oxidised in water vapour [37, 39]. This effect is less for stainless steel since the iron doping into the oxide might weaken its p-type property [37, 39]. From these results, we may consider that, for the oxidation of chromium or the chromia-forming stainless steel in water vapour, the hydroxyl ion might be incorporated into the oxide in the substitutional site and diffuses via the oxygen vacancy, or it might be incorporated into the oxide grain boundary and further diffuses to the metal/scale interface. Bamba et al. [37, 40] oxidised Fe−15Cr−0.5Si in Ar−10%H2O at 850∘C and found the presence of OH in the oxide. Furthermore, Chandra-ambhorn et al. [5, 37] oxidised AISI 441 and AISI 444 stainless steels in O2 and H2−2%H2O at 800∘C and took these samples to assess the scale adhesion using the tensile test as explained in Section 2.3 in Chapter 2. They found that the oxide formed in H2−2%H2O is better adhered to the steel substrate [5] as shown in Fig. 6.4 in Chapter 6. The similar trend was also observed for the thermal scale on Fe-15.7Cr-8.5Mn formed in humidified oxygen compared with the one formed in oxygen [33]. The better adhesion indicates the possibility that the hydroxyl ion diffuses inwardly to form oxide at the metal/scale interface and therefore promotes the scale adhesion.

5.4 Characteristics of the Scale Formed in Humidified Atmosphere

Section 5.3 reviews the defect structure in M2O3 formed in water vapour. It represents the simple case of chromia which is a major oxide thermally grown on stainless steel. However, in the real case of stainless steel, the situation is more complex since the other oxides can also be formed and water vapour can also affect the other properties of the metal-scale system. These aspects will be reviewed in the present section.

Issartel et al. [41] reported the oxidation mechanism of 4509 ferritic steel ( Fe−18Cr−Nb−Ti ) at 1000∘C in air-7.5% H2O. They presented that, under water-containing oxidising condition, chromia and Cr−Mn spinel with the composition of Mn1.5Cr1.5O4 were formed while chromium species volatilisation occurs simultaneously. The oxidant, H+and/or H2O, inwardly transported through oxide scale and reacted not only with chromium to form chromia but also with iron at the chromium-depleted zone to form wüstite (FeO). Chromia internally formed owing to chromium oxidation meanwhile it was consumed by reacting with FeO to form FeCr2O4. FeO could possibly react with H2O which diffuses internally through scale and forms Fe3O4. At the same time, iron
outwardly diffused and externally reacted with H2O and Fe2O3 is then formed as a nodule which was a sign of catastrophic or breakaway oxidation.

Many works reported that the catastrophic oxidation of stainless steels occurs by the formation of Fe-rich nodules (haematite nodules) on the top surface [41-43]. Breakaway is highly dependent on the partial pressure of water vapour [44]. In the presence of water vapour, the breakaway can be reached faster than that in dry condition that means water vapour accelerates the onset of breakaway oxidation kinetics [45, 46]. There are many researches describing how Fe-rich nodule grows. As the increasing of scaling growth rate, the thicker scale is formed on the alloy exposed to wet condition rather than to dry condition. So, the cracks and spallation occur faster than that exposed to dry condition. Since the cracks occur, Fe2O3 could be thermodynamically reproduced and this oxide is indeed observed to grow [39]. Ehlers et al. [47] pointed out that the transformation of protective chromia into non-protective iron-rich oxide is governed by the ratio of H2O/O2 rather than the absolute level of H2O. The competitive adsorptions of H2O and O2 at the metal surface governs whether breakaway oxidation occurs [47]. Therefore, the occurrence of breakaway oxidation is determined by H2O/O2 ratio and it was found experimentally that a critical condition for breakaway is pH2O/pO2>1 [47].

The breakaway oxidation can also be enhanced by other effects like the chromium species volatilisation and acidity of the surface. Segerdahl et al. [48] oxidised 11% Cr steel (CrMoV11 1) in dry and wet O2 with the humidity of 10% or 40% at the temperature range of 450−700∘C. In wet atmosphere, either protective oxide or non-protective oxide was formed on the steel while chromium species volatilisation occurred in both cases. The oxide formed on the steels remains protective whenever the chromium loss by volatilisation and chromium supply by solid state diffusion to the oxide are still balance. If the rate of chromium loss is higher than the chromiumsupplying rate, the oxide transforms from protective internal chromium-rich ( Cr,Fe)2O3 and external iron-rich (Cr,Fe)2O3 to non-protective internal (Fe,Cr)3O4 and α−Fe2O3 due to the depletion of chromium resulting in the feasibility of breakaway oxidation. In addition, the presence of water vapour is more corrosive to alloy comparing with dry atmosphere by promoting higher chromium species volatilisation rate, and then the breakaway oxidation is reached. As for the acidity of the oxide surface, due to more acidity of Fe2O3 than Cr2O3, when Fe2O3 grows on the existing Cr2O3, the difference of surface acidities of these oxides arises, so the faster catastrophic oxidation occurs [43]. In order to prolong the lifetime of alloys, sufficient chromium must be presented in the alloy for selectively oxidising and forming a protective chromia scale. If alloyed chromium is not enough, iron-rich nodule, rather than the desired chromia, co-exists on the alloys [49].

Water vapour cannot only lead to the catastrophic oxidation, it can also change the properties of the oxide-metal system. Chromium exposed to water-containing atmosphere tends to internally oxidise more than that exposed to dry condition due to hydrogen dissolved in alloys [46, 50,51]. This may affect one or more following factors: oxygen solubility, oxygen diffusivity in alloys, alloy interdiffusion coefficient and parabolic rate constant [51]. Ani et al. [52] also described that, even though the dissolved hydrogen does not influence the diffusion coefficient of chromium, it increases the oxygen permeability in α−Fe in wet condition by a factor of 1.4 comparing with that in dry condition. This results in the increase in minimum concentration of chromium to form an external scale. Additionally, the dissolved hydrogen changes the oxide shape from discrete spherical particle in dry atmosphere to spike-like precipitates leading to the enhancement of oxygen transport along the metal/oxide precipitate interface [52]. Although many studies reported that water vapour extensively affects the scaling rate of oxide, it is rather different from the recent research by Cheng et al. [53] indicating that water vapour accelerated the formation of oxide scale but it was not significant. The difference in water vapour percentage does not affect the composition of the scale. However, water vapour causes the oxide grain refinement comparing with the oxide formed in dry condition which is consistent with the works of Othman et al. [49] and Jacob et al. [54].

For alumina-forming alloys, many researches confirmed that water vapour had little effect on their oxidation kinetics whenever the α-alumina was extensively formed on the alloy surface
[55-58]. However, according to the study of Janakiraman et al. [59] demonstrating the cyclic oxidation of PWA 1480, PWA 1484, CMSX 4, diffusion aluminide coatings on PWA 1480 and PWA 1484, and Co−24Cr−10.5Al−0.3Y in wet and dry atmospheres at 1100∘C, it was found that, in cases that alumina scale was cracked and spalled in dry air, the presence of water vapour caused the increasing in degradation rate by a factor of 2 . On the contrary, if there was no crack or spallation of alumina in dry air, water vapour had no effect on the oxidation behaviour. Furthermore, it was proposed that water vapour caused stress corrosion cracking at alumina/alloy interface during cyclic oxidation.

Nevertheless, there is a research proved that water vapour affects the early stage oxidation of alumina-forming alloy. Götlind et al. [60] isothermally oxidised FeCrAl alloy Kanthal AF at 900∘C in dry O2 and O2−40%H2O. They found that the oxidation rate was significantly faster in wet atmosphere than in dry one during the first 72 h . The development of alumina in dry and wet O2 at 900∘C was proposed to be as follows: stage (1), γ−Al2O3 and the mixture of oxides corresponding to the composition of the alloy nucleate on the surface. After prolonged exposure, a chromium-rich band within the alumina scale is observed. At stage (2), γ−Al2O3 outwardly grows with the high rate leading to rapid mass gain. Also, γ−Al2O3 provides outward mass transport paths for divalent cations. Meanwhile, α−Al2O3 nucleates and grows slowly inwards and laterally. Pores form in γ−Al2O3 layer close to the inner oxide interface. After 1 h of exposure, the size and distribution of the pores are essentially the same in dry and wet O2. Then, stage (3), a continuous layer of inwardgrowing α−Al2O3 forms. This layer restricts the supply of cations for the growth of outer γ−Al2O3 layer resulting in a much slower oxidation rate. After that, the mechanisms in dry and wet oxidations are different. In dry O2, the outer oxide changes into a mixture oxide phases resulting in the cease of outward oxide growth in which the α−Al2O3 is formed as the outer layer. The γ−Al2O3 continues its slow outward growth leading to the formation of an uneven oxide while the porosity decreases. The slower oxidation in dry O2 is due to an increase in the effective thickness of the α−Al2O3 layer by transformation of γ−Al2O3 and by pore filling. On the other hand, in wet O2, γ−Al2O3 does not transform to α−Al2O3 due to the stabilisation of γ−Al2O3 by H2O forming a hydroxylate surface. The pores have no sign to disappear and the scale formed in wet condition is thicker than that formed in dry condition. This makes an alloy exposed to wet atmosphere more susceptible to spallation than that exposed to dry atmosphere.

5.5 Summary

This chapter reviews the role of water vapour on the high temperature corrosion of stainless steels, starting from the scenario when water vapour is in the bulk gas to the situation when it is at the solid/gas interface and in the oxide.

When water vapour is in the bulk gas, it can be dissociated giving oxygen and hydrogen. The oxygen uptake for oxidation of metals is not mainly from the oxygen dissociated from water vapour but from the water molecule itself. The dissociation of water can change the ratio of the partial pressure of water vapour to the partial pressure of hydrogen ( pH2O/pH2 ), which could further affect the stability of the oxide formed. The variation of this ratio when the oxidation kinetics are linear and parabolic is derived.

Once water vapour adsorbs on the surface, it can oxidise the metal or volatilise the formed oxide. In the case that water vapour cannot significantly enter the oxide and change its transport properties, the surface reaction especially the dissociation of the adsorbed OH to the adsorbed O can be a rate determining step. In that case, the more acidic oxide, as indicated by the less negative value of the cation hydration enthalpy, tends to have the faster linear oxidation rate. The surface reactioncontrolled kinetics cannot be observed in the chromium oxidation in water vapour but can be found when stainless steel is oxidised in carbon-containing atmosphere.

Water vapour can also volatilise the chromia to the volatile species. The volatilisation rate depends on parameters such as the linear velocity of the gas in the atmosphere. At relatively slow
linear velocity, the volatilisation rate increases with the velocity, indicating that the rate is controlled by the mass transfer of the volatile species from the solid/gas interface to the bulk gas. When the velocity is higher, the volatilisation rate can be constant which indicates that the rate is controlled by the surface reaction. Coating or surface modification e.g. the pre-oxidation technique can be used to reduce the rate of chromium species volatilisation.

Many works reported that water vapour can be incorporated into the oxide primarily in the form of hydrogen interstitial with positively unit effective charge ( Hi• ). Based on Norby's work [35], the construction of the Brouwer diagram of M2O3 formed in humidified oxygen is derived, and the graphical representation of the role of water vapour on metal and oxygen vacancies concentrations is proposed. In the regime of relatively high water vapour, increasing water vapour increases the metal vacancy concentration which could contribute to the increased rate of the formation of the metal-deficient oxide. Hi• can also be bounded with the oxygen in the oxide resulting in the substitutional hydroxyl defect (OHO•). The hydrogen dissociation may also give OHO• and Hi′. The linear combination of the mass transfer due to metal vacancy and metal interstitial together with the hydrogen dissociation giving of the hydrogen defects ( OHO• and Hi′ ) were proposed by Hänsel et al. [32] to describe the defect situation in the chromia formed in water vapour. The hydroxyl ion could also help facilitate the inward oxygen transport due to its smaller ionic radius and charge compared with those of O2−. For the chromia formed in water vapour, the oxide is found to be n -type and in this case the hydroxyl ion may transport via the oxygen vacancy or along the oxide grain boundary to the metal/oxide interface. The improved adhesion of scale on stainless steel formed in the atmosphere containing water vapour supports the possibility of this hydroxyl ion inward diffusion.

Apart from chromia, other oxides can be formed during the oxidation of stainless steel in water vapour and in many cases can lead to the catastrophic oxidation. This breakaway oxidation relates to many effects such as the pH2O/pH2. The chromium species volatilisaion or high acidity of the oxide can also provoke the breakaway oxidation. Water vapour can also affect the other properties of stainless steel or the oxide scale. It can increase the oxygen permeability in the ferrite matrix or refine the oxide grain boundaries. For the alumina-forming alloys, many works reported that water vapour has little effect on the oxidation kinetics. However, water vapour can affect the oxide formation on the alumina-forming stainless steel. It can stabilise the γ-alumina and form a thick and porous oxide leading to susceptibility to the oxide spallation.

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Solid State Phenomena, ISSN: 1662-9779, Vol. 300, pp 135-156

doi: 10.4028/www.scientific.net/SSP.300.135

© 2020 The Author(s). Published by Trans Tech Publications Ltd, Switzerland.

Submitted:
2019-05-27

Revised:
2019-06-28

Accepted:
2019-08-08

Online:
2020-02-12

Sébastien Chevalier 1,a*, Lionel Combemale 1, b, Ioana Popa 1,c, Somrerk Chandra-ambhorn 2, d, Walairat Chandra-ambhorn 3,e, Piyorose Promdirek 2,f and Patthranit Wongpromrat 3, g

The original version of this paper is available on https://www.scientific.net/SSP.300.135.pdf

1 Laboratory ICB UMR 6303 CNRS, Université de Bourgogne Franche-Comté, Dijon, France

2 High Temperature Corrosion Research Centre, Department of Materials and Production Technology Engineering, Faculty of Engineering, King Mongkut's University of Technology North Bangkok, 1518, Pracharat 1 Road, Bangsue, Bangkok, 10800, Thailand

3 Department of Chemical Engineering, Faculty of Engineering, King Mongkut's Institute of Technology Ladkrabang, Chalongkrung 1 Road, Lat Krabang, Bangkok, 10520, Thailand

a  sebastien.chevalier@u-bourgogne.fr, blionel.combemale@u-bourgogne.fr, c  ioana.popa@u-bourgogne.fr, d  somrerk.c@eng.kmutnb.ac.th, e  walairat.ch@kmitl.ac.th, f piyorose.p@eng.kmutnb.ac.th, 9 patthranit.wo@kmitl.ac.th

Keywords: solid oxide fuel cells, interconnect, stainless steel

Abstract

The chapter introduces components and working principle of solid oxide fuel cells (SOFCs). It is followed by the explanation on the choices of materials focussing on ferritic stainless steels. The review is further made on the required properties of these steels, i.e. low oxidation rate, low chromium species volatilisation rate, high electrical conductivity and good scale adhesion. For the oxidation aspect, the behaviour of stainless steel interconnect in cathode, anode (hydrogen and biogas), and dual atmospheres are described. Surface modification by pre-oxidation and coatings to improve the oxide electrical conductivity and to reduce chromium species volatilisation is finally reviewed.

CHAPTER 6

The original version of this paper is available on https://www.scientific.net/SSP.300.135.pdf

Development of SOFC Interconnect Stainless Steels

The original version of this paper is available on https://www.scientific.net/SSP.300.135.pdf

6.1 Introduction

As described by Kofstad and Bredesen, solid oxide fuel cells (SOFCs) are electrochemical reactors for direct conversion of the chemical energy of gaseous fuel and oxygen to electrical energy at high temperatures [1]. Among the various types of fuel cells, SOFCs operate at high temperature, typically between 600 and 1000∘C, and have advantages in terms of high conversion efficiency [2,3].

Fig. 6.1 schematises the SOFC components and the working principle [4]. A single SOFC cell is generally composed by a porous anode (e.g. YSZ/Ni), a porous cathode (typically a perovskite mixed conductor, e.g. LaxSr1−xMnO3,LaxSr1−xCoyFe1−yO3, etc.) separated by a gas imperviousness ion-conducting electrolyte, such as yttria-stabilised zirconia (YSZ). For the operating principle of SOFCs, oxygen in air on the cathode side received electrons supplied from the anode through the external circuit thus giving oxygen ions ( O2− ). The oxygen ions transfer from the cathode side through the electrolyte to the anode side. At the anode, hydrogen reacts with oxygen ions giving water and electrons. The flow of electrons through the external circuit gives then electricity.

For the atmospheres on the cathode side, Antoni [4] reported that when air is flowed into this side the partial pressure of oxygen at the inlet is 0.2 bar ( 10−0.7 atm ). The oxygen partial pressure can be reduced to about 10−4 atm at the gas outlet. It is noted that water vapour can be presented in air and can further affect behaviour of stainless steel interconnect at high temperatures [5, 6]. At equilibrium between liquid and gaseous water vapour, the water vapour content in the atmosphere is in the range of 2.7−3.6% at 20−25∘C calculated using thermodynamic data compiled by Barin [7]. For the anode side, the oxygen partial pressure is considerably low. It can be about 10−18 atm at the
gas inlet and increases to be about 10−8 atm at the gas outlet. It is noted that in the anode side, the fossil fuel might be used instead of hydrogen since at high temperatures the fossil fuel can be reformed giving hydrogen gas [2,8]. However, the use of fossil fuel can dramatically change the atmosphere in the anode side compared with the typical one, which mainly contains hydrogen. For example, in the case of methane, gas composition at the inlet is methane but gradually changes to a gas mixture composed of H2+H2O and CO+CO2 [1]. At the gas outlet, the partial pressure ratio of CO2 to CO can be 9:1 and the oxygen partial pressure at 900∘C is 6.7×10−15 atm [1]. Furthermore, it was reported that the water vapour content at this location can be as high as up to more than 0.5 atm [1].

[picture]

Fig. 6.1. SOFC components and working principle. Reproduced from L. Antoni, Mater. Sci. Forum 461-464 (2004) 1073-1090 [4].

Fig. 6. 1. SOFC components and working principle. Reproduced from L. Antoni, Mater. Sci. Forum 461-464 (2004) 1073-1090 [4].

What described above is about a single cell. However, to enhance the voltage output, multiple cells are "stacked" in series. Therefore, an "interconnect" for providing electrical connection between an anode of one individual cell and a cathode of the neighbouring one is required. It also acts as a barrier to avoid any contact between anode atmosphere (hydrogen or fossil fuel) and cathode atmosphere (oxygen or air), while at the same time it helps maintain the structural integrity of the stack.

In order to perform their intended functions, interconnects should have the following characteristics [9-12].
(1) Thermal expansion coefficient (TEC) should match with those of ceramic cell components and seal materials (around 10.5×10−6 K−1 ), so that the thermal stresses developed during starting up and shutting down could be minimised.
(2) Oxidation and corrosion resistances should be good enough during simultaneous exposure to both cathode and anode atmospheres for 40000 h .
(3) Volatility of constituents that can degrade cell performance should be minimal.
(4) Electrical conductivity, in particular through the oxide scale thermally grown on the metallic interconnect surface, should be high. The acceptable area-specific resistance (ASR) level is considered to be below 100 mΩ cm2 [13].
(5) Scale adhesion should be good so that the separation of scale from the steel substrate, which could lead to electrical resistance, does not occur.
(6) Chemical compatibility with other cell components in contact with interconnects such as seals and electrodes should be good.

In addition, the materials must be cost effective in terms of raw materials and manufacturability. Therefore, the development of suitable interconnect materials is one of challenges in improving the performance and cost-effectiveness of SOFCs.

6.2 Selection of Materials for the Use as an Interconnect

When operating at 900−1000∘C, the leading candidate material for the interconnect was doped lanthanum chromite, LaxSr1−xCrO3 [4, 14]. However, these materials are quite expensive; moreover, it is difficult to obtain high density at reasonable sintering temperatures and the tendency of the chromite interconnect to be partially reduced at the fuel gas/interconnect interface, causing the component to warp and the peripheral seal to break [2, 15]. The recent trend in developing lower temperature device (650-800 ∘C ), with micrometer thin electrolytes [16], new electrolytes with improved conductivity [17] and/or improvement of the cathode/electrolyte interface reaction (i.e. triple phases boundaries (TPB) to internal diffusion (ID) mechanism) makes it feasible for lanthanum chromites to be supplanted by metals or alloys as the interconnect materials. Indeed, they offer a number of advantages over ceramic materials. They are easier, and therefore cheaper to fabricate, they are less brittle and, in general, they have higher electrical and thermal conductivities compared to the ceramic interconnects [9, 10, 12, 18, 19].

To be reliable, the alloy must develop a thermally grown oxide scale on its surface, which generally has high electrical resistivity. It was reported that electrical conductivity of chromia is about 10−1 S cm−1 at 750−800∘C while that of alumina is about 10−10 S cm−1 at 700∘C [20]. The relatively high electrical conductivity of chromia can be considered as a "compromise" and lead to the choice of chromia-forming alloys. Among the different groups of chromia-forming alloys (nickel-based, chromium-based and iron-based alloys), the bcc base alloys (chromium-based alloys and ferritic stainless steels) have been the most widely investigated category of alloys for interconnect applications because they have a thermal expansion coefficient ( 12×10−6 K−1 ), which matches well with that of ceramic cells (11×10−6 K−1)[18,19,21].

Chromium-based alloys, like Plansee Ducralloy, Cr5Fe1Y2O3, were the early favourites as a replacement for doped lanthanum chromites because they demonstrate excellent oxidation resistance, good electrical properties and high temperature mechanical strength at high temperatures ( >900∘C ) SOFCs [22, 23]. However, these alloys are very expensive, particularly in comparison with Fe−Cr alloys. In contrast, Fe−Cr ferritic stainless steels (FSS), designated as 400 series in the AISI system, not only demonstrate good oxidation resistance and the ability to match the thermal expansion coefficients of cell components, but also are among the most cost effective alloys [18, 21]. Ferritic stainless steels (e.g. AISI 430, AISI 444, Crofer 22 APU, ZMG 232) were proposed and tested as interconnect materials in stacks in the late 1990s [24].

Traditionally, oxidation resistant alloys have been designed with an emphasis on surface and structural stability, but not electrical conductivity of the scale, which is equally important for SOFC interconnect applications. Thus, alloying practices used in the past to enhance surface and structural stability may not be compatible with high electrical conductivity. For example, silicon that is often a residual element in alloy substrate, may help improve alloy oxidation resistance through formation of silica subscale along the metal/scale interface. This subscale is electrically insulating and therefore increases electrical resistance. It can also cause scale spallation due to its low thermal expansion coefficient, as indicated by studies on ferritic stainless steels such as AISI 446 [25]. Such chromia-forming stainless steels form, at high temperatures, a Cr2O3-based thermally grown oxide (TGO). Unfortunately, the formation of volatile gaseous Cr(VI) species such as CrO3 or CrO2(OH)2 can cause serious degradation of the fuel cell performances by poisoning the cathode and can cause the lost of the protectiveness of chromia [5,6,26,27] especially in presence of water vapour. The formation of a protective, single-phase chromia layer requires chromium content of approximately 17−20%, depending on temperature, on surface treatment and on minor alloying additions [18, 21]. With respect to oxidation behaviour, the most important minor alloying additions in commercial
ferritic steels are manganese, titanium, silicon and aluminium; their concentration commonly being a few tenths of a percent. Manganese was added to obtain external spinel formation, which is expected to decrease the formation of volatile chromium-species and to enhance the electrical conductivity of the scale [28-31]. Titanium was also added to obtain fine internal oxide precipitates of titania leading to strengthening of the near-surface region, thus reducing the tendency of surface wrinkling caused by relaxation of oxidation-related stresses during thermal cycling [24]. The amount of aluminium and silicon must be kept low to prevent the formation of their insulating oxides, alumina and silica. Furthermore, SiO2 can also cause scale spallation due to its low thermal expansion coefficient, as indicated by studies on ferritic stainless steels [32, 33]. Elements such as tungsten or molybdenum can also be added to match better thermal expansion coefficient of the alloys to those of other fuel cell components. Among newly developed alloys, the one that has received the most attention is Crofer 22 APU, a ferritic stainless steel developed by Quadakkers et al. at Jülich and commercialised by thyssenkrupp of Germany. Crofer 22 APU, which contains ~0.5% manganese, forms a scale composed of a (Mn,Cr)3O4 spinel top layer and a chromia sublayer. Besides the "clean" form of ferritic stainless steels that require extra process such as vacuum melting, recent practice has focussed on fabrication of ferritic stainless steels by alloying approach to avoid the negative effects of silica sublayer formation. A small amount of niobium is added to the ferritic substrate and this element ties up residual silicon, mitigating the risk of formation of the insulate silica layer at the metal/scale interface. The niobium-containing ferritic stainless steel, as AISI 441, demonstrates an electrical performance comparable to Crofer 22 APU [34, 35].

Table 6.1 exemplifies commercial ferritic stainless steel grades for interconnect application [5, 36-39]. It can be seen from the table that the basic grade is AISI 430, which mainly consists of iron with 16.5wt.% chromium. The higher grades in this series are AISI 441 and 444, which contains higher chromium content of about 18wt.%. In general, the chromium contents are in the ranges of 17.5-18.5 wt.% and 17-20 wt.% for AISI 441 and 444 respectively [40]. The AISI 444 is also stabilised by titanium but additionally alloyed with molybdenum by about 2wt.%. It was reported that the molybdenum addition could help increase the pitting corrosion resistance [41], the stress corrosion resistance [42], the cyclic oxidation and high temperature strength of the steel [43-45]. However, as will be shown in Section 6.6, the molybdenum addition could form Fe-Mo intermetallic compound, which deteriorates the scale adhesion at high temperatures [36].

Table 6.1 Chemical compositions (wt.%) of ferritic stainless steels for SOFC interconnect application [5, 36-39].

Grade | Cr | Other elements | Ref.
430 | 16.5 | 0.4Mn–0.35Si | [5]
441 | 17.8 | 0.24Mn–0.6Si–0.006Al–0.13Ti–0.55Nb | [36]
444 | 17.6 | 0.42Mn–0.42Si–2.04Mo–0.16Ti–0.28Nb | [36]
Crofer 22 A | 22 | 0.5Mn–0.1Si–0.1Al–0.1Ti–0.1La | [37]
Crofer 22 APU | 22 | 0.5Mn–0.02Si–0.02Al–0.1Ti–0.1La | [37]
Crofer 22 H | 22 | 0.5Mn–0.25Si–0.02Al–0.1Ti–0.1La–2W–0.5Nb | [37]
ZMG 232 | 22 | 0.5Mn–0.4Si–0.21Al–0.22Zr–0.04La–0.26Ni | [37]
ZMG 232L | 21.8 | 0.46Mn–0.08Si–0.05Al–0.19Zr–0.05La–0.34Ni | [38]
ZMG 232 G10 | 23.6 | 0.27Mn–0.05Si–0.08Al–0.32Zr–0.06La–0.51Ni–1.99W–0.96Cu–0.02N | [39]
Sanergy HT | 22.3 | 0.24Mn–0.02Si–0.03Al–0.77Mo–0.03Ti–0.6Nb–0.1Zr | [37]
ITM | 25.7 | 0.04Mn–0.03Si–0.04Al–1.9Mo–0.21Ti–0.24Y | [37]
E-bright | 26 | 0.05Mn–0.2Si–1Mo–0.1Nb–0.15Ni | [37]

The next ferritic grades with higher chromium content are the grades with about 22wt% chromium. The major grades specially developed for this application are Crofer and ZMG. As previously mentioned, the Crofer grade has been developed by Jülich Research Centre and commercialised by thyssenkrupp. Niewolak et al. reported that Crofer 22 A stainless steel mainly consists of iron with 22wt.% of chromium, also with 0.1wt.% of lanthanum as a reactive element and 0.1wt.% of titanium. Silicon and aluminium contents in this grade are equally 0.1wt.%. The Crofer 22 APU still contains the same contents of chromium, lanthanum and titanium. However, the silicon and aluminium contents of this grade are reduced to be 0.02wt.%. Furthermore, in order to increase the creep strength of this stainless steel, tungsten with the content of 2wt.% is added to Crofer 22 APU with slight modification of the chemical compositions of the other elements, giving rise to the Crofer 22 H grade [37]. The ZMG grade has been developed by Hitachi Metals. This grade also contains chromium content of about 22wt.% but it is added by the reactive elements i.e. zirconium and lanthanum. It is noted that the silicon and aluminium contents of this steel are relatively high i.e. 0.4 and 0.21wt%, respectively [37]. The contents of these two elements were reduced to be 0.08 and 0.05wt% for the improved ZMG 232 L [38]. Recently, new ZMG 232 G10 has been developed still with the reduced silicon and aluminium contents but additionally with about 2wt% of tungsten, 1wt% of copper and 0.02wt% of nitrogen [39].

The ferritic grades with higher chromium content of 25−26wt.% are such as ITM and E-bright as shown in Table 6.1. Similar to the steel with lower chromium contents, these grades tend to be produced by lowering the silicon and aluminium contents and alloying molybdenum, titanium, niobium or the reactive elements such as yttrium. However, the steels with chromium content higher than 25wt.% are prone to precipitate the sigma phase even at the high temperature of 800∘C if some other elements, e.g. molybdenum and silicon, are added to the steel [46]. Singheiser et al. [46] then recommended that the chemical composition of the high-chromium steel must be carefully adjusted to the actual environment to avoid the sigma phase formation, which could cause the embrittlement of the steel.

6.3 Oxidation Behaviour

6.3.1 Oxidation on the Cathode Side

During high temperature exposure in air, chromium contained in the alloy substrate is preferentially oxidised, forming an oxide scale on the alloy surface to protect it from further environment attack. Some studies revealed that, in most cases, outward chromium transport is the dominant growth process in the chromia scale, whereas the contribution of oxygen transport is substantially smaller [47, 48]. However, stainless steel does not contain only chromium but also manganese. The latter element can play a role on oxidation kinetics. Park et al. [49] have recently studied the air oxidation mechanism of Crofer 22 APU and found that the oxidation mechanism composes of two steps. In the first step, Cr−Mn spinel and chromia are grown with the faster rate for the spinel formation because of the rapid outward diffusion of manganese [49]. Later, when oxygen penetrates through the spinel layer and reaches the chromia, the growth rate of chromia is changed abruptly and becomes faster than the growth rate of spinel [49]. This mechanism finally leads to the formation of chromia as a major scale with the relatively thin Cr−Mn spinel formed on the top of chromia.

Furthermore, the outward chromium diffusion results in the formation of voids and cavities at the metal/scale interface leading to poor scale adherence. Thus, after exposure over thousand hours, the chromia layer can exhibit localised spallation and thus creates an important electrical power loss at the electrode/interconnect interface [21,50,51]. It is well known that the addition of small amounts of reactive elements (lanthanum, zirconium, yttrium, neodymium, etc.), to Cr2O3-forming alloys greatly enhances their resistance to high temperature oxidation [47, 48, 52]. This effect, known as reactive element effect, is reviewed in Chapter 4. However, it may be briefly described here that their main role consists of changing the diffusion mechanism during the chromia scale formation
from outward diffusion of the chromium ions to inward diffusion of the oxygen ions. The change of the diffusion mechanism allows for the decrease of the corrosion rate and can significantly increase the oxide scale adherence. The reactive element effect can be produced by a wide variety of reactive element addition methods, such as microalloying [21], dispersion of oxides [25] and surface coatings with a thin layer of the reactive element oxide [23,53]. The enhancement of the electrical properties by addition of reactive element was also observed [23,53].

6.3.2 Oxidation on the Anode Side 6.3.2.1 Oxidation in H2−H2O

In comparison with the oxidising, cathode side environment, the reducing, anode side environment is more complex, particularly when a hydrocarbon fuel is used. However, the understanding of the oxidation behaviour in anode environment ( H2−H2O ) is very important and it is still lacking. The literature data dealing with the effect of water vapour on chromia-forming alloys show that water vapour plays a huge role on the corrosion phenomenon. The water vapour seems to be either beneficial or detrimental. Concerning the corrosion rate, some papers are contradictory [54]. Brylewski et al. [55] reported similar values of kp for oxidation of AISI 430 in air and in 94%H2/6%H2O but the electrical resistance was higher in the anode gas than in the cathode gas. On one hand, Quadakkers et al. [56] reported that the chromia scales formed on ferritic stainless steels in simulated anode gas at 800∘C were thinner than those formed in air. On the other hand, although the oxygen partial pressure was lower on the anode side, Quadakkers et al. [56] also found that the oxidation rate of chromium-based ODS alloy at 950−1050∘C in a H2/H2O mixture (equilibrium oxygen partial pressure approximately 10−15bar ) was higher than in the environments with high oxygen partial pressure (air, Ar/O2). England and Virkar [57] oxidised nickel-based alloy (Haynes 230) in H2−H2O and in air between 700 and 1100∘C. The growth rate of the chromia scale at 800∘C in H2−H2O was higher than that observed in air; moreover, the electrical resistance of the scale was higher in H2−H2O than in air. Liu [58] also found that, in a reducing atmosphere of Ar5%H2−3%H2O, the electrical property of the oxide scale was worse. Finally, Fontana et al. [53] found that in a reducing atmosphere H2−10%H2O, the corrosion rates of Crofer 22 APU and Fe− 30Cr were higher than those in air atmosphere. Due to the technical difficulty to measure electrical resistance in H2−H2O, the electrical property of the thermally grown oxide in H2−H2O is not well understood. Guillou et al. [59] evidenced in solid oxide electrolyser (SOE) conditions and by using marker experiments with D2O that it was the hydrogen coming from water vapour that was the main cause of the drastic decrease of the electrical conductivity of chromia scale. One beneficial effect of water is that it can improve the adherence of the scale [36, 60], especially during thermal cycling, by reducing the amount of porosity at the alloy/scale interface. The severe buckling of the scale formed on chromia-forming alloy without reactive element observed by many authors during air oxidation does not occur during oxidation in H2−H2O. For example, oxide scale formed on chromia-forming alloy in H2−H2O showed good adhesion during cyclic oxidation at 950∘C, whereas scale formed in air in the same condition showed spallation [24]. Conversely, on the outer part of the scale, water vapour was shown to promote the formation of less protective microporous scale [24, 61].

6.3.2.2 Oxidation in Biogas

As mentioned in Section 6.1, one of the advantages of SOFC is that the fossil fuel might be used since it can be reformed at the SOFC working temperature thus providing hydrogen gas for the anode side. One of the promising types of renewable fuel is biogas. The biogas issued from fermentation typically contains about 50−70% methane, 30−50%CO2 with H2 S as impurity [62, 63]. To understand the high temperature corrosion mechanism of stainless steel under biogas, the corrosion mechanism of stainless steel under CO2 was first explored and then followed by CO2 mixed with methane which is in fact the biogas.

For the AISI 441 stainless steel oxidation in dry CO2, Promdirek et al. [64] found that at 700−850∘C the oxidation kinetics was linear, while at higher temperatures i.e. 925−1000∘C the oxidation was firstly linear followed by parabolic at longer times [64]. At the SOFC working temperature of 800∘C, they found that the relationship between the linear rate constant ( k1 ) and CO2 partial pressure ( pCO2 ) was in the following form:

k1=ApCO21+BpCO2(6.1)

where A and B are constant. Furthermore, they found from the photoelectrochemical experiment that the oxide formed in this atmosphere was n-type [64]. Based on these results, the oxidation mechanism in CO2 was proposed as follows.

First, CO2 adsorbs on the metal surface i.e.

CO2( g)+s↔CO2−s(6.2)

where s is the available site on the surface and CO2−s is the adsorbed CO2 [64].
Second, the adsorbed CO2 dissociates giving the adsorbed oxygen and CO according to the following reaction [64]:

CO2−s↔O−s+CO(g)(6.3)

If it is considered that VO•• is a dominant defect in this condition, this defect could be formed at the metal/scale interface by the following reaction:

2<Cr>↔2CrCr+3 VO••+6e′(6.4)

where < Cr> is the chromium in stainless steel and CrCr is chromium on the chromium site in chromia. The formed VO•• diffuses to the external interface. Such defect is filled by the adsorbed oxygen in Reaction 6.3 giving

VO••+O−s↔OO+s(6.5)

In the case that Reaction 6.3 is a rate determining step, the reaction rate ( r ) is as follows [64]:

r=k6.3θ(6.6)

where k6.3 is the rate constant of Reaction. 6.3 and θ is the coverage ratio of CO2 on the surface. If equilibrium of Reaction 6.2 is assumed, we can write the equilibrium constant as

K6.2=θpCO2(1−θ)(6.7)

From Eqs. 6.6 and 6.7, we can obtain that

r=k6.3 K6.2pCO21+K6.2pCO2(6.8)

This reaction rate is in the same form of the relationship between the linear rate constant and the CO2 partial pressure, as experimentally found in Eq. 6.1. This congruence supports the feasibility of the proposed mechanism [64].

Further, when methane is added to CO2 by 30−70% at 700−900∘C, Promdirek et al. [63] found that the corrosion kinetics was linear and the linear rate constant increased with the increased ratio of methane in the biogas. They proposed that the high temperature corrosion of AISI 441 stainless steel in biogas was due to two parallel mechanisms: the oxidation by CO2 and the carburisation by methane. They further reported that the carbon deposition and carbide formation were very rapid in pure methane. This phenomenon should contribute to the faster corrosion rate in biogas with the higher proportion of methane as observed from the experiment [63]. From the experiment, they further found that the scale was compact at 700 and 800∘C and could act as a barrier for carbon penetration, while at the higher temperature of 900∘C the scale was porous and carbide was formed near the metal/scale interface which could lead to the embrittlement of the steel. The SOFC working temperature of 800∘C seems then to be the maximum temperature for the use of AISI 441 stainless steel in this biogas atmosphere [63].

6.3.3 Oxidation in Dual Exposure Conditions

When simultaneously exposed to air at the cathode side and to H2−H2O at the anode side during SOFC service, interconnects experience a hydrogen partial pressure or chemical potential gradient from the fuel side to the air side. The oxidation behaviour in dual environment is very important and it is still lacking. Yang et al. [65, 66] observed that the oxidation behaviour of an alloy exposed to wet hydrogen on one side and air on the other side was different from that exposed to air on both side. The scales formed in air on AISI 430 exposed to a dual atmosphere contained iron-rich spinel or Fe2O3 nodules. In comparison, there was no haematite phase formation on the air/air sample, on the hydrogen side of the air/hydrogen sample, or on the hydrogen/hydrogen sample. This anomalous oxide (haematite instead of chromia) in air side was described in literature in dual atmosphere [67-74]. The potential detrimental effect of the dual exposure appeared to be dependent on the alloy composition, in particular chromium percentage in the Fe-Cr substrate. For Crofer 22 APU with 22wt.% chromium, no haematite phase formation was observed under the same test conditions as for AISI 430. The mobility of iron in chromia is higher than that of chromium [75], but the oxidation behaviour in a dual atmosphere suggests that hydrogen may further increase the mobility of iron. Measurements of the permeation of hydrogen through ferritic stainless steels [76] indicate that hydrogen can diffuse through the sample. However, the permeability of hydrogen through chromia is low, so once a protective chromia scale forms, the amount of hydrogen penetrating through oxidised sample is low. This anomalous behaviour in dual condition was found to be emphasised by electrical current in SOE conditions [77]. The mechanism of this phenomenon is not very clear and further work is required to gain further insight into the oxidation behaviour under dual environment and its effects on long term metallic interconnect stability.

6.4 Chromium Species Volatilisation

On the cathode side where the oxidising atmosphere is humidified, the chromia formed on stainless steel does not only oxidise but could also volatilise. At temperatures lower than 1500∘C, the dominant reaction of the volatilisation of chromia by oxygen and water vapor was reported to be the following one [5, 6, 78]:

12Cr2O3( s)+34O2( g)+H2O(g)↔CrO2(OH)2( g)(6.9)

In the case of stainless steel which contains not only chromium but also manganese, the addition of manganese could lead to the formation of Cr−Mn spinel on the top of chromia layer. In this case, the chromium volatilisation from the Cr−Mn spinel could be as follows [79, 80]:

(1+x)Cr2MnO4( s)+3xH2O(g)+5x2O2( g)→Cr2−xMn1+xO4( s)+3xCrO2(OH)2( g)(6.10)

It was reported that the CrO2(OH)2 volatile species could diffuse through the porous cathode to the interface between cathode and electrolyte and reduced there giving chromium oxide, possibly by the following reaction [81]:

2CrO2(OH)2( g)+3 VO••+6e′→Cr2O3( s)+3OO+2H2O(g)(6.11)

The precipitation of this oxide blocks the active sites for oxygen reduction at the cathode/electrolyte interface, thus degrading the cell performance. In the case that the cathode is LSM, the oxide (Cr,Mn)3O4 can also be formed at the interface between the cathode and electrolyte [81, 82]. This problem can be reduced by forming the oxide scale that exhibits low chromium volatilisation rate, which can be done by adding some alloying elements, by pre-oxidation before using the alloy in service, or by appropriate coating. The two first methods will be described in the following while the last method will be described in Section 6.7.

As for the effect of the alloying element, the important element for suppressing the chromium volatilisation is manganese. Galerie calculated the partial pressure of the chromium volatile species i.e. CrO2(OH)2 from chromia (Cr2O3) and Cr−Mn spinel (MnCr2O4) in N2−20%O2−15%H2O at 900−1500∘C [40]. Thermodynamic data for Cr2O3,MnCr2O4 came from Opila [78] and Jung [83], respectively. It was found that the partial pressure of CrO2(OH)2 from Cr2O3 is in the range of 1.6×10−7 to 3.7×10−6 bar, while the partial pressure of CrO2(OH)2 from MnCr2O4 is lower in the range of 5.2×10−9 to 3.5×10−7 bar. It indicates the possibility that the chromium volatilisation rate from the scale containing Cr−Mn spinel is lower than that containing only chromia.

For the chromium volatilisation from alloys, Niewolak et al. [37] compiled the results from literature [84, 85] and empirically found that the increase of manganese in various alloys up to about 0.5wt.% could lead to the reduced chromium species volatilisation rate. It is widely accepted that the formation of Cr−Mn spinel could suppress the chromium species volatilisation from stainless steel [5, 37]. However, the protectiveness from chromium species volatilisation due to Cr−Mn spinel is not only from the formation of continuous layer but also the chromium-to-manganese ratio in this oxide. In this matter, Falk-Windish et al. [80] reported that the ratio of chromium to manganese in Cr−Mn spinel i.e. (Cr,Mn)3O4 can vary from 2:1 to below 1:2. In the early period of oxidation, this ratio is relatively high, but it gradually decreases due to the rapid outward manganese diffusion and approaches a constant value. The decrease in this ratio implies the lower chromium activity and thus the lower chromium species volatilisation rate.

For the pre-oxidation technique, stainless steel might be oxidised in a reducing atmosphere to give the n-type chromia [6, 64]. In the case that the dominant defect at the external part of the chromia is a p-type metal vacancy, the hydroxyl ion cannot jump into the metal vacancy site [6]. On the contrary, if the dominant defect in chromia is an n-type oxygen-vacancy, the hydroxyl ion could be easily incorporated into the oxide, as shown in Fig. 6.2 [6]; in such a case that the chromium species volatilisation could be suppressed. Such possibility to reduce chromium species volatilisation is congruent with the experiment conducted by Wongpromrat et al. [6]. They preoxidised AISI 441 stainless steel in 50% CO- 50%CO2 atmosphere at 850∘C to let the formation of n-type oxide, and found that the pre-oxidation could help reduce the chromium specie volatilisation rate of the steel in O2−5%H2O at 800∘C [6]. For example, at the linear gas velocity of 4 cm s−1,
the chromium species volatilisation rate of the bare AISI 441 is about 2μ g m−2 s−1. It reduces to about 0.25μ g m−2 s−1 when the mentioned pre-oxidation was conducted [6].

[picture]

Fig. 6.2. Effects of metal vacancy (left) and oxygen vacancy (right) in the oxide on the chromium species volatilisation. Redrawn from Wongpromrat et al., Corros. Sci. 106 (2016) 172-178 [6].

Fig. 6. 2. Effects of metal vacancy (left) and oxygen vacancy (right) in the oxide on the chromium species volatilisation. Redrawn from Wongpromrat et al., Corros. Sci. 106 (2016) 172-178 [6].

6.5 Area Specific Resistance

Electrical resistivity of the oxide ( ρo ) follows the basic relation ρo=RA/lo where R is the electrical resistance of the oxide, A is the surface area of the oxide that electric current passes through, and l0 is the oxide thickness. From this relation, the area specific resistance (ASR) is defined as follows [86, 87]:

ASR=ρolo=RA(Ω cm2)(6.12)

This parameter is widely used to assess the electrical property of thermally grown oxide on metal because it simultaneously takes into account the proper electrical property (i.e. the resistivity of the oxide, ρo ) and the oxidation kinetic parameter (i.e. the oxide thickness, lo ) that is increased during oxidation [86, 87]. This parameter can be obtained by measuring R and A from the experiment. In the ASR measurement, if the interfacial electrical resistances between metal and electrodes are negligible, the electrical resistance measured from the experiment is attributed to the resistance of oxide and metal. The resistance of metal is significantly lower than that of the oxide as Huang et al. reported that the resistivity of Fe−26Cr−1Mo at 900∘C was 0.0141 mΩ cm2 [87]. Thus, the measured electrical resistance can be considered as the electrical resistance of the oxide (R) [87].

Fig. 6.3 shows ASRs of the bare ferritic stainless steels after oxidation at 800∘C in air [88-94]. The first and second bars from the left present, respectively, the results after 100 and 200 h of exposure. The other bars present the results after 300 h of exposure. After the 300 h exposure, it can be seen that stainless steel (Fe-17Cr) with 1wt.% of aluminium exhibits the highest ASR of about 20000 mΩ cm2 [91]. The ASR of AISI 441 is about 30 mΩ cm2 [91, 92], similar to that of ZMG 232 [91]. The model alloys in a family of FeCrMn(La/Ti) leading to the development of Crofer 22 APU have ASRs in the range of 10−15 mΩ cm2 [91] lower than the ASR of Crofer 22 APU which is in the range of 35−50 mΩ cm2 as reported by Hosseini et al. [93] and Ebrahimifar and Zandrahimi [94]. At shorter oxidation period i.e. 100 h , the ASR of Crofer 22 APU was also reported to be
about 50 mΩ cm2 [88]. For AISI 430, the ASR was reported to be about 35 mΩ cm2 after the exposure for 200 and 300 h [89, 90].

Many methods can be done to decrease the ASR of ferritic stainless steels. The first method is to dope the elements that could lead to the generation of electron in the conduction band of the oxide. For example, the titanium in titania may be doped into chromia by the jumping of titanium into the site of chromium in chromium oxide giving TiCr• and Cri−−∘. The formation of these defects are accompanied by the generation of electron presumably by Reaction 6.13, thus increasing the electrical conductivity of the oxide as experimentally observed in the titanium doping into the oxide thermally grown on AISI 441 [86]. However, it is noted that the titanium doping could not only increase the electrical conductivity of the oxide but also enhanced the steel oxidation rate [86]. The other methods are further looked for to decrease the ASR.

TiO2+CrCr↔TiCr•+Cri•••+4e′+O2(6.13)

[picture]

Fig. 6.3. Area specific resistance of ferritic stainless steel interconnects in air at 800∘C [88-94].

Fig. 6. 3. Area specific resistance of ferritic stainless steel interconnects in air at 800 ∘ C [88-94].

The second method to decrease ASR is to add the elements to promote scale adhesion. This is to avoid the separation of the scale from steel substrate which further results in the formation of air gap which has an infinite electrical resistivity. However, this addition must also be optimised with the possible altered oxidation rate. The roles of element on altering the scale adhesion will be described in the next section. The last method consists in covering the surface by a conductive coating that also enables to limit chromia formation.

6.6 Scale Adhesion

Scale adhesion can be affected by steel chemistry and gas composition in the atmosphere. For the steel chemistry effect, some elements e.g. lanthanum, molybdenum, titanium and niobium are added to stainless steels as seen in Table 6.1 and could affect the scale adhesion. Many testing methods have been developed to asses the scale adhesion test (as reviewed in Chapter 2). In this section, the adhesion test using the tensile method is described. The studied sample is placed in a micro-tensile
machine which is situated in a SEM chamber. The sample is strained in the SEM chamber at room temperature and the evolution of scale failure is observed during straining. At a given strain after spallation takes place, the spalled area is measured and normalised by the total area of the taken image, giving the spallation ratio. This parameter is used to assess the scale adhesion in the manner that higher spallation ratio indicates worse scale adhesion.

Fig. 6.4 shows the surface of AISI 441, 444 and Crofer 22 APU oxidised in simulated atmospheres at cathode side (synthetic air) and anode side (H2−2%H2O) of SOFCs at 800∘C for 200 h , after straining by the similar strain [36, 86]. The tensile loading was applied in the horizontal direction for the samples shown in this figure. It can be seen that the scale on Crofer 22 APU is best adhered to the steel substrate because the scale is failed only by the transverse crack without spallation on the major part of the oxide. This should be from the effect of reactive element i.e. lanthanum added to the steel. The scale on AISI 444 is badly adhered to the steel substrate as the high spallation ratio is observed in both atmospheres. After spallation in humidified hydrogen, Fe3Mo2 intermetallic compound was observed on the metal side at the internal interface between scale and the AISI 444, congruent to the thermodynamic prediction [36]. The formation of this compound is believed to be responsible to the worse scale adhesion.

[picture]

Fig. 6.4. Top view of the scales on AISI 441, AISI 444 and Crofer 22 APU oxidised at 800∘C in synthetic air (first row) and H2−2%H2O (second row) after tensile testing using the similar imposed strain. Reproduced from S. Chandra-ambhorn et al., J. Power Sources 171 (2007) 688-695 [36] and S. Chandra-ambhorn, PhD Thesis, INP Grenoble, France, 2006 [86].

Fig. 6. 4. Top view of the scales on AISI 441, AISI 444 and Crofer 22 APU oxidised at 800 ∘ C in synthetic air (first row) and H 2 − 2 % H 2 O (second row) after tensile testing using the similar imposed strain. Reproduced from S. Chandra-ambhorn et al., J. Power Sources 171 (2007) 688-695 [36] and S. Chandra-ambhorn, PhD Thesis, INP Grenoble, France, 2006 [86].

As for the effect of gas composition in the atmosphere on scale adhesion, Fig. 6.4 shows the surface of AISI 441 and 444 oxidised in different atmospheres (synthetic air and H2−2%H2O ) after straining by the similar strain. It can be seen that water vapour can promote the scale adhesion as for each grade the spallation ratio is less when it is oxidised in humidified atmosphere. It was proposed that water vapour could be dissociated to be hydroxyl ion [36, 95]. This species dissolves in the oxide and rapidly diffuses to the internal steel/scale interface because the ionic radius of the hydroxyl ion is less than that of the oxygen ion [36, 95]. As a result of the hydroxyl ion inward diffusion, scale is formed at the steel/scale interface, thus promoting the adhesion [36, 95].

6.7 Influence of Coatings

As an alternative approach to the bulk modification of alloys, a metallic interconnect can be surface-modified via a protective oxide layer. The protective layer is particularly important on the cathode side due to the oxidising environment and the susceptibility of SOFC cathodes to chromium poisoning. The protection layer is intended first to serve as a barrier to both chromium cation outward and oxygen anion inward diffusion, so that it can decrease the alloy oxidation kinetics and mitigate or even prevent chromium migration and formation of Cr (VI) oxide and oxide-hydroxide. Secondly, the protection layer helps minimise the area specific resistance between electrodes and interconnect by limiting the growth of the chromium-based oxide scale, which has relatively low conductivity. The promising coating materials for these purposes are at least perovskite or spinel coatings.

Perovskite has a formula of ABO3. A is a trivalent rare earth cation e.g. lanthanum, strontium and yttrium, while B is a transition metal cation such as chromium, manganese, cobalt, iron, copper and nickel [96]. It was reported that the compositions of iron, copper, nickel could be adjusted to increase the oxide electrical conductivity and TEC compatibility [10, 96]. The examples of perovskite compounds are lanthanum chromite (LaCrO3), lanthanum strontium chromite ( La1−xSrxCrO3 ), lanthanum strontium manganite ( La1−xSrxMnO3 ), lanthanum strontium cobalite (La1−xSrxMnO3) and lanthanum strontium ferrite (La1−xSrxMnO3). [96]. It was reported that the addition of strontium to the LaCrO3,LaCrO3 or La(Co,Fe)O3 could significantly increase the electrical conductivity of these oxides [97]. However, even though the doped LaCrO3 with p-type conducting property exhibits high electrical conductivity in the oxidising atmosphere [96, 98], its conductivity could be significantly reduced in a reducing atmosphere [10, 98]. Shaigon et al. [99] also commented that the perovskite coatings could allow the oxygen to transport to some extent and that it was difficult to develop a fully dense layer. Reactive element oxides could be coated to provide the reactive element effect including the reduced oxidation rate of the alloys. Fontana et al. [100] demonstrated that the addition of a nanometric reactive element oxide ( La2O3,Y2O3 ) layer applied by Metal Organic Chemical Vapour Deposition (MOCVD) drastically improved both corrosion rate and electrical properties in air of Crofer 22 APU and Haynes 230 alloys even after 20 months exposure at 800∘C in air. During aging in SOFC conditions, La2O3 transformed into LaCrO3. This underlying LaCrO3 perovskite layer contributed to better electrical behaviour, but was able to dope the chromia layer grown at the alloy surface, providing then the well known "reactive element effect" [101, 102].

For the spinel coating, this compound has a formula of AB2O4. A and B could be di-, tri- or quadri-valent cations in octahedral or hexahedral sites, while O is an oxygen anion located at face-centre-cubic lattice sites [96]. The types of elements and compositions of A and B can be adjusted to give the spinel with TEC compatible to the chromia and high electrical conductivity. Petric and Ling [103], and Paknahad et al. [104] measured these two properties when A and B were various elements as shown in Figs. 6.5 and 6.6.

Fig. 6.5 presents the TEC of AB2O4 spinel at 800∘C where the elements shown in x -axis are B . Different bars at a certain B element x -axis present the TEC of spinels consisting of different A elements. All spinels in the figure have a formula of AB2O4 except Cu−Mn spinel which has a formula of Cu1.3Mn1.7O4. It can be seen that most of the spinels which B is Fe tend to have the TEC in the range of about 11×10−6 to 12.5×10−6 K−1, compatible to chromia which has the TEC of about 9.6×10−6 K−1 at 25−900∘C [10] and Fe−20Cr ferritic stainless steel which has the TEC of about 12.5×10−6 K−1 at 800∘C [105]. The exception is ZnFe2O4 which its TEC is relatively lower i.e. 7×10−6 K−1. However, it is noticed that some other spinels i.e Cu1.3Mn1.7O4,MnCo2O4 and CuCo2O4 also have the TECs of 12.2×10−6,9.7×10−6,11.2×10−6 K−1, respectively, close to those of chromia and Fe-20Cr ferritic stainless steel.

[picture]

Fig. 6.5. Thermal expansion coefficients of different AB2O4 spinels at 800∘C[10,103−105].

Fig. 6. 5. Thermal expansion coefficients of different A B 2 O 4 spinels at 800 ∘ C [ 10 , 103 − 105 ] .

Figs. 6.6 (a) and (b) present electrical conductivities of the spinels. All spinels in the figures have a formula of AB2O4 and their conductivities were measured at 800∘C. The exception is Cu−Mn spinel which has a formula of Cu1.3Mn1.7O4 and its conductivity was measured at 750∘C. It can be seen that the electrical conductivities of spinels in Fig. 6.6(a) are less than 1.5 S cm−1. In Fig. 6.6(b), it is seen that Cu1.3Mn1.7O4 exhibits the highest value of the electrical conductivity of about 225 S cm−1 even though the test temperature was lower than those of the other spinels. At 800∘C,MnCo2O4 and CuCo2O4 exhibit the relatively high values of the electrical conductivities of about 60 S cm−1 and and 27.5 S cm−1 respectively. As mentioned in the former paragraph, these three spinels also have the TECs close to those of the chromia and Fe-20Cr ferritic stainless steel. Cu1.3Mn1.7O4, MnCo2O4 and CuCo2O4 are then promising to be used as coating material for the present application.

As for the examples of the spinel coatings, Yang coated Mn1.5Co1.5O4 on Crofer 22 APU [97, 99, 106]. After the exposure at 800∘C in air, ASR of the uncoated material was 40 mΩ cm2 after 400 h [97, 106], similar to the results shown in Fig. 6.3. The coating could reduce the ASR to 10 mΩ cm2 even after the exposure for 1000 h [97, 106]. After the longer exposure for 6 months under thermal cycling, the coating was found to be well adhered to the substrate and the sharp, thin profile of chromium across the interface between the subscale and spinel coating was observed [99, 106].

However, in some cases the spinel coating could also increase the chromium species volatilisation rate, such as the case of Cr−Mn spinel coating on AISI 441 conducted by Wongromrat et al. [5]. The discussion was made that the formation of chromium volatile species according to Reaction 6.9 might be from the non-elementary steps as shown in Reactions 6.14-6.16. In this mechanism, oxygen gas in the atmosphere may adsorb on the solid surface according to Reaction 6.14 or directly react with chromia to form CrO22+ according to Reaction 6.16. The adsorbed oxygen gas is reduced giving hydroxide ion according to Reaction 6.15. It was reported that the Mn−Co spinel is one of the good electrocatalysts for oxygen reduction found in the alkaline battery [5, 107, 108]. It is then possible that Reaction 6.15 is enhanced. If Reaction 6.15 is a rate determining step for the overall chromium species volatilisation, the increased rate of Reaction 6.15 then increases the overall rate of the chromium species volatilisation.

[picture]

Fig. 6.6. Electrical conductivities of different AB2O4 spinels at 800∘C [103, 104].

Fig. 6. 6. Electrical conductivities of different A B 2 O 4 spinels at 800 ∘ C [103, 104].

12O2( g)+s→12O2−s12O2−s+H2O(g)+2e′→2OH−+s12Cr2O3( s)+14O2( g)→CrO22++2e′

(6.14)(6.15)(6.16)

Recently, new strategies have been proposed in order to increase the oxidation resistance as well as the electrical conductivity. This new evolution in coating approach consists in associating several types of coatings in order to obtain positive properties from each of them. As an example, Balland et al. [109] pre-covered a Crofer steel by La2O3, performed an annealing treatment in order to form LaCrO3 perovskite coating prior to deposit metallic (Co,Mn) films. During aging of the coated samples, the metallic (Co,Mn) film transformed into (Co,Mn)3O4 which guaranteed good electrical conductivity. These duplex coatings inhibited both chromia sublayer formation (compared to the single (Co,Mn)3O4 layer) and the formation of iron oxide in dual atmosphere exposures (compared to the single La2O3 ).

6.8 Conclusions and Perspectives

The SOFC working temperature has been aimed at reducing to 800∘C or lower. Thus the application of stainless steel as an interconnect becomes possible. A ferritic stainless steel is selected for this application because of the compatibility of its thermal expansion coefficient (TEC) with a single cell. Chromia-forming stainless steel is further selected instead of the alumina-forming one to avoid the formation of insulating alumina during oxidation. The required properties of ferritic stainless steel for this application are not only the TEC compatibility and high electrical conductivity of the oxide formed at high temperatures, but also its low oxidation rate, low chromium species volatilisation and good scale adhesion in the SOFC environment.

The steel grades for this application start from stainless steel with chromium of about 16wt.% (AISI 430) and up to the grades with higher chromium content of about 18wt.% (e.g. AISI 441) and 22wt% (e.g. Crofer 22 APU and ZMG 232). Stabilisers and reactive elements can be added to the steel such as titanium and niobium in AISI 441, titanium and lanthanum in Crofer 22 APU, or zirconium and lanthanum in ZMG 232. Coating is recommended to be applied on steel surface in order to reduce the oxidation rate, avoid chromium species volatilisation and increase electrical conductivity of the steel. The promising coating materials are such as spinels e.g. Mn−Co,Cu−Mn or Cu−Co spinels, as well as peroveskites.

The tendency of SOFC development is to operate it at lower temperatures e.g. towards 600∘C. The possibility to use the fossil fuel in SOFCs requires the basic understanding in reactivity of interconnect materials in such environments, followed by the development of protective coatings.

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Keywords Index

A

Adhesion Energy Mechanical Behaviour of Thermal Oxide Scales on Stainless Steels

Alloying Element 3a4ab4d3-f5fe-4972-8701-6abfa04373a0; Effect of Water Vapour on the High Temperature Oxidation of Stainless Steels

Alloying Elements Thermodynamics and Kinetics of the High Temperature Oxidation of Stainless Steels

Alumina 3a4ab4d3-f5fe-4972-8701-6abfa04373a0; Effect of Water Vapour on the High Temperature Oxidation of Stainless Steels

C

Chromia 3a4ab4d3-f5fe-4972-8701-6abfa04373a0; Effect of Water Vapour on the High Temperature Oxidation of Stainless Steels

Chromium Volatilisation Characterisation of Thermal Oxide Scales on Stainless Steels

E

Energy Dispersive X-ray Spectroscopy Characterisation of Thermal Oxide Scales on Stainless Steels

F

Focused Ion Beam Characterisation of Thermal Oxide Scales on Stainless Steels

H

High Temperature Oxidation Thermodynamics and Kinetics of the High Temperature Oxidation of Stainless Steels; 3a4ab4d3-f5fe-4972-8701-6abfa04373a0; Effect of Water Vapour on the High Temperature Oxidation of Stainless Steels

I

Interconnect b7a883ce-6d33-436f-a9e9-08e0b160ec2c

Interface Mechanical Behaviour of Thermal Oxide Scales on Stainless Steels

Inverted Blister Test Mechanical Behaviour of Thermal Oxide Scales on Stainless Steels

K

Kinetics Thermodynamics and Kinetics of the High Temperature Oxidation of Stainless Steels

M

Mechanical Failure Mechanical Behaviour of Thermal Oxide Scales on Stainless Steels

Mode-mixity Mechanical Behaviour of Thermal Oxide Scales on Stainless Steels

Modified 4-point Bending Test Mechanical Behaviour of Thermal Oxide Scales on Stainless Steels

P

Photoelectrochemistry Characterisation of Thermal Oxide Scales on Stainless Steels

Plane Stress Analysis Mechanical Behaviour of Thermal Oxide Scales on Stainless Steels

R

Raman Spectroscopy Characterisation of Thermal Oxide Scales on Stainless Steels

S

Scanning Electron Microscope Characterisation of Thermal Oxide Scales on Stainless Steels

Solid Oxide Fuel Cells b7a883ce-6d33-436f-a9e9-08e0b160ec2c

Stainless Steel 3a4ab4d3-f5fe-4972-8701-6abfa04373a0; Effect of Water Vapour on the High Temperature Oxidation of Stainless Steels; b7a883ce-6d33-436f-a9e9-08e0b160ec2c

Stainless Steels Thermodynamics and Kinetics of the High Temperature Oxidation of Stainless Steels

Strain Mechanical Behaviour of Thermal Oxide Scales on Stainless Steels

Stress Mechanical Behaviour of Thermal Oxide Scales on Stainless Steels

T

Tensile Testing Mechanical Behaviour of Thermal Oxide Scales on Stainless Steels

Thermodynamics Thermodynamics and Kinetics of the High Temperature Oxidation of Stainless Steels

Thermogravimetric Method Characterisation of Thermal Oxide Scales on Stainless Steels

X

X-ray Diffractometer Characterisation of Thermal Oxide Scales on Stainless Steels

X-ray Photoelectron Spectroscopy Characterisation of Thermal Oxide Scales on Stainless Steels

Authors Index

B

Braccini, M. Mechanical Behaviour of Thermal Oxide Scales on Stainless Steels

C

Chandra-ambhorn, S. Thermodynamics and Kinetics of the High Temperature Oxidation of Stainless Steels; Mechanical Behaviour of Thermal Oxide Scales on Stainless Steels; 3a4ab4d3-f5fe-4972-8701-6abfa04373a0; Effect of Water Vapour on the High Temperature Oxidation of Stainless Steels; b7a883ce-6d33-436f-a9e9-08e0b160ec2c

Chandra-Ambhorn, W. Effect of Water Vapour on the High Temperature Oxidation of Stainless Steels; b7a883ce-6d33-436f-a9e9-08e0b160ec2c

Chevalier, S. b7a883ce-6d33-436f-a9e9-08e0b160ec2c

Combemale, L. b7a883ce-6d33-436f-a9e9-08e0b160ec2c

H

Hayashi, S. 3a4ab4d3-f5fe-4972-8701-6abfa04373a0

L

Latu-Romain, L. 3a4ab4d3-f5fe-4972-8701-6abfa04373a0

N

Nilsonthi, T. Mechanical Behaviour of Thermal Oxide Scales on Stainless Steels

P

Parry, V. Mechanical Behaviour of Thermal Oxide Scales on Stainless Steels

Pascal, C. Thermodynamics and Kinetics of the High Temperature Oxidation of Stainless Steels

Popa, I. b7a883ce-6d33-436f-a9e9-08e0b160ec2c

Promdirek, P. b7a883ce-6d33-436f-a9e9-08e0b160ec2c

S

Srisrual, A. Characterisation of Thermal Oxide Scales on Stainless Steels

T

Thublaor, T. Thermodynamics and Kinetics of the High Temperature Oxidation of Stainless Steels; Characterisation of Thermal Oxide Scales on Stainless Steels; Effect of Water Vapour on the High Temperature Oxidation of Stainless Steels

W

Wongpromrat, P. Characterisation of Thermal Oxide Scales on Stainless Steels; 3a4ab4d3-f5fe-4972-8701-6abfa04373a0; Effect of Water Vapour on the High Temperature Oxidation of Stainless Steels; b7a883ce-6d33-436f-a9e9-08e0b160ec2c