High Temperature Corrosion of Advanced Materials and Protective Coatings 0444889701, 9780444889706, 9780444599278, 0444599274

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HIGH TEMPERATURE CORROSION of Advanced Materials and Protective Coatings

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HIGH TEMPERATURE CORROSION of Advanced Materials and Protective Coatings Proceedings of the Workshop on High Temperature Corrosion of Advanced Materials and Protective Coatings Tokyo, Japan, December 5-7 1990

as part of the International Symposium Solid State Chemistry of Advanced

on Materials

Editors: Y. SAITO B.ΦNAY T. MARUYAMA Tokyo

Institute

ofTechnology

Japan

1992 NORTH-HOLLAND AMSTERDAM · LONDON · NEW YORK - TOKYO

ELSEVIER SCIENCE PUBLISHERS B.V. Sara B u r g e r h a r t s t r a a t 2 5 P.O. Box 211,1000 AE A m s t e r d a m , T h e Netherlands Distributors for the United States a n d Canada: ELSEVIER SCIENCE PUBLISHING COMPANY INC. 655 Avenue of the A m e r i c a s N e w York, N.Y. 10010, U.S.A.

Library

of Congress C a t a l o g 1 n g - 1 n - P u b l i c a t i o n

Data

W o r k s h o p on H i g h T e m p e r a t u r e C o r r o s i o n of A d v a n c e d M a t e r i a l s a n d Protective Coatings (1990 : Tokyo, Japan) Hiçjh t e m p e r a t u r e c o r r o s i o n of a d v a n c e d m a t e r i a l s a n d p r o c t e c t i v e c o a t i n g s : p r o c e e d i n g s of the W o r k s h o p on H i g h T e m p e r a t u r e C o r r o s i o n of A d v a n c e d M a t e r i a l s a n d P r o t e c t i v e C o a t i g s , T o k y o , J a p a n , D e c e m b e r 5-7 1990 as part of the I n t e r n a t i o n a l S y m p o s i u m on S o l i d S t a t e C h e m i s t r y of A d v a n c e d M a t e r i a l s / e d i t o r s , Y. S a i t o , B . Ö n a y , T. Maruyama. ñ. cm. I n c l u d e s b i b l i o g r a p h i c a l r e f e r e n c e s a n d index. ISBN 0-444-88970-1 1. C o r r o s i o n a n d a n t i - c o r r o s i v e s — C o n g r e s s e s . 2. H e a t r e s i s t a n t alloys—Corrosion—Congresses. 3. P o w e r - p l a n t s — E q u i p m e n t a n d supplies—Corrosion—Congresses. 4. P r o t e c t i v e c o a t i n g s — C o r r o s i o n -Congresses. I. S a i t o , Y. ( Y a s u t o s h i ) II. Önay, Â. I I I . M a r u y a m a , T. ( T o s h i o ) , 1 9 5 0 - . IV. I n t e r n a t i o n a l S y m p o s i u m on S o l i d S t a t e C h e m i s t r y of A d v a n c e d M a t e r i a l s (1990 : T o k y o , J a p a n ) V. T i t l e . T A 4 6 2 . W 6 2 3 1990 620. V 1 2 2 3 — d c 2 0 91-39147 CIP

I S B N : 0 444 88970 1 © 1992 ELSEVIER SCIENCE PUBLISHERS B.V. All rights reserved. No part of t h i s p u b l i c a t i o n m a y be r e p r o d u c e d , stored in a retrieval s y s t e m or t r a n s m i t t e d in any f o r m or by any m e a n s , electronic, m e c h a n i c a l , p h o t o c o p y i n g , recording or o t h e r w i s e , w i t h o u t the prior w r i t t e n p e r m i s s i o n of the publisher, Elsevier Science Publishers B.V., C o p y r i g h t & Permissions D e p a r t m e n t , P.O. Box 521,1000 A M A m s t e r d a m , T h e Netherlands. Special regulations for readers in t h e U.S.A. - This p u b l i c a t i o n has been registered w i t h t h e C o p y r i g h t Clearance Center Inc. (CCC), S a l e m , Massachusetts. I n f o r m a t i o n can be o b t a i n e d f r o m t h e CCC a b o u t c o n d i t i o n s under w h i c h p h o t o c o p i e s of parts of t h i s p u b l i c a t i o n m a y be m a d e in t h e U.S.A. A l l o t h e r c o p y r i g h t q u e s t i o n s , i n c l u d i n g p h o t o c o p y i n g o u t s i d e of t h e U.S.A., s h o u l d be referred t o t h e c o p y r i g h t owner, Elsevier Science Publishers B.V., unless o t h e r w i s e specified. No responsibility is a s s u m e d by t h e p u b l i s h e r f o r any i n j u r y and/or d a m a g e t o persons o r p r o p e r t y as a m a t t e r of products liability, negligence or o t h e r w i s e , or f r o m any use or o p e r a t i o n of any m e t h o d s , p r o d u c t s , instructions or ideas contained in t h e material herein. pp. 29-38, 51-60, 213-220, 227-232: C o p y r i g h t not transferred. Printed in The Netherlands

PREFACE

The papers in this volume were presented at the Workshop on High Temperature Corrosion of Advanced Materials and Protective Coatings held in Tokyo, Japan, December, 5-7, 1990 as part of the International Symposium on Solid State Chemistry of Advanced Materials. The workshop was organized by a committee chaired by Professor Yasutoshi Saito of the Tokyo Institute of Technology and attended by scientists from universities, research organizations and the industry. It was sponsored by numerous science foundations and industrial organizations whose names appear in a separate section in this volume. The workshop was the first, widely-attended international meeting held in Japan on High Temperature Corrosion since the symposium organized by the Japan Institute of Metals on the same topic in 1982. More than 100 scientists attended the workshop. Foreign attendees, most of whom were invited speakers, came from North America, Europe, Asia, and Australia. The workshop provided a valuable opportunity for foreign and Japanese scientists doing basic and applied research in the field to exchange views and hold discussions on topics related to the high temperature corrosion of engineering alloys, ceramics and protective coatings. The workshop was also intended to celebrate the 70th birthday of Keizo Nishida, Emeritus Professor of the Hokkaido University who has made valuable contributions to the field of high temperature corrosion through his scientific research and academic guidance. The papers were presented in different sessions at the workshop and therefore, are collected accordingly in this volume. The sessions included: (1) Fundamental Studies on High Temperature Corrosion, (2) High Temperature Corrosion of Engineering Alloys, ( 3 ) Hot Corrosion of Engineering Alloys and Corrosion of Nuclear Energy-Related Materials, (4) High Temperature Corrosion of Protective Coatings and Intermetallics, and (5) High Temperature Corrosion of Ceramic Materials. Following each session, a poster session was held for further discussion of the presented papers. Some of the poster-only papers submitted to the workshop are also included in this Proceedings. All the papers appearing in this volume are peer-reviewed to meet the high standards of a scientific publication. We would like to thank all scientists who reviewed the papers. We would also like to thank all the participants and sponsors of the workshop for their invaluable scientific and financial contributions.

Yasutoshi S Α Π Ό BülentΦNAY ToshioMARUYAMA

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vii

WORKSHOP ON HIGH TEMPERATURE CORROSION OF A D V A N C E D MATERIALS A N D PROTECTIVE COATINGS ORGANIZING COMMITTEE Chairman: Secretary:

Yasutoshi SAITO Biilent Ö N A Y Hisao FUJIKAWA YoshioHARADA Hiroshi ICMMURA Shigemitsu KIHARA Makoto KIKUCHI Toshio MARUYAMA Hiroshi NAGAI Toshio NARITA Rikio ÍÅÌÏÔÏ Mikio OBAYASHI Masaru OKUYAMA Yutaka SHINATA Sokichi TAKATSU Atsushi TAKEI Shigeji TANIGUCHI Mikio YAMANAKA Masayuki YOSHEBA

Tokyo Institute of Technology Tokyo Institute of Technology Sumitomo Metal Industries Ltd TOCALO Corp., Ltd. Nippon Rare Earth Co., Ltd. Ishikawajirna-Harima Heavy Industries Co., Ltd. Tokyo Institute of Technology Tokyo Institute of Technology Osaka University Hokkaido University Nippon Yakin Kogyo Co., Ltd Toyota Central Research and Development Laboratory, Inc Oyama Technical College Akita University Toshiba Tungaloy Co.,Ltd. National Research Institute of Metals Osaka University Nippon Steel Corp. Tokyo Metropolitan University

INTERNATONAL ADVISORY BOARD M. J. BENNETT D.L. DOUGLASS Teiichi HONMA Tatsuo ISHKAWA P. KOFSTAD Tatsuo KONDO K . L . LUTHRA Ohmi MTYAGAWA Taishi MOROISM Kazuyoshi ÍÐ Akira OHTOMO

(U. K . ) (U.S.A.) (JAPAN) (JAPAN) (NORWAY) (JAPAN) (U.S.A) (JAPAN) (JAPAN) (JAPAN) (JAPAN)

F. S. PETTIT R. A RAPP M. SCHÜTZE Toshio SHIBATA Takanobu SHNOHARA W. W. SMELTZER F.H. STOTT R. STREIFF J. STRINGER Ryohei T A N A K A W.L.WORRELL

(U. S. A . ) (U.S.A.) (GERMANY) (JAPAN) (JAPAN) (CANADA) (U.K.) (FRANCE) (U. S. A . ) (JAPAN) (U.S.A.)

viii

SPONSORS OF WORKSHOP ON HIGH TEMPERATURE CORROSION OF ADVANCED MATERIALS A N D PROTECTIVE COATINGS Tokyo Ohka Foundation for the Promotion of Science and Technology The Asahi Glass Foundation Nippon Sheet Glass Foundation Nippon Mining Co., Ltd. Sumitomo Metal and Mining Co., Ltd. TDK Corp. Japan Tobacco Inc. Nippon Steel Corp. Sumitomo Metal Industries Ltd. Ishikawajima - Harima Heavy Industries Co., Ltd. Nippon Yakin Kogyo Co., Ltd. Toyo Engineering Corp. Toshiba Corp. Yoshida Kogyo K.K. Toray Industries, Inc. Seiko Instruments, Inc. TOCALO Corp., Ltd. Bando Chemical Industries Ltd. Chichibu Cement Co., Ltd. Ebara Research Co., Ltd. Fuji Electric Corporate Research and Development Ltd. Fujikura Ltd. Hitachi Chemical Co., Ltd. JGC Corp. Kawasaki Steel Corp. Kobe Steel Ltd. MAC SCIENCE, Co. Matsushita Research Institute Tokyo, Inc. Mitsubishi Heavy Industries, Ltd. Mitsubishi Kasei Corp. Mitsubishi Materials Corp. Mitsubishi Steel Mfg., Co., Ltd. NEC Corp. NGK Insulators, Ltd. NIDAK Corp. Nihon Cement Co., Ltd. Nippon Metals Industry Co., Ltd. Nippon Rare Earths Co., Ltd. Nippon Yttrium Co., Ltd. NKK Corp. Rigaku Corp. Shimadzu Corp. Sinku - Riko, Inc. The Japan Steel Works, Ltd. Toshiba Tungaloy Co., Ltd. Tosoh Corp. Toyota Central Research & Development Laboratory, Inc. Ube Industries, Ltd. Uchida Rokakuho Publishing Co., Ltd.

IX

T A B L E OF C O N T E N T S Preface

í

List of Committees

vii

List of Sponsors

viii

FUNDAMENTAL STUDIES ON HIGH TEMPERATURE CORROSION OF ADVANCED MATERIALS On the Use of the Wagner Model in Oxidation in Mixed Reactants P. KOFSTAD and R. BREDESEN

3

Some Principal Mechanisms in the Simultaneous Erosion and Corrosion Attack of Metals at High Temperature D. RISHEL, F. PETTIT and N . BIRKS

13

Grain Boundary Segregation in Ionic Solids and Its Effect on High Temperature Heterogeneous Kinetics J. N O W O T N Y

29

Failure of Oxide Scales on Advanced Materials Due to the Presence of Stresses M. SCHÜTZE

39

Real Time Studies of Scale Development and Failure M. J. BENNETT

51

Electrochemical Determination of Sulfide Growth on Iron in Sulfur Environments T. A . R A M A N A R A Y A N A N , J. D. MUMFORD, H. VEDAGE and S. N . SMITH

61

IR - RAS Study of Oxide Film Formed on SUS430 Stainless Steel in Moist Atmospheres K. HONDA, T. A T A K E and Y . SAITO

67

HIGH TEMPERATURE CORROSION OF ENGINEERING ALLOYS Oxidation Behavior of Fe-Al-Si Alloys at 1073 and 1173 Ê S. GUAN, J. CORKUM and W . SMELTZER

75

High-Temperature Oxidation of Iron - Aluminum Alloys R. PRESCOTT, D.F. MITCHELL, G. I. SPROULE, R. J. HUSSEY and M . J. G R A H A M

83

÷

Oxidation of Fe-Cr-Mn-Al Stainless Steels K. KUROKAWA, Y . MIZUTA and H. TAKAHASHI

91

Structure and Oxidation Behavior of the Scale Formed on Al-Containing Ferritic Stainless Steel S. S A S A Y A M A and T. K A M I Y A

97

Cyclic Oxidation Behavior of Microcrystallized CoCrAl Alloy Film F. W A N G , H. LOU and W . W U

103

High Temperature Oxidation of Ni - Cr Alloys T. A M A N O and O. MOMOSE

Ill

Effects of Oxygen and Water Vapor Pressures on Oxidation of Iron - Chromium Alloys at 573 Ê T. TSUJI, S. KOBAYASHI, M . O D A and Ê. NAITO

115

High Temperature Oxidation of Heat Resistant Stainless Steels in COG Combustion Environment B. G. SEONG, J. H. SONG, S. Y . H W A N G and Ê. Y . K I M

123

Internal Nitridation of Ni-Cr-Al Alloys R. P. RUBLY and D. L. DOUGLASS

133

Internal Bromine Corrosion of Dilute Ni-Cr Alloys B. Ö N A Y , Y . SAITO and T. A T A K E

147

Sulfidation Properties of Low Alloy Steels in H2S-H2 Atmospheres T. NARITA, T. ISHIKAWA and M . N A K A M O R I

155

High Temperature Oxidation of Ni - Cr Alloys with Small Additions of Si and Ce T. A M A N O and O. MOMOSE

163

Effect of Lanthanoid on Oxidation Behavior of Fe-Cr-Al Foil K. OHMURA, M . Y A M A N A K A , M . F U K A Y A and H. A B O

167

Formation of a LaCrOß Particle Dispersed Alloy Layer as a Means for Improving Oxidation Resistance H. KONNO and R. FURUICHI

177

HOT CORROSION OF ENGINEERING ALLOYS A N D CORROSION OF NUCLEAR ENERGY - RELATED MATERIALS The Role of Applied Creep Stress on Hot Corrosion Behavior of a Nickel-Base Superalloy M. YOSHIBA

187

xi

Corrosion Resistance of Advanced Tube Materials in Coal-Fired Boilers S. KIHARA, K. N A K A G A W A , W. WOLOWODIUK, J. L. BLOUGH and W.T. B A R K E R

197

Hot Corrosion of Commercial Tube Steel Materials in a Japanese Waste Incinerator Environment N. OTSUKA and T. KUDO

205

Corrosion Performance of Ultra-High-Purity Chromium Sheet in High Temperature Aggressive Environments Y . SHIMIZU, M . K A T O , T. S A K A K I and M . YOSHIBA

213

AC Impedance and Electrochemical Techniques for Evaluating Hot Corrosion Resistance C.X. W U , A . N I S H I K A T A and T. TSURU

221

Characterization of Corrosion of Solid Metals in Flowing Liquid Metals T. SUZUKI and I. MUTOH

227

High Temperature Corrosion of Hastelloy X R in HTGR Helium Environment T. TSUKADA, M . SHINDO, T. SUZUKI, H. N A K A J I M A and T. KONDO

233

Oxidative Vaporization of Simulated Fission-Produced Noble Metal Alloys T. MATSUI, T. HOSHIKAWA and K. N A I T O

243

HIGH TEMPERATURE CORROSION OF PROTECTIVE COATINGS A N D INTERMETALLICS The Influence of Oxidation on the Sliding Friction and Wear of Ceramic-Coated Steel at High Temperature F. H. STOTT, D. R. G. MITCHELL and G. C. WOOD

255

Hot Oxidation of Arc Ion-Plated Nitride Coatings A . K A W A N A and H. ICHIMURA

267

Corrosion Resistance of Iron Coated by Plasma Spray Ceramic Coatings to Molten Fluoride M. O K U Y A M A , M . K A W A K A M I and Ê . ITO

275

High Temperature Corrosion Behavior of Air-Cooled Alloys in Combustion Gas and Protective Effect of MCrAlY-VPS Coatings M. N A K A M O R I

285

Silico-Aluminized Coating on Molybdenum and Its High Temperature Oxidation Resistance T. M A R U Y A M A , X . F. BI and K. N A G A T A

291

xii

Sulfate Induced Corrosion Behavior of FeAl Intermetallic Compound in O 2 -0.5%(SO2+SO 3 ) Atmosphere at 600-870°C W. WU, Y . NIU, J. GUO and Y . ZHANG

301

A Chromia-Pack Pretreatment for Improving the Oxidation Resistance of TiAl S. TANIGUCHI, T. SHIBATA and S. SAKON

309

Aluminized Coatings on Titanium Alloys and TiAl Intermetallic Compound A. TAKEI and A . ISHIDA

317

Effect of Ternary Alloying Elements on the Oxidation Behavior of Ti-Al Intermetallic Compound Y . SHIDA and H. A N A D A

325

HIGH TEMPERATURE CORROSION OF CERAMIC MATERIALS Effect of High Temperature Oxidation on Flexural Strength of Sintered S13N4 K. IMAI, H. H I R A T A and N . OKABE

335

High-Temperature Oxidation of CVD-SiC in CO-CO2 Atmosphere T. NARUSHIMA, T. GOTO, Y . IGUCHI and T. HIRAI

345

Reaction of Graphite with Gaseous Silicon Monoxide H. IMAI, K. FUJII and S. N O M U R A

351

V205"Na2S04-NaCl Molten Salt Corrosion Behavior of Various High Temperature Structural Ceramics H. W A D A and M . YOSHIBA

355

Hot Corrosion of Reaction-Sintered S13N4 in Molten Na2S04 Y. SHINATA, M . HARA, T. N A K A G A W A and C. SHIMIZU

365

List of Participants

373

Author Index

377

Subject Index

379

FUNDAMENTAL STUDIES ON HIGH TEMPERATURE CORROSION OF ADVANCED MATERIALS

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High Temperature Corrosion of Advanced Materials and Protective Coatings Y . Saito, B. Önay and T. Maruyama (Editors) © 1992 Elsevier Science Publishers B.V. All rights reserved.

3

ON THE USE OF THE WAGNER MODEL IN OXIDATION IN MIXED REACTANTS

Per Kofstad, Department of Chemistry, University of Oslo, P.B.1033 Blindem, 0315 Oslo 3, Norway Rune Bredesen, Center for Industrial Research, P.B.124 Blindem, 0314 Oslo 3, Norway

The Wagner theory has been one of the most important contributions to our understanding of high temperature oxidation of metals by providing a fundamental understanding of the basic features of growth of dense scales. The applicability of the theory has been demonstrated for the reaction of some high-purity metals in single reactants (e.g. O2 and S2). However, for reactions in mixed reactants under conditions where only one reaction product is expexted to be formed (e.g. oxidation of Fe in CO+CO2 and H2+H2O mixtures), the parabolic rate constants may deviate significantly from that predicted by the theory and as observed in the single reactant (e.g. oxidation of Fe in O2). Possible reasons for this, and particularly effects of mixed gases on defect conditions at interfaces, are discussed.

1. INTRODUCTION The Wagner theory has been one of the most important contributions to the field of high

where the applicability can be tested, and - as will be discusssed later - oxidation of high purity cobalt in oxygen probably serves as a good example.

temperature corrosion and generally serves as a basis for interpretations of protective oxidation of 1

In interpretations of high temperature oxidation in

metals and alloys . However, the model is an ideal

mixed reactants, e.g. in CO+CO2 or H2+H2O it is

one and in its original form it was limited to growth

generally assumed that it is only necessary to

of dense scales controlled by lattice diffusion of

consider the activity of oxygen in the mixed gases in

atoms or ions of the reactants or by electron

the use of the Wagner model. By way of example,

transport through the scales. Over the years it has

for oxidation of cobalt in CO+CO2 mixtures, it is

been increasingly realized that grain boundary

accordingly only necessary to consider the oxygen

diffusion also may be an important, if not the

activity of the gas mixture in predicting the

predominant, diffusional transport mechanism in

temperature and gas pressure dependence of the

growth of continuous scales on many metals,

parabolic rate constant. The aim of this contribution

particularly for highly protective scales and at

is to discuss this aspect and the validity of this

reduced temperatures. The following discussion will be confined to reactions for which lattice diffusion predominates.

assumption in more detail. But as a start some basic features and a brief summary of the theory will be presented.

Furthermore, it will be illustrated by metal-oxygen reactions, but in principle it also applies to other gas-metal reactions.

2. THE WAGNER OXIDATION T H E O R Y

1 -3

The basic assumption of the original theory is

The applicability of the Wagner theory has

that lattice diffusion of the reacting atoms or ions or

generally been tested for corrosion of high purity

transport of electrons through dense scales is rate-

metals in single reactants. As regards metal-oxygen

determining in the overall reaction. As lattice

reactions there are relatively few such systems

diffusion takes place because of the presence of

4 2

point defects, the transport processes may

rate constant in c m / s e c , the expression for Wp is

alternatively be considered in terms of transport of

given by

defects.

ο

As lattice diffusion or electron transport is rate determing, reactions at the interfaces (outer and

k p=

\

j

(gjO

M

+ D o ) d l n p 02

(2)

inner surfaces of the scales) are considered to be rapid, and it is assumed that thermodynamic equilibria are established at the interfaces, i.e.

where D M and D o are the random self-diffusion

between oxygen gas and the oxide at the outer

coefficients of the metal and oxygen ions in the

surface and between the metal and the oxide at the

oxide, respectively, z c and z a the valences of the

inner surface. Furthermore, local equlibria are

cations and anions, and PQ? and ρ ' θ2 are the

assumed to exist within the scales.

oxygen pressures (activities) at the outer ( o ) and inner (i) interfaces of the scale.

This type of reaction is parabolic with time, and

The validity and applicability of the model may be

the differential and integrated forms of the rate

tested by comparing directly measured values of the

equation are given by

parabolic rate constants with values calculated from Eq.2 from independently measured values of the

dx df

=

.k 1 px

x2 = 2 k pt + C = k pt + C

. 1. a. < >

self-diffusion coefficients in the oxide. In order to

(1b)

oxygen pressure dependence of the self-diffusion

evaluate the integral, it is necessary to know the coefficients. In the following the applicability will be

where χ denotes the oxide thickness, kp ( = 2 k p) is

tested by first considering oxidation of an

the parabolic rate constant, and C the integration

appropriate metal in 0 2 a n d then in mixed reactants,

constant. The oxide thickness χ may alternatively be

e.g. CO+CO2 and H2+H2O mixtures.

expressed as oxygen uptake (weight gain) per unit surface area of the scale surface.

3. REACTIONS IN SINGLE REACTANTS 3.1 Oxidation of Co to CoO.

In his theory Wagner derived an expression for the parabolic rate constant by making use of the fact

As mentioned above, high temperature oxidation of high purity cobalt in oxygen gas serves as an

that the growth rate rate of the dense scale is the

example where the validity of the basic features of

sum of the fluxes of the metal ions (cations) and the

the model may be tested.

oxygen ions (anions) through the scale and that no

The defect structure and defect-dependent

net current flows through the scale. The rate

properties of CoO have been extensively studied by

constant may then be expressed in terms of the

many investigators . The oxide is metal-deficient

electrical conductivity and the transport numbers of

and it is concluded that the predominant defects at

1

the ions and electrons in the scales or alternatively -

near atmospheric oxygen pressures are singly

for electron conducting scales - in terms of the self-

charged cobalt vacancies. In terms of the Kröger-

diffusion coefficients of metal and oxygen ions in the

Vink notation cobalt vacancies are written, V Q 0 ,

oxide. In the latter case and when the growth rate of

where a, which is the effective charge of the defects,

the oxide is expressed in cm/sec and the parabolic

is equal to one for CoO.

5

The self-diffusion coefficient of cobalt in CoO,

TEMPERATURE

Deo. is several orders of magnitude larger than that

Ο ο

Ο δ

iû

of oxygen, Do, i.e. D c o » D o - Furthermore, as z c=

m

I

, ι

ο

ο ο

ο

Δ V •

P a ï d a s s i et a l . K r ü g e r et a l . S n i d e et a l .

Ο

B r i d g e s et al ( i n t e r p o l . val.)

°x

Po2

kp=~

D Co d l n p 0 2

j

(3)

N o 2

C

^

L_,_

|z a| for this system, Eq.2 is then simplified to ο

c

t^J

2

L

'C

Ο δ

O x i d a t i o n of Co in air.

Ε

Po2

At near-atmospheric oxygen pressures experimental results show that the oxygen pressure dependence of the cobalt tracer self-diffusion coefficient, Dco> can be approximately expressed b y tr D Co

nt r , 1 D

1

£ ο ζ ο

=> < LL er u_

Λ.

β

-

1 0

C o - t r a c e r d i f f u s i o n in h

o,

CoO . air.

α ^ _j ce ο

LU CO Ο
« 1200 °C) where lattice diffusion may be

surface layer compared to that in the bulk, and then

expected to dominate, the ratio of kVDJ^'j is similar

a reasonable agreement is obtained between

0

to that for cobalt. Again, this may be taken as a

measured and calculated values of the parabolic

confirmation of the basic aspects of the validity of

rate constants as shown above for high temperature

the Wagner model.

oxidation of cobalt and and nickel. However, if one

However, at reduced temperatures the values of k p for oxidation of nickel are orders of magnitude

oxidizes these metals in mixed gases where two (or more) components may interact with the surface

higher than that expected for lattice diffusion

and significantly affect the defect concentration in

controlled oxidation. These high values of k p are

the surface layer, more significant effects may

concluded to reflect that grain boundary diffusion of

possibly be obtained. In this respect it is of interest

nickel in NiO is the predominant mode of diffusional

to consider oxidation of cobait and iron in gas

transport in growth of NiO scales at reduced

mixtures such as C O + C 0 2 a n d H2+H2O mixtures.

temperatures. These aspects are of great importance in the field of high temperature corrosion, but we shall in the following limit the

4.1 Oxidation of Co in CO+CO2 mixtures. When Co is oxidized in CO+CO2 mixtures,

discussion to scale growth controlled by lattice

thermodynamic properties of bulk materials predict

diffusion.

thai CoO is the only stable phase for C02:CO ratios larger than about 96:4 in the temperature range 1000-1200 °C. Thus at correspondingly higher

4. REACTIONS IN MIXED REACTANTS/INTERACTANTS In this use of the Wagner model an important implicit assumption has been made. When one

C 0 2 : C O ratios Co is oxidized to CcO. The bulk solubility of carbon is in all probaiity small (< a few p p m )

10

and much smaller than the

integrates Eq.3 over the scale surface from the outer

native concentration of cobalt vacancies, and

to the inner interface and calculates the vaiue of the

accordingly it is reasonable to assume that carbon

parabolic rate constant, one makes use of the self-

dissolved in the bulk has no significant effect on

diffusion coefficient measured in the bulk oxide and

equiiibrium vacancy concentrations in the bulk, at

implicitly assumes that the self-diffusion coefficient

least at near-atmospheric oxygen pressures. It is not

in the surface or surface layer is exactly the same as

uncommon to assume that the same applies to

in the bulk. (In Eq.6 it is seen that the parabolic rate

interfacial regions and thus that the presence of CO

constant (for growth of metal-deficient scales) are

and/or C 0 2 molecules (or carbon) have no

7

significant effects on the defect structure of CoO

ι

surfaces exposed to CO+CO2 mixtures and that the

ι

«

ι

ι

ι

«

I

ι

I

I

ι

ι

Τ"

1000°C

only parameter of importance is the oxygen activity

Co-^CoO

of the gas and in the surface layer. Such an assumption may be correct, but in view of the large




K

dt=

ξ

"

k

kc

F I G U R E 21 Erosion Enhanced Corrosion - T y p e I I . This r e g i m e is characterized by a corrosion rate constant greater than that o f corrosion only, and suggests that erosion increases the scaling rate by processes distinct f r o m but concurrent w i t h scale thinning.

where:

< k'

C o r r o s i o n p r o d u c t t h i c k n e s s is s m a l l , but affects p r o p e r t i e s of the m e t a l or alloy s u r f a c e v i a c o m p o s i t e s c a l e formation. Metal Recession

χ

= »k= t

F I G U R E 23 Oxidation Affected Erosion. This r e g i m e can be characterized as w h e r e the corrosion product thickness is small. H o w e v e r its influence on the metal or alloy surface properties, via composite scale formation is likely to be large; for this r e g i m e , it is appropriate to quantify material degradation in terms o f metal recession.

27

REFERENCES 1.

B a r k a l o w , R . H . , G o e b e l , J. A . and F . S. Pettit, Materials P r o b l e m s in F l u i d i z e d - B e d C o m b u s t i o n Systems: H i g h T e m p e r a t u r e E r o s i o n - C o r r o s i o n b v H i g h V e l o c i t y (200 m / s ) Particles. Pratt & W h i t n e y A i r c r a f t G r o u p , M a y 1980, E P R I C S - 1 4 4 8 , P r o j e c t 9 7 9 - 4 , Final R e p o r t .

2.

K a n g , C . T . , Pettit, F.S. and N . Birks, "Mechanisms in the Simultaneous E r o s i o n - O x i d a t i o n A t t a c k o f N i c k e l and C o b a l t at H i g h T e m p e r a t u r e s , " M e t . Trans. Α . . V o l . 18, ( 1 9 8 7 ) , p p . 1785.

3.

Proceedings o f the 6th International C o n f e r e n c e on Erosion b v L i q u i d and Solid I m p a c t . "On the C o m b i n e d E f f e c t s o f C o r r o s i o n and E r o s i o n , b y S. H o g m a r k , A . Hammersten and S. Soderberg," ( U n i v e r s i t y o f C a m b r i d g e , 1983, p p . 37-1 through 37-8).

4.

Wright, L G . , Nagarajan, V . and J. Stringer, O b s e r v a t i o n s o f the R o l e o f O x i d e Scales In H i g h Temperature Erosion-Corrosion of Alloys, "Oxidation o f M e t a l s . V o l . 25, N o s . 3/4. ( 1 9 8 6 ) , p p . 175-199.

5.

T e d m o n , C S . Jr., " T h e E f f e c t o f O x i d e V o l a t i l i z a t i o n on the O x i d a t i o n K i n e t i c s o f C r and F e - C r A l l o y s " L E l e c t r o c h e m . S o c . V o l . 113. ( 1 9 6 6 ) . DP. 766.

6.

L i u , Y . Y . and K . Natesan, M e t h o d o l o g i e s f o r Predictionof Metal Oxidation-Vaporization-Erosion. A N L F E - 8 8 - 2 , ( A r g o n n e , Illinois: A r g o n n e N a t i o n a l L a b o r a t o r y , U n i v e r s i t y o f C h i c a g o , D e c e m b e r 1987).

7.

Srinivasan, Y . and K . Y e d u l a , e d . , C o r r o s i o n and Particle Erosion at H i g h T e m p e r a t u r e . " T h e E r o s i o n Corrosion B e h a v i o r o f N i c k e l in M i x e d O x i d a n t A t m o s p h e r e s , b y D . M . R i s h e l , F . S. Pettit and N . Birks" ( W a r r e n d a l e , Pa: T h e M i n e r a l s , M e t a l s and Materials Society, 1989), p p . 265-314.

8.

L e v y , Α . V . , S l a m o v i c h , E. and N . Jee, "Elevated T e m p e r a t u r e C o m b i n e d E r o s i o n - C o r r o s i o n o f Steels," W e a r . V o l . 110, N o . 2 , (July 1986), pp. 117-150.

9.

Chang, S. L . , Pettit, F.S. and N . Birks, "Effects o f A n g l e o f I n c i d e n c e on the C o m b i n e d E r o s i o n O x i d a t i o n A t t a c k o f N i c k e l and Cobalt," O x i d a t i o n o f Metals in p r i n t ) , ( 1 9 8 9 ) .

10.

M r o w e c , S. and K . P r z b y l s k i , "Transport Properties o f Sulfide Scales and Sulfidation o f M e t a l s and A l l o y s , " O x i d a t i o n o f M e t a l s . V o l . 23, N o s . 3 / 4 , ( 1 9 8 5 ) , pp. 107-139.

11.

H a f l a n , B . and P. K o f s t a d , " T h e R e a c t i o n o f N i c k e l w i t h S 0 2 + O z / S O j at 500-900°C," C o r r o s i o n S c i e n c e . V o l . 23, N o . 12, ( 1 9 8 1 ) , p p . 1333-1352.

12.

H o c k i n g , M . G . and P. S. S i d k y , "The H o t C o r r o s i o n o f N i c k e l - B a s e d T e r n a r y A l l o y s and Superalloys f o r

Gas T u r b i n e A p p l i c a t i o n s - I I . T h e M e c h a n i s m o f Corrosion in SO2/O2 A t m o s p h e r e s , "Corrosion Science. V o l . 27, N o . 2 , ( 1 9 8 7 ) , pp. 205-214. 13.

E v a n s , A . G . , C r u m l e y , G . B . and R . E . D e m a r a y , "On the M e c h a n i c a l B e h a v i o r o f Brittle Coatings and L a y e r s , " O x i d a t i o n o f M e t a l s , v o l . 20, no. 516, ( 1 9 8 3 ) , p p . 193-216.

14.

Birks, Ν . and G . H . M e i e r , Introduction to H i g h Temperature Oxidation o f Metals. (London: Edward A r n o l d , 1983), p p . 172-122.

15.

L e v y , A . and Y . F. M a n , "Elevated T e m p e r a t u r e E r o s i o n - C o r r o s i o n o f 9 C r - l M o Steel," W e a r . V o l . I l l , ( 1 9 8 6 ) , pp.135-159.

16.

Srinivasan, Y . and K . V e d u l a , e d . , C o r r o s i o n and Particle Erosion at H i g h T e m p e r a t u r e s . "Mechanisms o f C o m b i n e d E r o s i o n - C o r r o s i o n o f Steels at E l e v a t e d Temperatures", b y A . V . L e v y ( W a r r e n d a l e , P A : T h e M i n e r a l s , Metals and Materials Society, 1989), pp. 207-230.

17.

Stephenson, D.J., N i c h o l l s , J. R . and P. H a n c o c k , "Particle-Surface Interactions D u r i n g the Erosion o f a Gas T u r b i n e M a t e r i a l ( M a r M 0 0 2 ) b y P y r o l y t i c C a r b o n Particles," C o r r o s i o n Science. V o l . 25, N o . 12, ( 1 9 8 5 ) , p p . 1181-1192.

18.

Stephenson, D.J., N i c h o l l s , J.R. and P. H a n c o c k , "The Interaction B e t w e e n Corrosion and Erosion D u r i n g Simulated Sea Salt Compressor Shedding In M a r i n e Gas Turbines," W e a r . V o l . I l l , ( 1 9 8 6 ) , pp. 15-29.

19.

V a n der Z w a a g , S. and J.E. F i e l d , T h e E f f e c t o f T h i n H a r d Coatings on the H e r t z i a n F i e l d , " Philosophical M a g a z i n e A . V o l . 46, N o . 1, 1982, p p . 133-150.

20.

Johnson, K . L . , Contact M e c h a n i c s . ( C a m b r i d g e : C a m b r i d g e U n i v e r s i t y Press, ( 1 9 8 5 ) , pp. 345

21.

Blazynski, T . Z . , e d . , Materials at H i g h Strain R a t e s . "Surface Response T o Impact" , b y J. E. F i e l d and I . M . Hutchings ( L o n d o n and N e w Y o r k : E l s e v i e r A p p l i e d Science, 1987, pp. 243-293.

22.

K o l s k y , H . , Stress Waves in Solids. ( N e w Y o r k : D o v e r Publication, Inc., 1963), pp. 2 4 - 3 8 .

23.

R i n e h a r t , J.S. and J. Pearson, B e h a v i o r o f Metals Under Impulsive Loads. ( N e w York: Dover Publication, Inc., 1965), pp. 28-44.

24.

R i n e h a r t , J.S., O n Fractures Caused by Explosions and Impact," Quarterly o f the C o l o r a d o School o f M i n e s . V o l . 55, N o . 4, ( O c t o b e r 1960).

25.

K h a n , A . S . and F. K . Irani, " A n E x p e r i m e n t a l Study o f Stress W a v e Transmission at a M e t a l l i c - R o c k Interface and D y n a m i c T e n s i l e Failure o f Sandstone,

28 L i m e s t o n e , and Granite," M e c h a n i c s o f M a t e r i a l s . V o l . , 6, ( 1 9 8 7 ) , pp. 285-292. 26.

A c h e n b a c h , J.D., W a v e Propagation in Elastic Solids. ( N e w Y o r k : A m e r i c a n E l s e v i e r Publishing C o m p a n y , Inc., 1973), pp. 165-198.

27.

K i n s l o w , R . , ed., H i g h V e l o c i t y I m p a c t P h e n o m e n a . " T h e o r y o f Impact on T h i n Targets and Shields and Correlation w i t h Experiment", b y J. W . G e h r i n g , Jr. ( N e w Y o r k : A c a d e m i c Press, 1970), p p . 463-514.

28.

M i l l e r , R . K . and H . T . T r a n , " R e f l e c t i o n , R e f r a c t i o n , and A b s o r p t i o n o f Elastic W a v e s at a Frictional Interface: Ρ and S V M o t i o n , " Trans. ASME A p p l i e d M e c h a n i c s . V o l . 48, ( M a r c h 1981), p p . 155160.

29.

B o w d e n , F . P. and J. E. F i e l d , " T h e Brittle Fracture o f Solids by L i q u i d Impact, b y Solid I m p a c t , and b y Shock," P r o c . R o v . Soc. L o n d o n Ser. Α . . V o l . 282, ( 1 9 6 5 ) , p p . 331-352.

30.

P r e e c e , C M . , ed., Treatise on Materials Science and T e c h n o l o g y . "Impact D a m a g e Mechanics: Solid Projectiles," b y A . G . Evans ( N e w Y o r k : A c a d e m i c Press, 1979), p p . 5 8 - 6 2 .

31.

Z a m b e l l i , G . and Α . V . L e v y , "Particulate Erosion o f N i O Scale," W e a r . V o l . 68, ( 1 9 8 1 ) , pp. 3 0 5 - 3 3 1 .

32.

C h a n g , S.L., Pettit, F.S. and N . B i r k s , "Some Interactions in the Erosion-Oxidation of A l l o y s , " O x i d a t i o n o f Metals ( i n press), ( 1 9 8 9 ) .

High Temperature Corrosion of Advanced Materials and Protective Coatings Y . Saito, B. Önay and T. Maruyama (Editors) 1992 Elsevier Science Publishers B.V.

29

GRAIN BOUNDARY SEGREGATION IN IONIC SOLIDS AND ITS EFFECT ON HIGH

TEMPERATURE

HETEROGENEOUS KINETICS Janusζ Nowotny Australian Nuclear Science and Technology Lucas Heights, NSW 22 34, Australia

Organisation,

Advanced

Materials,

The defect chemistry of interfaces such as external surfaces and grain boundaries of non-stoichiometric oxides differs from that of the bulk phase as a result of segregation. The defect structure of the interface region usually involves enrichment in both cation and anion vacancies as well as interstitials. Consequently, segregation results in the formation of both chemical and electrical potential gradients in the boundary layer. The gradients have a strong effect on the diffusion in the interface layer. Both mechanism and kinetics of the diffusion within the boundary layer differ substantially from that in the bulk phase. This paper will discuss segregation in binary metal oxides and its effect on the high temperature transport kinetics both along and across interfaces such as grain boundaries.

1. INTRODUCTION Interface

may result in different transport mechaprocesses

such

as

grain

nisms than those

boundary diffusion play an important role

determination

in high temperature

kinetics

and

alloys.

oxidation

Therefore,

of metals

better

under-

related

is the subject

netics may lead to more efficient

ties.

An

better

sion.

properties

ported

diffusion

data

in the literature

have

been r e -

for many

ionic

to

local

the

transport

microstructure

of substantial

awareness

bition of high temperature gaseous corroBulk

the

and microchemistry of the boundary

standing of grain boundary transport k i inhi-

in the bulk phase. The

of

is

understanding

of

is important

layer

difficul-

growing grain

that

a

boundary

for correct in-

terpretation of grain boundary

transport

kinetics. So far, this knowledge

is r e -

solids. In contrast little is know about

stricted to approximate models which are

grain

not

boundary

transport

nonstoichiometric studies

kinetics

compounds.

in

Current

in this matter are limited to a

few oxide systems and nickel oxide

such as alumina [1-4] [5-9]. Modelling of the

adequate

nature

of

disorder

to

both

describe

the

crystalline

within

the grain

complex

and

boundary r e -

gion. Detailed knowledge of defect istry of the boundary

defect

layer

chem-

is required

grain boundary transport kinetics is the

for quantitative considerations of diffu-

subject

sion processes in this layer.

of conflicting

reports

even for

these binary compounds. It

is

expected

grain boundary defects,

diffusion

however,

This paper will discuss the effect

that

surface

and

occur by point

segregation-induced

structural changes in the boundary

layer

of segregation on the defect structure of the

grain

boundary

nonstoichiometric

region

of

oxides as well as the

impact of the defect structure on the in-

30

terface displacement of charged de-fects.

undoped CoO is different from that in the

The

bulk phase

transport

kinetics

both

along

and

across interfaces will be considered. The effect

phase boundary,

of

segregation-induced

chemical

potential

interface

gradients

and

the

in

[12]. It has been

that

the

vicinity

of

documented

the

C o O / C o 30 4

after the enrichment

in

Co vacancies surpasses a certain critical

corresponding electric fields on gas/so-

value,

lid equilibration processes will be con-

cobalt interstitials which have not been

the

defect

sidered in more detail.

observed for the bulk phase

2. INTERFACE DEFECT CHEMISTRY

ture

also been observed

It has been documented that defects of

ionic

solids

as a result

segregate to

of

the

excess

interfaces

of

interface

is

formed

structure

involves

[13]. It has

that the C o 30 4 struc-

at

the

surface

of

CoO

grains within the stability range of the CoO phase. Fig. 1 represents the phase

diagram

near

the

boundary

exhibits a gradient of defects and corre-

the CoO phase stability, in which a C o 30 4-

sponding chemical composition. The gradi-

type overlayer is formed [12].

tions

in a wide range of

between

the

defects

changes

in

defect

the

900 CoO PHASE

continuous

mobility

within

interac-

within

boundary region. Accordingly,

the range,

phase

energy. Concordantly, the interface layer

ent results

illustrating

T-p(0 2)

C o O / C o 30 4

within

the ^

interface layer can be expected [10, 1 1 ] .

PHASE

When the concentration of defects in the

interface

critical

layer

value

surpasses

a

structural

certain

reordering

4.0

takes place resulting in the formation of low

dimensional

structures

l o g Po 2tPo 2in P a l

[11]. These

structures exhibit extraordinary properties not displayed by bulk phases. Tran-

45

Fig. 1. The T-p(0 2) phase diagram for the near-surface layer of undoped CoO [12]

sport properties of these structures may be entirely different bulk

phase.

effect

Little

of the

structures

known

diffusion its role

in

the

about

low dimensional

on

consequently,

from those of

is

the

interface

solids

in high

and,

tempera-

The effect of the interface defects chemistry and the corresponding

chemical

and electrical gradients on the transport kinetics

along

and

across the

interface

will be analyzed in this paper.

ture oxidation of metals and alloys. Segregation may

conditions for

of

one

type

lead to establishment the

of

of

in the grain boundary

transport

kinetics

ally assumed that both surface and inter-

structure

of

which

boundary

has been described in several fundamental publications [10, 1 4 ] . It has been gener-

studies

defects

Grain

BOUNDARIES

are

properties

of

region

3. DIFFUSION ALONG GRAIN

not stable within the initial structure. Recent

formation

defect

favourable

surface

electrical

have

shown

that

the

within

the

boundary

defect

layer

of

face diffusion diffusion.

Fig.

are more rapid than 2 illustrates the

matic representation

bulk sche-

(in the form of the

31

type of polycrystalline

a)

NiO

formed as a

scale of oxidized Ni while ceramic materials prepared by sintering do not exhibit enhanced grain boundary diffusion [6]. Recent

studies of Moya et al.

shown that

the grain

[10] have

boundary

transport

enhancement is independent of the preparation procedure of NiO ceramic material. It has also been argued that the lack of experimental diffusion b)

evidence

studied

of

grain

either

by

boundary

measurement

of radiotracer penetration profiles or by the

autoradiography

considered

method

cannot

as a corroboration

of a

be lack

of interface diffusion enhancement [10]. It has been assumed that the diffusion

mechanism

involves along

a

the

surfaces

along

rapid

grain

transport

interface and

layer

subsequent

boundaries of

species

or

internal

slow

lattice

diffusion into the bulk phase. By assum-

τ-'

ing that the thickness of the grain boundary region is very small and limited to

Fig. 2. Schematic illustration of different diffusion coefficients in a polycrystalline material (a) and their representation in the Arrhenius plot (b)

1-2

atomic

Suzuoka

Arrhenius-type diagram) of the relation-

et

solutions enable

layer al. of

one

both [17,

Whipple 18]

Fick's

to

have

second

determine

the

sion

responding to different processes such as

enrichment factor. Extensive

along

lattice grain

diffusion

diffusion boundaries

( D L) ,

diffusion

D

( gb) ,

surface

(D s) and diffusion across grain

boundaries

and

proposed law

which

product

of

the grain boundary thickness, the diffu-

ship between diffusion coefficients corbulk

[16]

coefficient

and

the

segregation description

of the grain boundary formalism has been reported by Peterson

[14]. The solutions

proposed

and

by

Whipple

Suzuoka

et

al.

[16-18], certainly valid for metals, have

(D s*) .

There have been conflicting

reports

also been applied

for

concerning the effect of the preparation

as nonstoichiometric

of NiO on diffusion

9].

[5-9]. Fig. 3 illus-

ionic solids oxide

Since the thickness

ceramics

of the

such [5-

boundary

trates the relationship between the grain

layer for metal oxides can be much larger

boundary diffusion, dislocation diffusion

than that of metals the validity

and bulk

lattice diffusion

according

solutions

Atkinson

and

It

Taylor

[5].

argued that the grain boundary enhancement

occurs

only

for

has

to

been

transport a

certain

derived

for

metals

of the

requires

verification. In the theoretical model assumed for grain boundary diffusion it has also been

32

TEMPERATURE [ ° C ] 1400 1000 700

assumed that the grain boundary diffusion 500

coefficient

( D ' ) , considered

sponding

very

to

the bidimensional

fast

as

corre-

transport

within

boundary

layer,

grain

is much higher than the lattice diffusion coeffient

( D ) . However, the grain bound-

ary of metal-deficient oxide is enriched in cation vacancies

and,

therefore,

the

lattice diffusion within the grain boundary region should be consistent with the local

defect

structure.

It

should

be

considered as a function of the distance from

the

gradual 6

8 10 4 -1 -1 10 Τ [ K l

12

Fig. 3. Bulk, grain boundary and dislocation diffusion in undoped NiO according to Atkinson and Taylor [5]

boundary.

change

coefficient

In

of the

within

consequence,

lattice

the

a

diffusion

boundary

layer

should be expected as illustrated in Fig. 4 for the grain boundary

structure pro-

posed by Duffy and Tasker

[19] for NiO.

Taking pects,

into

which

nonstoichiometric

account are

the

above

specific

compounds,

the

asto

appli-

cability of the proposed solutions of the diffusion equation for the determination of the grain boundary cient of

diffusion

ionic compounds

coeffi-

should be con-

sidered with great caution.

4. D I F F U S I O N A C R O S S T H E

INTERFACE

Diffusion across interfaces involves lattice transport normal to the interface along

electrical

and

chemical

potential

gradients in the boundary layer (Fig. 5 ) . CONCENTRATION

POTENTIAL

l= a

Fig. 4. The Duffy and Tasker [19] grain boundary structure of NiO and expected changes in the defect mobility within the grain boundary region

Fig. 5. Schematic representation of the diffusion across the boundary layer

33

This diffusion plays an important role in

4

I

all gas/solid equilibration processes. In contrast

to diffusion

along

interfaces,

which has received major attention in the literature,

very

little

is

known

3

about

i

diffusion across interfaces. In

the

consideration

of

*

i f

gas/solid

/

2

heterogeneous processes it has been generally assumed that transport across the interface is relatively fast and, therefore, that the reaction kinetics is pure-

1 migration effect inegative)

ly bulk diffusion controlled. Most of the

/

/

/

/

diffusion data reported in the literature have

been

determined

based

on

this

as-

sumption.

-

4

-

3

-

2

-

1

0

1

2

3

4

Studies of the equilibration kinetics,

performed

for

Cr-doped

NiO

as

a

model material, have shown that the segregation-induced

electric

field

within

Fig. 6. The dependence between the kD/, Fe-26Cr-4Al > Fe-15A1 >

suitable for 0 / S I M S transport studies. Scales

18

Fe-25A1. The curves are characterized by an

formed at 1100°C were somewhat more uniform.

initial rapid oxidation changing to a much slower

Fig. 2a shows a layer of C C - A I 2 O 3 completely

rate, particularly when OC-AI2O3 forms a

covering the Fe-15A1; oxide ridges or convolutions

continuous layer. The presence of OC-AI2O3 after

are better seen at higher magnification in Fig. 2b.

oxidation of Fe-15A1 and Fe-25A1 at 1000° and 1100°C was confirmed by RHEED.

Time, h

FIGURE 1 Oxidation kinetics of Fe-Al alloys at 1000° and 3 1100°C in 5xl0" torr oxygen 2 ( ^ g / c m r 5 n m α-Α1 20 3).

FIGURE 2 SEM micrographs of oxide formed on Fe-15A1 after 2h oxidation at 1100°C.

86

scale has spalled off revealing the underlying alloy whose surface is quite uneven with ridges, depressions, steps and grooves. 3.2

Auger Depth Profiles

Fig. 4 shows Auger depth profiles through the oxide formed at 1100°C on Fe-25A1 (a) and Fe15A1 (b). The level of Fe in the scale on Fe-25A1 is just about at the Auger detection limit (-0.1%); the presence of Fe in small amounts (0.04-0.10%) was confirmed by SIMS. The Fe content is higher (-1% as deteraiined by Auger and SIMS) in the outer oxide on Fe-15A1, and as seen in Fig. 4b the content increases towards the oxide/metal interface indicating the presence of an Fecontaining oxide phase. The oxide/metal interface on both alloys is broad due to the non-uniformity of the oxide. 100

80

L

0

Fe-25A1, 1100°C

10

20

40

30

50

Sputter Time, min 100

1

Fe-15A1, 1100°C

80

FIGURE 3 SEM micrographs of oxide formed on Fe-25A1 after 2h oxidation at 1100°C.

Small oxide nodules are also observed in this micrograph, both on the ridges and on the depressions between the ridges. The oxide morphology on Fe-25A1 is similar to that on Fe15A1 in that ridges and nodules are also present (Fig. 3a). The scale appears to have cracked (presumably on cooling) along the top of the ridges. Often flakes of scale detached completely from the substrate; Fig. 3b shows areas where

ο

20

40

60

80

Sputter Time, min

FIGURE 4 Auger depth profiles of oxides formed after 2h oxidation at 1100°C on Fe-25A1 (a) and Fe-15A1 (b). Sputtering was by 4 keV argon.

87

3.3

significant outward cation diffusion and

SIMS Analysis and Oxide Growth

18

significant ingress of 02 via cracks, fissures or

Fig. 5 shows SIMS profiles of oxide

interconnected porosity, and would indicate that

formed on Fe-25A1 at 1100°C. The oxide was 16

produced in two stages - firstly, in 0 2 for l h and 18

then in 0 2 for l h with the specimen maintained at temperature during the gas changeover.

inward oxygen diffusion is the major transport process. In the absence of other data, one could 18

conclude that the 0" profile shows predominantly lattice diffusion of oxygen and that 18

the 0 ' shoulder at the oxide/alloy interface illustrates a minor short-circuit (eg. grain boundary) diffusion component (10,19). However, it may be difficult to make the distinction between lattice and grain boundary oxygen diffusion when the oxide grain size is small. In the present work the

OC-AI2O3 grain size is

only -100-300 nm, as

indicated by SEM and confirmed by TEM, and so atoms in the short-circuit paths could exchange with those in the lattice and the result would be difficult to distinguish from lattice diffusion if the oxygen self diffusion coefficient is S 1 0

- 14

2

cm /sec. 18

Further, the data in Fig. 5 represent average 0" values obtained over a fairly large area (-300 μπι χ 300 μπι) of a somewhat non-uniform oxide (Fig. 3).

Sputter Time, min

Clearly, it would be advantageous to obtain SIMS data on a finer scale. Using a liquid metal ion gun it is now possible to analyze by SIMS areas as

FIGURE 5

2

SIMS profiles of Fe-25A1 oxidized at 1100°C first in 16 2 0 2 for l h (oxygen uptake, 60 μg/cm ) followed by 18 2 lh in 0 2 (oxygen uptake, 8 μg/çm , Shown are 18 2 the !60- and 0"(x5)signals and thç S6Fe 7Al+ profile which indicates the position of the alloy/oxide interface.

small as 1 μπι . Fig. 6 shows SIMS images of the same specimen as in Fig. g obtained with a Ga ion gun. 18

A sequence of 0" images are presented at a series of depths through the scale (Fig. 6a - j). In Fig. 6a, near the outer oxide surface, the white patches represent - 50%

Shown are 5 6

2 7

Fe Al

+

16

18

18

0", the white lines 18

- 25% 0", the grey areas - 2% 0", and the

18

0 " and 0 " (x5) profiles and the

signal which provides an indication of

the oxide/metal interface. As seen in the figure, 18

the 0 " profile falls from an initial average value of -20% at the outer oxide surface to a low limiting

black areas are where patches of oxide have spalled off. The sequence of photos shows that as we move towards the alloy surface the white rich patches disappear, the amount of

18

0~-

18

0"

value at the oxide/alloy interface. There is a

associated with the lines remains constant

noticeable increase in the signal as the oxide/alloy

(although some lines disappear) and the low level

interface is approached. (The area under the

18

0"

18

of 0" in the grey areas is maintained. In 18

profile corresponds to -11% of the total oxide in

addition, white dots containing - 50% 0" appear

agreement with the weight gain data). The

as the oxide/alloy interface is approached. (Also

general shape of the profile would exclude

shown in Fig. 6 are

16

0" images taken near the

88

outer oxide surface (Fig. 6i) and near the oxide/alloy interface (Fig. 6j). The images are 18

complementary to the 0~ images in Fig. 6b and h, respectively.) It can be concluded from the images that the

18

0~

1 8

0 distribution is localized and

non-uniform. The

18

0 " - rich lines appear to

correlate with oxide ridges readily observed in 18

SEM stereo pairs. The 0 " dots appearing near the oxide/alloy interface are consistent with oxygen short-circuit diffusion and the formation of new oxide grains at the alloy surface. The outer 18

white 50% 0 " patches may represent an average 18

of 100% 0 " at the oxide grain boundaries and a low value over the grains of the fine-grained (100300nm) alumina. SIMS profiles through the white patches to the oxide/alloy interface are similar to that in Fig. 5. Why some patches of 18

oxide contain 50% 0 " and some only 2% is not clear. It may have something to do with localized loss of contact of the oxide during growth. Clearly, the growth mechanism of OC-AI2O3 is more complex than simply oxygen grain boundary diffusion. Stress effects and failure of the oxide on a microscale are obviously important. Work is 18

continuing to try to correlate the 0 " images with finer structural features of the oxide, and to obtain All8ol60polyatomic SIMS (12,15) images to distinguish between oxygen diffusion and exchange processes. 4.

SUMMARY Protective OC-AI2O3 scales have been formed

on high Al-content Fe-Al alloys at 1000 and 1100°C.

18

0 / S I M S analysis indicates that the

scales grow by oxygen transport. SIMS images of 2

FIGURE 6 Series of SIMS images for the same sample as 1 6 Fig. 5, i.e. Fe-25A1 oxidized first in 0 2 and then 18 1 8 in θ 2 · (a-h) are 0 " images taken at the following percentage through the oxide film towards the oxide/alloy interface: a, 3%; b, 15%; c, 25%; d, 35%; e, 45%; f, 60%; g, 70%; h, 80%. i and j 1 6 are 0 " images taken at 8% and 85%, respectively.

small areas (a few μιη ) show that the

1 8

0

distribution within the scale is localized at particular regions; nearly half of the surface has 18

not oxidized at all in θ 2 · The results demonstrate that OC-AI2O3 growth is more complex than SIMS profiles would suggest.

89

ACKNOWLEDGEMENT

14.

M. J. Graham, J. I. Eldridge, D. F. Mitchell and R. J. Hussey, Mater. Sei. Forum, 42, 207 (1989).

15.

R. J. Hussey, D. F. Mitchell and M. J. Graham, Werkstoffe und Korros. 2S> 575 (1987).

16.

R. J. Hussey, P. Papaicovou, J. Shen, D. F. Mitchell and M. J. Graham, 'Corrosion and Particle Erosion at High Temp.' Ed. V . Srinivasan and K. Vedula, The Minerals, Metals and Materials Soc, ρ 567 (1989).

17.

M . J. Graham, Proc. Conf. 'Microscopy of Oxidation', Cambridge, U.K., March 1990. The Institute of Metals (in press).

18.

M . J. Graham and M . Cohen, J. Electrochem. Soc. 112, 879 (1972).

19.

A. Atkinson, Rev. Mod. Phys. £Z, 437 (1985).

The authors thank J.W. Fraser for his assistance with the S E M examination.

REFERENCES 1.

J. K. Tien and F. S. Pettit, Metall. Trans. 2, 1587(1972).

2.

T. A. Ramanarayan, R. Ay er, R. PetkovicLuton and D. P. Leta, Oxid. Met. 22, 445 (1988).

3.

J. Jedlinski and S. Mrowec, Mater. Sei. and Eng. SZ, 281 (1987).

4.

E. W . A. Young, H. E. Bishop and J. H. W . De Wit, Surf, and Interface Anal. 2, 163 (1986).

5.

E. W . A. Young and J. H. W . De Wit, Solid State Ionics lfi, 39 (1985).

6.

E. W . A. Young and J. H. W . De Wit, Oxid. Met. 2fi, 351 (1986).

7.

G. B. Abderrazik, G. Moulin, A. M . Huntz, E. W . A. Young and J. H. W . De Wit, Solid State Ionics 22, 285 (1987).

8.

A. M. Huntz, G. B. Abderrazik, G. Moulin, E. W . A. Young and J. H. W . De Wit, Appl. Surf. Sei. 28, 345 (1987).

9.

Κ. P. R. Reddy, J. L. Smialek and A. R. Cooper, Oxid. Met. 12, 429 (1982).

10. W . J. Quadakkers, H. Holzbrecher, K. G. Briefs and H. Beske, Oxid. Met. 32, 67 (1989). 11. W . J. Quadakkers, W . Speier, H. Holzbrecher and H. Nickel, Proc. Conf. 'Microscopy of Oxidation', Cambridge, U.K., March 1990. The Institute of Metals (in press). 12. J. I. Eldridge, R. J. Hussey, D. F. Mitchell and M. J. Graham, Oxid. Met. 3Q, 301 (1988). 13. J. I. Eldridge, D. F. Mitchell, R. J. Hussey and M . J. Graham, MRS Intl. Mtg. on Adv. Mats. Vol 4, ρ 377. The Materials Research Soc. (1989).

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High Temperature Corrosion of Advanced Materials and Protective Coatings Y . Saito, B. Önay and T. Maruyama (Editors) © 1992 Elsevier Science Publishers B.V. All rights reserved.

91

OXIDATION OF Fe-Cr-Mn-Al STAINLESS STEELS

Kazuya KUROKAWA, Yoshio MIZUTA and Heishichiro TAKAHASHI Metals Research Institute, Fuculty of Engineering, Hokkaido University, Sapporo 060, Japan

The oxidation behavior of Fe-10Cr-Mn-3Al alloys containing up to 25% manganese was studied at temperatures ranging from 873 to 1173 Κ in 1%02~He gas mixture. The reacted specimens were characterized by X-ray diffraction and SEM-EDX. In all the alloys used in the present study, mass gain values enough to measure were obtained only at 1173 K. The oxidation at the temperature roughly obeyed a parabolic rate law, and the oxidation rates were significantly increased with the manganese content. However, the addition of aluminum caused a remarkable improvement in the oxidation resistance of Fe-Cr-Mn alloys. The increased oxidation resistance is due to the reduced transport rate of cations through the oxide, alumina or aluminate spinel, resulting in the formation of the spinel oxide, MnFe2Û^, without formation of wustite which has high diffusivity of cations.

1. INTRODUCTION Fe-Cr-Ni stainless steels such as Type 316

4.0% manganese indicated that the oxide scales

are attractive as candidate materials for first

were generally similar to those formed on the

wall and structural components for fusion reac-

parent metals and remarkable voids were formed

tor applications.

in the substrate.

However, the calculated

radioactivity decay curves for various pure

Douglass et a l . ^ ' ^ studied

the oxidation of Fe-9.5Cr-17.8Mn and Fe-19.6Cr-

elements after exposure to neutron flux in a

1$.1Mn ternary alloys.

fusion reactor indicates that elements such as

low Cr alloy, in which extensive wustite scale

nickel and molybdenum are converted into long-

was formed, had insufficient chromium to form a

-1

lived radioactive isotopes .

In order to sim-

They showed that the

protective Cr^O^ scale.

In the high Cr alloy,

plify waste disposal for fusion reactor compo-

no wustite formation occured and spinel oxide

nents, the development of reduced-activation or

whose growth rate was rate controlling was

fast induced-radioactivity decay steels is in

formed at temperatures ranging from 973 Κ to

progress.

1273 K.

The replacement of nickel in steels

with manganese would produce steels which show

In this investigation, in order to improve

an acceptable fast induced-radioactivity decay.

the oxidation resistance of low Cr steel con-

However, oxidation behavior of these steels

taining up to about 25% manganese the addition

must also be considered in the evaluation of

of 3% aluminum was made.

As will be discussed,

the compatibility of Fe-Cr-Mn alloys with com-

a considerable improvement in oxidation rate

mercial grade helium-coolant with anticipated

was observed.

low concentrations of impurities such as O2 and H2O in ordinary operations.

Furthermore, pos-

sible effects of accidental exposure of a Fe-

2. MATERIALS AND METHODS Four Fe-Cr-Mn-Al alloys were used in this

Cr-Mn alloy-structure to air must be also

investigation.

considered.

the phases present in each alloy are given in

A previous study

on the oxidation behavior

of iron-based binary alloys containing up to

The chemical composition and

Table 1, where the concentration of each element is given in weight %.

The phase structure

92

TABLE 1 Chemical composition and phase structure of Fe-Cr-Mn-Al alloys. Alloy

Ni

A B C D

0.01 0.01 0.13 0.01

9.88 15.03 24.73

0.002Λ 0.0018 0.25 0.002Λ

Ni equiv.

C

Cr

Al

Si

0.003 0.005 0.10 0.004

10.08 10.22 11 . 7 7 9.85

2.99 2.93 2.37 3.07

0.01 0.01 0.12 0.01

N

Mn

Phase

Cr equiv.

5.10 7.72 16.85 12.55

Ferrite Ferrite + Austenite Ferrite + Austenite Ferrite + Austenite

26.55 26.36 25.05 26.76

Figure 1 shows a conventional Schaeffler diagram with all the alloys used in the present work.

In this figure, M, F and A indicate

martensite, ferrite and austenite-stable regions,'respectively.

The phases observed

for each alloy agree with the Schaeffler diagram.

In other words, only the Fe-10Cr-1OMn-

3A1 alloy is a ferritic steel, and the others consist of ferritic and austenitic phases. The alloys were cut into samples measuring approximately 5 mm by 7 mm by 1 mm, which were 0

10 Cr

20

30

polished to a 1 micron meter diamond finish. The oxidation tests were carried out at

e q u i v a l e n t / m a s s °/o

temperatures ranging from 873 Κ to 1173 Κ for 2.16x10^ seconds in 1%0 2-He gas mixture.

FIGURE 1 Schaeffler

diagram o f Fe-Cr-Mn-Al

alloys.

The

specimen was hung in a quartz reaction tube from a microbalance with a Pt wire, next to a thermo-couple.

was determined by X-ray diffraction.

The con-

The gas mixture was introduced

into the reaction tube after evacuation.

The

centrations of chromium and aluminum as fer-

oxidation test was started by elevating the

ritic stabilizers are approximately fixed in

electric furnace, which was kept at the desired

all the alloy.

temperature until the sample was in the hot

Namely, the amount of chromium

is about 10% and that of aluminum is about 3%.

zone.

The composition of manganese as an austenite

in about 1.8x10

stabilizer varies from 10% to 25%.

mass was automatically recorded.

Here, the

nickel and chromium equivalents were calculated according to the following relationships-^

The prescribed temperature was reached sec.

The change in sample

The identification and observation of scales formed on the alloys were synthetically examined by X-ray diffraction and scanning elec-

Ni

equiv.=(Ni)+(Co)+0.5(Mn)+0.3(Cu)+25(N)+30(C)

Cr

equiv.=(Cr)+2(Si)+1.5(Mo)+5(V)+5.5(Al)+ 1.75(Nb)+1.5(Ti)+0.75(W)

tron microscopy with EDX.

3. RESULTS AND DISCUSSION 3.1. Scaling kinetics

where the concentration of the respective elements given in parentheses is in weight %.

The oxidation kinetics obtained at 1173 Κ are shown in Figure 2.

The oxidation of each

93

Oxid temp : 1173 Κ

F(?-9.85Cr-24.73Mn-3.07Al L Oxidation time : 21.6 ks

9,8ΘΜη-2.99ΑΙ

0

10

900

20

Oxidation time / ks

1000

1100

1200

1300

Temperature / Κ

FIGURE 2

FIGURE /,

Scaling kinetics of some Fe-10Cr-Mn-3Al alloys at 1173 Κ in flowing 1%02-He gas mixture.

Temperature dependence of mass gain of Fe-10Cr25Mn-3Al alloys after oxidation for 21.6 ksec.

oxid.temp. : 1173 Κ significantly increases with the Mn content. -7

For example, Kp value of the 25% Mn alloy is about two orders of magnitude higher than that of the 10% Mn alloy.

Ο

Ε

However, even for the 25

% Mn alloy, a mass gain was little observed at

\ α -10

temperatures less than 1073 K, as shown in

/° /

Figure Λ ·

The temperature in first wall and

structural components of fusion reactors would 1

i

10

15

Mn

.

be kept at temperatures less than 973 K.

1

20

25

content / mass °/o

Therefore, it is recognized that the oxidation resistance of the present alloys containing 3% aluminum is quite high, especially at lower

FIGURE 3 Variation of the parabolic rate constant at 1173 Κ in flowing 1%02-He gas mixture with Mn content in Fe-10Cr-Mn-3Al alloys.

temperatures. As seen in Figures 2-4-, there were marked differences in oxidation resistance of the a l loys used in the present study, although the respective concentrations of chromium and alu-

alloy proceeded according to a parabolic rate

minum were approximately identical in all the

law, and the amount of mass gains of the a l -

alloys.

In general, the volume fraction of the

loys strongly depended on the Mn content.

ferritic phase in Fe-Cr-Mn alloys increases

The relationships between the Mn content and

with decreasing manganese content, and the dif-

the parabolic rate constant at 1173 Κ are shown

fusivities of chromium and aluminum in the

in Figure 3·

ferritic phase are much greater than those in

The parabolic rate constant

94 the austenitic phase.

Thus, it is expected

1273K

1173K

1073K

973K

that corrosion resistance decreases with the volume fraction of the austenitic phase.

In

the present study, the alloy having the highest fraction of austenitic phase is alloy C, as seen in figure 1.

On the other hand, the alloy

which showed the highest oxidation rate is a l loy D. Therefore, it is concluded that the oxidation resistance of Fe-Cr-Mn-Al alloys strongly depends on the manganese content rather than the amount of a given phase in the alloy. Figure 5 shows a comparison between the

-10

oxidation behavior of the alloys containing _n ι

aluminum and that of the alloys without aluminum.

7.5

In this figure, the parabolic rate con-

•

.

.

.

80

8.5

90

9-5

T"

stants measured in the present study are plotted only for 1173 K, because mass gains at temParabolic rate Fe-10Cr-Mn-3Al

peratures less than 1073 Κ were too small to measure parabolic rate constants.

1

/

. — ι 100

10.5

1öV

FIGURE 5 c o n s t a n t s f o r o x i d a t i o n o f some alloys.

As seen in

this figure, there is a large difference in the

formed on the present alloys at 1173 K.

corrosion resistance of the alloys with and

the 10% Mn alloy, the oxidized surface was

without aluminum, although the respective con-

covered with a thin layer, possibly of alu-

centrations of chromium and manganese are

mina, and by small oxide protrusion.

similar.

protrusions were believed to be of manganese

Namely, the addition of aluminum to

Fe-Cr-Mn alloys provides high oxidation resis-

oxide from the result of EDX analysis.

tance, especially in low-manganese alloys. 3.2.

In

These

Such a

structure was observed also in Fe-Mn-Al alloys for which the growth mechanism was already dis-

Scale morphology

Figure 6 shows scanning electron micro-

cussed by Jackson and Wallwork^.

In alloys

with higher Mn concentration, the results of

graphs of the external surfaces of the scales

FIGURE 6 Scanning e l e c t r o n micrographs o f t h e e x t e r n a l a f t e r o x i d a t i o n f o r 2 1 . 6 k s e c . a t 1 1 7 3 K.

surfaces

o f the scales

formed on Fe-10Cr-Mn-3Al

alloys

95

Cross-sectional micrographs o f the scale k s e c . a t 1 1 7 3 K.

FIGURE 7 formed on t h e Fe-2.8Cr-36Mn a l l o y

after

oxidation f o r 21.6

X-ray diffraction and EDX analyses showed that MnFe2Û^ and/or Fe^O^ layer was formed at the top surface of scales. Figure 7 shows the cross-section of the Fe2.8Cr-36Mn ternary alloy after oxidation at 1173 K.

This external scale mainly consists

of FeO-MnO solid-solution.

It should be

noticed that the thick external scale and many pores formed in the substrate beneath the scale were significant.

These pores may be intercon-

nected, and the formation may be attributed to the high vapor pressure of manganese. Moreover, Jackson and Wallwork suggested that the vapor phase transport of metal from the substrate to the scale could be an important factor in the ρ development of scale .

1

On the other hand, in the alloys containing 3% aluminum, such a thick external scale and

c

C

iL

pores in the alloy matrix were not formed. Figure 8 shows the cross-sectional micrograph

>> Β

rü

and the characteristic X-ray intensities of elements at indicated points in the scale

-Li

formed on the Fe-1OCr-25Mn-3Al alloy at 1173 K. MnFe2Û^ and Fe^O^ were identified in the scale by X-ray diffraction taken directly from the scale.

AI

CrMrvCr —

— Fe

10(eV) Fe.Mn

These oxide phases are believed to

exist in regions D and E, respectively. Moreover, according to the result of the EDX analysis, the marked enrichment in aluminum is found in the region C adjacent to the alloy

FIGURE 8 C r o s s - s e c t i o n a l micrograph and c h a r a c t e r i s t i c X - r a y i n t e n s i t i e s o f e l e m e n t s a t some p o i n t s i n the scale formed on t h e Fe-10Cr-25Mn-3Al a l l o y a f t e r o x i d a t i o n f o r 2 1 . 6 k s e c . a t 1 1 7 3 K.

96

substrate.

Therefore, a layer of alumina or

alloys

were

carried

out a t temperatures

ranging

aluminate spinel such as M n A ^ O ^ may be formed

f r o m 873 Κ t o 1173 Κ i n 1 % 0 2- H e g a s m i x t u r e .

in this region.

The

The increase in oxidation resistance and

oxidation rates

w i t h t h e manganese

significantly content.

increased

The a d d i t i o n o f

the suppression of pore formation in the alloy

aluminum

substrate by the addition of aluminum to Fe-Cr-

improvement

Mn alloys are thought to be due to the reduced

m a i n c o r r o s i o n p r o d u c t w a s MnFe2Û^.

transport rate of cations through the alumina

dation resistance

t o Fe-Cr-Mn a l l o y s

of the alloys

t h o s e w i t h l o w manganese

of the spinel oxide MnFe20^ without the forma-

high,

particularly

1073 K.

sivity of cations.

REFERENCES

The previous investigation^ of the oxidation

a

remarkable

i n the oxidation resistance.

or aluminate spinel, resulting in the formation

tion of wustite layer which has high diffu-

caused

The

The o x i -

especially

c o n t e n t was q u i t e

a t temperatures

less

than

1.

F.W. W i f f e n a n d R . T . S a n t o r o , P r o c . o f T o p i c a l C o n f . o n F e r r i t i c A l l o y s f o r Use i n Nuclear Energy Tech. (1983) 193.

2.

P.R.S. Jackson and G.R. Wallwork, 20(1983) 1 .

3.

D . L . D o u g l a s s , F . Gesmundo Oxid. Met. 25(1986) 235.

4.

D . L . Douglass and F. R i z z o - A s s u n c a o , Met. 29(1988) 2 7 1 .

because wustite phase is destabilized by chro7 mium . It was expected that aluminum has ef-

5.

H. S c h n e i d e r , F o u n d r y 108(1960) 563.

fects similar to those of chromium, and this

6.

P.R.S. Jackson and G.R. Wallwork, 21(1984) 135.

7.

C E . B i r c h e n a l l , O x i d . M e t . A l l o y s (ASM, Metals Park, 1971) 177.

of a low-Cr steel, Fe-9.5Cr-17.8Mn, demonstrated that the steel was oxidized too rapidly for a serious consideration of the alloy as a structural component in fusion reactors.

The

high oxidation rate of this alloy was attributed to extensive wustite formation.

Suppres-

sion of the formation of the wustite scale can be achieved by increasing the Cr content^",

expectation has been confirmed in the present study.

Λ.

CONCLUSIONS Oxidation tests of the Fe-10Cr-(lO-25)Mn-3Al

Trade

Oxid. Met.

a n d C. de A s m u n d i s ,

Oxid.

Journal

Oxid. Met.

High Temperature Corrosion of Advanced Materials and Protective Coatings Y . Saito, B. Önay and T. Maruyama (Editors) © 1992 Elsevier Science Publishers B.V. All rights reserved.

97

STRUCTURE AND OXIDATION BEHAVIOR OF THE SCALE FORMED ON AL-CONTAINING FERRITIC STAINLESS STEEL

Shin-ichi SASAYAMA and Takeshi KAMIYA Technical Research Center, Research and Development Division Nippon Yakin Kogyo Co.,Ltd. 4-2 Kojima-cho Kawasaki-ku, Kawasaki Japan

The oxidation scales formed on Al-containing ferritic stainless steel exposed to oxidizing environment having different oxygen partial pressure at 850-1200°C have been characterized using TEM, SEM, X-ray diffraction and other microstructural techniques. The scales formed above 1000°C are composed of alpha alumina and the surface of these scales is generally smooth. The scales formed at 850-950°C are predominantly composed of theta alumina and delta alumina which are called transition alumina. These forms of alumina grow as whisker morphology which are 3/»m in length, 0.5j»m in width and 0.1/mm in thickness. These whiskers consist of theta alumina crystal and delta alumina single crystal. The oxidation rates of these scales are much faster than those of scales which are composed of alpha alumina. These facts mean the growth mechanism of the transition alumina formed on Al-containig ferritic stainless steel is quite different from that of the alpha alumina in the scales. Once the whiskers formed, these whiskers can be transformed to the alpha alumina crystal structure without changing the morphology by heating above 1000°C. The oxidation rates of the alpha alumina whiskers are slower than that of the transition alumina whiskers.

1. Introduction

effect on oxidation kinetics were investigated.

There have been many investigations of oxide whiskers

grown

researchers

over

have

metal

studied

whiskers on steel

surfaces.

formation

surfaces

of

Several

at 400-850°C

in air 1 2

or oxygen atmospheres containing water vapor. CuO and ZnO whiskers were also observed when the base metals were

oxidized

2. Experimental

at high

tempera-

3 4

Figure 1 schematically shows the experimental procedure

shown

the

of

alumina

al.^observed perature and

alloys

whiskers.

alumina

oxidation

Bornstein

NiAl

used.

The

base

^

of

whiskers

during

Fecralloy

observed

were

P.T.Moseley

alumina

oxidized

high

steel.

with

et

whisker

in Table

1, was

coated

formation of alumina whiskers

substrate.

layers after

has not

The

Al

content

the diffusion

with

Al

and

a as

was

are

of

the

surface

approximately

8

mass%.

Table 1

Chemical composition of the base metal

when

(itX) c Si Mn

explanation

is

Al could diffuse into

vapor

present at high temperature.

However, a satisfactory

which

tem-

Smeggil

NaCl

metal,

Fe-Cr-Al alloy with the chemical composition

heated in vacuum so that

On the other hand, there are a few investigations

procedure

hematite

for

the

yet

been

proposed.

Ρ

S Ni Cr AI Ti

Ν

0.0100.330.290.026 0.001 0.5217.723.0Ï0.17im balance iron Samples were cut into approximately 50 χ 10 χ

In the present study, the formation of alumina whiskers on ferritic stainless steel and

its

0.2mm sizes and

ultrasonically

degreased

acetone. They were then subjected

to a

by in

two-step

98

oxidation

treatment

referred

here

the

"whisker

The surface of one of the samples was covered by

treatment". The conditions of the whisker treat-

a

ment are given in Fig.l and were found necessary

(Fig.2(B)).

to grow whiskers over the whole surface of the

given the "whisker

samples.

long, 0.5fim wide and 0.1pm

smooth

oxide

scale

However,

with

for

nodular

the

sample

features which

was

treatment", whiskers of

3μm

thick were

observed

over the whole surface of the sample.(Fig.2(A)) Samples were characterized using TEM, SEM and X-ray

diffraction.

were

also

under

carried

Weight-gain out

two different

for

measurements

samples

conditions

of

preoxidized the

first

oxidation-step. One sample has been covered

by

whiskers and the other sample has been covered by a smooth oxide scale during

the

tests. TEM

was operated at 200KV using scale pieces removed from

the whiskers

covered

over

surface

of

the

scale. For X-ray diffraction, CuKoC radiation at 35KV and 30mA was used.

Base metal

Whisker morphology

Nodule morphology and smooth oxide scale

°Cx1 min First oxidation-step, Q9C 8 (A) Po2=2xi 0~ a tm Second oxidation-step,910°Cx16hr atm Ρ ο 2 = 0.21

First oxidation-step, 11008Cx10min (B) Po 2=2xi0" atm C Second ox i dat i on-step,900C x16hr Po 2 = 0.21atm

Fig.2

SEM micrographs of oxide scales

Figures 3 and

toing Al ac

C

conditions

of

the

4 show

the effects

first

and

second

Heat treatment WCJhr or2x1(Tatm for Al diffusiP on 8

the

1

900°C,16hr Second oxidation-stepPor0.21atn

first

oxidation-step,

oxidation-step

the

conditions

duration

necessary

to decrease with

to

of

form

increasing

temperature. On the other hand, whiskers were

Whisker treatment

•

1200 \

1100 -

• TEM,SEM analysis •X-ray diffraction analysis •Measurement of weight-gain curves Fig.l

second

whiskers was found

Imιη

First oxidation-stepMxf'ati

the

the

oxidation-

steps on whisker growth. For a fixed of

of

1000

900 -

Ο

800 -

Experimental procedure

0.1

3. Results

Fig.3

#

ο \ G Θ

_

Second oxidation-step 9 0 0°Cx 16hr in air (Po2=Q.2l atm)

\ Θ

ι 10 Time(min)

• • •

•no wh t s k e r g r οw t h G

w h i s k e r growth

Ο L 100

Effect of the time and temperature of the first oxidation-step on

3-1 Morphology of scales

whisker growth

Surface morphologies of two samples oxidized under different conditions are shown

in

Fig.2.

not observed when the second oxidation-step was

99

carried

out

at

temperature

higher

than

about

ed

that the oxidation

in vacuum.

sample with the smooth oxide scale.

• • • •

1100

ρ

1000

CL Ε

CD

I—

• • • · • • •

0

ο

900

• 9

•

First oxidation-step 9 0 0 °C Χ 1 m i η8 in vac. (Ρ ο 2 = 2 Χ10 ~ a t m) • no whisker g r ο w t r

ο whι s k er growth

scale with whiskers

ο

Ο

Ο 0

G

different

0

G

oxide

regions

scale

and

whisker-covered 5

50

10

100

Τ i me(h r)

Fig.4

from

Figure 6 is a cross sectional

G Ο

different

that

sample of

the

3-3 X-ray diffraction analysis

ο

800 -

was

of the

with

1200

whiskers

behavior

1000°C following the initial oxidation at 900°C

as observed corresponding

base

metal

were

view of by

SEM.

to

the Three

whiskers,

visible.

The

(top portion of the scale) was

scraped off the base metal using a knife edge. Both the whiskers and

the base metal with

Effect of the time and temperature of

remaining

were

the second oxidation-step on whisker

diffraction.

oxide

scale

analyzed

by

the

X-ray

growth

3-2 Kinetics

Weight gains of two samples oxidized at 900°C in air are shown in Fig.5. The sample which had whiskers gained weight much more than the sample with a smooth oxide scale without whiskers.

g 0. 20 CD Ε

^ 0.15

·-··· sample without whisker -o- sample with whisker

Microstructure of the cross-section of the sample with whiskers

The X-ray diffraction patterns of the whisk-

5 0.10

ers and

c

the base metal with oxide

scale

are

shown in Fig.7 and Fig.8, respectively. The best

CO ι—ι

Fig.6

0.05

25

50 75 Time(hr)

identification

of

the

alumina,

is

thermodynamically

which

whisker

is

not

alpha

the

most

stable phase, but theta and delta alumina even though some peaks are absent. On the other hand, alpha alumina was observed as the major phase in

Fig.5

Effect of scale morphology on weight

the scale left on the sample. Theta and

gain of preoxidized samples further

alumina were observed

oxidized at 900°C in air

fore, it was concluded that the whiskers mainly consisted

These morphological and kinetic results suggest-

delta

in small amounts. There-

of

theta

and

delta

alumina

and

the

oxide scale mainly consisted of alpha alumina.

100

that

the whisker

was made

of a theta

alumina

single crystal and a delta alumina single crystal _ Λ

Λ .

C? —A J2 2 O3

s

i

- A

!

almost

λ

Α.,

|| ,| ||

I,I

ΛΑ !Λ

and

the interface the delta

between

alumina

coherent. The growth

the

theta

crystals

directions

was

of the

theta alumina and the delta alumina crystals in 9

in,

»Li. I

O2 3

a-ki

and that

alumina

2

e

t

h

whisker were [Oil] and [111], respectively.

n,i

2Ο3 a:-Fe

Fig.7

X-ray diffraction pattern of the whisker-containing top portion of the oxide scale scraped off the sample

0. 5K

0.25

r"

10

30

0 - A £ 2O 3

•

ll.Ml

1

03

S - A £

2

a-ki

2Ο3

50

IL

I

11

Ι

Ι.

I

111

90 1

1 M

Θ

Fig.9

L

1 I , il

ι, ,

I

I

ör-Fe Fig.8

l.I

70

Lattice image of a whisker-shaped crystal and the corresponding

L

diffraction pattern

I

X-ray diffraction pattern of the oxide scale left the surface and

4. Discussion

the base metal 4-1 Growth mechanism of the whiskers 3-4 TEM observation

Three possible line defect models have been proposed

Figure 9 shows a high resolution TEM image of

to explain the growth mechanism of 1 2 3 . These are the hollow

hematite whiskers

a fragment of a whisker. The diffraction pattern

tunnel model, internal grain boundary model and

showed

screw

(022) reflection

(110) reflection indicated

that

from

from theta alumina and delta

the theta

alumina

alumina

and

and

it

delta

dislocation

model.

However,

neither

of

them can be used to explain the growth mechanism of alumina whiskers.

alumina had the same c axis. The streak of the diffraction whisker

pattern

also

indicated

had faults. Therefore,

that

the

it is presumed

Alumina consisted

whiskers of

theta

observed alumina

in this

and

delta

study alumina

101

phases. The theta alumina structure formed

from

In fact, it was observed that whiskers trans-

the delta alumina structure and these had almost

formed

same lattice spacings except along

changing

It

is, however,

presumed

that

the c axis.

the

distance

between atoms of the theta alumina and the delta

above

into

alpha

alumina

the morphology

1000°C

The

by

structure heating

transformation

without

the

sample

into

alpha

alumina was non-reversible.

alumina at the interface of them will be larger than

the

distance

between

atoms

of

the

theta

alumina or between those of the delta alumina.

Furthermore, weight was

heat

treated

at

gain of a sample which 1200°C

for

Ihr

after

the

"whisker treatment" was studied. The sample had P.T.Moseley

et a l . ~* showed

that

alumina

whiskers formed on Fecralloy steel composed of a highly

defective

transition

alumina

phase

(probably theta alumina). Faults in the crystal

whiskers

with

gain

of

almost

this

smooth alumina

the

suppression

results except

crystal

were

structure.

similar

that Moseley

to

These

those

of

authors

our

et a l . pointed

study

out

that

alumina

structure.

whisker-containing

the same as that of

can occur during ordering of the cation sites in alumina

alpha

Fig.10

shows the result of this test. Since the weight

gain of

was with

confirmed

that

s c a l e , it was

of weight

sample

the sample

the whisker-con-

taining samples during oxidation could be accomplished if the whiskers contain alpha alumina.

gamma alumina phase ordered in theta alumina.

Growth mechanism of the alumina whiskers can, therefore, be explained Al

ion

diffuses

by the fact that

through

the

interface

since

between

the theta alumina crystal and the delta alumina crystal which

exists

parallel

to

the

growth

J CD

~

-•-no whisker (A-AI0j) -ο-with w hisker (fl + i Al0 ) 0. 20 -Δ-with w hisker ( i + J-*ff transformation)^ 0.15 2

23

0.10

direction of the whisker, the whiskers grow one direction and Al ion also diffuse rapidly at the interface because of the larger distance between atoms at the interface.

4-2

Suppression

03 CD

0. 05

CD

of weight gain

of

whisker-con-

0

taining samples

The reason why samples with whiskers and delta alumina) had

(theta

weight gains much higher

Fig.10

25

50 75 Time(hr)

100

Effect of crystal structure on weight gain of the samples which

than samples with only smooth oxide scales with

show whisker morphology but

nodular

different crystal structure during

features

(alpha

alumina)

was

probably

the faster diffusion of Al ion along the defective interface in the whisker. This

further oxidation at 900°C in air

hypothesis

also implies that suppression of weight gain of the whisker-containing can be accomplished

samples during if

the w h i s k e r s

oxidation contain

5. Conclusion

alpha alumina. Alumina whiskers were grown on the surface of

102

Al-containing

ferritic stainless steel by using

a 2-step oxidation treatment called the "whisker treatment". The aim of our study was to characterize the alumina whisker and understand of its growth mechanism. The following conclusions were made.

1)

Whiskers

stainless

formed

steel

on

Al-containing

consist

of

theta

ferritic

and

delta

alumina. The interface between the theta alumina and the delta alumina crystals is almost coherent .

2) Weight

gains

consisting

theta

higher

than those

of

the sample

and of

delta the

with

alumina

whiskers were

sample with

a

much

smooth

oxide scale of alpha alumina. However, suppression

of

the

weight

gain

of

the

containing sample can be accomplished

whiskerby trans-

forming the whisker structure into alpha alumina without

changing

the

whisker

morphology

by

heating the sample above 1000°C.

References (l)D.A.Voss, E.P.Butler and T.E.Mitchell Trans. TMS-AIME vol.l3A (1982) 929 (2)R.L.Tallman and E.A.Gulbransen chem. Soc. vol.115 (1968) 770

:

: J. Electro-

(3)R.Takagi: J. Phys. Japan vol.12(1957)1212 (4)J.A.Sartell, T . L . J o h n s o n and vol.215(1959)420

R.L.Stokes, S.H.Bendel, C.H.Li : T r a n s . TMS-AIME

(5)P.T,Moseley, K.R.Hyde, B.A.Bellamy and G.Tappin : Corrosion Science vol.24(1984)547 (o)J.G.Smeggil and E.A.Bornstein chem. Soc. vol.125 (1978)1283

: J.

Electro-

High Temperature Corrosion of Advanced Materials and Protective Coatings Y . Saito, B. Önay and T. Maruyama (Editors) © 1992 Elsevier Science Publishers B.V. All rights reserved.

103

CYCLIC OXIDATION BEHAVIOUR O F MICROCRYSTALLIZED CoCrAl ALLOY FILM Fuhui WANG, Hanyi LOU and Weitao WU Corrosion Science Laboratory, Institute of Corrosion and Protection of Metals, Academia Sinica, Shenyang 110015, China The cyclic oxidation resistance of normal-grained cast Co-30Cr-5Al alloy and its microcrystallized film was investigated at 1000 °C in air. The results indicated that normal-grained CoCrAl alloy exhibited very poor resistance against cyclic oxidation. After only 2 cycles, the initially formed A l 2 0 3 scale severely spalled off, and then C r 2 0 3 and C o C r 20 4 spinel were formed as a surface layer, while A 1 2 0 3 only existed in the form of internal oxides. Acoustic emission measurements showed that the A 1 2 0 3 scale formed on the normal-grained Co-30Cr-5Al alloy would crack and spall off only during cooling, and cracking and spalling at about 400°C. On the contrary, the microcrystallized film showed very excellent cyclic oxidation resistance. No AE single was detected during cooling from 1000°C even down to room temperature, which was consistent with the surface morphologies observed by SEM. This denoted that the microcrystallization can significantly improve cyclic oxidation resistance of Co-30Cr-5Al alloy.

2.EXPERIMENTAL

1.INTRODUCTION The improvement

of oxidation resis-

Sputtering technique is used to produce

tance of CoCrAl-type alloys by small

microcrystallized

amounts of reactive elements such as

terial, substrate alloy(IN738) and sput-

films. The target ma-

yttrium, hafnium or cerium has been con-

tering process were the same as those

firmed by many investigators. One of the

given in reference2. The thickness of

beneficial effect of these reactive

microcrystalline layer is about 25 pm.

elements on the oxidation

resistance

The grain size is smaller than 0.5 pm,

is known to enhance the oxide adhesion.

which is about 2 orders of magnitude

Various mechanisms has been proposed to

smaller than that of normal-grained cast

explain the beneficial effect^, in which

alloy. Both cyclic and isothermal oxida-

a very important mode is mechanical

tion tests were carried out at 1000°C

keying due to the formation of oxide

in air. For cyclic oxidation test, the

pegs rich in reactive elements which

samples were kept at designed

anchor the oxide scale to the substrate.

for 1 h and then rapidly pulled out of

However, recent studies on sputtered Co-

the furnace, and cooled down to room

CrAl coatings without reactive

elements

temperature

temperature for 10 min in a cycle. The

indicated that reduced grain sizes, like

mass changes of samples for cyclic oxi-

the addition of RE to the alloys, can

dation do not include the part of the

greatly improve the oxidation resistance .

spalled scales. For isothermal oxidation

This paper discribes the effect of m i -

test,the samples were kept at 1000°C for

crocrystallization

on the cyclic oxida-

tion resistance of Co-30Cr-5Al

alloy.

10,

20, 40, 60 and 100 h respectively

and cooled to room temperature for mass measurement.

104

Specimens after oxidation

exposure

microcrystallization

is in improving of

were examined by various techniques in-

scale/substrate adhesion, and this is

cluding an acoustic emmision

more critical during thermal cycling

optical metallography, microscopy, and X-ray

technique,

scanning

electron

diffraction.

test than isothermal exposure. Fig.lc shows the isothermal kinetics of normalgrained Co-30Cr-5Al alloy oxidized at

3.RESULTS

1000°C in air. The mass change is much

3.1. Cyclic oxidation

kinetics

lower than that for cyclic oxidation

Fig.l shows the cyclic oxidation ki-

and is nearly the same as the mass change

netics at 1000°C in air for both normal-

of microcrystallized

grained Co-30Cr-5Al alloy and its micro-

that the microcrystallization

greatly

crystallized

improves the cyclic oxidation

resistance,

films. Since the scale

formed on normal-grained

alloy

cracked

and partially spalled off, the recorded

film. This indicates

but shows no obvious effect on the isothermal oxidation.

mass change was only apparent. So the

3.2. Oxide morphology

total mass gain for normal-grained alloy

A typical cross section of the nor-

is larger than that given in Fig.la.

mal-grained Co-30Cr-5Al alloy after 100 cycles of oxidation at 1000°C in air is

1.5

illustrated

in Fig.2a. There is severe

internal oxidation. The outer part of the internal oxidation zone consists of A 1 20 3,

and besides A l 2 0 3 ,

the inner part

contains A1N. The degradation about 50 μτα,

layers is

where ß-CoAl phase was de-

pleted. However, as compared with normal-grained CoCrAl alloy, no

internal

oxidation was found in microcrystalline CoCrAl film 50 Cycling Times

100

(Fig.2c). Fig.2b shows the

cross section of normal-grained

CoCrAl

alloy after isothermal oxidation

(h)

for

100 h at 1000°C in air. Internal oxidaFIGURE 1 Mass change of samples oxidized at 1000 °C in air: (a) and (b) for normal-grained alloy and its microcrystallized films respectively during cyclic oxidation, and (c) for normal-grained alloy during isothermal oxidation.

tion was not found either. This result is consistent with the mass change

Fig.3 shows a comparison of the surface morphologies of oxide scales formed on the normal-grained

For CoCrAl microcrystalline film,

the

(Fig.

lc) .

and microcrystal-

line CoCrAl alloys. After 2 cycles, the

total mass gain is the same as that

scale formed on the normal-grained

show in Fig.lb, because no scale spalling

tended to spalling from the substrate

was

observed.

As indicated earlier, the main effect of

(Fig.3a). The remained alumina

alloy

scale

was very dense. After the spalling

105

FIGURE 2 Cross section of a l l o y s a f t e r 100 h o x i d a t i o n at 1000°C i n a i r : (a) and (b) of n o r m a l - g r a i n e d a l l o y f o r c y c l i c p x i d a t i o n and i s o t h e r m a l o x i d a t i o n , r e s p e c t i v e l y ; (c) f o r m i c r o c r y s t a l l i z e d f i l m f o r c y c l i c o x i d a t i o n .

FIGURE 3 Surface c y c l e s ,

of

the

i n i t i a l l y

formed a

m o r p h o l o g i e s : ( a ) a n d ( b ) of n o r m a l - g r a i n e d a l l o y a f t e r 2 and 100 r e s p e c t i v e l y ; and (c) of m i c r o c r y s t a l l i z e d f i l m a f t e r 100 c y c l e s .

as

a

surface

i n t e r n a l

l a y e r

i n t e r d i f f u s i o n system^.

The

v i t i e s

were

C o C r 20 l i z e d cycles scale

a l l o y

are 4

mainly

s p i n e l .

f i l m ,

the in

cycles

on

a

l o t

the

CoCrAl

of

scale

formed

sound

( F i g . 3 c ) .

observed.

X-ray

scale

a f t e r

s p a l l i n g

d i f f r a c t i o n

ca-

l i n e

CoCrAl

3.3.

scales was

r e c e n t l y used

to

d u r i n g

high

the

of

the

m a l - g r a i n e d

oxide

t h i s

was

80

was

was

0.68

w i t h

3

. measurement

emission

o x i d e s c a l e

c r a c k i n g

scales a l l o y

formed

and

been

o x i d a t i o n 4 .

technique

the

has

c r a c k i n g

was

In

used

behaviour on

both

nor-

m i c r o c r y s t a l l i n e

f i l m s .

N e i t h e r

represented

0

m i c r o c r y s t a l -

emission

t h i s

and

CoCrAl

2

on

temperature

study

i n v e s t i g a t e

In

A l

acounstic

d e t e c t

present

100

formed

Acounstic

t o

m i c r o c r y s t a l -

the

nor

that

C r 2Ü 3

remained c r a c k i n g

of

on

s u r f a c e .

showed

the

CoCrAl

100

composed

For

the

t h a t

alloy

scale

and

r e s u l t

than

low

oxide

a f t e r

d i s t r i b u t e d

alumina

r a t h e r

of

of

( f i g . 3 b )

d i f f r a c t i o n

oxides

oxide

c o e f f i c i e n t

normal-grained coarse

s c a l e ,

because

s t r u c t u r e

became

X-ray

formed

dB

the

t e s t , and

V.

a

The

AE

t o t a l

f i x e d

samples

f u r n a c e .

The

system

t h r e s h o l d were

c o o l i n g

gain v o l t a g e

c o o l i n g r a t e

i s

106

about 2 orders of magnitude lower than

may correspond to the appearence of the

that for cyclic oxidation. Fig.4 illu-

mass crackings and the beginning of the

strates the AE counts and

scale spallation. However, for microcry-

temperature

vs cooling time curve for normal-grained

stalline films, no AE signal was detec-

CoCrAl alloy. The first signal appeared

ted during cooling, even down to room

at about 600°C, which may correspond

temperature, which was cosistent with

to

the initiation of cracking at the edge

the SEM observation of the surface scale

of the specimen. After then a large num-

morpholoy. This denoted that no crack

ber of signals occurred at 400°C, which

was generated during

250

cooling.

1000

800

600

u -μ

400

to U

% 200

20

40

Cooling Time

60

80

(min)

FIGURE 4 AE counts detected during furnace cooling of normal-grained CoCrAl alloy after 100 h oxidation at 1000°C in air

of reactive elements, the oxide scale

4. DISCUSSION High temperature alloys exhibit their oxidation resistance by forming a thin

formed on the surface showed in general very poor adhesion. It was sugested

oxide scale on the surface. This oxide

that the poor scale adhesion was caused

scale should remain adherent to the al-

by several main factors: (1) large com-

loy during both isothermal and

thermal

pressive growth stress developed by the

cycling exposures, which is considered

reaction between the inward diffusion of

to be prerequisite to a protective

anion

scale. In MCrAl system without

oxide

additions

(oxygen) in the grain

boundaries

and outward diffusion of cation

(alumi-

107

nium),

r e s u l t i n g

oxides

w i t h i n

produces layer^;

the

ide

i n t e r f a c e

growth

a

bonding oxide

the

a

stress

of

chromina

ing

C r 20 3 ,

c i e n t s

of

and

are

the

higher from

in

the

crack

metal

several C r 20 3

i t

change

be

the

o x i d a t i o n

c l e a r l y

throughout

time.

X-ray

the

composed the

oxides h

d u r i n g I t

on

of

A 1 20 3. and

the

This

only

mass

showed af-

mainly that

is

d i f f i c u l t

to

calculate On

the

the other

i s

ΔΤ

sample^.

by

the

the

scale

oxide

f o l -

on

l i n e

growth hand.

on

f i l m s

explain

hesion

i s

(1)

the

the

of

improvement

the

Since

mechanism

of

oxide

c l e a r .

f o l l o w i n g the

of

n e i t h e r

m i c r o c r y s t a l -

the

f u l l y

to meas-

ad-

The explaina-

improvement

i n

adhesion. For

g r a i n

s i t e s

8.

before

appeared.

begun,

the

2

observation

j u s t

f o r

4

y i e l d

kg/cm

t h a t

has

not

-4.1xl0

adherent

SEM

s p a l l i n g

be

stress

emission

out

o x i d a t i o n

account

oxide

or

would

m i c r o c r y s t a l l i n e

very

p o i n t e d

proposed

to

on

.

cooling

reached.

A c o u s t i c

t i o n

on

i s

f r a c t u r e

is

2

k g / c m

thermal

(1)

be

order

6

4

formed

authors

the

A 1 20 3

- 2 x l 0

could

s t i l l

of

3 . 8 x l 0

The

Pos-

modulus.

temperature

is

c o o l i n g

nor

study

i n

is

1

s

compressive

a l l o y

scales

i s

equation

might

s u b s t r a t e .

urement

Ε

-0.75%.

the

stress

) ,

1

Youg

t y p i c a l l y

A l 20 3

oxide

The

the

t h e o r e t i c a l l y .

of

Co-30Cr-5Al

occurred

c o o l i n g .

stresses

the

is

of

room

of

w h i l e

thermal

to

whole

to

from

2

k g / c m ,

i s

Ε

generated

c a l c u l a t e d

the

a l l o y

between

the

the

given

( = Δ α · ΔΤ

and

7

order

c r a c k i n g

is

d i f f e r e n c e

and

on

the is

(1)

value

1000°C

the

a l l o y ( F i g . the

of

s t r a i n

k i n e t i c s

i n d i c a t e s

s p a l l i n g

Cr

However,

d i f f r a c t i o n

o x i d a t i o n

The

So

Co,

that

Co-30Cr-5Al

isothermal

cracking

as

i s

s t r a i n

the

time, form-

is

value 0.3,

during

magnitude

Co-30Cr-5Al

of

s t r e n g t h

coef-

small

100

ox-

higher

is

t e r

the

e x h i b i t

AI2O3.

r a t i o s ,

growth

d i f f u s i o n

of

e

s i o n ' s

from

of

is

Δα

change

I f much

f i l m

c o e f f i c i e n t

-

Here,

of

case

such

seen

a l l o y

off

is

s u b s t r a t e , oxide

metal,

s t r e s s

σ = ε·

The

co-

alumina

ions

i n

a l -

was

the the

data.

f i l m

Ε

to

c e r t a i n

In

the

c a l c u l a t e d

r e l a t i o n :

of

scale

s p a l l

o x i d a t i o n that

and

orders

isothermal

can

and

would

than

due

i n

the

lowing

and

weak

the

a f t e r

a l l o y

normal-grained

l c ) ,

thermal

The

of in

, where

· ΔΤ

be

oxide

expansion

temperature

s t r e s s

form.

because

the

than

stresses,

would

change,

r o l e

by Δα

thermal

the

expansion

oxide

that

T>e

Co-30Cr-5Al

growth

o x i d a t i o n

and

A l

the

would

i s o t h e r m a l

the

given

than

can

c o e f f i c i e n t

the

generated

m e t a l / o x -

s p a l l i n g

c o o l i n g

thermal

s p a l l i n g

I f

source

on

important

and

do.

scale

in

between

more

cracking

of

stress

s t r a i n

of

associated

formation

and

and

thickness

smaller

in

stresses

expansion

oxide

mechanism^.

normal-grained

e f f i c i e n t s plays

mass

non-

the

the

the

s t r e s s

c r a c k i n g

d i f f e r e n c e

ide

the

at

i n

large

o x i d a t i o n

large

causes

thermal

main

voids

oxide occurs

scales.

In

the

of

by

thermal

the

between

(3)

the from

d i f f e r e n c e

and

caused

new

which

e n t i r e which

c o e f f i c i e n t s

d u r i n g

of

of

the

K i r k e n d a l l - t y p e

combination

loy,

to

of

oxide,

s t r e s s

a l l o y s ;

areas

and

w i t h

due

and

f o r m a t i o n

growth

expansion

oxides

contacted

the

e x i s t i n g

thermal

c o o l i n g

thermal the

l a t e r a l (2)

d u r i n g

i n

the

of

m i c r o c r y s t a l l i n e boundaries

oxides

f u r t h e r

are

d u r i n g

o x i d a t i o n

the

CoCrAl

o x i d a t i o n , along

films,

n u c l e a t i o n

g r a i n

and boun-

108

daries into coatings forms micro-pegs

scale formed on microcrystalline film

which anchor the scales to the coatings.

would be nearly the same as in that formed

This beneficial effect of microcrystal-

on normal-grained

alloy.

lization is similar to that of reactive 5.

elements. (2) The oxide scales formed on micro-

CONCLUSION Both isothermal and cyclic oxidation

crystalline films have finer structure

tests were carried out in air at 1000°C

than those formed on normal-grained al-

on normal-grained Co-30Cr-5Al alloy and

loys. The deformation rate of diffusional

its microcrystallized film. The follow-

creep of a polycrystalline material (ε)

ing conclusions were reached.

is

given by reference^ ΒJ D σΩ 2

where

1

d

d kT ο is the tensile stress,

The normal-grained CoCrAl alloy exhibited very good isothermal oxidation re-

(2)

sistance, but very bad cyclic oxidation

Ω the

resistance. Just only 2 cycles, the alu-

atomic volume, d the average grain size,

mina scale formed on normal-grained al-

Bi and B 2 are numerical constants, D va n d

loy cracked and spalled off. During fur-

D b are the volume and grain boundary dif-

ther oxidation the chromina formed as a

fusivities, kT has the usual meaning and

surface layer, while alumina existed in

δ is the thickness of the boundaries.

the form of internal oxides. Since the

At low temperature where boundary diffu-

protectiveness of chromina is not so

sion predominates, one obtains

good as that of alumina, the normal-

e=B2aClôOh/d

kT

According to equation

grained CoCrAl alloy exhibited high mass (3)

change during cyclic oxidation.

( 3 ) , the diffusional

On the contrast, the microcrystal-

creep rate of a polycrystal may be en-

lized CoCrAl film possesed not only ex-

hanced by reducing the crystal size,d,

cellent isothermal but also cyclic oxi-

and by increasing the boundary diffusi-

dation resistance. This denoted that the

vity, D b · So the plastic deformation of

microcrystallization can significantly

fine-grained oxides formed on microcrys-

improve cyclic oxidation resistance.

talline films would be much easier than that of coarse-grained ones formed

on

normal-grained alloys, because the grain

ACKNOWLEDGMENTS This work was supported by the Na-

boundary sliding was the major deforma-

tional Natural Science Foundation of

tion mechanism in the s c a l e ^ . On the

China. The authors gratefully

other hand, the fine-grained oxide scale

ledge M r . Li Meishuan for conducting

may block the outward diffusion of metal

the acoustic emission tests.

ions and prevent void

acknow-

formation.

The above mechanisms were proposed from the view point of improvement of bonding at the metal/scale interface and of releasing the stress during cooling, because the thermal stress in alumina

REFERENCES 1. D.P. Whittle and J. Stringer, Trans. Royal S o c , London, Ser.A, 309

295(1980)

2. F.Wang and H.Lou, Mater. Sei. Eng., 129A (1990) 279.

109

F.H. Stott, G.C.Wood, and M.G. Hobby, Oxid. Met., 3 (1971) 103.

7. K.L. Luthra and C.L. Briant, Oxid. Met., 26 (1986) 396.

T. Li and M. Li,Mater. Sei. Eng., 120A (1989) 235; 239.

8. Y.I. Kuraki, Ceramic Handbook, translated by D.liu and Sh.Chen, Light Industry Press, Beijing, (1984) 584 (in C h i n e s e ) .

F.A. Golightly, F.H. Stott and G.C. Wood, Oxid. Met., 10 (1976) 163 V. Provenzano, K. Sadananda, N.P. Louat and J.R. Reed, Surf. Coat. Technol., 36 (1988) 61

9. J.Karen, R.Birringer and H.Gleiter, Nature, 330(10) (1987) 556. 10.J.K.Tien and F.S.Pettit, Metall. Trans., 3 (1972) 1587.

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High Temperature Corrosion of Advanced Materials and Protective Coatings Y . Saito, B. Önay and T. Maruyama (Editors) © 1992 Elsevier Science Publishers B.V. All rights reserved.

Ill

HIGH TEMPERATURE OXIDATION OF Ni-Cr ALLOYS

Tadaaki AMANO and Osamu MOMOSE* Shonan Institute of Technology, Fujisawa 251, Japan *School of Engineering, Tokai University, Hiratsuka 259-12, Japan

The oxidation behavior of Ni-20, -40, -60 and -80wt%Cr alloys was studied at 1373 and 1473K for 14.4ks in O2 by thermogravimetric analysis, X-ray diffraction and scanning electron microscopy. The weight gain of the alloys increased with increasing time of oxidation. Increasing Cr content tended to increase the weight gain of the alloys. The spalling of surface oxide was observed at the oxide/alloy interface for all the alloys. The surface oxide for Ni-20Cr, Νi-40Cr and Ni-60Cr alloys spalled partially. On the other hand, the spalling of the surface oxide for Ni-80Cr alloy occurred over the entire surface. Oxide phases, morphology of surface oxide and oxidation kinetics will be presented.

1. INTRODUCTION

copper hearth in an argon atmosphere. The

The oxidation of Ni~Cr alloys has been

compositions of alloys are Ni~20Cr, Ni-40Cr,

studied in detail because of its importance as

Ni~60Cr and Ni-80Cr.

a heat-resistant alloy.

dimensions of about 10mm χ 10mm χ 1mm. All

It has been found that

The buttons were cut to

when sufficient Cr is present in the alloy

specimens had a 800 grit SiC surface finish, and

(i.e., near 20wt%) a protective scale of Cr2Û3

then were ultrasonically rinsed with ethyl

is formed which results in low oxidation

alcohol.

rates.

1 7-

Mechanisms which describe the growth

The oxidation process was carried out

by a thermogravimetric method at 1373 and 1473K

and spalling of oxide scales on Ni-Cr alloys are

for 14.4ks in O2.

complex and the effects produced by various

were identified by X-ray diffraction. The

contents of Cr in the alloys are not clearly

morphology of oxides on the alloys was measured

understood.

by scanning electron microscopy.

In the present study, high

The oxides on the surface

temperature oxidation behavior of Ni-(20, 40, 60, 80wt%)Cr alloys is studied after oxidation at 1373 and 1473K for 14.4ks in 0 2,

in order to

obtain a better understanding of the oxidation

3. RESULTS AND DISCUSSION 3.1. Thermogravimetric analysis Figures 1 and 2 show the weight gain of the

rates and morphology of the surface scales on

alloys oxidized at 1373 and 1473K, respectively.

the alloys.

The weight gain of the alloys increased in the order of N i - 2 0 C r < N i - 4 0 C r < N i - 6 0 C r < N i - 8 0 C r

2. EXPERIMENTAL

at both temperatures of oxidation.

Nickel (99.9%) and chromium(99.99%) were used to produce the alloys.

Buttons weighing 200g

were prepared by arc-melting on a water-cooled

Hodgkiess

8>

has shown that increasing the chromium content of Ni-Cr alloys above 20wt% increases the oxidation rate.

Therefore, our data are in good

112

Fig. 3 Parabolic plots of the isothermal oxidation at 1373K for 14.4ks in O2.

Fig. 4 Parabolic plots of the isothermal oxidation at 1473K for 14.4ks in O2.

Fig. 2 Isothermal oxidation of Ni-(20, 40, 60 80)Cr alloys at 1473K for 14.4ks in 0 2. agreement with previous results.

Figures 3 and

4 show parabolic plots of the isothermal oxidation at 1373 and 1473K for 14.4ks in 0 2, respectively.

Data for all the alloys followed

approximately a parabolic law.

The result

suggested that the oxidation kinetics of the alloys were controlled by ionic diffusion in the surface oxide. 3.2. X-ray diffraction Table 1 shows the phases identified by X-ray

113

Table 1 Summary of X-ray diffraction analysis of oxides formed on Ni-(20, 40, 60, 80)Cr alloys oxidized at 1373 and 1473K for 14.4ks in 0 2. Alloy Ni-20Cr Ni-40Cr Ni-60Cr Ni-80Cr Ni-20Cr Ni-40Cr Ni-60Cr Ni-80Cr W'weak,

Temperature/K

Oxides C r 20 3 (w) C r 20 3 (w) 1373K C r 20 3 (m) Cr 2Û3(m) C r 20 3( s ) N i 0 ( m ) Cr 203(m) 1473K C r 20 3 (m) C r 20 3 (m) mrmedium, s : strong.

diffraction of the surface oxides on the alloys. The surface oxide C r 20 3 for all the alloys was detected at both temperatures of oxidation. The formation of NiO was also observed after oxidation at 1473K of a Ni-20Cr alloy.

After

oxidation at 1373K the peak strength by X-ray diffraction increased with increasing weight gain of the alloys.

This fact suggested that

the spalled amount of surface oxide on the alloys was small and was independent of chromium content of the alloys.

On the other hand, after

oxidation at 1473K the peak strength of C r 20 3 on the alloys was independent of the weight gain. The peak strength of C r 20 3 on Ni-20Cr was strong

Fig. 5 Scanning electron micrographs of surface oxide of Ni-(20, 40, 60, 80)Cr alloys oxidized at 1373 and 1473K for 14.4ks in 0 2 . (a) Ni-20Cr. (b) Ni-40Cr. (c) Ni~60Cr. (d) Ni-80Cr.

compared with the other alloys which showed a similar peak strength.

This result showed

that the spalled amount of surface oxide

recognized along the center of swollen surface oxide of these alloys.

The fine granular oxide

on Ni~20Cr alloy was smaller than that of

formed on the alloys was l ~ 1 0 / / m

surface oxide on Ni-40Cr, Ni-60Cr and Ni-80Cr

was nearly the same after oxidation at both 1373

al loys.

and 1473K.

3.3. Scanning electron microscopy Figure 5 shows the surface oxide on Νi-(20, 40, 60, 80)Cr alloys formed at 1373 and 1473K after 14.4ks in 0 2. After oxidation at 1373K, the surface oxide on Ni-20Cr alloy was almost planar.

On the other hand, convoluted-

morphologies were observed at surfaces on N i 40Cr, Ni-60Cr and Ni-80Cr alloys and cracks were

in size, and

Figure 6 shows regions spalled of

surface oxide of Ni-(20, 40, 60, 80)Cr alloys oxidized at 1373 and 1473K for 14.4ks in 0 2. Surface oxide of all the alloys spalled partially after oxidation at 1373 and 1473K. After oxidation at 1373K, voids were observed along the grain boundaries.

The number of voids

increased with increasing Cr content of the alloys.

The voids were l ~ 5 / z m in size for all

114

grain size on alloys oxidized at 1373K was ~ l / z m and that of oxide formed at 1473K was about 3/zm.

4. CONCLUSIONS 1. The oxidation rate of Ni-Cr alloys containing more than 20wt%Cr increased with increasing Cr content of the alloys. 2. The main surface oxide was Crs03 Î NiO was detected for Ni-20Cr alloy after oxidatin at 1473K for 14.4ks. 3. Spalling of surface oxide was observed for all the alloys.

The amounts of spalled oxide

tended to increase with increasing Cr content of the alloy. 4. Spalling of surface oxide was attributed to the formation of voids which formed at grain boundaries and the oxide/alloy

interface.

ACKNOWLEDGEMENT The authors would like to thank Prof. K. Okazaki at Shonan Institute of Technology for Fig. 6 Scanning electron micrographs of regions of surface oxide of Ni-(20, 40, 60, 80)Cr alloys oxidized at 1373 and 1473K for 14.4ks in 0 2. (a) Ni-20Cr. (b) Ni-40Cr. (c) Ni-60Cr. (d) Ni-80Cr.

his continued interest and encouragement, and H. Homma and Y. Hirano, Tokai University for their assistance in the experimental work. REFERENCES

the alloys.

Vacancies accumulating at grain

I.

C. S. Giggins and F. S. Pettit, Trans. Met. Soc. AIME 245 (1969) 2495.

2.

C. S. Giggins and F. S. Pettit, Trans. Met. Soc. AIME 245 (1969) 2509.

3.

Β. Chattopadhyay and G. C. Wood, J. Electrochem. Soc. 117 (1970) 1163.

4.

D. L. Douglass and J. S. Armijo, Oxid. Met. 2 (1970) 207.

However,

5.

C. E. Lowell, Oxid. Met. 7 (1973) 95.

the number of voids for Ni~80Cr alloys decreased

6.

G. M. Ecer and G. H. Meier, Oxid. Met. 13(1979)159.

7.

T. Amano, S. Yajima, T. Kimura and Y. Saito, Corros. Eng. (Boshoku Gijutsu) 24 (1975)19.

8.

T. Hodgkiess, Ph.D. Thesis, University of Manchester (1967).

boundaries can coalesce to nucleate voids at the 6

alloy/oxide interface . The voids observed at grain boundaries and beneath the rapidly growing Cr2Û3 surface oxide

resulted in the

separation of the oxide from the alloy over large areas.

After oxidation at 1473K, voids

were also observed for all the alloys.

compared with the alloy oxidized at 1373K.

This

result is attributed to a phase change in Ni~Cr accompanied by oxidation for 14.4ks at 1473K. From Fig. 6, it is seen that the

oxide

6

High Temperature Corrosion of Advanced Materials and Protective Coatings Y . Saito, B. Önay and T. Maruyama (Editors) © 1992 Elsevier Science Publishers B.V. All rights reserved.

115

EFFECTS OF OXYGEN AND WATER VAPOR PRESSURES ON OXIDATION OF IRON-CHROMIUM ALLOYS AT 573 Κ

T o s h i h i d e TSUJI, S h i n j i KOBAYASHI, Masashi ODA and K e i j i Department o f Nuclear E n g i n e e r i n g , Chikusa-ku, Nagoya 464-01, Japan

NAITO

F a c u l t y o f E n g i n e e r i n g , Nagoya U n i v e r s i t y , Furo-cho,

The o x i d a t i o n o f iron-chromium a l l o y s ( 0 , 1 and 5 wt%Cr) was i n v e s t i g a t e d by means o f Rutherford b a c k s c a t t e r i n g s p e c t r o s c o p y a t 573 Κ i n t h e oxygen p a r t i a l pressure range from 10 t o 10 Pa w i t h o u t and w i t h water vapor a t 10 Pa o f water vapor p r e s s u r e f o r 1 - 48 h. The o x i d a t i o n o f iron-chromium a l l o y s ( 0 - 5 wt%Cr) obeyed t h e p a r a b o l i c r a t e law a f t e r the r a p i d o x i d a t i o n i n t h e f i r s t s t a g e . When iron-chromium a l l o y s were o x i d i z e d under v a r i o u s oxygen p a r t i a l p r e s s u r e s w i t h o u t w a t e r vapor, the p a r a b o l i c r a t e constant decreased w i t h i n c r e a s i n g chromium c o n c e n t r a t i o n s a t constant oxygen p a r t i a l pressure probably due t o t h e d e c r e a s e o f i r o n d i f f u s i o n through chromium o x i d e , and t h e mechanism o f o x i d a t i o n was d e v i d e d i n t o t h r e e r e g i o n s : Below 10 Pa t h e d e c r e a s e o f the o x i d a t i o n r a t e w i t h i n c r e a s i n g oxygen p a r t i a l pressure i s due t o the formation o f dense h e m a t i t e . Between 10 and 10 Pa the p r o t e c t i v e h e m a t i t e c o v e r a g e o y e r m a g n e t i t e kept the p a r a b o l i c r a t e c o n s t a n t s low r e g a r d l e s s o f oxygen p a r t i a l p r e s s u r e . At 10 Pa an i n c r e a s e o f the p a r a b o l i c r a t e constant was observed probably due t o t h e simultaneous growth o f hematite and m a g n e t i t e . The remarkable i n c r e a s e o f the p a r a b o l i c r a t e constant f o r the o x i d a t i o n o f iron-chromium a l l o y s by water vapor was observed a t lower oxygen p a r t i a l p r e s s u r e s . T h i s f a c t may be e x p l a i n e d by vapor-phase d i f f u s i o n o f v a l a t i l e F e ( 0 H ) 2 from the m e t a l - o x i d e i n t e r f a c e t o t h e o u t e r l a y e r c r y s t a l s . The c o r r o s i o n b e h a v i o r o f iron-cnromium a l l o y s in gas phase c o n t a i n i n g w a t e r vapor may be c l o s e l y r e l a t e d t o t h a t i n h i g h temperature and h i g h - p r e s s u r e w a t e r around 573 K.

1. INTRODUCTION The radioactivity buildup around the primary coolant system of nuclear water reactor is one of the important problems from the point of view of the personal exposure 1 • In order to reduce radioactivity buildup in boiling water reactor (BWR), the injection of some amounts of oxygen (20 - 60 ppb) into the coolant water has been carried out in order to protect the corrosion2 . Videm3 also showed that the addition of small amounts of oxygen gas into the coolant water reduced the corrosion of carbon steel in the simulated BWR environment effectively due to the formation of protective oxide film. The details of the behavior of the corrosion, however, have not been understood quantitatively yet. Water as a coolant shows such low electrical conductivity that may make difficult the occurrence of wet corrosion based on the local cell. Thus an approach from the viewpoint of dry corrosion study is considered

to be helpful to understand the corrosion behavior in nuclear reactor coolant water. In the previous papers by the present authors 4 ,5, the oxidation of iron was investigated by Rutherford backscattering spectroscopy at 573 K in the oxygen partial pressure range from 10- 1 to 105 Pa without and with water vapor at 103 . 15 Pa of water vapor pressure for 3 - 24 h. By comparing the oxidation rate of iron in gas phase with the corrosion rate in water, it was found that the oxidation behavior of iron in gas phase containing water vapor is closely related to corrosion behavior of carbon steel in high-temperature and high-pressure water around 573 K. The oxidation of iron in gas phase below 673 K has been studied by several investigators mainly under reduced pressures and the results are briefly summarized in our previous paper 4 . Studies on the effect of water vapor on the 6 oxidation of iron below 773 K are very few ,7.

116

Surman7 studied the oxidation of iron in controlled hydrogen-water vapor mixtures at 773 K and proposed the vapor-phase diffusion model. The oxidation behavior of iron-chromium alloys below 873 K has been studied by many ' au th ors 8-12 by uSIng sur f ace sensl't'Ive techniques such as X-ray photoelectron spectroscopy, Auger electron spectroscopy, Raman spectroscopy, etc. Gardiner et al. 8 have reported from Raman spectroscopy that the oxide film formed on iron-chromium alloys less than 5 wt%Cr consisted of two layers of magnetite and hematite. It has been found by Tjong et al. 9 and Tjong 10 that for the Cr addition of 3 wt%, the oxide consisted mainly of iron, and the addition of Cr from 9 up to 18 wt% resulted in the formation of layered oXides, i.e. Fe 0 2 3 existed in the outer oxide/gas interface and Fe _xCr x04 in the inner region with the 3 predominant chromium oxide next to the substrate. However, studies on the effects of the oxygen partical pressure and water vapor pressure on the oxidation of iron-chromium alloys have not been carried out yet. In this study, Rutherford backscattering spectroscopy (RBS) method was used to measure the thickness of oxide film formed, and the oxidation of annealed specimen of iron-chromium alloys (0 - 5 wt%Cr) was investigated at 573 K in the oxygen partial pressure range from 10- 1 to 105 Pa without and with water vapor at 10 4 Pa of water vapor pressure. 2. EXPERIMENTAL The iron-chromium alloys (0,1 and 5 wt%Cr) were prepared by cutting a cold-rolled sheet (about 0.6 rom in thickness) in a size of 15 x 15 rom, and polished mechanically with 100 and 600-grit silicon carbide papers successively. After degreasing with ethyl alcohol, the specimens were polished chemically for 5 min in a mixture of oxalic acid (10 g), hydrogen

peroxide (20 ml) and water (270 ml) at 315 K, followed by electropolishing at 20 volt in ar acetic acid (190 ml) and perchloric acid (10 ml) solution for one minute. These specimens were annealed in vacuum for 12 h at 1273 K and then for 2 h at 1023 K. The chemical- and electro-polishings were carried out again just before use. The apparatus is shown schematically in Fig. 1. The flowing gas was purified by passing through 5A molecular sieves column cooled with a dry ice-ethanol solution as a cold trap13. The control and measurement of oxygen partial pressure were described in the previous papers by the present authors 14 ,15. The control of water vapor pressure was made by passing through a water saturator which was maintained at 323 K. The oxidation of iron-chromium alloys (0 - 5 wt%Cr) was carried out at 573 K in the oxygen partial pressure range from 10- 1 to 105 Pa without and with water vapor at 10 4 Pa of water vapor for 1 - 48 h. The thickness of the oxide film oxidized repeatedly in the same oxygen partial pressure and water vapor pressure was determined by RBS, where a 1.5 MeV 4He+ and 1.0 MeV 1H+ ion beams were impinged vertically on the specimen for thinner and thicker oxide films than 500 nm, respectively.

~---------l

FIGURE 1 Schematic diagram of experimental apparatus.

117

F i n a l l y the surface o f the o x i d e was analyzed by means o f X-ray d i f f r a c t o m e t r y and e l e c t r o n microscopy

3.

scanning

(SEM).

RESULTS 3.1.

E f f e c t o f oxygen p a r t i a l

p r e s s u r e on

o x i d a t i o n o f iron-chromium a l l o y s F i g u r e 2 shows t h e t y p i c a l backscattering

Rutherford 4 + s p e c t r a when a 1.5 MeV He beam

was impinged on t h e o x i d e formed on i r o n and 1 wt%Cr a l l o y a f t e r 6 h and on 5 wt%Cr a l l o y 4

a f t e r 66 h a t 573 Κ and a t 1 0 ' oxygen p a r t i a l

pressure.

3

Pa o f the

As shown i n f i g . 2 ,

the edge o f i r o n i s around 400 channel,

and t h e

energy width between t h e m i d p o i n t s o f the s l o p e s can correspond t o t h e t h i c k n e s s 4 .

The

boundary between o x i d e l a y e r s was unable t o be found i n t h e s p e c t r a ,

although t h e o x i d e was

supposed t o c o n s i s t o f two l a y e r s ,

magnetite

and hematite containing some amounts of chromium. The thickness of the oxide film formed on iron-chromium alloys at 573 K and at 102 Pa of the oxygen partial pressure without water vapor is plotted against square root of time in fig. 3. It is seen from the figure that the thickness of oxide formed on various iron-chromium alloys (0,1 and 5 wt%Cr) increased linearly with square root of time during the measured oxidation time, which indicates that the oxidation obeyed the parabolic rate law. The parabolic rate constants obtained from the slope in fig. 3 are shown as a function of oxygen partial pressure in fig. 4 (a). In this figure, the previous results of the oxidation of iron by the present authors4 are also shown. As seen in fig. 4 (a) the t/h

Fe edge

n , '|

'PITTING'

15

'-

"PASSIVATION'

. -

- 1 0

.

. .

.

1 2 3 log(P0 / P a )

L

. i 1 — I

4

,i

L.

5

FIGURE 7 Comparison between c o r r o s i o n r a t e i n w a t e r and p a r a b o l i c r a t e constant i n gas phase w i t h o u t and w i t h water vapory 15 X Videm e t a l . 588 Κ i 1 f v T a c k e t t e t a l . a t 573 Κ ? Q i y O V r e e l a n d e t a l . a t 588 K . 1g qi Corrosion r a t e i n w a t e r ^ ' ' O x i d a t i o n o f 0( ) , 1( ) and 5( ) wt%Cr a l l o y s i n t h e absence o f w a t e r vapor O x i d a t i o n o f 0 ( 0 ) , 1 ( Δ ) and 5 ( G ) wt^Cr a l l o y s i n t h e presence o f water vapor

r a t e i s estimated t o be the a v e r a g e v a l u e between t h e beginning (10 - 100 h) and f i n a l (1,000 h - ) o f the c o r r o s i o n .

The i n c r e a s e o f

the c o r r o s i o n r a t e w i t h d e c r e a s i n g oxygen c o n c e n t r a t i o n i n water i s found t o be t h e same trend as t h e oxygen p a r t i a l

pressure dependence

o f the p a r a b o l i c r a t e constant i n t h e o f water v a p o r .

presence

The i n c r e a s e i n gas phase

o x i d a t i o n a t around 10^ Pa j u s t corresponds the boundary between t h e ' i n h i b i t i o n ' 'pitting'

i t i s suggested

regions in water c o r r o s i o n .

to

and t h e Although

5. CONCLUSION The oxidation of iron-chromium alloys (0,1 and 5 wt%) was investigated by Rutherford backscattering spectroscopy at 573 K in the oxygen partial pressure range from 10- 1 to 10 5 Pa without and with water vapor at 10 4 Pa of water vapor pressure for 1- 48 h, and the following conclusion was obtained. (1) When iron-chromium alloys was oxidized under various oxygen partial pressures without water vapor, the parabolic rate constant decreased with increasing chromium concentrations at constant oxygen partial pressure probably due to the decrease of iron diffusion through chromium oxide, and the mechanism of oxidation was divided into three regions (a) The parabolic rate constant decreased with increasing oxygen partial pressure from 10- 1 to 10 Pa due to increasing coverage of hematite over magnetite. (b) In the oxygen partial pressure range from 10 to 10 3 Pa, hematite could cover magnetite so uniformly that grain boundary diffusion would be reduced much. (c) At 105 Pa an increase of the parabolic rate constant occurred due to the simultaneous growth of hematite and magnetite. (2) The remarkable increase of the parabolic rate constant for the oxidation of ironchromium alloys (0 - 5 wt%) by water vapor was observed at lower oxygen partial pressure.

122

T h i s enhanced o x i d a t i o n o f iron-chromium

alloys

6.

P. L. Surman and J. E. C a s t l e , S e i . , 9 (1969) 7 7 1 .

7.

P. L. Surman, C o r r o s .

by water vapor may be e x p l a i n e d by vapor-phase diffusion

o f v o l a t i l e F e ( O H ) 2 from the

oxide interface

t o the outer l a y e r

( 3 ) The c o r r o s i o n behavior o f

metal-

crystals.

iron-chromium

a l l o y s in gas phase c o n t a i n i n g water vapor may be c l o s e l y r e l a t e d

t o that in high-temperature

and h i g h - p r e s s u r e water around 573 K.

P r o f . Y . Hosoi o f

Nagoya U n i v e r s i t y f o r supply o f

iron-chromium

a l l o y s samples and the use o f SEM. Prof.

13 (1973) 113.

8. D. J. Gardiner, C. J. L i t t l e t o n , Κ. M. Thomas and Κ. N. S t r a f f o r d , O x i d a t i o n o f M e t a l s , 27 (1987) 57. 9 . S. C. Tjong, J. E l d r i d g e and R. W. Hoffman, A p p l . Surface S e i . , 14 (1982-83) 297. 10. S. C. Tjong, Mat. Res. B u l l . , 157.

ACKNOWLEDGMENT The authors are indepted

Sei.,

We a l s o thank

K. Morita o f Nagoya U n i v e r s i t y f o r

the

Corros.

18 (1983)

11. G. H u l t q u i s t , M. Seo and N. Sato, o f M e t a l s , 25 (1986) 363. 12. H. J. Mathieu and D. L a n d o l t , 26 (1986) 547.

Oxidation

Corros.

Sei.,

use o f RBS.

13. Κ. N a i t o , T. T s u j i , T. Matsui and K. Une, J. Nucl. S e i . T e c h n o l . , 11 (1974) 22.

REFERENCES 1. R. A. Shaw, Nucl. T e c h . , 44 (1979) 97.

14. K. N a i t o , T. T s u j i and S. Watanabe, S t a t e I o n i c s , 1 (1980) 509.

2 . Y . Meguro, S. Sakai, K. Wakamatsu and T. Yamamoto, Genshiryoku Kogyo ( N u c l . E n g . ) , 25 (1979) 5. 3. K. Videm, 3rd Geneva Conf., V o l . 8 (1964) 600. 4 . H. Sakai, T. T s u j i and K. N a i t o , J. S e i . T e c h n o l . , 21 (1984) 844.

Nucl.

5. T. T s u j i , K. Okumura and K. N a i t o , P r o c . Sym. on High Temperature M a t e r i a l s Chemistry-IV, eds. by Z. A. Munir, D. C u b i c c i o t t i and H. Tagawa (The E l e c t r o c h e m i c a l S o c , I n c . , 1988) p . 156.

15. K. N a i t o and T. T s u j i , I n d . ) , 27 (1976) 518.

Solid

Kagaku Kogyo (Chem.

16. M. Hansen, C o n s t i t u t i o n o f Binary A l l o y s (McGraw-Hill Book Company, New York, 1958). 17. D. D. Wagman e t a l . , Nat. Bur. Stand. Tech. Note 270-3 ( 1 9 6 8 ) , Washington, D. C. 18. D. E. T a c k e t t , P. E. Brown and R. T. Esper, WARD-LSR(c)-134, B e t t i s P l a n t ( 1 9 5 5 ) . 19. D. C. V r e e l a n d , G. G. Gaul and W. L. Corrosion 17 (1961) 269.

Pearl,

High Temperature Corrosion of Advanced Materials and Protective Coatings Y . Saito, B. Önay and T. Maruyama (Editors) © 1992 Elsevier Science Publishers B.V. All rights reserved.

HIGH TEMPERATURE OXIDATION COMBUSTION ENVIRONMENT

OF

HEAT

123

RESISTANT

STAINLESS

STEELS

IN COG

B. G. SEONG, J. H. SONG, S. Y. HWANG, AND Κ. Y. KIM* Research Institute of Science and Technology, P. 0. Box 135, Pohang, 790-600, Korea (S) . *Pohang Institute of Science and Technology, P. 0. Box 125, Pohang, 790-600, Korea (S) In this study, the high temperature corrosion features of heat resistant stainless steels in COG (Coke Oven Gas) combustion environment are presented. These stainless steels were modified Mo-Re 1 and HK-40. They were used as radiant tubes in continuous annealing lines at cold rolled steel plants. These steels were severely corroded in a COG combustion environment. According to exposed temperatures of the alloys, these alloys exhibited various morphologies. Under the penetrated area due to severe corrosion, there was a large scale deposit. Considering a few possible modes of degradation, it was concluded that spalling of oxide layers was mainly responsible for the severe degradation. 1.

INTRODUCTION In a continuous annealing line (CAL) of cold rolled steels, radiant tubes have been used as an indirect heating method of steel strips in recent years. These radiant tubes consist of burners and tubes attached to the burner. Coke oven gas(COG) is used as a fuel for the burners. Table 1 shows the chemical compositions of COG and exit gases after combustion. These radiant tubes are usually made into a number of shapes. In the facility of concern, W-shape radiant tubes are used. Fig. 1 shows a schematic diagram of a radiant tube that is removed from service. The outer diameter of the tube is 19 0mm, and the thickness of the tube is 8 mm. Table 2 shows the alloy compositions of the first and second straight sections of the radiant tubes. While the first straight section alloy

is Modified MO-RE-1, straight section alloy HK-40.

the second is Modified

Concerning the corrosion of radiant tubes, holes were observed at the radiant tubes after 1.5 years even though these radiant tubes were expected to last for 3-5 years without perforation. Due to this problem, engineers measured the temperatures at various locations along the radiant tubes. (Table 3) Ideally, it has been known that a temperature should be maximum at the end section of first straight tube according to a manufacturer. However, the test result showed that a maximum temperature was exhibited at the initial part of the second straight tubes. Therefore, it has been thought that flames hit the location No. 4 of Fig. 1. Though the flame

124

Table 1.

Chemical Composition of COG and Exit Gases after COG Combustion.

1. Chemical Composition of COG

CH4

CXHY

CO

26.6%

2.9%

8.4%

H

2

56.4%

co2

N

°

2

3.1%

2

2.3%

0.3%

2. Chemical Composition of Exit Gases after (Vol. %)

N

o

74%

3.6%

2

CO

2

292ppm

co2

H 20

ΝΟχ

so x

9.5%

1 2%

R

R

Combustion

R: Residual Amount

Table 2.

Analyzed Chemical Compositions of Radiant Tube Alloys. (1st Straight Tube Section and n Straight 2 d Tube Section)

Alloys

Cr

Ni

C

Si

Mn

Mo

1st Tube

25

33

0.38

1.08

1.01

0.068

1.25

Bal.

2n d

26

20

0.39

1.01

0.8

-

-

Bal.

Tube

characteristics were under investigation by other researchers in our institute, this paper concentrated on high temperature corrosion mode of radiant tubes from service. 2.

PREPARATION OF SAMPLES.

Samples of radiant tubes were collected at various locations along radiant tubes. Fig. 1 showed the locations of samples that were collected. The samples were cut and epoxy mounted. These were polished with SiC papers and alumina. These samples were examined with SEM (Scanning Electron Microscopy) and WDS(Wave Dispersive Spectroscopy).

Length

3.

RESULTS

W

Fe

AND DISCUSSIONS.

First, visual observations of corroded radiant tubes and chemical analysis of deposits were made. The radiant tubes that were not exposed to high temperatures showed some voids at the inner part (Fig. 2 ) . It has been known that this kind of surface defects can induce local and severe corrosion. For the radiant tubes that were in service, the size of the perforated holes varied from 1 to 3 cm in diameter. Around the holes, the tubes showed a gradient of thicknesses. When the tubes were sectioned at the beginning of the 2nd straight section, there were large amounts of scales that were accumulated.

125

No.l

No. 2

1st Tube (Modified Mo-Re 1)

Perforated Part No. 4

Scale Deposit

it If FIGURE 1 Positions of specimens at radiant tubes.

Microstructures of straight section)

FIGURE 2 centr ifugally cast

The deposit was taken and analyzed via X-ray diffraction and chemical analysis. Table 4 showed the result of wet chemical analysis. This indicated that the deposits contained most of the alloying elements and were not protective scales. The scale was FesNiOß and Cr2Û3 according to X-ray diffraction. With this information, it

radiant

tube

(first

was concluded that the alloy did not form a protective scale and was in the propagation stage of oxidation. Fig. 3 shows the microstructures examined by optical microscopy. These specimens were collected from the 1st and 2nd straight tube. Specimen #1 was collected from the beginning part of the radiant tube. Its microstructure

126

Table 3.

Measured Temperatures at Various Locations of the Radiant Tube. The Location Numbers are the Same as Those in Fig. 1. Location

Temperature (°C)

Table 4.

No. 1

No. 2

No. 3

No. 4

No. 5

No. 6

991

1006

1014

1027

1011

1008

Chemical Composition of Deposit in Radiant Tube (Second Straight Tube)

Element wt.%

Cr

Fe

Ni

26.2

49.0

23.2

was close to the one that was as-cast. Eutectic carbides were present along grain boundaries, and the carbides were finely dispersed. Specimen #2 and #3 showed microstructures that were exposed at higher temperatures. The carbides present along the grain boundaries became enlarged, and carbides present in the grains became coalesced. Specimen #4, #5, #6 were collected from the 2nd straight part of the radiant tubes. Especially, Specimen #4 was collected from the area where the tubes were perforated. The specimen showed primary carbides and secondary carbides which became quite enlarged and agglomerated. To find the composition of the precipitates of Specimen #5, EDS (Energy Dispersive Spectroscopy) and WDS was used. Fig. 4 showed some chromium precipitates formed at the corrosion front. In Fig. 5, nitrides and carbonitrides were shown at inner part of Specimen #5.

Μη 1.3

Na 0.031

Si 0.26

C 0. 05

s 0.02

However, at the middle of thickness of Specimen #5, only carbides were shown. To check nitride formation on the outer part of the tubes. X-ray diffraction was performed on the part after slight polishing. The result indicated that there was a substantial formation of Cr2N precipitates. Therefore, internal nitrides formed the inside part and the outside part of tubes. Also, specimen #6 showed carbides that were somewhat enlarged. To observe the Cr depletion in the matrix of the alloy, EDS was used to measure the approximate concentration of Cr in the matrix of the alloy. Table 5 showed the measured results of Cr concentrations. Specimen #4 showed that the Cr concentration fell below 15 wt. % across the alloy. Especially, the Cr concentration of the inner surface fell to 10.5 wt. %. Also, there was a substantial depletion of Cr in the outer part of the tube. The Cr depletion is thought to be

127

Microstructures months.

of

radiant

FIGURE 3 tube at various

locations

after

32

FIGURE 4

SEM and x-ray maps of bottom part of radiant tube after 32 months use (No.5 specimen) a)SEM, b)-h)X-ray Maps, b)Fe, c)Cr, d)Ni, e)Si, f)Mn, g)0, h)S

129

SEM a n d specimen)

WDS

profile

of

FIGURE 5 bottom part

of

radiant

tube

(No.5

130

Table 5.

Cr Concentrations of Matrix at Various Locations. No. 1

No. 2

No. 3

No. 4

No. 5

No. 6

Outer

26.1

17.7

15.1

7.8

13.9

15.5

Middle

25.1

20.1

19.0

11.2

17.6

18.4

Inner

25.0

10.6

12.8

10.5

11.0

13.9

Specimen

related to the formation of carbides and nitrides. When the alloy became exposed to high temperatures, secondary carbides precipitated. This would reduce the Cr depletion. Also, the formation of nitrides would have the same effect. From the above results, it could be concluded that the depletion of Cr due to the formation of nitrides and carbides did not affect the corrosion resistance of the alloy because this alloy still could form protective scales after extensive formation of nitrides as shown in the outer part of the tube. Also, as far as the protective scale did not spall, the alloy seemed protective for a long time. Consequently, the spalling of the oxide layers is thought to be the most significant factor for the perforation of radiant tubes from the inside because there was a heavy scale deposit. It has been known that spalling induces rapid degradation of alloy. The spalling of oxide layers could be induced due to a thermal cyclic effect of a flame end. The thermal cyclic effect induced stresses at a scale/metal interface due to different thermal expansion coefficients. It has been known that

there are substantial thermal cycles at the flame end. This induces the thermal cyclic effect of metals. The thermal cyclic effect is well shown by 1 When S. Dils and Follansbee. D.(standard deviation) of a thermal cycle of flame is 139° C, the resultant S. D. of metal thermal fluctuation is 1.9° C for a FeCrAlY component. Generally, typical 2xS.D. of metal temperature cycle is ±5.6° C They showed that this kind of thermal cycles induced severe oxidation. In the radiant tube, a thermal cycle of flame was about 1300±100° C. We expected that this thermal cycle certainly induced severe oxidation when the flame hit the radiant tubes. According to the literature, there were carbur izat ion and sulfidation 2 5 processes at radiant tubes. " However, no evidence of severe carburization and sulfidation showed in this study. In Fig. 4, there was a trace of sulfur at the corrosion front. In the literature, sulfide formation was possible at the corrosion front 6 7 where P o 2 was low. ' However, rapid corrosion phenomenon due to sulfidation was possible when formation of liquid phase or substantial cation diffusion through sulfides or substantial removal of protective elements from scales

131

occurred. Since the amount of sulfides

important

observed in Fig. 4 appeared very small

of

for

possible

the

thought

above

mechanisms,

it

was

that rapid degradation due to checking

the possibility

of

vaporization Cr2C>3 via CrOß (g) , it was concluded

that

vaporization known

that

highly The

of

chromia

volatile

Cr

from

there

content

total

was

metal

forms

(0^03)

Cr0 3 of

little

It has been

(^203.

above

the

a

900o

scale

weight

radiant

was

C.

8

deposit

for

the

tubes.

perforation

Though

degradation

(vaporization

Cr 203,

of

nitridation,

sulfidation was not likely. After

the

factors

carburization)

a few modes

internal

sulfidation, were

and

discussed,

these

modes of degradation were not thought to be likely.

To elongate the life of

the radiant tubes, the repositioning of the

burners

and

adding

some

oxygen

active elements to the alloys could be used.

compared

with the Cr content of the alloy.(Table

ACKNOWLEDGEMENT

4)

Authors would like to thank Pohang Iron

The result showed that Cr remained

mostly in the deposit.

This indicated

that there was little vaporization of

and

Steel

Company,

Ltd.

(POSCO)

for

providing a funding for this research.

the chromium oxide in this case. As

a

summary,

the

severe

local

degradation sequence was thought to be as follows.

REFERENCES

In the beginning of the 1.

R. P. Dils and P. S. Follansbee, Corrosion, 33 (1977) 385.

2.

S. S. Pani, et al., Tool and Alloy Steels, 18 (1984) 117.

3.

Metals Hand Book, 9th Vol. 11, (1986) 292.

4.

B. S. Starokozhev, Met. Sei. Heat Treat (USSR) 22, (1980) 661.

alloying elements.

5.

G. A. Slisarenko, Met. Sei. Heat Treat. (USSR), 27 (1985) 324.

4. CONCLUDING REMARKS AND FUTURE WORK.

6.

P. Singh and N. Birks, Oxidation of Metals, 19, (1983) 37.

7.

F. S. Pettit, et al., Corrosion Science, 9 (1969) 903.

8.

N. Birks and G. H. Meier, Introduction to High Temperature Oxidation of Metals, (Edward Arnold Ltd., 1983) 80.

operation, a protective oxide layer may have

formed.

But

due

to

highest

temperature and high thermal cycles at the end of the flame, the scale spalled continuously. spalling,

the

protective spalled

In

alloy

scales. scale

this

degradation discussed.

After no As

a

longer a

contained

paper,

the

of

radiant

the The

number

formed

result, most

severe

highest

of a of

local

tubes

was

temperaure

exposure and spalling of scales due to the thermal cyclic effect of the flame end were thought to be the most

Edition,

This page intentionally left blank

High Temperature Corrosion of Advanced Materials and Protective Coatings Y . Saito, B. Önay and T. Maruyama (Editors) © 1992 Elsevier Science Publishers B.V. All rights reserved.

INTERNAL ΝI TRI DAT I ON OF N i - C r - A I

R. P. RUBLY*

and

*AIIied-Signal CA, 9 0 5 0 9 - 2 9 6 0 , **Materials

133

ALLOYS

D. L DOUGLASS**

Aerospace

Co., Ai Research,

Los A n g e l e s

Division,

Torrance,

USA.

Science

and E n g i n e e r i n g

Department,UCLA,Los

Angeles,

CA, 9 0 0 2 4 - 1 5 9 5 ,

USA.

The internal nitridation of Ni-10Cr-5AI and Ni-20Cr-5AI alloys was studied over the range of 700-900'C in ammoniahydrogen mixtures. The reaction kinetics followed the parabolic rate law. There appeared to be a two-stage process in which initial rapid kinetics were followed by slower kinetics. Activation energies for the rate constants were 50.3 and 42.5 Kcal.mol for Ni-10Cr5AI and Ni-20Cr-5AI, respectively. Both CrN and AIN formed at 800 and 900*C, whereas only CrN formed at 700'C. No C ^ N formed under any conditions. The precipitates were extremely fine at 7 0 0 Ό but increased markedly in size with increasing temperature. The precipitate number density and size varied with position within the reaction zone, the density decreasing with distance from the surface, and the size increasing with distance.

CrN tended to form small spheroidal particles near the surface, whereas AIN formed Widmanstätten

plates and/or blocky particles away from the surface toward the reaction front.

The morphology of AIN was checked by

nitridation of Ni-5AI at 900'C for 48 hours. Ni-rich surface protrusions formed in all cases, the size and number density of the protrusions increasing with increasing alloy content and time. The formation of the protrusions is attributed to stresses induced by the volume change due to precipitation, which in turn causes pipe-diffusion-controlled creep as a stress-relief mechanism. This mechanism, suggested by Guruswamy, et al, results in nodule formation on the surface. Results are compared to previous work on internal nitridation of Ni-Cr binary alloys and to the internal oxidation/external oxidation transition of these same alloys. In general, higher solute concentrations are required for the transition during nitriding compared to oxidation. Reasons for this behavior are discussed.

1.

INTRODUCTION

Although nitridation of steels has been commercially

nitridation has involved iron-base alloys, whereas very little

utilized for about 100 years, the phenomenon of internal

work has been reported on nickel-base alloys. In general, internal

nitridation has been studied very sparsely compared to internal

nitridation rates are considerably less in nickel-base alloys

oxidation. Furthermore, the major emphasis in studies of internal

compared to iron-base alloys due to the much lower nitrogen

134

solubility and diffusivity in nickel. The present authors recently

to the test samples.

reported work on the internal nitridation of Ni-Cr binary alloysC).

Bottled commercial-purity gases were used for nitriding

The work reported here concerns internal nitridation in Ni-Cr-AI

and for purging the system before and after nitriding runs.

alloys in which both Cr and Al form stable nitrides. In addition,

Nitriding was performed using an ammonia/hydrogen mixture

one binary Ni-AI alloy (5 w/o Al) was investigated at one

while purging was performed using pure nitrogen. Gas flow rates

temperature to discern the nature of AIN formed during internal

and mixtures were controlled using independent flow meters. A

nitriding. The main thrust of the research was to determine the

total gas flow rate of 200 ml/miη was used for nitriding.

kinetics of the reaction and to study the morphology of

composition of the nitriding gas for all runs was 10% ammonia

precipitates forming within the reaction zone.

(by volume).

It was of interest

In order to minimize oxygen and moisture contamination

also to compare the nitriding behavior to internal oxidation and to

within the system, all gasses were purified before introduction

discern the nature of the observed differences.

into the reaction tube. 2.

The

packed

EXPERIMENTAL PROCEDURES

with Drierite

Nitrogen was passed through a tube desiccant.

Ammonia was

passed

successively through potassium hydroxide, calcium oxide and

Sample Preparation Samples were obtained from an earlier investigation and were in the form of arc-melted buttons. All samples were in the

Drierite. Hydrogen was passed successively through an oxygen getter containing titanium shavings at 700'C and Drierite. Test samples were suspended on a platinum wire

solution-annealed condition. Small slices (approximately 1.5 χ 1.0 cm) were removed

attached to a magnet and lowered into the top of the reaction

The samples were

tube. The magnet and wire were held in place using a second

ground through 600 grit paper, polished with 6-micron diamond

magnet on the outside of the tube. A small crucible, containing

paste, ultrasonically cleaned in methanol, rinsed in acetone, and

approximately 1 gram of polyvinyl chloride (PVC) powder to act

blown dry immediately prior to nitriding.

as a fluxing agent, was also attached to the wire just below the

Test Apparatus and Procedures

samples.

from the buttons (1.0 to 2.0 mm thick).

Nitriding experiments were performed using a laboratory 2

After sealing the tube, the system was purged with

nitrogen, evacuated with a roughing pump, purged with nitrogen

nitriding system described by Chen and Douglass.* ) The system

again, and evacuated a second time. The tube was then purged

consists of a quartz reaction tube, the center portion of which

with the ammonia/hydrogen

passes through a tube furnace.

Following the purging cycles, the inlet and exhaust ports were

The tube and furnace are

gas mixture for 30 minutes.

oriented in the vertical position. The nitriding or purging gases

closed, and the samples and crucible were lowered to an

enter the bottom of the reaction tube, flow up through the tube,

intermediate position at which point the PVC was volatized.

and are exhausted at the top end. Temperature

The samples were held in the PVC vapor for 10 minutes.

measurement

The

was accomplished using a thermocouple placed within the

PVC

reaction tube in the hot zone of the furnace immediately adjacent

mixture introduced, and the samples were lowered into the hot

vapor was then exhausted, the

ammonia/hydrogen

135

zone of the furnace commencing the nitriding run.

typically observed below the Widmanstätten structure. In some

Sample Characterization After nitriding, a portion of each sample was removed and prepared for metallographic examination.

Optical and

scanning electron microscopy (SEM) were used for evaluation. Semiquantitative chemical evaluation was performed using energy dispersive x-ray (EDX)

analysis.

Microhardness

traverses, using a Vickers indentor, were performed across the nitrided regions on representative samples.

X-ray diffraction

(XRD), using copper Κ α radiation, was performed to identify nitrides on the exposed surfaces and within the samples using a step-grinding procedure.

3.

RESULTS

Microstructure: Ni-10Cr-5AI Micrographs of cross-sections of internally-nitrided zones in Ni-10Cr-5AI samples nitrided at 700,800 and 900'C are shown in Figs. 1 & 2.

XRD scans identified both CrN and AIN in

representative samples for the 800 and 900°C exposures, whereas only CrN was detected at 700'C. e

The 900 C samples exhibited a relatively coarse Widmanstätten structure composed of discrete, plate-like precipitates.

A secondary dispersion of smaller, irregularly-

shaped precipitates was observed interspersed within the Widmanstätten structure to a depth of approximately 2/3 the depth of the plate-like precipitates.

In some areas, the

Widmanstätten structure gave way to a very coarse dispersion of large blocky precipitates. The plate-like precipitates extended inward in a direction generally normal to the exposed surface, and tended to exist in colonies of parallel plates with one or two preferred orientations evident within most colonies. The large, blocky precipitates were

Fig. 1.

Microstructure of internal-nitridation zone in Ni-10Cr-5AI nitrided 96 hours at various temperatures.

136

TABLE I Matrix Alloy Compositions: Ni-10Cr-5AI Composition,

Area

Al

Shallow (1) Deep (2) Below Reaction-front Base Alloy (4)

1.0 0.1 3.0 3.2

(1 ) Fig. 2.

Variation of precipitate morphology in Ni-10Cr-5AI (2)

nitrided at 900*C for 48 hours.

(3) (4)

(2)

w/p

Cr 28 9.5 9.4 9.6

Ni 962 90.5 87.6 872

Shallow region of internally-nitrided zone where both CrN and AIN formed. Deep region of internally-nitrided zone where only AIN formed. Base alloy just below reaction-front. Base alloy well below reaction-front.

areas, the blocky precipitates were observed at shallower depths, but always in the absence of the plate-like precipitates. EDX analysis indicated the plate-like and the blocky

For the 800'C samples, the internally-nitrided zones were similar but much more finely dispersed as compared to the

precipitates to be rich in aluminum, while the interspersed

900'C samples.

precipitates were found to be rich in chromium. These results

observed, however the individual precipitates were much smaller

are in agreement with the observations from back-scattered-

and the dispersion much finer. The plate-like precipitates were

electron (BSE) images. As previously noted, XRD indicated the

visible only in the deeper regions of the internally-nitrided zones.

presence of CrN and AIN on representative samples.

The shallower regions consisted of a fine, unresolved dispersion.

The

Discrete, plate-like AIN precipitates were

metallographic observations combined with the EDX and XRD

The 800'C samples were found to have a dual-layer

results indicate the plate-like and the blocky precipitates to be

appearance in some areas, especially for the 24-hour exposure

AIN, and the interspersed rounded precipitates to be CrN.

sample as shown in Fig. 3. The dual-layer structure appears to

Additional EDX evaluation was performed on the matrix alloy within and just below the internally-nitrided zone.

be due to the formation of both AIN and CrN.

Visual

Results

observations indicated the inner layer to consist primarily of AIN.

are shown in Table I. The composition of the base alloy was

The outer layer was not resolvable microscopically, but XRD

found to be close to the normal Ni-10Cr-5AI composition. Within

results indicated both AIN and CrN on and just below the

the internally-nitrided zone in the area where both CrN and AIN were present, the matrix was completely denuded of aluminum and partially denuded of chromium. In the deeper region where only AIN formed, the matrix was again completely denuded of aluminum, but there was no measurable chromium depletion.

exposed surface. EDX analyses were performed on the two layers.

Due to the fine nature of the dispersions, it was not

possible to analyze individual precipitates, so area scans were made in both layers.

Results of the area scans indicated the

overall compositions to be nearly the same as the base alloy

137

temperature. EDX analysis indicated the surface protrusions to be rich in nickel with only minimal indication of chromium and aluminum.

For the 900'C samples, a thin external layer was

observed also along the exposed surfaces, including the protrusions. Windowless EDX analysis was used to evaluate the protrusions and the thin external layer. This technique has the capability to qualitatively identify low-atomic-number elements including nitrogen and oxygen. The outer layer was found to have a significant oxygen peak indicating the layer to apparently be an oxide. The protruded nodule below the layer was rich in nickel but showed no indication of oxygen or nitrogen. Standard Fig. 3.

Dual-layer nature of internal-nitridation zones of Ni1 0 O 5 A I nitrtided for 24 hours at 800'C.

EDX analyses of typical protruding nodules confirm the composition of the nodules to be essentially pure nickel and the

outer layer to be rich in nickel with a

slight

chromium content (on

the order of 10 wt%).

composition in both of the layers. For the 700'C samples, the internally-nitrided zones

Microstructure: Ni-20Cr-5AI

consisted of a very uniform, featureless dispersion. Individual

Typical photomicrographs of cross-sections through the

plate-like precipitates were barely resolvable in the deepest

internally-nitrided zones for the Ni-20Cr-5AI samples nitrided at

areas.

XRD scans on the exposed surfaces and after step-

700, 800 and 900'C are shown in Fig. 4. XRD scans identified

grinding on representative samples resulted in weak indications

both CrN and AIN in representative samples for 800 and 900'C

of CrN only. There was no indication of AIN on the surface or

exposures, but at 700'C only CrN was detected.

within the nitrided zones for any of the 700'C samples examined.

The general microstructural features for the Ni-20Cr-

Reaction-fronts for the Ni-10Cr-5AI samples nitrided at

5AI samples were found to be very similar to those for the Ni-

900'C were fairly planar in most areas.

The reaction-fronts

10Cr-5AI samples.

The internally-nitrided zones of the 900'C

were very uniform for the 700 and 800'C samples. There was

samples consisted of a Widmanstätten structure composed of

no evidence of intergranular nitride formation observed for any

large plate-like precipitates, a secondary dispersion of smaller

of the samples.

precipitates in the shallow regions, and a coarser dispersion of

For most of the Ni-10Cr-5AI samples, protrusions were

large, blocky precipitates in the deep regions.

observed along the exposed surfaces as can be seen in cross

The 800'C samples once again were seen to have a

section in Figs. 1-2. The protrusions tended to increase in size

dual-layer appearance similar to the Ni-10Cr-5AI samples. The

and in number density with increasing exposure time and

24-hour exposure sample had a distinct dual-layer appearance

138

Fig. 5.

Ni-20Cr-5AI nitrided for 24 hours at 800'C showing dual layers in the internal-nitridation zone.

The structure of the Ni-20Cr-5AI samples nitrided at 700'C consisted of a very uniform, unresolvable dispersion similar to that observed for all samples nitrided at 700'C.

XRD

indicated the presence of CrN only, both on the exposed surface and within the internally-nitrided zones. BSE imaging was again used to highlight contrast differences between the various phases in samples nitrided at 800 and 900'C.

As with the Ni-10Cr-5AI samples, the visual

observations from the BSE images, combined with EDX and XRD results, indicated the plate-like and the large blocky precipitates to be AIN, and the smaller, interspersed precipitates in the shallow regions of the internally-nitrided zones to be CrN. XRD scans through the internally-nitrided zones were Fig. 4.

Microstructure of internal-nitridation zones in Ni-20Cr5AI nitrided 96 hours at various temperatures.

again performed by step grinding. XRD scans of the as-exposed

surface and after very light grinding indicated the presence With

both

as shown in Fig. 5. The plate-like precipitates were visible in the

only CrN.

inner layer only.

detected, with the latter becoming more predominate with

The outer layer consisted of a very fine,

unresolved structure. As shown in Fig. 5, separation between the two layers was observed in some areas.

additional grinding,

of

CrN and AIN were

additional grinding steps. Energy dispersive x-ray maps of the internally-nitrided

139

zone for the sample nitrided at 900'C for 48 hours are shown in Fig. 6. The x-ray maps revealed a chromium-rich region along the exposed surface. Discrete chromium-rich areas were seen to extend from the exposed surface to a depth of approximately 1/2 of the internally-nitrided zone.

Aluminum-rich areas were

seen to extend from just below the exposed-surface to the reaction-front. Semiquantitative EDX results for the composition of the matrix alloy within and just below the internally-nitrided zone are shown in Table II. The results show that the composition of the base alloy is close to the nominal Ni-20Cr-5AI composition. Within the internally-nitrided zone, the matrix was completely denuded of aluminum and partially denuded of chromium in the shallower region where both AIN and CrN formed.

In the deeper region

where only AIN formed, the matrix was completely denuded of aluminum but was not depleted in chromium. Protrusions were observed again along the exposed surfaces of the Ni-20Cr-5AI samples, increasing in size and number with increasing alloy content. For the sample nitrided at

TABLE II Matrix Alloy Compositions: Ni-20Cr-5AI

Composition,

w/o

Area

Al

Cr

Ni

Shallow (1) Deep (2) Below Reaction-front (3) Base Alloy (4)

0.3 0.9 4.0 3.1

5.3 19.0 19.4 19.6

94.4 802 76.6 77.3

(1 ) (2) (3) (4)

Shallow region of internally-nitrided zone where both CrN and AIN formed. Deep region of internally-nitrided zone where only AIN formed. Base alloy just below reaction-front. Base alloy well below reaction-front.

Fig. 6.

EDX X-ray maps of Ni-20Cr-5AI nitrided for 48 hours at 900'C.

140

900'C for 96 hours, the surface protrusions formed a nearly continuous layer.

A thin external layer was observed also

similar to that observed on the Ni-10Cr-5AI sample nitrided at 900'C. The exposed surface of the sample nitrided at 900'C for 24 hours was examined in plan view in the SEM.

The surface

was found to be covered with a closely-packed array of rounded nodules as shown in Fig. 7.

EDX analysis was performed on

these nodules and compared with the underlying surface.

The

results indicated the nodules to be essentially pure nickel, while the composition of the underlying surface was found to be similar to the base-alloy composition.

Fig. 7.

SEM micrograph of nodules formed on the surface of Ni20Cr-5AI nitrided for 24 hours at 800'C.

Micrpstructure; Ni-SAI A single Ni-5AI sample was nitrided at 900'C for 48 hours to evaluate the characteristics of AIN formation in the absence of CrN. The internally-nitrided structure for this sample is shown in Fig. 8. The microstructure consisted of a relatively coarse dispersion of short, plate-like precipitates. In contrast to the AIN precipitates observed on the Ni-Cr-AI ternary alloys, which formed in parallel colonies oriented generally normal to the surface, the AIN precipitates for the binary alloy were randomly oriented with respect to the surface. The precipitates still had a plate-like morphology, but were short and wider than those observed in the ternary alloys. The exposed surface was again found to have a relatively dense distribution of protrusions.

Fig. 8.

Microstructure of Ni-5AI internally nitrided for 48 hours at 900'C.

Kinetics Kinetics plots, Figs. 9 and 10, are given in parabolic form, i.e., reaction zone thickness vs t

1 2/

for both ternary alloys at

positive values of the ordinate, whereas the 900'C curve for Ni-

three temperatures. The straight-line fit of the data shows that

10Cr-5AI extrapolates through a positive value on the abscissa.

the parabolic rate law was followed. However, none of the lines

The likely explanation for positive values of the ordinate is that

extrapolate through zero. All of the curves for Ni-10Cr-5AI and

a two-stage process occurred, involving a rapid initial stage

two for Ni-20Cr-5AI (700 an 800'C) extrapolate through some

followed by a slower second stage which is seen in Figs. 9 and

141

0 1

2

3

4

5

6

7

8

9

10

11

11 12

-9 -9.5 A

NÎ-20C

-10 CL

0 1

2 3 4 5 6 7 8 9 10 Exposure Time (square root hours)

11 12

Ο

-10.5

ο

20Cr-5Al

-11 A Fig. 9.

Parabolic plots of internal-nitride growth in Ni-10Cr-5AI

-11.5J

at various temperatures.

-12 0 1

2

3

4

5

6

7 8

9

10

9 10 1/T χ 10,000

11 12

11

Fig. 11. Arrhenius plot of internal-nitridation kinetics.

Microhardness 0 1

2 3 4 5 6 7 8 9 10 Exposure Time (square root hours)

11 12

Microhardness traverses of Ni-20Cr-5AI nitrided at three different temperatures are shown in Fig. 12. Several features of these plot are noteworthy.

First, there is no abrupt change in

hardness at the reaction fronts, the location of which are shown Fig. 10. Parabolic plots of internal-nitride growth in

by arrows on each plot. Second, the hardness is greater at the

Ni-20Cr-5AI at various temperatures.

10 for the times studied. The positive abscissa value is usually associated with an incubation period.

TABLE III Experimentally-Determined Activation Energies

The temperature dependence of the rate constants is shown in Fig. 11. This plot includes data for Ni-1 OCr and Ni1

Alloy

Q (kcal/mole 'Κ)

20Cr from previous work.( ) The rate of nitriding decreased, as

Ni-1 OCr

41.1

expected when 5% Al is added to the binary alloys. Activation

Ni-20Cr Ni-10Cr-5AI Ni-20Cr-5AI

27.4 50.3 42.5

energies for all four alloys are given in Table III.

142

desirable for abrasion resistance, then the Ni-20Cr-5AI alloy

700°C

nitrided at 7 0 0 Ό would provide the hardest case.

4.

DISCUSSION

Precipitate Morphology A marked difference existed in the morphologies of

Ο

10 20 30 40 50 60 70 80 90 100 Depth (microns)

precipitates formed in Ni-Cr binary alloys versus those found in Ni-Cr-AI ternary alloys. There was also a significant difference

800 700 ^600

I

800°C

between the morphology of AIN formed in Ni-5AI and that of CrN formed in either Ni-10Cr or Ni-20Cr.

CrN formed small

500

precipitates that were generally

spheroidal, or at

least

CO

S 400 200 100

approaching spheroids. On the other hand, AIN formed either Widmanstätten or blocky precipitates.

0

10 20 30 40 50 60 70 80 90 100 Oepth (microns)

Ternary alloys formed

both types of precipitates with the spheroidal CrN particles dominating near the surface, and the AIN platelets dominating near the reaction front. AIN is more stable thermodynamically

700 600 CO «500

900°C

within the reaction zone, and this was indeed the case.

υ

CO cu "§300 χο 200

100

than CrN, thus it is expected that AIN only would exist deep The

exterior portion of the zone, close to the surface, contained primarily CrN, whereas only AIN formed at the front.

40

60 80 100 Depth (microns)

120 140

The

shallow regions containing both CrN and AIN were exposed to a nitrogen activity sufficiently high to form both nitrides, but the lower nitrogen activity near the front was too low for CrN

Fig. 12. Microhardness profiles in Ni-20Cr-5AI nitrided for 96 hours at various temperatures. Arrows indicate the reaction fronts.

formation but still high enough for AIN to form. This effect was more noticeable in Ni-20Cr-5AI than in Ni-10Cr-5AI.

It is

apparent also that the alloy having 20Cr was very close to the critical value for the transition from internal nitridation to lower nitriding temperature and decreases significantly with

continuous-film formation, and the near-continuous nitride film

increasing nitriding temperature. Third, the hardness is greater

may inhibit nucleation of AIN which is known to have nucleation

in the alloy containing 20 wt% Cr compared to the 10-Cr alloy.

problems in Fe-base alloys, particularly at the lower

This is expected in view of the greater volume fraction of

temperatures.( )

precipitates in the higher-Cr alloy. If the high surface hardness is

3

143

There appeared to be two distinct morphologies for AIN:

large volume change associated with internal-oxide formation

Widmanstätten or blocky. However, these may very well be the

which in turn produced a stress gradient between the stress-free

same. If the Widmanstätten platelets are sectioned parallel to

surface and the reaction front. It was suggested that silver

the plane of the platelet rather than transversely, they will

transported to the surface by pipe-diffusion-controlled creep in

appear as blocky particles.

Thus the orientation of the

order to allow stress relief to occur. The nodules resulted from

precipitates with respect to the plane of polishing may determine

the excess silver arriving at the external surface. Nitridation

the appearance of the particles.

also results in a significant volume change which will create a

The precipitate size increased with increasing distance

stress gradient. Thus, Ni, the unnitrided element in the alloys,

from the surface, but the number density of particles decreased

diffuses to the surface via dislocation pipes.

with increasing distance. As noted in many other studies of

enables sufficient creep to occur so that the stress gradient is

internal oxidation and nitridation, this is consistent with

eliminated.

nucleation dominating near the surface and growth becoming

Kinetics

more important toward the reaction front. Nucleation of new particles depends on the degree of supersaturation in advance of

This process

The rate of thickening of the internal-reaction zone for 6

the limiting case of ΫΝ

is given by( )

(

Ν Β'

existing particles. This is determined by the relative magnitudes of nitrogen and solute fluxes. Near the surface, the nitrogen flux is high, and nucleation is relatively easy. As the front advances,

dt

=

the nitrogen flux decreases, and nucleation of new particles becomes increasingly difficult. Thus, growth of existing particles

where N ^ =

mole fraction of nitrogen at

becomes more favorable, resulting in larger, more-widely-spaced precipitates.

(1)

(0

2 v N B) t

the surface DN

=

diffusivity of nitrogen in the

t

=

time

ν

=

ratio of nitrogen to metal in

substrate

Surface Protrusions Nearly pure Ni surface nodules formed on all samples,

the nitride

increasing in size and density with increasing solute concentration and/or exposure time. Other investigators have

N

mole fraction of solute in the alloy

4

noted similar behavior. For example, Stott and Wood( ) noted

The parabolic rate constant, kp, is expressed as

nodules of nickel virtually depleted in solute during internal 5

oxidation of Ni-AI alloys. Guruswamy et al( ) found nodules of Ag on Ag-ln alloys after internal oxidation. The appearance of

S

kp

t

2NN( ) DN ν NB(°)

(2)

the protrusions on Ag-ln alloys was nearly identical to the

The temperature dependence of the Arrhenius plot can

protrusions observed as a result of internal nitridation in this

be attributed to two factors-the variation of nitrogen solubility

study. Guruswamy et al attributed the nodule formation to a

with temperature and the temperature variation of the nitrogen

144

diffusivity. The solubility of nitrogen in Ni is reportedly very small and shows little variation.with temperature. Wriedt and 7

TABLE IV Comparison of Some Diffusivities of Interstitial Elements

Gonzalez* ) measured nitrogen solubilities in Fe-Ni alloys ranging 2

Diffusion Coefficient, cm /sec

from pure Fe to pure Ni. The very low solubilities in high-Ni alloys is within the range of experimental scatter, but if data for

Ν in Fe-20Ni (Ref8)

T.'C

Ν in Ni-Alloys (This Studvl

700

9.5x10-^xlO"

have a decreasing solubility of nitrogen with increasing

800

3.2x10-β-δ,δχΙΟ-

temperature, whereas Ni-rich alloys show increasing solubilities

900

1.4x10- -4.Ox1O"

more dilute alloys are considered, it is seen that Fe-rich alloys

8

8

1.17X10"

8 7

8

3.86X10" 7

7

1.4x10"

Oin Ni (Ref 12)

Cin Ni (Ref 9) 9

3.19x10"

8

1.47Χ10"

7

0.55X10

11

7.48x10'

10

5.05x10"

9

2.38x10"

with increasing temperature. The temperature dependence is quite small, about 3 Kcal/mol for Fe-40Ni. No diffusivity data of nitrogen in Ni are available,

attributed to the plate-like morphology of AI2O3 precipitates

however some approximations can be obtained from the values

which enabled boundary diffusion to occur between the particles

of the N n D n product and of N n mentioned above. The lowest

and the matrix.

value of N n reported by Wriedt and Gonzalez for nearly pure

An Arrhenius plot of log N n D n and N q D o obtained from

nickel is 0.0001. Using this figure, one obtains Dn values ranging

internal nitridation, internal oxidation, and No Do values

9

8

8

8

from 9.5 χ 10' to 2.3 χ 10" at 700'C, 3.2 χ 10' to 8.5 χ 10" at 7

7

calculated from solubility and diffusivity data in the literature is

800'C, and 1.4 χ 10* to 4.0 χ 1 0 at 900'C, depending upon the

given in Fig. 13.

particular alloy. Comparison of these results with other relevant

permeabilities for nitriding are significantly greater that those for

data is shown in Table IV. The values obtained from the N n D n

internal oxidation, even though rapid interfacial diffusion of

8

product compare very closely to nitrogen diffusion in Fe-20Ni( ) 9

Several features are obvious.

First, the

oxygen in Ni-AI alloys is reportedly orders of magnitude greater 10

and for carbon diffusion in nickel.* ) The lattice parameter of Fe-

than lattice diffusion.* ) Second, there was little difference

20Ni is about 1.7% larger than that of pure nickel, thus it is

between the various alloys for internal nitridation with the

expected that nitrogen diffusion in Fe-20Ni should be slightly

exception of Ni-10Cr-5AI which appears to have lower

faster than in pure nickel due to the larger interstitial sites. On

permeabilities at all temperatures. On the other hand, the alloy

the other hand, nitrogen is slightly smaller than carbon, thus the

with 5AI but a higher Cr content, Ni-20Cr-5AI, appeared to have

diffusivity of nitrogen should be slightly higher than carbon.

permeabilities about the same as the binary Ni-Cr alloys except

It is interesting to compare values of N n D n to NoDo in 1011

for the datum point at 800'C.

Third, although only one

Ni-base alloys. Stott and co-workers* ' ) have studied internal

temperature was studied, NÎ-5AI had a slightly lower

oxidation in various Ni-base alloys, including Ni-AI alloys. There

permeability for nitrogen than for oxygen at 900'C. Although

was virtually no difference (with one exception) in their values

AIN formed as Widmanstätten platelets in all cases which were

of No Do among the various alloys, the exception being that

very similar to those formed during internal oxidation, the

much larger values were found in Ni-AI alloys.

presence of the interfaces apparently did not result in enhanced

This was

145

Temp.

10

" ' V

(°C)

fastest and the largest species slowest. This same behavior

—

Ni-20Cr-5A ! (int. nitrd. ) :

was noted between carbon and oxygen in FCC γ-Fe as reported

V

13

by Hauffe.( ) The only explanation, which is unsubstantiated, is that oxygen is trapped in the FCC lattices of Fe and Ni. Transition from Internal reaction to Continuous-Film Formation It was observed previously^) that much higher solute levels appeared necessary for continuous films to form during nitriding than during oxidation. The critical mole fraction of Cr, Ncn to form CrN on Ni-Cr alloys is between 0.3 and 0.4, whereas the amount required to form

&2Ο3

during oxidation is

about 0.15, depending upon temperature and oxygen pressure. -j 8.0

,

9.0

1/T x10

Fig.

13.

1

,

4

1

1 10.0

1 11.0

(Ί/Κ)

Arrhenius plot of permeabilities,

NnDn

Oxidation in air of the two ternary alloys reported in this study resulted in continuous-film formation, as seen in Fig. 14. The

and

NoDq

difference in the values of Ncr for oxidation and nitridation is

determined from internal nitridation, internal oxidation,

attributed to the large difference in the permeabilities as already

and measured values for oxygen.

noted. Wagner proposed( ) that the transition occurs from

14

internal to external oxidation when the solute concentration is interfacial diffusion during internal nitriding, unlike the behavior

sufficient to form some critical volume of oxide particles at the

noted in internal oxidation.

reaction front. Assuming that diffusion of the oxidant in the

On the basis of permeabilities it is not surprising that

oxide particles is slow compared to diffusion in the matrix, one

internal nitridation in Ni-base alloys is more rapid than internal

concludes that diffusion can occur only in the channels between

oxidation. What is surprising, however is the much higher

particles. Growth of particles at the reaction front can proceed

permeability of nitrogen compared to oxygen. It appears that

by sidewise growth, or new particles may nucleate. Sidewise

the oxygen solubility is slightly higher than the nitrogen solubility

growth of the particles would favor continuous-film formation,

12

on the basis of Park and Alstetter's results( ) and those

whereas rapid diffusion of the oxidant between the particles

mentioned previously by Wriedt and Gonzalez.^) Thus, the

would cause supersaturation in advance of the front and

diffusivity of nitrogen appears to be nearly two orders of

nucleation of new particles. Thus rapid diffusion of the oxidant

magnitude greater than that of oxygen, as noted in Table IV.

favors formation of new particles, but slower diffusion of the

The diffusivity of carbon in nickel is likewise nearly two orders

oxidant enables sidewise growth of existing particles to occur.

of magnitude greater than that of oxygen. The respective

Because nitrogen diffusion in Ni and Ni-base alloys is so much

atomic radii of carbon, nitrogen, and oxygen are 0.77, 0..70, and

more rapid than oxygen diffusion, a higher concentration of

0.61 Â. It is expected that the smallest species would diffuse

solute is necessary to achieve the critical particles density for

146

continuous-film formation.

4

WagnerO ) also notes that if the

10.

F. H. Stott, G. C. Wood, D. P. Whittle, B. D. Bastow, Y. Shida and A. Martinez-Villafane, Solid State Ionics, iZ (1984) 365.

11.

F. H. Stott, A. Martinez-Villafane and G. C. Wood, Proceedings, Internal Congress on Metallic Corrosion, Toronto 1984, National Reserach Council of Canada, Vol. Ill, 317.

12.

Jong-Wan Park and Carl J. Alstetter, Met. Trans A, 18A (1987) 43.

13.

K. Hauffe, Oxidation of Metals. Plenum Press, N.Y., (1965) 45.

14.

Carl Wagner, Corr. Sei., 5, (1965) 751.

oxygen diffusivity is decreased (low partial pressure of O2 for Ag-ln alloys), the outward diffusion of solute becomes significant and leads to a higher particle density and easier transition from internal oxidation to continuous-film formation. Rapid diffusion of the oxidant inward causes precipitation of the solute before it (a much slower diffuser than the oxidant) can diffuse outward. In some respects, the slower diffusivity of an oxidant is analagous to "internal gettering" proposed by Wagner to explain continuousfilm formation on ternary alloys in which the concentration of the first solute is less than the critical value for transition. The second solute, which must have an intermediate affinity for oxygen to that of the first solute and the solvent, forms its oxide which reduces the oxidant level sufficiently so that the first solute may form its own oxide.

REFERENCES 1.

R. P. Rubly and D. L. Douglass, Internal Nitridation of Ni-Cr Alloys, Oxid. Met., 35, (1991) 269.

2

I. Chen and D. L. Douglass, The Internal-Nitriding Behavior of 310 Stainless Steel with and without Al and Ti Additions," Oxid. Met., 34, (1990) 473.

3.

Η. H. Podgurski and Η. E. Knechtel, Trans. Met. Soc. AIME, 245, (1969)1595.

4.

F. H. Stott and G. C. Wood, Mat. Sei. and Tech., 4, (1988)1072.

5.

S. Guruswamy, S. M. Park, J. P. Hirth and R. A. Rapp, Oxid. Met., 25, (1986) 77.

a

R. A. Rapp, Corrosion, 21, (1965) 382.

7.

H. A. Wriedt and 0. D. Gonzalez, Trans., AIME, 221 (1961)532.

a

H. J. Grabke and Ε. M. Peterson, Scripta Met., 12 (1978) 1111.

9.

P. L. Gruzin, Y. A. Polikarpov and G. B. Federov, Fiz. Metal, i Metalloved., 4 (1) (1957) 94.

Fig. 14. Microstructure of scales formed during oxidation of Ni10Cr-5AI and Ni-20Cr05AI for 48 hours at 900'C, showing continuous scales and no internal oxidation.

High Temperature Corrosion of Advanced Materials and Protective Coatings Y. Saito, B. Önay and T. Maruyama (Editors) © 1992 Elsevier Science Publishers B.V. All rights reserved.

147

INTERNAL BROMINE CORROSION OF DILUTE Ni-Cr ALLOYS

BULENT ONAY , YASUTOSHI SAITO

and

TOORU ATAKE

Research L a b o r a t o r y o f Engineering Materials, Tokyo I n s t i t u t e o f Technology, 4259 N a g a t s u t a , M i d o r i - k u , Yokohama 227, Japan. * »Research L a b o r a t o r y f o r Nuclear Reactors, Tokyo Institute o f Technology, 2-12-1 O-okayama, Meguro-ku, T o k y o 152, Japan.

In o r d e r t o g e n e r a t e b o t h b a s i c and e n g i n e e r i n g d a t a o n bromine corrosion of a l l o y s at e l e v a t e d t e m p e r a t u r e s , internal bromidation of Ni alloys containing 1 and 5 mass % Cr was s t u d i e d . A t 1073 K, u n i f o r m r e a c t i o n z o n e s w e r e o b s e r v e d i n Ni-5Cr a l l o y s , whereas g r a i n b o u n d a r y precipitates f o r m e d in N i - l C r a l l o y s . The b r o m i n e p e r m e a b i l i t y p r o d u c t c a l c u l a t e d was in t h e same o r d e r as t h o s e reported f o r f l u o r i n e and c h l o r i n e in t h e l i t e r a t u r e . This r e s u l t suggested that bromine diffuses substitutional^ in N i - C r a l l o y s due to its relatively larger atomic s i z e . The results also provided evidence that engineering a l l o y s used at elevated temperatures in bromine-containing gaseous environments may suffer from intergranular corrosion.

1. INTRODUCTION

Furthermore,

Recently, research

a

University

group

has

thermochemical and

Oxygen

chemical

reactions oxides

and

bromides

containing

bromide

and

must by

resist the

data alloys are

corrosion

the

in

halogen

temperatures.

Because

as of

Fe-base authors'

has

not

that some

in

been

conducted

on

very

*Present address: T e c h n o l o g y , 2-12-1

in

studied

as

a

research

corrosion

has

of

on

been

the

dilute

of

elevated

In

of Ni-

initiated this

Internal

Ni-Cr

in

paper, bromine

alloys

are

presented.

in

EXPERIMENTAL PROCEDURE Internal

conducted

fluorine-

containing

studies

has have

corrosion

alloys.

2.

Although

environments

bromine

Fe-base

alloys

oxygen.

few

behavior at

purposes,

laboratory.

and

temperature

been

alloys

corrosion

are are

interpretation

bromine

obtained

metals

corrosion

alloys

bromine

however,

service

the

the

results

the

to

in

data,

the

these on

and

reliable

engineering

investigated,

and

high

For project

of

caused

be

diffusion

during

reactor

atmospheres

been

Ni-

gas

a

bromine-containing

chlorine-containing

common

alloys

related

Such

for

prediction

the

may

and

and

1

environment,

However,

necessary

Hydrogen

data

available.

iron

engineering

gases

corrosion

that

corrosion

of

extensively

for

solubility not

via

of

required.

corrosion

and

materials

process

for

produce

under

vapor.

a

water

Bromine,

water

constructional

to from

between

mixture

Tokyo

proposed

process

Hydrogen

calcium

of

of

samples

corrosion using 1

was

sealing

in

5

cut

electropolished sample

Ni-Cr

and

were

binary

mass from

quartz

Research Laboratory for Nuclear Reactors, O - o k a y a m a , M e g u r o - k u , T o k y o 152, J a p a n .

with

%

were alloys

Cr.

Small

annealed

sheets.

cleaned a

experiments

and

Each

alloy

acetone

before

tube

Tokyo

with

a

mixture

Institute

of

148

of Ni and

NiBr 2 powders. The

prepared

in

a

high-purity bromides quartz argon

argon

are

tube and

flushed

period

test,

was

filled

because

with metal

hygroscopic. with

The

high-purity

Upon

in a preheated

predetermined

and

Figure

experiments.

2 shows

zone

developed

1073K

after

21

the Ni-NiBro

the in

internal

a

hours

Ni-5Cr of

corrosion alloy

reaction

at with

pack.

for

a

of time. Following

analyses

optical

internal corrosion

sealing, it furnace

microstructural

compositional using

gas

very

was

mixture

box

evacuated.

was placed the

glove

were

and

performed

scanning

by

electron

microscopes. This

"Rhines

pack"-type

was used to generate in

the

tube

bromine alloy

to

arrangement

a bromine

sufficient

react

with

constituent;

the

less

chromium,

forming a chemical c o m p o u n d the base

activity

enough

for noble

without

(scale) with

metal.

3. RESULTS Figure 1 shows a

typical

Ni-Cr

the microstructure alloy

sample

used

of for

the experiments.

FIGURE 2 (A) Microstructure of a Ni-5Cr alloy sample reacted with the Ni-NiBr« mixture for 21 hours at 1073 Κ, (B) ^different area from the same sample. The precipitates were observed non-spherical. to be

higher

Their

close

density

to

surface.

Relatively

continious

precipitates

The sample

Fig. 2, a continious

along visible was

upon

light

single-phase

second-phase

large and

grains etching. did

not

precipitates

which The

were

to

grain

were

metallic

over

any

reaction

with

the

kind of a layer was observed

the

alloy

and

observed

boundaries. As

the

alloy

larger,

observed

alloy

contain

prior

alloy

to be found

the original

FIGURE 1 Microstructure of a Ni-5Cr alloy sample before the internal corrosion experiment. (Etched) had

was

seen

layer

sample

Ni-NiBr2

after

pack. for

in was

This

all the

149

Ni-5Cr samples used

in this study. Parts

of

deformed

this

away

layer

from

before

were

the

the

samples

and

during

microscopical

broke

polishing

observation

(Fig.3). Figure

3 shows

morphology zone

of

the

developed

sample

after

both

the

internal

in a n o t h e r

50

extent

and

corrosion

Ni-5Cr

alloy

hours.

FIGURE 4 Grain boundary precipitate formed in the Ni-lCr sample at 1073 Κ after reaction with the Ni-NiBr2 mixture for 16 hours.

FIGURE 3 Microstructure of a Ni-5Cr alloy reacted with the Ni-NiBr 0 mixture Δ hours at 1073 Κ. The was

morphology

observed

observed times.

for

to

samples

However,

precipitates deep

in

sample.

this

lower

Cr

%),

be

significantly

samples, alloy

large

grain

Figures

4

grain

boundary reaction scans

The

both

obtained

bromine

precipitate

was

at

(Fig.

and

with

found

formed

along

1073

shown of

in the

increased K.

one the

chromium

5 B,C).

to

these

amount

from

of

In

precipitates time

higher

samples

as

larger precipitates s h o w e d of

was

different.

5.

the

original

morphology

zone

precipitates

with line

alloy

boundaries

and

of

the

alloy

that

shorter

amount

the

the internal corrosion

to

for

from

the

of the

(1 mass

precipitates

reacted

the

into

case

the

similar

extending

surface

In the

of be

sample for 50

X-ray of

the

presence in

the

FIGURE 5 ( A ) Grain boundary precipitates formed in the N i - l C r sample at 1073 Κ after reaction with the N i - N i B r 2 mixture for 25 hours. X-ray line scans across a precipitate; ( B ) for Cr, ( C ) for Br.

150

4. DISCUSSION Both

the

compositional samples

morphological analysis

yielded

this

study.

expected

the

reacted

evidence

internal bromidation in

of

of the

Such

a

and

for

the

alloys

used

result

was

thermodynamically.

For

the

reaction

alloy. This

dependency

related

the

to

diffusivity By

=

CrBr 2 (s)

(1)

bromine the zone

a function in

is

given

as:

the

that

necessary

AG°(

= RT lniaçj. . P , ^ )

1)

(2)

the

formation

of pure

activity. Equation to

give

N Cr

CrBr2

(2) can be

at

unit

rearranged

as:

is

f

fraction) of Y£

and

Cr

as

Br^

activity

NiBr2

pack,

about

5.7xl0"

6

alloy

equilibrium by

ppm).

Nçr with

By for

the Ni-

calculated

5.7

with

to react

the

Y C r( 2 . 7 ) ,

and ten

is (

mole

data '

established

Nçr

alloys

expected

in

(in

coefficient.

J/mole) p

taking

Ni-Cr

amount

available

A G ° ( 1 )( - 1 9 8 5 4 7 bromine

the

dissolved

the

and

where

is the activity

using

>

to

discussed 4 -6

internal

a

diffusion-

expressed

which

the

by

the

1 /2

5.7

ppm

are

to

form

higher

than

Therefore,

the

contained

this

critical

bromine-rich precipitates in the

s

_ exp( η

x

φ

b

Ν

δ χ

the

surface,

samples

the

reactive

precipitates

was

found of the

N°

B

fraction

nitrogen,

the

alloy

function

φ = D X/ D B of the

mole

like) dissolved

error

of the

the

1/2 θΓίοίηψ )

φ)

the

complementary

to be dependent on the Cr content

from

(5)

(oxygen,

is thermodynamical ly justified. H o w e v e r , the

conditions, the

calculated

2

is X

the

2-5,

Upon

equation

erf η

β χ Ρ( η

and

Figures

)

1/2 n°

as

in

time.

2 N

of

formation

the

diffusion

diffusion

be

internal

equation

much

study

can

the

dimensionless

oxygen

(t)

Fick's

η

parameter

Cr

this

and

of a

for the proper b o u n d a r y

oxidant

in

^he

the

(4)

depth

(^Q)

be

Therefore,

bromine

solving

alloys

morphology

is

be

(f|)

coefficient

Cr

observed

are

that

zone,

parameter,

where

amount.

the

for

( £ ) is the

oxidation

C r B r 2 precipitates in the alloy. The Niused

and

satisfy

relation

2 Nq

which

obtained

literature.

£ = 2 η (D 0 t )

where

as

system

process can

the

gaseous

be

parameters

alloys

controlled

of

the

may

showed of

kinetics

alloys

time, diffusivity for

alloy-gas

in the

Wagner oxidation

for

depth

alloy the

conditions

extensively

Ni-Cr

be and

in the sample

data

the experimental Ag°q|

in

of reaction

solubility

provided

to

solubility

later.

internal reaction

species B r 2 (g)

low

measuring

and

Çr (s) +

of

as discussed

is believed

very

in the mole

and

where

oxidant

and

alloy

erfc

erf is

function.

Dx

is the

DR

the

the

fraction

component,

error

of

bromine,

is

at of the the

Also,

diffusivity diffusivity

151

FIGURE 6 Concentration profiles for the oxidant, X, and the reactive alloy component, B, for internal reactions (after ref.4). Refer text for details. of

the

reactive

Equation

(5) can

special Figure For

(N

limiting

followed by the laser irradiation.

Electroplated Ni

80

The laser

irradiation is useful to convert the precursor Fe

films and fuse them with the alloy surface, though the power used here was too large to

ίΟ

60

form only the dispersed layer. Fused layer Alloy >
> 40

be distinguishable in the oxidation tests, as the structure and composition of the formed La(x2) Cr

20

layers were similar as described above.

The

performance of the layer was mainly dependent

(Ο

Ni

Β

on the number of imperfections. After the oxidation at 1000°C in a 0.20 atm

Ο

I

0

1 ι

ι

ι

1

ι

ι

50 Distance / μΐη

H 20 - a i r , the surface treated Type 430 specimens showed a very small mass gain and no accelerated oxidation, as shown in Fig. 8.

The mass

gain after 24.5 h is mostly due to an oxide nodFIGURE 6 Composition profiles in the depth direction cross section of Type 304 specimen (by XMA)

ule formed on a corner (Fig. 8(B)), and if this part is excluded the total mass gain is esti2

mated to be less than 2 g / m .

The mass gain was

so small that accurate oxidation curves were not obtained with the electric balance used. The oxide scale is composed of C r 20 3 and spinels (mostly MnCr2Ûi+) .

The apparent parabolic

oxidation rate constant for 8 h oxidized speci5 2

4

men (Fig. 8(A)) is about 3 χ 1 0 " g / ( m s ) whereas the initial rate constant for untreated speci3

2

men is about 4 χ 1 0 ~ g / ( m ^ s ) ^ : untreated Type 430 stainless steel shows accelerated oxidation after 76 min (average) and the mass gain exceeds 20 g/m

2

within 2 h , the formed oxides are FesOi^. 6

and F e C r 20 i + . When L a C r 0 3 has been coated on stainless steels by pyrolysis, decomposition of L a C r 0 3 is not observed by XRD even after prolonged oxidation tests'^, but here the results of XRD indi50 Distance / ym

FIGURE 7 Composition profiles in the depth direction cross section of Type 430 specimen (by XMA)

cate that a part of the dispersed L a C r 0 3 particles decompose slowly (Fig. 8 ) .

This seems un-

usual, since L a C r 0 3 is reported to be stable at least to the oxygen partial pressure, p n , of 7 1Q-I5pa m S 0 f a r> the decomposition mechanism is unknown, but it may occur through intermedi-

182

(1 cm / division) FIGURE 8 Surface photographs and XRD patterns for the surface treated Type 430 specimens after oxidation a t 1000°C in 0.20 atm H 20 - air. ate compounds which probably correspond to unidentified peaks in Fig. 8,

and eventually

L a 20 3

having different crystalline systems are formed. With the coated layer, there was no indication that La ions were in full play as reactive eleg ments , but the L a C r 0 3 layer functioned as a barrier hindering outward diffusion of metal ions and leading to a low oxidation rate"*: this is similar to the spontaneously formed protective C r 20 3 layer, though the L a C r 0 3 layer is much more protective. After 24.5 h oxidation at 1000°C in a 0.20 atm H 20 - air atmosphere, the composition profiles in the depth direction of the laser treated Type 430 specimen changed as shown in Fig. 9.

From

the mass gain, the thickness of the newly formed oxide layer should be no more than 1.5ym, but 50 Distance / μιη

it appears much thicker.

face roughness. FIGURE 9 Composition profiles in the depth direction cross section of the surface treated Type 430 specimen after oxidation a t 0.20 atm H 20 - air for 24.5 h (by XMA)

This must be attrib-

uted to the low resolution of XMA and the surIt can be observed that Cr is

concentrated in the outermost part and that it is followed by Cr depleted zone.

Small amounts

of Si and Mn also concentrate in the outermost

183

part.

There is a random Cr enrichment coupled

the additional effects of L a 20 3 particles form-

with Fe deficiency, indicating that internal

ed by the decomposition.

grain boundary oxidation is taking place.

however, reduces electric conductivity.

These

The decomposition, If con-

concentration profiles and oxidation behaviour

version of the precursor and sintering can be

are similar to those for the surface coated and

carried out at the very surface, the formed

oxidized specimens^.

L a C r 0 3 layer will not decompose during oxida-

The dispersion of L a C r 0 3

particles into the surface layer, however, is

tion.

more effective to improve oxidation resistance

tion condition used here was too powerful.

In this aspect, again, the laser irradia-

than surface coatings: this may be ascribed to

ditions which enable L a C r 0 3 to sinter at the

Con-

surface must be established. The results of cyclic oxidation at 1100°C for 60 min and cooling to room temperature in air are shown in Fig. 10.

The photographs were

taken after the final cycle for each specimen. With Type 304 specimens, excellent oxidation resistance was observed as shown in the photograph (A) which is well contrasted with the violent spalling of the blanV: specimen (B) . 1—I—I—I—I

I

I

I

I

I

Γ

The

final mass gain after 10 cycles was 3.2 g/m

2

(D)

which corresponds to the apparent parabolic oxi-

Type 430 blank

dation rate constant of 2.8 χ 1 0

-t L

2

g /(m^s) : this

value is extremely low for Type 304 stainless steel, even though the mass loss due to the sublimation of chromium species must be taken into account.

With Type 430 specimens, the mass gain

was much smaller than with the blank specimen, but fairly large due to pinhole imperfections at the surface and on the sides, as can be seen in

Type 430 treated

the photograph (C).

ß (C)

(A)

4. CONCLUDING REMARKS

Type 304 treated

The L a C r 0 3 particles dispersed alloy layers are formed by electrodeposition of LaOHCrO^

Type 304 blank

ηH20

films followed by laser irradiation, which enables the precursor films to convert to L a C r 0 3

Ο -150 (Β) J

I

I

I

I

I

5

I

I

L

10

Number of cycles

particles and to fuse with the alloy surface. It is obvious that the L a C r 0 3 particle dispersed alloy layer is very effective to improve oxidation resistance of alloys.

FIGURE 10 Results of cyclic oxidation test a t 1100°C for 6 0 m i η and cooling to room temperature in air (Photographs were taken after the final cycle for each specimen)

For high temperature

oxidation tests, however, the whole area of the specimen must be perfectly treated, since a single imperfection leads to large mass gains after prolonged oxidation.

A high yield of per-

184

Singhal (Electrochemical Society, Pennington, N. J., 1989).

fectly treated specimens has not been achieved by the present method.

This is mainly due to

(i) the difficulty of laser irradiation of the sides of thin specimens, and (iL) the ununiform-

2. H. Konno, M. Tokita, S. Kitazaki and R. Furuichi, J. Surf. Finish. Soc. Jpn., 40 (1989) 825.

ity of the precursor film created during handling and laser irradiation.

We believe it possi-

3. H. Konno, M. Tokita and R. Furuichi, J. Electrochem. S o c , 137 (1990) 361.

ble to overcome these technical problems. 4. M. Tokita, H. Konno and R. Furuichi, Ext. Abs. 82nd Ann. Meeting Surf. Finish. Soc. Jpn., (1990) 272.

ACKNOWLEDGEMENT A part of this work was carried out under the Visiting Researcher's Programme of the Institute for Materials Research, Tohoku University.

The

authors wish to express their appreciation to Prof. K. Hashimoto and Mr. T. Sato for their cooperation in the laser experiments.

5. H. Konno, S. Kitazaki and R. Furuichi, Boshoku Gijutsu (Corr. Engr.), 39 (1990) 544. 6. H. Konno, I. Saeki and R. Furuichi, Boshoku Gijutsu (Corr. Engr.), 37 (1988) 338. 7. T. Nakamura, G. Petzow and L. J. Gauckler, Mat. Res. Bull., 14 (1979) 649. 8.

REFERENCES 1.

Solid Oxide Fuel Cells, ed. S. C.

D. P. Whittle and J. Stringer, Phil. Trans. R. London, A295 (1980) 309.

HOT CORROSION OF ENGINEERING ALLOYS AND CORROSION OF NUCLEAR ENERGY - RELATED MATERIALS

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High Temperature Corrosion of Advanced Materials and Protective Coatings Y. Saito, B. Önay and T. Maruyama (Editors) © 1992 Elsevier Science Publishers B.V. All rights reserved.

187

THE ROLE OF APPLIED CREEP STRESS ON HOT CORROSION BEHAVIOR OF A NICKEL-BASE SUPERALLOY

Masayuki YOSHIBA Department of Mechanical Engineering, Faculty of Technology, Tokyo Metropolitan University 2 - 1 - 1 Fukazawa, Setagaya-ku, Tokyo 1 5 8 , Japan

In order to clarify the role of an applied creep stress on the hot corrosion behavior of nickelbase superalloys, the creep rupture tests were performed of the wrought nickel-base alloy 751 at 800°C in the hot corrosive environment simulated both by coating with a given amount of 9 0 % N a 2S 0 i + 10%NaCl synthetic salt mixture and by controlling the oxygen partial pressure ( p 0 2) of gaseous atmosphere in a wide range from 0 to 1 atm. The metallographic surface loss as a measure of a general corrosion was found to increase in accordance with a quasi-parabolic kinetics, independently of an applied stress. On the contrary, a preferential intergranular attack was much stimThen the simultaneous ulated under the creep stress, particularly in the higher p 0 2 atmospheres. action of an applied stress and the hot corrosion associated with high p 0 2 atmosphere tends to bring about the most significant reduction of the creep rupture life. From the metallographic examinations using a X-ray microanalysis, it was revealed that both a general corrosion and an intergranular attack propagate through an essentially similar reaction process of a successive sulfidation-oxidation mechanism. Furthermore an enhanced corrosion rate in the higher p 0 2 atmospheres should be attributed to a rapid propagation of the Ni-rich sulfides drived into an alloy interior by the following oxidation. An applied creep stress should play a primarily important role in developing an active short circuit path such as a microchannel or a crack along the grain boundary across the thick surface scale layer so as for a substantial amount of the corrosive species such as S and 0 to be introduced directly from the molten salt environment into an alloy interior by the manner different from a simple grain boundary diffusion.

1.

INTRODUCTION

degradation has been found to become more seri-

High temperature materials such as nickel-

ous in general in the case associated with a

base superalloys are usually subjected to the

localized attack, for instance with an inter-

simultaneous effect of both thermomechanical

granular penetration of sulfides followed by

damages due to creep and/or fatigue and thermo-

oxides, rather than a general corrosion-dominat-

chemical damages mainly due to hot corrosion.

ing case, depending strongly on the environ-

Thus it has been often pointed out that from a

mental conditions.

practical viewpoint the high temperature strength

Inversely, an applied stress also seems to

properties should be appropriately evaluated in

affect the hot corrosion behavior in the aspects

such aggressive environments together with in

of corrosion kinetics and morphologies, which in

the basic environments such as air, since an

turn brings about a change both in the lifetime,

interaction between the mechanical stress and

with a general importance for the life shorten-

corrosive environment is able to become a life

ing, and in the failure mode of the components.

limiting factor for the hot section components

However its phenomenology is hardly clarified

of various heat engines such as gas t u r b i n e s

1 - 6

.

consistently because of the variety and complex-

In fact many numbers of studies including a se-

ity of the affecting factors to be considered,

ries of the authors studies have been reported,

although several studies have been recently

referring to a significance of the corrosion-

carried o u t »

induced strength degradation, particularly in the 7

11

creep rupture p r o p e r t i e s " .

According to

these studies, a corrosion-induced

strength

12

1 3

.

In the present study, both kinetics and morphologies of hot corrosion were investigated mainly from a microstructural viewpoint for a

188

nickel-base superalloy simultaneously

sustaining

Table 1

Chemical composition of Inconel 751. (mass %)

a mechanical creep stress, and the role of the applied stress in the hot corrosion behavior was discussed.

C Si Mn Ni Cr Ti AI Fe Cu Nb+Ta 0.04 0.17 0.17 Bal. 15.42 2.14 1.03 7.87 0.11 0.99

2. MATERIAL AND EXPERIMENTAL PROCEDURES The wrought nickel-base superalloy Inconel 751 was used in this study. try is shown in Table 1.

This alloy chemisThis alloy is known

After the corrosion test, the metallographic examinations were made for both the cross- and longitudinal-sections of the ruptured specimens by means of an optical microscopy, a scanning

to be poor in the hot corrosion resistance be-

electron microscopy, and a X-ray microanalysis

cause of an insufficient Cr content of approxi-

(ΕΡΜΑ) with the careful preparation of specimens

mately 1 5 %

l i +1 6

~ .

Heat treatment adopted was

10

for an a n a l y s i s .

the simple solution and aging treatment as follows :

3. RESULTS

1200°Cx2h+WQ + 750°Cx24h+AC.

3.1. Hot corrosion kinetics

The smooth bar specimens with 5mm in a dia-

Figure 1 shows the hot corrosion kinetics in

meter and 30mm in a gage length were machined

terms of a surface loss at 800°C under different

from the heat-treated rods, and were emery-

levels of the creep stresses.

polished through 500 grit.

was determined from a metallographic measurement

Thereafter they

were cleaned ultrasonically in aceton. In order to simulate an actual hot corrosive

A surface loss

at the cross-sections of the ruptured specimens 9/

1 7

, and the testing time corresponds to the

environment in this laboratory test, the speci-

rupture time.

mens were coated with a given amount of the

with a creep deformation was negligible in this

A reduction of area associated

synthetic salt mixture composed of 90%Na2S0^-10%

study, since the tested alloy has ruptured in a

NaCl.

markedly brittle manner in hot corrosive environ-

The amount of salt precoated was 40 2

mg/cm , and the same amount of salt mixture was

ment with a rupture elongation less than 3% at

repeatedly coated at every 200h to keep a cor-

most.

rosive action for the prolonged test duration.

fairly large scatter bands, it can be seen that

Gas atmospheric condition also was controlled

Although the kinetics data are in the

a general corrosion tends to proceed in

by using different compositions of N2-O2 gas mixtures including pure N2 and O2 gases, because

S t r e s s X P, o

(MPa)\(atm)

it should be one of the important environment

250 200 150 100

factors affecting the hot corrosion behavior. Then the oxygen partial pressure (pÛ2) of gas

0 0.05 0.2

1

Ο Φ

m

φ-

•

•

3

Œ

Ψ τ

P o 2= 1 a t m

τ •

atmosphere was varied widely from 0 to 1 atm, under a given total pressure 1 atm, as follows: p 0 2 = 0, 0.05, 0.2, 1 atm A creep load was applied on the specimens

0.1 P o 2< 0 . 0 5 a t m

coated with salt mixture in the stress range of 100 to 250MPa, using the single lever type creep testing machines.

0.03

The hot corrosion test under

the applied stress was carried out at 800°C until the specimen is ruptured.

10

100 Time (h)

Fig. 1

1000

Hot corrosion kinetics in terms of a surface loss at 800°C.

189

accordance with a q u a s i - p a r a b o l i c r a t e l a w . Such a l a r g e data s c a t t e r

appears t o be a t t r i b -

uted p a r t l y t o the ununiformity of a g e n e r a l c o r r o s i o n through the specimen gage l e n g t h , d e pending on the s i t e of s e c t i o n i n g .

Furthermore

the c o r r o s i o n r a t e i s found t o depend on the pCh of the gas atmosphere; i n c r e a s i n g in the p Û 2 r e s u l t s in an enhanced surface l o s s , of an a p p l i e d s t r e s s .

regardless

Fig.

2

However the prolonged

c o r r o s i o n t e s t s a s s o c i a t e d w i t h bearing the r e l a t i v e l y low s t r e s s r e s u l t e d i n a s i m i l a r e x t e n t of surface l o s s even in d i f f e r e n t p Û 2 of

Micrograph o f a p r e f e r e n t i a l i n t e r granular a t t a c k observed i n the specimen creep-ruptured f o r 21.7h under an a p p l i e d s t r e s s 200MPa i n the hot c o r r o s i v e environment w i t h 1 atm p0 2 gas atmosphere. Stress axis i s v e r t i c a l .

the gas atmosphere, probably due t o a consumption of the molten s a l t .

So f a r as a surface

loss

i s concerned, anyhow, i t can be said t h a t i t hardly promoted by an a p p l i e d creep s t r e s s , it

is since

a l s o has been confirmed i n the p r e v i o u s study

by the author using the same s a l t mixture as i n t h i s study but heating i n a i r t h a t a surface

Type A

loss

was almost e q u i v a l e n t whether a creep s t r e s s was

Fig.

3

9

a p p l i e d or n o t . 3.2.

Feature o f a g g r e s s i v e i n t e r g r a n u l a r attack

I t has been found by a number of s t u d i e s

that

Table 2

s t r e n g t h degradation i s a s s o c i a t e d i n many cases granular a t t a c k s

7 - 1 1

.

In p a r t i c u l a r

interi t has been

that only a l i m i t e d number of such an

(MPa)\(atm)

inter-

granular a t t a c k can propagate so r a p i d l y t h a t the most predominant one i s capable o f causing d i r e c t l y a premature f r a c t u r e ner

9 - 1 1

.

C l a s s i f i c a t i o n of t h r e e t y p e s of creep rupture morphologies observed i n hot c o r r o s i v e environments w i t h d i f f e r e n t p0 2 of gas atmospheres. Numerals in the parentheses i n d i c a t e the a v e r a g e v a l u e s of the rupture l i f e i n hour.

Stress\Po2

r e v e a l e d from the p r e v i o u s s t u d i e s by the author

Type C

Schematic drawing of t h r e e t y p e s of the creep rupture morphologies observed i n the hot c o r r o s i v e environment w i t h d i f f e r e n t p 0 2 gas atmospheres (from ref. 11).

a c o r r o s i o n - i n d u c e d s e r i o u s creep rupture with an occurrence of the p r e f e r e n t i a l

Type Β

0

0.05

0.2

1 C(14)

200

A (93)

Β (40)

C(37)

150 100

A(161)

C(91)

C(93)

C(28)

A (193)

C(265)

C(218)

C(219)

in a b r i t t l e man-

Furthermore the g r a i n

boundaries

rupture morphology was c l a s s i f i e d i n t o

three

undergone such an a g g r e s s i v e a t t a c k i s l i m i t e d t o

t y p e s , as i l l u s t r a t e d

those l y i n g roughly normal t o the s t r e s s a x i s ,

Type A suggests t h a t an i n t e r g r a n u l a r

as shown i n F i g . 2 .

hardly promoted by an a p p l i e d s t r e s s because

However i t i s very

c u l t t o obtain the l o n g i t u d i n a l s e c t i o n

diffisuitable

f o r a q u a n t i t a t i v e measurement of such an a g g r e s sive intergranular imen.

a t t a c k from the ruptured s p e c -

In order t o e v a l u a t e q u a l i t a t i v e l y a s i g -

n i f i c a n c e of an a g g r e s s i v e i n t e r g r a n u l a r in t h i s study,

attack

c o n s e q u e n t l y , the c o r r o s i o n creep

schematically in F i g . 3 attack

1 1

.

is

the creep rupture morphology i s e s s e n t i a l l y s i m i l a r t o i n the n o n - a g g r e s s i v e environment such as i n a i r .

On the c o n t r a r y ,

suggests f o r an i n t e r g r a n u l a r

type C

attack

t o be s i g -

n i f i c a n t l y enhanced by an a p p l i e d s t r e s s so that it

causes d i r e c t l y a premature f r a c t u r e

prior to

190

an initiation of the creep-induced internal grain boundary cracks.

Type Β also is essentially

sectional surface zone of the specimens undergone a general corrosion in the different p 0 2 gas at-

similar to type C, which suggests a fairly

mospheres of 0, 0.05 and 1 atm, respectively.

stress-enhanced intergranular attack.

The result of Al is eliminated in these figures

The result of the metallographic observation on the longitudinal section of the specimen rup-

in spite of an analysis carried out, since only 1% of Al was found to occupy no important posi-

tured in hot corrosive environment is summarized

tion in hot corrosion behavior.

in Table 2, in associating with the p 0 2 of gas

from Figs. 4-6 that a general corrosion proceeds

atmosphere and an applied stress level.

It is

It can be seen

by almost similar process regardless of the p 0 2

clear that an aggressive intergranular attack is

of gas atmospheres; a sulfide formation precedes

apt to play more important role in the 0 2-bearing

an oxide formation.

atmospheres rather than in the 0 2-free (p0 2= 0

an image analysis using the ΕΡΜΑ color mapping

atm) one.

Furthermore, a combination of high

p 0 2 atmosphere and high creep stress is found to

It has been confirmed by

that Ni sulfides; perhaps the low melting NiN i 3S 2 eutectic, were formed at the front of the

result in the most significant reduction of the

general corrosion, although they tend to be con-

creep rupture life.

verted quickly to Cr sulfides probably due to a

This suggests for an aggres-

18

sive intergranular attack to be significantly

thermodynamical s t a b i l i t y .

enhanced by an applied stress so as to cause

a potential for the oxide scale formation is

directly a premature fracture.

strongly dependent on the p 0 2 of the gas atmos-

3.3. Morphology and thermochemistry of hot corrosion Figures 4-6 show both the backscattered electron and characteristic X-ray images at the cross-

Fig. 4

On the other hand,

phere; increasing in the p 0 2 leads to the more intensive and thick oxide scale formation toward the internal sulfide layer.

Furthermore, a

nodule-like scale morphology is characteristic

Backscattered electron and characteristic X-ray images at the cross-sectional surface zone of the specimen undergone a general corrosion in the hot corrosive environment with the 0 atm p 0 2 gas atmosphere. (applied stress :250MPa, time to rupture : 24.2h)

191

'ig. 5

Backscattered electron and characteristic X-ray images at the cross-sectional surface zone of the specimen undergone a general corrosion in the hot corrosive environment with the 0.05 atm p 0 2 gas atmosphere. (applied stress :200MPa, time to rupture : 39.6h)

'ig. 6

Backscattered electron and characteristic X-ray images at the cross-sectional surface zone of the specimen undergone a general corrosion in the hot corrosive environment with the 1 atm p 0 2 gas atmosphere. (applied stress :150MPa, time to rupture : 12.3h)

f the hot corrosive environment with relatively

ward along the alloy surface to form a continu-

ow p 0 2 gas atmosphere; for instance of 0.05 atm

ous scale layer.

s shown in Fig. 5, which appears to grow after-

Figures 7 and 8 show both the secondary

192

electron and characteristic X-ray images at the

sion.

Since an image analysis has revealed for

tip portion of an aggressive intergranular attack

the preceding sulfides to be associated with Ni,

developed in the specimen undergone a hot corro-

a penetration of Ni sulfides should always pre-

sion in the 0.05 and 1 atm p 0 2 of gas atmospheres,

dominate the intergranular attack process.

respectively.

was also found that an occurrence of such an ag-

No aggressive intergranular at-

tack was observed for the specimen corroded in the 0 atm p 0 2 gas atmosphere.

It can be seen

It

gressive intergranular attack is limited only beneath the scale grown sufficiently, including

that such an aggressive intergranular attack also

a nodule-like scale as in the case of the rela-

propagates by a successive

tively low p 0 2 gas atmosphere.

sulfidation-oxidation

process similar to the case of a general corro-

Fig. 7

Secondary electron and characteristic X-ray images at the tip of an aggressive intergranular attack developed in the specimen undergone a hot corrosion in the 0.05 atm p 0 2 gas atmosphere, (applied stress :200MPa, time to rupture : 39.6h, stress axis is horizontal)

Fig. 8

Secondary electron and characteristic X-ray images at the tip of an aggressive intergranular attack developed in the specimen undergone a hot corrosion in the 1 atm p 0 2 gas atmosphere, (applied stress :150MPa, time to rupture : 12.3h, stress axis is horizontal)

193

11

the Ni s u l f i d e s .

4. DISCUSSION It was shown that hot corrosion principally through a successive

An applied creep stress was found to affect

proceeds

sulfidation-

significantly on the aggressive

intergranular

oxidation process whichever forms of a general

attack behavior.

corrosion and an intergranular attack.

role in the two stages of intergranular attack

This

situation can be illustrated schematically shown in Fig. 9.

as

Such a successive sulfidation-

process.

It should play an important

The first is at an initiation process

of the predominant intergranular attack capable

oxidation process seems to be accounted for by a

of causing directly a premature fracture,

modification of the early sulfidation model pro-

corresponding to at an onset of the rapid propa-

posed by Simons et a l .

1 9

, as also shown in Fig.

gation of it.

In the present study, the pre-

9, although for a general corrosion the other

dominant intergranular attack showed a morpholo-

mechanisms such as an electrochemical

gy characteristic of sustaining a creep stress,

20

may be also o p e r a t i v e .

reaction

Furthermore the rather

in particular at the portion of scale layer.

enhanced corrosion rate in hot corrosive environ-

Figure 10 shows a typical microstructure of the

ment with the increased p 0 2 gas atmosphere should

predominant intergranular attack together with

become reasonable by taking into account both of

the surface scale layer.

a difference in the penetration (diffusion) rate

a grain boundary crack is formed across the

It can be seen that

between S and 0, and of a driving effect by the

scale layer to provide a microchannel

following oxidation on the inward propagation of

between a molten salt environment and the alloy

N 2- 0 2

connecting

Atmosphere

Proposed "Sulfidation-Oxidation" Reaction Process Molten Salt

N a 2 S 0 4 - N a 2 0 + S 0 3 (or S 0 2 + 1 / 2 0 2) - * N a 20 + S + 3 0

Oxide Layer 3Ni +

Sulfide Layer

_l

ll Ni +

I

2 S - N i 3S 2 N i 3S 2— N i - N i 3S 2( e u t e c t i c )

N i - N i 3S 2 + 2 x C r - 2 C r x S + 4 N i

t

—

i

=

=

i

together with 0 2 " from atomosphere

—

2 C r xS + 3 x 0 ^ x C r 2 0 3 + 2 S

D°: Ο Volume Diffusion, Dg b: Ο Grain Boundary Diffusion s D : S Volume Diffusion, D | b: S Grain Boundary Diffusion Fig. 9

Schematic illustration showing a typical corrosion morphology together with the corrosion products associated, the affecting factors and the proposed reaction process in hot corrosion behavior of nickel-base superalloy under a creep stress.

194

a) by simple grain boundary diffusion Fig. 10

Scanning electron micrograph of the cross-sectional surface scale layer with a short circuit path for an aggressive intergranular attack. ( p 0 2: 1 atm, applied stress :150MPa, time to rupture : 12.3h, stress axis is horizontal)

Fig. 11

b) by direct penetration

;

( in the present study ) Schematic illustration showing a difference in the morphology of the intergranular attack developed by two ways of processes.

an aggressive penetration also takes place by the mechanism essentially different from a

interior.

Therefore it is reasonable to con-

simple grain boundary diffusion.

Provided an

sider that a substantial amount of the corrosive

intergranular attack occurs mainly by a grain

species such as S and 0 can be introduced direct-

boundary diffusion mechanism, the rather network-

ly and so rapidly through such a short circuit

like corrosion morphology must be developed, by

path by the mechanism different from a simple

virtue of only a chemical concentration gradient

grain boundary diffusion, as shown schematically

of the corrosive species as an effective driving

in Fig. 11.

force .

An applied stress should be capable

of causing such a grain boundary crack in the

16

Then a stress-enhanced intergranular

attack should be attributed to some thermomechan-

scale layer by some ways; for example either by

ically induced mechanism.

a brittle cleavage fracture or by a rather duc-

be possible for the grain boundary microcracks

tile grain boundary sliding, although it remains

to be introduced by a grain boundary sliding so

unclear which mechanism is dominating.

that along such a defective region the corrosive

The second is at a propagation process of such a predominant intergranular attack.

As already

For instance, it may

species are able to penetrate rapidly into an alloy interior.

However this is unrealized in

shown in Fig. 2, the morphology of the predomi-

the present study by the fact that an aggressive

nant intergranular attack is characterized by

intergranular attack has propagated too rapidly

having an extremely preferential orientation for

prior to reaching a steady-state creep stage to

a rapid propagation; in particular favoring the

cause the creep-induced m i c r o c r a c k s .

grain boundary in the nearly normal direction

it may be reasonable to consider that the grain

with respect to the stress axis, in a macroscopic

boundary dislocations induced by the creep defor-

scale.

mation can provide the penetration path available

This suggests that a propagation of such

11

Instead,

195

for the corrosive species, being associated with

study was competently carried out by M.Eng. T.

a sort of pipe diffusion, although its certainty

Mizoguchi, Messrs. H. Enomoto, F. Nakao and H.

1 2 2 1

has to be p r o v e d ' .

Anyhow, an applied

stress should be essential to both a premature

Tanabe, who were formerly undergraduates at Tokyo Metropolitan University.

initiation and a rapid propagation of the aggressive intergranular attack.

5. CONCLUDING REMARKS (1) A general corrosion was found to be hardly enhanced by an applied creep stress, but is increased its rate as the p02 of gas atmosphere becomes higher in hot corrosive environment. It increased in general in accordance with a

REFERENCES 1. P. Hancock, Summary of the Mechanisms of Hot Corrosion in Marine Gas Tubines and the Role of Scale Failure, in: Proc. 1974 Gas Tubine Materials in the Marine Environment Conference, eds. J. W. Fairbanks and I. Machlin (MCIC 75-27, Columbus, 1975), pp.225-236. 2. V. Guttmann and M. Merz (eds), Corrosion and Mechanical Stress at High Temperatures (Elsevier, London, 1981).

quasi-parabolic kinetics. (2) Aggressive intergranular attack was found to be significantly enhanced not only in the higher p02 of gas atmosphere, but also under the applied creep stress.

3. H. W. Grünling, W. Hartnagel, R. Bürgel and R. Bauer, Hot Corrosion Effects on Creep and Fatigue, in: Proc. 9th International Congress on Metallic Corrosion, Vol.2 (NRCC, Ottawa, 1984), pp.54-63.

Then a combination of the

higher p02 atmosphere and higher creep stress resulted in the most serious reduction of the creep rupture life by virtue of the most enhanced intergranular attack. (3) Both a general and an intergranular attacks were confirmed to proceed by the almost similar reaction process of a successive sulfidation-oxidation accompanied by the low melting Ni sulfide formation at the corrosion front. (4) An applied creep stress is essential to develop a microchannel as a short circuit path across the thick scale layer so as for a substantial amount of the corrosive species such as S and 0 to be introduced directly and rapidly along there. (5) An applied stress appears to be also effective to enhance the propagation of such an

4. V. Guttmann and M. Schutze, Interaction of Corrosion and Mechanical Properties, in: High Temperature Alloys for Gas Turbines and Other Applications 1986, eds. W. Betz, R. Brunetaud et al. (D. Reidel Publ., London, 1986), pp.293-326. 5. B. F. Dyson and S. Osgerby, Mater. Sei. and Technol., 3 (1987), 545. 6. M. Schutze, Mater. Sei. and Eng., A121 (1989), 563. 7. H. Huff and F. Schreiber, Werkstoffe Korrosion, 23 (1972), 370. 8. Kh. G. Schmitt-Thomas, H. Meisel and H. J. Dorn, Werkstoffe Korrosion, 29 (1978), 1. 9. M. Yoshiba, 0. Miyagawa, T. Sakaki and H. Fujishiro, J. Iron Steel Inst. Japan, 68 (1982), 120. 10. M. Yoshiba, 0. Miyagawa, H. Mizuno and H. Fujishiro, Trans. Japan Inst. Metals, 29 (1988), 26.

aggressive intergranular attack at the rate more than that of a simple grain boundary diffusion.

ACKNOWLEDGMENTS The author wish to acknowledge Mr. T. Masaki of Shimadzu Corp. for the valuable analyses by ΕΡΜΑ.

The experimental work in the present

11. M. Yoshiba and 0. Miyagawa, Environmental Effects of Hot Corrosion on the Creep Rupture Properties of Nickel-Base Superalloys, in: High Temperature Materials for Power Engineering 1990, Part II, eds. E. Bachelet et al. (Kluwer Academic Publ., Dordrecht, 1990), pp.1215-1224. 12. J. K. Solberg and H. Thon, Met. Trans., 14A (1983), 1213.

196

13. Β. Pieraggi, Mater. Sei. and Eng., 88 (1987), 199. 14. R. Viswanathan, Corrosion, 24 (1968), 359. 15. G. C. Clark, Sulfidation Corrosion of NickelBase Exhaust Valves, SAE Paper 750044 (SAE, Warrendale, 1975). 16. M. Yoshiba, 0. Miyagawa and H. Fujishiro, J. Iron Steel Inst. Japan, 67 (1981), 996. 17. P. A. Bergman, C. T. Sims and A. N. Beltran, Development of Hot-Corrosion-Resistant Alloys for Marine Gas Turbine Service, in: Hot Corrosion Problems Associated with Gas Turbines, ASTM STP 421 (ASTM, Philadelphia, 1967), pp.38-60.

18. G. J. Danek, Naval Eng., 77 (1965), 859. 19. E. L. Simons, G. V. Browning and H. A. Liebhafsky, Corrosion, 11 (1955), 505t. 20. F. S. Pettit and C. S. Giggins, Hot Corrosion, in: Superalloys II, eds. C. T. Sims, N. S. Stoloff and W. C. Hagel (J. Wiley & Sons., New York, 1987), pp.327-358. 21. B. Ralph, Grain Boundaries in Engineering Materials, in: Grain Boundary Structure and Kinetics (ASM, Ohio, 1980), pp.181-208.

High Temperature Corrosion of Advanced Materials and Protective Coatings Y . Saito, B. Önay and T. Maruyama (Editors) © 1992 Elsevier Science Publishers B.V. All rights reserved.

197

CORROSION RESISTANCE O F ADVANCED TUBE MATERIALS IN COAL-FIRED

BOILERS

S.Kihara*, K.Nakagawa*, W.Wolowodiukt, J.L.Blought, and W.T.Bakker§

To select alloy f o r superheater tube of advanced coal fired boilers, laboratory coal-ash corrosion tests were done for candidate alloys and coatings. Following conclusions were obtained; (1) T h e corrosion resistance is significantly improved when the steels contain more than 2 0 % Cr. Some highly alloyed austenitic steels have both of excellent high temperature strength and high coal-ash corrosion resistance, and are promising f o r superheater tube of advanced boilers. (2) Isocorrosion diagrams, which show corrosion rates as a function of % SÛ£ in gas and % alkali sulfate in ash deposit, c a n be used to estimate the approximate corrosion rate for each alloy. 1. INTRODUCTION Under fossil

program.

the recent fuel

difficult

supply

preservation,

situations of

and

the trend

environment

in

fossil

power

tube

high

efficient

units

and

conditions:

operated

at super-high

temperatures

(ex. steam

31 M P a , 593°C)

surpassing 1

present levels (ex. 16.6 M P a , δ Β β ^ ) . » To

realize

the advanced

units,

the

2

paper

alloys

done

for the EPRI

oxidation

accompanied

with

ash

deposited

o n the tube

corrosivity

of e n v i r o n m e n t

quality

t h e coal

are

required

pressure

and temperature

for superheater

a n d reheater

1985 the Electric

Institute program

of

laboratory

(EPRI) (RP1403)

embarked 3

Power on an

extensive

from

Previous

operating

at

advanced

steam

conditions.

fireside

like

corrosion

of

superheater and

reheater areas of coal-fired boilers, extensive of materials

resistant

to

coal-ash

corrosion is being conducted, as part of this

kinds

in

The corrosion

sulfates

in coal 6 8

surface. " depends

being tests

were

The

o n the

burned.

of coals

studies high

have

have

strength

17Cr-14Ni-Cu-Mo

steels

Since coal-ash corrosion is most significant

testing

corrosion

various

conventional

in

sulfidation

done

The under

to be used in

actual boilers.

components for improved coal-fired power plants of

to develop

Research

materials a n d

capable

are

conditions simulating the atmospheres resulting

tubes. In

4

is mainly governed by the concentrations of SO2

to

higher

program

molten alkali iron trisulfates.^ in combustion gas and alkali

materials

of

Coal-ash corrosion is considered catastrophic

which exhibit strength and corrosion resistance withstand

the results

mainly reported.

generation technology is toward the adoption of pressure

In this

laboratory corrosion tests for various kinds of

a

indicated austenitic

that steels

and 16Cr-12Ni-Mn-Mo-Nb-V

poor

coal-ash

corrosion

resistance and high corrosion resistance steels like

25Cr-20Ni

temperature

steel

strength.^

have

a

poor

Therefore,

high the

utilization of co-extruded tubes (high strength tube

cladded

* Ishikawajima-Harima Heavy Industries Co., Ltd. Tokyo, Japan t Foster Wheeler Development Corporation, Livingston, NJ § Electric Power Reserarch Institute, Palo Alto, CA

with

high

corrosion

resistant

198

alloy) been

and coated considered,

(chromized

etc.) tubes has

as

as

well

the

Ash composition:

alloys

Alkali sulfate: 2.5, 5, and 10 w t %

specially developed for advanced boilers.

( N a 2S 0 4/ K 2S 0 4= l : l ) Balance:

2. EXPERIMENTAL

PROCEDURE

2.1 Materials Tested Three

groops

S 0 2: 0.05, 0.1, 0.25, and 1.0 v o l %

of materials

were

tested;

tube alloy (can be used as a stress 2) cladding material

alloy

for

(can be used

co-extruded

(chromizing).

lMo steel) were

0 2: 2 and 4 v o l %

component)

3)

(347H

steel

selected

1)

coating

stainless)

2

The

amount of ash coated was 40 m g / c m .

gas

flow

to each

and T91 (Mod.9Cr-

the

as base material for

100

hrs exposure,

chemically

cladding

descaled.

alloys

tested

here

the test

The

one test After

coupon

The descaling

was

treatment

Primary; Boiling in 18% Na0H+3% K M n 0 4 solution

are shown in

Secondary; Boiling in

Table 1.

10%

Ammonia

Citrate

solution

2.2 Test Procedure coupons

synthetic

2.3 Evaluation of corrosion

(15*25x3

ash were

mm) coated

placed

in

a

with the

report

at the test temperatures.

The test

the

weight

variables

occurred

under

materials

Metal temperatures: 600, 650, 700, and 750°C

^^\Element

descaling

all

tested,

from

the

Since general

test

conditions

the weight

represents the corrosion loss.

initial

corrosion

loss

for all properly

To estimate the

Identification and chemical composition of alloys tested

C

Si

18-12-Nb (347H)

0.08

17-14-Cu-Mo

0.11

16-12-Mn-Mo-Nb-V 22-35 (Alloy 800H)

Mn

Ρ

S

0.49

1.48

0.025

0.45

0.75

0.025

0.11

0.41

6.00

0.08

0.55

1.17

Fe

Ti

Al

Nb

Co

V

-

-

0.78

-

-

Bal

15.5

2.41

0.21

-

0.41

-

Bal

12.2

16.1

1.19

-

-

0.87

-

0.22

Bal

-

32.7

21.1

0.50

0.52

-

-

-

Bal

0.49

-

-

Cr

0.006

-

12.6

18.4

0.004

2.97

14.0

0.021

0.005

-

0.019

0.001

22-35-Mo-Nb

0.06

0.52

1.20

0.004

0.001

-

35.7

22.8

1.22

-

-