Science of Ceramic Interfaces II

SCIENCE OF CERAMIC INTERFACES II

ELSEVIER SCIENCE B.V. Sara Burgerhartstraat 25 P.O. Box 211,1000 AE Amsterdam, The Netherlands

ISBN: 0-444-81666-6 © 1994 Elsevier Science B.V. All rights reserved. No part of this publication may be reproduced, stored in a retrieval system or transmitted in any form or by any means, electronic, mechanical, photocopying, recording or otherwise, without the prior written permission of the publisher, Elsevier Science B.V., Copyright & Permissions Department, P.O. Box 521, 1000 AM Amsterdam, The Netherlands. Special regulations for readers in the USA - This publication has been registered with the Copyright Clearance Center Inc. (CCC), Salem, Massachusetts. Information can be obtained from the CCC about conditions under which photocopies of parts of this publication may be made in the USA. All other copyright questions, including photocopying outside of the USA, should be referred to the copyright owner, Elsevier Science B.V., unless otherwise specified. No responsibility is assumed by the publisher for any injury and/or damage to persons or property as a matter of products liability, negligence or otherwise, or from any use or operation of any methods, products, instructions or ideas contained in the material herein. This book is printed on acid-free paper. Printed in The Netherlands

MATERIALS SCIENCE MONOGRAPHS, 81

SCIENCE OF CERAMIC INTERFACES I! Edited by

Janusz NOWOTNY

Australian Nuclear Science & Technology Organisation Advanced Materials Program Lucas Heights Research Laboratories Menai, NSW 2234 Australia

1994 ELSEVIER Amsterdam

- Lausanne

- New York-

Oxford

- Shannon

- Tokyo

This Page Intentionally Left Blank

FOREWORD

This collection

International

Lucas Heights, on

of papers constitutes

Workshop on Interfaces Australia,

February

the Proceedings

of Ceramic Materials

1-5,

1993.

of the

held at

The objective of the Workshop was to discus research progress

the

chemistry

aspects.

of

The Workshop

ceramic

interfaces

program

and

related

included mainly

review-type

focussed on both basic and applied aspects of ceramic

mainly

of

included

This

functional

ceramics.

in this volume. Workshop

aimed

at

Some

technical

providing

industrial

interfaces,

papers

answers

papers

to

are

some

also

basic

questions concerning ceramic interfaces and also at formulating new questions

this area. The

which may serve as guide posts

extent

well

of

ceramics

as

as

ceramic

interfaces.

specific their

problems

importance

for further research

related

for

to

modern

interfaces

day

results in the formation of a new scientific discipline: experimental interfaces

and

science,

several

character

scientists solid

of

conceptual

overlaps

multidisciplinary gathered

Because

approaches areas

of

representing

state

temperature chemistry.

the

of

ceramic areas

extensive the

science

interfaces

science of

of

Due

this

such as ceramics,

electrochemistry,

metallurgy

of

technology

range

expertise.

in

of

both

to

this

ceramic

Workshop

surface

and

high

The participants represented countries such

as Japan, USA, UK, Holland, Denmark, Germany, France and Australia. The

present

collection

is a continuation

of

the previous

book published under the same title [Materials Science Monographs,

vol.

75,

engineers solid

Elsevier,

working

state

It

is

addressed

field of materials

electrochemistry

also to students I would

1991].

in the

in materials

and

high

temperature

of papers.

scientists

and

chemistry

and

science and engineering.

like to thank all the participants

in the preparation

to

science,

metallurgy,

for their efforts

The

Workshop

was

supported

by

the

Australian

Government

through the Department of Industry, Science and Technology

(DIST).

Substantial support was also obtained from the host organisation:

The Australian Nuclear

Science and Technology Organisation.

This

support is gratefully acknowledged.

The Workshop was organised under auspices of the Australasian

Ceramic Society.

Thanks are also due to Members of the International Advisory

Board, especially Yasutoshi Saito, Yusuke Moriyoshi, Woflfang Gust, Bruce Wagner, Keith Reeve and John Bannister for their cooperation.

I would like also to express thanks to my co-workers from ANSTO's

Advanced

Committee

Materials made

this

Program,

Workshop

whose

hard

possible.

work I

on

would

the

like

Organising to

thank

especially David Alexander, Sue Ansmit, Bruce Begg, Mark Blackford,

Mike

Colella,

Mike

LaRobina,

Kath

Smith,

Lue

Vance,

Debbie

Tagliaferro, Robyn Thornburn, Zhaoming Zhang and many other persons

whose

friendly

help

was

so

essential

for

the

success

of

the

help

and

Workshop. I would like also to thank Adam Jostsons, Director of the Advanced

Materials

Program

of ANSTO,

for

his

personal

encouragement since the inception of the project.

Janusz Nowotny Sydney,

1994

vii

CONTENTS

FOREWORD CONTENTS

vii xxiii

L I S T OF A U T H O R S N O N S T O I C H I O M E T R Y AND R E L A T E D CERAMIC INTERFACES

PROPERTIES

OF

J. N o w o t n y Abstract i. I n t r o d u c t i o n 2. A s p e c t s of M a t e r i a l s C h a r a c t e r i s a t i o n 2.1. C h e m i c a l C o m p o s i t i o n 2.2. S t r u c t u r e 2.3. N o n s t o i c h i o m e t r y 3. N o n s t o i c h i o m e t r y and D e f e c t S t r u c t u r e 3.1. B u l k P h a s e 3.1.1. B i n a r y M e t a l O x i d e s 3.1.1.1. E f f e c t of p(O2) on Composition 3.1.1.2. E f f e c t of P(O2) on the M o b i l i t y of E l e c t r o n i c Charge Carriers 3.1.1.3. E f f e c t of A l i o v a l e n t Ions 3.1.2. T e r n a r y M e t a l O x i d e s 3.2. I n t e r f a c e L a y e r 3.2.1. S u r f a c e C h a r g e N e u t r a l i t y Requirements 3.2.2. B i n a r y M e t a l O x i d e s 3.2.3. T e r n a r y O x i d e s 3.3. B i d i m e n s i o n a l S u r f a c e S t r u c t u r e s 3.4. C o n c l u s i o n s

9

i0 II 15 15

viii . Effect 4.1.

5.

6.

7.

8.

9.

of I n t e r f a c e s on T r a n s p o r t E f f e c t of S e g r e g a t i o n - I n d u c e d Electric Fields 4.2. E f f e c t of t h e L o c a l D o p i n g of t h e Interface layer 4.3. S u r f a c e E q u i l i b r a t i o n K i n e t i c s 4.4. C o n c l u s i o n s S e m i c o n d u c t i n g P r o p e r t i e s of I n t e r f a c e s 5.1. E f f e c t of A l i o v a l e n t I o n s 5.2 . T h i n F i l m s 5.3 . C o n c l u s i o n s A p p l i ed A s p e c t s 6.1 . E f f e c t on S i n t e r i n g 6.2 . C e r a m i c G a s S e n s o r s 6.3 . N o n l i n e a r E f f e c t s 6.4 . C a t a l y s t s 6.5 . H i g h T c S u p e r c o n d u c t o r s 6.6. M e t a l l i z a t i o n of C e r a m i c s 6.7. C o n c l u s i o n s Interface Engineering Q u e s t i o n s to be A n s w e r e d Summary Acknowledgements References

S T U D I E S OF I N T E R F A C I A L B E H A V I O R MICRODESIGNED INTERFACES A. M.

IN C E R A M I C S

VIA

15 16 16 17 17 20 20 21 22 22 22 23 24 25 25 25 25 25 28 29 30 30

33

Glaeser Abstract I. I n t r o d u c t i o n 2. E x p e r i m e n t a l P r o c e d u r e 2.1. O v e r v i e w 2.1.1. M a s k P r e p a r a t i o n 2.1.2. P h o t o l i t o g r a p h y 2.1.3. E t c h i n g 2.1.4. H o t P r e s s i n g 2.1.5. M o d e s of O b s e r v a t i o n 2.2. L i m i t a t i o n s 2.3. E x t e n s i o n s 3. A p p l i c a t i o n s 3.1. G r a i n B o u n d a r y M i g r a t i o n in D e n s e A l u m i n a 3.2. P o r e - B o u n d a r y I n t e r a c t i o n s a n d S u r f a c e T r a n s p o r t in A l u m i n a

33 33 36 36 37 37 40 40 43 43 45 45 46 48

3.3.

High Temperature Crack Healing 3.3.1. B a c k g r o u n d 3.3.2. C r a c k H e a l i n g in U n d o p e d Sapphire 3.3.3. C r a c k H e a l i n g in D o p e d Sapphire . S u m m a r y and C o n c l u s i o n s Acknowledgements References I N T E R F A C E S IN Z I R C O N I A B A S E D E L E C T R O C H E M I C A L S Y S T E M S A N D T H E I R I N F L U E N C E ON E L E C T R I C A L P R O P E R T I E S S.P.S.

Badwal

50 50 52 56 65 66 67

71

and J. D r e n n a n

Abstract i. I n t r o d u c t i o n 2. G e n e r a l C o m m e n t s 3. E l e c t r o d e R e a c t i o n s at E l e c t r o d e / E l e c t r o l y t e Interfaces 3.1. E l e c t r o d e M o r p h o l o g y 3.2. R e l a t i v e l y C l e a n I n t e r f a c e s 3.3. E x i s t e n c e of I n t e r p h a s e due to S e g r e g a t i o n 3.4. E x i s t e n c e of I n t e r p h a s e due to R e a c t i o n / D i f f u s i o n 3.5. I n t e r f a c e s B e t w e e n C o m p o n e n t s of a C o m p o s i t e E l e c t r o d e 4. I n t e r f a c e s w i t h i n the E l e c t r o l y t e G r a i n s 4.1. P r e c i p i t a t i o n of a S e c o n d P h a s e w i t h i n the G r a i n s 4.2. C o m p o s i t i o n a l V a r i a t i o n s 4.3. S e c o n d P h a s e I n c l u s i o n s 4.4. M i c r o d o m a i n F o r m a t i o n and O r d e r i n g 5. I n t e r f a c e s B e t w e e n G r a i n s (Grain B o u n d a r i e s ) 5.1. P h a s e Free B o u n d a r i e s 5.2. I n t e r m e d i a t e P h a s e F o r m a t i o n F r o m the M a t r i x 5.3. I n t e r m e d i a t e P h a s e s of the Impurity Type 5.4. I n c l u s i o n s 6. C o n c l u s i o n s 7. A c k n o w l e d g e m e n t s 8. R e f e r e n c e s

71 71 72 73 75 78 79 82 86 88 88 93 97 97 i00 i00 I00 102 103 105 106 106

A P P L I C A T I O N OF L O W E N E R G Y TO O X I D I C S U R F A C E S H.H.

Brongersma,

P.A.C.

ION S C A T T E R I N G

Groenen

and J.-P.

113 Jacobs

Abstract I. I n t r o d u c t i o n i.i. An I n t r o d u c t i o n to LEIS on O x i d e s 1.2. P r i n c i p l e of LEIS 1.3. A p p l i c a t i o n of LEIS 2. E x p e r i m e n t a l 2.1. I n t r o d u c t i o n 2.2. E n e r g y A n a l y z e r s for S c a t t e r e d Particles 2.3. E x p e r i m e n t a l F a c t o r s C o m p l i c a t i n g the A n a l y s i s 2.4. I n t e r p r e t a t i o n of E n e r g y S p e c t r a 2.5. C h o i c e of E x p e r i m e n t a l C o n d i t i o n s 2.6. F i t t i n g of LEIS S p e c t r a 2.7. C h a n g e of the S u r f a c e C o m p o s i t i o n by the A n a l y s i s 2.8. C o m p o s i t i o n a l D e p t h P r o f i l i n g 3. Q u a n t i f i c a t i o n of S u r f a c e C o m p o s i t i o n 3.1. I n t r o d u c t i o n 3.2. Ion F r a c t i o n 3.3. Q u a n t i f i c a t i o n and the P r e s e n c e of M a t r i x E f f e c t s 3.4. I n f l u e n c e of C o n t a m i n a t i o n 3.5. C a l i b r a t i o n A g a i n s t O t h e r M e t h o d s 3.6. Q u a n t i f i c a t i o n by the D I S C M e t h o d 3.7. S u r f a c e R o u g h n e s s 4. S u r f a c e S t r u c t u r e of S i n g l e C r y s t a l s 4.1. I n t r o d u c t i o n 4.2. L o c a l A t o m i c S t r u c t u r e 4.3. S u r f a c e D e f e c t s 4.4. Site L a b e l i n g 5. S u r f a c e S t r u c t u r e of N o n - S i n g l e C r y s t a l s 5.1. I n t r o d u c t i o n 5.2. The E n e r g y M e t h o d 5.3. S p i n e l s and the I m p o r t a n c e of the I n f o r m a t i o n D e p t h 5.4. S i g n a l as F u n c t i o n of L o a d i n g 6. A p p l i c a t i o n s of LEIS to S u r f a c e S e g r e g a t i o n 6. i. I n t r o d u c t i o n 6.2. S u r f a c e S e g r e g a t i o n 6.3. S u r f a c e S e g r e g a t i o n and S p i n e l s

113 114 114 116 118 121 121 121 124 128 130 131 131 132 134 134 134 136 140 141 141 142 142 142 143 145 147 148 148 148 149 152 154 154 154 160

6.4. 6.5.

Segregation During Oxidation I n f l u e n c e of Ion B o m b a r d m e n t on S u r f a c e S e g r e g a t i o n 7. G r o w t h and W e t t i n g 7.1. O x i d a t i o n 7.2. O x i d e s on O x i d e s 7.3. O x i d e s on M e t a l s and M e t a l s on O x i d e s References I N T E R F A C I A L P H E N O M E N A IN Y203-ZrO2-BASED A SURFACE SCIENCE PERSPECTIVE A.E.

CERAMICS:

160 162 164 164 168 170 172

183

Hughes

183

Abstract i. I n t r o d u c t i o n 2. I n t e r f a c e s in ZrO 2 C e r a m i c s 2.1. O r i g i n of I m p u r i t y P h a s e s 2.2. S e g r e g a t i o n M o d e l s 2.3. I n t e r f a c i a l D e v e l o p m e n t During Sintering 2.4. G r a i n B o u n d a r i e s in the Fully Dense State 3. Some C o m m e n t s on S u r f a c e Analytical Techniques 3.1. Core level B i n d i n g E n e r g y S h i f t s 3.2. D e f e c t S t r u c t u r e s 4. Y203-ZrO2 4.1. P h a s e D i a g r a m 4.2. S i n g l e C r y s t a l C u b i c S t a b i l i z e d Y203-ZrO 2 4.3. P o l y c r y s t a l l i n e F u l l y S t a b i l i z e d Y203-ZrO2 4.4. P o l y c r y s t a l l i n e T e t r a g o n a l Y203-ZrO2 4.4.1. S o l u t e P a r t i t i o n i n g and Grain Growth 4.4.2. S u p e r p l a s t i c i t y 4.5. Low T e m p e r a t u r e D e g r a d a t i o n 4.5.1. M o d e l s 5. CeO2-Y203-ZrO 2 6. AI203-Y203-ZrO2 7. C o n c l u s i o n s 8. R e f e r e n c e s

183 183 185 186 187 191 198 199 199 201 203 203 206 209 213 214 218 220 222 228 231 232 233

xii

I M P O R T A N T R O L E OF THE I N T E R F A C E S HIGH TEMPERATURE SUPERCONDUCTORS S.X.

Dou and H.K.

IN THE

Liu

Abstract i. I n t r o d u c t i o n 2. W e a k L i n k s at G r a i n B o u n d a r i e s 2.1. I m p u r i t y in the B o u n d a r i e s 2.2. G r a i n M i s o r i e n t a t i o n 2.3. C h a r g e D i s t r i b u t i o n in the B o u n d a r i e s 2.4. C o m p o s i t i o n C h a n g e in the B o u n d a r i e s 2.5. P h a s e C h a n g e N e a r the B o u n d a r i e s 3. G e o m e t r i c a l M o d e l s for B o u n d a r i e s 3.1. F l u x P i n n i n g B o u n d a r i e s 3.2. G e o m e t r i c a l M o d e l for L a t t i c e Match 4. G r a i n B o u n d a r i e s in T e x t u r e d Bi-Pb-Sr-Ca-Cu-O 4.1. H i g h Jc in A g / B i - B a s e d H T S C T a p e s 4.2. M a g n e t i c F i e l d D e p e n d e n c e of Jc for t h e s e T a p e s 4.3. 'Brickwall' M o d e l 4.4. No E v i d e n c e of W e a k L i n k s of A g / B i - B a s e d T a p e s at 77 K 4.5. G r a i n B o u n d a r y S t r u c t u r e in Textured BSCCO 4.6. G r a i n A l i g n m e n t 5. Flux P i n n i n g M e c h a n i s m in B S C C O 5.1. I n t r i n s i c P i n n i n g 5.2. D e f e c t P i n n i n g 6. S p e c i a l R o l e of A g / B S C C O I n t e r f a c e on M e c h a n i c a l P r o p e r t i e s 7. S u m m a r y Acknowledgements References THE R O L E OF I N T E R F A C E S M.

239

IN N U C L E A R

TECHNOLOGY

239 240 241 241 241 241 242 243 244 244 244 246 246 247 249 250 253 254 256 256 257 259 262 263 263 267

Yamawaki Abstract i. I n t r o d u c t i o n

267 267

xiii

2. F i s s i o n R e a c t o r Fuels 3. F u s i o n R e a c t o r P l a s m a - F a c i n g M a t e r i a l s 4. F u s i o n R e a c t o r B l a n k e t B r e e d e r M a t e r i a l s References SOME A S P E C T S E.G.

OF G R A I N B O U N D A R Y

DIFFUSION

IN O X I D E S

268 271 272 276 277

Moya Abstract i. I n t r o d u c t i o n 2. I n t e r f a c e D i f f u s i o n and S o l i d S t a t e Reactions 2.2. O x i d a t i o n of M e t a l s 2.3. M e t a l - C e r a m i c B o n d 3. G r a i n B o u n d a r y D i f f u s i o n and D e f e c t s in O x i d e s 3.1. S p a c e C h a r g e and I n t r i n s i c Segregation 3.2. E x t r i n s i c S e g r e g a t i o n 3.3. P r e c i p i t a t i o n A l o n g G r a i n B o u n d a r i e s 4. G r a i n B o u n d a r y D i f f u s i o n : D i r e c t Measurement Methods Remarks 5. R e v i e w s on A1203, Cr203, N i O and CoO C o m m e n t s and D i s c u s s i o n 6. C o n c l u s i o n s 7. R e f e r e n c e s

C H E M I C A L A N D S T R U C T U R A L A L T E R A T I O N IN T H E S U R F A C E L A Y E R S OF O X I D E S A N D S U L P H I D E S R.St.C. Smart, P. Arora, R. Hayes, C. P r e s t i d g e and J. R a l s t o n

B-S.

277 277 278 278 279 281 281 281 284 284 287 295 304 307 307

311

Kim,

Abstract i. I n t r o d u c t i o n 2. S o l u t i o n I n d u c e d R e s t r u c t u r i n g of O x i d e S u r f a c e s 3. F o r m a t i o n of S u r f a c e S i l i c a t e S t r u c t u r e s in O x i d e F i l m s 4. S o l u t i o n - I n d u c e d R e s t r u c t u r i n g of S u l p h i d e S u r f a c e s

311 312 316 323 328

xiv

5. O x i d a t i o n of S u l p h i d e S u r f a c e s : E f f e c t s of I m p u r i t y S i t e s 6. S u m m a r y a n d C o n c l u s i o n s 7. R e f e r e n c e s

INTERFACES K.

IN C E R A M I C

SUBSTRATES

331 336 337

341

Niwa

i. 2. 3. 4. 5. 6.

Abstract Introduction Aluminium Nitrate Low Temperature Sintering Magnesia Substrate Conclusions References

DIFFUSION-INDUCED GRAIN BOUNDARY IN M E T A L S A N D O X I D E C E R A M I C S E. R a b k i n ,

C.Y.

Ma

a n d W.

Ceramics

PHENOMENA

341 341 342 345 349 349 351

353

Gust

Abstract i. I n t r o d u c t i o n 2. G e n e r a l D e f i n i t i o n s 3. D i f f u s i o n - D r i v e n Interface P h e n o m e n a in N o n m e t a l s 3.1. D I G M in Z r O 2 - B a s e d C e r a m i c s 3.2. O b s e r v a t i o n s in o t h e r C e r a m i c Materials 4. W h a t is S t i l l U n c l e a r in D I G M ? 5. M o d e l b a s e d on t h e G r a d i e n t T e r m in t h e E x p r e s s i o n for the F r e e E n e r g y of an A l l o y 5.1. S t a t i o n a r y GB M o t i o n 5.2. I n i t i a l S t a g e s of D I G M 6. F o u r n e l l e ' s V a c a n c y M e c h a n i s m of D I G M 7. C o n c l u s i o n s 8. A c k n o w l e d g e m e n t s 9. R e f e r e n c e s

353 353 354 356 356 358 359 360 361 365 367 368 368 369

XV

T H I N F I L M S ON G R A I N B O U N D A R I E S IN M E T A L S A N D C E R A M I C S A N D T H E I R I M P O R T A N C E FOR THE P R O P E R T I E S OF THE M A T E R I A L S E.

371

Rabkin Abstract i. I n t r o d u c t i o n 2. G r a i n B o u n d a r y W e t t i n g 2.1. B i c r y s t a l in C o n t a c t w i t h the M e l t 2.2. W e t t i n g in M u l t i p h a s e S y s t e m s 3. S t a b i l i t y of T h i n I n t e r g r a n u l a r F i l m s 3.1. T w o - P h a s e R e g i o n 3.2. S i n g l e - P h a s e R e g i o n 4. 9R P h a s e in the S t r u c t u r e of T w i n s in Cu and Ag 5. G r a i n B o u n d a r y A m o r p h i z a t i o n D u r i n g Plastic Deformation 6. G r a i n B o u n d a r i e s at the B u l k Order-Disorder Transformations 7. G r a i n B o u n d a r y R o u g h e n i n g 8. H o w G r a i n B o u n d a r y F i l m s A f f e c t on the P r o p e r t i e s of M a t e r i a l s 8.!. M e c h a n i c a l P r o p e r t i e s 8.2. G r a i n B o u n d a r y M i g r a t i o n 8.3. D i f f u s i o n - A c t i v a t e d S i n t e r i n g 9. C o n c l u d i n g R e m a r k s I0. A c k n o w l e d g e m e n t s ii. R e f e r e n c e s

S Y S T E M A T I C U N D E R S T A N D I N G OF C E R A M I C AND RELATED INTERFACIAL PHENOMENA K. U e m a t s u

and Y.

PROCESSING

372 377 380 380 382 386 388 389 392 393 393 395 396 396 396 397

399

Zhang

Abstract i. I n t r o d u c t i o n 2. The L i q u i d I m m e r s i o n T e c h n i q u e 2.1. The P r i n c i p l e of the M e t h o d 2.2.

371 371 372

S t r u c t u r e of G r a n u l e s P r e p a r e d by S p r a y - D r y i n g 3. S y s t e m a t i c U n d e r s t a n d i n g on F e a t u r e s from G r a n u l e s to C e r a m i c s

399 399 400 400 403 406

xvi

3.1. 3.2.

S t r u c t u r e of G r e e n B o d y S t r u c t u r e of P a r t i a l l y Sintered Body 3.3. M o r p h o l o g i c a l C h a n g e of P r o c e s s i n g Defect During Densification 3.4. S t r u c t u r e of F u l l y S i n t e r e d B o d y 4. B i n d e r S e g r e g a t i o n D u r i n g C e r a m i c P r o c e s s i n g A P o s s i b l e S o u r c e of M e t a l in G r e e n B o d y 4.1. E x p e r i m e n t a l C o n f i r m a t i o n of Binder Segregation 4.2. M e c h a n i s m of B i n d e r S e g r e g a t i o n in C e r a m i c P r o c e s s i n g 4.3. R e s u l t s 4.4. E f f e c t of A d s o r p t i o n on B i n d e r S e g r e g a t i o n in S p r a y - D r i e d G r a n u l e s 5. C o n c l u s i o n s 6. R e f e r e n c e s I N T E R F A C E P H E N O M E N A IN SYNROC, A TITANATE-BASED NUCLEAR WASTE

CERAMIC

406 411 413 414 417 419 422 424 427 428 428

431

E.R. Vance, C.J. Ball, M.G. B l a c k f o r d , R.A. D a y G.R. Lumpkin, K.L. Smith, K.P. Hart, P. M c G l i n n and G.J. T h o r o g o o d i. 2. 3. 4. 5. 6.

Abstract Introduction Radiation Damage Effects I n t e r g r a n u l a r Films L e a c h i n g at the W a t e r - S y n r o c I n t e r f a c e C o n c l u s i o n s and Final R e m a r k s References

C H E M I C A L AND P O L A R N A N O D O M A I N IN R E L A X O R - T Y P E L E A D S C A N D I U M L.A.

Bursill,

J.L.

FLUCTUATIONS TANTALATE

431 431 433 436 437 438 439

441

P e n g and H. Q i a n

Abstract I. I n t r o d u c t i o n 2. E x p e r i m e n t a l 3. H R T E M R e s u l t s 3.1. I n t r o d u c t i o n to the T e c h n i q u e 3.2. H R T E M R e s u l t s

441 442 443 444 444 444

xvii 4. A n a l y s i s of H i g h R e s o l u t i o n I m a g e s 4.1. I n t r o d u c t i o n 4.2. C o m p u t i n g F a c i l i t y 4.3. C o m p u t e r S i m u l a t i o n T e c h n i q u e s 4.4. S t r u c t u r a l B a c k g r o u n d 4.5. C h e m i c a l M i c r o d o m a i n S t r u c t u r e 4.5.1. S t r u c t u r e I m a g e of O r d e r e d Microdomains 4.5.2. D e t e r m i n a t i o n of L o c a l Polarization Characteristics 4.5.3. S u m m a r y 5. D a r k - F i e l d Image R e s u l t s 5.1. H i e r a r c h y of D o m a i n T e x t u r e s 5.2. P o l a r D o m a i n F l u c t u a t i o n s in D i s o r d e r e d PST 6. M o n t e C a r l o and N e x t - N e a r e s t - N e i g h b o r S i m u l a t i o n s of C h e m i c a l M i c r o d o m a i n T e x t u r e s 6.1. I n t r o d u c t i o n 6.2. S t r u c t u r a l B a c k g r o u n d 6.3. The O r d e r P a r a m e t e r s 6.4. M o n t e C a r l o S i m u l a t i o n R e s u l t s 6.5. NNNI R e s u l t s 6.5.1. I n t r o d u c t i o n 6.5.2. A l g o r i t h m s and P r o c e d u r e 6.5.3. The E v o l u t i o n of C h e m i c a l Microdomain Textures 7. D i s c u s s i o n 7.1. M i c r o d o m a i n B e h a v i o r at t h e Ferroelectric/Paraelectric Phase T r a n s i t i o n of B a T i O 3 7.2. P o l a r D o m a i n F l u c t u a t i o n s in PST Acknowledgements References COPPER AND NICKEL ULTRATHIN CRYSTAL SURFACES

F I L M S ON M E T A L - O X I D E

446 446 446 446 447 447 447 451 453 455 455 455 458 458 460 460 460 463 463 464 465 468 468 469 470 470

473

P. J. M ~ l l e r I. 2. 3. 4.

Abstract Introduction Experimental Methods Electron-Beam Induced Charging C o p p e r and N i c k e l on R o c k s a l t M e t a l - O x i d e Structures

473 473 474 475 477

xviii

5.

6.

7. 8.

9.

4.1. Cu on M g O (i00) and M g O (iii) 4.2. Ni on M g O (i00) 4.3. Cu on CaO (i00) 4.4. Ni on N i O (i00) C o p p e r a n d N i c k e l on R u t i l e S t r u c t u r e s 5.1. Cu on T i O 2 (ii0) 5.2. Ni on T i O 2 (ii0) 5.3. Ni on T i O 2 (i00) C o p p e r a n d N i c k e l on C o r u n d u m S t r u c t u r e s 6.1. Cu on AI203 (0001) 6.2. Ni on ~-AI203 (0001) ...... 6.3. Cu on ~-Fe203 (0001) 6.4. Cu on ~-Fe203 (1012) C o p p e r on P e r o v s k i t e S t r u c t u r e s 7.1. Cu on S r T i O 3 (I00) C o p p e r a n d N i c k e l on W u r t z i t e S t r u c t u r e s 8.1. Cu on ZnO (0001) 8.2. Cu on ZnO (0001) 8.3. Cu on ZnO (i010) 8.4. Cu on ZnO (1120) 8.5. Ni on ZnO (0001) a n d ZnO (0001) C o p p e r on F l u o r i t e M e t a l - O x i d e S t r u c t u r e s Acknowledgements References

T H E S U R F A C E C H E M I S T R Y OF T I N (IV) DEFECTS, DOPING AND CONDUCTIVITY R.G.

OXIDE:

478 480 480 482 484 487 496 498 498 499 504 506 507 5O8 5O9 511 512 513 514 517 519 520 522 522

527

Egdell Abstract i. I n t r o d u c t i o n 2. T h e S t r u c t u r e of T i n (IV) O x i d e S u r f a c e s 2.1. S i n g l e C r y s t a l S u r f a c e s 2.2. P o l y c r y s t a l l i n e Surfaces 3. E l e c t r o n i c S t r u c t u r e a n d D e f e c t S t a t e s 3.1. P h o t o e m i s s i o n a n d B u l k B a n d s S t r u c t u r e 3.2. O x y g e n D e f i c i e n t S u r f a c e s : S t a t e s in the B u l k B a n d g a p 4. n - T y p e D o p i n g of T i n (IV) O x i d e 4.1. D o p i n g in P o l y c r y s t a l l i n e and Thin Film Material 4.2. S u r f a c e S t u d i e s of S b - D o p e d S n O 2 4.3. D o p i n g by Ion I m p l a n t a t i o n 5. S u r f a c e V i b r a t i o n a l S p e c t r o s c o p y

527 527 529 529 536 537 537 543 547 547 550 555 559

xix

563 564 565 565

6. S u r f a c e C o n d u c t i v i t y 7. C o n c l u d i n g R e m a r k s 8. A c k n o w l e d g e m e n t s 9. R e f e r e n c e s

I M P A C T OF G R A I N B O U N D A R I E S ON P R O P E R T I E S M U L L I T E AS A S O L I D E L E C T R O L Y T E K. J.

Yamana, Nowotny

M.

Miyamoto,

K.

Doi,

OF

T. A r a h o r i

571 and

Abstract i. I n t r o d u c t i o n 2. G e n e r a l C o n s i d e r a t i o n s 3. B a s i c P r o p e r t i e s 3.1. P h a s e D i a g r a m in t h e AlzO3-SiO z S y s t e m 3.2. C r y s t a l S t r u c t u r e 3.3. M i c r o s t r u c t u r e 3.4. M e c h a n i c a l P r o p e r t i e s 3.5. T h e r m a l S h o c k R e s i s t a n c e 4. M u l l i t e as an O x y g e n C o n d u c t o r 5. C o n c l u s i o n s 6. A c k n o w l e d g e m e n t s 7. R e f e r e n c e s

H I G H T E M P E R A T U R E E M B R I T T L E M E N T OF C E R A M I C MATRIX COMPOSITES - INTERFACE EFFECTS J.L.

571 572 572 574 574 574 577 577 581 584 590 590 590

593

Cocking Abstract i. I n t r o d u c t i o n 2. E x p e r i m e n t a l 2.1. N o m i n a l C o m p o s i t i o n 2.2. M e c h a n i c a l T e s t i n g 2.3. I n s t r u m e n t a l T e c h n i q u e s 3. R e s u l t s 3.1. R o o m T e m p e r a t u r e M e c h a n i c a l B e h a v i o r 3.2. H i g h T e m p e r a t u r e M e c h a n i c a l B e h a v i o r 3.3. C h a r a c t e r i s a t i o n of t h e R o o m Temperature Composite

593 593 594 594 594 595 596 596 596 597

xx 3.4.

4. 5. 6. 7.

Characterisation Composite Discussion Summary References Acknowledgements

MULTILAYER H. K i s h i

CERAMIC

of the

1000°C

CAPACITORS WITH NICKEL

ELECTRODES

H.

613

and N. Y a m a o k a

Abstract i. I n t r o d u c t i o n 2. D i e l e c t r i c M a t e r i a l s 2.1. C o m p o s i t i o n and P o w d e r P r e p a r a t i o n 2.2. E q u i l i b r i u m E l e c t r i c a l C o n d u c t i v i t y 2.3. E l e c t r i c a l C o n d u c t i v i t y at the Cooling Stage 2.4. D i e l e c t r i c P r o p e r t i e s 2.5. M i c r o s t r u c t u r a l O b s e r v a t i o n 3. M u l t i l a y e r C e r a m i c C a p a c i t o r s 3.1. F a b r i c a t i o n of M L C 3.2. E l e c t r i c a l P r o p e r t i e s of M L C s 4. C o n c l u s i o n 5. R e f e r e n c e s CORROSION COATINGS

600 607 610 611 612

PROPERTIES

Ichimura

OF ION P L A T E D T i N A N D C r N

613 613 615 615 615 617 618 620 622 622 622 626 626

629

and A. K a w a n a

Abstract i. I n t r o d u c t i o n 2. E x p e r i m e n t a l D e t a i l s 3. R e s u l t s and D i s c u s s i o n 3.1. H i g h T e m p e r a t u r e O x i d a t i o n 3.2. H i g h T e m p e r a t u r e O x i d a t i o n 3.3. A q u e o u s C o r r o s i o n B e h a v i o r 4. C o n c l u s i o n s 5. R e f e r e n c e s

of T i N of CrN

629 629 630 632 632 636 639 643 643

xxi

D I R E C T E N E R G Y C O N V E R S I O N IN S I N G L E G A S M I X T U R E BY A N O X I D E S O L I D E L E C T R O L Y T E E L E C T R O C H E M I C A L C E L L A N D ITS A P P L I C A T I O N S IN M U L T I - G A S S E N S I N G D.Y.

Wang Abstract i. I n t r o d u c t i o n 2. T h e o r y 3. E x p e r i m e n t a l 4. R e s u l t s a n d D i s c u s s i o n s 4.1. O x i d a t i o n T y p e of G a s 4.2. R e d u c t i o n T y p e of G a s 5. S u m m a r y 6. A c k n o w l e d g e m e n t s 7. R e f e r e n c e s

TRENDS J.

645

OF R E S E A R C H

ON C E R A M I C

Dopants Dopants

INTERFACES

645 645 646 648 649 649 655 660 660 661

663

Nowotny Abstract i. I n t r o d u c t i o n 2. S y s t e m s 3. M a t e r i a l s 3.1. O x i d e M a t e r i a l s 3.2. E f f e c t of I m p u r i t i e s on P r o p e r t i e s 3.3. L o w D i m e n s i o n a l S y s t e m s 3.4. S u r f a c e A c t i v e C e n t r e s 4. M e t h o d s 5. E x p e r i m e n t a l M a t e r i a l 6. G e n e r a l C o m m e n t s 7. A c k n o w l e d g e m e n t s 8. R e f e r e n c e s

SUBJECT

INDEX

663 663 665 666 666 667 668 669 669 671 672 673 673

675

This Page Intentionally Left Blank

xxiii

LIST

OF AUTHORS

T. Arahori (571)* Sumitomo Metal Industries, Ltd. Advanced Technology Research Laboratories 16, Sunayama, Hasaki, Ibaraki, 314-02, Japan P. Arora (311) Particle and Surface Technology Research Group University of South Australia The Levels, South Australia 5095, Australia S.P.S.

Badwal

(71)

CSIRO Division of Materials Science and Technology Clayton, Victoria 3168, Australia C.J. Ball (431) ANSTO Advanced Materials Program Lucas Heights Research Laboratories Menai, NSW 2234, Australia M.G. Blackford (431) ANSTO Advanced Materials Program Lucas Heights Research Laboratories Menai, NSW 2234, Australia H.H. Brongersma (113) Faculty of Physics and Schuit Institute of Catalysis Eindhoven University of Technology PO Box 513 5600 MB Eindhoven, The Netherlands L.A. Bursill (441) School of Physics The University of Melbourne Parkville, Victoria 3052, Australia J.L. Cocking (593) Ship Structures and Materials Division Materials Research Laboratory DSTO PO Box 50, Ascot Vale, Victoria 3032, Australia

*/ Numbers in parantheses indicate Author' s contributions begin

the

pages

on

which

the

xxiv

R.A. Day (431) ANSTO Advanced Materials Program Lucas Heights Research Laboratories Menai, NSW 2234, Australia J. Drennan (71) CSIRO Division of Materials Science and Technology Clayton, Victoria 3168, Australia K.

Doi (571) Sumitomo Metal Industries, Ltd Advanced Materials Division 1-3, Ote-machi, l-chome, Chiyoda-ku Tokyo i00, Japan S.X. Dou (239) School of Materials Science and Engineering University of New South Wales Kensington, NSW 2033, Australia R.G. Egdell (527) Inorganic Chemistry Laboratory South Parks Road Oxford, OXl 3QR, UK A.M.

Glaeser

(33)

Department of Materials Science and Mineral Engineering University of California Lawrence Berkeley Laboratory Berkeley, CA 94720, USA P.A.C. Groenen (i13) Faculty of Physics and Schuit Institute of Catalysis Eindhoven University of Technology PO Box 513 5600 MB Eindhoven, The Netherlands W. Gust (353) Max-Planck-Institut fur Metallforschung Seestr. 75 70174 Stuttgart, Germany K.P. Hart (431) ANSTO Advanced Materials Program Lucas Heights Research Laboratories Menai, NSW 2234, Australia R. Hayes

(311) Particle and Surface Technology Research Group University of South Australia The Levels, South Australia 5095, Australia

XXV A.E. Hughes (183) CSIRO Division of Materials Science and Technology Normanby Road Clayton, Victoria 3168, Australia H. Ichimura (629) Sumitomo Metal Mining Co., Ltd. Central Research Laboratory 18-5, 3-Chome, Nakakokubun Ichikawa,

Chiba,

272, Japan

J.-P. Jacobs (113) Faculty of Physics and Schuit Institute of Catalysis Eindhoven University of Technology PO Box 513 5600 MB Eindhoven, The Netherlands A. K a w a n a (629) Sumitomo Metal Mining Co., Ltd. Central Research Laboratory 18-5, 3-Chome, Nakakokuban Ichikawa,

Chiba,

272, Japan

B-S. Kim (311) Particle and Surface Technology Research Group University of South Australia The Levels, South Australia 5095, Australia H. Kishi (613) Taiyo Yuden Co., Ltd. Takasaki, Gunma, 370, Japan H.K. Liu School of Materials Science and Engineering University of New South Wales Kensington, NSW 2033, Australia G.R. Lumpkin (431) ANSTO Advanced Materials Program Lucas Heights Research Laboratories Menai, NSW 2234, Australia C.Y. Ma (353) Max-Planck-Institut fur Metallforschung Seestr. 75, 70147 Stuttgart, Germany P. McGlinn (431) ANSTO Advanced Materials Program Lucas Heights Research Laboratories Menai, NSW 2234, Australia M. Miyamoto (571) Industrial Research Institute of Ishikawa Ro-l, Tomizu-machi, Kanazawa, Ishikawa 920-02, Japan

xxvi

P.J. Moller (473) Department of Chemistry University of Copenhagen Universitetsparken 5 DK-2100 Copenhagen, Denmark E.G. Moya (277) Laboratoire de Metallurgie URA CNRS 443 Faculte des Sciences et Techniques de St. Jerome 13397 Marseille Cedex 20, France K. Niwa (341) Fujitsu Laboratories I0-i, Morinosato-Wakamiya Atsugi,

243-01 Japan

J. Nowotny (I, 571, 663) ANSTO Advanced Materials Program Lucas Heights Research Laboratories Menai, NSW 2234, Australia J.L. Peng (441) School of Physics The University of Melbourne Parkville, Victoria 3052, Australia

C. Prestige (311) Particle and Surface Technology Research Group University of South Australia The Levels, South Australia 5095, Australia H. Q u a n (441) School of Physics The University of Melbourne Parkville, Victoria 3052, Australia

E.I. Rabkin (353, 371) Max-Planck-Institut fur Metallforschung Seestr. 75, 70174 Stuttgart, Germany J. R a l s t o n (311) Particle and Surface Technology Research Group University of South Australia The Levels, South Australia 5095, Australia

R.St.C. Smart (311) Particle and Surface Technology Research Group University of New South Wales The Levels, South Australia 5095, Australia K.L. Smith (431) ANSTO, Advanced Materials Program Lucas Heights Research Laboratories Menai, NSW 2234, Australia

xxvii

G.J. Thorogood (431) ANSTO, Advanced Materials Program Lucas Heights Research Laboratories Menai, NSW 2234, Australia K. Uematsu (399) Department of Chemistry Nagaoka University of Technology Kamitomioka 1603-1, Nagaoka, Niigata, Japan 940-21 E.R. Vance (431) ANSTO, Advanced Materials Program Lucas Heights Research Laboratories Menai, NSW 2234, Australia Da Yu Wang (645) GM, NAO R & D Warren, MI 48090-9055, K.

Yamana

USA

(571)

Industrial Research Institute of Ishikawa Ro-l, Tomizu-machi, Kanazawa, Ishikawa 920-02, Japan N. Y a m a o k a

(613)

Taiyo Yuden Co. ,Ltd. Takasaki, Gunma, 370, Japan M. Yamawaki (267) Nuclear Engineering Research Laboratory Faculty of Engineering University of Tokyo Hongo, Bunkyo-ku, Tokyo 113, Japan Y. Zhang (399) Department of Chemistry Nagaoka University of Technology Kamitomioka 1603-1, Nagaoka, Niigata, Japan 940-21

xxviii

PHOTOGRAPH OF THE WORKSHOP PARTICIPANTS 1st row (sitting) from left: J. Bruce Wagner, Jr.; Richard Hannink, Kath Smith, Eliette G . Moya, Yasutoshi Saito, Zhaoming Zhang, Wolfgang Gust, Hidde H. Brongersma, Les A. Bursill, Janusz Nowotny, Preben Moller, Koichi Niwa, Keith Reeve, Yusuke Moroyoshi, Michio Yamawaki, (2nd row) from left: E.R. (Lou) Vance, Jim Woolfrey, Ray L. Frost, Janis Cocking, Roger St.C. Smart, Da Yu Wang, Andrew Glaeser, S.X. Dou, Ron Hutchings, Hiroshi Ichimura, Seshu B. Desu, Russell G . Egdell, Kazuo Yamana, Keizo Uematsu, Michael La Robina, Antony E. Hughes, Sukhvinder P.S. Badwal.

Science of Ceramic Interfaces II J. Nowotny (Editor) 9 1994 Elsevier Science B.V. All rights reserved.

NONSTOICHIOMETRY

AND

RELATED

PROPERTIES

OF

CERAMIC

INTERFACES

Janusz Nowotny Australian Nuclear Science and Technology Organisation, Advanced Materials Program, Lucas Heights Research Laboratories, Menai, NSW 2234, Australia

ABSTRACT The paper considers several aspects of nonstoichiometry and related charge neutrality conditions in the bulk phase and in the interface region of metal oxides. The effects of segregation on local properties of the interface layer of ceramic materials are discussed involving the development of concentration gradients, related electric fields and structural deformations. Applied aspects of ceramic interfaces are also briefly discussed. Several questions have been formulated with respect to the effect of interfaces on processing and properties of ceramic materials.

[. INTRODUCTION Bulk properties of compounds are different to those of the interface layer. The difference is mainly caused by segregation m d adsorption. Segregation results in the formation of composition [radients within the interface region. These gradients and 'elated electric fields have a strong impact on local properties uch as transport and, consequently, on reactivity. The ompositional changes within the interface region result in ocal structural deformations. Thus formed low dimensional nterface structures exhibit outstanding properties which are ot displayed in the bulk phase [i].

Polycrystalline materials are composed of individual gra ~ or crystallites. Consideration of their properties requires t! different concepts are applied to the bulk phase and I boundary layer (Figure i).

Figure I. Schematic illustration of a polycrystalline materi involving individual grains composed of the bulk phase and t boundary layer (BL) The properties of the bulk phase are relatively we described in the literature for most compounds. It has be realized, however, that boundary regions (Figure i) ha entirely different properties in respect to both composition a structure. Figure 2 illustrates the effect of Cr on bu electrical properties determined by thermopower and the surfa E F [3]. Therefore, the characterisation of bulk properties not sufficient to explain the properties of polycrystalli materials involving interfaces. In other words full character sation of polycrystalline materials requires that chemic composition, structure and nonstoichiometry be determined f the interface layer in addition to the bulk phase. This paper will consider nonstoichiometry and relat defect disorders for nonstoichiometric oxides independently f the bulk phase and for the boundary layer. The bulk vs. bounda layer defect disorders will be discussed in terms of properti of ceramic oxide materials such as electrical properties.

Many properties of ceramic materials, such as dielectric properties, nonlinear characteristics of resistance and catalytical properties are determined by the interfacial properties rather than by bulk properties. This strong impact of interfaces on the functions of ceramic materials has resulted in an increasing interest in the determination of the properties of interfaces such as grain boundaries and gas/solid interfaces. The purpose of this paper is to consider basic aspects of nonstoichiometry of ceramic oxides as well as related appl~ed aspects. The effects of defect segregation on nonstoichiometry and related defect chemistry of the interface layer will be discussed in more detail.

2. ASPECTS OF MATERIALS

CHARACTERISATION

Routine studies of solids, based on nonstoichiometric compounds, involve the determination of the following properties: (i) chemical composition, (2) structure, and (3) nonstoichiometry. 2.1. Chemical Composition The basic chemical composition essentially involves the host lattice elements in both cation and anion sublattices. The properties of compounds also depend on the concentrations of foreign ions which are present either in a form of intentional solutes (dopants) or un-intentional additions (impurities). The effect of the impurities on properties may be substantial even if they are present at the ppb level [i, 2]. 2.2.

Structure Structural and phase relation studies aim at the determination of the stability range of the materials within a given range of temperature and gas phase composition. 2.3. Nonstoichiometry The studies on nonstoichiometry involve the determination of the extent of the deviation from stoichiometry in all sublattices and related defect disorders. In the case of metal oxides oxygen nonstoichiometry plays an important role. This nonstoichiometry is determined by oxygen activity of the gas phase surrounding the specimen during annealing at elevated temperatures.

3. NONSTOICHIOMETRY AND DEFECT STRUCTURE 3.1.

Bulk

Phase

Defect chemistry has been widely applied to explain properties of non-stoichiometric compounds such as semiconducting and transport properties.

1.1

Ni0- CrzO3

'/

1.0 >

1.1. ILl

,-fi,, o.9 ,,>, J

1.[CFN~] =IVan] 2"[CrNi] = 2[VN'i ] 3.[Cr'Ni] + [h" ] :2[VN' i ] + [VNi] 4.EXPERIMENTAL RESULTS

0.8 Z

I

-5

I

-4 ln[Cr]

I

-3

I

-2

I ! I

I

-1

[Cr in at.%]

Figure 2. Fermi energy level, EF, of Cr-doped NiO as a function of Cr concentration. Curves I, 2 and 3 correspond to the theoretical dependencies determined using the mass action law and different charge neutrality conditions. Black points (curve 4) correspond to experimental data determined by thermopower [3 ] The concentrations of the predominant lattice defects in most of the binary metal oxides have been determined as functions of temperature and oxygen partial pressure [4].

The most reliable data on the defect structure of the bulk phase are based on properties of single crystals. In considering p o l y c r y s t a l l i n e materials one should realize that a m e a s u r e d property always involves the bulk component and the interface component [i]. Evaluation of these two components will be a major task of future studies on correct c h a r a c t e r i s a t i o n of ceramic materials. In the following section the basic rules for deriving defect disorder models will be considered for both binary and ternary metal oxides. 3.1.1. Binary Metal Oxides 3.1.1.1. Effect of p(O2) on Composition Bulk defect disorders for binary metal oxides will be considered for metal deficient oxides such as CoO and NiO. The p r e d o m i n a n t defects in these oxides are metal vacancies and electron holes which are formed as a result of oxygen interaction with the oxide crystal. The formation of these defects may be r e p r e s e n t e d by the following equilibria: i/2 0 2

~

V~'

x

+ 2h'+ O O

(1)

where V'' denotes a doubly ionised cation vacancy, h" is an electron hole and OX0 denotes oxygen in its lattice site. The lattice charge neutrality condition requires that Ch'] = 2CV~']

(2)

where the brackets denote concentrations. Eq. (2) represents a local charge neutrality requirement within the entire bulk phase, when the crystal is in thermodynamic equilibrium. Based on condition (2) one may derive simple relationships between the c o n c e n t r a t i o n of defects and oxygen partial pressure, p(O2): [h'] =

P(O2 )I/n exp {-AHf/RT}

(3)

where AHf is the change of the formation enthalpy of metal vacancies, and n is a parameter related to valency of the defects. Changes in nonstoichiometry, corresponding to shifts in equilibrium (i), may be monitored by measurements of electrical properties such as electrical conductivity: a = q [h'] ~h

(4)

where q is the elementary charge and #h is the mobility of electron holes. According to Eq. (i)-(3) the deviation from stoichiometry in equilibrium is determined by temperature and oxygen activity in the gas phase. Therefore: y = f{T, p(O2)}

(5)

where x + [v~] + [vM]

y = z [v~,]

- 1 0 [ C~

2~ o-40

'/~

1473 K

-30

(6)

-10

/,

Co0

,.y

//Vc,o 9

- 3.0

f

o

"

-40

-60

-6.0 I/

I

I

I

I

I

I

1 3 5 7 9

log po[ po:~n Pcl ]

Figure 3. The stoichiometry, concentration undoped C00 as p(O2) at 1473 K

deviation from y, and related of defects in a function of [5]

-"

- [D']

-

-5.0 -

-5 -3 -1

//

/

-

-50

11

/7/

1473 K

i

I

I

-5 -3 -1

I

1

1

I

I

1 3 5 7 9

log po[ Po2in Pcl ]

Figure 4. The deviation from stoichiometry and the concentration of defects in CoO doped with a donor as a function of P(O2) at 1473 K [5]

CoO 1473K

-1.0

-2.0

_ 0.1%

,#~

[A']

-3.0

~ o

-4.0

tvc~

i

-5.0

/ -6.0 ,

,/,

-5-3-I

,

1

,

3

,

5

,

79

,

[og Po2[Po~nPo.]

Figure 5. The deviation from stoichiometry and the concentration of defects in CoO doped with an acceptor as a function of p(Oz) at 1473 K [5]

where [V'.] and [VX,] denote the concentration of singly ionised and neutral cation vacancies. Figure 3-5 illustrate the total nonstoichiometry, y, as well as the c o n c e n t r a t i o n of particular defects for donorand acceptor-doped CoO [5]. As seen both donors and acceptors have a substantial effect on the defect structure already at the level of 0.i at %. In both cases the effect of p(O2) on y is observed above about 105 Pa while below this value the nonstoichiometry is determined by the concentration of the addition. In equilibrium the activity of defects is the same within the entire crystal (except the boundary layer). Then Eq. (2) represents the local bulk lattice electroneutrality requirement.

3.1.1.2. Effect of p(O2) on the Mobility of Electronic Charge Carriers A c c o r d i n g to Eq. (4) the electronic component of the electrical conductivity depends on both the c o n c e n t r a t i o n of charge carriers, such as [h'] and [e'], and their mobility, ~h" It has been a general assumption that a change of oxide n o n s t o i c h i o m e t r y results in a change of its electrical conductivity via the concentration of electron holes while the mobility term, ~h, remains constant [4]. Recent study of electrical conductivity and thermopower of undoped CoO performed by Kowalski et al. [6] has shown that change in P(O2) results in changes of not only the c o n c e n t r a t i o n term, [h'], but also the mobility term, #h, of Eq. (4).

'_.~

CoO 1273 K

("4 ,.--,,

9SINI3LECRYSTAL

~0./,

o POLYCRYSTAL

oi"

/

o

_.1 0 -r Z

o0.3 u

o

w !1 0

. . ~ 6"-

0

>- 0.2

l.--

I

2

10g p(02)

m O Z

3 [p(02)in Pal

/+

5

,.---,, ,,-/

> 0.5 15

C00

1373 K

9SIN6LECRYSTAL o POLYCRYSTAL

,--..,

i-.

:<,

I/+73 K

I./3 l.J.l

" 0.~

a SIN6LE CRYSTAL

O nZ O

L_J uJ u_ 0

i~I

/

/

/

~

o

~

03 o

>I.-__1 rn 0

log p(02) [p(02) in Pal

Figure 6. The mobility of electron holes in CoO, ~h, as a function of equilibrium p(O2), at 1273, 1373 and 1473 K, for both single crystal and polycrystalline specimens [6] In other words it has been shown that change in oxide nonstoichiometry results in change of the mobility of electronic charge carriers. This effect was observed for both single crystals and polycrystalline materials but it is still not clear what is the effect of grain boundaries on the character of this dependence (Figure 6).

3.1.1.3. Effect of Aliovalent Ions Incorporation of aliovalent ions into the lattice positions of a metal deficient oxide, M1.yO, forming acceptors (A) and donors (D), may be represented by the following equilibria for mono- and tri-valent ions, respectively: x

A20 + 1/2 02 - 2A~ + h" + 200

(7)

x

(8)

D203 - 2D M + V~' + 300 The following requires that

[AI~I] + [ e ' ]

general

lattice

electroneutrality

+ 2 [ V ~ ' ] = [DM] + [ h ' ]

condition

(9)

According to the condition (9) one may control the concentration of the electronic charge carriers [e' ] and [h'] by doping the crystal with aliovalent ions. 3.1.2. Ternary Metal Oxides The predominant defects in model ternary oxides, such as BaTiO3.y , are oxygen vacancies (V0"), cation vacancies (VTi'''' and Vsa'' ) electronic charge carriers and impurities [7]. The full lattice electroneutrality condition in the crystalline bulk requires that: [h']+2[V O " ]+[DTi ] = [e']+4[ V'''' Ti ]+2[ V'Ba' ]+[A~i ]

(10)

where [D'Ti] and [A'Ti ] denote the concentration of singly ionised donors and acceptors, respectively, in the Ti sublattice. 3.2. Interface Layer 3.2.1. Surface Charge Neutrality Requirements Construction of any defect model for the bulk phase must be based on the local lattice electroneutrality condition which requires the balance of charge at any point within the crystal. In contrast to this requirement, local charge neutrality for the surface layer is not required. Then an excess of the surface charge may be compensated by the space charge in the boundary layer as is schematically illustrated in Figure 7.

]0

+++ +++ +

+

w L9 ~r nu

I I

Z U.l l--

~P

DISTIANCE FROMTHE SURFACE

DISTANCE FROM THE SURFACE

Figure 7. Surface models of nonstoichiometric compounds involving both potential and charge distribution within the surface and space charge layers 3.2.2. Binary Metal Oxides At elevated temperature the lattice defects have a tendency to segregate resulting in an enrichment (or impoverishment) of the boundary layer in certain elements. Segregation of defects results from their different formation energies in the bulk and in the boundary layer. In consequence, segregation leads to the formation of strong concentration gradients within the boundary layer. Accordingly, the concentration of defects within the boundary layer is a function of both T and p(O2) as well as position (the distance from the interface {): y = {T, p(O2), {}

(11)

11

"•VER /~-i(T, Po2,

LAYER STRUCTURE

~)

y- f ( T, Po2)

Ml_+yO

SUBLAYER BULK

MOl-y

Figure 8. Illustration of the nonstoichiometry the surface layer

gradient within

In the case that segregation results in an enrichment in defects their concentration increases with the decrease of the distance, ~. Then strong interactions between the defects result in the formation of a surface low dimensional structures (Figure

8).

Assuming that in metal-deficient oxides metal vacancies segregate to interfaces then the interface layer is enriched with these defects resulting in the formation of a negative charge at the interface. The surface electroneutrality requires that the segregation-induced negative surface charge within the outer layer is compensated by a positive charge within the space charge layer (Figure 7). This model becomes more complicated when aliovalent cations are present in the lattice (see section 3.3.). Figures 9 and I0 illustrate the surface model of undoped and Cr-doped CoO. As seen introduction of tri-valent ions into the lattice of CoO (and also NiO [3]) results in a change of the surface model involving an inversion of the surface charge and the space charge. 3.2.3. Ternary Oxides The studies of Zhang et al. [8] have shown that segregation leads to an enrichment of the BaTiO 3 surface with both Ti and Ba vacancies, however, different formation energies of these two defects result in a Ti/Ba ratio greater than unity within the surface layer. It has been observed that the surface enrichment in BaTiO 3, expressed by the Ti/Ba ratio, depends on oxygen nonstoichiometry as well as the rate of cooling (Figures ii and 12) [8]. The results obtained are consistent with a surface model involving a negative surface charge and a positive charge in the space charge layer (Figure 13).

12

~

"~

I,,

i i I

i

"'-

SPACE CHARGE

I -

i I

I --

I I I

I

I

|

.

.

.

|

.

.

.

.

~_J

0_

,~ ~ 9

,

,

~

,

,,',

' ,

|

i

DISTANCE FROM THE SURFACE

Figure

9.

Surface

model

of u n d o p e d

CoO

0

__I ,.--,.

l--

o~

. ~o >

U 0

SPACE CHARGE

I = ! !

Z

l..ul

_3 4: l--l.ul _3 Lul

§

|

|

|

|

§

§ §

~

,,<~,

!

, I !

ugln

....

J

DISTANCE FROM THE SURFACE

Figure

I0.

=

Surface

model

of C r - d o p e d

CoO

~

13

Estimated 2.0 i

1.5

1.o5-1

,-

2.5 9 i

,

Probed

Depth

3.0 i

,

3.5 i

--

].oo

(nm) 4.0 i

9

Angular-Dependent Oxidised

9~

4.5 I

,

,

i.o5

XPS BaTiO 3

1.00

0.95

0.95

0.90

0.90

0.85

0.85 Rapidly

0.80

-

,

0.3

Cooled

,

Samples

,' 0.5

'

0.4

".

, 0.6

.

, 0.7

~

~-0.8

. -

, 0.9

,-

0.80 1.0

sin0

Figure ii. The Ti/Ba XPS intensity ratio vs. probed depth rapidly cooled BaTiO 3 [8]

Estimated 1.o5

.5

2.0

~ ~ , , ,

Probed

2.5

I

I

Depth

3.0

I

' ....

(nm)

3.5

l

'

--I

4.0 --'

l

Angular-Dependent

1 .oo

for

4.5 '

"I

"

,- 1.05

XPS 1.00

. ~ 0.95

0.95

0.90

0.90

b-, 0.85

0.85 Slowly

0.80 i 0.3

,

Cooled ,

o14

~ 0.5

Samples ,.

, o16

' " ' 0.7

~ 0.8

"'

,0.9

r

0.80 1.0

sin0

Figure 12. The Ti/Ba XPS intensity ratio vs. probed depth slowly cooled BaTiO 3 [8]

for

14

,.., ~. ~ ] ~ "~

Space Charge Layer

Bulk

Vo", h"

o N

-

+

-:-+ \ \

+

+++

Distance from the Surface

Figure

13. Surface model of undoped BaTiO 3 [8]

It has g e n e r a l l y been assumed that barium titanate involves only oxygen v a c a n c i e s (BaTiO3.z). The studies of Zhang et al. [8] have shown that a deficit in both cation sublattices of BaTiO 3 must also be c o n s i d e r e d within the surface layer. Accordingly, a correct formula for the boundary layer of barium titanate should be w r i t t e n assuming deficit in all sublattices: Bal_xTil_yO3_ z

(12)

where x, y and z are a function of p o s i t i o n w i t h i n the b o u n d a r y layer, ~, as well as of t e m p e r a t u r e and P(O2). In the boundary layer all of the q u a n t i t i e s x, y and z assume substantial values while in the bulk phase both x and y is very small because their t r a n s p o r t from the surface into the bulk is e x t r e m e l y slow. One may expect that the observed s u r f a c e - i n d u c e d e n r i c h m e n t of the surface layer of BaTiO 3 in B a - v a c a n c i e s has an effect on the electrical p r o p e r t i e s such as t h e r m o p o w e r (S). In fact S assumes d i f f e r e n t values for a single crystal, Ssc , and a polyc r y s t a l l i n e material, Sp.

15 It was documented that the S,c/Sp ratio assumes 1.72. This value, which is independent of temperature, coincides with the ratio of the theoretically determined thermopower value using the band model, St, to Sp [9]. This agreement suggests that the value of the ratio, Ssc/Sp=l.72 , is determined by interfaces. 3.3. Bidimensional Surface Structures Segregation results in a much higher concentration of defects in the interface layer than in the bulk phase. It was shown that the interface layer may accommodate defects which are not displayed in the bulk phase such as Co interstitials in CoO [i]. The enrichment of the interface layer in defects results in substantial interactions between the defects within the interface layer resulting in the formation of defect complexes and clusters which are not present in the bulk. When the segregation-induced concentration of defects in the boundary layer exceeds a certain critical value then a local structural reordering takes place resulting in the formation of bidimensional interface structures. Presence of these structures are limited to the outermost layer. These structures have outstanding properties not displayed in the bulk phase. Their identification is of strategic importance for the preparation of materials with enhanced properties. 3.4. Conclusions Correct understanding of the properties of the interface layer requires determination of the local nonstoichiometry and related defect disorders involving (i) the host lattice defects, (ii) intentional solutes, and (iii) impurities (unintentionally added solutes). Thus urgent studies are needed to accumulate an empirical data base on the effect of segregation on defect chemistry of the interface region and related properties. Enrichment of the interface layer of ionic crystals in defects results in local structural deformations which lead to the formation of bidimensional interface structures with outstanding properties. Studies are needed to characterise these structures. 4. EFFECT OF INTERFACES ON TRANSPORT In considerations of gas/solid kinetics it has been generally assumed that reactions taking place at the gas/solid interface are very fast and that bulk transport is rate controlling. This assumption has resulted in appropriate interpretation of the equilibration kinetic data which, according to the above assumption, have been considered as corresponding to the bulk transport kinetics.

16 It was shown that the above assumption is not always valid and that the gas/solid equilibration kinetics may be controlled by a slow process at the gas/solid interface. It was shown that the gas/solid interface may have a substantial effect on the local transport kinetics within the interface layer [i, i0]. This effect results either from the segregation-induced electric fields or structural deformation of the interface layer. In the following sections several examples of interface reactions, which are rate controlling for the gas/solid kinetics will be considered. 4.1. Effect of Seqreqation-Induced Electric Fields The electric field, which is formed along segregationinduced concentration gradients within the interface layer, may have a strong effect (either retarding or accelerating) on transport of defects across these fields [i0]. In the first case segregation results in the formation of a local surface diffusive resistance which is rate controlling the gas/solid kinetics. Therefore, diffusion data obtained by using solutions of the diffusion equation for bulk-controlled kinetics result in apparent diffusion coefficients which involve two components: the bulk diffusion component and the component corresponding to the surface diffusion resistance. Verification of the available diffusion data from the viewpoint of the surface component is required for obtaining real values of the transport numbers in the bulk phase and in the interface layer. However, application of appropriate analytical solutions of the diffusion equation, involving the segregation-induced diffusion resistance, requires knowledge of the segregation-induced concentration profiles within the surface layer [I0]. 4.2. Effect of the Local Dopinq of the Interface Layer It was shown that equilibration of the oxygen/BaTiO 3 system is very fast. As seen from Figure 14 the isothermal changes of electrical conductivity for an undoped BaTiO 3 single crystal at 1310 K during a re-equilibration is limited to about 30 min. After this time the electrical conductivity assumes a new constant value which corresponds to a new equilibrium state. Incorporation of a donor into a surface layer of BaTiO 3 results in a substantial change of the transport within this layer and, consequently, in the gas/solid kinetics. As seen from Figure 15 the incorporation of Nb into a surface layer results in an increase of the equilibration time from about 30 min for a single crystal of undoped BaTiO 3 to several months for a polycrystalline material of Nb-doped specimen.

17

0.019 ,'T

l.ul Z

w_<

0.016 0.013

l_j

,.o

=

Q Z O

u

t--o-.o-<)-..o-,,o p02= 3.9" 1012 Pa

0.003

[~~

p02= 3.9.1012 Pa I

I

TIME [h] Figure 14. Change of the conductance of undoped BaTiO 3 single crystal during equilibration (reduction run) at 1310 K [9] Figure 16 illustrates the model of Nb-doped BaTiO 3 which involves a strong gradient of Nb within the surface layer of BaTiO 3 (about 0.1-0.2 ~m thick). The donor doping results in a substantial reduction of the concentration of the mobile oxygen vacancies and the formation of cation vacancies which exhibit very low mobility. Concordantly, the formation of the surface layer doped with Nb results in blocking the transport of oxygen and, consequently, in retarding the gas/solid kinetics. 4.3.

Surface Equilibration Kinetics Equilibration kinetics of the surface layer may be monitored by measurements of work function. Figure 17 illustrates changes in the contact potential difference (CPD) between yttria-doped zirconia and a Pt electrode (these CPD changes are determined by work function changes of the zirconia specimen) during consecutive runs of oxidation and reduction at 1054 K [2]. The absolute CPD values as well as their slow changes within the time of annealing, ACPD, are determined by segregation kinetics of impurities such as Ca and A1 [2]. 4.4. Conclusions Nonstoichiometry within the interface layer of undoped oxides may be considered in terms of segregation of intrinsic lattice defects such as cation and oxygen vacancies, cation interstitials and defect complexes.

18

70

60

Z

~ 50 rr

30 ~

P02 = 220 Pa

I

1

30

I

60

i

90

1

120

TIME [ DAYS ]

Figure 15. Changes of resistance of a polycrystalline of Nb-doped BaTiO 3 during isobaric equilibration [9]

.,,-,.

I

--.i: Z

>

CK I-Z

o

w "la.

~

o,%

DISTANCE FROM THE SURFACE

Figure 16. Surface

layer model of Nb-doped BaTiO 3

specimen

19

I-\

~' eR

R

04 >

~ J

C3 13..

o

03

02 -

I

0.1 - R-REDUCTION O-OXIDATION ~

I

20

- -

( 1.74-10 Pa ) (4.3 lOZ'Po 9 )

"--.2

i

i

40

2

I

60

TIME

80

"L'~--~ Z_./

9

100

f ft

0

,00

120

[h ]

Figure 17. Changes in CPD between y t t r i a - d o p e d zirconia and Pt during p r o l o n g e d annealing at 780~ (increase of CPD c o r r e s p o n d s to increase in work function of zirconia) [2] However, assuming that all m a t e r i a l s contain less or more impurities we have to realize that the picture of the segregat i o n - i n d u c e d c o m p o s i t i o n gradients should be c o n s i d e r e d as a s u p e r i m p o s i t i o n of intrinsic defect c o n c e n t r a t i o n s and extrinsic defect c o n c e n t r a t i o n s (impurities). Lateral interactions between all these defects results in the formation of low d i m e n s i o n a l interface structures. It is e x t r e m e l y difficult to determine e x p e r i m e n t a l l y the s e g r e g a t i o n of intrinsic defects resulting from n o n s t o i c h i o m e t r y because most of the available surface t e c h n i q u e s are not sensitive enough for d e t e c t i o n of changes in crystal stoichiometry. In contrast to these intrinsic defects most of the reports on s e g r e g a t i o n in oxide m a t e r i a l s concern solutes. Impurities are also termed as u n i n t e n t i o n a l dopants present at a very low concentration. Despite this low c o n c e n t r a t i o n their effect on m a t e r i a l s properties can be significant. Traces of a l i o v a l e n t impurities added to an insulating crystal, such as Li in NiO, result in its t r a n s f o r m a t i o n into a s e m i c o n d u c t o r or even a good conductor. The impressive effect of traces of MgO on s i n t e r i n g of alumina may serve as an example of a substantial impact of additions on p r o c e s s i n g and resulting p r o p e r t i e s of the materials.

20 A c c o r d i n g to the d i s c u s s i o n presented above s e g r e g a t i o n of impurities in n o n s t o i c h i o m e t r i c compounds must be c o n s i d e r e d along with intrinsic defects and their mutual interactions in the b o u n d a r y layer. Therefore, the picture of s e g r e g a t i o n of an impurity is well defined when the spectrum of all the impurities and the conditions d e t e r m i n i n g the equilibrium, i.e. t e m p e r a t u r e and the gas phase composition, are well defined. The observed substantial effect of e q u i l i b r i u m p(O2) on the depth p r o f i l e of solute ions, such as Cr in NiO [ii], confirms a strong effect of oxide n o n s t o i c h i o m e t r y within the interface layer on the segregation p r o f i l e of the ions. C o m p e t i t i v e segregation of several elements of d i f f e r e n t driving forces may result in the formation of a s a n d w i c h - t y p e surface layer as has been reported for y t t r i a - d o p e d ZrO 2 [2]. In this case the outer surface layer is p r e d o m i n a n t l y e n r i c h e d with one elements and the sub-surface layer is enriched with another one. The enrichment coefficient (surface/bulk concentration) of defects may assume several orders of m a g n i t u d e (104 - l0 s) [i, 2]. Therefore, the s e g r e g a t i o n - i n d u c e d c o n c e n t r a t i o n s in the interface layer may assume very high values even if the segregating elements are present in the bulk phase at a very low level. 5. S E M I C O N D U C T I N G 5.1.

PROPERTIES OF INTERFACES

Effect of A l i o v a l e n t Ions Fermi energy is the basic quantity of s e m i c o n d u c t o r s which is sensitive to the density of states (donors and acceptors) and their p o s i t i o n in the band model. Thus m e a s u r e m e n t s of the Fermi energy are very sensitive to the defect structure and related s e m i c o n d u c t i n g properties. It was d o c u m e n t e d that the m e c h a n i s m of i n c o r p o r a t i o n of a solute in the bulk phase may be different to that corresponding to the interface layer [i, 3]. Parallel m e a s u r e m e n t s of thermopower, which is a bulk sensitive property, and work function, w h i c h is a surface sensitive property, have shown that Cr acts as a donor when incorporated into the bulk of NiO while in the bulk Cr forms acceptor centres (above 0.2 at%). These studies also indicate that tri-valent ions, such as Cr, result in changes of the surface polarity from negative for u n d o p e d NiO to p o s i t i v e for Cr-doped NiO. Similar model is valid for CoO [12].

2]

5.2. Thin Films Electrical conductivity and thermopower are essentially bulk sensitive properties. However, decreasing the thickness of thin films result in increasing the interface component of the measured property. Figure 18 illustrates the reciprocal of the oxygen pressure exponent of electrical conductivity for undoped CoO as a function of film thickness [12].

4,2 4,1

0

"~

O')

o .. D o., O r,.,..

'

II

4.0 3,9 3,8 3,7 3,6

3,5

. - ~

3,4 3,3

.... 900

, .... 950

, .... 1000

, .... 1050

TEMPERATURE

Figure 18. Reciprocal thin films [12]

the P(O2)

o []

7.95 I~m 1.05 pm

A

0.375 p.m

, .... 1100

. 1150

.... 1200

[~

conductivity exponent

for CoO

As seen this quantity decreases with decrease of the film thickness to values much below four which is a critical value for ideal defect model of CoO [4]. This effect indicates that the interactions between defects in the interface region of CoO are much stronger than those in the bulk phase. These interactions may be considered in terms of Co interstitials, which form preferentially in the boundary layer and, along with Co vacancies. Both these defects result in the formation of the spinel-type overlayer structure. These structural changes within the interface layer lead, in consequence, to the formation of a Schottky-type barrier within this layer (Figure 19).

22

Figure 19. A Schottky-type

barrier at the CoO surface

5.3. Conclusions Defect disorder of the boundary layer and related semiconducting properties may be entirely different to those of the bulk phase. Also the mechanism of incorporation of solute ions may be different as well. Accordingly the mechanism of doping which has been established for the bulk phase of compounds [4] is not valid for the interface layer. The determination of the effect of doping on the local defect chemistry of the interface layer requires an independent experimental approach either by studying surface sensitive properties, such as work function, or (2) by measuring bulk properties for specimens of different grain size or (3) by measuring properties of thin films. 6. APPLIED ASPECTS 6.1.

Effect on Sintering

The effect of MgO on sintering of alumina has been known since 1956 [13]. Since then intensive studies have been carried out on the determination of surface and grain boundary properties of Mg-doped alumina [14-16]. The main difficulties in the determination of Mg-segregation in alumina were produced by a very strong segregation of Ca.

23 Only very recently it has been found that Ca selectively segregates only to the prism crystallographic plane (i010) resulting in enrichment by a factor of 103 while the enrichment of the basal (0001) plane in Ca is only 50 [16]. On the other hand Mg segregates effectively to both planes, however, its presence at the external surface is limited because of substantial evaporation. 6.2. Ceramic Gas Sensors The sensing signal, which usually involves change of an electrical property of the sensor material, is generated at the gas/solid interface (Figure 20). Therefore, in the preparation of gas sensors with desired sensitivity and selectivity the main attention should be focussed on surface technology.

C02 o-0-o

0--o CO

H20> m

COATING

BULK MATERIAL

Q

~

Figure 20. Schematic illustration of a gas sensor based on the determination of electrical conductivity In chemical gas sensors the sensing signal is generated during the charge transfer between the adsorbed (or absorbed) gas molecules and the surface of a semiconducting sensor material. Accordingly, the desired sensitivity of sensors towards a particular gas phase component may be achieved by preparation of surface which exhibits possibly a high reactivity with this component. The selectivity may also be increased by decreasing the surface reactivity with other gas phase components. Modification of the surface reactivity may be achieved by changes of surface chemical composition e.g. via local dGping and coating (Fig. 16).

24 Segregation of defects, mainly of aliovalent ions, plays an important role in the modification of surface properties. One of the most sensitive ways of measuring the sensing signal at the gas/solid interface is based on the measurement of surface potential using a high temperature Kelvin probe [17]. It appears that this probe is extremely sensitive to changes in surface composition and, therefore, may serve as a sensor for the determination of traces of gases.

Figure 21. Illustration of the PTCR effect within a single grain and cross two grains [18] 6.3. Non-linear Effects Awareness is growing that interfaces have a substantial effect on properties. This effect can be well illustrated by the positive temperature coefficient of resistance (the PTCR effect) which has been reported for BaTiO 3. This effect, however, is restricted to polycrystalline material and is not displayed by a single crystal (Figure 21 [18]). Studies of Mizutani et al. [18] have shown that at least one grain boundary is required for the PTCR effect to be observed. Attempts to understand the PTCR effect and also other properties of ceramics have resulted in an increasing interest in studies of interfaces of ceramic materials.

25 6.4. Catalysts Catalytic properties of solids, such as activity and selectivity towards the formation of desired reaction products, are determined by the properties of the surface and related active centres which are formed during the preparation of the catalyst and its subsequent activation process. 6.5. High T c Superconductors There have been several reports indicating that segregation-induced changes of the local grain boundary composition result in the formation of weak links in conducting properties of oxide cuprates [19, 20]. One of the reasons for the weak links is the decrease of the local T c within the grain boundary region. Alternative reports have shown that the local T c in grain boundary regions may assume a much higher value than in the bulk phase [19]. 6.6. Metallization of Ceramics Metallization resulting in strong adhesion at the metal/ceramic interface is an important issue in micro-electronics. The adhesion is determined by properties of individual surfaces such as chemical composition and microstructure. One may expect that segregation of lattice defects and its effect on the local interactions at the metal/ceramic interface may be an important aspect of interface engineering. 6.7. Conclusions Interfaces have a controlling effect on the properties of industrial ceramic materials such as sensors [21, 22], dielectrics [23] and non-linear resistors [18]. So far, however, the preparation of these materials is empirically based rather than based on knowledge of the local properties of interfaces. Therefore, there is an urgent need to increase our understanding of interface properties and the relationship between the local properties of interfaces and the properties of bulk materials.

7. INTERFACE

ENGINEERING

Awareness is growing that in order to meet the tough requirements of new technologies, development of advanced materials of enhanced properties will be required. In the case of functional ceramic materials the main effort should be focussed on engineering of interfaces, such as surfaces and grain boundaries, which control the properties of functional

26 ceramic materials. The engineering of interfaces should be based on through understanding of their local properties and better understanding of the relationship between the interface properties and conditions of processing such as gas phase composition, temperature, time of annealing and rate of cooling. Appropriate engineering also involves appropriate post preparative treatment of materials such as local doping of the interface layer, coating and etching. An important strategy in the development of interface engineering involves better understanding of the properties of low dimensional systems such as thin films, multilayer systems, fine grain ceramics and heterogeneously dispersed systems. An important area involves the effect of impurities on interface composition and related properties. Correct characterisation of materials for their impurity level, especially at low levels, is very difficult. It has been realised that even traces of impurities in the bulk phase may segregate during a high temperature treatment, resulting in very high enrichment at interfaces. In developing basis of interface engineering composition we have to understand the principles of segregation for compounds. Better knowledge of the segregation phenomenon may help in removing undesired impurities from the interfaces and to bring up the defects which result in the desired composition and, consequently, desired properties. We are still a long way from being able to control the interface composition via appropriate processing and, concordantly, most industrial materials are still fabricated on an empirical basis rather than through understanding of interface phenomena such as segregation and interface transport kinetics. Development of a basis of interface engineering requires the following items to be addressed: i.

There is a need for better understanding of interface nonstoichiometry and related local properties of compounds at elevated temperatures and under controlled gas phase composition.

2.

Basic studies are needed to evaluate the driving forces of segregation and to determine segregation equilibria in order to predict the effect of both un-intentional solutes (impurities) and intentional solutes (dopants) on interface chemistry. Again the segregation equilibria may be modified in a wide range by adjusting the appropriate gas phase composition during all steps of processing.

27 3.

It becomes increasingly important to control nonstoichiometry of the interface region. As in the bulk phase this may be performed either by change of partial pressure of one of the lattice component in the gas phase during annealing or by doping with aliovalent ions. In the case of doping one should realize that the bulk concentration of the dopant should be lower than its desired concentration in the boundary layer by the factor which is equal to the enrichment coefficient. When the segregation enrichment is very high then introduction of traces of elements may be sufficient to modify entirely the interface composition. Also modification of nonstoichiometry in the boundary layer in a controlled way requires knowledge of segregation equilibria.

4.

One should evaluate the relationship between interface chemistry and materials properties. In order to address this issue there is a need to accumulate a data base on chemical composition of interfaces and related properties such as electric properties and dielectric properties.

5.

The low dimensional interface structures, which are formed as a result of defect segregation, must be characterised.

6.

There is a need to develop a strategy of modification of interface properties in a desired way e.g. by forming interface structures in a controlled manner. We also need to develop processing techniques which result in the preparation of reproducible interface structures.

7.

Item #5 & 6 may be effectively addressed using more sophisticated surface sensitive techniques for 'in situ' monitoring of local interface properties of compounds during processing at elevated temperatures and under controlled gas phase composition. Therefore, there is a need to develop new surface techniques which will meet the above requirements.

8.

It is very important that characterisation of ceramic specimens involves routine determination of the concentration of impurities at the level of several ppm or even at lower level. High temperature processing may result in substantial enrichment of these impurities at interfaces. It is impossible to control interface chemistry without knowledge of the level of this un-intentional doping especially with aliovalent ions.

28 9.

Awareness is growing that cooling results in substantial changes in properties of the interface layer. One source of these changes is the formation of concentration gradients in the interface layer as a result of the incorporation of elements coming from the gas phase, such as oxygen. Another effect results from local changes in segregation equilibria. Superposition of these effects may result in a complicated picture of the interface layer. By varying the cooling rate and the gas phase composition one can impose the concentration gradients in a controlled way. So far, these changes in industrial materials have been imposed in an empirical manner rather than based on knowledge of segregation and gas/solid reactivity. Therefore, there is a needed to evaluate the changes that occur at interfaces during cooling and to develop an understanding of the kinetics of interface phenomena that allow one to impose desired changes during cooling.

8. QUESTIONS TO BE ANSWERED The purpose of this section is to formulate some questions which may arise concerning the effect of interfaces and interface segregation on the properties of ceramic materials. i.

What is the segregation-induced enrichment of the interface region in both intrinsic and extrinsic defects? Knowledge of the bulk concentration and the enrichment factor allows one to evaluate the surface composition. In this regard it is important to evaluate the predominant driving force(s) of segregation and the resulting enrichment factor which may reach several orders of magnitude.

2.

How does the segregation-induced enrichment of the boundary layer depend on bulk composition? One may expect that multisegregation results in either attractive or repulsive lateral interactions within the boundary layer. Concordantly, the extent of segregation of ions which are undesired to be present at the interface may be decreased by dissolution of secondary ions which results in repulsive interactions thus leading to decrease of the enrichment in the undesired ions.

3.

How does the segregation enrichment depend on the composition of the gas phase?

29 .

.

.

What is the depth of the segregation profile and how does it depend on crystal properties? What is the effect of segregation on the defect chemistry of the interface region and how may the defect structure be represented? What is the difference between the bulk and the interface miscibility limit and what is the impact of this upon phase relations?

7.

What is the effect of segregation on ionic reconstructions within the interface region and the formation of low-dimensional interface structures? What are the properties of these structures?

8.

What is the effect of the interface structures on material properties?

9.

What is the effect of segregation on the formation of fast diffusion pathways both along and across the interface?

i0.

What is the impact of the interface region on the reactivity and processing of ceramic materials?

Ii.

What is the effect of interface properties on the electrical signal which is generated at the gas/solid interface for sensor-type materials?

12.

What is the impact of segregation on catalytically active surface centers?

13.

What is the role of segregation on the formation of the electrical grain boundary barrier in nonlinear resistors?

14.

How may the interface layer be engineered, using the effect of segregation in order to achieve the desired properties and functions of ceramic materials?

the

formation

of

9. SUMMARY Chemical composition and nonstoichiometry of the interface layer may differ substantially from those of the bulk phase as a result of segregation and related phenomena.

30 The segregation-induced compositional gradients may lead to structural deformations of the outermost interface layer resulting in the formation of bidimensional structures with outstanding properties 9 These properties may be modified, by changes of grain size or by varying the thickness of thin films 9 The effect of impurity segregation, even if the impurities are present at very low concentration in the bulk phase, may have a substantial effect on the entire picture of segregation 9 Therefore, the data concerning equilibrium segregation should correspond to well characterized materials with respect to their impurity concentrations 9 Studies are needed to engineer interface properties by imposing the desired chemical composition through an appropriate combination of processes such as segregation of desired elements, annealing and programmed cooling as well as other processes such as coating, implantation and CVD. The gas phase composition is an important processing parameter since the gas phase components interact with the ceramic body during all stages of the process 9 The reactivity depends on temperature 9 However, even at low temperatures the gas/solid reactivity may have a substantial effect on chemical composition of the interface layer and, consequently, on processing and properties of the ceramic material 9

ACKNOWLEDGEMENTS Support of the Commonwealth of Australia through the Department of Industry, Science and Technology (Grant # C91/02826) is gratefully acknowledged. This paper was kindly reviewed by Dr. Cliff Ball. His comments are sincerely appreciated. REFERENCES

1 2. 3. 4. 5.

J Nowotny, in: 'Science of Ceramic Interfaces' J Nowotny, Ed., Elsevier, Amsterdam, 1991, p. 79 J. Nowotny, M. Sloma and W. Weppner, Solid State Ionics 283o

(1988)

144s

J. Nowotny and M. Rekas, Solid State Ionics 12 (1984) 253 P. Kofstad, 'Nonstoichiometry, Diffusion and Electrical Conductivity in Binary Metal Oxides', Wiley-Interscience, New York, 1972 J. Nowotny and M. Rekas, J. Am. Ceram. S.c., 72 (1989) 1199

31 6. 7. 8. 9. i0. ii. 12. 13. 14. 15. 16. 17. 18. 19.

20. 21. 22. 23.

K. Kowalski, Thesis, U n i v e r s i t y of Nancy I, Faculty of Sciences, 1994 J. N o w o t n y and M. Rekas, Solid State Ionics 49 (1991) 135 Z. Zhang, P.J. Pigram and J. Nowotny, Intern. C e r m . M o n o graphs 1 (1994) 455 J. N o w o t n y and M. Rekas, Ceram. Intern., 20 (1994) 265 Z. A d a m c z y k and J. Nowotny, J. Phys. Chem. Solids 47 (1986) ii W. Hirschwald, I. Sikora and F. Stolze, Surf. Interface A n a l y s i s 3 (1985) 157 A. Bernasik, K. Kowalski and J. Nowotny, unpublished results R.L. Coble, J.Appl.Phys., 32 (1961) 787 H.L. M a r c u s and M.F. Fine, J.Am. Ceram. Soc., 55 (1972) 568 W.C. J o h n s o n and D.C. Stein, J.Am. Ceram. Soc., 58 (1975) 485 S. Baik, J.Am. Ceram. Soc., 69 (1986) CI01 J. Nowotny, M. Sloma and W. Weppner, A d v a n c e d in Ceramics, 23 (198v) is9 N. Mizutani, u n p u b l i s h e d results J. Nowotny, M. Rekas, D.D. Sarma and W. Weppner, in: 'Surface and N e a r - S u r f a c e Chemistry of Oxide Materials', J. N o w o t n y and L.C. Dufour, Eds., Elsevier, Amsterdam, 1988, pp. 669 S.X. Dou, this volume, p. 239 D.Y. Wang, this volume, p. 645 K. Yamana, M. Miyamoto, K. Doi, T. Arahori and J. Nowotny, this volume, p. 571 H. Kishi and N. Yamaoka, this volume, p. 613

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Science of Ceramic InterfacesII J. Nowomy(Editor) 1994 ElsevierScience B.V.

33

S T U D I E S OF INTERFACIAL B E H A V I O R IN C E R A M I C S VIA MICRODESIGNED INTERFACES Andreas M. Glaeser Department of Materials Science and Mineral Engineering, University of California, and Center for Advanced Materials, Lawrence Berkeley Laboratory, Berkeley, CA 94720 USA ABSTRACT Submicron- to millimeter-scale, controlled-geometry, controlled-crystallography pore or defect structures can be produced at internal interfaces in controlled misorientation bicrystals, single-crystal/polycrystal, and polycrystal/polycrystal ensembles by utilizing a technique that combines lithography, ion beam milling, and hot pressing. Such microdesigned interfaces provide a new tool for studying interfacial properties and behavior. The capabilities and limitations of the technique are summarized. Model experiments investigating a broad range of phenomena including pore coarsening and pore elimination, pore boundary interactions during sintering, evolution of pore networks during sintering, and high temperature crack healing have been developed. These new model experiments provide a framework for improving our understanding of the thermodynamic and kinetic characteristics of surfaces and grain boundaries, their sensitivity to impurities, and more generally, the role that interfaces play in the sintering of ceramics. Recent advances in the understanding of poreboundary separation and of the role of crystallography and impurities in hightemperature crack healing achieved through the use of microdesigned interfaces will be highlighted.

1. INTRODUCTION Typically, crystalline ceramics are fabricated from fine particle size powders, powders with an average particle size in the range of a few tenths to several microns. As a result of the free particle size, the specific surface area of the powders, and of compacts formed from them, is typically of the order of several m2/g. Pore structures in well-packed compacts also have a characteristic dimension of the order of a micron or less, and thus, during sintering, the evolution of micron or submicron scale void networks and voids is of interest. More generally, the thermodynamic properties and kinetic characteristics of surfaces will play an important role in the evolution of the structure. Multimodal packing and the use of nanometer-scale powders drive the size scale of interest even lower. The interfacial area to volume ratio scales inversely with a characteristic dimension of the particles (e.g., a particle diameter) or of the microstructure (e.g., the grain size). A trend to processing with finer particle size powders will further increase the importance of (solid-vapor) surface-related phenomena in processing and sintering. If densification occurs without a substantial increase in grain size, the use of finer particle size powders will also

34 increase the specific grain boundary area at a given density, and thereby increase the relative importance of solid-solid interfacial phenomena. The microstructural changes that occur during sintering and hightemperature use of ceramics are generally complex. In real systems, the path of microstructural evolution reflects the outcome of a competition between many concurrent processes. For sintering, particle size, particle size distribution, and particle packing play a major role in determining the pore size, pore size distribution, and pore spacing in the green compact. These characteristics determine the thermodynamic driving forces for both coarsening and densification processes. The rate and manner in which these topologically dictated driving forces are exploited depends upon the absolute and relative values of kinetic factors such as the transport coefficients for grain boundary, lattice, and surface transport, or interfacial reaction rate constants, or both. The kinetic factors will play a major role in determining the relative fluxes associated with competing mass source-sink pairs, the dominant transport mechanism, and ultimately, the nature and direction of microstructural evolution. During the early stages of sintering, surface diffusion and evaporationcondensation can contribute to particle coarsening and interparticle neck growth without densification. Lattice diffusion and grain boundary diffusion lead to neck growth, but also produce particle-particle approach. Thus, the relative rates of mass transport by surface diffusion, evaporation-condensation, lattice diffusion and grain boundary diffusion will influence the path of microstructural evolution. During the intermediate and final stages of sintering, densification is commonly accompanied by grain growth. Cannon, Yan, and Chowdhry [ 1] have developed a theory of simultaneous grain growth during densification. The modelling assumes that densification is the result of material transport from the grain boundary to pores by lattice or grain boundary diffusion. Grain growth is assumed to be pore-drag-limited, and therefore, the grain growth rate is controlled by the pore mobility. The pore mobility is in turn limited by either vapor transport, lattice diffusion, or surface diffusion. This theoretical framework allows consideration of various combinations of densification and coarsening processes. Analytical relations have been derived for several limiting cases including: 1) densification controlled by grain boundary diffusion and grain growth limited by either surface diffusion or evaporation and condensation, and 2) densification controlled by lattice diffusion and grain growth controlled by surface diffusion. These relations provide a means of assessing the effects of changes in the particle size, green density, and the controlling transport mechanism on the path of microstructural evolution as expressed by a trajectory in a grain size-density plot. Thus, these relations and plots provide a useful tool for identifying controlling sintering mechanisms. If the competing sintering mechanisms are identified, these plots and relations can suggest changes in the processing conditions that will improve sinterability. Several strategies for manipulating the coarsening to densification ratio favorably are discussed in a review by Brook [2]. With progressive reduction of the particle size, scaling law arguments indicate that the relative contribution of grain boundary diffusion to densification will increase, and similarly, that the relative contribution of

35 surface diffusion to coarsening will increase [3]. Thus, the limiting case of grain growth controlled by surface diffusion and densification controlled by grain boundary diffusion will be relevant. The analysis by Yan et al. shows that for this limiting case, the grain size-density trajectory is independent of the initial grain size, but is strongly affected by the initial density and the ratio of the diffusivities, as expressed in the parameter A. A is defined as

176SbDbYs A = 35sDs~b

(1)

where Ds and Db are the surface and grain boundary diffusivities, 5s and bo are the effective widths for surface and grain boundary diffusion, and Ts and Tb are the surface and grain boundary energies, respectively. To illustrate the sensitivity of the predicted density-grain size trajectories to the value of A, when A - 0.2, substantial grain growth is predicted, and in practice, high densities may not be reached. For values ofA > 1, high densities with minimal grain growth are predicted. Clearly, relatively small shifts in the 5bDb/5~s ratio can have significant impact on the path of microstructural evolution. If densification occurs, the nature of pore-boundary interactions during the final stages of sintering can dete _rTnine whether complete densification is possible. During the initial and intermediate stages of sintering, the pore network is continuous and intersected by grain boundaries. As a result, grain boundary transport can contribute (and may be the dominant contributor) to densification. In the final stages of sintering, the pore network is no longer continuous. If discrete pores remain attached to grain boundaries, densification due to grain boundary diffusion can continue. However, if premature poreboundary separation occurs, grain boundary diffusion ceases to be an active contributor to pore shrinkage, and complete densification may be precluded. As a result of its importance, considerable attention has been given to modelling poreboundary interactions during sintering [4-7]. The existence of a maximum or peak pore velocity has emerged, and the importance of maintaining the grain boundary velocity below the peak pore velocity has been realized. Consequently, there is great interest in identifying additives (dopants) and processing conditions that forestall pore-boundary separation. Thus, simplified models exist that can guide the selection of proper processing conditions and can help to identify effects of impurities that would be desirable. However, extracting the relevant information from conventional experiments can be difficult, because it is olden difficult to identify the controlling processes, the systems being studied may differ geometrically from those that have been modelled, and the additions of dopants can have many effects on the process. Intentional and inadvertent impurities can affect all transport rates (e.g., lattice, surface, and grain boundary diffusivities) and surface and interfacial properties (e.g., the grain boundary mobility, surface and grain boundary energies, the surface and grain boundary energy anisotropy). In some cases, impurities may lead to the formation of an intergranular glassy phase, and the emergence of new transport paths and mechanisms. Collectively, these complications impair our ability to study and improve our fundamental understanding of microstructural development. As an example, in the nearly thirty years that have passed since Coble first demonstrated that MgO additions

36 allow sintering of alumina to full density, approximately seventy studies have sought to identify the role of MgO [8]. What has emerged is a fuller appreciation of the m a n y ways in which MgO modifies (and other dopants may modify) the thermodynamics and kinetics of microstructural evolution during sintering of alumina. In this context, it is sometimes advantageous to study phenomena other than sintering as a means of gaining insight and trying to understand the importance of certain variables on sintering behavior. For example, similarities exist between the morphological changes that occur during sintering and those that occur during high temperature crack healing [9]. Similar transport mechanisms can control or affect sintering and high temperature crack healing. As a result, crack healing studies are an alternative vehicle to more conventional studies of sintering for improving our understanding of sintering. As will be discussed, measurements of crack morphology evolution provide a convenient and powerful means of assessing the effects of impurities on surface properties and transport processes. In summary, our efforts have been guided by the belief that our understanding of the basic science underlying microstructural development in ceramics could be furthered by experiments that: 1) facilitate the isolation of specific processes that contribute to or are closely related to those that occur during sintering, 2) improve control over the microstructure under investigation, and 3) control the nature and amount of impurities present. This philosophy has led us to design model experiments that utilize microfabrication methods and diffusion bonding to create microstructures with controlled crystallography, geometry, grain boundary structure, and chemistry, so-called "microdesigned" interfaces. We have used these microdesigned interfaces as a vehicle for studying and improving our understanding of the behavior of ceramic surfaces and interfaces at elevated temperature. This paper will describe the experimental procedure that has been developed for generating microdesigned interfaces, and present results from recent and on-going studies utilizing this approach. 2. EXPERIMENTAL PROCEDURE 2.1 O v e r v i e w Lithography has been used by researchers to pattern external surfaces for many years. For example, Huang et al. studied the morphological evolution of patterned sapphire surfaces at high temperature in the 1970's [10]. Gilmore pioneered the use of lithographically produced surface patterns as the basis for nondestructive evaluation standards [ 11]; the patterned surface was pressed against (but not atomically bonded to a plate to simulate an internal defect). The method has also been applied to the fabrication of controlled-geometry internal defects. McKellar and Wardlaw used lithography to prepare two-dimensional glass micromodels of pore systems with >40 ~m pore diameters to study fluid flow in porous media [ 12]. Related techniques have also been developed to generate large stable void structures for various purposes. Kahn et al. [13, 14] introduced patterned macrovoids (with dimensions of the order of several hundred microns) into PZT ceramics using fugitive ink and tape technology to increase the hydrostatic pressure-sensing response of the ceramic. Burger et al. [15] used photolithographic techniques to introduce flaws into Sb/Al203

37 interfaces, which were 5-200 ~m wide and 2 ~m deep. The flaw morphology did not change during a 2 h anneal at 1700~ Refinements of the basic lithographic method have made micron and submicron scale structures accessible, and studies of their evolution feasible [ 16, 17]. As a result, the lithographic method has become the basis of a new set of research tools for investigating a broad range of processes that occur during ceramic fabrication or during use at high temperature [ 16-27]. This tool is also applicable to a broad range of materials. Combining photolithographic methods, ion beam etching, and hot pressing provides the ability to define and introduce surface features with a controlled geometry and location, and to subsequently transform these surface features into internal features. Although the studies that will be described subsequently have focused primarily on alumina and sapphire single-crystals, the method has been applied successfully to a variety of glasses [28], glass-bonded alumina [29], lithium fluoride [24], and silicon [28]. There are five basic processing steps in the production of a sample with a microdesigned interface: resist coating, exposure, etching, sample assembly, and bonding. The procedure used to generate a microdesigned interfacial defect structure is shown in Figure 1. The ensuing discussion will detail some of the important aspects of the processing steps. A more in-depth description of the experimental methods has been published elsewhere [16, 17, 23].

2.1.1 Mask Preparation Masks are used to transfer the desired pattern, generated using a computer-based layout system, to the surface of a sample. Chrome-oxide-coated glass disks are coated with a photoresist layer. The masks are prepared from these disks using a pattern generator (e.g., GCA Corporation, Bedford, MA) that selectively exposes the photoresist through a rectangular aperture. The aperture size adjusts in response to the information generated by the CAD program. The photoresist is developed, and the exposed oxide is removed by chemical etching. The resulting patterned mask is used to selectively expose photoresist-coated substrates. The smallest features that can be generated on the mask are 2 ~Lm• 2~m squares. More elaborate or larger features on a mask are created by combining rectangular elements with edge lengths that are multiples of 2 ~m. With experience, smoothly curved surfaces can be generated on the substrate. Examples of mask patterns that have been generated for crack healing studies are shown in Figure 2. 2.1.2 Photolithography To minimize the potential for unintentional contamination of surfaces, all processing of surfaces is performed in Class 100 (less than 100 particles >0.5 ~tm in diameter per cubic foot of air space) in the Microfabrication Laboratory at the University of California at Berkeley. Wafer surfaces are subjected to several cleaning procedures, a first cleaning in a solution of I part NH4OH, I part H202, and 5 parts H20 (hereafter called SC1) to remove organic residue from the surface, and a second cleaning in 1 part HC1, I part H202 and 6 parts H20 (SC2) to remove trace metals. Wafers are then rinsed in purified water and baked out at 150~ for 20 min. Any remaining dust particles are blown off with a nitrogen gun. A drop of positive photoresist (e.g., Shipley 1400:13, Shipley Co., Inc., Santa

38

Sample Preparation

Figure 1

Schematic illustration of processing steps used to prepare microdesigned internal interfacial structures. Top to bottom: photoresist coated substrate; selective exposure of the photoresist; etching of patterned photoresist; microdesigned interface.

39

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Clara, CA) is then placed onto the wafer and spun at 5000 rpm for 30 sec to produce a uniform 1.8 ~m thick photoresist layer. The spin rate and photoresist type can be modified to produce thinner or thicker photoresist layers for high resolution features with small etch depths, or a deep etch, respectively. For the 1.8 ~m thick Shipley photoresist, wafers are baked out at 90~ for 20 min on a hot plate to dry the resist film, prevent tackiness, and improve adhesion to the substrate. In the next processing step, the mask pattern is transferred to the photoresist-coated wafer by exposure with 436 nm radiation. This produces photochemical changes in the resist that result in differential solubility. Exposure is performed either using a 4:1 projection printer (Canon, Santa Clara, CA) that reduces the size of features in the mask by a factor 4 (which must be considered in the mask design stage) or using a 1:1 contact printer. The exposure conditions that yield optimum resolution must be determined in advance. In the final step prior to etching, the photoresist is developed for 1 min, rinsed with water for 1 min, blown dry with a nitrogen gun, and given a postbake at 95~ for 10 min. The developer and post-bake conditions depend upon

40 the specific photoresist used, and the values given are appropriate for the Shipley photoresist. 2.1.3 E t c h i n g Although crystallographicaUy selective wet etching procedures for silicon exist [30], dry etching is generally necessary for ceramics due to their good resistance to chemical attack. For the majority of the ceramics that have been studied, etching was performed with an argon ion beam in an ion mill (Veeco Microetch System, Veeco Instruments Inc., New York, NY) under perpendicular incidence. For low thermal conductivity substrates, it is necessary to take steps that facilitate removal of the heat heat generated during ion bombardment. This is necessary to prevent photoresist deterioration during etching. The substrate is rotated continuously during etching. The ion beam is uniform over a diameter of up to --7 cm (in the absence of the heat shield). Thus, several samples can be etched concurrently under uniform conditions. Sections or areas of a wafer can be masked (covered) for differing periods of time to produce a single wafer that has a locally constant etch depth (e.g., w i t h i n a specific mask area) but incorporates several etch depths. One wafer can then be annealed and used to assess the effects of a change in feature voh!me on subsequent defect evolution rates and characteristics. In some cases, it may be advantageous to operate under conditions that cause a systematic variation in etch rate within a specific set of features. For example, a spatially nonuniform etch rate can be used to produce a wedge-shaped cracklike defect. After etching, the residual photoresist is removed with acetone, and the wafers are cleaned with purified water. To remove all organic residue (from both the bombarded photoresist and the emulsion oil), wafers are typically cleaned with solutions SC1, SC2, heat-treated for a prolonged period in air at an elevated temperature (1000~ to 1200~ for sapphire), and then again cleaned with SC1 and SC2. Failure to completely remove the organic residue can result in extensive interfacial pore generation in bonded assemblies [17]. The flexibility and reproducibility of the procedure is apparent in Figure 3, which depicts pore structures intended for several different studies. The structure illustrated in Figure 3a is suitable for studies of pore drag/pore separation or pore elimination studies. A bimodal pore array in which the freer pores have submicron width is shown in Figure 3b, and is suitable for studies of pore coarsening or pore elimination. Pore channels with varying aspect ratios, as illustrated in Figure 3c, serve as a starting point for studies of pore channel break-up, cavitational creep, the development of equilibrium pore shapes, and fracture of porous solids and composite structures. In recent work, masks containing "starburst" patterns have been generated, as illustrated in Figure 4, that allow an assessment of the effects of crystallographic orientation on the morphological evolution of channel-like and cracklike defects within a single sample. The morphological evolution of tailored nonequilibrium pore shapes, such as the chevrons in Figure 3d, can be examined to probe the development of metastable pore geometries.

2.1.4 Hot pressing The general procedure for hot pressing involves a final cleaning of the surfaces to be bonded with a nitrogen gun or clean room compatible paper,

41

Figure 3

SEM micrographs of accessible surface structures: a) monomodal pore array, b) bimodal pore structures, c) pore channels and d) chevrons.

sample assembly, and then diffusion bonding under vacuum conditions. The final cleaning is designed to remove particles that adhere to the surface and may impair bond formation. Our efforts to study the effects of airborne contaminants on subsequent crack healing behavior emphasized instead the strong correlation between particle-free surfaces and ultimate bond quality. At its simplest level, sample assembly entails placing an etched wafer onto an unetched wafer (under clean room conditions), then placing the ensemble between two chemically nonreactive (e.g., high-purity BN) spacers, and finally, loading this assembly into a high-purity graphite die. For sapphire/sapphire assemblies, bonding temperatures ranging from as low as 925~ to as high as 1370~ a bonding pressure of 15 MPa sustained for i h at the bonding temperature with a vacuum of 2.6 x 10-3 Pa was typical. When transparent substrates are used, the success of the bonding procedure is readily assessed. For sapphire, conditions within the indicated range of hot pressing conditions generally resulted in continuously bonded assemblies. Some effects of sapphire orientation on the ease of bonding were noted, with basal plane sapphire having the most favorable bonding characteristics. To apply the method to other

42

mn

Figure 4

mm

mm

mn

mm

mn

mm

mn

Example of a mask containing features (from the top down) that are designed to produce 3 ~m wide channels, 6 ~m wide channels, channels of varying aspect ratio and 100 x 200 ~m cracks.

materials, the bonding conditions must be determined by trial and error. We note that this bonding method has produced standards with well-defined internal defect patterns for calibration of NDE systems, and in turn, NDE has provided a means of assessing the quality of the bond achieved when opaque materials such as Si have been diffusion bonded [31]. There is tremendous flexibility in the choice of materials that are used to form the ensemble for bonding, and this leads to a broad range of sample configurations that provide the basis for studying a wide range of processes involving ceramic interfaces. As a result of the technique's flexibility, the potential exists of isolating the effects of defect geometry, crystallographic orientation or grain boundary misorientation, and chemistry on subsequent

4] evolution of the microdesigned interfacial structure. Specimens prepared by bonding an etched single crystal to a polished single crystal of identical orientation provide a means of simulating defects in single crystals, intragranular flaws. Controlled misorientation bicrystals can be used to study intergranular defects. Single crystal-polycrystal ensembles and polycrystalpolycrystal ensembles provide an opportunity to examine the effects of a wider range of misorientation, and to simulate and study defects in typical polycrystalline materials and laminates. Doped polycrystalline samples can be made by conventional routes. Recent experiments have demonstrated that ion implantation can be a useful method of introducing controlled levels of impurities, and assessing their effect on feature evolution. For some materials, e.g., Ti-sapphire, well characterized single crystals are commercially available. A wide variety of dissi_mi!ar materials can also be bonded, allowing the extension of the method to heterophase interfaces. 2.1.5 Modes of O b s e r v a t i o n Once surface features have been transferred to an internal interface, several methods can be used to study their evolution. Optical microscopy provides a convenient nondestructive mode of observation when optically transparent materials such as sapphire are used. Using this method, a single sample can be given repeated anneals, and the time evolution of specific defects can be studied. Optical hot stage microscopy provides the potential for truly continuous observation of the evolution. When resolution of the freest features is required, or when opaque material is used, fracture surfaces can be examined using optical or scanning electron microscopy. Figure 5 illustrates a fracture surface showing a series of channel-like defects exhibiting varying degrees of instability. Lithographically introduced precracks along the perimeter of the features of interest can be used to promote failure along the interface. Cross sections perpendicular to the bond surface are also valuable, and can provide useful complementary information on the evolution of the defect shape. 2.2 L i m i t a t i o n s Virtually every conceivable defect geometry can be generated on a ceramic surface and then transferred to an interface by hot pressing. As illustrated, pores or defects with rectangular, triangular, and circular cross sections, as well as defects with varying aspect ratios can be defined and introduced. Pore arrays containing pores of almost any size, spacing, and number can be produced. There are only a few limitations. The number of features that can be generated on a mask depends upon the characteristics of the particular pattern generator used. Since a mask can be exposed multiple times on a specific surface, and the translation of the mask can be controlled with great precision and accuracy, it becomes possible to generate several hundred thousand features, which is certainly adequate for most purposes. Pore size and spacing are limited by the wavelength of the radiation used to expose the photoresist. This results in minimum pore sizes and spacings that are of the order of 1 to 2 ~m. This feature size (width) in conjunction with the limitations imposed on the minimum etch depth normally establishes a minimum pore volume that is equivalent to that for a 0.3 to 0.4 ~m radius spherical pore.

44

Figure 5

SEM micrograph of bicrystal fracture surface (basal plane sapphire) with array of 5 channels of varying size. Sample had been heat_ treated for 25 hours at 1800~ The directions e and g are [ 1120] and [1100], respectively.

The upper and lower limits on the feature depth are defined by two factors. The etch depth should exceed the characteristic roughness of the substrate. Sapphire substrates with a surface roughness of the order of 0.025 rum_ are commercially available, and similar finishes can be produced by polishing. Defects with depths as low as 0.09 rum__have been successfully transferred to an internal interface, and it is likely that shallower defects could also be transferred. An upper limit on the etch depth arises because both the resist and the substrate are etched. Thus, the maximum achievable etch depth hinges upon the relative etch rate, and the thickness of the resist. Etch depths of up to a few microns are in principle achievable. Since for many of the studies of interest etch depths of only a few tenths to one micron are desired, this limitation is again not of serious consequence. If noncubic materials or dissimilar materials are to be bonded, thermal expansion mismatch induced cracking becomes a factor. This is a particular concern in alumina, wherv thermal expansion coefficients along the a-axis and caxis differ by as much as 10% [32]. A full range of basal plane twist boundaries can be produced, but the range of misorientations accessible when prismatic plane crystals are to be bonded is much more restricted. Symmetric tilt boundaries are of course possible. Finally, preserving the interfacial defect structure requires consideration of the overall arrangement of the surface structures on the substrate. If closely spaced etch pits are distributed uniformly over the entire surface, hot pressing will promote pore elimination as a result of vacancy condensation along the interface. Uniform vacancy condensation requires a vacancy flux that increases linearly with distance from the midpoint between two equisized pores. This in turn requires a parabolic vacancy concentration profile between the two pores if the concentration gradient is the driving force for the vacancy flux. To reduce the

45 vacancy flux, and therefore the pore shrinkage rates, the interpore spacing must be increased. Fortunately, this need not be done uniformly and throughout the sample; isolating arrays of specific pore size and spacing with pore-free regions several hundred microns wide is sufficient to reduce pore shrinkage during both hot pressing/bonding and subsequent annealing significantly. If the pore-free ligament is retained, it allows controlled studies of pore coarsening. If the porefree regions can be removed, pore elimination can be studied. 2.3 E x t e n s i o n s Although most of the features that have been shown and discussed are discrete features, masks can be designed that include continuous defect structures, i.e., a continuous pore network. Thus, there is an opportunity to increase the geometric complexity of the interfacial defect structure. In some cases, these patterns can be generated easily, by simply changing the nature of the photoresist used from a positive resist to a negative resist. This change reverses the nature of the areas that are selectively removed. As a result, the isolated circular cavities illustrated in Figure 3a would become circular mesas, and these circular regions would become the contact regions during bonding. Alternatively, a negative of a mask can be made. In the situations that have been described, the defects are cavities, and used to simulate pores and cracks. However, it is also possible to increase the chemical complexity of the interfacial structures. Film deposition techniques are well-developed in the semiconductor industry, and the use of these techniques to "fill" the etched surfaces provides an opportunity to microdesign ceramic-metal composite interfaces. With continuous pore networks, infiltration of the structure, and the removal of phases from the network can be studied. A growing interest in nanoscale materials is driving an effort to extend the accessible feature size downward to the nanoscale range. E-beam lithography provides this capability, and the National Nanofabrication Facility at Cornell has facilities that provide 25 nm resolution [33]. Thus, the fabrication of nanoscale interfacial structures is possible, if substrates of sufficiently high surface quality can be prepared. Very recent developments have significantly reduced the cost of a high-resolution lithography capability. McIntyre et al. report that by interfacing a personal computer with an SEM, a system allowing the fabrication of e-beam lithography resolution features can be assembled for under $90,000 [34]. The interface system in essence converts images drawn using graphics software to an instruction set that controls the movement of the SEM's electron beam. The possibility of fabricating 5 nm features is suggested. With features of this size, the temperature range over which the evolution of structures could be studied would be extended significantly. However, detection of the defects will be much more challenging.

3. APPLICATIONS The lithographic method provides a powerful tool for isolating particular aspects of microstructural development in ceramics. The objective of the ensuing sections is to illustrate the application of the method to studies of grain boundary migration, pore-boundary interactions, and high temperature crack healing. Grain boundary migration and pore-boundary interactions were

46 reviewed in a previous paper [35], and as a result these topics are given lesser coverage in this paper. Instead, greater attention will be devoted to the results of recent crack healing and pore channel evolution studies. However, for all topics, the emphasis will be on showing how lithography can be used to advantage. The direction of on-going and planned research will be indicated. 3.1 G r a i n b o u n d a r y migration in dense a l u m i n a The preservation of a fine grain size in a ceramic may have several benefits. In the processing of a dense ceramic, a decrease in the grain growth rate can help to avoid pore-boundary separation, and thereby allow the processing of theoretically dense materials. A fine grain size may have beneficial effects on properties. For materials that will be shaped by superplastic deformation subsequent to densification, a fine grain size increases the creep rate, and facilitates forming. An understanding of the factors that lead to anisotropic grain growth can also be of value, potentially suggesting a means of producing microstructures containing columnar or platelike grains that may reinforce crack bridging in the wake of a crack. An understanding of the chemical or microstructural conditions that induce abnormal grain growth is important in its own right, and because abnormal grain growth and poreboundary separation are often linked. Abnormal grain growth, although generally undesirable during sintering, provides a convenient way of characterizing the grain boundary velocity-driving force (Vb-Fb)relationship, and thereby determining the grain boundary mobility, Mb. By comparing Mb, the velocity per unit driving force, for the growth of an abnormal grain into both an undoped and a doped theoretically dense polycrystalline matrix, the effect of the impurity on Mb can be dete _rmined. A quantitative assessment requires that the "average" grain growth rate of the matrix grains in the undoped and doped material be determined, and from this the "average" Mb for matrix grains. Comparisons between the effect of the dopant on Mb of the abnormal and the "average" matrix grain can then be made. The experimental situation is illustrated schematically in Figure 6a. Our work has focussed on a model ceramic system, A1203 with and without MgO dopant, and has examined the migration behavior of a model interface, the interface between single sapphire and undoped or MgO-doped alumina [21]. Lithography has been used to introduce a periodic array of immobile reference markers to mark the initial interface position, as illustrated in Figure 6b, and permit highly accurate measurements of the seed displacement. Measurements of average grain growth rates showed that 250 ppm of MgO reduced MbTbby a factor of--4 relative to undoped alumina at 1600~ and resulted in a more uniform and equiaxed microstructure. The ability to use seeds of controlled orientation revealed seed orientation specific solute-boundary interactions. For a basal plane (0001) seed, MbTbwas not affected by MgO additions; velocity differences were due to differences in G, or equivalently, the driving force for seed migration in the undoped and doped materials. Migration of a (1120) seed into undoped alumina was very irregular at the scale of the matrix grain size, suggesting strong effects of local grain boundary misorientation on MbYo.The addition of MgO reduced the spatial variability in

47

Figure 6

a) Schematic illustration of abnormal grain growth in which a seed grows with velocity Vb into a polycrysta__lline matrix of average linear intercept L, average grain size G, and average grain boundary energy ~ . b) Micrograph illustrating the use of lithographic reference markers to facilitate interfacial displacement measurements.

displacements, i.e., migration was more uniform, and displacements were smaller, suggesting MbTb was reduced. In principle, this method of characterizing microstructures may help identify dopants that induce anisotropic grain growth and promote the formation of elongated grains. In addition, it was determined that the grain size and time dependencies of the Mb~ products for normal and abnormal grains differed. Abnormal grains developed a progressively larger Mb~ advantage with increasing matrix grain size (anneal time). This behavior may play a role in sustaining abnormal grain growth, and possibly in the initiation of abnormal grain growth. Current efforts are focusing on extending this method to study the effect of other cation impurities in alumina, specifically, Ca and Ti. There is direct as well as indirect evidence to suggest that at least Ca segregates anisotropically to surfaces and grain boundaries in alumina [36-39]. It would be of interest to determine whether this contributes to anisotropic grain growth, and to better understand how these additives impact microstructural evolution. A parallel effort seeks to determine whether the method can be adapted to studying the effect(s) of controlled thickness and composition glassy films on grain boundary migration in alumina [29]. Opportunities exist for studying the growth behavior of single crystal seeds spanning a broader range of surface orientation, thereby

48 providing information on the interrelationship between local grain boundary misorientation, solute-boundary interactions, and resulting grain growth anisotropy. In summary, such experiments provide a tool for assessing the migration characteristics of grain boundaries at high temperatures and evaluating the effects of dopant additions.

3.2 Pore-Boundary Interactions and Surface Transport in A l u m i n a Abnormal grain growth can also provide a vehicle for studying poreboundary interactions. Figure 7a depicts an abnormal grain growing into an idealized porous matrix. The pore size plays a key role in determining the pore mobility, Mp, and the maximum drag force per pore, Fp. The pore spacing dictates the areal density of pores N, the total drag force per unit area exerted by pores, NFp, and thus, the net driving force for grain boundary migration. As indicated previously, the migration characteristics of pores relative to those of grain boundaries play a key role in the late stages of densification. The net driving force Fb - NFp and seed mobility determine the seed growth velocity; the product MpFp determines the pore velocity. Hsueh et al. have shown that for a pore attached to a migrating grain boundary, there is a peak or maximum pore velocity, Vp When pore migration is surface diffusion controlled Vp - l'lSsDs}'s (17.9

-

kTr 3

6.2~,)

(2)

where ~ is the atomic volume, r the radius of a spherical pore of equivalent volume to the interfacial pore, and ~ is the dihedral angle formed at the poregrain boundary intersection. Our research again focused on the migration of a model interface, that between single crystal sapphire and undoped and MgO-doped alumina [22]. Using lithography, reference markers and arrays of pores of controlled size and spacing (as in Figure 3a) were introduced at the sapphire/alumina interface. The pore size was freed (r --0.67 pm) to fix the peak pore velocity, and the pore spacing was varied systematically (center to center spacings of 4, 6, 8, and 10 ~tm) to adjust N, and thereby the net driving force acting on the grain boundary. Since the interfacial defect geometry could be duplicated precisely, it was possible to isolate and study the effect of the MgO addition. During annealing at elevated temperature (1600~ in this study), the pores, which originally have a disklike shape, Figure 7b, become more equiaxed. As the effective areal density of the pores decreases, the net driving force on the boundary increases, the boundary and pores accelerate, eventually reaching a steady-state velocity. During growth of the sapphire into the polycrystal, pores either migrate with the interface (remain attached as shown in Figure 7c) or separate from the interface as shown schematically in Figure 7a. Whether pores separate or remain attached can be assessed by microstructural evaluation of sapphire/alumina cross sections. The effects of changes in G and N (i.e., pore spacing) on pore-boundary interactions can thus be assessed. Several key results emerged from this study. From measurements of the pore velocity just prior to separation, it was possible to determine Vp

49

9

9

I

o

9

9

-

Figure 7

a) Illustration of seed growth at velocity Vb into a porous polycrystalline matrix. The attached pores reduce the driving force acting on the boundary, and exert a "drag" force that depends upon the pore size and spacing, b) Illustration of as-bonded sapphire/alumina interface, c) Lithographically introduced pores migrating with the sapphire/polycrystal interface.

experimentally, and using the analysis of Hsueh et al.,~a value of the 5sDs product could be extracted. Comparisons of values of VD in undoped and MgOdoped alumina indicated that 5sDs was increased by a t~actor --4 at 1600~ by the addition of 250 ppm MgO. Thus, used in this way, the technique is capable of providing information on the transport properties of surfaces at elevated temperature and allows an assessment of the effect of dopants on the surface diffusivity. The method is particularly attractive because the diffusivity is inferred from measurement of a process of direct interest, rather than a microstructurally dissimilar process.

50 In addition to providing a quantitative measurement of 5sDs, measurements of lTp also lend themselves to a prediction of pore-boundary separation conditions. Pore-boundary separation occurs when the velocity of the pore-laden boundary, Vb(p), exceeds l/p. For single crystal seed growth into a polycrystalline matrix with average grain size G, separation is predicted when

(3) where f is the center-to-center pore spacing, and Vb is the velocity of the porefree boundary into the polycrystalline material. The critical grain size for poreboundary separation can thus be related to the product of two terms, one dependent upon geometric parameters only, and the other dependent upon a ratio of measurable kinetic terms. Of the parameters, f and r can be controlled precisely. G-, Vb and Vp are experimentally measurable. Thus, all parameters can be controlled or measured. For the case of sapphire growing into undoped and MgO-doped polycrystalline alumina, good agreement was obtained between predicted and experimentally observed conditions of separation and attachment. MgO additions broadened the range of conditions within which pore attachment was maintained. Efforts have been initiated seeking to extend this method to studying the effect of Ca and Ti impurities on microstructural evolution in alumina. Experiments will emphasize measurement of the SsDs product and 17p We ultimately hope to use the method to study the effects of intergranular glassy films on grain growth, and to characterize the interactions between dopants and glassy films. More generally, this approach is applicable to any material that is available in single crystal form that can be bonded to a polycrystal. As a result, the approach should be useful for investigating and assessing the effects of dopants on pore-boundary separation, and on surface transport at elevated temperature in a wide range of ceramics. One recently discussed research opportunity entails depositing patches of metal at a polycrystalline/single-crystal ceramic interface. If the metal particles migrate with the interface, information on volume or heterophase interfacial diffusion could be obtained. If migration of the interface isolates the metal particles in the ceramic matrix, the opportunity for determining the equilibrit, m shape of the metal particles arises. This in turn would provide a means of determining the Wulff plot for this two-phase ceramic-metal system.

3.3 High temperature crack healing 3.3.1 Background There has been considerable interest in both high temperature crack healing [9, 19, 20, 25, 26, 40-63], and in other microstructurally similar phenomena [64-76] for many years. Crack healing is of direct importance for structural ceramics because the diffusive healing processes that occur at elevated temperature can reduce the deleterious effects of cracks on strength,

51 allowing partial or complete recovery of strength in cracked or machined ceramics [9, 50-63]. Crack healing phenomena is also of more general interest and utility because the microstructural changes that occur have strong parallels with those that take place during sintering [20,44,46,65,66]. As a result, studies of crack healing have the potential to contribute to an improved understanding of sintering. Crack healing in oxides has been reviewed most recently by Gupta [9]. Most studies of crack healing indicate that there are several geometrically distinct stages to the process. The initial stage is characterized by crack tip regression and blunting; regression and blunting are in some cases accompanied by facetting [58]. Discontinuous crack pinching can also initiate at cleavage steps or other irregularities along the crack front, or on the crack face [48], and in some systems and for some crystallographies this m a y dominate healing. Crack regression leads to the formation of a cylindrical ring around the crack periphery. This cylindrical rim is unstable to periodic axial perturbations. The second stage of crack healing is marked by the formation of "cylindrical" pore channels. Collapsed or healed portions of the crack perimeter propagate towards the crack interior, producing pore channels or ligaments whose axes are often perpendicular to the original crack edge. These pore channels are in turn subject to Rayleigh instabilities[68], leading to the formation of discrete spherical pores. Subsequent to the formation of isolated pores, pore elimination and pore coarsening can occur. Several important variables that affect crack healing and strength recovery have been identified. The crack opening displacement can affect both the morphological characteristics [48], and the rate of healing [61]. The crack face crystallography can modify the healing behavior. If the surface is a low index low energy surface, surface energy anisotropy can stabilize a planar face and impede "cylinderization'. If the surface is not part of the Wulff shape, roughening to form lower energy facets is energetically favorable, and this roughening can assist local crack closure. M a r u y a m a and Komatsu have demonstrated the sensitivity of the morphological evolution of cracks to grain boundary misorientation [47]. Crack face microstructure can also be important; Singh and Routbort have suggested that hillock formation accompanying grain boundary grooving can induce bridging of crack faces, and thereby contribute to strength recovery [63]. Finally, both crack healing studies [62] and measurements of surface diffusivities[77] in alumina have revealed a sensitivity of the healing and transport rates to impurities. This raises the possibility that by identifying and quantifying the effects of particular impurities on crack healing, insight could be gained on their effects on transport during sintering. Although general behavioral trends are evident, and some of the important variables have been identified, there are also pronounced system-tosystem and even sample-to-sample variabilitiesin behavior whose origins are not well understood. The difficultyin identifying the cause of these differences stems in part from a difficulty in controlling and isolating the effects of specific variables on the healing behavior. In view of the demonstrated importance of crack geometry, crack face crystallography, crack face microstructure, and sample chemistry/impurities on crack healing behavior, clarification and quantification of these effects seemed warranted. Thus, our initial efforts focused on applying lithographic methods to the study of crack healing, and

52 subsequently assessing the i m p o s e of these variables on crack healing in sapphire. More recent work has focused attention on the role of intentional impurity additions on evolution, and has sought to provide a framework for extracting information on solute-surface interactions and the effect on the Wulff plot from measurements of pore channel evolution.

3.3.2 Crack h e a l i n g in undoped sapphire Initial experiments entailed studies of crack healing in high-purity sapphire [19]. Figure 8 provides an example of a typical cracklike flaw introduced using lithography viewed in a cross section taken perpendicular to the crack plane. The flaw dimensions are 100 pm by 200 ttm by 180 nm. The morphological evolution of geometrically similar cracks lying parallel to the basal (0001) and prismatic (11~0) planes was characterized to assess the effects of a difference in crack tip (edge) and crack face crystallography on the healing behavior. Up to 200 identical cracklike flaws were generated in each experiment. Most anneals were conducted at 1800~ (0.86 Tin) in a vacuum of =1.3 x 10-3 Pa. Optical microscopy was used to study the morphological evolution of specific cracks within a particular interface. The flaw in Figure 8a was etched into a basal plane, and after bonding, the sample was annealed for 10 min at 1800~ The pronounced difference in curvature along the crack perimeter induces mass redistribution, which leads to recession of the crack perimeter, and the formation of an annular ring, which is seen more clearly in the enlargements of the edge geometry, Figures 8b, 8c. Some facetting along the crack perimeter is evident, suggesting that surface energy anisotropy continues to be important even at the high anneal temperature used. Figure 9 presents optical micrographs taken using transmitted light, illustrating the morphological evolution of a pair of cracks oriented parallel to the basal plane, originally 200 pm by 100 pro, and with a depth of 180 nm. The directions e and g are [1120] and [1Y00], respectively. Figure 9a illustrates the as-prepared crack geometry. The same two cracks are shown in Figure 9b after an anneal of 60 rain at 1800~ The annular ring, evident in the cross section in Figure 8, is unstable to growth of sufficiently long wavelength perturbations. These perturbations cause a variation in the local crack opening displacement. A decrease in crack opening displacement increases the crack regression rate, and thus, points at which the annulus is thinner will recede more rapidly. This, coupled with longitudinal mass redistribution, causes what appears as local pinch-off of the annular ring. Once closure occurs, healing via intrusion is rapid; Figure 9c shows the interface after a total of 90 min of annealing at 1800~ and a considerable fraction of the original crack has been healed. Figure 9d shows the crack after 135 min total annealing time. At this stage, the crack contains several cylindrical segments. These are subject to Rayleigh instabilities. Although one can study the evolution of features such as these, that have evolved from larger defects, it is more convenient to investigate this later stage of healing using lithographically introduced features like those shown in Figure 3c [20]. The healing sequence presented is qualitatively consistent with the general observations summarized by Gupta [9].

53

Figure 8

SEM micrograph of a cracklike flaw lying parallel to the basal plane, viewed in a cross section taken perpendicular to the crack plane after 10 rain at 1800~ e = [1120] and f = [0001]

A comparison between the healing behavior of the cracks in Figure 9, and those of other basal plane cracks of identical geometry and crystallography shows that differences in healing behavior are related to the relative extents of crack regression before perturbation initiates. If edge instability occurs readily, a fine scale structure, and relatively rapid healing results. If instead, the edges are more resistant to instability, crack regression dominates; this reduces the healing rate and produces a much coarser healed structure. It is interesting to note that even in cases where highly controlled flaws are introduced, variabilities exist. One can anticipate a reduced level of reproducibility in the flaw geometry in real mechanically induced flaws, and thus, quantitative kinetic data derived from such studies should be treated with an appropriate level of caution.

54

Figure 9

Optical micrographs of a pair of 100 pro_x 200 ~m x 0.18 ~m deep cracklike flaws, after a) 0 min, b) 60 min, c) 90 min, and d) 135 min anneals at 1800~ The directions e and g are [11~0] and [ 1T00], respectively. Edge perturbation dominates and leads to a fine-scale healed structure.

Dramatic differences in morphological evolution can occur when the crystallography of the crack edges and faces is altered. One type of healing behavior for cracks oriented parallel to the prismatic plane is illustrated in Figure 10. Figures 10a and 10b show two initially 120-nm deep cracks oriented parallel to the (1120), after 5 and 210 min at 1800~ Crack healing occurs via the growth of "pillars" or columns that ultimately connect the crack faces. During continued annealing at elevated temperature, the columns expand, and thereby heal considerable portions of the crack. This mode of healing has the potential to produce a significant increase in the healing rate relative to that due to crack regression alone, and provides a mechanism for healing cracks at positions away from the crack perimeter. Hickman and Evans have observed similar behavior in calcite [48]. Not all prismatic plane flaws healed in this manner. In other cracks healing involved the formation of pore channels that were parallel to the original crack front. The healing mechanism is distinct from that for basal plane cracks,

55

Figure 10

Two cracks, initially 120 nm deep, oriented parallel to the (1120) after a) 5 and b) 210 minutes at 1800~

and does not appear to involve instabilities along the crack perimeter, but instead closure parallel to and away from the crack perimeter. The potential effects of grain face microstructure on crack healing became evident when the stability of flaws at single crystal-polycrystal interfaces was studied. For flaws that intersect multiple grains, grain boundary grooves will form. When the dominant transport mechanism for grooving is either surface diffusion [78,79] or volume diffusion [80], a shoulder or hillock develops adjacent to the groove, that extends above the original surface. These hillocks, if they become sufficiently large, can bridge the crack. This hillock-induced healing was observed, and bridging appeared imminent at other sites. The degree to which this mechanism can contribute to healing will be sensitive to the crack opening displacement variation with distance from the crack perimeter. The grain size will influence the spacing of the bridging sites, thus introducing a potential microstructural sensitivity to healing behavior. Profound effects of inadvertent impurities on the rate of healing were observed. Isolated cracks that we believe were contaminated with dust particles healed at rates that were estimated to be two to three orders of magnitude higher than those of "clean" cracks. This observation encouraged us to explore crack healing in doped or ion-implanted crystals as a means of studying the effects of specific crack face contaminants on transport during healing. It was

55 our intent to determine whether such studies could provide useful insights on the role that these impurities play in affecting microstructural development during sintering [81]. Studies of the morphological evolution of lithographically introduced pore channels (as shown in Figures 3c and 5) provided evidence of Rayleigh instabilities. However, important and reproducible differences existed between the observed characteristics of evolution, and those predicted for isotropic materials [20]. While this is hardly surprising- alumina has anisotropic surface energies - it suggested that similar measurements performed on doped crystals might provide information on how dopants modify the surface energy anisotropy. Determining whether parallels existed between the effects of dopants on sintering and crack healing became an objective. The possibility that crack healing experiments involving highly controlled crystallography cracks might provide evidence for anisotropic segregation [36-39] on processes whose kinetics are surface diffusion controlled emerged.

3.3.3 Crack healing in doped sapphire Work investigating crack healing in doped sapphire focused on three issues: 1) assessing the effects of controlled dopant additions on the morphological and kinetic characteristics of initial stage crack healing (cylinderization), 2) developing a framework for assessing the influence of surface energy anisotropy on the kinetics of pore channel evolution, and 3) using this framework to deduce changes in the degree of surface anisotropy from detailed experimental measurements of pore channel evolution. The dopants selected for investigation were Mg, Ca, and Ti. Ca and Mg were chosen because numerous studies investigating the surface segregation characteristics of Ca and Mg in alumina have been conducted [36, 82-84]. The results of these studies suggest that Mg segregates isotropically, whereas Ca segregates anisotropicaUy. Of specific interest is the observation that Ca does not appear to segregate to the basal plane, but does segregate to other low index planes. These results suggest that Mg and Ca additions would modify the surface energy anisotropy of sapphire in different ways. Ti is reported to segregate to grain boundaries in alumina [85], however, no experimental study of Ti segregation to free surfaces was found. The effects of Mg additions on surface diffusion have been investigated by several groups using a variety of methods [22, 86-89]. Some studies show a decrease in the surface diffusivity [87], while others suggest little change [86] or a modest increase in the diffusivity [22, 88, 89]. To our knowledge, no systematic study investigating the effect of either Ca or Ti impurities on surface diffusion in alumina has been conducted. Numerous comparative studies of microstructural development in undoped and MgO-doped alumina have been performed [8, 90], and changes in intragranular pore shapes [8] and in grain-boundary/surface dihedral angle distributions [90] suggest that MgO has a homogenizing effect on surface and grain boundary energies. Several recent studies have focused on assessing the effects of Ca impurities, either acting alone or in conjunction with Si impurities, on microstructural evolution [37, 90, 91]. Baik et al. have examined microstructural evolution and grain boundary chemistry in alumina containing only 100 ppm CaO-dopant and in samples co-doped with 300 ppm of MgO [37].

57 The MgO additions appeared to homogenize and reduce the level of Ca detected at intergranular fracture surfaces. Subsequent work indicates that Ca additions lead to a more facetted grain structure than is evident in similarly heat-treated undoped and pure alumina, and more anisotropic pore shapes than in undoped or MgO co-doped samples [92]. Moreover, CaO doping alone was not able to prevent pore-boundary separation. Several studies indicate that Ti additions promote the sintering of alumina, increasing the initial stage sintering rate significantly [93, 94], and facilitate the fabrication of dense compacts. A defect model [95] attributing the enhancement to the formation of aluminum vacancies as the charge compensating species for substitutional titanium ions has been proposed. A more recent defect model proposed by KrSger [96] requires a substantial concentration of acceptors. Morgan and Koutsoutis have suggested an entirely different "defecC model [97]. Their results suggest that the enhanced densification may be the result of liquid phase formation. They argue that commercial aluminas used may have contained sufficient Na impurity to cause liquid phase formation at the sintering temperatures explored. In view of the sensitivity of crack healing behavior to changes in crack geometry, crystallography, and impurity content, one of the major goals of sample preparation was to fabricate specimens that were as identical as possible in every respect except the dopant used. To achieve this goal, ion implantation was used in combination with the lithographic processing techniques described previously. Ion implantation was selected as the doping method because an extremely well-defined quantity of dopant can be introduced into a material, the resulting concentration profile can be predicted accurately, and the profile can be modified or reproduced as needed. Ion implantation provides a means of exploring the effects of a broad range of impurity additions, and thus, identifying dopants that merit the preparation of more costly melt-grown crystals. Two stage implants were performed to produce a more uniform impurity distribution than would be possible using a single-stage implant. The doses were well below those necessary to amorphize sapphire, and thus, this type of implantation damage was not a concern. Figure 11 shows the predicted low energy implant, high energy implant, and total dopant profiles for Mg-implanted samples; implant conditions for Ca and Ti were adjusted to produce similar total implant profiles. Peak concentrations are =300-350 ppm. The mask illustrated in Figure 2 was exposed four times on the surface of the samples, resulting in the simultaneous introduction of cracklike and channel-like features into both doped and implanted basal plane sapphire. After bonding, samples were annealed at 1700~ A qualitative comparison of the healing behavior of the larger defects was performed, and quantitative measurements of the instability characteristics of the channel-like defects were made. The results indicate that the three dopants induce major differences in the healing behavior. For Mg implanted sapphire, crack healing followed a morphological pattern similar to that observed in undoped sapphire. Edge to edge differences in the resistance to perturbation were evident, suggesting that surface energy anisotropy remains an important factor in the initial stages of crack healing, even in Mg-doped material. Compared to undoped sapphire, the cracks healed

58 Mg Implant Depth Profile 10 z~ u

O r

4-1

10 ~

~ - ~"

1018

0

o .m 4-J

I t

l

t~ 4-J

10 4 ~

\

1016

il

o

.......... M g - 50 keV

~ .....

~9

!

M g - 140 keV 10 -s

" M g - Total 1014 0

50

100

150

200

250

300

Depth (nm) Figure 11

Predicted low and high energy implant Mg impurity profiles resulting from two-stage implant, and the total Mg impurity profile.

in approximately half the time, consistent with the suggestion that Mg increases the rate of surface diffusion. Ca had an unexpectedly large accelerating effect on the rate of crack healing. Whereas =260 min of annealing at 1700~ was required to transform virtually all the 0.15-0.16 ~m deep cracks in Mg-implanted sapphire completely into isolated pores and low aspect ratio pore channels, this same degree of healing was achieved in _<40min in Ca-implanted sapphire. Isolated observations of crack bridging by facetted protrusions suggest that the "high" surface concentration of Ca induced by etching the implanted material may destabilize the (0001) surface and induce surface roughening. This is qualitatively consistent with the failure to observe Ca segregation to the (0001) surface, and evidence of strong segregation to other low index surfaces. This may in turn make a facetted nonplanar surface energetically favorable in comparison to the doped (0001) plane. Subsequent work involving samples with lower Ca surface concentrations shows no evidence of such rapid healing. In unimplanted samples, crack edges (and particularly pore channels) lying parallel to the [ 1120] direction were relatively more resistant to instabilities than those lying parallel to the [1100] direction. Ti additions reverse this relative stability, and stabilize crack edges lying parallel to the [ 1100] direction to such an extent that healing due to growth of edge instabilities was never observed. Rectangular cracks with long edges lying parallel to the [1100]

59 direction healed entirely by edge regression. Figure 12 shows cracks with similar geometry in Mg-, Ca-, and Ti-implanted sapphire, respectively, aider similar duration anneals at 1700~ The comparatively rapid healing of Ca-implanted sapphire, and the effect of [ 1T00] edge stabilization in Ti-implanted sapphire are evident. In summary, each dopant has a strong effect on the healing process, and the effects are reproducible and different in each case. These differences are also evident when the morphological evolution of pore channels is examined and behaviors are compared. Measurements of the morphological evolution of pore channels have the potential to provide a broad range of useful and fundamental information relating to the energetic and kinetic properties of ceramic surfaces, and their sensitivity to impurities. Analyses of the morphological instability of continuous phases, such as those presented by Plateau [98], Rayleigh [68], Nichols and Mullins [70], and others subsequently, focus on the stability of a cylindrical body with isotropic surface free energy to sinusoidal perturbations of infinitesimal amplitude. Two key quantities emerge from such analyses. The first quantity imposes a thermodynamic limit on the process, while the second focuses on the kinetics of evolution. Plateau was the first to determine the thermodynamic minimum or critical wavelength ~cr/t that must be exceeded if there is to be a driving force for the amplitude to increase [98]. In an isotropic system, infinitesimal perturbations with wavelength ~ < 2~R (the cylinder radius) are predicted to decay. Only perturbations with a wavelength >kcrit, i.e., >2uR, increase in amplitude. The magnitude of the interfacial area (and energy) reduction increases with increasing ~. Rayleigh subsequently analyzed the mode of maximum instability, and thus, determined the perturbation wavelength that would dominate the breakup [68] of liquid jets (9.02R) and of hollow cylindrical jets in an inviscid fluid (12.98R) [99]. Nichols and Mullins extended the method of Plateau and Rayleigh and evaluated the mode of maximum instability for solid cylindrical rods [71], and for the breakup of cylindrical voids in a solid via surface or lattice diffusion [70]. The results for the cylindrical void are pertinent to crack healing. For a cylindrical void with isotropic surface energy, both the perturbation growth rate and the dominant wavelength ~.maxdepend upon the relative contributions of surface and lattice diffusion to breakup [70]. For surface diffusion dominated breakup ~max is 8.89R(=~'~crit); lattice diffusion dominated breakup has a ~max of 12.96R The ultimate pore spacings are expected to reflect the relative contributions of these transport mechanisms to pore channel breakup. Surface diffusion is expected to dominate the evolution of sufficiently fine scale features [3]. In dete _rmining ~.max, it is commonly assumed that the perturbation amplitude is small. Thus, the predictions of the kinetic analyses are strictly valid only for the initial stages of perturbation growth. However, since the perturbation amplitude increases exponentially with time, Rayleigh and others have assumed that the wavelength that is initially dominant will dominate

Figure 12

Comparison of crack healing behavior in a) M g , b) Ca-, and c) Ti-implanted sapphire. The annealing time at 1700°C was 140 min for the Ca- and Mg-implanted sample, and 200 min for the Ti implanted sample. Directions are e = [llzOl and g = 11TOOI.

61 throughout breakup. Therefore one expects the final spacing of discrete pores (or particles) to reflect ~,max. For surface diffusion controlled breakup one would thus expect the normalized perturbation wavelength and normalized pore spacing, k]R, to assume a value of 8.89. A numerical model treating the breakup of rods in directionally solidified eutectic composites by diffusion along the rod/matrix interface supports this view. The model predicts that for interfacial diffusion, rods are most likely to pinch off at a characteristic wavelength times the rod circumference [ 100]. (The solution for interfacial diffusion in such a solid-solid system is expected to produce the same result as an analysis of surface diffusion for the solid-vapor system.) A more recent analysis by Hackney [101] suggests that even a perturbation with ), = ~.max is unstable to a periodic distortion of the wavelength; this provides a mechanism for broadening and modifying the node spacing distribution and the ultimate pore spacing distribution. In the aforementioned analyses, the surface free energy 7 is independent of surface orientation, as illustrated schematically in Figure 13a. As a result, one can focus on the effects of the perturbation induced shape changes on curvature, local chemical potential, and surface area. In anisotropic systems, changes in surface orientation result in local changes in 7, and as a result, modify the breakup behavior, as illustrated in Figure 13b. Cahn [74] evaluated the stability of single crystal rods with a specific surface energy 7 that is isotropic in the plane transverse to the cylinder axis (the R-0 plane), but a function of ~ = (~R/~z) where z is the axial coordinate. The analysis indicates that surface energy anisotropy can have either a stabilizing effect (kcr/t > 2xR) or a destabilizing effect (kcr/t < 2=R). Whether kcr/t increases or decreases relative to 2~R depends upon the sign of the second derivative of 7 with respect to r Specifically, 1

(4)

rt > JJ

When (D2Tfj~2) is positive, as illustrated in Figure 13b, the perturbation increases the "average" surface energy per wavelength segment relative to that of the unperturbed cylinder. Thus, a longer wavelength perturbation, one that reduces the total surface area sufficiently to compensate for the increase in "average" surface energy, is necessary to allow amplitude growth. In this situation, kcr/t will exceed 2xR. Conversely, when (~2TD~2) is negative, ~,crit will be <2~R. Since ~,max must exceed kcrit, parallel shifts in the kinetically dominant wavelength are to be expected. In recent work by Stflken [ 102], the effects of surface energy anisotropy on the kinetics of evolution of the same radially isotropic, axially anisotropic system considered by Cahn were determined. The kinetically dominant wavelength for surface diffusion controlled evolution is 1

(5)

52

Figure 13

a) Illustration of an unperturbed reference cylinder, and b) a perturbed cylinder with isotropic surface free energy. Changes in the surface orientation, as represented by the changing orientation of the surface normal, results in no change in T from the value characterizing the mean surface orientation, c) Illustration of a perturbed cylinder with axially anisotropic surface free energy. For the case shown, y increases with deviations of the surface orientation from the mean orientation, and as a result, ~ increases relative to the value in an isotropic system.

63 Thus, the same ~ relationship between the kinetic maximum and thermodynamic minimum wavelength found in isotropic system is predicted to apply to this simple anisotropic system. Experimentally, it is possible to measure periodicities or spacings that reflect the kinetic maximum by measuring the perturbation wavelength prior to the formation of isolated pores or particles, or the ultimate pore or particle spacing, or both. These spacings can be normalized by the channel radius, and the resulting ratio can be compared to theoretically predicted values. Our experience with undoped sapphire indicates that values of k/R always exceed the value of 8.89 predicted for breakup of pore channels by surface diffusion in isotropic solids. For pore channels lying in the basal plane of undoped sapphire with the cylinder axis oriented parallel to the [ 1120] and [1T00] directions, the normalized values of ~/R are of the order of 135 and 45, respectively. It was therefore of interest to determine whether these values would change as a result of dopant additions, and whether the results were consistent with the findings of studies focusing on their effect on sintering. The results indicate that Mg doping does indeed have a homogenizing effect on pore channel evolution. As indicated in Figure 14, the addition of Mg leads to a much more nearly isotropic behavior; cusps are softened. Interestingly, the values do suggest a residual effect of surface energy anisotropy, consistent with the observations made during the earlier stages of crack healing in Mgimplanted material. In addition, breakup occurs more rapidly. Thus, there appears to be a good correlation between the findings of sintering studies and other work assessing the effect of Mg additions. The observation that the energetics and rate of evolution are changed by the dopant addition points to a potential problem in interpretation of the change in the surface diffusivity. The fluxes that are measured experimentally reflect a diffusivity-driving force product. Clearly, attributing a change in a flux to solely a change in the diffusivity can lead to erroneous conclusions. Ca was also effective in reducing the barriers to breakdown arising from surface energy anisotropy, reducing the value of k/R for channels oriented parallel to the [1120] direction to --60-65, as illustrated in Figure 15. However, in this case, a stronger residual effect of surface energy anisotropy is indicated. The evolution of equivalent size features in Ca-implanted sapphire is more rapid than that in Mg-implanted sapphire. Ignoring changes in the energetics of evolution, this would indicate a higher diffusivity. Since the energetic barrier to evolution is greater in Ca-implanted material, the results indicate this even more strongly. It would thus be interesting to use the techniques described for measuring peak pore velocities to determine whether an enhancement in the pore mobility is evident. Ti had the effect of reducing the stability of defects oriented parallel to the [11~0], evidenced by a reduction ofk/R from values of=135 to values of--80. At the same time, channels oriented parallel to the [1100] direction were stabilized to such an extent that as was the case for cracks with edges oriented along this direction, the channels never developed any periodic perturbations and never broke up into discrete pores. In the context of the simplified model that has been

64 proposed, for channels of this orientation, segregation of Ti must result in the development of a sharp cusp. We acknowledge that the model of pore channel evolution in anistropic solids used to interpret the observations represents an oversimplification in that the channels introduced will form facetted cross sections. Studies of crosssections of channels prior to the onset of instability are in progress. It is our hope that we will be able to relate changes in stability, to changes in the (facetted) cross section. Such results may help to identify specific surfaces that appear to be stabilized by the addition of specific impurities. These results could be a useful companion to computer calculations of segregation behavior, and may provide a basis for testing predictions of surface-specific segregation tendencies. More detailed measurements of the orientation dependence of crack and channel stability are also in progress.

I

20

I

I

I

Mean = 1 7

15

[1 i-00] Mean = 2 6

10

[1120]

I

0

-

I

!

!

10

20

30

-

I

I

40

50

Normalized Pore Spacings ()dR) Figure 14

Comparison of normalized pore spacings derived from 2 ~m wide channels in Mg-implanted sapphire. Data from channels parallel to the [1100] and [1120] directions show much more similar behavior than observed for channels of identical orientations in unimplanted sapphire and in zero misorientation bicrystals. E

65

Normalized Pore Spacing (~/R) Figure 15

Normalized pore spacings for 2 pm x 0.16 pm (effective radius R = 0.32 ttm) channels in Ca-implanted oriented parallel to the [ 1120] direction. Values represent frequency of values within the indicated data collection interval. Smooth curve rendering of data is also shown.

SUMMARY AND CONCLUSIONS This paper has focused on describing the potential for using lithographybased experimental methods to study both the energetics and transport properties of solid-vapor, solid-liquid, and solid-solid interfaces. As in a prior paper devoted to this general topic [35], an effort has been made to highlight some of the more recent contributions to the literature that have brought new theoretical insights to the problem of microstructural evolution in ceramics, or that have provided new tools for studying specific aspects of microstructural evolution with more rigor than had previously been possible. Grain growth in theoretically dense undoped and MgO-doped polycrystalline alumina was studied and average grain boundary migration rates were determined. Through the use of lithographically introduced reference markers, it became possible to measure the migration rates of a-plane and cplane sapphire into the same undoped and MgO-doped materials with great precision and accuracy. A grain size dependent grain boundary mobility grain boundary energy product, MbTb is indicated, and the grain size dependencies of the MbTb products for seed and matrix grains differ. The ability to independently

66 control and vary the orientation of abnormal grain and to control and vary the chemistry of the polycrystalline matrix was used to isolate the effect of seed crystallography on migration behavior. Results suggest that as in metals, soluteboundary interactions are interface-structure dependent. The method may facilitate the development of a more complete understanding of solute-boundary interactions in ceramics. Microdesigned interfacial pore structures were utilized to study pore drag and pore-boundary separation in alumina. This approach allows the creation of pore arrays containing pores of controlled size and spacing at well-defined single crystal seed/polycrystalline matrix interfaces, and enables experimental determination of the peak pore velocity. From the peak pore velocity, values of the surface diffusion coefficient pertinent to sintering can be extracted. Doping with 250 ppm MgO resulted in an --4x enhancement of the surface diffusivity at 1600~ Quantitative agreement between theoretically predicted and experimentally observed separation/attachment conditions was obtained. This encourages the extension of the method to the study of other ceramics, other dopants, or both. In crack healing studies, the ability to control the crack geometry and crystallography was exploited to isolate the effects of these variables on crack behavior. The importance of these variables was demonstrated in crack healing studies performed on undoped sapphire. More recent research has exploited the ability to modify the surface chemistry using ion implantation to isolate the effects of specific impurities. Results clearly indicate that Mg, Ca, and Ti impact crack healing in distinct ways. A simplified analysis of pore channel evolution in anisotropic materials provides a framework for evaluating the effects of these impurities on the morphological evolution. Crystallographic characterization of pore cross-sections is in progress in the hope of linking the observed changes in breakdown characteristics to changes in the facet structure of pore channels. This may provide valuable information on the segregation characteristics of impurities to specific surfaces. It is hoped that the experimental method will provide a useful means of probing the energetic and kinetic characteristics of surfaces at elevated temperature.

ACKNOWLEDGEMENTS J. W. RSdel, J. D. Powers, J. S. StSlken, A. S. Nickles, and H. D. Ackler performed the research described in this paper, and their contributions are gratefully acknowledged. Research support for the work of J. W. RSdel and H. D. Ackler was provided by the Director, the Office of Energy Research, Office of Basic Energy Sciences, Materials Sciences Division of the U.S. Department of Energy under Contract No. DE-AC03-76SF00098. Research performed by J. D. Powers, J. S. StSlken and A. S. Nickles was supported by the National Science Foundation under Grant No. DMR-8821238. The work described would not have been possible without the assistance of the technical staff of the Microfabrication Laboratory in the Department of Electrical Engineering and Computer Science at the University of California at Berkeley. I have benefitted greatly from numerous discussions with too many colleagues to list individually. Discussions with Rowland Cannon over a period of many years have been particularly stimulating and beneficial.

57 REFERENCES ,

2. 3. 4. ,

6. .

8.

.

10.

11. 12. 13. 14. 15. 16. 17. 18. 19. 20.

21. 22. 23. 24. 25. 26. 27.

M. F. Yan, Mater. Sci. and Eng., 48, [1], 53-72 (1981). R. J. Brook, Proc. Br. Ceram. Soc., 32, 7-24 (1982). C. H. Herring, J. Appl. Phys., 21, [4], 301-03 (1950). C. H. Hsueh, A. G. Evans and R. L. Coble, Acta Metall., 30, [7], 1269-79 (1982). R. J. Brook, J. Am. Ceram. Soc., 52, [ 1], 56-57 (1969). R. J. Brook, pp. 331-64 in Treatise on Materials Science and Technology, Vol. 9, Ed. F. F. Y. Wang (Academic Press, New York, 1976). F. M. A. Carpay, J. Am. Ceram. Soc., 60, [1-2], 82-83 (1977). S. J. Bennison and M. P. Harmer, pp. 13-49 in SINTERINGOF ADVANCED CERAMICS, edited by C. A. Handwerker, J. E. Blendell and W. A. Kaysser, The American Ceramic Society, Columbus, OH (1990) T. K. Gupta, pp. 750-66. in ADVANCESIN CERAMICS, Vol. 10, W. D. Kingery, Ed., The American Ceramic Society, Columbus, OH (1984). F. H. Huang, R. A. Henrichsen and Che-Yu Li, "A Study of Capillarity and Mass Transport on the A1203 Surface," pp. 173-86 in Mat. Sci. Res., Vol. 10, Sintering and Catalysis, Edited by G. C. Kuczynski, Plenum Press, New York, 1975. R. S. Gilmore, K. C. Tam, J. D. Young and D. R. Howard, Phil. Trans. R. Soc. Lond. A, 320, 215-235 (1986). M. McKellar and N. C. Wardlaw, J. Can. Petr. Technol., 21, [4], 39-41 (1982). M. Kahn and B. Kriese, in Electronic Packaging Materials Science II, Mat. Res. Soc., 72, 35-40 (1986). M. Kahn, A. Dalzell and B. Kovel, Adv. Ceram. Mat., 2, [4], 836-40 (1987). K. Burger, W. Mader and M. Rtihle, Ultramicroscopy, 22, [ 1], 1-14 (1987). J. RSdel and A. M. Glaeser, J. Am. Ceram. Soc., 70, [8], C172-C175 (1987). J. RSdel, Ph.D. thesis, University of California, Berkeley (1988). J. RSdel and A. M. Glaeser, Materials Letters, 6, [10], 351-55 (1988). J. RSdel and A. M. Glaeser, J. Am. Ceram. Soc., 73, [3], 592-601 (1990). J. RSdel and A. M. Glaeser, pp. 243-257 in SINTERING OF ADVANCED CERAMICS, edited by C. A. Handwerker, J. E. Blendell and W. A. Kaysser, The American Ceramic Society, Westerville, OH (1990). J. RSdel and A. M. Glaeser, J. Am. Ceram. Soc., 73, [11], 3293-301 (1990). J. RSdel and A. M. Glaeser, J. Am. Ceram. Soc., 73, [11], 3302-12 (1990). J. RSdel and A. M. Glaeser, J.Ceram. Soc. Japan, 99, [4], 251-65 (1991). J.W. Bullard, A.M. Glaeser and A. W. Searcy, Mater. Res. Soc. Proc., 238, 247-52 (1992). J. D. Powers and A. M. Glaeser, J. Am. Ceram. Soc., 75, [9], 2547-58 (1992). J. D. Powers and A. M. Glaeser, in press, J. Am. Ceram. Soc.. J. D. Powers and A. M. Glaeser, to be submitted to J. Am. Ceram. Soc..

58 28. 29. 30. 31 32. 33. 34. 35. 36. 37. 38. 39.

40. 41. 42. 43. 44. 45. 46. 47. 48. 49.

R. W. Gilmore, H. D. Ackler and A. M. Glaeser, unpublished research. H. D. Ackler and A. M. Glaeser, unpublished research. R. A. Colclaser, MICROELECTRONICS: PROCESSINGAND DEVICE DESIGN, John Wiley and Son, New York, New York (1980). R. S. Gilmore, A. M. Glaeser and J. C. Wade, to appear in NDE FOR ENGINE COMPONENTSAND RELATED MATERIALS(Proceedings of the 1993 A S M E TurboExpo to be held in Cincinnati, OH; May 24-28th.) E. DSrre and H. Htibner, ALUMINA, Springer Verlag Berlin 1984. G. Galvin, MRS Bull., 13, (7), 31-34 (1988). B. McIntyre, R. Perman, and B. Gold, R&D, 44-48, (December, 1989). A. M. Glaeser, pp. 287-322 in THE SCIENCE OF CERAMIC INTERFACES, edited by J. Nowotny, Elsevier Press, Amsterdam, Holland (1991). S. Baik and C. L. White, J. Am. Ceram. Soc., 70, [9], 682-88 (1987) S. Baik and J. H. Moon, J. Am. Ceram. Soc., 74, [4], 819-22 (1991). R. C. McCune, W. T. Donlon and R. C. Ku, J. Am. Ceram. Soc., 69, [8], C 196-C 199 (1986). a) M. J. Davies, P. R. Kenway, P. J. Lawrence, S. C. Parker, W. C. Mackrodt and P. W. Tasker, J. Chem. Soc. Faraday Trans. 2, 85, [5], 55563 (1989); b) W. C. Mackrodt and P. W. Tasker, J. Am. Ceram. Soc., 72, [9], 1576-83 (1989). A. J. Forty and C. T. Forwood, Trans. Brit. Ceram. Soc., 62(a), 715-25 (1965). S. E. Bronisz and D. L. Douglass, Phys. Stat. Sol., 29, K95-97 (1968). C. F. Yen and R. L. Coble, J. Am. Ceram. Soc., 55 [10], 507-09 (1972). B. J. Hockey and B. R. Lawn, J. Mater. Sci., 10, 1275-84 (1975). T. K. Gupta, pp. 354-65 in MICROSTRUCTURES'76, R. M. Fulrath and J. A. Pask, eds., Westview Press, Boulder CO, (1977). S. M. Park and D. R. O'Boyle, J. Mater. Sci., 12, 840-41 (1977). T. K. Gupta, J. Am. Ceram. Soc., 61 [5-6], 191-95 (1978). O. Maruyama and W. Komatsu, Ceramurgica International, 5 [2], 51-5 (1979). S. H. Hickman and B. Evans, Phys. Chem. Minerals, 15, 91-102 (1987). J. R~del and A. M. Glaeser, pp. 485-90 in INTERFACIALSTRUCTURES, PROPERTIES AND DESIGN, (Mater. Res. Soc. Proc., 122, Pittsburgh, PA

1988). 50. 51. 52. 53. 54. 55. 56. 57. 58. 59.

A. H. Heuer and J. P. Roberts, Proc. Brit. Ceram. Soc., 6, 17-27, (1966). F. P. Mallinder and B. A. Proctor, Proc. Brit. Ceram. Soc., 6, 9-16, (1966). L. M. Davies, Proc. Brit. Ceram. Soc., 6, 29-35, (1966). F. F. Lange and T. K. Gupta, J. Am. Ceram. Soc., 53, [1], 54-55 (1970). F. F. Lange and K. C. Radford, J. Am. Ceram. Soc., 53 [7], 420-21 (1970). J. T. A. Roberts and B. J. Wrona, J. Am. Ceram. Soc., 56 [6], 297-99 (1973). R. Dutton, J. Am. Ceram. Soc., 56 [12], 660-61 (1973); ibid., 59 [1-2], 88 (1976). T. K. Gupta, J. Am. Ceram. Soc., 58 [3-4], 143 (1975). R. Raj, W. Pavinich and C. N. Ahlquist, Acta Metall., 23 [3], 399-403 (1975). T. K. Gupta, J. Am. Ceram. Soc., 59 [5-6], 259-62 (1976).

69 60. 61.

62. 63. 64. 65. 66.

67. 68. 69. 70. 71. 72. 73. 74. 75. 76. 77. 78. 79. 80. 81. 82.

83. 84. 85. 86.

87. 88. 89. 90.

G. Bandyopadhyay and J. T. A. Roberts, J. Am. Ceram. Soc., 59 [9-10], 415-19 (1976). G. Bandyopadhyay and C. R. Kennedy, J. Am. Ceram. Soc., 60 [1-2], 48-50 (1977). A. G. Evans and E. A. Charles, Acta Metall., 25 [8], 919-27 (1977). R. N. Singh and J. L. Routbort, J. Am. Ceram. Soc., 62 [3-4], 128-33 (1979). F. F. Lange and D. R. Clarke, J. Am. Ceram. Soc., 65 [ 10], 502-06 (1982). W. C. Carter and A. M. Glaeser, J. Am. Ceram. Soc., 67 [6], C 124-C 127 (1984). M. D. Drory and A. M. Glaeser, J. Am. Ceram. Soc., 68 [1], C14-C15 (1985). C. Scott and V. B. Tran, Am. Ceram. Soc. Bull., 64 [8], 1129-31 (1985). Lord Rayleigh, Proc. London Math. Soc., 10, 4-13 (1879). C. S. Smith, Trans. A.I.M.E., 17511], 15-51 (1948). F. A. Nichols and W. W. Mullins, Trans. A.I.M.E., 233 [10], 1840-48 (1965). F. A. Nichols and W. W. Mullins, J. Appl. Phys., 36 [6], 1826-35 (1965). A. D. Brailsford and N. A. Gjostein, J. Appl. Phys., 46 [6], 2390-97 (1975). F. A. Nichols, J. Mater. Sci., 11 [6], 1077-82 (1976). J. W. Cahn, Scripta Metall., 13111], 1069-71 (1979). W. C. Carter and A. M. Glaeser, Acta Metall., 35 [1], 237-45 (1987). A. M. Glaeser, Materials Letters, 5 [7/8], 239-245 (1987). W. M. Robertson and F. E. Ekstrom, pp. 273-83 in Materials Science Research, Vol. 3, edited by W. W. Kriegel and Hayne Palmour, III, (Plenum, New York, 1966). W. W. Mullins, J. Appl. Phys., 28, [3], 333-39 (1957). W. M. Robertson, J. Appl. Phys., 42, [1], 463 (1971). W. W. MuUins, Trans. A.I.M.E., 218, 354-61 (1960). J. E. Blendell and C. A. Handwerker, J. Cryst. Growth, 75, [2], 138-60 (1986). S. Baik, D. E. Fowler, J. M. Blakely and R. Raj, J. Am. Ceram. Soc., 68, [5], 281-86 (1985). S. Baik, J. Am. Ceram. Soc., 69, [5], C101-C103 (1986).\ S. M. Mukhopadhyay, A. P. Jardine and J. M. Blakely, J. Am. Ceram. Soc., 71, [5], 358-62 (1988). C. W. Li and W. D. Kingery, pp. 368-78 in ADVANCESIN CERAMICS, Vol. 10, W. D. Kingery, ED., The American Ceramic Society, Columbus, OH (1984). J. M. Dynys, R. L. Coble, W. S. Coblenz and R. M. Cannon, pp. 391-404 in SINTERING PROCESSES (Materials Science Research, Vol. 13), G. Kuczynski, Ed., Plenum Press, New York, New York, (1980). C. Monty and J. Le Duigou, High Temp.-High Pressures, 14, [6], 709-16 (1982). K. A. Berry and M. P. Harmer, J. Am. Ceram. Soc., 69, [2], 143-49 (1986). S. J. Bennison and M. P. Harmer, J. Am. Ceram. Soc., 73, [4], 833-37 (1990). C. A. Handwerker, J. M. Dynys, R. M. Cannon and R. L. Coble, J. Am. Ceram. Soc., 73, [5], 1371-77 (1990).

70 C. A. Handwerker, P. A. Morris and R. L. Coble, J. Am. Ceram. Soc., 72, [1], 130-36 (1989). S. S. Kim and S. Baik, in Proceedings of the International Symposium on 92. the Science and Technology of Sintering, July 24-27, 1991, UBC, Vancouver Canada, Trans. Tech. Publ. Ltd., Aedermannsdorf, Switzerland, 1992. R. D. Bagley, I. B. Cutler, and D. L. Johnson, J. Am. Ceram. Soc., 53, [3], 93. 136-41 (1970). M. Harmer, E. W. Roberts, and R. J. Brook, Trans. J. Brit. Ceram. Soc., 94. 78, [1], 22-25 (1979). R. J. Brook, J. Am. Ceram. Soc., 55, [2], 114-15 (1972). 95. 96. F. A. KrSger, J. Am. Ceram. Soc., 67, [6], 390-92 (1984). P. E. D. Morgan and M. S. Koutsoutis, J. Am. Ceram. Soc., 68, [6], C15697. C158 (1985). Plateau as cited in S. Chandrasekhar, HYDRODYNAMICAND 98. HYDROMAGNETICSTABILITY, Dover Publications, Inc., New York (1981). Lord Rayleigh as cited in S. Chandrasekhar, HYDRODYNAMICAND 99. HYDROMAGNETICSTABILITY, Dover Publications, Inc., New York (1981). 100. T. F. Marinis, Ph.D. thesis, Carnegie-Mellon University, 1981. 101. S. A. Hackney, Scripta Metall. et Materialia, 25, [4], 799-804 (1991). 102. J. S. StSlken and A. M. Glaeser, Scripta Metall. et Materialia, 27, [4], 44954 (1992). 91.

Science of Ceramic Interfaces II J. Nowotny (Editor) 9 1994 Elsevier Science B.V. All rights reserved.

71

INTERFACES IN ZIRCONIA BASED ELECTROCHEMICAL SYSTEMS AND THEIR INFLUENCE ON ELECTRICAL PROPERTIES S.ES. Badwal and J. Drennan CSIRO, Division of Materials Science & Technology Private Bag 33, Rosebank MDC, Clayton 3169, Victoria, Australia Ceramic Fuel Cells Ltd Monash Science and Technology Park 710 Blackburn Road, Clayton, Victoria 3168 Abstract The interfaces play a critical role in establishing both electrical and electrochemical properties of zirconia based electrolyte cells. Various types of interfaces which exist in such systems have been described. Apart from gas/electrode and electrode/electrolyte interfaces, the role of interfaces within the grains (coherent growth of precipitates, compositional variations within grains and second phase inclusions) and at grain boundaries (intermediate phase formation from the matrix, intermediate phase of the impurity type, and inclusions and pores) of polycrystalline electrolyte materials has been elucidated. The effect of interfaces on electrical and electrochemical properties has been discussed. 1. INTRODUCTION Zirconia based materials, apart from their use in mechanical applications, are commonly used in solid oxide fuel cells for power generation; in oxygen sensors for combustion, pollution and process control; oxygen pumps for generation of oxygen or oxygen atmosphere control; and electrochemical reactors for production of chemicals [1-5]. The interfaces play a crucial role in establishing both electrical and electrochemical properties of zirconia-based electrolyte cells. In broad terms, processes contributing to overall charge transport through a solid electrolyte device consist of electrode reactions at both electrodes and ionic transport within the solid electrolyte. The electrode reactions take place at the electrode/electrolyte interface. The types of interfaces which contribute to electrode reactions are: gas/electrode, electrode/electrolyte, gas/electrolyte and three phase boundary between the gas phase, electrode and electrolyte. Impurity segregation can significantly modify the external surface of the electrolyte. The efficiency of electrode reactions apart from material properties and electrode morphology therefore depends significantly on the state of external surface of the electrolyte. Following the oxygen exchange reaction, ionic transport takes place within the solid electrolyte. For practical reasons, polycrystalline materials are invariably used and this further increases the number of interfaces within the system. Ionic transport can be divided into two main components: through the

72 grains and across grain boundaries. The grains in the ceramic may have inclusions, pores, precipitates of other phases, and compositional inhomogeneities. They all create additional interfaces which can influence ionic transport within the grains. Grain boundaries may have secondary phases, inclusions, pores and enhanced segregation of matrix components. Depending upon the conducting properties of these phases and their quantity and location, the effect on the electrical or electrochemical performance of the device may be quite significant. Any interface which hinders mass transport as a result of applied potential is detrimental to the performance of an electrochemical device. In this paper, the role of different types of interfaces which exist in zirconia based electrolyte cells has been considered and their influence on electrical and electrochemical properties has been discussed. The main sections under which interface properties have been discussed are: electrode reactions at the electrode/electrolyte interface; oxygen-ion transport within the solid electrolyte (interfaces within grains); and grain boundaries.

2. GENERAL COMMENTS A number of figures in this chapter are impedance diagrams. For readers not familiar with the topic a simple explanation is given below. In an electrochemical cell, as discussed above, overall mass transport consists of a number of different processes at the electrode/electrolyte interface and within the solid electrolyte. At the electrode/electrolyte interface, typical electrode processes are adsorption/dissociation, diffusion and charge transfer and within the electrolyte the main processes involve oxygen-ion migration through the lattice and across grain boundaries. Each of these processes has a different time constant associated with it and occurs over a different time or frequency domain and can be simulated by a network of resistors and capacitors. Figure 1 shows a simple electrical equivalent circuit of a solid electrolyte cell. The response of such a system in the frequency domain when represented in the impedance plane consists of a number of arcs each corresponding to a different process. This is shown schematically in Figure 1. Electrode processes, in general, have a large time constant and invariably relax over the low frequency domain followed by ion migration across grain boundaries and bulk diffusion of ions in the crystal lattice. One or more arcs are usually observed for electrode reactions in the complex impedance plane as shown in Figure 1 [6-10]. For a polycrystalline electrolyte, two arcs are commonly observed. The low frequency electrolyte arc is associated with the migration of oxygen ions across grain boundaries and build-up of space charge layers near the grain boundary surface region. In a single crystal or grain, this arc is not observed. The high frequency electrolyte arc is due to dipole relaxation within the bulk of grains. The difference of intercepts of each arc on the real impedance or x-axis defines the resistance or rate of that particular process. The time constant for each process can be determined from the frequency at the top of each arc. Due to selection of measurement conditions and availability of the frequency range or instrumentation, all arcs due to electrode or electrolyte processes may not be seen or only partly seen in the impedance plane. However, this does not imply that a process is not contributing to the overall mass transport in a solid electrolyte cell. Measurements over a wide frequency, temperature and gas concentration are required to fully elucidate the role of rate limiting processes. For more detailed explanation of impedance or admittance diagrams, the readers, for example, should read [7].

73 (a)

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T~ Electrolyte

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Z Figure 1. A simple schematic impedance plane representation of an electrochemical cell. The cell response is shown as consisting of two electrode (adsorption (Rd, C~d) and charge transfer (Re1, C1)) and two electrolyte (grain boundary (R_, C .) and lattice (Rt Cd) processes and is not a true representation for every electrode/~ectr~yte system. 3. ELECTRODE REACTIONS AT ELECTRODE/ELECTROLYTE INTERFACES

In zirconia based electrolyte systems the oxygen reduction reaction take place at the cathode/electrolyte interface. The oxygen ions thus produced transport through the electrolyte and discharge at the anode/electrolyte interface. The discharge reaction may be simple evolution of oxygen (in the oxygen pumping mode) or it may involve complex mechanisms where discharging oxygen ions react directly or indirectly with fuel species (fuel cell operation) or chemicals (electrochemical reactor) [1-4, 11-15]. The state of both cathode/electrolyte and anode/electrolyte interfaces and physical, chemical and thermal compatibility of the electrode and electrolyte interfaces are critical for long term and satisfactory performance of the device. At the cathode oxygen reduction takes place through the normal processes of gaseous diffusion, adsorption/dissociation at the electrode or electrolyte surface, diffusion of adsorbed electroactive species to the reaction site and charge transfer reactions. The adsorption/dissociation reaction may proceed via simultaneous transfer of charge at the gas/electrode interface. At the anode/electrolyte interface, oxygen ions discharge releasing electrons to the electrode in the normal mode of operation (e.g. oxygen pumping mode). However, in the presence of oxidisable gases (e.g. CO, HE, CH 4, other gaseous chemicals), the reaction mechanism is complex and may involve adsorption/dissociation of fuel species at active sites on the electrode or electrolyte surface [14,15]. The electroche-

74 mical processes taking place in a solid electrolyte cell are shown schematically in Figure 2. The nature of the electrode material, its electronic and ionic conductivity, surface area and porosity, play a significant role. For many electrode materials which have a low solubility for oxygen or which have low oxygen-ion mobility, the role of the three phase boundary (tpb) between the gas phase, electrode and electrolyte is critical [16,17]. The constraints imposed by the existence of tpb reduce electrode/electrolyte contact area and requires the electrode microstructure to be carefully optimised. For this reason electrode materials with high oxygen-ion conductivity are preferred as the requirement of tpb is considerably reduced. Also the adherence of the electrode material to the electrolyte is critical for better contact between electrode and electrolyte. The electrodes must have minimum thermal expansion mismatch with the electrolyte and other cell components. A large mismatch can lead to delamination or cracking of electrode layers causing loss of the contact area and poor definition of the interface. The electrode morphology also plays a

Figure 2. A schematic of electrochemical processes occurring in a solid electrolyte cell. significant role in defining reaction rates and can not be ignored. The electrode pore structure must allow reactive gases to diffuse through to reaction sites and allow unreacted gases and reaction products to escape freely without any hindrance. The main types of interfaces, baring the physical appearance (microstructure, porosity and surface area of the electrode) are: relatively clean interfaces (no impurity segregation or an intermediate phase formation); existence of an interphase due to segregation; existence of an interphase due to reaction/diffusion between cell components; and in the case of multi phase electrodes, interfaces between components of a composite electrode. Various types of interfaces which can have an affect on the electrode performance are schematically depicted in Figure 3. The role of electrode morphology and different types of interfaces

75 have been discussed below.

3.1 Electrode Morphology The microstructure, porosity and surface area of the electrode define the contact area between gas/electrode and electrode/electrolyte interfaces and play a crucial role in defining electrode kinetic rates both in the presence and absence of an interphase. The electrodes which have low oxygen-ion conductivity or low solubility for oxygen require optimum porosity for gas diffusion so that the maximum number of triple phase boundaries between gas, electrode and electrolyte are available. Noble metals (Ag, Au, Pt, Pd), especially Pt, are good catalytic electrodes for oxygen reduction, exhibit good stability in both oxidising and reducing environments and have been commonly used in oxygen sensors, oxygen pumps and fuel cells. However, these metals undergo sintering and grain growth when exposed to temperatures in the vicinity of 700 - 1000~ This leads to a substantial reduction in the gas/electrode interface and three phase contact area between gas/electrode/electrolyte (essential for oxygen charge transfer reaction) [18,19]. Figure 4 shows the microstructure of Pt paste and Pt electrodes sputter deposited on zirconia-yttria electrolyte and

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Figure 3. Various types of interfaces in solid electrolyte cells which affect electrode performance. (1) Segregation of electrolyte components; (2) impurity segregation at the surface of the electrolyte; (3) reaction products between electrode and electrolyte phases; (4) non-contact region (dead zone) for air or fuel electrode reactions; and (5) gas/electrode and clean electrode/electrolyte interfaces. given heat treatment at different temperatures in air. The grain growth and reduction in the porosity and three phase contact area with increasing temperature is clearly obvious and has a substantial effect on the interfacial electrode resistance. The effect of the microstructure on the electrode resistance is shown in Figure 5 for platinum electrodes which have low solubility for oxygen and require the presence of a large three phase contact boundary area between electrode, electrolyte and the gas phase. With decreasing electrode surface area and three phase boundary (due to the high temperature treatment), both the electrode

76

Figure 4. Microstructure of sputter deposited Pt: (a) & (b), and Pt paste: (c) & (d), electrodes on zirconia-yttria substrate and heat treated in air at 600~ (50h): (a) & (c) or 900~ (2h): (b) & (d).

resistance and the time constant associated with the electrode processes increase and different processes become rate limiting as shown in Figure 5 by the appearance of another arc on the low frequency end of the impedance spectrum. This low frequency arc is a function of the electrode surface area and is generally attributed to slow adsorption/dissociation of

77 oxygen at the electrode surface [20,21]. In general, oxide electrodes are more refractory and undergo less sintering and grain growth compared with metals. Nevertheless degradation in the electrode performance does occur with time. The common cause of degradation in oxide electrodes with low oxygenion conduction, in the absence of a chemical reaction with the electrolyte phase, is also sintering and grain growth leading to a reduction in the three phase boundary contact area. For example, Sr doped LaMnO~ is used as the air electrode in solid oxide fuel cells. Although the electronic conductivxty is high (about 190-200Scm 1 for (Lao ~Sro 2)MnO3 at 1000~ the oxygen-ion diffusion is extremely slow and the presence-ofthree phase boundary region and large porosity are crucial to low overpotential losses [22,23]. However, if oxygen electrodes (cathodes) were available with high bulk oxygen-ion diffusion rates as well as fast surface exchange kinetics at the gas/electrode interface, the role of three phase boundary and porosity would not be as critical and the contact area between electrode and electrolyte can be enhanced by using dense electrode materials. For such electrodes, oxygen exchange with the gas phase would take place at the gas/electrode interface followed by oxygen-ion transport within the electrode and transfer at the electrode/electrolyte interface. Several materials with the perovskite and related structures have high oxygen-ion diffusion rates and high oxygen-ion conductivity [24] and for these materials, the rate of the oxygen transfer reaction is expected to be established by the surface exchange properties of the electrode material. Carter etal [25] and Steele etal [26]

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78 have measured bulk diffusion coefficients for oxygen ions in the electrode and surface exchange coefficients at the gas/electrode interface for both lanthanum manganite and lanthanum cobaltite perovskite electrodes doped at A- and B-sites and have shown that for doped LaCoO 3 electrode materials, oxygen-ion diffusion rates are 4-6 orders of magnitude higher than those observed in lanthanum manganites with similar doping levels. These authors have demonstrated that although the surface exchange coefficients in cobaltite based perovskites are higher by one to three orders of magnitude compared with the manganites, the oxygen flux through the material is limited by the oxygen exchange coefficient at the gas/electrode interface. In lanthanum cobaltites, the net oxygen flux through the material (or the electrode reaction rate) is controlled by the surface exchange reaction whereas in manganites the process is bulk diffusion controlled. Thus it appears that optimisation of the microstructure is required for both types of materials. Similarly for Ni/zirconia-yttria cermet, which is the common fuel electrode material used in solid oxide fuel cells, the electrode microstructure plays a dominant role in determining the electrode performance at the interface. Not only fuel gases (H 2, CO) need to diffuse in to the electrode/electrolyte interface, but products of reaction (steam and CO~) are required to diffuse out as well. About 50% or higher porosity is required and pore size and distribution are crucial to the performance of the electrode. Zirconia-yttria is added to Ni to reduce the thermal expansion mismatch at the interface and to slow down sintering and grain growth of the Ni particles. Therefore, the distribution and the amount of Ni and zirconia-yttria phases need to be carefully controlled in the electrode to avoid performance degradation of the cell at the anode/electrolyte interface. The thermal expansion of Ni metal is significantly higher than that of zirconia-yttria. This combined with poor wetting at the Ni/electrolyte interface can lead to debonding or delamination at the interface, poor microstructure and higher overpotential losses [27,28]. 3.2 Relatively clean interfaces The electrode/electrolyte interfaces free of impurities and other phases can truly define the electrode kinetic behaviour of the system and are generally preferred unless the formation of another phase somehow facilitates electrode reactions (see below). However, for all practical purposes clean surfaces or interfaces are rarely encountered. Even in the absence of segregated impurities and secondary phases, particulates in the gas stream, unwanted reaction products (coke deposition in fuel ceils in carbon containing atmospheres) and impurities in the gas phase (sulphur containing compounds, metal vapours etc) can significantly affect the performance of electrodes. Hydrogen embrittlement occurring in metallic electrodes is another common problem affecting the electrode behaviour. For example, the adsorption of very small (ppm levels) quantities of fuel gases even in oxidising environments (in the reference air stream for example) is known to poison the electrode and has been shown to adversely affect the performance of oxygen sensors especially at lower temperatures [29-31]. In potentiometric sensors, large deviations from the Nernstian behaviour are commonly observed for Pt electrodes below about 400 - 500~ as shown in Figure 6. This type of behaviour is related to the presence of trace levels of CO and hydrocarbons in the gas stream and their adsorption on the surface of Pt electrodes giving rise to mixed potentials [29,31,32]. Sensors prepared with platinum electrodes with a low electrode surface area (heat treated at high temperatures: above about 800~ show much larger non-Nernstian behaviour [30] possibly due to high sensitivity of the reduced surface area electrodes to trace levels of combustibles in the gas stream. When Pt electrodes are replaced with materials such as nonstoichiometric oxides (e.g. (U,M)O2• ~) with apparently low adsorption for these gases or which facilitate reaction between trace mapur-

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3.3 Existence of interphase due to segregation The starting powders used to make zirconia ceramics invariably contain small amounts of impurities. These impurities are either deliberately added as sintering aids, or are present as contamination picked up during the processing or are present because of the excessive cost of cleaning the raw materials. During the sintering of zirconia ceramics, complex chemical reactions occur between adsorbed or segregated impurities and often involve bulk matrix components leading to the formation of glassy phases [33-35]. These phases are quite mobile (some with a melting point below 1000~ and move around rapidly in grain boundaries and tend to migrate to the external surface during sintering [36-37]. Even if the surfaces of zirconia ceramics are cleaned free of impurities by chemical etching or a mechanical grinding process, during subsequent heat treatments, annealing or use of the ceramics at high temperatures (800-1200~ more impurity phases migrate to the external surface from the grain boundary network. Hughes and Badwal, from x-ray photoelectron spectroscopy studies performed on zirconia-yttria ceramics with different impurity content have shown that most of the impurities which are present in the starting powders migrate to the external surface [38,39] during sintering or subsequent annealing. In addition, enhanced segregation of dopant (yttrium) is commonly observed at the sintered surface. Figure 7 shows the impurity segregation at the external surface of a 3 tool% Y.O 3 - ZrO. ceramic (previously sintered at 1500~ ground and polished) annealed at 1200~ for 50h. ZDepend_ ing on the quantity and type of impurities, significant surface layers can build up. Ceramic electrolytes prepared from commercial powders have been shown to have texture at the

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S p u t t e r i n g Time, Sec Figure 7. Yttrium and impurity segregation measured by x-ray photoelectron spectroscopy at the external surface of a 3 tool% Y203 - ZrO 2 ceramic (the sintered surface cleaned by grinding and polishing after sintering at 1500~ annealed at 1200~ for 50 h in air. OY/Zr; A- Si/Zr; O- Na/Zr; v- Ti/Zr ratios. Figure courtesy of A.E. Hughes. sintered surface which differs from the bulk [40]. Figure 8 shows scanning electron micrographs taken of the sintered (1500~ 4h) surface of two 8 tool% Y20. - ZrO 2 ceramics. The presence of the glassy phase at the external surface is clear. Tl~e amount of the glassy phase at the external surface is a function of the impurity content, types of impurities present and sintering conditions (temperature, time, gas atmosphere). During the sintering process as the impurities migrate to the external surface, some of these (e.g. alkali metal oxides) have high vapour pressures at sintering temperatures and volatilise rapidly from the external surface. Ciacchi et al [40] have reported the precipitation of Y203 at the as-sintered surface of some zirconia-yttria ceramics especially those with a high yttna content. This was attributed to the migration of a Y-rich glassy phase to the external surface followed by its decomposition on volatilisation of some components. Similar impurity migration to the external surface has also been reported by other authors [41,42]. For example, Chaim et al [41] reported migration of an iron-rich silicate phase which also contained significant amounts of yttrium, to the external surface of ZrO 2 - 4 wt% Y203 (a tetragonal phase) on annealing pre-sintered ceramic specimens. This glassy phase following migration to the external surface reacted with surface grains leading to dissolution of the tetragonal zirconia phase grains and re-precipitation of rather larger YEO3-rich cubic zirconia grains containing about 10 wt% Y203. The glassy phase was then nearly depleted of Y203, C a t and Na20 and was very much enriched in Fe203 but was reported to be present in large quantities at the external surface coveting the heat treated surface. The depletion of Na20 occurred from the glassy phase probably due to volatilisation and those of Y203 , C a t due to their dissolution in the zirconia grains [41]. Drennan and Hannink [43] have also reported the migration of the glassy phase to the external surface in the case of MgO- partially stabilised zirconia, a tough and strong

81

Figure 8. Scanning electron micrographs of the external surface of two sintered (1500~ 4h) ceramics of 8 tool% Y203 - ZrO 2 composition showing the glassy phase segregation (dark areas). ceramic with high wear resistance properties. They have discussed the beneficial effects of this migration in terms of bulk mechanical properties. The existence of SiO2-based impurity phase at grain boundaries leads to the formation of MgzSiO4 thus depleting stabiliser from the surface of grains resulting in precipitation of excessive amounts of monoclinic phase at the grain boundaries which like the glassy phase itself has poor mechanical properties. On addition of a small quantity of SrO (in the form of SrCO 3) to starting powders, a low melting glassy phase was formed at interfaces between grains (grain boundaries) which during the sintering of the ceramic migrated to the external surface leading to cleaning of the grain boundaries and counteracting the formation of MgESiO4. Significant improvements to bulk mechanical properties were reported [43] due to control of the grain boundary chemistry. One of the striking feature of this impurity migration to the external surface is the grain growth and the phase redistribution they can cause at the external surface. Thus the microstructure and phases present at the external surface of the ceramic may be substantially different from those in the bulk. Chaim et al [41] reported that migration of the impurity phase to the external surface caused substantial grain growth of the external grains in addition to redistributing its contents. As shown in Figure 9, exaggerated grain growth has been reported in a zirconia-yttria ceramic (with 3 tool% YzO3) heated in contact with a glassy phase [44]. The constituents of the glassy phase were the same as that present in the bulk of the ceramic (equivalent of impurity contents in the starting powder). The glassy phase apparently helps redistribute the solute. The larger grains in zirconia-yttria ceramics are usually associated with the presence of a cubic phase with higher amounts of Y203 and smaller tetragonal grains with a lower concentration of Y.O 3 [45,46]. Thus the grains at the external surface (in Figure 9 and in the results reported" by Chaim et al [41]) are rich in YzO3 and are close to the cubic phase although the bulk phase in both cases has a tetragonal structure. The segregated phases often have low conductivity, minimum activity to oxygen charge transfer reactions and therefore impede electrode reactions. Depending upon the amount of

82

Figure 9. Scanning electron micrographs showing exaggerated grain growth in a zirconiayttria ceramic induced by heating the sintered ceramic in contact with the glassy phase at 1400~ in air). (a) Surface (--,) directly in contact with the glassy phase; and (b) microstructure in the bulk of the ceramic. Note different magnification for (a) and (b). segregated phases present, the effect on the electrode kinetic behaviour can be substantial. Up to an order of magnitude higher electrode resistance has been observed as a result of contamination of the external surface during sintering [47]. Since the segregation of impurity phases to the external surface of the electrolyte continues for a long period of time, stability of the electrode/electrolyte interface and long term stability of the zirconia- based solid electrolyte devices (e.g. fuel cells, oxygen pumps) is questionable if electrolyte materials with high impurity contents are used. Another source of impurity segregation at the external surface of zirconia ceramics is the substrate which is placed in contact with green density or sintered ceramics during sintering and subsequent heat treatments. Any impurities present in the substrate material can segregate on the external surface of zirconia ceramics. This is a common problem in technology. Steele [48] has reported that during the sintering of zirconia wafers, SiO2 contamination occurred from alumina boards (96% Al~O3 with significant quantity of SiO2) placed on top of zirconia wafers to keep them flat. St.eele reported more than an order of magnitude higher electrode resistance at the electrode/electrolyte interface for contaminated electrolyte sheets. 3.4 Existence of interphase due to reaction/diffusion Chemical stability of electrode materials and electrode/electrolyte interfaces in the operating environment is of paramount importance in solid state electrochemical devices. A stable interface is desirable for long term stability of the device. However, if the interface corrodes as a result of oxidation or reduction of the electrode material or reactions between electrode and electrolyte materials or interdiffusion of species from one phase into the other, the consequences for the interfacial electrochemical reactions may be severe. The nature of the intermediate phase and its amount and distribution will then determine the electrode kinetic properties. Several examples have been reported in the literature where formation of intermediate layers between the electrode and electrolyte either during the cell

83 operation or as a result of cell fabrication have lead to serious degradation or poor cell performance. Some of these examples are discussed below. Many metals are used as electrode materials in zirconia-based solid electrolyte cells. The interface between a metal and the zirconia electrolyte is characterised by the presence or absence of an intermediate layer of a metal oxide which in turn depends on the thermodynamic stability of the electrode and the intermediate layer. As mentioned previously, Pd has high catalytic activity to oxygen charge transfer. However, formation of a thin intermediate layer of PdO at the interface between Pd and stabilised zirconia has a large influ-

PdO Pd

4.0 3.53.0-N

-r

o

O

2.52.0-

_J

1.51.0-

6o0

I

1

I

J

7oo 800 900 ~o0o TEMPERATURE,~

Figure 10. Low frequency (10 Hz) impedance of a Pd/YSZ/Pd cell as a function of temperature in 100% oxygen.

ence on the oxygen exchange kinetics and leads to almost blocking of the oxygen reduction reaction. The formation of the PdO layer is dictated purely by thermodynamic requirements and is dependent on temperature and oxygen partial pressure. For example above 870~ in pure oxygen Pd exists as a noble metal but below this temperature Pd oxidises to PdO. Figure 10 shows electrode impedance (monitored at a low frequency) versus temperature in pure oxygen. The temperature at which a sudden decrease in the electrode impedance occurs corresponds to the thermodynamic decomposition temperature of PdO in a given oxygen partial pressure [49,50]. Another example worthy of discussion is the stability of the interfaces between perovskite based electrode systems and the zirconia electrolyte during cell fabrication and subsequent device operation in constrained thermodynamic environments. Perovskites of the formula .ABO3+x with substitution at A-site or both A- and B-sites are commonly used as electrodes in oxygen sensors, oxygen separation membranes, steam electrolysis and solid oxide fuel cells [51-54]. Sr doped lanthanum manganite (LaMnO 3) materials are commonly used as electrodes on the air side because of their close thermal expansion match with the electro-

84 lyte, high electronic conductivity, mixed electronic/ionic conduction, ease of fabrication and the potential they offer to tailor their properties to specific needs. Although these cathode materials are reasonably stable in contact with zirconia-based electrolytes at the solid oxide fuel cell operating temperature of 1000~ on heat treatment of the (La,Sr)MnO3/doped zirconia interface above about 1250-1300~ an intermediate layer of a pyrochlore La2Zr207 and/or SrZrO 3 is formed [55-61] The conductivity of La2Zr207 or SrZrO 3 is several orders of magnitude lower than that'of the electrode. The intermediate phase not only acts as a barrier layer for oxygen transfer reaction, it also contributes to resistive voltage losses at the interface. Thus formation of an intermediate layer during cell fabrication or during cell operation over a period of time will lead to cell performance degradation. Yokokawa et al [58,59] have performed thermodynamic calculations on the phase stability of several perovskites in contact with zirconia electrolytes. The most probable reaction in the case of LaMnO 3 is: LaMnO 3 + (2x)ZrO 2 + (3x/2)O 2 = Lal.2xMnO3 + (x)La2Zr207. In this reaction a small amount of La20. reacts with ZrO 2 to form La2Zr207 and as the La deficiency at A-site increases the drivin~ force for the formation of the pyrochlore phase subsides. According to Yokokawa et al [58,59], LaMnO 3 sufficiently deficient (10-15% deficiency) at the A-site should be inert towards zirconia. However, as the La deficiency increases at the A-site, the Mn activity increases at the B-site. Dissolution of Mn into the zirconia matrix under these conditions has been discussed. However, solubility of Mn 3+ and Mn 4+ is relatively low in zirconia in air or pure oxygen and Milliken et al [60] have observed an extensive manganese diffusion along grain boundaries in zirconia once the interface is heated above 1200~ For SrO doped LaMnO 3, the A-site deficiency limit is affected by the La/Sr ratio and depending on the doping level, both SrZrO 3 and La2Zr.O 7 can exist in equilibrium with the 4 perovskite electrode [58,59]. The schematic of interfaclal reaction between perovskite electrodes and zirconia electrolytes is given in Figure 11. A number of other electrode materials with doping at the B-site have also been proposed as potential electrodes for the oxygen reduction reaction. Sr doped LaCoO 3 is a much better electrode in an electrochemical sense than Sr doped LaMnO 3 but it reacts even more vigorously with zirconia. Ivers-Tiff6e et al [61] have reported that for (La,Sr)(Mn,Co)O 3 electrodes, the amount and ratio of the secondary phases (SrZrO 3, La2Zr207 and cobalt oxide) formed between mixtures of electrode and electrolyte powder compacts (heat treated at 1300~ was a function of La:Sr and Mn:Co ratios. Co-free materials showed a small quantity of SrZrO 3 phase and the amount of this phase increased with the SrO content in the perovskite. For Co-containing compositions, both SrZrO. and La..Zr207 phases were 0 ,Z . observed in large quantities with the amount of the former increasing w~th increased Sr doping at the A-site. For (La,Sr)CoO 3, more than 50% of the electrode consisted of S r Z r O 3, La2Zr207 and cobalt oxide. Ivers-Tiff6e et al [61] have also reported that at the electrode/electrolyte interface, the overpotential losses increased dramatically with increasing heat treatment temperature of the interface for Co-containing electrode compositions whereas the effect of increasing heat treatment temperature (from 1150 to 1300~ on electrode overpotential losses was relatively small. Nevertheless, the electrochemical characteristics of the interface can change if reasonable quantities of one or more secondary phases are formed. The electrode reaction then will be dictated by the physical and electrochemical nature of the intermediate layer.

85

Trends for intermediate phase formation for AyBO 3 (A = La, B = Mn, T > 1200oC)

y > 0.9

La2Zr207 (I')

y

La2Zr207 ($) La2Zr207 (1')

<

0.9

Mn ~ in zirconia

Sr substitution at A-site

Sr:La ratio ( t )

SrZrO 3 ( f )

Co substitution at B-site

C o : M n ratio ( t )

La2Zr207 ( f ), SrZrO 3 ( t ), CoO ( t ).

Figure 11. A schematic of the LSM/zirconia interface and possible interface reactions for different conditions.

Another example of interface reactions is that of electrodes based on stabilised uranium oxide. These nonstoichiometric oxides of the general formula (U,M)O2~:x (M = Sc, Y, Dy, Pr) and having the fluorite-type structure are good electrodes for use m oxygen sensors and fuel cells [62-65]. However, some of these solid solutions when heated above about 1200~ in contact with zirconia-based electrolytes react to form intermediate phases such as (U, Zr)O2~_x and/or (U,M,Zr)O... These intermediate phases may have different catalytic activity~to oxygen exchange ~dnetics and once formed at the interface can alter the rate of oxygen transfer reaction. Badwal et al [66] have reported that the formation of an intermediate phase between (U,M)O z electrodes and zirconia-based electrolytes is a function of the type of dopant used to sta:~]lise uranium dioxide. In the case of (U, Sc)O2• x electrodes heated above 1200~ in contact with a zirconia-yttria electrolyte, the formation of an intermediate phase was clearly observed as shown in Figure 12. On the contrary (U,Y)Oz• x electrodes of similar composition showed much higher stability under similar heat treatment conditions [66] The (U, Sc)O_ fluorite solid solutions show very high volattlity for U in the form of UO 3 and the mecffanism for the intermediate phase formation appears to be through vapour phase transport rather than through the direct solid state reaction between electrode and electrolyte phases. In the case of (U, Sc)Oz_ solid solutions, the loss of uranium in (U, Sc)O~_ also leads to the precipitation ~J~ Sc203 and decomposition of the fluorite phase, and~Xer cause for the loss of catalytic activity of the electrode. ;l:X

.

"

2

X

86

Figure 12. (a) Scanning electron micrograph of the interface between zirconia-yttria (1) and (U, Sc)O2§ x (3) heated at 1400~ for 15 h in air showing the formation of a dense intermediate phase (2) at the interface. (b) X-ray line scan across the electrolyte/electrode interface showing diffuse region. 3.5 Interfaces between components of a composite electrode Composite electrodes are commonly used in solid state electrochemical devices based on zirconia for several different reasons. These may include enhancing the catalytic activity of the electrode at the interface, change in the thermal or mechanical properties of the material, control of sintering and grain growth kinetics. For example Badwal et al [67,68] have reported that composite electrodes consisting of a noble metal (such as Pt) and some semiconducting metal oxides show much superior electrode kinetic behaviour for the oxygen exchange reaction compared with either Pt or the metal oxide electrode alone. About an order of magnitude reduction in the values of the time constant and the resistance associated with the electrode process have been reported for mixtures of Pt with either (U, Sc)O2§ x or CrNbO 4 materials [67,68]. This is shown in Figure 13 for mixtures of Pt and CrNbO 4 electrodes. Such electrodes gave a very good performance in potentiometric sensors ana have enabled the operating temperature to be lowered to below 400~ As mentioned earlier, zirconia doped with yttria or a similar dopant is added to the Ni fuel electrode to reduce thermal expansion mismatch between the composite electrode and the electrolyte, to reduce sintering and grain growth of Ni particles and to control electrode morphology [15,27,28,69]. It is not entirely clear if the addition of zirconia to Ni directly enhances electrode reaction rates by increasing the electrode/electrolyte/gas contact area boundary. The method of Ni/zirconia composite electrode preparation and distribution and size of both Ni and zirconia particles are important to achieve high conductivity, lower overpotential losses as well for long term stability of the electrode performance [70-72]. In

87 CrNb04 +Pt Electrode 4.0

~ 600o C

0.5

CrNb04

+ Pt Electrode

~

600Oc

3.5 o

"

9

2.5

~

2.0

9-0.5

1.5 1.0

I 0

I 20

I 40

I 60

Wt% Pt02

1 80

I .. 1O0

-1.0

I 0

1 20

I 40

I 60

I 80

1 1O0

wt% Pt02

Figure 13. The effect of Pt addition on the electrode resistance and time constant of CrNbO 4 electrode. The data were recorded at 600~ PtO 2 decomposes to Pt on heat treatment at 600~

any case, the interface between Ni and zirconia plays a critical role and the existence of an oxide near the Ni particles appears to facilitate the hydrogen oxidation reaction. Contradictory mechanisms suggesting the dominant role of the electrode or the electrolyte have been proposed. It has also been reported that the presence of a small amount of water vapour (2-3%) in the hydrogen fuel considerably reduces overpotential losses at the Ni/zirconia interface [ 15]. The hypotheses proposed to explain this behaviour are numerous and unsatisfactory. Recently Jiang and Badwal [73] have suggested that the preferred mechanism for hydrogen oxidation is via its adsorption on the Ni surface which is facilitated in the presence of water vapour. However, in dry hydrogen, hydrogen adsorption takes place predominantly at the electrolyte surface and electron migration on the electrolyte surface is the rate limiting step. Other composite electrodes proposed for the fuel oxidation reaction are (i) a mixture of Ru and zirconia which is a better electrode than Ni/zirconia cermet because of the higher melting point of Ru compared with Ni (low sintering and grain growth), higher steam reforming activity and lower carbon deposition rates [74]; and (ii) zirconia-yttria-titania mixed conducting materials [75]. Mixed conducting electrodes prepared by the addition of a transition metal to zirconia-based electrolytes have attracted interest because of the possibility of an increase in the reaction interface between the electrode and the electrolyte and a consequent reduction in the polarisation losses. Also the stability of such an electrode/electrolyte interface would be good due to the expected low sintering and grain growth kinetics for the refractory oxide and closely matching thermal expansion coefficient (dewetting similar to that observed at the Ni/zirconia interface is less likely to occur) with other cell components [75-77].

88 4. INTERFACES WITHIN THE ELECTROLYTE GRAINS In this section interfaces within the electrolyte grains have been considered. In homogeneous single phase zirconia-based electrolyte materials, in the absence of inclusions or precipitates of another phase, the ion transport properties are determined by the level of dopant (vacancy concentration), type of dopant, interaction between vacancies and dopant cations and the temperature range over which measurements are made. In such materials ion transport takes place freely without any hindrance from internal interfaces. However, in practical electrolyte systems, the diffusing anions must negotiate a series of barriers as it transports through the zirconia-based electrolyte grains. Depending on the system under study, the types of barrier can be broadly categorised into four distinct groupings:

(a) (b) (c) (d)

Precipitation of a second phase within the grains. Compositional variations across grains. Second phase inclusions. Microdomains/ordering of cation or anion sublattice.

In each of the cases, the interfaces described produce a deviation from the idealised picture of an electrolyte consisting of a homogeneous single phase material with a high concentration of randomly distributed vacancies providing diffusion pathways. 4.1 Precipitation of a second phase within the grains In zirconia based systems it is the general case that the optimum electrolyte which combines both high ionic conductivity with high mechanical strength is not a single phase material but is found in the two phase field where it is possible to produce the desired combination of phases and microstructure. Moreover, the maximum conductivity (above 800~ in zirconia-based systems occurs for a dopant concentration close to the phase boundary in the two phase field. The maximum in ionic conductivity has been reported at about 8 tool% M203 in ZrO 2 - M 2 O. (M= Y, Yb, Sc) and 12 -13 mo1% CaO in ZrO 2 CaO (3,78,79]. In all such systems, the thermodynamic requirements of phase diagrams would require some growth of precipitates of another phase. Additionally, in order to produce a homogeneous solute distribution, high temperatures well above the operating temperature of the cell are often required for processing. However, the subsequent cooling invariably results in the samples passing into two phase regions where even though the kinetics are such that cation redistribution is insignificant, very often over very localised regions of a few atomic units evidence of ordering or minor inhomogeneities can be seen. Further to this is the operating temperature of the solid electrolyte device. Although the kinetics of phase redistribution are slow, the temperature regime in which the system finds itself is usually well inside a two phase region and invariably phase separation begins albeit on an extremely fine scale due to thermodynamic constraints. Figure 14 schematically shows this situation for an idealised system in which the optimum aliovalent cation content for which maximum ionic conduction occurs is close to the phase boundary between the single and two phases regimes (for detailed phase diagrams see Yoshimura [80]). The developing precipitates in a homogeneous matrix of electrolyte grains will definitely contain very little stabilising dopant. However, if some of the dopant remains in the developing precipitate, this will lead to some conductivity which is expected to be much lower than that in the surrounding matrix An example of phase redistribution in a ZrO 2 - Y..O3 system sintered in the single phase field and subsequently annealed in the two phase fielt~ is

89

3000

o

0

(1) I--

2000 c

Sintering

L._

(9 O.

t~rrTp~r~tdr~

E

1000

_ _ _Fuel cell op_er_ati_on

Technologically important

composition range

4

I

I

I

2

4

6

I 8

_

I 10

1 12

I 14

I 16

mol % stabiliser

Figure 14. A schematic of parts of the phase diagram showing the regions of technological interest. shown in Figure 15. The so called 'tweed structure' has developed in a specimen of 8 tool% YzO3 stabilised zirconia after prolonged heat treatment at 1000~ This structure is caused by the development of coherent precipitates of the tetragonal variant of zirconia within the cubic grains. There is a resultant decrease in the conductivity of this material as a consequence of this so-called "ageing" because of the precipitation of the poorly conducting tetragonal zirconia. The conductivity degradation as a function of time at 1000~ is shown in Figure 16 for 8 tool% YzO. - ZrO 2. A further example was reported by Steele et a~ [81] for the system CaO-ZrO z. In this case the developing microstructure as a function of aging time at specific temperatures (i.e. solid solution - > tweed structure - > tetragonal precipitates - > monoclinic precipitates) could be clearly followed, even in this very complicated system, by conductivity measurements. The varying departure from the idealised case of randomly distributed vacancies for this system leads to a continually deteriorating conductivity as the nature and intensity of interfaces between various phases within the grains of the electrolyte changes. The situation is even further complicated in some zirconia-based systems by the thermodynamic requirements of the phase rule and the slow kinetics of phase redistribution. The necessary redistribution of dopant cations dictated by the development of a two phase region in combination with low diffusion rates of the cation species can lead to inhomogeneous accumulations of cations at precipitate interfaces. The consequence of this will be a reduction in the concentration of vacancies available for conduction. An extreme example of this is in the case of Mg-PSZ (magnesia partially stabilised zirconia) in which the development of coherent tetragonal precipitates gives rise to solute rich regions around these precipitates which have been shown to accumulate into ordered regions of ~5-phase [82]. Figure 17 shows the dark field electron micrograph in which the 6-phase (an ordered phase

90

Figure 15. The "tweed structure" observed in a yttria-stabilised zirconia after prolonged ageing. Electron diffraction patterns recorded from this sample show tetragonal symmetry. 0-18 i

8 mol% Y203 - ZrO 2

1000~

0.17 k

I E

,-0 IC

0.16 0.15

0.14 0.15 0.12 L 0

.

. . . 1000

-'~176176 2000

I

3000

|

.,

I

4000

,

5000

Time, min Figure 16. Conductivity degradation in 8 mol% Y203 - Z r O 2 at 1000~ as a function of time resulting from the composition being in the two phase field. Diagram courtesy of ET. Ciacchi.

91

Figure 17. Dark field electron microscope image using the ~-phase reflections observed in electron diffraction patterns of this area. The light regions show the development of ~phase domains at the interface of the growing tetragonal precipitates. (Micrograph courtesy of R.H.J. Hannink). with a higher solute content) precipitates are highlighted. The evidence for this occurring in Mg-PSZ comes from the analysis of diffuse scatter in electron diffraction experiments in combination with analytical electron microscopy [83,84). Diffuse scatter in similarly aged specimens of other stabilised zirconias is often observed, i.e. Sc~O3-ZrO 2 [85] and Y O ZrO 2 [86] and although these systems have not been fully examined as in the case o~2~g PSZ, it is not an outlandish assumption made that a similar phenomenon is occurring. Some ordering is taking place, and with increased formation of ordered phases with a higher dopant content more vacancies become unavailable for conduction. In the Y203 -.ZrO 2 [86-88] and Sc203 - ZrO 2 [85] systems, which have the most important properties m connection with their use in the solid oxide fuel cell, it has been shown that rapidly quenching specimens which have been fired within the cubic phase field produces a metastable phase known now as the t'-phase. This phase contains a high concentration of solute and the rapid quench prevents the partitioning of the material into the equi-

92 librium two phase condition. The t'- phase is characterised by the occurrence of large twins running through the grains of the material. The conductivity of this phase is high but it rapidly deteriorates with annealing at moderate temperatures [89]. The t'-phase is seen to decompose into a cubic phase with precipitation of the equilibrium tetragonal phase which is usually observed. The tetragonal phase is very low in solute and hence will have a low concentration of vacancies available for conduction. This is illustrated in Figure 18 which shows the changing microstructure of a 7.75 mol% SCzO3 - ZrO z as a function of annealing and the corresponding deterioration in the conductivity. A further example of phase changes which can effect the conducting properties of these materials occurs when the microstructure of the ceramic itself prevents the system attaining its equilibrium state. In such cases the system is metastable and external factors such as mechanical stress can trigger a transformation to the lower energy state with detrimental effects on the conductivity. An example of this is the doped - TZP (Tetragonal Zirconia Polycrystals) materials which have a very well defined microstructure. For the material to remain tetragonal and not transform to the monoclinic state, the grain size must not exceed a critical size (about 0.5~m). The grain size nevertheless is dependent on the solute content, type and amount of impurities present, solute homogeneity and sintering conditions. If some damage is externally introduced then the transformation to the monoclinic phase can spontaneously occur with an accompanying volume expansion. The monoclinic phase is intrinsically a poor conductor and in addition an accompanying volume expansion gives rise to the formation of insulating microcracks. The sequence of events can be seen in Figure 19 in which the tetragonal grain is seen to have transformed in this case as a consequence of beam heating in the electron microscope (Figure 19). ZrO z doped with 2-3 mol% Y20. has tetragonal structure, and is one of the most widely studied potenttal electrolyte material for use in oxygen sensors and fuel cells [90,91]. For Y-TZP flexural bending strength in excess of 1 GPa has been reported [2]. Although these materials have lower conductivity than the 8 mol% YzO3 - ZrO.. composition at 1000~ at lower temperatures (below about 400 - 500~ the conductivi~ty is comparable or better [78]. However, precipitation of poorly conducting phases has also been reported in 3 tool% YzO3 - ZrO 2 ceramics [92]. This composition is in the two phase field over a wide temperature range [80] and on annealing sintered and rapidly cooled ceramics in the 8001200~ temperature range, solute redistribution leading to the presence of several variants of the tetragonal phase within the same grain and precipitation of monoclinic zirconia was observed [90]. The development of this type of microstructure leads to a substantial reduction in the intragrain conductivity as shown in Figure 20 [92]. Doped tetragonal zirconia (TZP) materials containing 2-3 tool% Y203 undergo a phase transformation to the less conducting monoclinic phase when annealed in the presence of moisture in the low temperature range of 100 - 400~ The transformation starts from the external surface and moves inwards and it is somewhat dependent on the microstructure of the ceramic. The phase transformation has serious consequences for the mechanical integrity of the ceramic [2]. Also Badwal and Nardella [93] have reported the formation of mZrO 2 on the anodic side (only) of the tetragonal zirconia electrolyte in complete cells on current passage (Figure 21). The amount of m-ZrO 2 formed is a function of the current density, time of current passage, the temperature of cell operation and the ceramic microstructure. No m-ZrO 2 was detected at the anodic side of the electrolyte at temperatures above 500~ This behaviour can have a significant effect on the electrode performance and was attributed [93] to the existence of space charge layers near the interface leading to substantial reduction in the oxygen vacancy concentration which otherwise is responsible for the stability of the tetragonal phase. 9

93

0.32 1000~ 0.30

", ~

TE ,..0 It2

%

0.28

-... %

0.26

"Oo

0.24

O0

OO

OOoOoQ 000Oo0 O

0.22

00 OIO@ @ OOOoo00oO e~176 ~176176

0.20

0

,

' 1000

I

2000

,

I

3000

J

I

4000

coo, t

5000

Time, min Figure 18. Bright field transmission electron micrographs of the sintered 7.75 tool% Sc20~ - ZrO 2 ceramic (prepared by coprecipitation) before (a) and (b) after annealing at 1000~ for 2000h in air. Also shown is the conductivity degradation for this material (freshly sintered) as a function of time change [3].

4.2 Compositional variations. Compositional variations can occur in zirconia based systems as a result of a number of situations which may include: processing difficulties which produce an inhomogeneous dis-

94

Figure 19. Bright field electron micrographs showing (arrowed) the partial transformation of a Y-TZP grain. tribution of cations; the tendency for produced metastable systems to disproportionate over a period of time at elevated temperature; and by the presence of contaminants which react preferentially with some component of the ceramic causing localised concentration gradients and creation of pseudo interfaces. These are not real interfaces in the sense that there is no precipitation of another phase. However, the system is not homogeneous with respect to solute distribution and regions of different composition have different electrical conductivities and overall ion transport through the grain involves oxygen ion exchange between these regions.

95 33000

3 tool% Y203-Zr02

350~

E 22000 0 AA

N I

11000

AAAAAAZ$&A

AA

AA~ OAOO O O O O O O OOo .AO Or,

0

0

11 bOO

22000

AAA&A

33bOO

44000

as000

Z' 9 f~cm

Figure 20. Impedance spectra of a 3 tool% Y.O 3 - ZrO 2 ceramic: (O) as-sintered (1500~ 4h in air) and quenched in air; (zx) after annealing at 1200~ for 50 h in air. The impedance data were recorded at 350~

60 [

I=1.4mA cm 2

9

50 4o

~

20 lo

0

25

50

75

100

125

Time, min

Figure 21. The effect of current passage in a electrode/electrolyte/electrode cell on mZrO 2 formation on the anodic side of a 3 mol% Y203 - ZrO z electrolyte.

96 Improvements in processing techniques associated with the improved powder technology have gone a long way to ensure that problems of inhomogenous mixing are rare. In zirconia based systems this was most important since inhomogeneities in the cation distribution are difficult to remove. It is a feature of the fluorite system that the cation sublattice is most stable in comparison with the mobile anion sublattice. Extreme examples of this problem have been reported in the Sc203 - ZrO 2 system [94-96] where the kinetics of cation redistribution are known to be very slow probably due to similarity in the size of Zr 4+ and Sc 3§ cations. The resulting microstructure is not optimised for maximum conductivity with localised regions of high resistivity causing blocking of the diffusing anions. Badwal and Drennan [97] in studying the Y20 - Sc203 - ZrO 2 system reported wide compositional variations (Sc/Y ratios) from grain t~ grain and also within the same grain. An example is shown in Figure 22. This type of inhomogeneous distribution produces additional interfaces for ionic transfer and materials with less than optimum ionic conductivity.

Figure 22. Scanning electron micrograph of 50 wt% (4.7 mo1% Sc203 - ZrO 2) + 50 wt% A1203 and electron probe micro analysis results showing inhomogeneous distribution of Sc within a zirconia grain. Sc/Zr = 7.0(1); 3.0(2); 3.0(3); 6.0(4); 5.5(5); 6.0(6); and

8.0(7). In a number of systems which are of technological importance in solid electrolyte devices the optimum compositions for maximum ionic conductivity, as reported earlier, are usually close to the phase boundary between single and two phase regions. Therefore, localised variation in composition may be brought about by the system having to be operated in

97 temperature regimes in which the equilibrium situation is a two phase mix. An excellent example of this is described by Chaim et al [98] in which disproportionation can occur and even if this is over atomically small regions the process removes available vacancies from the conducting process. A further possibility first outlined by Heyne [99] and later discussed by Steele and Buffer [100] is the formation of space charge regions developing at the surface or the grain boundary of the ceramic. These form as a result of the large difference in mobility between the cation and anion species and the resulting difficulty the system has in rapidly attaining the equilibrium situation. Local variations in the concentration of vacancies will effectively produce barriers to the diffusing anion. Contamination of the starting powders by glassy impurities can produce situations where the microstructure can be dramatically altered. In general, impurities are found to accumulate at the grain boundary of the ceramic and will in general react preferentially with one of the components of the system. Examples of this can be seen in the yttria-zirconia [101105] system where it has been shown that accumulations of secondary grain boundary phase rich in S i t . aid in the redistribution of the solute in the system giving rise to inhomogeneous grain g~rowth and large compositional gradients. An example is shown in Figure 23. These large variations in microstructure once again interfere with the continuity of conducting pathways, producing additional interfaces which result in inferior materials for use in electrochemical devices.

4.3 Second phase inclusions. Generally speaking, any restriction of the pathway of the diffusing anions through the ionic conductor will result in the degradation of the conductivity. Second phase inclusions of a non-conducting species will have two detrimental effects. Firstly, the second phase effectively dilutes the volume of conducting matrix and secondly, the insulating phase constricts the direct pathway for the diffusing ions. However, subtleties can arise. It has been reported that the addition of alumina to yttria stabilised zirconia can have a beneficial effect on the total conductivity of the ceramic at low temperatures [106-108]. The proposed mechanism involves the gettering of the contaminants such as silica so that the addition effectively cleans the grain boundary regions opening up the pathways for conduction (see section 5.4). Nevertheless the effect of insulating phase addition is normally detrimental to the intragrain conductivity. It should be noted that the dilution of the conducting species has still occurred and more restricted pathways have been introduced by the addition. As a result of second phase addition, the net intragrain conductivity (ignoring grain boundary resistivity effects), usually is lower than can be calculated taking into account the volume fraction of the insulating phase due to isolation of some grains of the conducting phase from the conducting pathways in addition to constriction of current lines. This is shown in Figure 24 [109]. One of the other features of second phase addition such as alumina is its effect on the grain size. The addition of alumina to ZrO 2 - Y203 leads to a significant reduction in the grain size with a consequent increase in the interfaces between grains (grain boundary surface area. As shown in Figure 25, the effect is clearly noticeable in the case of fully stabilised ZrO 2 - Y203 [110]. In general, if defect sizes are minimised in combination with a reduction in the grain size, an increase in the mechanical strength is observed. 4.4 Microdomain formation and ordering A subtle effect is observed in stabilised zirconia systems which manifests itself in the appearance of diffuse scatter in diffraction experiments. Although these effects have been

98

Figure 23. Bright field micrographs showing the inhomogeneous grain growth in a specimen of Y-TZE The yttria content of the large grain was observed to be close to 9 mol%, approximately three times the yttria content of the surrounding grains.

observed and studied for over thirty years, there still exists doubt as to the exact cause of the effect. All workers agree that some ordering process is taking place and with the ordering comes the inevitable reduction in ionic conductivity. For the system CaO-ZrO 2 (within the composition region Ca Zr. O. where x lies in the range 0.1-0.2) Allpress and Rossell [111] provided very ~lea~Xev~lence that the observed diffuse scattering was as a result of the formation of microdomains of the order of 3nm in diameter within the grains of the material. The microdomains were shown to have clear structural integrity corresponding to the ordering observed in the compound which occurs in the CaO-ZrO 2 system, CaZraO 9. In the case of Y203-ZrO. the situation is less clear. A microdomain model has been suggested [ 112,113] but the difficulty with this model is that diffuse scatter can be observed over a wide range of stoichiometry but no compound exists to which the microdomain model can be applied. More recently alternative suggestions [114, 115] invoke the concept of the modulated structure where intermediate compositions between the end members of the solid solution range show a regular commensurable modulation in very

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specific crystallographic orientations. Irrespective of the mechanism, an ordering process is clearly taking place which is limiting the availability of anions for the conduction processes.

Figure 25. The effect of alumina addition on the grain size distribution in 8 mol% Y203 ZrO z ceramic sintered at 1500~ for 4h. (a) No alumina addition and (b) with 10 wt% alumina addition.

100 5. INTERFACES BETWEEN GRAINS (GRAIN BOUNDARIES) In most applications polycrystaUine materials are invariably used. The polycrystalline nature of ceramic materials means that the interfacial region between the grains have a profound influence on the physical properties of the material. In terms of the electrical properties, the boundary region provides a barrier to the transport of oxygen ions in zirconia - based solid electrolytes. The grain boundary interfaces can be divided into a number of different categories similar to the interfaces within the grains discussed in the previous section. Four major types of grain boundaries encountered are" (a) Phase free (or clean) boundaries. (b) Intermediate phase formation at grain boundaries from the matrix components. (c) Intermediate phases of the impurity type. (d) Inclusions. 5.1 Phase free boundaries In terms of the electrical properties, the boundaries which are free from any secondary phases or segregation of impurities should be the ones with minimum hindrance to the ionic conducting species. However, this would require complete matching of the atomic interfaces which in practice is not observed. Phenomena such as lattice mismatching, dislocations and residual stresses observed commonly in metals are also observed in ceramic systems although their analysis and characterisation is still in the early stages of study due to difficulties in preparing representative samples and the added crystallographic complexities. However, imperfections, lattice mismatch, and dislocations are all observed in ceramic systems and no doubt will receive more attention as the structures are better understood. Nevertheless, these are some of the interfaces present in clean materials which will influence the ionic transport properties across grain boundaries. Moreover, the presence of pores at grain boundaries or imperfect contacts between grains will cause constriction of current pathways leading to a decrease in the overall conductivity. Also enhanced segregation of bulk phase matrix components near the grain boundary region can alter the nature of space charge layers. In zirconia-based systems, it is extremely rare to come across grain boundaries free of secondary phases and/or segregated impurities. Apart from single grains or crystals with low angle grain boundaries, there is no demonstrated case of socalled clean grain boundaries. 5.2 Intermediate Phase Formation from the Matrix In the case of complex zirconia - based systems which are maintained at temperatures in which phase changes are taking place, the grain boundary region can provide the nucleation site for the development of secondary phases. The classic example of this is the MgO-PSZ system which when maintained below 1400~ develops a decomposition product around the perimeter of each grain. The decomposition reaction is simply the MgO-PSZ decomposing into zirconia (monoclinic form ) and free MgO. This reaction nucleates at the boundaries and develops towards the centre of the grains as can be clearly seen in Figure 26 which shows the grains outlined by the decomposition region. Obviously the effect on the conductivity of the formation of this insulating layer is dramatic. With increase in the amount of m-ZrO 2 at grain boundaries, a substantial increase in the grain boundary resistance is observed as shown in Figure 27 (78). Attempts to reduce the amount of decomposition with time have been successful by controlling the chemistry of the grain boundary phase. The addition of SrO to the system has

101

Figure 26. Optical micrograph of the decomposition region around the grains of over aged Mg-PSZ. The white regions consist of monoclinic zirconia and free MgO.

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102 been shown by Drennan and Hannink [43] to favourably alter the grain boundary decomposition with subsequent improvements to changes to the mechanical properties of the ceramic. Although MgO stabilised zirconia are not suitable materials for most applications requiting high ionic conduction apart from their possible use as immersion sensors for molten metals, measurements of conductivity have been made [116] and the influence of SrO additions examined. These appear-to support the idea that the addition effectively reduces grain boundary decomposition. Although this is an extreme case of intermediate phase formation, it serves as an example of the situation which may occur when nucleation points such as grain boundaries and pores provide the regions where solid state reactions can rearrange the local chemistry of these metastable systems.

5.3 Intermediate phases of the impurity type Important advances have been made in the preparation of zirconia powders. Contamination levels are controlled down to the 10's of ppm and additives deliberately added to aid in sintering are controlled precisely. In general most contaminants or sintering additives accumulate in the grain boundary region. The control of the location and chemistry of this grain boundary phase becomes paramount in optimising the ionic conduction of the ceramic. Two competing phenomenon take place. Homogeneous, high density, flaw free ceramics are necessary properties of any electrolyte since with present device designs the electrolyte must maintain integrity whilst being thin enough to have minimum resistance. A means of obtaining these properties is to use sintering aids which in tum produce liquid phases to facilitate homogeneous grain growth and assist in the elimination of porosity. The difficulty arises when the liquid phase formed remains along the grain boundary providing a most effective insulating barrier. In such cases the beneficial effects of the presence of ,the liquid phase on the densification processes are negated by the insulating grain boundary phase. Examples of the type of liquid phase location are shown in Figure 28. In Figure 28 (a), the grain boundary phase can be seen outlining the grains of yttria stabilised zirconia, totally wetting any interface region. Altematively, in (b) the triple point regions show accumulations whilst in (c) the total dewetting of the boundary results in isolated accumulations along the boundary. The type of grain boundary phase formed is a function of the sintering temperature, gas atmosphere and chemical nature of impurities [39]. The electrical behaviour of these different grain boundary microstructures, as expected, is quite different. The grain boundary phases which wet the grain surfaces well have the most detrimental effect on the conductivity across grain boundaries. Figure 29 shows examples of different contributions to the grain boundary resistivity for different grain boundary microstructures. Attempts have been made to influence the ionic conductivity of the grain boundary phase so that minimum disruption to the conduction is obtained. For example Verkerk et al [117] added Bi203 to yttria stabilised zirconia to aid in densification but also to influence the conducting properties of the boundary since bismuth oxides are known to have high ionic conductivity at elevated temperature. The addition succeeded in assisting densification, however, the conductivity was seen to decrease as a result of unwanted interfacial reactions. So far attempts to enhance grain boundary conduction by adding Bi203 have failed. The grain boundary phases are quite mobile and their location can be altered by postsinter heat treatments. Some examples are discussed below. Altering the cooling rate of the fired specimen of yttria stabilised zirconia has been found by Badwal and Drennan [ 118, 119] to affect the grain boundary resistivity. They suggested that the effects observed were a result of the changing nature of the wetting behaviour of

103

Figure 28. A series of micrographs showing the various distributions of grain boundary phase in yttria zirconia. (a) Distribution evenly along the grains; (b) in the triple point regions; (c) in pockets along the boundary. the grain boundary phase. At elevated temperatures the liquid phase was observed to dewet the ceramic boundary and rapid cooling effectively froze-in this structure. Alternatively slow cooling allowed the grain boundary phase to rewet the boundary region consequently reducing the total conductivity. External pressure [ 120] on the sintered ceramic at elevated temperatures (in the vicinity of 1200~ can also force the grain boundary phase to move away from the conduction paths in the direction of pressing and move into parallel paths thus occupying regions it would not under normal conditions. In effect the application of anisotropic pressure to force the liquid phase into regions parallel to the applied force has lead to increase in the ionic conductivity along the pressure direction. The grain boundary resistivity was also seen to decrease as a function of increased applied stress. Similarly heat treatments of the sintered ceramic lead to migration of the grain boundary phase to the external surface thus decreasing their impact at grain boundaries [47]. This shows that the grain boundary microstructure of ionic conducting materials can be engineered and the nature of interfaces changed to optimise conditions. 5.4 Inclusions Since the present generation of commercially supplied powder materials are of a high quality we will ignore the cases of accidental contamination by foreign items. It is sufficient to say that any non conducting inclusion present along grain boundaries will in general impede ionic transport. However, the exception has been found to occur with the addition of A1203 to yttria stabilised zirconia. The effect here is subtle. The addition of the nonconducting species does reduce the ionic conductivity through the grains as would be

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Z' f~cm Figure 29. Impedance spectra showing different contribution of the grain boundary resistivity (arc on the right hand side) for different grain boundary microstructure. (a) Relatively clean grain boundaries; (b) grain boundary microstructure with grain boundary phase present at a number of interfaces between grains; and (c) extremely wetting grain boundary phase microstructure. The data were recorded at 350~ expected but for certain conditions, the total conductivity of the system is increased because the alumina changes the grain boundary conductivity. Figure 30 shows impedance spectra of two samples of Y-TZP with differing amounts of added alumina. A clear reduction in the total conductivity of the system is seen to occur for the specimen containing 10 wt% AlzO3. The explanation for this is that the wetting of the grain boundary phase at the AlzO3/yttria-zirconia interface is favoured over the yttria-zirconia/yttria-zirconia interfaces (Figure 31). Evidence for this comes from detailed microstructural studies combined with electrical measurements [106,121]. The actual amount of A1203 will depend on the distribution of the alumina and the relative size of the Y-TZP and the alumina particles. As the concentration of A1203 increases above a critical amount then the blocking nature of the alumina becomes dominant and the pathways for conduction are greatly reduced.

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Z', ~, cm Figure 30. Impedance diagrams (recorded at 400~ for a Y-TZP ceramic (a) without and (b) with the addition of 10 wt% alumina during powder processing. A substantial decrease in the grain boundary resistivity and a small increase in the intragrain resistivity has occurred. Some recent work [122] has shown that inclusions, pores and even microcracking can be investigated using complex impedance measurements and information on how these particular phenomena interfere with the ionic conductivity is becoming available. Previously we have cited the example of how the use of external pressure during sintering has been used to influence the location of the grain boundary phase. A further interesting case is where deliberately added second phase inclusions provide another means of influencing the electrical properties of the material when external pressure is applied during sintering. This has been shown by Drennan et al [123] for Y-TZP/A1203 composites where the application of external pressure by hot pressing gives rise to a microstructure in which the A1203 particles appear to accumulate in rafts perpendicular to the pressing direction. As a consequence of the formation of this anisotropic microstructure impedance measurements show a 25-30% increase in grain boundary resistance measured parallel to the pressing direction and a 9% increase in the bulk conductivity. 6. CONCLUSIONS We have described a number of the interfacial phenomena which occur in systems which incorporate zirconia based electrolytes and have attempted to categorise them according to the specific region of activity in solid state electrochemical cells. The understanding of the

106

Figure 31. Micrographs showing the interfacial region between Y-TZP grains and alumina particles where accumulations of a secondary grain boundary phase has occurred. processes that occur at various interfaces and their effect on the electrical and electrochemical properties of the materials becomes an important part of the development of successful devices especially those that must continue to perform over long periods of time. By obtaining an understanding of the processes it is possible to engineer the various micro structures involved so that optimum properties can be achieved. The study of interfacial properties will continue to be a most important field as we move into the age of the fuel cell and associated electrochemical de~,ices.

7. ACKNOWLEDGEMENTS Authors wish to thank Nathasha Rockelman and Kylie Crane for assistance with preparation of micrographs and Dr M.J. Bannister for reviewing this manuscript.

8. REFERENCES S.C. Singhal and H. Iwahara (eds), Proceedings of Third International Symposium on Solid Oxide Fuel Cells, The Electrochemical Soc. Inc., Pennington, NJ, 1993. S.P.S. Badwal, M.J. Bannister and R.H.J. Hannink (eds), Science and Techno-

107

10 11 12 13

14 15 16 17

18 19 20 21 22 23 24 25

logy of Zirconia V, Technomic Publishing Co. Inc., Lancaster, USA, 1993. S.P.S. Badwal, in Materials Science and Technology, A Comprehensive Treatment, eds., R.W. Cahn, P. Haasen and E.J. Kramer, Vol 11, "Structure and Properties of Materials, M.V. Swain, volume editor, (VCH, Weinheim, 1994) p 567. R.W. Spillman, R.M. Spotnitz and J.T. Lundquist Jr., Chemtech, 14 (1984) 176. M.J. Murray and S.P.S. Badwal, Mater. Sci. Forum, 34-36 (1988) 213. J.E. Bauerle, J. Phys. Chem. Solids, 30 (1969) 2657. J. R. Macdonald, ed. in Impedance Spectroscopy (John Wiley & Sons, New York, 1987). E.J.L. Schouler, in Proceedings of Solid State Ionic Conductors III, eds., LB. Goodenough, J. Jensen and A. Potier (Odense University Press, 1985) p. 16, M. Kleitz, H. Bernard, E. Fernandez and E. Schouler, in Science and Technology of Zirconia, Advances in Ceramics, Vol. 3, eds, A.H. Heuer and L.W. Hobbs (The Amer. Ceram. Soc., Columbus, Ohio, 1981) p. 310. S.P.S. Badwal, in Proceedings of International Seminar on Solid State Ionics Devices, eds., B.V.R. Chowdari and S. Radhakrishna (Woeld Science Publishing, Singapore, 1988) p. 165. C.S. Vayenas, Solid State Ionics, 28-30 (1988) 1521. T. Setoguchi, K. Okamoto, K. Eguchi and H. Arai, J. Electrochem. Soc., 139 (1992) 2875. I.V. Yentekakis, S.G. Neophytides, A.C. Kaloyiannis and C.G. Vayenes, in Proceedings of Third International Symposium on Solid Oxide Fuel Cells, eds., S.C. Singhal and H. Iwahara (The Electrochemical. Soc., Inc., Pennington, NJ, 1993) p. 904. B.C. Nguyen, T.A. Lin and D.M. Mason, J. Electrochem. Soc., 133 (1986) 1807. N.Q. Minh, J. Amer. Ceram. Soc., 76 (1993) 563. Y. Takeda, R. Kanno, M. Noda, Y. Tomida and O. Yamamoto, J. Electrochem. Soc., 134 (1987) 2656. J. Mizusaki, H. Tagawa, T. Saito, K. Kamitani, T. Yamamura, K. Hirano, S. Ehara, T. Takagi, T. Hikita, M. Ippommatsu, S. Nakagawa and K. Hashimoto, in Proceedings of Third International Symposium on Solid Oxide Fuel Cells, eds., S.C. Singhal and H. Iwahara (The Electrochemical. Soc., Inc., Pennington, NJ, 1993) p. 533. S.P.S. Badwal and F.T. Ciacchi, Solid State Ionics, 18/19 (1986) 1054. W.C. Maskell, N.M. Sammes and B.C.H. Steele, J. Phys. D: Appl. Phys., 20 (1987) 99. S.P.S. Badwal and H.J. de Bruin, J. Electrochem. Soc., 129 (1982) 1921. P. Fabry and M. Kleitz, J. Electroanal. Chem., 57 (1974) 165. J. Mizusaki and H. Tagawa, in High Temperature Electrode Materials and Characterization, eds., D.D. Macdonald and A.C. Khandkar (The Electrochemical. Soc., Inc., Pennington, NJ, 1991) p. 75. J. Mizusaki, H. Tagawa, K. Tsuneyoshi and A. Sawata, J. Electrochem. Soc., 138 (1991) 1867. B.C.H. Steele, Solid State Ionics Symposium Proceedings, E-MRS Conference on Advanced materials (ICAM 1991), Strasbourg, France, May 27-31, 1991, Mater. Sci. & Eng. B, 13 (1992) 79. S. Carter, A. Selcuk, R.J. Charter, J. Kajda, J.A. Kilner and B.C.H. Steele, Solid State Ionics, 53-56 (1992) 597.

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49 50 51 52 53 54 55

B.C.H. Steele, S. Carter, J. Kajda, I. Kontoulis and J.A. Kilner, in Proc. Second International Conference on Solid Oxide Fuel Cells, eds. F. Grosz, P. Zegers, S.C. Singhal and O. Yamamoto (Commission of the European Communities, Brussels, 1991) p. 517. A. Khandker and S. Elangovan, Denki Kagaku Kogyo 58 (1990) 551. S. Majumdar, T. Claar and B. Flandermeyer, J. Amer. Ceram. Soc., 69 (1986) 628. J.E. Anderson and Y.B. Graves, J. Electrochem. Soc., 128 (1981) 294. S.P.S. Badwal, M.J. Bannister and W.G. Garrett, J. Phys. E: Sci. Instrum., 20 (1987) 531. W.J. Fleming, J. Electrochem. Soc., 124 (1977) 21. D.M. Haaland, J. Electrochem. Soc., 127 (1980) 796. B.V.N. Rao and T.P. Schreiber, J. Amer. Ceram. Soc., 65 (1982) C44. R.G. Simhan, J. Non-Cryst Solids, 54 (1983) 335. G.S.A.M. Theunissen, A.J.A. Winnubst and A.J. Burggraaf, J. Mater. Sci. Lett., 8 (1989) 55. A.E. Hughes and S.P.S. Badwal, Solid State Ionics, 46 (1991) 265. A.E. Hughes and S.P.S. Badwal, Materials Forum, 15 (1991) 261. A.E. Hughes and S.P.S. Badwal, Solid State Ionics, 40/41 (1990) 312. S.P.S. Badwal and A.E. Hughes, J. European Ceram. Soc., 10 (1992) 115. F.T. Ciacchi, K.M. Crane and S.P.S. Badwal, Solid State Ionics, submitted, 1994. R. Chaim, D.G. Brandon and A.H. Heuer, Acta Metall., 34 (1986) 1933. P.J. Whalen, F. Reidinger, S.T. Correale and J. Marti, J. Mater. Sci., 22 (1987) 4465. J. Drennan and R.H.J. Hannink, J. Amer. Ceram. Soc., 69 (1986) 541. S.P.S. Badwal and S. Rajendran, Solid State Ionics, in print (1994). T. Tsukuma, Y. Kubota and T. Tsukidate, in Advances in Ceramics, Vol 12, Science and Technology of Zirconia II, eds., N. Claussen, M. RUhle and A. H. Heuer, The Amer. Ceram. Soc., Columbus, Ohio, 1984) p. 382. M. Matsui, T. Soma and I. Oda, in Advances in Ceramics, Vol 12, Science and Technology of Zirconia II, eds., N. Claussen, M. Rtihle and A. H. Heuer, The Amer. Ceram. Soc., Columbus, Ohio, 1984)p. 371. S.P.S. Badwal and A.E. Hughes, in Proc. Second International Conference on Solid Oxide Fuel Cells, eds., F. Grosz, P. Zegers, S.C. Singhal and O. Yamamoto (Commission of the European Communities, Brussels, 1991) p. 445. B.C.H. Steele, in Science and Technology of Zirconia V, eds., S.P.S. Badwal, M.J. Bannister and R.H.J. Hannink (Technomic Publishing Co., Inc., Lancaster, USA, 1993)p. 713. H.J. de Bruin and S.P.S. Badwal, J. Solid State Chemistry, 34 (1980) 133. S.P.S. Badwal and H.J. de Bruin, Aust. Chem. Eng., 20(6) (1979) 9. Y. Ohno, S. Nagata and H. Sato, Solid State Ionics, 9/10 (1983) 1001. Y. Teraoka, H. Zhang, K. Okamoto and N. Yamazoe, Mater. Res. Bull., 23 (1988) 51. O. Yamamoto, Y. Takeda, R. Kanno and M. Noda, Solid State Ionics, 22 (1987) 241. H.U. Anderson, Solid State Ionics, 52 (1992) 33. O. Yamamoto, G.Q. Shen, Y. Takeda, N. Imanishi and Y. Sakaki, in High Temperature Electrode Materials and Characterization, eds., D.D. Macdonald and

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A.C. Khandkar (The Electrochemical. Soc., Inc., Pennington, NJ, 1991) p. 158 S. Elangovan, A. Khandkar, M. Liu and M. Timper, in High Temperature Electrode Materials and Characterization, eds., D.D. Macdonald and A.C. Khandkar (The Electrochemical. Soc., Inc., Pennington, NJ, 1991) p. 191. A. Khandkar, S. Elangovan and S. Liu, Solid State Ionics, 52 (1992) 57 H. Yokokawa, N. Sakai, T. Kawada and M. Dokiya, Solid State Ionics, 52 (1992) H. Yokokawa, N. Sakai, T. Kawada and M. Dokiya, in Science and Technology of Zirconia V, eds., S.P.S. Badwal, M.J. Bannister and R.H.J. Hannink (Technomic Publishing Co., Inc., Lancaster, USA, 1993) p. 752. C. Milliken, D. Tucker, S. Elangovan and A. Khandkar, in Fuel Cell Seminar, Phoenix, Arizona, November 1990. E. Ivers-Tiff6e, M. Schiel31, H.J. Oel and W. Wersing, in Proceedings of Third International Symposium on Solid Oxide Fuel Cells, eds., S.C. Singhal and H. Iwahara (The Electrochemical. Soc. Inc., Pennington, NJ,1993) p. 613. S.P.S. Badwal and ET. Ciacchi, J. Appl. Electrochem., 16 (1986) 28. S.P.S. Badwal and D.J.M. Bevan, J. Mater. Sci., 14 (1979) 2353. S.P.S. Badwal, D.J.M. Bevan and J.O'M. Bockris, Electrochimica Acta, 25 (1980) 1115. S.ES. Badwal, M.J. Bannister and M.J. Murray, J. Electroanal. Chem., 168 (1984) 363. S.ES. Badwal, ET. Ciacchi and D.K. Sood, J. Mater. Sci., 21 (1986) 4035. S.ES. Badwal, J. Electroanal. Chem., 202 (1986) 93. S.ES. Badwal, ET. Ciacchi and J.W. Haylock, J. Appl. Electrochem., 18 (1988) 232. T. Kawada, N. Sakai, H. Yokokawa, M. Dokiya, M. Mori and T. Iwata, J. Electrochem. Soc., 137 (1990) 10. D.W. Dees, T.D. Claar, T.E. Easler, D.C. Fee and F.C. Mrazek, J. Electrochem. Soc., 134 (1987) 2141. S. Murakami, Y. Akiyama, N. Ishida, T. Yasuo, T. Saito and N. Furukawa, in Proc. Second International Conference on Solid Oxide Fuel Cells, eds., E Grosz, E Zegers, S.C. Singhal and O. Yamamoto (Commission of the European Communities, Brussels, 1991) p. 561. I. Yasuda, T. Kawashima, T. Koyama, Y. Matsuzaki and T. Hikita, in Proc. International Fuel Cell Conference, Makuhara, Japan (NEDO, 1992) p. 357. S.E Jiang and S.ES. Badwal, 185th Electrochemical Society Meeting, San Francisco, May 22-27, 1994. M. Suzuki, H. Sasaki, S. Otoshi and M. Ippommatsu, in Proc. Second International Conference on Solid Oxide Fuel Cells, eds., E Grosz, E Zegers, S.C. Singhal and O. Yamamoto (Commission of the European Communities, Brussels, 1991) p. 585 K.E. Swider and W.L. Worrell, in Proc. Second International Conference on Solid Oxide Fuel Cells, eds., E Grosz, E Zegers, S.C. Singhal and O. Yamamoto (Commission of the European Communities, Brussels, 1991) p. 593. W.L. Worrell, Solid State Ionics, 52 (1992) 147. M.T. Colomer, J.R. Jurado, R.M.C. Marques and EM.B. Marques, in Proceedings of Third International Symposium on Solid Oxide Fuel Cells, eds., S.C. Singhal and H. Iwahara (The Electrochemical. Soc., Inc., Pennington, NJ, 1993) p. 523. S.ES. Badwal, Solid State Ionics, 52 (1992) 23.

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90 91

92 93 94 95 96 97 98 99 100 101 102 103 104 105 106 107

R. Manner, E. Ivers-Tiff6e and W. Wersing, in Proc. Second International Conference on Solid Oxide Fuel Cells, eds., F. Grosz, P. Zegers, S.C. Singhal and O. Yamamoto (Commission of the European Communities, Brussels, 1991) p. 715. M. Yoshimura, Ceram. Soc. Bull., 67 (1988) 1950. B.C.H. Steele, J. Drennan, R.K. Slotwinski, N. Bonanos and E. P. Butler, Science and Technology of Zirconia, Advances in Ceramics, 12 (1984) 286. H.J. Rossell and R.H.J. Hannink, Science and Technology of Zirconia II, Advances in Ceramics, 12 (1984) 139. R.H.J. Hannink, D.M. Maher and G. Cliff, J. Aust. Ceram. Soc., 20 (1984) 38. H.J. Rossell, Science and Technology of Zirconia, Advances in Ceramics, vol. 3, eds., A.H. Heuer and L.W. Hobbs (The Amer. Ceram. Soc., Columbus, Ohio, 1981) p. 47. S.P.S. Badwal and J. Drennan, Solid State Ionics, 53-56 (1992) 769. V. Lanteri, R. Chaim and A. H. Heuer, J. Amer. Ceram. Soc., 69 (10) (1986) C258. A.H. Heuer, R. Chaim and V. Lanteri, in Science and Technology of Zirconia III, Advances in Ceramics, vol. 24A, eds., S. Somiya, N. Yamamoto and H. Yanagida (The Amer. Ceram. Soc., Columbus Ohio, 1988) p. 3. M.Shibatu-Yanagisawa, M. Kato, H. Seto, N. Ishizawa, N. Mitzutani and M. Kato, J. Amer. Ceram. Soc., 70 (7) (1987) 503. ET. Ciacchi and S.P.S. Badwal, J. European Ceram. Soc., 7 (1991) 197. Y. Akiyama, T. Yasuo, N. Ishida, S. Taniguchi and T. saito, in Proceedings of Third International Symposium on Solid Oxide Fuel Cells, eds. S.C. Singhal and H. Iwahara (The Electrochemical. Soc., Inc., Pennington, NJ,1993) p. 724. W. Weppner, Solid State Ionics, 52 (1992) 15. S.P.S. Badwal, E T. Ciacchi and R.H.J. Hannink, Solid State Ionics, 40/41 (1990) 882. S.P.S. Badwal and N. Nardella, Appl. Phys. (A), 49 (1989) 13. E. Summerville, Ph.D. Thesis, The Flinders University South Australia (1973). S.P.S. Badwal and J Drennan, J. Aust. Ceram. Soc., 20 (1984) 28. S.P.S. Badwal, J. Mater. Sci., 18 (1983) 3117. S.P.S. Badwal and J. Drennan, Mater. Forum, 16 (1992) 237. R. Chaim, A.H. Heuer and D.G. Brandon, J. Amer. Ceram. Soc., 69 (1986) 243. L. Heyne, in Mass Transport in Solids, ed E Beniere and C.R.A. Catlow, Plenum Press, New York (1983) p. 425. B.C.H. Steele and E.P. Buffer, Electrical Ceramics, British Ceramics Proceedings, 36 (1985) 45. M.L. Mecartney, J. Amer. Ceram. Soc., 70 (1987) 54. Yung-Jen Lin, P. Angelini, and M.L. Mecartney, J. Amer. Ceram. Soc., 73 (1990) 2728. M. Rtihle, N. Claussen, and A.H. Heuer, in Advances in Ceramics, Vol 12, Science and Technology of Zirconia II, eds., N. Claussen, M. Rtihle and A. H. Heuer (The Amer. Ceram. Soc., Columbus, Ohio, 1984) p. 352. T. Stoto, M. Nauer and C. Carry, J. Amer. Ceram. Soc., 74 (1991) 2615 S.P.S. Badwal and J. Drennan, J. Mater. Sci., 22 (1987) 3231. J. Drennan and E.P. Buffer, Science of Ceramics, ed., P. Vincenzini, Ceramurgia, 12 (1984) 267. E.P. Buffer and J. Drennan, J. Amer. Ceram. Soc., 65 (1982) 474.

111 108 109 110 111 112 113 114 115 116. 117 118 119 120 121 122 123

S. Rajendran, J. Drennan and S.P.S. Badwal, J. Mater. Sci. Letts., 6 (1987) 1431. S.P.S. Badwal, J. Mater. Sci., 18 (1983) 3230. F.T. Ciacchi, K.M. Crane and S.P.S. Badwal, unpublished work. J.G. Allpress and H.J. Rossell, J. Solid State Chem., 15 (1975) 68. S. Suzuki, M. Tanaka, and M. Ishigame, Jap. J. Appl. Phys., 24 (4) (1985) 401. J.G. Allpress, H.J.Rossell and H.G. Scott, J. Solid State Chem., 14 (1975) 264. R.L.Withers, J.G. Thompson and P. Barlow, J. Solid State Chem., 94 (1991) 89. T.R. Welberry, R.L. Withers, J.G. Thompson and B.D. Butler, J. Solid State Chem., 106 (1993) 461. EW. Poulsen, J.B. Bilde-Sorensen, K. Ghanbari-Ahari, G.G. Knab and M. Hartmanova, Solid State Ionics, 40/41 (1990) 947-951. M.J. Verkerk, A.J.A. Winnubst and A.J. Burggraff, J. Mater. Sci., 17 (1982) 3113. S.P.S. Badwal and J. Drennan, J. Mater. Sci., 24 (1989) 88. S.P.S. Badwal, Appl. Phys. (A), 50 (1990) 449. S.P.S. Badwal, ET. Ciacchi, M.V. Swain and V. Zelizko, J. Amer. Ceram. Soc., 73 (1990) 2505. J. Drennan and S.P.S. Badwal, in Science and Technology of Zirconia III, Advances in Ceramics, vol. 24B, eds., S. Somiya, N. Yamamoto and H. Yanagida (The Amer. Ceram. Soc., Columbus Ohio, 1988) p. 807. M. Kleitz, C. Pescher and L. Dessemond, in Science and Technology of Zirconia V, eds. S.P.S. Badwal, M.J. Bannister and R.H.J. Hannink (Technomic Publishing Co. Inc., Lancaster, USA, 1993) p. 593. J. Drennan, M.V. Swain and S.P.S. Badwal, J. Amer. Ceram. Soc., 72 (1989) 1279.

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Science of Ceramic Interfaces II J. Nowotny (Editor) 9 1994 Elsevier Science B.V. All rights reserved.

113

APPLICATION OF LOW ENERGY ION SCATTERING TO OXIDIC SURFACES

H.H. Brongersma, P.A.C. Groenen z and J.-P. Jacobs

Faculty of Physics and Schuit Institute of Catalysis, Eindhoven University of Technology, P.O. Box 513, 5600 MB Eindhoven, The Netherlands

Abstract Low-energy (0.1-10 keV) ion scattering (LEIS) is used for the determination of the atomic composition and structure of oxide surfaces. Its extreme surface sensitivity enables the selective analysis of the outermost atomic layer. It is precisely this layer that is largely responsible for many chemical and physical properties of oxides. Lowering of the surface energy provides a strong driving force for segregation to the surface. Surface segregation is very important at temperatures that are high enough to enable diffusion. Since the surface enrichment is often restricted to the outermost atomic layer, the unique capability of LEIS to analyze this layer selectively, is one of the main applications of the technique. Other applications pertain to studies of the mechanism of oxidation of metals, semiconductors and to that of the growth of oxides on itself and on other materials. For equilibrated surfaces of spinels (powders or sintered) both LEIS and chemical methods show that, while cations in octahedral sites are accessible, cations in tetrahedral sites are not. This suggests that the equilibrium surfaces of these oxides are dominated by only one or two crystallographic planes. For single crystals, blocking, shadowing and focussing effects of the incident ions or scattered ions have been used to determine the location of surfce atoms as well as the presence and annealing characteristics of surface defects. Although the number of such studies are very limited, the results show great promise for the understanding of oxide surfaces.

2Present address: University of Twente

114

1. Introduction

1.1 An introduction to LEIS on oxides

Elastic binary collisions of keV rare gas ions with surface atoms provide energy spectra that are characteristic for the distribution of the masses of the surface atoms. Already in the early sixties, Panin [1] and Walther and Hintenberger [2] had demonstrated that for inert gas ions a clear correlation exists between the energy loss of a scattered ion and the identity of the surface atoms, and the work of Smith [3,4] gave a strong impulse to the use of low-energy ion scattering in surface science. While the energy spectra of the scattered ions are mainly used for composition analysis, shadowing and blocking effects enable one to use angular dependent studies to determine the location of surface atoms on the surface of single crystals. For general applications the Low-Energy Ion Scattering technique is called LEIS or ISS (Ion Scattering Spectroscopy). Here, LEIS is preferred since the emphasis is on the low energies and ISS does not make that distinction. For the more specific applications, especially for various types of LEIS in structure analysis of surfaces, many other abbreviations exist (ICISS [5], NICISS [6], NBISS [7], LENS [8], TOF [9], TOF-SARS [10]). The most important feature of LEIS is its extreme surface sensitivity when inert gas ions are used. This feature finds its origin in the high neutralization probability and scattering cross sections for these ions at energies of at most a few keV. The high neutralization probability ensures that almost all ions that penetrate beyond the first atomic layer are neutralized. When energy analyzers are used that only detect ionized particles, one thus effectively eliminates scattered particle contributions from deeper layers. In contrast to other commonly used surface analytic methods such as Auger Electron Spectroscopy (AES), X-ray Photoelectron Spectroscopy (XPS) or Electron Spectroscopy for Chemical Analysis (ESCA) and Secondary Ion Mass Spectrometry (SIMS) that probe several atomic layers, it was shown by Brongersma and Mul [11,12] that LEIS can be used to selectively study the outermost atomic layer of a surface. It is this atomic layer that plays an important role in many applications of oxides. Typical examples of such applications are: catalysts identity of active sites or of the species causing poisoning spreading of the active phase over the support quantification of the composition of the active phase ceramics surface active species promoting sintering oxygen transport in fuel cells oxygen sensors cathodes origin of low workfunctions adhesion promotion or prevention corrosion

115 protection of metals Especially for single crystals and sintered ceramics, which have a low specific surface area, segregation of impurities or dopants may dominate the surface properties. The driving force for segregation of alkalis is often so strong that after moderate heating a large fraction of the surface is covered by alkalis (even if the bulk concentration is only in the ppm or sub-ppm range; see section 6.2) and thus the surface properties are altered significantly. Since thermodynamics often restricts the enrichment to the outermost atomic layer, LEIS is the ideal technique to detect and quantify segregation. In most applications of LEIS the main reason for using LEIS is its monolayer sensitivity. In general, the reason for a composition analysis by LEIS is related to the following features: - selectivity for 1st atomic layer samples can be conductors as well as insulators in-situ studies at high temperatures straightforward quantification - very rough surfaces possible detection of isotopes. Some of these features are particularly advantageous for studies of ceramics and catalysts and will be discussed in more detail below. -

-

-

-

Surface sensitivity In the early days of LEIS full advantage of its surface sensitivity was not always taken, because the ion beams were often so intense that the first monolayer was destroyed before the analysis was completed. Nowadays, the fluence required for a composition analysis can be as low as 3 . 101~ ions/cm 2 [13], thus allowing many measurements before the damage becomes detectable (so-called Static LEIS). This also allows one to follow processes such as surface segregation in real time.

Quantification; rough surfaces Quantification of the atomic composition by calibration against standards is now well established. This means that, in general, the sensitivity for a given element is not affected by the identity of the neighboring atoms ("no matrix effects") as long as the element is present in the outermost atomic layer. Some exceptions to this rule exist, but these can be avoided by proper choice of the ion and energy. The quantification can even be carried out for very rough surfaces, such as encountered for supported catalysts.

Isotopes The scattered ion energy depends on the mass of the target atom. Thus, at least in principle, it is possible to distinguish the isotopes of elements. Use is being made of this property for studies of oxidation mechanisms [14] of metals, quantification of oxygen diffusion in ceramics for sensor applications [15] and to distinguish oxygen surface sites via selective oxygen exchange [16].

Insulators The relative ease of producing electrons of a few eV and the fact that in contrast to most SIMS experiments no draw-out fields are being used, makes it simple to compensate charging effects by the incoming ion beam. LEIS is also not very sensitive to surface charging of a few volts. This enables one to investigate insulators as well as conductors.

High temperature Since high temperatures of the sample do not affect the analysis, it is possible to follow, for instance, surface segregation in ceramics in-situ. A restriction is, of course, that the

116

vapor pressure of the material under investigation is low enough to allow for undisturbed passage of the ion beam (pressure lower than 10 4 mbar) and, more importantly, to prevent significant contamination of the equipment.

1.2 Principle of LEIS When inert gas ions are used for the analysis, neutralization is very high. Thus, the probability that these particles are still in an ionized state after interaction with more than one target atom is strongly reduced. If only ionized particles are detected, as is often the case, singly scattered ions are strongly favored over doubly- and multiply-scattered ions. For the singly scattered particles, a simple two-body collision model is sufficient to describe the interaction of an ion with a solid target. If one considers an ion with incident energy Ei and mass Mio, which is scattered through an angle O by a target atom of m a s s Mat , the final energy Ef of the ion will be E f -- E i

(c~

~] l+r

2 = E i 9k

(1.1)

where the positive sign applies to r z 1 and both the + and - are both solutions if 1 z r [sinO[. The mass ratio r = Mat / Mio., and k is the so-called kinematic factor. The equation is simply obtained by solving the equations for conservation of energy and momentum. During a given experiment, the energy spectrum of the scattered ions at a particular scattering angle is recorded for a fixed energy of the incident ions. The energies of the scattered ions are then characteristic of the masses of the target atoms. Fig. 1.1 illustrates this for 1360 elastic scattering of He* ions by an oxygen exchanged Sr-Sm-Co-oxide [15]. The peaks for Sr, Sm, Co, O t8 and O t6 are clearly observable. In addition, ions are also observed with a most probable energy near 0 eV. These particles are ions sputtered from the target. Since light ions, such as H t, contribute to this peak, an intense peak at low energies is generally a strong indication of a contaminated (water, hydrocarbons) surface. Since the angles of the incoming and outgoing ion with the target surface (ixi and (zf, respectively) do not enter eq.l.1, the inclination of the target surface with respect to the ion beam does not influence the energy of a binary collision peak. This explains why LEIS can still be used for very rough surfaces (see section 3.7). At very low angles (0 ~ 200) of the incoming or outgoing beam, shadowing and blocking will, however, influence the intensities of the peaks. This dependency is used in structure analysis. In a number of LEIS studies, scattering of non-inert gas ions, such as H § Li § and Na § has been used. Since neutralization no longer limits the information depth, a very wide, featureless spectrum of scattered ion energies is obtained. In such cases it is important to restrict the studies to single crystals and to use shadowing and blocking directions to further restrict the scattered ions. Although such studies can be quite successful in special cases, very few applications to oxides or oxidation [17,18] are known. The present paper is, therefore, restricted to the use of inert-gas ions. For composition analysis of the surface by LEIS, a large scattering angle (1400 - 180 ~ is favored to obtain a good mass resolution, although the scattering cross section is at its

117

1500

.>,

Sm

Sr

1000

e,-

500

O

.

0

.

.

.

I

500

,

i

i

i

I

I

1000

9

Final energy (eV)

'

1500

Figure 1.1 Energy spectrum of 1500 eV 4He* scattered from 160/z80 exchanged Sm0.sSr0.2CoO3 [15].

lowest here. Both the incoming (often perpendicular to the surface) and outgoing trajectory of the ions make large angles (preferably at least 400) with the surface of the sample (see fig. 1.2a) to reduce additional neutralization, shadowing and blocking effects that complicate a quantitative analysis (section 3.5). In backscattering, all atoms having masses greater than the lightest inert gas ion (3He+) can be detected (eq. 1.1). Thus, apart from hydrogen, it should be possible to detect all elements of the periodic system. Although this is essentially true, the signals for the elements with a mass close to that of the probing ion (Li, B, Be, C and N) are quite low. The difficulties of detecting the light elements present a potential danger to the quantification of the surface composition. Although these elements cannot be detected, even atoms as small as hydrogen on the surface do neutralize or physically shield the incoming or outgoing ions and thus prevent the detection of the underlying atoms. At low angles with the surface, the local atomic structure dominates the scattering process (fig. 1.2b). To simplify the interpretation, one of the beams (generally the incoming beam) is often at larger angles with the surface, while the other (outgoing) beam is glancing. Another attractive choice is to use 1800 scattering (Impact Collision Ion

118

Figure 1.2 Illustration of incoming and outgoing ion beams for a. composition analysis b. structure analysis (TOF-SARS) c. structure analysis (ICISS)

Scattering Spectroscopy, ICISS) so that the incoming and outgoing beam coincide (see fig. 1.2c for an illustration). The angular distribution (azimuthal and or polar) gives direct information on the atomic structure, defects, etc. of a surface (see section 4). If the binary collision is such that a target atom is ejected (recoil particle), the energy E r of this particle is given by E r = E i

4r (l+r)Z

cos2~

(1.2)

where ~ is the recoil angle. For the light elements the detection as a recoil particle is generally much more sensitive than in scattering. This is especially true for hydrogen that, since it is lighter than any inert gas ion, cannot even be observed in backscattering. Rabalais and coworkers [10] have developed a technique to perform surface and absorbate structural analysis using time-of-flight scattering and recoiling spectrometry (TOF-SARS).

1_3 Applications of LEIS The specific advantages of LEIS for surface investigations of ceramic materials were recognized from the early start of the technique. Smith [4] observed that when AI is oxidized to A1203 the reduction in the AI signal corresponded precisely with the expected reduction of the A1 concentration (no matrix effect), see also ref. [19]. Honig and Harrington [20] used LEIS in combination with sputtering for depth profiling of a Ta205 film on Ta. Since the beginning of LEIS some 300 papers have now been published on the use of various types of LEIS for studies of oxides, or oxidation. In table 1.1a and 1.1b a selection of these papers is classified according to material and subject.

119 Most studies concentrate on the composition of the first atomic layer. The main application has been for supported catalysts. It seems, however, that many of the findings in this area should have important consequences in other fields. There are only a few studies of the surface structure of monocrystals of oxides, although some of them are very detailed (see section 4). Review papers on the combination of LEIS and oxides have been given by Honig [20], McCune [21], Carver et al. [22], Horell and Cocke [23], Brongersma and van Leerdam [24] and by Taglauer [25]. In most studies discussed in the present review, various types of LEIS are combined with other surface techniques. Such combinations are extremely important since these techniques often provide essential complementary information. For instance, the larger information depth of XPS, AES and SIMS provide in combination with LEIS, information on the first and the deeper atomic layers. XPS and SIMS also provide chemical information. Low-Energy Electron Diffraction (LEED) is essential for studies of single crystal surfaces to establish the occurrence of long range periodicity and the dimensions of the unit cell. On the other hand, LEIS provides, in a relatively easy and straight forward way, the short range order within the unit cell. In the present paper the emphasis is on the possibilities and achievements of LEIS in surface studies. For the sake of conciseness the other techniques used by the authors are only occasionally mentioned, although it is realized that this often does not do proper justice to the full extent of the arguments presented in a paper. SiO2; glass

A1203

MgO

ZrO2

VO,

SnO2

,

composition

234, 240, 104, 158, 275, 33, 291, 292, 301, 101,321

104, 210, 19

221,239, 165, 280 99

structure

61, 33, 291,292

81, 125, 259, 279, 296

221,229, 165 239, 99, 6, 118

115

235, 135, 16, 120

diffusion segregation

169, 232, 158, 275, 33, 291, 167, 170, 171

201, 153, 278, 297

6, 153, 36, 278

165, 40

115, 34, 303

41

growth

188, 187, 271, 180, 57

212, 184, 185, 215, 178, 217, 216

40, 179, 263, 57, 187, 188

140, 194

181

adsorption desorption

232, 260, 104, 271

104

187, 179

303

181

depth profile sputtering

292, 293, 101, 70

246, 247

neutralization reionization

61, 104, 286

61, 104, 110

99, 265

230, 57

118

165

235

165, 179, 40

198, 181

anneal

260

81

oxidation reduction

271, 240

217

dispersion

24, 130, 192, 189, 260

81, 138, 137, 22, 37, 64, 140, 136, 190, 189

Table 1.1a

221, 229

214, 235, 135, 242, 198

165, 246, 247, 263

A selection of papers on LEIS and oxidic surfaces.

115, 34

120

FeO,

YBCO

i ceramics and i other oxides

catalysts

Si

Me~s

I

composition

222, 246

structure

262, 264, 165

224, 238

112, 130, 51, 19, 226, 231,233, 235, 270, 299, 295, 100, 21, 312, 22, 55, 122, 57, 121, 61

209, 226, 235, 236, 220, 135, 246, 248, 249, 234, 253, 149, 259, 258 136, 268, 143, 202

H.

|

|

224, 116,' 125, 126, 128, 100, 238 129, 235, 56, 114, 316, 116, 16, 120

269, 273, 283, 285, 234, 284, 288, 296, 297, 251 298, 300, 304, 304, 138, 308, 313, 81, ,235, 135, 248, 253, 136, 268, 134, 25, ,24, 23, 82, 141, 142, 143, 144, 129 i |

diffusion segregation

19, 183, 185, 184, 218, 250, 248, 252, 256, 257, 276, 282, 325, 323, 311 212, 184

i

145, 182, 39, 147, 302, 165, 102, 286, 149, 201

126

258, 277

237, 274, 311, 182, 322, 328

growth

107, 200, 264

187, 188, 108, 228, 189

81,249, 261, 137, 180, 269, 285, 284, 300, 277 139

212, 213, 107, 219, 227, 108, 183, 22, 107, 195, 196, 198, 199, 200, 130

adsorption desorption

107, 264

128,208

273, 283

212, 219, 185, 254, 272, 274, 177, 317

302, 315, 243, 245, 287, 293, 20

80, 84, 309

depth profile sputtering neutralization reionization

224

anneal

116

oxidation reduction

246, 149

dispersion

211

239

123, 230, 108, 286, 80, 142, 232, 241 294, 56, 314, 55, 61, 52, 320, 324, 47, 329 145, 235 81, 235, 259, 137, 267, 269, 273, 131, 313 39, 48, 49 273, 297, 298, 246, 180, 149, 195, 255 277

189, 270

269, 281, 81,225, 236, 253, 37, 137, 266 ,.,

Table 1.1b

180, 223

A selection of papers on LEIS and oxidic surfaces

218, 250, 252, 276, 289, 323, 328 108

207

184, 218, 219, 185, 227, 228, 183, 2, 244, 19, 173, 174, 175, 177, 178, 179, 181,252, 254, 272, 274, 306, 307, 310, 315, 318, 326, 327 203, 204, 205, 206

121

2. Experimental 2.1 Introduction An analysis of the composition of a surface is based on the application of eq. 1.1. For a well-defined ion beam (Ei, Mio~ the mass resolution would thus be fully determined by the precision with which the scattering angle O and the energy of the scattered ions Ef are determined. In reality this is not completely true since inelastic processes (giving an extra energy loss) and vibration of the surface atoms (giving a Doppler type of broadening in the energy [26]) will also influence the energy of the scattered ions. However, for most experimental conditions the inaccuracy in O and the energy resolution dominate the results. At present, most analyses of oxides by LEIS use composition rather than structure analysis. The emphasis in section 2.2 will, therefore, also be on experimental arrangements for studies of the composition. Also, most materials that are used in applications, do not have a structure with long range order.

2.2 Energy analyzers for scattered particles Two methods are commonly used to determine the energy of the scattered ions.

2.2a Time-of Flight (TOF) method The principle of the time-of-flight method is illustrated in fig. 2.1 [27]. The primary ion beam is chopped to produce short (order of 20 ns) ion pulses. The scattered particles are energy analyzed by measuring their time-of-flight (TOF) to a detector that is 0.4 - 1 m removed from the target. When the detector is mounted in the forward direction, one cannot only observe scattered but also recoil particles ( target atoms or ions that are emitted with an appreciable fraction of the incident energy after being hit by a primary ion). These particles are detected irrespective of their charge state. The detection of recoil particles provides an unique possibility to detect light atoms, such as hydrogen, on a surface. Large angle scattering, preferably 180 ~ gives a good mass resolution for the backscattered ions. Rabalais and coworkers [28] have developed a very sophisticated TOF system with a rotatable detector enabling scattering angles from 0 - 1650 to be studied. The capability of rotating the detector is very valuable, since the energies of the scattered and recoil particles depend differently on the scattering (recoil) angle. Aono [29] developed a technique called Impact Collision Ion Scattering Spectroscopy (ICISS) where ions are scattered through a fixed angle which is essentially 180 ~ The setup is illustrated in fig. 2.2. The interpretation of the TOF (energy) spectra is then fully unambiguous, since recoil particles (only forward scattering) are not detected. Also, since the ingoing and outgoing ions move along the same trajectory, shadowing and blocking effects are essentially the same thus facilitating the interpretation and enhancing the signal. By post-accelerating the scattered charged particles (ion acceleration tube in fig. 2.2), it is possible to distinguish them from the scattered atoms. The analysis of the ions has the

122

Figure 2.1 Principal of time-of-flight, Cu8sPdls(ll0) using 1.5 keV Ar § [27].

illustrated

using

a

TOF-spectrum

of

advantage of being surface sensitive, while the neutrals (large majority) give much higher signals and allows the investigation of deeper layers, interfaces, etc. Both TOF systems have proven to be extremely valuable in studies of the structure of the surface and deeper layers [10,29] of single crystals. 2.2b Electrostatic energy analyzers In an electrostatic energy analyser the charged particles are deflected in an electrostatic field. The deflection depends on the energy of the ions and the energy can thus be determined. In the early years of LEIS, 1270 analyzers [30,31] 1800 Hemi-Spherical Analyzers (HSA or CHA) [32] and 900 electrostatic analyzers [33] were used. In modem times, a cylindrical (Cylindrical Mirror Analyzer, CMA) or a hemi-spherical electrostatic analyzer with improved transmission has been used. These analyzers are used for LEIS studies of oxides (CMA: e.g. [34,35,36,37] HSA: [38, 39,40,41]). The CMA has the advantage that when the ion beam is coaxial, the scattering angle is the same for the full 3600 of the azimuth. For composition analysis, for which no azimuthal information is required, one can thus use the full 3600 of the azimuth of the analyzer. This greatly increases the detection efficiency. In commercial systems the ion source is incorporated in the inner cylinder of the analyzer. In the first CMA for LEIS analysis, developed by Brongersma et al. [42], a somewhat different design with a ringdetector was used. This enables one to separate the ion source from the analyzer and

123

ion microchannel acceleration plates deflector 9 tube beam scanning oi....ol \ grounded[ chopping / einzel , aenector ",~" \ mesh l anode aperture / lens

,.sa~p,e

_w/ .

\

,r

\

~~-~--'x~-~

Nx..l__// / ~~f~"

/

~'

~

t~

ch_o~p_m_ g .... ion pL~c~ com mator source

.....j.._~/ ../

~ ~

,

_"Z'~/ILr --~ ~

/

/~

| lensl ~microchannel I , I ~ ~l lv~ beam scanning I voltage i

[

/ P late v~

acce .~ration [ tube voltage !

I l

[ preamplifier 1

t__

t _

stop

start

Figure 2.2 Schematic of ICISS. A pulsed-beam low-energy ion source and a time-offlight (TOF) energy analyzer are placed coaxially (CAICISS) so as to take the experimental scattering angle at 1800 [29].

thus to differentially pump and mass-filter the primary beam. The exclusion of ions such as H § and C § from the primary ion beam results in lower background in the spectra (such ions could still leave the surface as an ion after many collisions in deeper layers; see [24]). An improved version of this set-up ("NODUS"), see fig. 2.3, is still used for most of our LEIS studies of oxides. In fig. 2.4 a cross-section of the ERISS (Energy Resolved Ion Scattering Spectrometer) is shown. The analyzer is rotationally symmetric around the primary ion beam. The scattered ions are energy separated in a double toroidal electrostatic analyzer. The configuration is such that a linear image in energy is obtained. A special 2-dimensional detector allows the simultaneous detection of 100 energy channels. Because of this simultaneous detection and some other improvements in the design, the sensitivity of this instrument is about a factor of 800 higher than that of conventional electrostatic analyzers. This allows one to use ion currents as low as 10 - 100 pA, thus reducing damage of the target surface by the ion to negligible amounts ("static" LEIS). ERISS is analogous to the earlier developed EARISS [43,44,45] but has been specially modified for composition analysis of the surface of insulators. Amongst other features, a special charge neutralizer compensates the surface charging (see section 2.3a).

124

Figure 2.3 Diagram of the CMA of the NODUS LEIS-machine.

2.3 Experimental factors complicating the analysis A brief discussion is given of the main experimental factors that may influence energy spectra and thus their interpretation. Some are related to the set-ups, others to the target preparation and the poor electrical conductance of most oxides.

2.3a Charging effects. The primary ion beam will positively charge the surface of insulators. For good insulators the charging continues until no further ion can reach the surface. For perpendicular incidence of the ions this means that the surface potential becomes equal to the accelerating potential applied to the ions. Subsequent ions are reflected by the field. Such reflected ions, which have not reached the target and thus did not lose any energy during scattering, are sometimes used for energy calibration of the analyzer or the ion source. When the charging is less severe, the scattered ion peaks will still be shifted to higher energies by the charging. Spraying the target with low-energy electrons can reduce the effects of the charging. This is illustrated in fig. 2.5 for a cobalt aluminate surface with and without proper charge neutralization. Surface charging not only shifts the peaks in a LEIS spectrum, it may also enhance the mobility of atoms [46] in the target and thus change the surface composition (especially for alkali ions). Even for a charged surface, peak identification is still possible. In a first approximation, charging to a potential Vch will have decelerated the incoming ion to E i - eVch when it

125

Figure 2.4 Diagram of the analyzer and detector of the ERISS.

reaches the target. According to eq. 1.1 the energy after scattering will be k (E 0 - eVc~). Upon leaving the surface the ions are post-accelerated by the surface charge to the detected energy :

E 7 = E/ + ( l - k ) e V h

(2.1)

In agreement with fig. 2.5 the effect of charging is thus largest for the lowest kinematic factor k (the lightest elements). For large scattering angles,such as encountered in a CMA with a coaxial ion gun (see section 2.2b), eq. 2.1 is a good approximation. A more accurate analysis should be used, however, to take account of the precise experimental

126

i|l

I

;, ;,

600-

I f !

c

9

m

_J

! I t ! I l

' '

4O0-

c "T r~

xl/3 t I ! t

I t I I l !

mm

I

!

'

'I

! '

! '

I

I

I I

! I

!

.!

i

200

O I

I

~,

,,

'

'

|

\ \

t

,

CO

AI

, /!

~!

I1

!~,~

-

pl I ,

I

;

',

;t,l/

'~__~

0

'

I

1000

'

I

2000

~.._

'

Final e n e r g y ( e V ) Figure 2.5 LEIS spectra (3 keV He § from proper charge compensation.

a

COEAIO4 catalyst with and without

conditions in a more general situation [47]. Surface charging will not only accelerate the scattered but also the sputtered ions. Especially for oxides, the sputtered ion signal is generally quite large. The most probable energy of these ions is just above 0 eV. The intensity of this peak thus decreases with increasing energy. Since the minimum kinetic energy of a sputtered ion that leaves the surface is 0 eV, these ions will attain a minimum energy of eVch after acceleration. This will lead to a sharp onset on the low-energy side of the spectrum. Although the precise energy of the onset will depend on the geometry of the set-up (on the dimensions of the front end of the analyzer, sample holder and beam spot), the exact value of eVch (and thus the energy correction for the scattered ions) can be easily be estimated for a given experimental set-up. A quantitative analysis of the surface composition becomes difficult if the surface charging is appreciable. This is especially the case for very rough surfaces, such as encountered with catalysts, where surface charging is very inhomogeneous. Fig. 2.6 illustrates this. A proper charge neutralization is then imperative. Since the primary ion current that has to be compensated is only of the order of 0.1 - 100 nA, a simple filament

127

Figure 2.6 Illustration of rough surface compensation by an electron beam

with

ion

beam

and

partial

charge

easily produces enough electrons to do this. Preferably these electrons should come from the same direction as the ions. For the ERISS set-up (fig. 2.4) this is approximately achieved by using a ring filament that surrounds the primary beam. A simple electrostatic deflection keeps the filament out-of-sight for the target. This prevents heating of the target by radiation and the contamination of the target by evaporation from the hot filament. The problems of charging can be circumvented by the use of a beam of neutrals [6], but this makes quantification much more difficult. Another possibility is to increase the electrical conductivity by heating the material. Griinert et al. [48,49] applied this method to zeolites, where the surface charging effectively could be reduced to enable the study of the surface by LEIS, UPS, XPS.

2.3b Target preparation The extreme surface sensitivity of LEIS makes proper target preparation a necessity for a quantitative composition analysis. Powders are pressed into pellets to reduce the macroscopic roughness and thus facilitate charge compensation (see above). It also provides a more reproducible density. The surface can be cleaned by sputtering or by chemical reactions. Sputtering generally affects, however, the atomic structure of the surface and subsurface layers. The most suited chemical reaction depends, of course, heavily on the target composition. This is especially important for oxides where the surface composition depends strongly on the oxygen partial pressure. Margraf et al. [37] use a H 2 beam having an effective pressure at the target of 10 .4 mbar during 1 hour at a target temperature of 600K. This effectively removes hydrocarbon contaminants from the surface. Remaining hydrogen on the surface will strongly neutralize the scattered ions and physically shield the underlying atoms. This affects the quantification (see section 3.4). It is believed however, that a small amount of sputtering will,

128 however, remove most of the hydrogen before sputtering of the target itself becomes too significant. The hydrocarbons and most of the hydroxyl-groups can generally also be removed by treating the target for 15 min. with 20 mbar of oxygen at 250~ After such a treatment the surface can be analyzed without any sputter cleaning.

2.3c Other experimental problems When the LEIS signals are calibrated against those of reference targets, reproducible scattering conditions, beam current and beam profile are important. The precise location of the target will influence the value of the scattering angle O somewhat. Especially for small scattering angles this can affect the interpretation. Helbig et al. [47] have given a more detailed description of this problem.

2.4 Interpretation of energy spectra The energy loss of the ion is mainly due to momentum transfer in an elastic collision between the ion and a surface atom. In addition to this, however, inelastic processes may further lower and broaden the energy of the "elastic" peak. Sometimes one can even distinguish a discrete inelastic loss of about 25 eV (see e.g refs [50,51]). Although the absolute value of this inelastic loss is not very high, it may complicate the interpretation of the spectra (especially at low ion energies). Wheeler (see [52]) reported, for instance, that for 500 eV 4He* scattering through 900 the inelastic energy loss shifted the peak for Ta (mass 181) to an even lower value than that of Ag (mass 108)! In general, however, there is no serious problem in assigning the peaks. This is especially true for impact energies of 1 keV and higher where the inelastic losses are relatively weak. When the forward scattering of heavy ions (2~ 4~247 is studied at low kinetic energies, multiple surface scattering is often observed and may even become the dominant feature in a spectrum. Under such conditions, the final energy for a given total scattering angle can be higher than that expected for single elastic scattering (eq. 1.1). For single scattering conditions, however, inelastic processes shift the peak to values slightly lower than that predicted by eq. 1.1. In practice, the high-energy onset of a peak is generally a good measure for the elastic value and should thus be used to calculate the mass of the surface atom. In fig. 2.7 a characteristic spectrum for 3 keV 4He* ion scattering by aluminum is given. In addition to the elastic scattering peak around 1700 eV, there is an intense background (tail) extending to lower energies. This tail results from ions that penetrate into the target, lose various amounts of energy in multiple collisions with target atoms and are finally scattered back into the direction of the analyzer. In general, the incoming ions are already neutralized during their first interaction with the target surface and will thus escape detection in an electrostatic energy analyzer (only charged particles are detected). It is known, however, that reionization can occur in binary collisions of certain ions with certain atoms (see ref. [50,53] for a list of combinations where reionization is observed). The tail results from ions that suffered such a collision upon leaving the target. Since reionization occurs as a result of the overlap of the inner electronic levels of the incident particle and the target atom, there is an energy threshold for this process. As the onset of the tail in fig. 2.7 illustrates, the threshold for He on AI is around 0.5 keV. De Wit et al. [54] observed for Ne § scattering from Cu that there is an increased

129

600 3 keV

4He+

AI

binarycollision

A ,,iBi

o

.--.,.

400

C m.-.

LU

m

reloilzedj~.~~~.~ '

200

__1

i

0

|

i

i

~

500

"

l

1000

i

multiplecolllslons l

|

i

1500

| ~

~

.

2000

|

I

|

2500

J

|

i

3000

Final energy (oV) Figure 2.7 LEIS-spectrum for He* scattering by AI

background on the low-energy side of the Cu elastic peak when oxygen is adsorbed. This was interpreted as a decrease of the neutralization probability. The importance and origin of the tails have already been described by Nelson [55]. Since the ions have lost additional energy while passing through the deeper layers, the tails are always on the lowenergy side of a peak. The background on the high-energy side of the peaks is thus always much lower and gives a good impression of the intrinsic noise level of the set-up. For LEIS studies of oxides, the reionization of He and Ar by oxygen is particularly relevant. Nelson [55] pointed out that the 2s in oxygen (23.7 eV) is closely matched with the ls level in helium (24.5 eV) and the 3s level in argon (25.3 eV). Indeed, tails are generally observed for He and Ar but not for Ne scattering from oxides. Tongson et al. [56] claim, however, in their experiments of inert-gas ion scattering from various gadolinium compounds, that the low intensity of the tails in the Ne § scattering are due to effective removal of the oxygen by sputtering. The presence of a tail seriously affects the detection of light elements. This is particularly serious since the elemental sensitivities (see section 3.3) for light elements are relatively low. At 3 keV He § scattering (O = 140~ surface concentrations of heavy elements can, therefore, still be detected for concentrations as low as 10 ppm, while a value of about 1 at.% holds for elements such as oxygen. One can improve the relative sensitivity for oxygen, however, by lowering the primary energy [21,52, 315].

130

Ions penetrating into thick targets may loose their full kinetic energy during straggling through the target. Thus scattered particles will have energies from zero up to the elastic binary collision peak. Since only ionized particles are detected, the onset of the tail is determined by the threshold for reionization. Martin and Netterfield [57] used 2 keV He + ions to follow the growth of zirconia on silica. It was found that the low-energy onset of the tail of the elastic Zr peak shifts to lower energies until a thickness of 70 ,~ is reached. Since the tail extends downward from the Zr peak, it is clear that scattering from Zr in deeper layers is involved. Van Leerdam et al. [58] used the tails in 3 keV He § scattering to obtain information on the composition of layers as deep as 60 .~,. Computer simulations of ion trajectories in the solid using the SISS-92 program [59] suggest that not only the atomic composition but also the crystalline structure can influence the intensity of a tail. For crystalline targets the ions are found to be deflected into channelling directions after passing only a few atomic layers. This reduces the chance of backscattering and thus the intensity of the tail in comparison to that of a truly amorphous material. It seems feasible to exploit this phenomena to obtain very unique information on the local crystallinity of a target. This may be very relevant information if the local crystallinity determines the oxygen diffusion through oxides.

2.5 Choice of experimental conditions In principle, any element that is heavier than the ion used for the analysis can be detected by LEIS. The optimal conditions result from a trade-off between scattering probability and neutralization. For large scattering angles, as encountered in a CMA, detection is easy for atoms that are at least 2.5 times heavier than the probing ion. Heavy ions and large scattering angles are advantageous to improve the mass resolution. For instance, with backscattering of Ne § and Ar § ions it is possible to determine the ratio of the isotopes in a Cu target, while this is very difficult with He + ions. A drawback of the use of heavy ions such as Ne § and Ar § is that their sputter rate is about ten times higher than that of He*. To benefit from both the light and the heavy ions, mixtures of ions are sometimes used [20,60]. For samples that are not too complex, it is still possible to assign the peaks to the proper incident ion - target atom combination. In fact, in set-ups that do not allow for differential pumping of the ion source, ion mixtures are often automatically present to a certain degree, because of contamination by inert gas from previous experiments. A wide range of energies (200-- 5000 eV) is used for composition analysis by LEIS. Since the sputter rate at these energies is almost proportional to the energy, many authors prefer the lower energies. Often, the argument is used that increased surface sensitivity is obtained at lower energies. For oxides it has been shown [61], however, that there is no detectable difference in the screening between ions of 500 and 3000 eV. Higher energies have the advantage of somewhat better control of the ion beam, while the effects of charging and contamination are less severe. In practice, energies between 500 and 3000 eV are commonly used and are a good compromise for the conflicting demands. Beam currents of a few pA to a few ~A are used with spot sizes of the order of a mm 2. Although sputter rates vary, of course, from target to target, a characteristic removal rate for 500 eV 4He* is 1 - 2 atomic layers [37] for a fluence of 2 . 1016 ions/cm 2.

131

2.6 Fitting of LEIS spectra As discussed above, the LEIS spectra can be divided in most cases into three contributions. The surface peak, which contains the surface sensitive information, and a background, which originates from reionized particles scattered from deeper layers. Generally, the backgrounds are relatively large in oxides. This greatly hampers the quantification of the LEIS signals, especially for the lighter elements. To extract the surface sensitive information from the spectra a number of methods have been put forward. However, all are based on (semi) empirical considerations since the neutralization and reionization processes are not well understood (see section 3.2). For materials were the tails in the spectra do not depend too strongly on the matrix (which is more the exception than the rule), Baun and Solomon [62] proposed subtraction techniques to analyze the spectra. A linear background subtraction has proven a very effective easy way to obtain peak areas, see for example refs [63,64] for a nice illustration. Also the use of peak heights can provide in many cases a reliable indication for the changes in surface composition. It has been found that the final results are identical for a number of cases [61,65] when peak heights are used instead of peak areas obtained after linear background subtraction. Empirical fitting procedures of background subtraction and peak identification have been introduced by Nelson [66] and were extended by Young and coworkers [67,68] This procedure allows for deconvolution of overlapping peaks but due to the number of parameters, convergence is sometimes difficult. Creemers et al. [60] describe a relatively simple but fast iterative procedure to obtain peak areas.

2.7 Change of the surface composition by the analysis The analysis itself may change the surface composition of a target. Firstly, vacuum is necessary to perform the LEIS measurements. The stability of the surface under such conditions is a necessary requirement to obtain relevant information. For oxides, however, the oxygen partial pressure is known to influence segregation equilibria [69]. A general requirement for LEIS measurements is that the pressure of the residual gas is low enough to prevent discharges between lens elements that are at high voltage. Another, generally less important, requirement is that collisions of the primary and scattered ions with the residual gas can be neglected. Unless special precautions are taken, the maximum pressure is around 10 .5 mbar. So, in principle, it is possible to vary the oxygen partial pressure between 10 s and 10 -1~ mbar. Up to now there are hardly any studies where this is actually done during the measurements. For a surface sensitive technique such as LEIS, preferential sputtering and damage by sputtering can also change the surface composition. For large enough fluences, preferential loss of oxygen is observed for most oxides. Pitts and Czanderna [70] found that even a very stable oxide such as SiO 2 is reduced by ion bombardment. A fluence of 7 mC/cm / (4.1016 He + ions/cm 2) of 1 keV 3He § reduces the oxygen content by 4 at.%. Both a collisional [71] and a thermal model [72] have been described (see also

132 [73,74]). The fluences used in LEIS (see section 2.5) are, however, generally too low to change the surface composition during an analysis. According to Saied et al. [75] one can reduce the damage by using beams of neutralized ions rather than the ions themselves. The quantification of the composition from the resulting scattered ions is, however, more complex for incident ions. In section 5.3 it is illustrated for alumina that damage can alter the surface composition of a material that does not show preferential sputtering.

2.8 Compositional depth profiling Besides the surface information, the in-depth distribution of the elements in a solid are of great interest. Usually depth profiling is performed combining a sputter ion beam with surface analysis by AES or SIMS. The bombardement time can be converted into a depth scale. It is clear that also LEIS could provide this information. Here the primary ion beam can also be used to sputter the material, but it has an advantage to use a separate ion beam for sputtering, as will be discussed below. In this section it is not the aim to give a detailed description of the processes that rule the sputter assisted compositional depth profiling, but a few pittfalls and the use of LEIS on oxides will be discussed. For a more thorough approach the reader is referred to ref. [76,77,78]. Depth profiling is mostly performed over large areas. The information will therefore be averaged over this area. The conversion of bombardement time into depth scale is not trivial. The sputter rate can be constant, when only one major component is present, but can also vary in depth. Assuming the rate is constant, measuring the crater depth after analysis can give the erosion rate. Another way is to measure the weight loss during profiling using a thickness monitor. However, in many cases no depth scale is given or the conversion is based on a rough estimate. When the analyzing and sputtered area are the same the signal from the crater edge will distort the depth profile. The fact that the ion-current distribution is not uniform can enhance this effect. Rastering of the ion beam and reducing the analysed area will overcome this problem. It is clear that when in LEIS the primary beam is also used for the sputtering the analyzing spot should be smaller than the diameter of the ion beam. The composition measured at the surface does not have to reflect the in-depth information due to a number of reasons. Firstly, as discussed in the previous section, preferential removal of one component is possible. While in the analysis this effect can be minimized by decreasing the ion dose, this problem is inevitable when performing ion-assisted depth profiling. Ion beam mixing is a well-known phenomena which can damage the structure, see ref. [79]. Radiation enhanced diffusion in particular of alkali metals can deplete the solid of this material, see section 6.5. Surface roughness, intrinsic or induced by the ion beam, can hinder accurate depth profiling. Nevertheless acurate depth profiling can be performed although for this SIMS and AES are most commonly used. For very shallow depth information, the sensitivity of LEIS for the outermost atomic layer gives a distinct atvantage. LEIS is, for instance, very well suited to determine the dispersion of a deposited species as is schematically shown in fig 2.8. For monolayer dispersion an exponential decay of the LEIS-signal is expected, while for a cluster the signal will stay constant until the top layers are sputtered away, after

133 which a transition to an exponential decay is expected. In catalysis, LEIS is frequently used in combination with depth profiling to obtain the near-surface compositonal profile. This is possible even for the highly porous very rough surfaces these catalysts mostly are. But, of course, the depth resolution is only small since I

1

>,~

>,,

.,,,.,

u~

c (D

._c d) i,,,

1.1,1 .,.I

I.IJ ..I

i

l

ion fluence

i

,

i

i

I

ion fluence

Figure 2.8 Schematic illustration of depth profiling in LEIS

the structure studied is not a nicely layered structure. Therefore, only a near-surface (few atomic layers) depth profiling on supported catalysts is meaningful Brinen et al. [80] used the combination of quantitative XPS and LEIS to study depth profiling in HDS catalysts. Van Leerdam et al. [81] (see also section 7.2) studied the stabilisation of y-Al203 by La. The groups of Taglauer and Kn6zinger have presented a number of studies where the spreading and dispersion of catalytically active oxides on their supports is rationalized, see for a review e.g. [82]. Also the SMSI (strong metalsupport interaction) effect, well-known in catalysis, of Rh on a titania model support could be identified by comparison of the near-surface depth profiles as a function of the calcination temperature [83]. Besides the porosity and roughness which will hinder the profiling of catalysts Den Otter et al. [84] showed that for the study of small clusters extra care should be taken. From molecular dynamics calculations on a small Rh cluster (148 atoms) on a support, they show that one probing ion can lead to complete explosion of the cluster. They state that only a very small dose can be used to study the surface of these clusters and that depth profiling of these clusters is not possible, since the sputtered particles do not necessary originate from the surface. This has also implications for the SIMS studies of these kinds

134

of systems. In LEIS, however, the interaction time with the cluster is so small that the probing ion is already far away from the cluster when the cluster begins to deform. As stated above, in LEIS the problem can be overcome by using low doses (< 7 x 1012 ions/cm 2 for 2 keV Ne*). This demand can not be fulfilled by the conventional LEIS machines but the new generation, like the ERISS, can provide analysis already for ion doses well below this limit.

3. Quantification of surface composition 3.1 I n t r o d u c t i o n

When inert gas ions are used to probe the surface, the peaks in the LEIS spectra only result from binary collisions of the ions with atoms in the outermost atomic layer of the surface. The signal Si of ions scattered from atom i is then given by S i -- I ' I V

i "c

"R

"Pi

+ 9 ..__.doi dr2

(3.1)

where I

=

N i

=

c

=

R

_.

Pi +

=

do i / dQ =

the primary ion current the number density of the surface atoms of element i an instrumental factor taking into account the acceptance efficiency of the analyzer correction factor for rough surfaces (R = 1 for a flat surface) ion fraction of projectiles after scattering from atom i differential cross section for scattering by atom i.

angle and

The differential scattering cross section can be calculated using the Moliere potential for the ion - atom interaction [85]. The instrumental factor c can be obtained from measurement of a well-known surface. The roughness factor R generally ranges from 0.5 for a very rough surface to 1.0 for a flat surface (see below). The main problem in using LEIS for the determination of the number density of surface atoms i depends on the knowledge of the ion fraction.

3 . 2 Ion f r a c t i o n

Various models have been developed that describe the ion fraction in terms of neutralization and reionization processes. Although much progress is being made in the understanding of these processes, it is still not possible to calculate and predict the ion fraction accurately. If only neutralization occurs, the ion fraction is determined by the interaction time t and

135

a constant "a" that depends on the identity of the collisions partners. Since neutralization will occur during both the incoming and outgoing trajectory of the projectile, the ion fraction can be written as P/* = exp ( - a

9t)

for E i ~ Eth: = exp - v c ( 1 / v i + 1/vf)

(3.2a)

where Et~ is the threshold energy for reionization, vi and vf are the incoming and outgoing velocities and vc is a constant that is characteristic for a given ion-atom combination (the "characteristic velocity"). For metals eq. 3.2a is sometimes written in terms of the perpendicular velocities (when one assumes that the conduction electrons determine the neutralization). Angular dependent studies of Ne § ion scattering indicate, however, that often the localized d-electrons play an important role [86]. Experimentally it has been found [50,53] that for certain ion - atom combinations the ion may first be neutralized during the incoming trajectory and is then reionized during the close encounter. Reionization only occurs above a threshold energy Eth that is characteristic for each ion - atom combination. For energies well above Etb it is generally assumed that reionization is complete; the term 1/v i in eq. 3.2a is then reduced to zero: for E i )) Eth: P/" = exp ( - v /

vr)

(3.2b)

For intermediate energies neutralization along the ingoing trajectory will reduce the ion fraction somewhat. This phenomelogical description holds for most ion - atom combinations. The velocity dependence of the ion fraction can be verified by using eq. 3.1 to calculate P+ and then plot In P+ against the reciprocal of the incident velocity or the square root of the incident energy (since x/Ei ~ Vi ~ Vf ). For infinite velocities (no interaction time for neutralization), the plot should extrapolate to In P+ = 0, since the ion fraction then becomes 1. This extrapolation to zero should, of course, hold for any material and thus provides an additional possibility for checking the validity of eq. 3.2 and the reproducibility of the measurements. Brongersma and coworkers [87,88] showed for Cu and two types of carbon that, within the accuracy of their measurements, the plots did indeed extrapolate to the same value of lnP +, which will be zero (corresponding P+ = 1). This is illustrated in fig. 3.1. It is somewhat surprising that such consistency checks are not common practice in LEIS. Plots such as those in fig. 3.1 are also important to establish in which energy range eqs. 3.2a,b hold, since for energies just above the threshold for reionization or that for an additional neutralization process becomes important, a deviation will occur [61]. Erickson and Smith [89] and Rusch and Erickson [90] have identified a few elements where an exact energy resonance between an occupied level of an atom of this element in the target and the vacancy in the He ls level lead to oscillations of the ion fraction as a function of the kinetic energy of the He + ion. These oscillations are particu-

136

larly large for scattering from Pb and Bi. In such cases a reliable quantification is still possible by using another noble gas ion for the analysis. We, therefore, restrict the discussion to ion - atom combinations where no (strong) resonance is observed. Several methods are used to obtain absolute surface concentrations: calibration against standards calibration against other methods the DISC method TOF methods, which enable the measurement of all scattered particles, but including particles from greater depths. 10 ~

~i, l

10"1

C 0

u lL . C

o

1 0 "l

10. s

1 0 -4.

0.00

o% o s ~___"X grapl;17ir +__ '~' /

~

-

,

0.50 llv,+llv,

carbidic carbon ,~,

,

, 1.00

. . . . 1.50

[10 4 s i r e ]

Figure 3.1 The ion fractions for copper, carbidic and three types of graphitic carbon as a function of the sum of the reciprocal velocities of the incident and scattered ion

[87].

3.3 Q u a n t i f i c a t i o n

and the presence of matrix effects.

The applicability of a calibration against standards depends completely on the availability of reliable reference samples. The ion fraction for scattering by an atom i in the surface of the target should be the same as that in the reference sample ("no matrix effects"). In techniques such as Secondary Ion Mass Spectrometry (SIMS), where one studies the target atoms that are sputtered in an ionized state, the precise composition and structure greatly influence the ion fraction. The reference sample should then closely resemble the target of interest. Especially for reactive surfaces, this is not an easy task. In contrast to SIMS, it is nowadays accepted that matrix effects only play a minor role in LEIS, at least when the proper scattering conditions are chosen (see also the discussion of "the energy method" in section 5.2). If there are no matrix effects, eq. 3.1 can be written as

S.t = rl~ - 0.!

9R ,

(3.3)

137

(s)

pure

Nb

sud pure T s

4OO

~

+

1.$ ItoV 4 Ho 40

8OO

i_o

o

2oo

el

j-a Z lOO

Nb 0

o

....

' ....

6

' ....

lO

' ....

16

20

Oxygen ,lOne0[~,1

Figure 3.2 Peak intensities of Nb and Ta versus the oxygen peak intensity for the pure elements Nb and Ta [173]. where rh = calibrated sensitivity factor of atom i 0i = fraction of the surface covered by atoms i. For a target that consists only of two elements (i, j) it follows that S~ = r l ~ - R

-S~-rl~/rly,

(3.4)

since 0i+0)=1.

(3.5)

When for various surface compositions, S i is plotted against Sj, eq. 3.4 shows that this should give a straight line. Fig. 3.2 demonstrates that this is indeed the case for the adsorption of oxygen on tantalum and niobium. A similar behavior has been found for various binary and ternary systems [91,92,93]. Sparrow has determined elemental sensitivities for 1380 scattering of 2 keV 3He § by 29 elements (see ref.6 in [80]). Taglauer [94] presented elemental sensitivities for 1370 scattering by 17 elements of 4He§ for 500 eV, 1000 eV and 2000 eV, while Swartzfager [95] even determined the sensitivity factor for (adsorbed) xenon. In fig 3.3 the relative elemental sensitivity factors for 1360 scattering at 1 keV for 3He§ (a) and 4He§ (b) for 13 elements, as measured by Mikhailov et al. [51] are presented. The two values indicated for Si are for two different assumptions of the surface densities, a polycrystalline Si surface and a Si(100) surface (Si'). The absence of matrix effects remains a matter of debate, although in most papers it is assumed without proper justification. A simple way to verify the absence of matrix effects it to carry out the composition analysis at more than one energy. If the characteristic velocities are really the same for scattering by atom i in the sample and in the reference

138 101

) 10-' Pd

10 0

Si

J~ L_ II c

Cu Ni I

AI~ / /

Rhe 1 0 -2 Q Mo

tt~-7

M

0

Si* 9

10-I

1 0 "11

O

I 0 -3

'

'

'

'

'

'

10

20

30

40

50

60

a

,

10-4

70

80

101

I0-I oPd Rh

Pt

9

9

Ni C u

10 0

1 0 -2 L--

A;t"

Si.

m

Nt

o

-1

10-t

O 1 0 .3

Si*

10-2 0

. . . . . . . . . . . . . 10 20 30 40 50 60

b 70

10-4 80

Z Figure 3.3 Relative elemental sensitivity factors (O=136 ~ 1 keV for 3He + (a) and 4He§ (b). The dotted line indicates the differential cross section [51].

sample, the determined composition will be independent of energy. Jacobs et al. [19] established this energy independence for a number of elements in widely differing environments. For instance, when a pure AI target is used as reference sample, He + ion scattering through 1420 gives for primary ion energies between 1 and 3.5 keV, concentrations of A1 atoms that are independent of the energy in an Ago.sA10.2 alloy, in sapphire (etA1203), in y-A1203, and in oxidized NiA1 . The same was found to be true for Ni in NiO, NiA1, Ni0.sPt0.2 and in a y-A120 a catalyst having a 14 wt% loading of Ni when pure Ni is used as a reference. In such measurements it is, of course, an absolute necessity that the ion dose is so low that preferential sputtering can be neglected at all energies. For the insulators a proper charge compensation is required. Also, while low angles of incidence or scattering are very useful in studies of the structure of the surface (see section 4), such low angles

139

should be avoided in a composition analysis study. A surface atom may otherwise shield several atoms within the same surface plane, which seriously complicates the analysis. The critical angle for such shielding effects depends on the surface structure, the energy and the type of ion. For He+ ions of 1 keV a rule of thumb is that the angles with the surface should exceed 300. In the past it was generally believed that an Auger process involving two conduction electrons was responsible for the neutralization. It is difficult to understand, however, that this process would be quantitatively the same for an ideal metal as aluminum and a wide band gap insulator as alumina. This observation, as well as the dependence of the ion fraction on the atomic number of an element [51], suggests that the electrons that are somewhat more tightly bound are mainly responsible for the neutralization. As an exception to the general rule, the AI signal of NiAI relative to that of pure A1 does vary with energy. So in this case no quantification is possible by calibration. Why there is a matrix effect for the AI signal in NiA1 is not clear at present. For AI a close resonance with a level in He (see above) is not expected. It is also not clear why alloying with Ni influences the ion fraction, while a much more drastic chemical change, such as oxidation, does not. The behavior of AI does emphasize, however, the importance of checking the absence of matrix effects. Matrix effects have also been claimed for other systems such as K/V6OI3 (001) [34,115]. This could be related to the very low work function of this material. This phenomenon is very well-known for the deposition of different alkali metals on metal surfaces [96,97,98] For such surfaces a resonant transfer of conduction electrons is possible into an excited state of He. The interpretation of the results of Landuyt et al. [115] is, however, not completely straightforward. In their extensive energy dependent measurements they use the 250 - 1000 eV energy range to obtain their neutralization constants. However, in this energy range their plots of In P+ do not extrapolate to the same value for vanadium and oxygen (as they should for a consistent interpretation; see section 3.2). If they had used their data in the 1 - 5 keV range, then extrapolation to infinite energy would have given the same value (within experimental error) of In P+. Even if one takes into account that the threshold for reionization by oxygen is around 700 eV, so that a transition from eq. 3.2a to 3.2b is expected, it is still surprising that the low energy part does not extrapolate to In P+ = 0. The strong reduction in the Mg peak during oxidation that Peng and Barteau [99] observe, but do not discuss, may be due to the very low energy of their He + ions (200 eV) and their scattering conditions. An important exception to the rule that matrix effects do not occur when the proper scattering conditions (see section 2.5) are used, was recently identified by Van den Oetelaar et al. [87] for He + ion scattering from carbon. Energy dependent measurements showed that the ion fraction for scattering from graphitic carbon is much lower than that from carbidic carbon. For 2 keV 4He+ scattering the ratio of the ion fractions is estimated to be a factor of 230. The origin of this extreme neutralization by graphite could be qualitatively explained by a quasi-resonant neutralization process [87].

140 3.4 Influence of contamination

The extreme surface sensitivity of LEIS implies that surface contamination greatly influences the signals. This has often confused the interpretation of LEIS. Especially, the presence of hydrogen atoms poses a problem. Since hydrogen has a lower mass than the lightest inert gas ion (He), conservation of momentum forbids backscattering of He by H atoms (although H can be observed in forward scattering or as recoil , see [10]). This means that although neutralization and scattering by hydrogen atoms may reduce the signals from the underlying atoms in a backscattering experiment, it cannot be detected itself. The application of eq. 3.5 is then disturbed by an unknown amount of hydrogen. The complication of the quantification by the presence of hydrogen is well-known [25,100,101,102,103,104,105,106]. The presence of an intense sputter peak (see section 1) is generally a good indication for a high concentration of hydrogen on the surface [104]. Hydrogen is a special problem for surfaces such as metallic Pd that dissociate molecular hydrogen or hydrocarbons which are generally present as residual gas in a vacuum system. This may thus selectively lower the signals for such elements, although spill-over reactions may also lead to adsorption of hydrogen on neighboring atoms, with corresponding decrease of their signals. Another important source of hydrogen atoms is the adsorption of water, which will form hydroxyl groups on the surface. Most of the hydrogen (as hydride or hydroxyl) is generally effectively removed by a treatment in oxygen at elevated temperatures (e.g. a 15 min treatment at 250~ in oxygen of 20 mbar). Since hydrogen can be sputtered quite efficiently, a small amount of sputtering will remove most of the hydrogen before the sputtering of the substrate becomes appreciable. In fig. 3.4 an example is given of a palladium surface on which hydrogen has been adsorbed. The surface composition has been analyzed in the EARISS instrument (see section 2.2b). The sensitivity of this instrument is so high that a very low intensity ion beam can be used to study the surface (static LEIS). The Pd signal is monitored while the hydrogen is slowly removed by a second ion beam. Without sputtering, the hydrogen completely masks the presence of the underlying Pd! This example underlines the extreme surface sensitivity of LEIS. Regular LEIS equipment is by no means static, even when very low ion energies are being used. In such experiments a much higher ion dose than in the EARISS is required to optimize the scattering conditions and obtain a first spectrum. At that time, some hydrogen will have been removed and some of the metal will already be exposed to the ion beam. Apart from hydrogen, all elements of the periodic system can be detected by backscattering of He § ions. Elements such as lithium, beryllium, boron and carbon have been studied by LEIS, but their elemental sensitivities are very low. The strong matrix effect observed for carbon complicates the detection of this type of carbon even more. When Li, Be, B or C are present on the surface in significant quantities, quantification is difficult, although not impossible. It is quite possible that many of the matrix effects observed in the early years of LEIS were in fact due to changing amounts of contamination on the surface. If none of these very light elements are present in significant amounts, it should be possible to obtain an accuracy of the order of 1 % for the surface concentrations of the main elements.

141 10000

8000

,

-!

.

.

.

.

.

.

.

,

. unsputtered (~---03.6E14 ions/cm2 [3-. --!:18.9E14 Ions/cm2 ~- - -~ 2.7E15 Ions/cm2 #

.0 L.

6000

/

\

v

"o ~-.

4000

r

z

2000

43-B

_0 . 0 . 0 - 0 - 0 - 0 " 0 0 1800

,

,

1850

O0-O49" O-o

. . . . . 1900

:

%, o,,

"0"0-0.0. ^ B'IE~

.......

~,~ tr~.r 2'i I v

1950

Final energy

2000

(eV)

2050

2100

Figure 3.4 Pd-signal as a function of sputterdose, 4 keV Ne §

3.5 Calibration

against

other

methods.

In special cases it is possible to calibrate the LEIS signals by comparison with other techniques such as LEED, AES and XPS. In fact, the quantitativeness of LEIS and its monolayer sensitivity make it attractive to do the opposite: other techniques are often calibrated against LEIS. In studies of growth of an overlayer on a solid, LEIS is a sensitive tool to determine when a full overlayer is obtained. This provides an excellent calibration for other techniques such as AES and XPS, which have a greater information depth, if layer-by-layer (Frank-van der Merwe) growth prevails. Vurens et al. [107] used LEIS in this way to calibrate the sensitivity of AES for Na in the adsorption of sodium oxide on platinum. Bardi [108] has shown that when a LEIS signal is plotted against the signal from a technique such as XPS or AES, one can take advantage of the difference in probing depth to determine the prevailing growth mechanism.

3.6 Quantification by the DISC m e t h o d . When two isotopes of the same inert gas ion having the same incident energy are scattered from the same target, the signal for the lighter isotope is generally significantly larger. This is particularly true when comparing the signals for scattering of 3He* and 4He*. The larger velocity of the 3He § for a given incident energy and the lower energy loss during the collision lead to a shorter interaction time with the target atom and thus to less neutralization. Ackermans et al. [109] have shown that when eq. 3.2a or 3.2b holds,

142 one can determine the characteristic velocity v c by comparing the signals for scattering of two inert gas ion isotopes. Since both measurements can be carried out immediately after one another, or even simultaneous when a gas mixture is used in the ion source, the results pertain to the same target with the same roughness and composition. Especially for samples where no good reference targets are available for all elements present, this Dual Isotope for Surface Composition (DISC) analysis is quite helpful to determine the ion fractions P§ for all elements present at the surface. From this the surface composition can be calculated. The accuracy is, however, generally not as good as can be obtained by calibration against standards.

3.7 Surface roughness. The influence of roughness on scattered ion signals has been described already by Nelson [110]. For tilted surfaces, physical screening will reduce the intensity, while the increased effective density will increase the signal. Jacobs et al. [19] have shown that even for the very rough surfaces of catalysts, such as ct-Al203 (5.5 m2/g) and y-A1203 (269 m2/g), the predicted total effect is at most a reduction of the signal by a factor of two. In the past, much larger reductions have been observed. Margraf et al. [111] report, for instance, values exceeding a factor of five. One of the main problems of rough insulating surfaces is proper charge compensation. In experiments where a very homogeneous and effective charge compensation was obtained by spraying low-energy electrons from all sides, the reduction for various high surface area materials as alumina and silica was only a factor of two [19,58], in agreement with what is expected from simple macroscopical screening. Berning and Niiler [112] have also described methods to construct the energy spectra of ion scattering from rough surfaces by a summing over an array of tilted flat segments. Although the methods were applied to high-energy ion scattering they are also applicable for low-energies.

4. Surface structure of single crystals 4.1 Introduction Low-energy ion scattering has been applied extensively to determine the atomic structure of single crystal surfaces of metals and semiconductors, as well as that of adsorbed overlayers [113]. The combination of ion scattering techniques with Low-Energy Electron Diffraction (LEED) has proven to be especially powerful. LEED is used to establish the presence of long-range order and to determine the dimensions of the surface unit cell. By studying shadowing and blocking effects in ion scattering, one obtains information on the relative positions of nearest neighbors (short range order). Next to the development of the ion scattering techniques to study the surface structure on an atomic scale, a major effort has been put into theory to be able to simulate the ion

143 trajectories and elucidate the surface structure. These calculations have proven to be very succesful in describing the geometry on the surface, but also to quantify small changes in the atomic location due to surface relaxation see refs [113,59] and references therein. For insulators both techniques (LEIS and LEED) become more difficult. Low-intensity beams, charge compensation (for ion scattering) or experiments at higher temperatures (higher conductivity of the samples) are used to overcome these charging problems, as discussed in section 2.3a.

4.2 Local atomic structure

Taylor and Ellis [114] used LEIS and LEED to study the structure of a UO 2 (100) c(2x2) surface. Their 500 eV He + ion beam was perpendicular to the axis of a masked double-pass CMA (total scattering angle variable from 900- 132~ The observed cut-off angles in the ion scattering from U, together with the diffraction features, showed that the outermost layer consists of oxygen atoms arranged in distorted bridge-bond, zig-zag chains along the < 100> directions. Landuyt et al. [115] used LEIS and LEED for their studies of potassium-doped (97 ppm) V60~3. The potassium segregates strongly to the surface. From the fact that no blocking or shadowing involving K was observed, it was concluded that the K must be situated in or near the outermost VO layer. At the point of saturation of the surface by potassium, the composition of the outermost layer is KV203. A most probable site was obtained for the K in the unit cell. Tanaka et al. [116] were successful in obtaining a crystalline film of the high Tc superconductor YBa2Cu307_x (100) ( l x l ) on Mg (100) by laser ablation. A special kind of CMA, using a LEED screen in combination with 2 slits, was used for the angle and energy analysis of the scattered ions. From a comparison of a computer simulation model with the experimental signals of O, Cu and Y relative to that of Ba at 500 ~ it is derived that, at this temperature, chains of Cu-O stabilize the surface. It is also evidence for the fact that a clean surface of this superconductor can be obtained. At 600 ~ where the orthorhombic - tetragonal transition occurs, the signals for copper and oxygen (relative to Ba) decrease; the oxygen desorbs. In the computer simulation it is assumed that the ion fractions of the scattered He are the same for scattering by all elements. This seems unrealistic. However, since the conclusions depend mainly on the changes of the signal ratios as a function of temperature, this assumption will hardly effect their conclusions. A perfect demonstration of the possibilities of ion scattering for structure analysis of ceramic surfaces has been given by Souda et al. [6]. The segregation of Ca 2§ ions to the MgO (001) surface is studied using an energetic neutral beam (1 keV He ~ instead of a He + ion beam. The He ~ atoms of this energy are effectively ionized in a close encounter with Ca or Mg, thus enabling the detection of scattered He+ ions. The high ionization probability enabled the use of extremely low fluences (less than 108 atoms/cm2). When using neutral beams, charging can only result from the scattered ions. Since this fluence is again much smaller than that of the incoming particles, charging effects are minimized. This circumvents the necessity of using electron bombardment or heating. A drawback of the Neutral Beam incidence Ion Scattering Spectroscopy (NBISS) technique is that only those elements can be detected where ionization occurs during the collision. It is found, for instance, that oxygen is not observed (and thus He ~ is not ionized by oxygen) under

144 these circumstances. This seems to be in contrast to the discussion in section 2.4, where it was suggested that He is reionized by oxygen. The threshold for reionization can be derived from the low-energy onsets of the tails in the energy spectra of oxides where the the ions have just enough energy to be reionized. This threshold for the emission of ions is around 700 eV. Since the energy of the incident He ~ was only 1 keV, the energy of the scattered He will be below this threshold, thus explaining that oxygen is not observed.

anion cation

-[110]

1001

Figure 4.1 The surface structure of MgO(100)

When a MgO crystal that is doped with 210 wt. ppm Ca is heated, Ca segregates above 800 ~ The surface structure of MgO (100) is illustrated in fig. 4.1. In fig. 4.2a,b it is shown how the scattered (ion) signal depends on the azimuth and the angle of incidence. The detection angle is perpendicular to the surface. This facilitates the interpretation of the spectra (no neutralization by neighboring atoms on the way out). Such data give very direct information on the atomic structure of the surface. Especially at low incident angles ct (down to 10~ with respect to the surface) strong shadowing of the Mg is observed in the <100> and <110> azimuths by neighboring 02. ions and Mg 2§ (or Ca 2§ ions, respectively. The fact that the azimuthal dependence of the Ca peak (fig. 4.2b) is similar to that of Mg (fig. 4.2a), but only pronounced at very low angles of incidence, suggests that the Ca 2§ ions have replaced some of the Mg 2§ ions but are protruding from the plane of the Mg 2§ and 02. ions. From these and other measurements the position of the Ca 2§ ions (0.4 +_. 0.1 ,~ above the plane of the Mg 2. ions) and the 02. ions (within 0.1 ,~ of the plane of the Mg 2§ ions) could be determined. The very low signal for scattering from Mg at an incident angle of 50 with the surface suggests that the surface is quite perfect (mutual shadowing of all atoms thus preventing large angle scattering). In fig. 4.3 the scattered ion intensities for Mg and Ca are plotted as a function of the angle et for the <110> azimuth and a fixed scattering angle of 1600 (Neutral Impact Collision Ion Scattering Spectroscopy mode). The Mg signal starts to rise significantly at ot = 12~ indicating that the Mg atoms then come

145 out of the shadows of neighboring atoms (see fig. 4.3). Around 20 o a focussing peak is observed. For Ca, all these features are shifted to lower angles, supporting the fact that Ca protrudes from the surface. It is good to realize that in such experiments, where the beam is lowered to very low angles with the surface, the relative magnitudes of the diameter of the incident beam and the acceptance of the analyzer may affect the spectra somewhat [117]. For instance, when the beam is wide, the analyzer "sees" more surface atoms at

(a) Mg peak

(b) Ca peak He + H e ~ .....

~o

90 ~ 'tP~

20 ~ N

%

Y. 5

Y.

20 ~

20 ~

Y.

" "p a,f

L5~

. 15 ~

5

_ 10 o I

I

I

I

[ 11-01

[ 100]

[ 11 O]

[01 O]

Azimuth (~)

."o

I

I

I

I

[ 110]

[ 100]

[ 110]

[010]

10 ~ 5~

Azimuth (~)

Figure 4.2 Intensities of He + ions from (a) M g 2§ and (b) Ca 2§ ions at the Ca 2+ segregated MgO(100) surface as a function of the glancing angle R and the azimuthal angle q~ [6].

low angles.

4.3 S u r f a c e def ect s

Low-energy ion scattering can also be used as a sensitive probe for surface defects. The low, but non-zero, intensity (fig. 4.3) of the Mg peak at angles around 10~ indicates that the surface of the Ca doped MgO (100) surface is not completely perfect. Since this onset at low angles also occurs for the surface before Ca segregation, it is due to scattering from

146 surface defects. It is interesting to note that in a conventional LEED system no broadening of the peaks or appearance of extra spots was observed. This illustrates the high sensitivity of (N)ICISS for surface defects. Nakamatsu et al. [118] investigated the defects on the MgO (100) surface more

[ 110] Azimuth Mg peak

9 +

/

O0

9

9

008 O

0 9 0000 O O

000

9

9 9

9 9 0 9

000

9

9 0000

9 00

00

9

O O 9

=

9

99 ~)

9

I

Ca peak

I

I

20

I

40

I

60

ct (deg) Figure 4.3 Intensities of He* ions from Mg 2§ and Ca 2§ as a function of the glancing angle ~ of the incident He ~ beam, scattering angle is fixed at 1600 in the <110> azimuth [6].

closely with ICISS. Fig. 4.4 shows the Mg peak for the <010> azimuth as a function of the angle of incidence. Due to the larger distance between the neighbouring atoms in the <010> azimuth as compared with that in the <100> azimuth, the critical angle ctc is displaced to higher values and the focussing is not visible anymore (compare fig. 4.4 with fig. 4.3). Analogous to Souda et al. [6] it was found that annealing of the target at 1000 ~ led to sharp spots in (conventional) LEED while the ICISS intensity at low (x indicated that many defects were still present. Careful analysis of the ICISS spectra as a function of the annealing temperature enabled Nakamatsu et al. [118] to show that the isolated vacancies disappear already at 1280~ while the clustering vacancies start disappearing at 1310 ~ In general, vacancy pairs and larger vacancy clusters are more stable than isolated vacancies [119]. Annealing at 1 3 8 0 - 1410~ results in an almost perfect crystal surface. Above 1430 ~ the intensity at low tx starts increasing again, due to the formation of thermal etch pits. The concentration of surface defects strongly influences the conductivity of oxides. SnO 2, for instance, derives it conductivity from bulk and surface oxygen vacancies. Cox et al. [16,120]have used LEIS in combination with LEED, X-ray (XPS) and ultraviolet photoelectron spectroscopy (UPS) as well as conductivity measurements in a study of the

147

r~ o,,,~

A o,-,i

A

11~%1I I ...... iii I

.........

(bl 0

.................. i0

I

I

20

30

!

40

a (deg) Figure 4.4 ICISS Mg-intensity variation with the incident angle (ct) of He ions for a sputtered surface heated at 1370 ~ (a) and 1280 ~ (b), a was changed in the (010) plane of MgO(100) [ 118].

effect of temperature and oxygen treatments on the occurrence of oxygen defects at the SnO 2 ( l l 0 ) - l x l surface. A well-oxidized nearly perfect SnO 2 (110) surface could be prepared by oxidation at 700 K in 1.0 Torr of O z for 3 min, as verified by LEIS. In these experiments an incident He + ion beam with a primary energy of 1 keV was used at an incident angle ct of 400 along the <001> azimuth and a scattering angle of 130 ~ Although annealing in vacuum at 1000 K still produces a ( l x l ) LEED pattern, LEIS shows that the ratio of oxygen to tin atoms is reduced from 1 to 0.17 [120]! The effect of the vacuum anneal on the XPS and UPS data is much less pronounced because of the larger information depths of these techniques.

4.4 Site labeling SnO z has the futile (TiO2) structure. This means that in the [110] direction the crystal consists of charge-neutral units, each containing three atomic planes: O 2, 2 Sn 4+ + 2 0 2, O 2. An ideal stoichiometric (110) surface will be terminated by such a unit to ensure charge neutrality. The outermost plane consists of oxygen atoms in bridging positions. Cox et al. [16] used LEIS to show that the bridging oxygen can be removed selectively by heating for 3 min. in vacuum at 700 K. Stripping away all bridging oxygen increases the number of Sn atoms that should be visible with LEIS by a factor of two, in agreement with the observed doubling of the height of the tin peak. The selective desorption and

148 adsorption of the bridging oxygen enables site-selective isotopic labeling. The 160 and '80 isotopes are easily distinguished in LEIS because of their large relative mass difference. The fact that it was found that selective labeling of the bridging oxygen is possible at 700 K shows that the mobility of oxygen is still very low at this temperature. The mass and surface sensitivity of LEIS thus provides an interesting possibility for studying the mobility and exchange of surface and bulk oxygen in metal oxides and could thus be important in studies of fuel cells and gas sensors, where these processes play a decisive role in their operation.

5. Surface structure of non-single crystals. 5.1 Introduction

In the previous section it has been discussed how LEIS can provide very detailed information on the atomic structure of surfaces. For amorphous and polycrystalline materials the information is averaged over all azimuthal angles, and for powders also over many angles of incidence. This makes techniques such as ICISS useless for these materials. Various authors have explored other possibilities to obtain structural information on such compounds.

5.2 The energy method

One method that has received attention (e.g. [21,22,55,57,121]) is based on the dependence of signal ratios on the incident energy. It was found that the energy dependence of the ratio C/A of the signal for scattering from a cation to that of an anion is very characteristic for a compound. McCune ([21,121]) made a careful study of these ratios for many compounds in the energy range of 150 - 5000 eV and was able to classify the oxides into two groups. One group, containing compounds such as SiO 2, A1203, Nb205 and Ta205 in which the atomic structure will be such that strong screening of the cation by oxygen is expected, showed a strong increase in C/A with energy. The other group, containing compounds such as MgO, CaCO 3 and ZnO where no or little screening is expected, showed no clear increase or even a decrease in C/A with energy. At that time the interpretation seemed straightforward. At higher energies, scattering cross sections will be lower while neutralization will also be less. Especially for well-shielded cations it seems, therefore, reasonable that the C/A ratio increases with energy. Van Leerdam and Brongersma did show [61,122], however, that the origin of the energy dependence of C/A is generally an intrinsic property of the elements involved and not related to shielding. When taking the ratio of the silicon signal from pure silicon (so Si cannot be screened) and the oxygen signal of B203, the same energy dependence of C/A was obtained as for SiO 2 (fig. 5.1, [61]). Also, for m1203 and ZnO it was found that the characteristic energy dependence of C/A is the same for the compounds as for the constituting elements. Mashhoff et al. [123]) confirmed the absence of shielding effects

149 on the energy dependence of the C/A ratio in TiO2.

~!

I

I

t i

I

3-

2

1

x Si 02 9 Si §

0

1 1

1 2

i 3

I t.

5

Ei{keVl

Figure 5.1 Energy dependencies of the Si/O ratio for SiO 2 (x) and the reference compounds (o) [61].

The explanation of the effect is related to the electronic properties of He and the cation involved. When a He + ion approaches a solid, the ion will be neutralized by processes such as (quasi) resonant and Auger neutralization. It has been found by Souda and Aono [50] and by Thomas et al. [53] that for a number of target atoms the He atoms may be reionized during the collision. The probability for reionization increases with energy (deeper penetration of the He into the target atom leads to stronger overlap of the relevant orbitals and thus, in general, to stronger Pauli excitation/ionization). Since only ions are detected in LEIS, this increases the signal. Elements such as Si, A1, and Ta have been found to reionize energetic He atoms, while Zn does not do this. Quantum mechanical calculations using the Anderson-Newns Hamiltonian could quantitatively explain the energy dependence of the signal for He scattering from silicon [124]. This implies that the correlation found by McCune and others is not related to the surface structure and thus cannot be used for the analysis of polycrystalline surfaces. The structure of the oxides is too open to shield the cations in the top layer.

5.3 Spinels and the importance of the information depth In low-energy ion scattering of powders, shadowing and blocking effects are averaged over many crystal orientations (various surface planes, all azimuths and many angles of incidence). At first sight one might, therefore, expect that for a homogeneous powder a LEIS spectrum would just give the bulk stoichiometry of this powder. This is, however, certainly not the case. Although LEIS is not able to detect the shielding by oxygen of

150 cations within the toplayer (see above), the neutralization, shadowing and blocking are very effective in shielding the deeper layers. This property can be used to obtain very valuable information on the surface structure of these powders and thus have a direct bearing on the understanding of sintering and chemical reactions at the surface. The importance of shielding in oxides is shown in fig.5.2 where the spectra for He + ion scattering from a- and ),- AI203 are compared [125]. The AI signal of the ),modification is only 75 + 5 % of that of the a-modification. Similar differences between the a- and ),-phase have been previously observed for Fe203 [126]. To compensate for differences in grain size and pore structure (a- and ),-alumina have specific areas of 5.5 and 280 g/m 2 resp.) the spectra have been normalized on the oxygen peaks (see also section 3.7). The same oxygen signals are expected since Q-alumina has the corundum structure with a hexagonal close packing of the oxygen ions, while )'alumina has a defected spinel structure with an almost perfect cubic close packing of the oxygen ions. The number densities of the oxygen ions in the topmost layer will thus be the same. For clarity the spectra in fig.5.2 have been shifted somewhat in energy. !

I

A1 a-A1203 O

600

_A120 3

1200

1800

2400

Ef (eV) Figure 5.2 LEIS-spectra of the surfaces of a-Al203 and ),-A1203. (3 keV 4He+) [125]. For clarity the ),-alumina spectrum is shifted somewhat in energy.

For the ),-alumina the AI and O intensities initially increase rapidly with bombardment time, after which they develop more gradually. If the powder is first heated at 400 ~ the initial AI intensity increases by only 35%. With higher pretreatment temperatures the initial AI and O signals are higher, but the final values are the same. It is generally known that hydroxyl groups are present at the surface of ),-alumina. It is, therefore, assumed that. the initial increase of the signals is due to the removal of these groups. Increasing pretreatment temperatures will lead to a higher degree of dehydroxylation. Moreover, as a result of the thermal treatment mainly bridging OH groups will remain [127], which hardly shield the underlying A1 atoms. After prolonged ion bombardment the AI signal slowly increases to the value of the a-

151 alumina, while the oxygen signals of both modifications decrease due to preferential sputtering. In the mid-seventies, Shelef et al. [100,128] used LEIS to investigate the relation between surface composition and chemical reactivity of normal spinels. These compounds have the general formula:

A 'et'(II)B2t(IIl)04

(5.1)

in which the formal valencies (roman numerals) and cation coordinations are indicated. Although the sensitivity of LEIS equipment in those days was not yet very high and although the authors stressed that their experiments were still of preliminary nature, their results suggest that tetrahedrally coordinated cations may not be accessible to LEIS. It is found that in the first spectrum of a fresh target (little damage by the ion beam) no Co or Zn are detectable in CoA1204 and ZnAI204, resp. [100]. The cations in tetrahedral surface sites are believed to be less stable and will, therefore, move to sites below the surface and thus be shielded more than the cations in octahedral sites. Another explanation is that only certain crystallographic planes may be preferentially exposed in spinels. Recently, the findings of Shelef et al. [100,128] have been confirmed in LEIS studies of various catalytically active spinels [129]. As illustrated in fig. 5.3 [130], the Zn signal in ZnAl204 is less than 1 % of that for ZnO; see also [131]. When comparing different Co-containing spinels it was found that for CoAIzO4,where almost all Co-cations are present as Co 2§ in the tetrahedrally coordinated sites, the Co signal is only 13 % of that for ZnCo204, where almost all Co-cations are present as Co 3§ in the octahedrally coordinated sites. The Co signals for Co304, where Co is present in both the tetrahedral in octahedral sites, show the same amount of cobalt on the surface as the ZnCo204. The surface composition as determined by LEIS was found to correlate directly with the catalytic activity [129]. The results strongly support the idea that in II-III spinels only the octahedral sites are occupied in the surface. For a series of Mn-containing spinels it was found that due to an oxidative transfer much more Mn was found on the surface [129]. LEIS showed that this model also applies to substituted ferrites (ZnxMgvxFe204) [132], which are inverse spinels. Inverse spinels have the same structure as normal spinels only their cation distribution is different: Bt~tr(III)A~176 in accordance with eq. 5.1. While Fe304 (inverse spinel) showed the same amount of iron on the surface as ZnFe204 (normal spinel), MgFe204 exhibited only half of the iron in comparison to the normal spinel, which is in complete agreement with the model where only octahedral sites are exposed. Furthermore, in ZnFe204 no Zn was detected, while in MgFe204, where the doubly charged Mg-cation is in the octahedrally coordinated interstices, Mg was clearly visible. In surface analysis techniques such as XPS, where the information depth is much larger, the cations in both octahedral and tetrahedral sites are detected. The tetrahedral and octahedral site occupancies are given for some spinels in ref. [133]. Since the pioneering work of Shelef, many other groups have combined LEIS with catalytic activity measurements. In particular the groups of Hercules and Houalla [134], Taglauer and Kn6zinger [25], Hoflund [135], Bertrand and Delmon [136], Bonnelle [137] and of Brongersma [24] have made important contributions; see also the reviews by Horrell and Cocke [23], Brongersma and Jacobs [130] and of

152

Aul 3000 (.I

i I

I I I I I I

e m

u~ 2000 c c

,~

,,,,J

looo

..

f

It

I

..-,,"

i_l I

..~

I

I

1000 1500 2000 Final energy (eV)

2500

Figure 5.3 LEIS-spectra of 3 keV He § scattering from ZnA1204 (solid line) and ZnO (dashed line) [130].

Taglauer and Kn6zinger [82]. In general the LEIS signals were found to correlate very nicely with catalytic activity, since both refer to the top atomic layer. When the necessary experimental precautions are taken, this correlation is even quantitative [130]. This also suggests that there is no change in the surface composition due to chemically induced segregation during the catalytic reaction. One should realize that such a correlation between measured surface composition and catalytic activity is unique to LEIS.

5.4 Signal as function of loading Another method that has found widespread application in studies of catalysts is to follow the ion scattering signals as a function of the loading of the catalyst. At low loadings (well below monolayer coverage) of the support it is not unreasonable that the adsorbed species ("the loading") do not cluster and are thus visible for LEIS. If this is the case, one expects a linear increase in the LEIS signal of the adsorbed species as a function of loading and a linear decrease of that of the support. Because of various experimental reasons (see section 2.3) signal ratios are generally used instead of the absolute signals. This may lead to erroneous interpretations. Zingg et al. [138] observed,for instance, for catalysts of a )'-AI203 support a clear increase in the Mo/A1 signal ratio at a loading of about 7 wt.%, when the molybdena loading was increased. It was assumed that at low loadings the Mo atoms occupy mainly tetrahedral sites and are thus less visible. Above 7 wt.% loading the additional Mo atoms were assumed to be in octahedral sites. This coverage-dependent site occupation was not observed by Kasztelan et al. [137]. Van Leerdam et al. [65] observed the same discontinuity in the Mo/AI signals as Zingg et al. [138]. The change occurred at somewhat higher loadings since their specific area of the ),-alumina was 270 ma/g instead of the 196 m2/g used by Zingg et al. [138]. However,

153 I

f

i

"i

i

i

MoO3/)'-Al203 "

2 - A1

-

_

\

~

1-

~

i

0

I. . . . .

10

I

1

20

I

_

I

30

M o O 3 - content (wt %) Figure 5.4 The LEIS-intensity (S) of Mo (squares) and A1 (circles) as a function of MoO 3 content [65].

they were able to determine both the Mo and AI signals as a function of surface loading (see fig. 5.4). The Mo signal is found to be strictly linear with M o O 3 coverage up to 20 wt.% (above this loading 3-dimensional structures of MoO 3 are formed). This suggests that up to this loading the Mo is always accessible to LEIS and thus not hidden in a tetrahedral site. The discontinuity in the Mo/AI signals is due to an abrupt change in the AI signal. At low loadings the AI is apparently not shielded by the Mo. At first sight this may seem surprising. The structure of a cation-deficient spinel such as 7-A1203 is, however, rather open. It is thus possible to find sites for the Mo where it is on top but does not shield any A1 atoms. At a loading of 10 wt.% these sites are full. When the loading is increased any further, all Mo atoms shift to new locations. Van Leerdam et al. [65] have indicated possible sites for the Mo atoms at low and high loadings. Both the constant AI signal at low loadings and the abrupt change around 10 wt.% are difficult to comprehend for a powder that exposes many different crystallographic planes. For this reason and others it has been suggested [65] that powders that have been annealed at high temperature for long periods may have reached a (kind of) thermodynamic equilibrium at which only one or two of their most stable surfaces are exposed. Whether such an equilibrium is actually reached may also depend on the structure of the starting material, surface impurities, etc. If the hypothesis of the existence of only one or at most a very limited number of crystallographic surface planes is true, it would enable the use of LEIS for atom location on powders. The research on the spinels (section 5.3) supports this hypothesis. In LEIS studies in the past, the support has been regarded as an inert base for the active compound, notwithstanding the fact that it is well-known in catalysis that there are strong metal-support interactions that completely change the chemistry. The LEIS signals were normalized to the A1 signal of the support. Since it was expected that if compounds like MoO3 were adsorbed on top they would always shield the underlying A1 atoms, one even

154 corrected for this shielding [134]. The results of van Leerdam et al. [65], however, clearly indicate the importance of measuring the signals themselves rather than signal ratios. The use of absolute signals also proved to be important in establishing the mechanism of growth of Ni on Al203 in Atomic Layer Epitaxy (ALE) [130,139]. It is presently believed that abrupt changes in the signal of the support as a function of loading are the rule rather than the exception (see also [140]). The changes in the support signal can be caused by rearrangements of the other atoms but also by phase transitions and other types of restructuring of the atoms in the support. It is quite possible that various results of LEIS experiments on bimetallic catalysts, where abrupt changes in the signal ratios of loading/support have been observed [141,142,143,144], are due to changes in the signal of the support.

6 Applications of LEIS to surface segregation 6.1 Introduction Soon after the development of LEIS, the technique was applied to the study of various aspects of oxidic materials. The sensitivity of LEIS for the outermost atomic layer was an important reason, although the ion beams were generally so intense that sputtering and damage largely nullified this feature. Another reason for using LEIS was the ease to cope with surface charging, thus making studies of insulators relatively easy (see experimental section). In the present section a survey is given of LEIS applications to oxidic materials. The work is classified according to their physical aspects.

6.2 Surface segregation. The composition of a surface is generally very different from that of the bulk. Apart from processes such as evaporation and adsorption from the surrounding atmosphere, the surface composition is often affected by surface segregation, i.e. enrichment or depletion at the surface of certain atoms or compounds from the bulk. The driving force is the establishment of thermodynamic equilibrium. Especially at higher temperatures, where diffusion rates become appreciable, the surface composition is generally radically different from that in the bulk. In many cases an impurity, of which the concentration in the bulk seems fully negligible, segregates to the surface and completely determines the surface properties of the material. This can strongly influence processes such as sintering. Since segregation is dependent on the surface plane, it will also stabilize surface planes differently. This will affect the morphology of small grains. The monolayer sensitivity of LEIS makes it an ideal tool to analyze surface segregation. Various applications are known, such as the use of LEIS to study the segregation of Pb impurities in garnets [145]. Fig.6.1 shows an example [102,146] of a very pure ZnO single crystal. A LEIS

155

(T~T)

analysis shows that after sputter cleaning the main constituents of the surface are indeed zinc and oxygen. At 100~ however, the surface becomes already strongly enriched in sodium, while Na is the main constituent at 550~ In this case, the bulk concentration of the Na impurity was only 8 ppm. Since Na is most stable at the oxygen terminated ('11"]') surface, the segregation is here particularly strong. However, Na was even found to dominate the Zn terminated (111) surface of ZnO. Creemers e t al. [147] confirmed these results and found for annealing temperatures of 250 - 400~ that even for a sub-ppm bulk concentration of Na this element dominates the (~1"]') surface of ZnO. This example also illustrates the importance of using complementary analytic techniques. The same sample had been studied at similar temperatures by Auger electron spectroscopy. Due to the proximity of the Na and Zn Auger peaks, however, no Na segregation had been inferred from the data, although an unusually low workfunction had been observed. Slightly Li doped ZnO did not show any segregation of impurities. The fact that Li is not observed may well be due to its very low sensitivity in LEIS. The absence of segregation of other alkali such as Na and K may be due to the mixed alkali effect which reduces the mobility of a given alkali ion by the presence of another alkali. The effect is used to tailor the refractive index profile in germanosilicate optical fibers [148]. In general, however, alkali segregation is commonly observed for oxides. In FischerTrops catalysts, where potassium is a well-known promoter, a strong enrichment of K in the outermost atomic layer has been observed with LEIS for Fe/Mn oxides containing potassium [149]. The thermodynamic theory for surface enrichment by segregation is reasonably well developed for alloys. For a regular solution the surface composition can be written as S

xi

S

_ B

xi

xj

. exp(-AH/kT)

(6.1)

B

xj

where x s and X B a r e the equilibrium mole fractions in the surface and bulk. AH is the heat of segregation. It can be split up in various contributions, such as that related to the lower coordination of surface atoms, heat of mixing and the enthalpy due to size mismatch of the atoms (strain). Other factors, such as surface reconstruction, can make significant contributions, however. In this model excess entropy effects, such as differences in vibrational energy of surface and bulk atoms, are neglected. The lower coordination of the surface atoms is generally the leading term. Since the lower coordination and strain relaxation are only significantly different in the outermost atomic layer of the surface, segregation is mainly restricted to this layer. Sometimes the effects extend over a few atomic layers. In those cases the enrichment may slowly decrease with increasing depth or enrichment and depletion may alternate. The very shallow depth over which surface segregation takes place makes LEIS ideally suited as analytic tool. In 1975 and 1986 this proved to be important to settle [150,151] a long standing controversy between physicists and chemists about the existence and sign of surface segregation in Cu-Ni alloys. In agreement with the observed chemical behavior, a very strong Cu segregation was found in the outermost atomic layer. Since segregation was

156 already almost absent in the 2nd layer, it could only be detected with a very surface sensitive technique. The Cu segregation was simultaneously confirmed by Helms [152] by careful analysis of his Auger data and by using the short mean free path of the lowenergy Auger electrons of Cu and Ni. LEIS has the advantage, however, that the contribution from the 2nd and deeper layers can be fully neglected, so no model for these contributions is needed. Wynblatt and McCune ([36,153,154] and [155] for a review) have extended and applied the regular solution model to metal oxides. One of the complicating factors in ionically bonded solids is the possibility that space charge layers are formed in the vicinity of surfaces [156]. This will influence the segregation of aliovalent solutes. For multivalent ions, segregation may depend on the valency of the ion and thus on the oxygen partial pressure. One should realize, of course, that accumulation of solutes in the sub-surface Debye layer cannot be detected by LEIS. Another complication is that when compounds are formed it is not always clear which units should be considered as solute and as solvent. For mixed Bi203-MoO 3 oxides, for instance, it was found [126] that strong Bi segregation takes place once the bulk concentration has an excess of Bi, while Mo segregation occurs when Bi is the minority. In this case, the stoichiometric oxide should thus be considered as the solvent, while the excess of Bi or Mo is the solute. The rule that a large negative heat of formation prevents segregation, thus only holds for the stoichiometric compound. McCune et al. [153] have tested the applicability of eq. 6.1 to oxides. Both Auger Electron spectroscopy (AES) and LEIS were used in-situ at the temperature of interest. For a CaO dopant in MgO, where Ca segregates strongly to MgO (100), the segregation was found to be reversible for the 1 1 0 0 - 1450 ~ temperature range. This indicates that thermodynamic equilibrium has indeed be reached. A plot of the logarithm of the ratio of the Ca/Mg surface concentrations versus 1/T did give, in accordance with eq. 6.1, a straight line. From the slope an enthalpy for segregation (-57 kJ/mol) was derived. This value was similar to that obtained from Auger electron spectroscopy (AES) data. It is quite seldom in the literature (but very valuable) that the reversibility of surface segregation is checked. For alloys it is well known that eq. 6.1 is followed closely by experimental results. Although differences in e.g. the vibrational amplitudes of surface as compared to bulk atoms may introduce an entropy contribution to the segregation, these effects are generally very minor. It is likely that this is also true for oxides. This would mean that when extrapolating the plot to infinite temperatures, one should obtain the concentration ratio of Ca and Mg in the bulk. This was indeed the case for the Auger data of McCune et al. [153], but not for their LEIS data. This would suggest that there was probably an error in the calibration of the concentrations of the LEIS results. It is believed that such an extrapolation to infinite temperatures could be very helpful as a check in future studies. It is of great importance to ensure that thermodynamic equilibrium has actually been reached. At low temperatures diffusion is too slow to establish equilibrium. Sputtering by the ion beam or the sample preparation will determine the surface composition. Also, ion bombardment will create defects and can thus enhance the diffusion near the surface (see section 2.8). The depth over which the enhancement occurs depends on the penetration depth of the ions (order of 10 nm in LEIS). This may lead to a local equilibrium of the outermost atomic layer with the nearby subsurface layers. A steady state can then be observed in the surface concentrations. It does not imply, however, that a full equilibrium

157 with the bulk has been reached. At high temperatures decomposition and evaporation may also cloud the issue. In studies Zn

t 1000eV He+ -->ZnO[lll][ / 0= 142~ ! 25~ after sputteringll "~ t ..... annealedat:550~ il Na

y.=

i[ Sputtered t[j ions il II /I /I

,!

N v

L \ I

I

I

I I I I

I I II II I I

I I%

.T I'~

,/

L ',

0

I I I I

II II

"..

-" a I

I

I

"r l

I

I

I

I

'

500 1000 Scattered Ion Energy (eV)

Figure 6.1 LEIS results from ZnO at room temperature and after annealing [102].

of MgO segregation in alumina Baik et al. [157] coped with the evaporation problem during equilibration by keeping two surfaces of the same material in close contact. In addition, it is well known that other impurities, having lower diffusion constants or concentrations, may become dominant at the surface at the higher temperatures. For a Ni doped MgO crystal containing a Ca impurity (870 wt. ppm Ni; 11 wt. ppm Ca) it was found that annealing at 1000 ~ leads to an initial enrichment in Ni (because of its larger concentration) but it is then replaced by calcium which has the greater driving force for segregation in the MgO lattice [153]. A typical example has been described by McCune [153] where a polycrystalline compact of AIzO3 containing 23, 83, 160 and 270 wt. ppm of Ca, Fe, Y and Mg resp. was studied by LEIS and AES. Calcium segregation was observed at 1600 ~ using AES. Due to evaporation the amount of Ca decreased in time and was replaced by yttrium. In fig.6.2 [153] the LEIS spectra are given for the doped specimen and compared with that of sapphire (0001). In LEIS, which has a higher sensitivity than AES, enrichment in Mg, Ca, Fe and Y was observed. After prolonged annealing at 1400 ~ the presence of AI at the surface is still only just detectable, illustrating again the high surface enrichment of trace elements that is generally obtained. In

158

0

4He+ 250 eV

Mg

Ca Fe Y

ra~ o~,,~

t~ ra~

4 Days 1400 ~ O

*w,.t

Mg A1

r,r

2 Days 1300 ~ Ar sputter cleaned

A I 2 0 3 (0001) (Ca) <40ppm wt

I

I

125

250

Scattered ion energy (eV)

Figure 6.2 LEIS-spectra of annealed and sputtered Y doped A1203 and comparisons to sapphire (0001) at primary beam energy of 250 eV [153].

159 these studies the analysis was carded out with He ions having a primary energy of only 250 eV. Another segregation study at elevated temperatures has been performed by De Gryse, Landuyt, Vandenbroucke and Vennik [34,115]. Despite the low bulk concentration of potassium (< 100 ppm) in their V~O~3 (001) sample, K segregates until a coverage of 1 potassium atom per unit cell is obtained. Kelso and Pantano have extended the surface analysis at elevated temperatures to its extremum, i.e. they have studied the melted surfaces of alkali-trisilicate glasses (Na20. 3 SiO2 and I(20. 3 SiO2 ) in the 350 - 1000~ interval, after the samples had been annealed in-situ at 1000~ [158]. The purpose of their study is to observe changes in the surface composition and structure during cooling down. In fig. 6.3 the LEIS signals for the alkalito-oxygen signal ratios are given for these glasses. The alkali surface concentration is

9 O 0

iO 2)

,i,,,a

~ 4

y ......... 9 ..............

..r . . . . . . . .

oO

--

10 ~o

----......... ' ~ 1 1 0

2)

r~ 2

0

9 350

'

a

550

I

i

750

I

i

950

0.5 Z

0

Temperature (~

Figure 6.3 Comparison of LEIS of alkali-trisilicate glass surfaces fractured in vacuum and heated in-situ [ 158, 275].

subject to reduction by evaporation and replenishment by bulk diffusion. Above the glass transition temperature the bulk diffusion is rapid such that a steady state alkali concentration is reached. Below the glass temperature (460 ~ for the sodium- and somewhat higher for the potassium-glass) the alkali concentration is depleted because replenishment from the bulk is too slow compared to evaporation and (probably at least as important) by sputtering during the analysis. Assuming reasonable bulk diffusion coefficients, it is unlikely that the steady state they observe already after only 8 hours annealing at 350 ~ is that for full thermodynamic equilibrium with the bulk. Also, the observed

160 alkali segregation for vacuum fractured glass will only be in equilibrium with the near sub-surface. The slow decrease of the Na concentration with temperature above 700 ~ seems a confirmation of thermodynamic equilibrium, in accordance with eq. 6.1.

6.3 Surface segregation and spinels In some studies, the formation of a spinel surface phase has been suggested as the driving force for segregation. Baik et al. [157] used Auger electron spectroscopy to study the segregation of an magnesia impurity to the basal (0001) plane of cx-A1203. Theoretical calculations confirm this although they are complicated by the fact that magnesium and aluminum are aliovalent. Mackrodt and Tasker [159], therefore, took advantage of the strong attraction between the magnesium impurity and the compensating vacancy by assuming a cluster of two magnesium ions and an anion vacancy. A sharp minimum in the total energy was found for the spinel composition (Mg~204). From neutron differential diffraction it is known [160] that MgA1204 is a normal spinel. The Mg 2§ ions thus occupy tetrahedral sites which, according to LEIS investigations (see section 5.3), are generally not present in the topmost atomic layer of the surface. Since Auger has been used by Baik et al. (which has an information depth of several atomic layers) the Mg will still be detected. It is surprising, however, that while the lower coordination of a surface atom is generally the driving force of surface segregation, the segregating Mg would be below the surface. Although unusual, such situations are known from alloy segregation where segregation of the solute atoms to the near surface area leads to stress or strain release while the surface energy of these atoms prevents them from segregating to the outermost atomic layer. An example is the segregation of Ni impurities to the subsurface layer in Ag [161]. The surface energy of Ag (1.2 J/m 2) is much lower than that of Ni (2.5 j/m2). Another possibility is that if the spinel layer is at most only a few atomic layers thick, the nearby presence of the underlying (0001) plane of a-alumina may influence the orientation of the spinel surface and thus stabilize a spinel surface that does contain tetrahedral sites in the toplayer. A similar situation seems to occur for y-alumina grown on top of NiAI (see section 7.1 and [19]). Poisoning of Co304 oxidation catalysts by segregation of alumina and magnesia has been reported by Wheeler and Bettman [162]. The segregation of alumina may be related to the formation of a CoA1204 surface spinel, since in this spinel the Co ions occupy the subsurface tetrahedral sites (see section 5.3). Although MgCo204 is an inverted spinel {Co(MgCo)O4}, this only partially explains the inaccessibility of Co since half of the Co ions would still occupy octahedral sites (see section 5.3). The formation of a surface spinel that "resembles" CuAI20 4 has been discussed by Strohmeier et al. [163] in studies of Cu/ml203 catalysts.

6.4 Segregation during oxidation The anneal atmosphere can strongly effect surface segregation of ambivalent cations. McHargue et al. [164] observed, in the case of Fe implanted in AI203 a profound redistribution of Fe towards the surface during annealing in an oxidizing atmosphere. The

161

Fe ~ v,,,4

0 0

Na Z

t,,d,, 250

400

550

I +Zr 700

850

1000

Kinetic energy (eV) Figure 6.4 LEIS spectra for the Fe implanted yttria-stabilized zirconia after oxidation at 400 ~ [165].

effect was not observed in a reducing atmosphere. A good demonstration of this effect is also given by van Hassel and Burggraaf [165]. They studied the use of mixed oxides on top of yttria-stabilized zirconia (YSZ) for use in oxygen sensors, oxygen pumps and fuel cells. For this purpose YSZ has been doped with either Fe or Ti and has been annealed in air. By LEIS and XPS it was found that the surface enrichment in Fe and Ti increased strongly upon oxidation in air at 400 ~ After such a treatment the outermost surface consisted mainly of Fe203 and TiO 2. In fig. 6.4 the LEIS result is given for the Fe implanted YSZ after oxidation. Some Na contamination is observed in addition to the Fe. A similar effect has been found by Geiger et al. [41]. They present an XPS study of the diffusion of palladium through SnO2, which is used as base material for chemical sensors. Besides small amounts of Na and CI, the main impurity is Fe. Using LEIS they show that annealing the SnO2 at 870 K in oxygen (p 02 >= 500 Pa ) results in Fe segregation that fully covers the SnO 2 (see fig. 6.5). It is also shown that after heating the sample in UHV for 10 min. at 970 K the Fe has dissolved again in the bulk. The Na and C1 contaminations that are clearly visible in fig. 6.5 could not be detected by XPS, due to its limited detection sensitivity.

162

LEIS

1 keV, He + 0.2 l.tA

Sn

O (a)

~

(c) v

K ~

\

r~

(b)

Fe

~

.Na . . . .

300

|

400

I

500

600

700

800

900

1000

Kinetic Energy (eV) Figure 6.5 LEIS-spectra (lkeV He+) of: (a) clean SnO2, (b) SnO 2 with high Fe contamination exposed to oxygen and (c) as (b) but treated in UHV (T=970 K, 10 min.) [41].

6.5. Influence of ion bombardment on surface segregation.

Ion bombardment of glasses is known to lead to alkali-metal diffusion. The bombardment will create an electric field in the glass surface over the penetration depth of the ions. This may lead to electromigration. Already in the early seventies McCaughan et al. [166,167] showed that ion bombardment mobilizes Na+ ions in silica films and causes drift of the sodium to the surface. Using a radioactive tracer technique Kushner et al. [167] could still detect the effect, even for bombardment with 500 eV Ar § with a dose of 10 ~2 ions/cm ~. More recently, Mazzoldi et al. (see e.g. [46,168]) have made extensive studies of the alkali depletion under ion bombardment, see fig 6.6. At low energies, where the stopping of the bombarding ions is mainly nuclear, an alkali depleted layer is formed which has a thickness of about twice the mean penetration depth of the ions (which is well known from experiments and trajectory simulations). The depletion is explained by radiationenhanced diffusion and segregation of the alkali elements to the surface, followed by preferential sputtering at the surface [46]. Some heating during the bombardment increases the mobility of the alkali ions, so the depletion is faster [169]. After formation of such a depletion layer, it is possible to restore the structure of the glass by annealing at a temperature that is still low enough to prevent thermal diffusion of the alkali metal. When the sample is at still higher temperatures, the return of the alkali to the surface can be monitored with LEIS. The monitoring is done with an ion dose that is much lower than the one used for creating the depletion. Since the thickness of the layer over which diffusion takes place is orders of magnitude smaller than in conventional set-

163

50 k e V - Ar +

o

10 _

o

o. i~,~"(i) ooo. . . . . .// m"

....... ~ " ..........

..........

o / /

-

8

-

i

/

9" "

/ /

##

."

/

Ir

41e"

9

..,"

.."

~.i

=,' ~

t

-o/ i x" , b N o.5 - ///,/"-" a v -'i i s t ..~ - / ~" ...~ '

#

, '~

~"~176

li~x").~

~ ~ "

0

Z 0

,,,, 100

9

9 ,

,~ ummplantedand

~x ,I

.s"

!

i

,,+,,.2 10 15 I~,,, 9 101~ z A 2.1016/cm 2 9 4.1016/cm 2 i 5.1016/cm2 X" !O1,7/cm2

o

260 Depth (nm)

. I-

...

'300

Figure 6.6 Experimental and theoretical sodium profiles after 50 keV-Ar irradiation at different doses illustrating alkali depletion [168].

ups for such measurements, it has proven possible to determine extremely low diffusion coefficients [169,170,171]. For sodium borosilicate glasses, which are currently under consideration as possible storage medium for radiotoxic nuclear fission products, one is e.g. interested in diffusion coefficients for cesium at the actual storage temperatures (400-750 K). The lowest temperature that proved feasible with conventional techniques was 825 K, which is still well above the glass transition temperature T 8 of this glass (786 K). Since structural changes are believed to occur at Tg, extrapolation of the measured coefficients from high temperatures to storage conditions is very risky. The very thin depletion layer (tens of nm) that can be formed by ion bombardment enables one to measure diffusion coefficients at temperatures as low as 723 K. In fig. 6.7 an Arrhenius plot is given of the results obtained by Sengers at al. [172] using the conventional technique and those of Ceelen et al. [170,171] using the depletion layer technique. There is good quantitative agreement between the results. The depletion layers were formed with He + for the thickest layer and with Ar + (smaller penetration depth) for the thinner ones. In this way, diffusion coefficients from only 2.6 ~ 10 .22 m2/s at the lowest temperature (723 K) up to 3.0 ~ 10 18 (849 K) could be measured. The fact that the temperatures above and below Tg fall on the same line implies that the activation energy for Cs diffusion does not change significantly for this glass at Tg. This may be due to the fact that while B203 and SiO 2 are network formers (for which significant changes occur around T~) Cs20 is only a network modifier.

164

o

SenSel"S

*

LBIS He

o

-80

A m

-85

LEI$ Ar

Tg o

-40

-45

-50

-56 1.00

I 1.10

, 1.20

1 .SO

1.40

1.50

1ooorr (l/K) Figure 6.7 Arrhenius plot of Cs diffusion in sodium borosilicate glass with both the results obtained by Sengers et al. [172] and those measured by Ceelen et al. using LEIS [ 170].

7 GROWTH AND WETTING 7.1 Oxidation The adsorption of oxygen on metals and semiconductors has been the subject of many LEIS studies. Since matrix effects are generally absent in LEIS of oxides, a quantitative composition analysis is possible (see e.g. [173] and section 3.3). Angular dependent low-energy ion scattering has been used to determine the location of oxygen atoms on surfaces such as Ag(ll0) [174,175], Ni(100)[176], P t ( l l l ) [177]. In some studies the transition of oxygen adsorption to oxide formation has been investigated. The behavior depends strongly on the metal and experimental conditions. Ocal et al. [178] exposed first an A I ( l l l ) surface at 80 K to 500 L of oxygen and then warmed for a certain time in vacuum at 400 - 700 K. The incorporation of the oxygen in the bulk can then be followed by LEIS via the decreasing oxygen or the increasing A1 signal. The time-dependence of the A1 signal showed that the oxygen incorporation is not a random-walk process (t t/z dependence), but indicates that the jump of a surface O atom to an unoccupied site in the second layer is the rate determining step. An activation energy of 0.50 +_ 0.05 eV was found for the jump. It was also found that the details of the growth procedure have a significant influence on the stoichiometry of the surface of the oxide. For the interaction of oxygen with zirconium, the influence of S and C1 on the oxidation properties has been discussed by Hoflund et al. [179]. The S and

165

CI were too difficult to detect here by AES and XPS. For Si (111) and (100) the reaction with atomic oxygen was found to be two-dimensional layer-by-layer growth [180] for substrate temperatures below 900K. When oxidized metals are compared with their bulk oxides, LEIS generally shows that the surface of the oxidized films contain less oxygen. Asbury and Hoflund [181] find hardly any oxygen on the surface for polycrystalline tin even after long exposure to oxygen , although XPS and AES indicate significant oxygen uptake. For tin, the oxygen apparently already penetrates beneath the outermost layers during the early stage of oxidation 9 A possible complication in these measurements is that the charge transfer from tin to He+ is almost resonant [90]. Peng and Barteau [99] compared oxidized Mg(0001) with MgO(100) by using very lowenergy (200 eV) He* ion scattering. Fig. 7.1 shows that there is a large difference between

He +, 200eV

0.55

I 0.42 "'" o"

~ v,,,,,,i

~-

I 9 .

9

.':

...

I

(c)

..

~J ~ o 5~

4 ~' ~-,: ~ v,,,,,,q

.o

%

."

:::

,b"

9

9

I

.o 9 r ".o -..

.

s -~

o

",,,

9

(b)

.'t'.,.'-" I. o..

-,b ~

-"I', I

1

x0.1 .~.~.,.~,~,v~-----

I

0

I

I

I I I I

0.2

I

(a)

/0.60\.

._.

I

I

I

I

0.4

I

0.6

I

I

I

I

0.8

I

I

I

l.O

Ee/E i

Figure 7.1 LEIS-spectrum of: (a) clean M g ( 0 ~ l ) , (b) oxidized Mg(0001), and (c) clean MgO(100) [99].

166 the LEIS spectra of these surfaces. From angular dependent studies it is concluded that the oxidized Mg(0001) surface resembles most the polar M g O ( l l l ) surface. The much lower (factor of ten) Mg signal for the oxidized Mg and the MgO surface, as compared to the pure Mg, is somewhat surprising. Since the MgO(100) surface contains about as many Mg atoms as a pure Mg surface, one would not expect a reduction of the Mg signal if MgO were formed (see also section 3.2). The very low incident energies and the special scattering configuration that are used in their experiments may lead to exceptional shielding by neighboring atoms. Higher energy data would have simplified the interpretation. Since the spectra of the oxidized magnesium and the MgO show a very intense sputter peak, this may be an indication of hydroxyl contamination which would also lower the intensities. A fundamental problem in the understanding of oxidation is the direction of diffusion: does the substrate diffuse outward or the oxygen inward during oxidation? Schmidt et al. [39,182] developed a special marker technique to differentiate between these two types of diffusion. Instead of using an isotope as tracer, an impurity that does not diffuse is introduced as marker. This type of marker should deposit uniformly on the surface and show negligible diffusion into the substrate and oxide. The surface concentration should be well-below a monolayer in order to affect the diffusion as little as possible. As illustrated in fig. 7.2, the marker will be buried after one or two monolayers of oxide growth if the metal diffuses outward. For oxygen diffusion, however, the marker will remain on the

metal

oxygen ambient ~ s i t e d

metal outward diffusion

oxygeninward diffusion

Figure 7.2 Schematic illustration of the marker distribution for metal oxide film growth. The dashed line represents the oxide-metal interface [182].

167 surface. Using; Sn as a marker, the UV-enhanced oxidation of GaAs was studied by LEIS. When a 10 A oxide layer was formed on the GaAs, this did not reduce the Sn signal. When the oxidation of Ni was studied, however, it was found that the Sn signal was reduced, indicating that here the Ni 2§ diffuses out of the metal through the oxide to the oxide/oxygen interface [182]. The high surface sensitivity of LEIS thus allows one to distinguish inward from outward diffusion for layers as thin as 10 A. Since LEIS can often resolve the isotopes of an element, one can in principle also use isotopic tracers. However, this is does not always have advantages over an inert tracer as Sn. Niehus and Comsa [177] showed with LEIS for P t ( l l l ) that O atoms are chemisorbed in the topmost layer, but move to deeper layers in the oxide for temperatures above 800 K. Thin oxide films on metals are becoming popular as supports for model catalysts. An advantage of such supports is that they are flat and somewhat conductive (due to tunneling through the thin oxide layer). This enables one to use techniques such as AES which are not suited to the analysis of high surface area insulating supports such as ),-alumina. A prerequisite to make the model studies useful is, of course, that the composition and atomic structure of the model surface is the same as that of the real support. The above studies show that this condition is not easily fulfilled by the oxidation of pure metals. Strong oxidation of a metal may lead to the same surface as the oxide, but the film will then be too thick to be conductive. The use of alloys or thin metal overlayers can be a more attractive starting material for the oxidation. Thin layers of aluminum on Mo or Ta have been suggested, while NiA1 alloys also have received much attention. Since LEIS can be used for real as well as for model catalysts, this provides an interesting possibility of checking the validity of the assumption. Bardi et al. [183,184] investigated the oxidation of Ni3AI (111) and (001) at high and low oxygen pressures. Exposure to 10 .5 Pa at temperatures above 800 K led to the formation of aluminum oxide islands of uniform thickness (about 5 ,~) up to the complete coverage of the surface. No evidence was found in LEIS for the presence of Ni in the outermost atomic layer. Their LEED data show periodicity only in the oxygen lattice. This would imply [184] that the AI atoms are randomly distributed over the octahedral and tetrahedral sites. A ),-type alumina is, therefore, proposed. With LEIS it should be possible to determine the octahedral site occupancy of the AI atoms. When the alloy is exposed to higher pressures of oxygen (order of 10 Pa) Ni appears at the surface [184]. Shen et a/.[185] confirmed the results of Bardi et al. et low oxygen pressures. Using angular dependent LEIS they clearly proved that at 700 ~ alumina islands are formed first in the initial stages of oxidation. For NiA1 (110) it has also been suggested [186] that a 5 .,~ alumina film is formed by oxidation. Quantification of the surface composition by LEIS [19] showed that there are more A1 atoms in the oxidized surface than in y-Al203 and even more than in a-ml203. This does not exclude the possibility that the surface is a kind of ),-alumina. It could be that the presence of the nearby alloy surface pins the lattice of the oxide layer. The crystal surface may thus differ from that exposed by the bulk oxide in thermodynamic equilibrium. In any case, the LEIS results show that the surface of the oxide grown on NiA1 (110) is completely different from that of ),-alumina and is thus not suitable as model support in studies of catalysts. Quantitative LEIS is required to check whether the oxide on NiA13 is a better candidate. However, one should keep in mind that since the alumina layer is only 5 ,~ thick, the alloy can still affect the properties of the surface, making it a unsuitable model support for small metal particles in catalysis.

168 7.2 O x i d e s

on oxides.

Growth of oxides on oxidic supports finds widespread application in the fabrication of catalysts and electronic devices. Whether the growth takes place via evaporation, decomposition of volatile compounds or via some liquid phase impregnation, in all cases the wetting properties of the oxides greatly influence the growth mode. White and coworkers [187,188] used LEIS to study the growth of SiO 2 films on ZrO2, T i O 2 and ZnO via decomposition of Si(OC2Hs) 4. The reduction of the LEIS signal of the substrate provides an easy tool to monitor the growth. The behavior was found to depend strongly on the nature of the substrate. While silica forms a monolayer film on ZnO, it forms a monolayer plus three-dimensional clusters on ZrO 2 and a mixed silicatitanate overlayer on TiO 2. Martin and Netterfield [57] grew ZrO 2 by evaporation on silica. When a monolayer of zirconia was reached, as determined by in-situ ellipsometry, the Si had completely disappeared (see fig. 7.3). The zirconia clearly wets the silica. Leyrer et a/.[189] have discussed the wetting of oxides by other oxides. Wetting is expected when ),. - ),, < ),,,~,

(7.1)

where ),, and )'s are the surface energies of the adsorbate and substrate, while the interface energy )', is the excess energy in the interface (thus),, is defined here as 0 for the interaction between the same oxides). When the reaction between the oxides is exothermic, )'~ is negative. While surface energies of a number of oxides are nowadays reasonably well known, interfacial energies are not. The use of this equation is, therefore, limited. Since interfacial energies are generally smaller than the surface energies, the surface energies often determine the outcome. Moreover, from the occurrence or absence of mixed oxides one can at least estimate the sign of the chemical contribution to the interfacial energy. For adsorbates such as MOO3, W O 3 and V205, which all have very low surface energies, the spreading on a high surface energy material as y-alumina [22,37,136,190] is thus understood [189]. From the low surface energy of titania it is also clear why silica did not spread, while it did spread on ZnO. Comparing the results of Martin and Netterfield [57] with those by White et al. [187,188], it is surprising that not only does silica spread on zirconia, the reverse is also observed. It may be that a strong interfacial interaction is responsible, since a stable ZrSiO4 is known [191]. The spreading of molybdena on ),-alumina, which has already been discussed in section 5.4, is well established. The spreading of molybdena on silica is more controversial. While Stencel et al. [192] used LEIS to demonstrate that under calcination conditions molybdena spreads on silica. Exposure to water at low temperatures leads agian to clustering of the molybdena. This can be done reversibly. In more recent LEIS experiments [24,130] spreading of molybdena was also clearly established. However, the spreading of M o O 3 is not expected on the basis of solid-solid wetting of M o O 3 o n SiO2, as was shown by Leyrer et al. [189]. In the higly dispersed Mo/SiO 2 catalyst, as studied by Stencel et a/.[192] and Brongersma and coworkers [24,130], the spreading can be assigned to the break up of so-called Keggin units (silicomolybdic acid, H2SiMo12040.2H20 ) which can be restored upon water treatment. This microspreading [193] is reversible. The growth of oxides on oxides may also lead to the formation of surface spinels Oust as in surface segregation; see section 6.3). In studies of Cu oxide impregnated ),-alumina,

169

r~

0

!

Zr

',,

ol

~" 1,4

l/I;

,.,'I~

d

I

I

0.5

1.0

Ef / E i Figure 7.3 LEIS spectra of (a) 0.1 nm ZrO 2 deposited on silica (dashed line), (b) 0.3 nm ZrO 2 (dash-dot line) and (c) the uncoated substrate (solid line)[57].

Strohmeier et al. [163] stress the importance of a surface spinel that resembles CuAI204. This spinel is formed for low loadings of copper ((4% Cu for a 100 m2/g support) and not too high temperatures (below 500 ~ On the basis of these and other studies they argue that, in contrast to the regular spinel, the Cu 2§ ions occupy predominantly (90%) the (somewhat distorted) octahedral sites. For bulk CuAI204 only 40% of the Cu 2§ ions are found in octahedral sites. At high Cu loadings bulk-like CuO is formed on the surface. In studies of lanthanum oxide on ),-alumina, van Leerdam et al. [81] deposited the lanthanum via a [La(EDTA)] complex. In fig. 7.4 [81] the dependence of the La/A1 signal ratio in LEIS is plotted as a function of the La content for a number of treatments. The linear increase of the signal ratio for the as-prepared sample indicates complete spreading (all signals for La are somewhat reduced due to neutralization by the complex). During calcination at 550~ the lanthanum oxide agglomerates to small 3-dimensional particles (especially at the higher loadings), while the lanthanum spreads again (formation of a lanthanum aluminate surface layer) during further calcination at 1050 ~ The distribution of the La during the process steps is schematically shown in fig. 7.5 ([81]). It has also been tried to follow the La distribution with XPS. Due to the much larger information

170

9

0.4 0

~ v.,.,a

I

!

I

1

2

3

60 ~

A 5h. 1050 ~ o 23h. 1050~ t 145h. 1050 ~

0.2

0

La-content (~t%) Figure 7.4 Dependence of the LEIS La/Al-ratio on the thermal treatments at various bulk La-contents of Lanthanum-added ),-alumina [81].

depth of this technique, the agglomeration and spreading cannot be seen. When V205 is impregnated in )'-A1203 [140], the vanadium species is optimally dispersed at loading up to 7 wt%. When the vanadium loading is further increased (7-11 wt%), both the V and the A1 signals remain constant. Apparently the extra V shields the previously deposited V. For still higher V loadings the surface becomes fully covered by vanadia. The detailed information is again characteristic of the monolayer information depth of LEIS. The findings of Jacobs et al. [140] have recently been confirmed by Eberhardt et al. [194].

7.3. Oxides on metals and metals on oxides.

Various groups have used LEIS to obtain detailed insight into the growth of metals on oxides (see e.g. [22,195,196,197,198]) and of oxides on metals (e.g. [107,199,200]). The growth of iron oxide on Pt(111) and Pt(100) was found to be a layer-by-layer growth by Vurens et al. [107,200] using LEIS, LEED and AES. The FeO orders on both substrates. Although for both substrates an FeO (111) overlayer is proposed, the iron signal is much more intense for the overlayer on the Pt(100) surface. The shielding by the oxygen is apparently more effective for the contracted layer on the Pt(111) surface. It is suggested that the open nature of the FeO(111) allows the Fe, which is located in the

171

60 ~

550 ~

"///////////////////////~'/////////~

V///////////////~~////~//~///7/,~

1050 ~

0 [-~

lanthanum EDTA support

Figure 7.5 Schematic representation of the distribution of lanthanum over the ),A1203 support after the different thermal treatments [81].

second layer in contact with the Pt substrate, to be still accessible to the impinging He + ions. Iron oxide on Pt is also used as a model system for studying sintering under the influence of alkali. The diffusion of metals through oxides can depend strongly on the stoichiometry of the oxide. Ocal et al. [201] studied diffusion of metals through thin aluminum oxide on its metallic support at a temperature of 80K. Using LEIS it is shown that deposited Au atoms diffuse through the oxide to the oxide - metal interface (loss of Au at the surface), only when the oxide is deficient in oxygen. Deposited potassium does not diffuse through the oxide. The results are interpreted in terms of the Cabrera - Mott mechanism for lowtemperature oxidation of metals. According to this mechanism a potential difference is created across the thin oxide. This acts as the driving force for anion diffusion through the oxide. The electron affinity of gold is high, so anions are formed that are driven into the oxide by the field. Potassium does not form anions and is thus not affected. The high diffusion barrier in stoichiometric oxides prevents diffusion at these low temperatures. Carver et al. [22] followed the clustering of Pt and/or platinum oxide on alumina as a result of heat treatments in oxygen. Skoglundh et al. [202] characterized two Pt/Pd catalysts on cordierite monoliths with and without prior hydrothermal treatment of the washcoat. From LEIS a surface enrichment of Pd was found in both catalysts. LEIS showed that the hydrothermal treatment increased the dispersion of the Pt. Jacobs et al. [139] combined catalytic activity measurements and surface analysis to model the growth mechanism of nickel in the preparation of Ni/AI203 by atomic layer epitaxy (ALE). Pitts et al. [196] used variable angle LEIS to determine the position of silver on tinsensitized silica. Hoflund et al. [195] used LEIS in combination with AES and XPS to establish that annealing in hydrogen of Pt/TiO 2 reduces the titania to a lower oxide which then moves over the Pt and covers it. The encapsulating titanium species migrate again off the Pt upon oxidation.

172 Madey and co-workers made a detailed study of the growth of ultrathin metal films on a TiO2 (110) surface using LEIS [203,204,205,206]. Due to the monolayer sensitivity of LEIS, one can clearly distinguish between growth modes where 3dimensional clusters are formed (Volmer-Weber) and there where initially a 2-dimensional monolayer is deposited (Stranski-Krastanov or Frank-van-der-Merwe). The wetting, where metals were deposited at room temperature by thermal vapor deposition, seemed to follow the reactivity of the 3d-metals with oxygen. Their results show that the wetting ability decreases as follows: Cr > Fe > Cu. Copper grows in 3D-clusters with the least spreading ability. Fe grows in flat 3D-islands while complete wetting is observed for Cr on TiO2 (110). SMSI effect could be observed: upon annealing Fe on TiO2 (110) a complete encapsulation by TiO2 was observed, while the Fe clusters remained in the metallic state

[207].

REFERENCES Ill [21 [31 [41 [51 [61 [7] [8] [91 [lt)] [ll] [12] [131 [14] [15]

[16] [17] [181 [19] [20] [21l [22]

B.V. Panin, Sov. Phys. JETP 15 (1962) 215. V. Walther and H. Hintenberger, Z. Naturforsch. 18A (1963) 843. D.P. Smith, J. Appl. Phys. 38 (1967) 340. D.P. Smith, Surf. Sci. 25 (1971) 171. M. Aono, Nucl. Instr. Meth. B2 (1984) 374 R. Souda, T. Aizawa and Y. Ishizawa, J. Vac. Sci. Technol. A8 (1990) 3218. R. Souda, M. Aono, C. Oshima, S. Otani and Y. lshizawa, Surf. Sci. 179 (1987) 199. H. Eschenbacher, A. Richard and V. Dose, Appl. Phys. A34 (1984) 19. T.M. Buck, G.H. Wheatly and L. Marchut, Phys. Rev. Lett. 51 (1983) 43. J.W. Rabalais, Surf. Sci. 300 (1994) 219 H.H. Brongersma and P.M. Mul, Chem. Phys. Lett. 14 (1972) 380. H.H. Brongersma and P.M. Mul, Surf. Sci. 35 (1973) 393. R.H. Bergmans, W.J. Huppertz, R.G. van Welzenis and H.H. Brongersma, Nucl. Instr. Meth. B64 (1992) 584. A.W. Czanderna, A.C. Miller, H.H.G. Jellinek and H. Kachi, J. Vac. Sci. Technol. 14 (1977) 227. I.C. Fullarton, J.-P. Jacobs, H.E. van Benthem, J.A. Kilner, H.H. Brongersma, P.J. Scanlon and B.C.H. Steele, to be published. D.F. Cox and T.B. Fryberger, Surf. Sci. 227 (1990) L I05. W. Englert, E. Taglauer and W. Heiland, Surf. Sci. 117 (1982) 124. S.H. Overbury and R.J.A. van den Oetelaar, Surf. Sci. 301 (1994) 313. J.-P. Jacobs, S. Reijne, R.J.M. Elfrink, S.N. Mikhailov, M. Wuttig and H.H. Brongersma, J. Vac. Sci. Techn. A I2 (1994) 2308. R.E. Honig and W.L. Harrington, Thin Solid Films 19 (1973) 43. R.C. McCune, J. Vac. Sci. Technol. 18 (1981) 700. J.C. Carver, S.M. Davis and D.A. Goetsch, in "Catalyst Characterization Science", Chapter 12, Ed. Publ. Co. American Chemical Society (1985)

173 [23] [24]

[251

[26] [27]

[28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38] [39] [40] [411 [42] [431

[44] [45] [46] [471 [481 [49] [501 [51] [52] [53]

[541

B.A. Horrell and D.A. Cocke, Catal. Rev.-Sci.Eng. 29 (1987) 447. H.H. Brongersma and G.C. van Leerdam in "Fundamental Aspects of Heterogeneous Catalysis Studied by Particle Beams", p. 283, Eds. H.H. Brongersma and R.A. van Santen, NATO ASI series B 265 (1991), Plenum press. E. Taglauer in "Fundamental Aspects of Heterogeneous Catalysis Studied by Particle Beams", p. 301, Eds. H.H. Brongersma and R. A. van Santen, NATO ASI B 265 (1991), Plenum press. J.B. Theeten and H.H. Brongersma, Rev. de Phys. Appl. 11 (1976) 57 (in French) R.H. Bergmans, M. van de Grift, A.W. Denier van der Gon, R.G. van Welzenis, H.H. Brongersma. S.M. Francis and M. Bowker, Nucl. Instr. Meth. B85 (1994) 435. J.W. Rabalais, H. Bu and C. Roux, Nucl. Instr. Meth. B64 (1992) 559. M. Aono, M. Katayama and E. Nomura, Nucl. Instr. Meth. B64 (1992) 29. R.F. Goff and D.P. Smith, J. Vac. Sci. Techn. 7 (1970) 72 W. Heiland, H.G. Schaeffler and E. Taglauer, Surf. Sci. 35 (1973) 381. H.H. Brongersma and P.M. Mul, Chem. Phys. Lett. 14 (1972) 380. J.F. Kelso, C.G. Pantano and S.H. Garofalini, Surf. Sci. 134 (1983) L543 R. de Gryse, J. Landuyt, L. Vandenbroucke and J. Vennik, Surf. Interf. Anal. 4 (1982) 168. J.F. Kelso and C.G. Pantano, J. Vac. Sci. Technol. A3 (1985) 1343 R.C. McCune and P. Wynblatt, J. Am. Ceram. Soc. 66 (1983) 111. R. Margraf, J. Leyrer, H. KnSzinger and E. Taglauer, Surf. Sci. 189/190 (1987) 842. C. Ocal, B. Basurco and S. Ferret, Surf. Sci. 157 (1985) 233. M.T. Schmidt, Z. Wu, C.F. Yu and R.M. Oswood, Jr., Surf. Sci. 226 (199 l) 199. P.E. West and P.M. George, J. Vac. Sci. Technol. A5 (1987) 1124. J.F. Geiger, K.D. Schierbaum and W. G/~pel, Vacuum 41 (1990) 1629. H.H. Brongersma, N. Hazewindus, J.M. van Nieuwland, A.M.M. Otten and A.J. Smets, Rev. Sci. Instr. 49 (1978) 707. G.J.A. Hellings, H. Ottevanger, S.W. Boelens, C.L.C.M. Knibbeler and H.H. Brongersma, Surf. Sci. 162 (1985) 913. G.J.A. Hellings, H. Ottevanger, C.L.C.M. Knibbeler, J. van Engelshoven and H.H. Brongersma, J. Electron. Spectrosc. RelaL Phenom. 49 (1989) 359. R.H. Bergmans, A.C. Kruseman, C.A. Severijns and H.H. Brongersma, Appl. Surf. Sci. 70/71 (1993) 283. A. Miotello and P. Mazzoldi, J. Phys. C16 (1983) 221 H.F. Helbig, P.J. Adelman, A.C. Miller and A.W. Czandema, Nucl. Instr. Meth. 149 (1978) 581 W. Griinert, R. Schl6gl and H.G. Karge, Surf. Interf. Anal. 20 (1993) 603. W. Griinert, R. Schl/3gl and H.G. Karge, J. Phys. Chem. 97 (1993) 8638. R. Souda and M. Aono, Nucl. Instr. Meth. B15 (1986) 114. S.N. Mikhailov, R.J.M. Elfrink, J.-P. Jacobs, L.C.A. van den Oetelaar, P.J. Scanlon and H.H. Brongersma, Nucl. Instr. Meth. B93 (1994) 149. W.L. Baun, Phys. Rev. A17 (1978) 849 T.M. Thomas, H. Neumann, A.W. Czanderna and J.R. Pitts, Surf. Sci. 175 (1986) L737. A.G.J. de Wit, R.P.N. Bronckers and J.M. Fluit, Surf. Sci. 82 (1979) 177

174

[55] [56] [57] [58] [59] [60] [61] [62] [63]

[64] [65] [66] [67] [68] [69] [70] [71] [72] [73] [74] [751 [76] [77]

[78] [79]

[801

[8~1 [82]

[83]

[841 [851

G.C. Nelson, J. Vac. Sci. Technol. 15 (1978) 702. L.L. Tongson, A.S. Bhalla and B.E. Knox, Mat. Res. Bull. 16 (1981) 775. P.J. Martin and R.P. Netterfield, Surf. Interf. Anal. 10 (1987) 13. G.C. van L~erdam, K.-M.H. Lenssen and H.H. Brongersma, Nucl. Instr. Meth. B45 (1990) 390. C.A. Severijns, G. Verbist and H.H. Brongersma, Surf. Sci. 279 (1992) 297. C. Creemers, D. Royer and P. Schryvers, Surf. Interf. Anal. 20 (1993) 233 G.C. van Le~rdam and H.H. Brongersma, Surf. Sci. 254 (1991) 153. W.L. Baun, J.S. Solomon, Anal. Chem. 48 (1976) 931 Ch. Lindsmeier, E. Taglauer and H. KnSzinger, in "Fundamental Aspects of Heterogeneous Catalysis Studied by Particle Beams", p. 413, Eds. H.H. Brongersma and R.A. van Santen, NATO ASI series B 265 (1991), Plenum press K. Josek, Ch. Lindsmeier, H. KnSzinger and E. Taglauer, Nucl. Instr. Meth. B64 (192) 596 G.C. van Leerdam, J.-P. Jacobs and H.H. Brongersma, Surf. Sci. 268 (1992) 45. G.C. Nelson. J. Vac. Sci. Technol. A4 (1986) 1567 V.Y. Young, G.B. Hoflund and A.C. Miller, Surf. Sci. 235 (1990) 60 V.Y. Young, N. Welcome, G.B. Hoflund, Phys. Rev. B48 (1993) 2981 J. Nowotny , "Interface defect chemistry and its impact on properties of oxide ceramic materials" in Science of Ceramic interfaces, Elsevier (1991) p. 79. J.R. Pitts and A.W. Czanderna, Nucl. Instr. Meth. B13 (1986) 245. J.B. Malherbe, S. Hoffmann and J.M. Sanz, Appl. Surf. Sci. 27 (1986) 355. R. Kelly, Surf. Sci. 100 (1980) 85. T. Choudhury, S.O. Saied, J.L. Sullivan and A.M. Abbot, J. Phys. D: Appl. Phys. 22 (1989) 1185. J.L. Sullivan, S.O. Saied, T. Choudbury, J. Phys.: Condens. Matter 5 (1993) A291 S.O. Saied, J.L. Sullivan, T. Choudhury and C.G. Pearce, Vacuum 38 (1988) 917. A. Benninghoven, F.G. Riidenauer and H.W. Wemer, in " Secondary Ion Mass Spectroscopy, Instrumental Aspects, Applications and Trends", Chemical Analysis vol. 86, John Wiley & Sons, New York, 1987 "Analysis of Microelectronic Materials and Devices", eds. M. Grasserbauer and H.W. Wemer, John Wiley & Sons, 1991 A.W. Czanderna, in "Ion Spectroscopies for Surface Analysis", p. 1l, Eds. A.W. Czanderna and D.M. Hercules, Plenum Press, New York, 1991. Y.-T. Cheng, S.J. Simko, M.C. Militello, A.A. Dow, G.W. Auner, M.H. Alkaisi and K.R. Padmanabhan, Nucl. Instr. Meth. B64 (1992) 38 J.S. Brinen, D.A. D'Avignon, E.A. Meyers, P.T. Deng, D.W. Behnken, Surf. Interf. Anal. 6 (1984) 295 G.C. van Le~rdam, H.H. Brongersma, I.I.M. Tijburg and J.W. Geus Appl. Surf. Sci. 55 (1992) 11. E. Taglauer and H. Kn/Szinger, in: "Surface Science, Principles and Applications", Eds. R.F. Howe, R.N. Lamb and K. Wandelt, Springer Proc. in Phys. 73, pp. 264, Springer-Verlag, Berlin, (1993). Ch. Linsmeier, H. Kntizinger and E. Taglauer, presented at EUROPACAT-I, September 12-17, 1993, Montpellier, France. W.K. den Otter, H.H. Brongersma, and H. Feil, Surf. Sci. 306 (1994) 215 G. Moliere, Z. Naturforschung 2A (1947) 133.

175 [861 [87]

[881 [89] [90] [91] [92] [93] [94] [95] [96] [97]

[98] [99] [100] [101] [102] [103] [104] [105] [1o6] [107] [108] [109] [ll0] [lll] Ill2] [113] [ll4] [115] [116] [117] [118] [119] [120]

J.M. van Zoest, C.E. van der Meij, J.M. Fluit and A. Niehaus, Surf. Sci. 152/153 (1985) 106. L.C.A. van den Oetelaar, S.N. Mikhailov and H.H. Brongersma, Nucl. Instr. Meth. B85 (1994) 420. S.N. Mikhailov, L.C.A. van den Oetelaar and H.H. Brongersma, Nucl. Instr. Meth. B93 (1994) 210. R.L. Erickson and D.P. Smith, Phys. Rev. Lett. 34 (1975) 297. T.W. Rusch and R.L. Erickson in "Inelastic ion-surface collisions", Eds. N.H. Tolk, J.C. Tully, W. Heiland and C.W. White, p. 73, Academic Press (1977) New York. H.H. Brongersma, G.C.J. van der Ligt and G.C. Rouweler, Philips J. Res. 36 (1980) 1. P.A.J. Ackermans, G.C.R. Krutzen and H.H. Brongersma, Nucl. Instr. Meth. B45 (1990) 384. H.H. Brongersma and J.-P. Jacobs, Appl. Surf. Sci. 75 (1994) 133. E. Taglauer, Appl. Phys. A38 (1985) 161 D.G. Swartzfager, Anal. Chem. 56 (1984) 55 M. Beckschulte and E. Taglauer, Nucl. Instr. Meth. B78 (1993) 29 R. Souda, T. Aizawa, K. Miura, C. Oshima, S. Otani, Y. Ishizawa, Nucl. Instr. Meth. B33 (1988) 374 M.J. Ashwin and D.P. Woodruff, Surf. Sci. 244 (1991) 247 X.D. Peng and M.A. Barteau, Appl. Surf. Sci. 44 (1990) 87. M. Shelef, M.A.Z. Wheeler and H.C. Yao, Surf. Sci. 47 (1975) 697. E. Haeussler and G.R. Sparrow, Silic. Ind. 43 (1978) 253. H.H. Brongersma and T.M. Buck, Nucl. Instr. Meth. B149 (1978) 569. J.R. Diehl, J.M. Stencel, C.A. Spitter and L.E. Makovsky, Surf. Interf. Anal. 6 (1984) 56 J.F. Kelso and C.G. Pantano, Surf. Interf. Anal. 7 (1985) 228. E. Taglauer, U. Beitat and W. Heiland, Nucl. Instr. Meth. 149 (1978) 605. F. Delannay, E.N. Haeussler and B. Delmon, J.Catal. 66 (1980) 469. G.H. Vurens, V. Maudce, M. Salmeron and G.A. Somorjai, Surf. Sci. 268 (1992) 170. U. Bardi, Appl. Surf. Sci. 51 (1991) 89. P.A.J. Ackermans, M.A.P. Creuwels, H.H. Brongersma and P.J. Scanlon, Surf. Sci. 227 (1990) 361 G.C. Nelson, J. Appl. Phys. 47 (1976) 1253. R. Margraf, H. KnSzinger and E. Taglauer, Surf. Sci. 211/212 (1989) 1083. P.R. Berning and A. Niiler, Nucl. Instr. Meth. B63 (1992) 434. H. Niehus, W. Heiland, E. Taglauer, Surf. Sci. Rep. 17 (1993) 213 T.N. Taylor and W.P. Ellis, Surf. Sci. 107 (1981) 249. J.P. Landuyt, L. Vandenbroucke, R. de Gryse and J. Vennik, Surf. Sci. 126 (1983) 598. S. Tanaka, T. Nakamura, M. Iiyama, N. Yoshida, S. Takona, F. Shoji and K. Oura, Appl. Phys. Lett. 59 (1991) 3637. G. Verbist, J.T. Devreese and H.H. Brongersma, Surf.Sci. 223 (1990) 323. H. Nakamatsu, A. Sudo and S. Kawai, Surf. Sci. 223 (1989) 193. E.A. Colboum and W.C. Mackrodt, Solid State Ionics 8 (1983) 221. D.F. Cox, T.B. Fryberger and S. Semancik, Phys. Rev. B38 (1988) 2072.

176 [121] R.C. McCune, J. Vac. Sci. Technol. 16 (1979) 1569. [122] G.C. van Leerdam and H.H. Brongersma, Proc. Symp. Surf. Sci., La Plagne (1990) 161. [1231 B.L. Maschhoff, J.-M. Pan, R.A. Baragiola and T.E. Madey, in: "Fundamental Aspects of Heterogeneous Catalysis Studied by Particle Beams", Eds H.H. Brongersma and R.A. van Santen, NATO-ASI series B265, pp. 419, Plenum Press, New York, (1991). [124] G. Verbist, H.H. Brongersma and J.T. Devreese, Nucl. Instr. Meth. B64 (1992) 572. [125] G.C. van Leerdam, PhD Thesis , Eindhoven University of Technology, The Netherlands, 1991. [126] H.H. Brongersma, L.C.M. Beirens and G.C.J. van der Ligt in "Materials Characterization using Ion Beams", p. 65, Eds. J.P. Thomas and A. Cachard, Plenum, (1978). [1271 H. KniSzinger and P. Ramasamy, Catal. Rev.-Sci. Eng. 17 (1978) 31. [128] H.C. Yao and M. Shelef, J. Phys. Chem. 78 (1974) 2490. [129] J.-P. Jacobs, A. Maltha, J.G.H. Reintjes, J. Drimal, V. Ponec and H.H. Brongersma, J. Catal. 147 (1994) 294. [130] H.H. Brongersma and J.-P. Jacobs, Appl. Surf. Sci. 75 (1994) 133. [131] B.M. Strohmeier and D.M. Hercules, J. Catal. 86 (1984) 266. [1321 M.R. Anantharaman, J.-P. Jacobs, H. Rosink, S. Reijne and H.H. Brongersma, unpublished [1331 R.W. Grimes, A.B. Anderson and A.H. Heuer, J. Am. Chem. Soc. 111 (1989) 1. [1341 F.M. Mulcahy, M.Houalla and D.M. Hercules, Anal. Chem. 62 (1990) 2232. [1351 J.E. Drawdy, G.B. Hoflund, S.D. Gardner, E.A. Yngvadottir and D.R. Schryder, Surf. Interf. Anal. 16 (1990) 369. [136] R. Prada Silvy, J.M. Beuken, J.L.G. Fierro, P. Bertrand and B. Delmon, Surf. Interf. Anal. 8 (1986) 167. [137] S. Kasztelan, J. Grimblot and J.P. Bonnelle, J. Phys. Chem. 91 (1987) 1503. [138] D.S. Zingg, L.E. Makovsky, R.E. Tischer, F.R. Brown and D.M. Hercules, J. Phys. Chem. 84 (1980) 2898. [139] J.-P. Jacobs, L.P. Lindfors, J.G.H. Reintjes, O. Jylh~i and H.H. Brongersma, Catal. Lett. 25 (1994) 315. [140] J.-P. Jacobs, G.C. van Leerdam and H.H. Brongersma in "Fundamental Aspects of Heterogeneous Catalysis Studied by Particle Beams", p. 399, Eds. H.H. Brongersma and R.A. van Santen, NATO ASI B 265 (1991), Plenum press. [141] B. HorreU, D.L. Cocke, G. Sparrow and J. Murray, J. Catal. 95 (1985) 309. [142] H. Jeziorowski, H. Kn6zinger, E. Taglauer and C. Vogdt, J. Catalysis 88 (1983) 286. [143] J.-M. Beuken and P. Bertrand, Surf. Sci. 162 (1985) 329. [1441 C.P. Li and D.M. Hercules, J. Phys. Chem. 88 (1984) 456. [145] S. Priggemeyer, H. Koschmieder, G.N. van Wijk and W. Heiland, Fresenius J. Anal. Chem. 341 ( 1991) 343. [146] H.H. Brongersma and T.M. Buck, private communication in H.H. Werner, Science of Ceramics 8 (1976) 73. [1471 C. Creemers, H. van Hove and A. Neyens in "Recent Dev. in Condens. Mater. Phys. 2 (1981) p. 363, Ed. J.T. Devreese, Plenum, New York.

177 [1481 H.H. Brongersma, C.M.G. Jochem, T.P.M. Meeuwsen, P.J.W. Severin and G.A.C.M. Spierings, Acta Electron. 22 (1979) 245. [1491 U. Lindner and H. Papp, Appl. Surf. Sci. 32 (1988) 75. [150] H.H. Brongersma and T.M. Buck, Surf. Sci. 53 (1975) 649. [1511 H.H. Brongersma, P.A.J. Ackermans and A.D. van Langeveld, Phys. Rev. B34 (1986) 5974. [1521 C.R. Helms, J. Catal. 36 (1975) 114. [153] R.C. McCune, Mat. Res. Soc. Symp. Proc. 48 (1985) 27. [1541 P. Wynblatt and R.C. McCune in "Surfaces and Interfaces in Ceramic and Ceramic -Metal Systems", p. 83-95, Eds. J.A. Pask and A.G. Evans, Plenum, New York (1981) [155] P. Wynblatt and R.C. McCune in "Surface and Near-Surface Chemistry of Oxide Materials, p. 247-279, Eds. J. Nowotny and L.-C. Dufour, Elsevier, Amsterdam (1988). [1561 W.D. Kingery, J. Am. Ceram. Soc. 57 (1974) 1 and 74. 1503. [1571 S. Baik, D.E. Fowler, J.M. Blakeley and R. Raj, J. Am. Ceram. Soc. 60 (1985) 281. [158] J.F. Kelso and C.G. Pantano, J. Vac. Sci. Technol. A3 (1985) 1343. [159] W.C. Mackrodt and P.W. Tasker, Mat. Res. Soc. Symp. 60 (1986) 291. [160] J.P. Beaufils and Y. Barbaux, J. Chim. Phys. 78 (1981) 347. (in French) [161] A. Rolland and B. Aufray, Surf. Sci. 162 (1985) 530 [162] M.A.Z. Wheeler and M. Bettman, J. Catalysis 40 (1975) 124. [1631 B.R. Strohmeier, D.E. Leyden, R. Scott Field and D.M. Hercules, J. Catal. 94 (1985) 514. [1641 C.J. McHargue, G.C. Farlow, P.S. Sklad, C.W. White, A. Perez, N. Kornilios and G. Marest, Nucl. Instr. Meth. B19/20 (1987) 813. [165] R.A. van Hassel and A.J. Burggraaf, Appl. Phys. A53 (1991) 155. [166] D.V. McCaughan and V.T. Murphy, J. Appl. Phys. 44 (1973) 2008. [1671 R.A. Kushner, D.V. McCaughan, V.T. Murphy and J.A. Heilig, Phys. Rev. BI0 (1974) 2632. [1681 G. Battaglin, G. Delia Mea, G. de Marchi, P. Mazzoldi and A. Miotello, Nucl. Instr. Meth. BI (1984) 511 [1691 O. van Kessel, H.H. Brongersma, J.G.A. Htlscher, R.G. van Welzenis, E.G.F. Sengers and F.J.J.G. Janssen, Nucl. Instr. Meth. B64 (1992) 593. [170] W.C.A.N. Ceelen, J.-P. Jacobs, H.H. Brongersma, E.G.F. Sengers and F.J.J.G. Janssen, NEVAC blad 4 (1992) 91. (in Dutch) [171] W.C.A.N. Ceelen, J.-P. Jacobs, H.H. Brongersma, E.G.F. Sengers and F.J.J.G. Janssen, Surf. Interf. Anal., submitted [172] E.G.F. Sengers, F.J.J.G. Janssen and H. de Waal, "Scientific Basis for Nuclear Waste Management XIII, Mat. Res. Soc. Symp. Proc. 176 (1990) 441. [1731 L.C.A. van den Oetelaar, J.-P. Jacobs, M.J. Mietus, H.H. Brongersma, V.N. Semenov and V.G. Glebovsky, Appl. Surf. Sci. 70/71 (1993) 79. [174] E. Taglauer and W. Heiland, Appl. Phys. 9 (1976) 261 [175] M. Canepa, P. Cantini, F. Fossa, L. Mattera and S. Terreni, Phys. Rev. B47 (1993) 15 823 [176] H.H. Brongersma and J.B. Theeten, Surf. Sci. 54 (1976) 519. [1771 H. Niehus and G. Comsa, Surf. Sci. 93 (1980) L147.

178 [178] C. Ocal, B. Basurco and S. Ferrer, Surf. Sci. 157 (1985) 233. [179] G.B. Hoflund, G.R. Corallo, D.A. Ashbury and R.E. Gilbert, J. Vac. Sci. Technol. A5 (1987) 1120. [1801 J.R. Engstrom, D.J. Bonser and T. Engel, Surf. Sci. 268 (1992) 238. [181] D.A. Ashbury and G.B. Hoflund, J. Vac. Sci. Technol. AS (1987) 1132. [182] M.T. Schmidt, Z. Wu and R.M. Oswood, Jr., Surf. Interf. Anal. 17 (1991) 43. [1831 U. Bardi, A. Atrei and G. Rovida, Surf. Sci. 239 (1990) L511. [1841 U. Bardi, A. Atrei and G. Rovida, Surf. Sci. 268 (1992) 87. [185] Y.G.Shen, D.J. O'Connor and R.J. MacDonald, Surf. Interf. Anal. 17 (1991) 903. [186] M. Wuttig, W. Hoffmann, R. Jaeger, H. Kuhlenbeck and H. Freund, Mat. Res. Soc. Symp. Proc. 221 (1991) 143. [187] T. Jin and J.M. White, Surf. Interf. Anal. 11 (1988) 517. [1881 S.K. Jo, T. Jin and J.M. White, Appl. Surf. Sci. 40 (1989) 155. [189] J. Leyrer, R. Margraf, E. Taglauer and H. Kn6zinger, Surf. Sci. 201 (1988) 603. [190] L. Salvatti,Jr., L.E. Makovsky, J.M. Stencel, F.R. Brown and D.M. Hercules, J. Chem. Phys. 85 (1981) 3700. [191] W.C. Butterman and W.R. Foster, Am. Mineralogist 52 (1967) 884. [192] J.M. Stencel, J.R. Diehl, J.R. D'Este, L.E. Makovsky, L. Rodrigo, K. Marcinkowska, A. Adnot, P.C. Roberge and S. Kaliaguine, J. Chem. Phys. 90 (1986) 4739. [193] M. de Boer, A.J. van Dillen, D.C. Koningsberger, J.W. Geus, M.A. Vuurman and I.E. Wachs, Cat. Lett. 11 (1991) 227. [1941 M.A. Eberhardt, M. Houalla and D.M. Hercules, Surf. Interf. Anal. 20 (1993) 766 [195] G.B. Hoflund, A.L. Grogan, Jr., and D.A. Asbury, J. Catal. 109 (1988) 226. [1961 J.R. Pitts, T.M. Thomas and A.W. Czandema, J. Vac. Sci. Technol. A4 (1986) 1653. [197] U. Bardi, G. Prelazzi, A. Santucci and G. Rovida, Catal. Lett. 5 (1990) 315. [198] S.D. Gardner, G.B. Hoflund, M.R. Davidson and D.R. Schryer, J. Catal. 115 (1989) 132. [199] R.J. Gorte, E. Altman, G.R. Corallo, M.R. Davidson, D.A. Asbury and G.B. Hoflund, Surf. Sci. 188 (1987) 327. [200] G.H. Vurens, D.R. Strongin, M. Salmeron and G.A. Somorjai, Surf. Sci. 199 (1988) L387. [2011 C. Ocal, B. Basurco and S. Ferrer, Surf. Sci. 163 (1985) 335. [202] M. Skoglundh, L.O. ld3wendahl, P.G. Menon, B. Stenbom, J.-P. Jacobs, O. van Kessel and H.H. Brongersma, Cat. Lett. 13 (1992) 27. [2031 Th.E. Madey, U. Diebold and J.-M. Pan, in: "Adsorption on Ordered Surfaces on Ionic Solids and Thin Films", Eds. E. Umbach and H.-J. Freund, Springer Proc. in Surf. Sci. 33, p 147, Springer-Verlag, Berlin, 1993 [204] U. Diebold, J.-M. Pan and T.E. Madey, Phys. Rev. B47 (1993) 3868. [2051 J.-M. Pan, U. Diebold, L. Zhang and T.E. Madey, Surf. Sci. 295 (1993) 411. [2061 J.-M. Pan and T.E. Madey, J. Vac. Sci. Technol. A l l (1993) 1667. [2071 J.-M. Pan and T.E. Madey, Catal. Lett 20 (1993) 269.

179 [2081 J.-M. Pan, B.L. Maschhoff, U. Diebold and T.E. Madey, J. Vac. Sci. Technol. AI0 (1992) 2470. [2091 D. Dissanayake, J.H. Lundsford, M.P. Rosynek, J. of Catal. 143 (1993) 286. [2101 B.G. Baker and M. Jasieniak, in: "Surface Science, Principles and Applications", Eds. R.F. Howe, R.N. Lamb and K. Wandelt, Springer Proc. in Phys. 73, p 279, Springer-Vedag, Berlin, (1993). [211] M. Salmeron, Stud. Phys. Theor. Chem. 77 (Comput. Chem., Pt.B) 55. [2121 Y.G. Shen, D.J. O'Connor and R.J. MacDonald, Surf. Interf. Anal. 18 (1992) 729. [213] M.C. Asensio, M. Kerkar, D.P. Woodruff, A.V. De Carvalho, A. Fernandez, A.R. Gonzalez-Elipe, M. Femanda-Garcia and J.C. Conesa, Surf. Sci. 273 (1992) 31. [2141 J.E. Drawdy, G.B. Hoflund, M.R. Davidson, B.T. Upchurch and D.R. Schryer, Surf. Interf. Anal. 19 (1992) 559. [215] R.P. Nettermeier, K.-H. Miiller, D.R. McKenzie, M.J. Goonan and P.J. Martin, J. Appl. Phys. 63 (1988) 760. [216] J.R. Pitts, T.P. Massopust, A.W. Czanderna and L.L. Kazmerski, J. V ac. Sci. Techn. A4 (1986) 241. [217] J.R. Pitts, S.D. Bischke, J.L. Falconer and A.W. Czanderna, J. Vac. Sci. Technol. A2 (1984) 10(~. [2181 A. Rossi, C. Calinski, H.W. Hoppe and H.H. Strehblow, Surf. Interf. Anal. 18 (1992) 269. [219] Y. Shen, D.J. O'Connor and R.J. MacDonald, Nucl. Instr. Meth. B67 (1992) 350. [220] S.V. Hattangady, M.J. Mantini, G.G. Fountain, R.A. Rudder and R.J. Markunas, J. Appl. Phys. 71 (1992) 3842. [2211 X.D. Peng and M.A. Barteau, Catal. Lett. 12 (1992) 245. [222] W. Mahdi, J. Schiitze, G. Weinberg, R. Schoonmaker, R. SchlSgl and G. Ertl, Catal. Lett. 11 (1991) 19. [223] A.H. A1-Bayatti, K.G. Orrman-Rossiter, D.G. Armour, J.A. Van den Berg and S.E. DonneUy, Nucl. Instr. Meth. B63 (1992) 109. [224] A.A. Dzhurakhalov, E.S. Parilis, N.Yu. Turaev, F.F. Umarov and I.D. Yadgarov, Izv. Akad. Nauk SSSR, Ser. Fiz. 55 (1991) 2405 (in Russian). [225] F. Arena, B.A. Horrell, D.L. Cocke, A. Parmaliana and N. Giordano, J. Catal. 132 (1991) 58. [226] S.D. Gardner, G.B. Hoflund, M.R. Davidson, H.A. Laitinen, D.R. Schryer and B.T. Upchurch, Langmuir 7 (1991) 2140. [227] U. Bardi, A. Atrei, G. Rovida and P.N. Ross, Surf. Sci. 2511252 (1991) 727. [228] V.D. Borman, E.P. Gusev, Yu.N. Devyatko, Yu. Yu. Lebedinskii, S.V. Rogozhkin, V.N. Tronin and V.I. Troyan, Vopr. At. Nauki Tekh. , Ser. Fiz. Radiats 2 (1990) 84 (in Russian). [229] X.D. Peng and M.A. Barteau, Langmuir 7 (1991) 1426. [230] R. Souda, W. Hayami, T. Aizawa and Y. Ishizawa, Phys. Rev. B43 (1991) 10062. [231] S.J. Hoekje and G.B. Hoflund, Thin Solid Films 197 (1991) 367. [232] A. Miotello, G. Cinque, P. Mazzoldi and C.G. Pantano, Phys. Rev. B43 (1991) 3831. [233] S.J. Hoekje and G.B. Hoflund, Appl. Surf. Sci. 47 (1991) 43. [234] A.H. AI-Bayati, K.G. Orrman-Rossiter, J.A. Van den Berg and D.G. Armour, Surf. Sci. 241 (1991) 91.

180 [235] S.D. Gardner, G.B. Hoflund, D.R. Schryer and B.T. Upchurch, J. Phys. Chem. 95 (1991) 835. [236] Y. Tong, M.P. Rosynek and J.H. Lunsford, J. Catal. 126 (1990) 291. [237] S.P. Jeng, P.H. Holloway, D.A. Asbury and G.B. Hoflund, Surf. Sci. 235 (1990) 175. [238] I.A. Zotov, I.Yu. Panichkin, S.P. Chenakin and V.T. Cherepin, Sverkhprovodimost: Fiz., Khim., Tekh. 3 (1990) 233 (in Russian). [239] X.D. Peng and M.A. Barteau, Surf. Sci. 233 (1990) 283. [240] A.M. Then and C.G. Pantano, J. Non-Cryst. Sol. 120 (1990) 178. [241] Q. Guo, L. Gui, H. Huang, B. Zhao and Y. Tang, J. Catal. 122 (1990) 457. [242] R. Cavicchi, M. Tarlov and S. Semancik, J. Vac. Sci. Technol. A8 (1990) 2347. [2431 B. Baretziqr NATO-ASI ser. E155 (1989) 329. [244] S.D. Gardner, G.B. Hoflund and D.R. Schryer, J. Catal. 119 (1989) 179. [245] F. Rummens, P. Bertrand and Y. de Puydt, NATO-ASI, ser. E155 (1989) 101. [246] U. Lindner and H. Papp, Fresenius' Z. Anal. Chem. 333 (1989) 540. [247] H. Bubert, H. Puderbach, H. Pulm and W.A. Roland, Fresenius' Z. Anal. Chem. 333 (1989) 304. [248] S. Houssenbay, S. Kasztelan, H. Toulhoat, J.P. Bonnelle and J. Grimblot, J. Phys. Chem. 93 (1989) 7176. [249] A. Bellare, D.B. Dadyburjor and M.J. Kelly, J. Catal. 117 (1989) 78. [250] G.B. Hoflund, M.R. Davidson, E. Yngvadottir, H.A. Laitinen and S. Hoshino, Chem. Mater., 1 (1989) 625. [251] S.V. Hattangady, G.G. Fountain, D.J. Vitkavage, R.A. Rudder and R.J. Markunas, J. Electrochem. Soc. 136 (1989) 2070. [252] S. Haupt and H.H. Strehblow, Corros. Sci. 29 (1989) 163. [253] D. Ouafi, F. Mauge, J.-C. Lavalley, E. Payen, S. Kasztelan, M. Houari, J. Grimblot and J.P. Bonnelle, Catal. Today 4 (1988) 23. [254] V.D. Borman, E.P. Gusev, Yu.Yu. Lebedinskii and V.I. Troyan, Poverkhnost 11 (1988) 138. [2551 E.S. Shpiro, B.B. Dysenbina, O.P. Tkachenko, G.V. Antoshin and Kh. M. Minachev, J. Catal. 110 (1988) 262. [256] P.F. Carcia, F.D. Kalk, P.E. Bierstedt, A. Ferretti , G.A. Jones and D.G. Swartzfager, J. Appl. Phys. 64 (1988) 1671. [257] S.M. Davis, Catal. Lett. 1 (1988) 85. [258] S.H. Corn, J.L. Falconer and A.W. Czanderna, J. Vac. Sci. Technol. A6 (1988) 1012. [259] P.J.C. Chappell, M.H. Kibel and B.G. Baker, J. Catal. 110 (1988) 139. [260] Y. Zhou, M. Nakashima and J.M. White, J. Phys. Chem. 92 (1988) 812. [261] C.T. Campbell, K.A. Daube and J.M. Whim, Surf. Sci. 182 (1987) 458. [262] J.M. Van Zoest, J.M. Fluit, T.J. Vink and B.A. Van Hassel, Surf. Sci. 182 (1987) 179. [263] K.H. Muller, R.P. Netterfield and P.J. Martin, Phys. Rev. B35 (1987) 2934. [264] T.J. Vink, J.M. Der Kinderen, O.L.J. Gijzeman, J.W. Geus and J.M. Van Zoest, Appl. Surf. Sci. 26 (1986) 357. [265] J.W. Rabalais and J.N. Chen, J. Chem. Phys. 85 (1986) 3615. [266] G. Kremenic, V. Cortes Coberan and J.M.D. Tascon, Surf. Interf. Anal. 9 (1986) 207.

181 [267] J.C. Carver, I.E. Wachs and L.L. Murrell, J. Catal. 100 (1986) 500. [2681 R.P. Silvy, F. Delannay, P. Grange and B. Delmon, Polyhedron 5 (1986) 195. [269] S. Kasztelan, E. Payen, H. Toulhoat, J. Grimblot and J.P. Bonnelle, Polyhedron 5 (1986) 157. [270] J.M.D. Tascon, P. Bertrand, M. Genet and B. Delmon, J. Catal. 97 (1986) 300. [271] R. Tromp, G.W. Rubloff, P. Balk, F.K. LeGoues and E.J. van I.x~enen, Phys. Rev. Lett. 55 (1985) 2332. [272] G.B. Hoflund, D.A. Asbury, P. Kirszensztejn and H.A. Laitinen, Surf. Sci. 161 (1985) L583. [273] J.M. Stencel, L.E. Makovsky, J.R. Diehl and T.A. Sarkus, J. Catal. 95 (1985) 414. [274] M. Erbudak and F. Stucki, Phys. Rev. B32 (1985) 2667. [275] J.F. Kelso and C.G. Pantano, J. Vac. Sci. Technol. A3 (1985) 1343. [276] S. Kannar and P. Mohazzabi, J. Mater. Sci. Lett. 4 (1985) 720. [277] A.E.T. Kuiper, G.C.J. Van der Ligt, W.M. Wijgert, M.F.C. Willemsen and F.H.P.M. Habraken, J. Vac. Sci. Technol. B3 (1985) 830. [278] R.C. McCune and R.C. Ku, Adv. Ceram. 10 (1984) 217. [279] H. Puderbach and H.J. Goehausen, Spectrochim. Acta 39B (1984) 1547. [280] B.H. Davis, Appl. Surf. Sci. 19 (1984) 200. [281] R.P. Silvy, J.M. Beuken, P. Bertrand, B.K. Hodnett, F. Delannay and B. Delmon, Bull. Soc. Chim. Belg. 93 (1984) 775. [282] G.C. Allen, Met. Sci. 18 (1984) 295. [283] J.M. Stencel, L.E. Makovsky, J.R. Diehl and T.A. Sarkus, J. Raman Spectrosc. 15 (1984) 282. [2841 L.E. Makovsky, J.M. Stencel, F.R. Brown, R.E. Tischer and S.S. Pollack, J. Catal. 89 (1984) 334. [285] B.R. Strohmeier and D.M. Hercules, J. Phys. Chem. 88 (1984) 4922. [286] H.W. Wemer and N. Warmoltz, J. Vac. Sci. Technol. A2 (1984) 726. [287] P. Varga and E. Taglauer, Nucl. Instr. Meth. B230 (1984) 800. [288] S. Kasztelan, J. Grimblot and J.P. Bonnelle, J. Chim. Phys. Phys.-Chim. Biol. 80 (1983) 793. [289] B.P. Loechel and H.H. Strehblow, J. Electroehem. Soc. 131 (1984) 713. [290] P. Bertrand, J.M. Beuken and M. Delvaux, Nucl. Instr. Meth. 218 (1983) 249. [291] J.F. Kelso, C.G. Pantano and S.H. Garofalini, Surf. Sci. 134 (1983) L543. [292] C.G. Pantano, J.F. Kelso and M.J. Suscavage, Mater. Sci. Res. 15 (1983)1. [293] I.S.T. Tsong, Mater. Sci. Res. 15 (1983) 39. [294] R. Kumar, J.A. Schultz and J.W. Rabalais, Chem. Phys. Lett. 97 (1983) 256. [295] H.J. Lewerenz, D.E. Aspnes, B. Miller, D.L. Maim and A. Heller, J. Am. Chem. Soc. 104 (1982) 3325. [296] R.L. Chin and D.M. Hercules, J. Catal. 74 (1982) 121. [297] B.C. Rodrigo, H. Jeziorowski, H. KnSzinger, X.Z. Wang and E. Taglauer, Bull. Soc. Chim. Belg. 90 (1981) 1339. [2981 J. Abart, E. Delgado, G. Ertl, H. Jeziorowski, H. KnSzinger, N. Thiele, X.Zh. Wang and E. Taglauer, Appl. Catal. 2 (1982) 155. [299] A. Shih and G.A. Haas, Appl. Surf. Sci. 8 (1981) 125. [3001 H. KnSzinger, H. Jeziorowski and E. Taglauer, Stud. Surf. Sci. Catal. 7 (1981) 604. [301] P.G. Fox, Glass Technol. 22 (1981) 67.

182 [302] S.V. Krishnaswamy, L.L. Tongson, N. Said and R. Messier, J. Vac. Sci. Technol. 18 (1981) 401. [3031 R. De Gryse, J.P. Landuyt, A. Vermeire and J. Vennik, Appl. Surf. Sci. 6 (1980) 430. [304] F. Delannay, E.B. Haeussler and B. Delmon, J. Catal. 66 (1980) 469. [305] W.L. Baun, Surf. Technol. 11 (1980) 385. [306] W.L. Baun, Surf. Technol. 11 (1980) 421. [307] H.H. Strehblow and B. Titze, Electrochim. Acta 25 (1980) 839. [308] F. Delannay, E.N. Haeussler and B. Delmon, Bull. Soc. Chim. Belg. 89 (1980) 255. [3091 F. Delannay and E.N. Haeussler, Ned. Tijdschr. Vacuumtechnol. 18 (1980) 38. [3101 A.C. Miller and A.W. Czanderna, Appl. Surf. Sci. 4 (1980) 481. [3111 W.L. Baun, Appl. Surf. Sci. 4 (1980) 374. [3121 R.C. McCune, J.E. Chelgren and M.A.Z. Wheeler, Surf. Sci. 84 (1979) L515. [3131 M. Wu and D.M. Hercules, J. Phys. Chem. 83 (1979) 2003. [3141 R.C. McCune, Anal. Chem. 51 (1979) 1249. [3151 P.H. HoUoway and G.C. Nelson, J. Vac. Sci. Technol. 16 (1979) 793. [3161 M. Nagasaka, R. Vasofsky and H.F. Helbig, J. Vac. Sci. Technol. 16 (1979) 151. [3171 W.P. Ellis and T.N. Taylor, Surf. Sci. 75 (1978) 279. [318] K. Akaishi, A. Miyahara and A. Sagara, J. Nucl. mater. 76/77 (1978) 378. [319] R. Schubert, J. Electrochem. Soc. 125 (1978) 1215. [320] W.L. Baun, Phys. Rev. AI7 (1978) 849. [321] A.T. DiBenedetto, T. Anthony and D.A. Scola, J. Colloid Interf. Sci. 64 (1978) 492. [322] A.W. Czandema, A.C. Miller, H.H.G. Jellinek and H. Kachi, Ind. Res. 20 (1978) 62. [323] A.C. Miller and A.W. Czandema, Vacuum 28 (1978) 9. [3241 D.L. Christensen, V.G. Mosotti, T.W. Rusch and R.L. Erickson, Nucl. Instr. Meth. 149 (1978) 587. [3251 G.C. Nelson, J. Electrochem. Soc. 125 (1978) 403. [326] N.T. McDevitt, W.L. Baun and J.S. Solomon, J. Electrochem. Soc. 123 (1976) 1058. [327] R.P. Frankenthal and D.L. Maim, J. Electrochem. Soc. 123 (1976) 186. [3281 A. Emmanuel, G.B. Donaldson, W.T. Band and D. Dew-Hughes, IEEE Trans. Magn., MAG 11 (1975) 763. [329] R.F. Goff, J. Vac. Sci. Technol. 10 (1973) 355.

Science of Ceramic Interfaces II J. Nowotny (Editor) 9 1994 Elsevier Science B.V. All rights reserved.

183

INTERFACIAL PHENOMENA IN Y203-ZrO2-BASED CERAMICS: A SURFACE SCIENCE PERSPECTIVE A.E. Hughes CSIRO, Division of Materials Science and Technology Locked Bag 33, Clayton, Victoria, Australia, 3168

Abstract Zirconia-based materials combine good thermomechanical properties and oxygen ion conductivity with chemical inertness and structural stability in a range of chemical environments. Polycrystalline materials, which are used in the bulk of applications, display a rich variety of segregation behaviour due to impurities which are invariably present either in the starting powders or inadvertently incorporated during processing. The mechanisms whereby, many, though not all, of these impurities and alloying components accumulate at the external or grain boundary interfaces are examined. The accumulation of impurities or components at the grain boundary can significantly change the grain boundary composition. Consequently, many of the macroscopic properties of these materials can depend critically on the chemistry of the grain boundary. An overview from both a thermodynamic and atomistic approach, of the development of the grain boundary network and the changing chemistry in the region of the grain boundary during the sintering process is presented. A qualitative appraisal of some current segregation and sintering theories are discussed. Surface analysis, by a variety of techniques, can probe the composition of interfaces within ZrOz-based ceramics giving valuable information of the distribution of impurities and components. A review of the contribution of surface analysis, particularly X-ray photoelectron spectroscopy, to the understanding of segregation phenomena and its effects on ceramic properties, is given. The review concentrates on Y203-ZrO2, CeO2-Y203-ZrO2and AI203-Y203ZrO2. The contribution of surface analysis to the understanding of low temperature degradation of yttria-tetragonal zirconia polycrystal is also examined.

1. INTRODUCTION ZrOz-based solid electrolytes are used in a broad range of technologies and offer the opportunity for improved process control and efficiency of energy conversion [1]. For example, oxygen sensors based on the Nernstian principle are already used in industrial processes, power generation and automobile emission to monitor and control gas environments. One important emerging application for solid electrolytes in general and Y203ZrO 2 materials in particular, is in the area of fuel cells where oxide solid electrolytes are now

184 seriously considered for application in small to medium sized power generation [2]. In other areas, ZrO2 has displayed activity as reforming [3] and oxidation catalysts [4], been explored as a substrate for semiconductor [5,6] and superconductor applications [7], been used as a thermal barrier coating on turbine blades [8] and displayed superplasticity. Interfaces of one form or another exist in all the materials used in the applications listed above. Examples of these interfaces include internal interfaces such as vain-vain interfaces and external ones such electrolyte-electrode or electrolyte-external environment interfaces It is the segregation of impurities, impurity phases and alloying components such as Y203, to these interfaces which is the prime concern of this article since the change in chemistry of the interface can significantly modify the ionic conducting and mechanical properties of the final product. Studies of segregation in single crystals can provide a good starting point for understanding segregation phenomena since segregation via lattice diffusion can be studied in isolation from the complicating effects of the grain boundary network. These types of studies are not merely of theoretical interest either, since single crystal Y203-ZrO2 has been used as a substrate for semiconductor and superconductor applications. For example, Si on Y203-ZrO2 has been investigated for radiation-hard high speed devices as a replacement for Si/sapphire technology [5,6]. Single crystals have also been used as substrates for the deposition of high Tc superconducting oxides, again focusing on device technology [7]. In both these applications the interaction between the Y-FSZ substrate and the material of interest is important. In the semiconductor application it is desirable to grow a thin SiO2 layer between the Si and the Y203-ZrO2 surface; the presence of segregated impurities at this interface may modify oxide growth and may even effect the Si semiconductor properties through doping. In superconductor applications, Si, which segregates to the external surface of Y-FSZ single crystals [9], quenches superconductivity in high Tc materials [10]. Polycrystalline sintered bodies are used in most applications, e.g. oxygen sensors [ 11], fuel cells and ceramic components. There are two classes of interfaces in polycrystalline materials which can have a significant influence on the macroscopic properties of the sintered material; these are the external and internal interfaces. The external interfaces are generally gas/solid or solid/solid interfaces. Gas/solid interfaces arise, for example, at the gas environment/ electrode interface when the Y203-ZrO~ is used as either an oxygen sensor or in a fuel cell. Examples of solid/solid interfaces include the electrode/electrolyte interfaces formed in fuel cells or oxygen sensors. These types of interfaces are the source of a rich variety of interfacial chemistry. Internal interfaces include the grain boundary network and other interfaces between different isomorphs of Y203-ZrO2 within the same grain e.g. antiphase boundaries, twin boundaries, i.e. extended defect sites within a grain. This presentation concentrates on segregation phenomena to the external surface and the grain boundary network. For Y203-ZrO2 ceramics both these types of interfaces exhibit some form of enrichment, notably of yttrium, cation impurities and anion vacancies. As demonstrated by Nowomy et al. [ 12] equilibration of the surface depends largely on the lattice diffusion of cations since the equilibration of vacancies is rapid by comparison. The driving force for Y enrichment is dominated by the electrostatic attraction of Y~ to the grain boundary [13] with ion size mismatch alone playing a lesser but significant role [9]. Y203ZrO 2 powders, however, invariably contain impurities or sintering aids added to achieve high density when sintering. Impurities (or additives) behave in a complex way and often form

185 glassy phases, particularly with yttrium, during sintering or subsequent heat treatments. The grain boundary network and the external surface of the ceramic are the preferred sites for the accumulation of these impurity phases. Segregation to these key interfaces within the ceramic can significantly alter both the electrical and mechanical properties of the ceramic. It is important, therefore, from a technological point of view, to understand interfacial phenomena in these systems. The details of the mechanisms whereby impurity phases are formed and migrate within the grain boundary network are not well understood. The lack of understanding is prirnarily due to a dearth of detailed information on the effect of ceramic processing parameters on the distribution of impurity phases. Without such information it is impossible to form predicative models of the sintering behaviour of Y203-ZrO2 ceramics. Characterization of ceramics has traditionally been by electron microscopy and X-ray diffraction among other techniques and has concentrated on microstructure. The determination of the chemistry in the region of interfaces is, however, extremely difficult with these techniques. On the other hand, surface analysis can provide a wealth of information about interfacial chemistry with comparative ease. A marriage between surface analysis, electron microscopy and other techniques provides the most powerful avenue for the characterization of ceramic systems. The prime aim of this article is to review the input of surface analysis, particularly electron spectroscopy, to the study of Y203-ZrO2 based ceramics. Electron spectroscopies such as X-ray photoelectron spectroscopy (XPS) and X-ray generated Auger spectroscopy (XAES) provide elemental identification quantitative analysis as well as chemical state information.

2 INTERFACES IN Z r O 2 CERAMICS

The two types of interfaces of concern here are the external surface and the grain boundary network. A schematic representation of these types of interfaces is given in Fig. 1. There are qualitative similarities between the two types of surfaces in enrichment of components and impurities and the development of impurity phases. However, there are also some fundamental differences, most notably, that the grain boundary is a confined region. In such a region, as pointed out by Clarke [14], attractive interactions between grains act to thin grain boundary phases. For thin grain boundary phases the grains are pushed apart by the structural disjoining pressure which has its origins in the epitaxial orientation of the grain boundary impurity phase on the surface of the grain which produces a type of steric hindrance. The equilibrium between these opposing forces changes with temperature. As the temperature increases, the viscosity of the impurity phases decreases, hence the thickness and the distribution of the grain boundary phase will change. In addition to grain boundary phases, there will also be a region near the surface of grains (both internal and external) where there will be a significant redistribution of cations and vacancies (Fig. 1). The cation distribution can be significantly enhanced in this region and is accompanied by a change in both the anion and vacancy concentrations. The development of an excess cation concentration at the grain surface may give rise to an electrostatic repulsion between grains, providing further stabilization of the grain boundary impurity phase. The origin of this redistribution lies in the ionic nature of Y203-ZrO2 and will be discussed in the

186 forthcoming sections.

Figure 1. Schematic representation of segregation within the grain boundary network and onto the external surface (Left). The thickness of the grain boundary impurity phase will be determined by the balance between the structural disjoining pressure and van der Waals' attraction between grains. There is a modified region near the surface of grains due to the segregation of cation impurities, vacancies and dissolution from the grain boundary impurity phase (Right).

2.1. Origin of impurity phases Most ZrO2-based ceramics contain some form of impurities. Impurities fall into two categories; those soluble in the lattice and those that are insoluble and likely to be introduced as processing impurities. In Fig. 2(a) the distribution of these two types of impurities is represented schematically. Soluble impurities can be either distributed throughout the bulk of the ZrO~ microcrystaUites or be present as other phases most likely adhered to the surface of the microcrystallites. Insoluble impurities, however, will probably be distributed throughout the powder as either surface or independent phases (salts or hydrated amorphous oxides) [ 15]. ZrO 2 forms solid solutions with many other oxides [ 16] hence there are a number of cations which can be present as soluble impurities within ZrO 2 ceramic powders. Some of the more common segregating impurities include Fe2+, Fe 3+, Ti4+ and Ca 2+. Other impurities may be introduced during processing or be deliberately added as, for example, a sintering aid. One of the most common impurities in ZrO 2 ceramics is SiO2 which has also been employed as a sintering aid. As an inadvertent impurity, SiO2, which is more likely in the form of silicate, probably arises as a residual impurity from the processing of ZrO 2 from zircon sands. Silicates are almost ubiquitous in ZxO2-based ceramics. Their morphology changes considerably during the sintering process from a thin film morphology coating the

187

/-

(a)

(b)

(c)

Figure 2. Schematic representation of the distribution of grain boundary impurity phases and surface enrichment. (a) Insoluble impurity phases (dark areas) may be adhered to or mixed with powder. Soluble impurities (represented by the dots) are dissolved into microcrystallites. (b) At intermediate sintering temperatures the impurity phases are distributed along the grain boundaries. Surface enrichment from bulk and/or from dissolution of the impurity phases may also occur. (c) At high sintering temperatures impurity phases move to lower energy sites such as triple points or the external surface.

grain boundary network to isolated pockets of glass within the body of the ceramic (Fig. 2(b)) [17]. Also represented in Fig. 2(b) is the surface enrichment of grains with impurity or alloying cations from either the bulk or from the dissolution of grain boundary impurity phases. Silicates have been observed in grain boundaries and at triple points of (Ca, Mg)-ZaO2 [ 18] and Y203-ZrO2 [ 19-22] by transmission electron microscopy (TEM). Silicates have also been observed on fracture surfaces of Y203-ZrO2, [17,23,24] and CeOz-Y203-ZrO2 using XPS [25]. These silicate phases invariably incorporate alloying components of the ceramic and other impurities within the system. Deliberate additions of SiO2 have been found to produce a significant improvement in the sinterability of Ca and Mg stabilized-ZrO2 [26]. These additions behave in a similar fashion to the impurity silica by forming silicate phases with alloying and impurity components of the ceramic.

2.2. Segregation Models Before discussing the process of sintering it is worthwhile examining segregation processes within individual grains. These processes are of extreme importance for sintering since they, to some extent, dictate the chemical composition of the surface of the grain. The chemistry of the grain surface plays a significant role in determining the type of sintering, i.e., liquid or solid state, and the mobility of the grain boundary, through impurity and solute drag mechanisms. Hence an understanding of segregation processes within individual grains will provide considerable insight into the sintering process discussed in the next section and may

188 one day lead to better engineering of the surface chemistry of ceramic powders. The description of equilibrium segregation phenomena in solids has been developed largely for binary alloys [27]. The surface cation mole ratio of solute i to solvent j in a two component system is described by [28]: l

nt

ns

-AG~r sz

(1)

where nx' and nBl are the equilibrium atom (mole or ion) fractions of the z~ (or j~) component in the interface and bulk phases respectively, AGsegis the free energy of segregation, T is the absolute temperature and R is the universal gas constant. Physically, AGses represents the flee energy associated with the exchange of a subsurface solute atom with a surface solvent atom and can be divided into entropy AS~g and enthalpy AH~g terms (AQ~g= AHs~g-TAS~g). The three contributions to the heat of segregation are:

AHs,g- AH, + AHw,, + AH,.

(2)

The interfacial energy contribution HI, is given by: An t

-

(V,

-

v elA

(3)

where 7i is the specific interfacial energy of species i (or j) and A is the interface area/atom. AH~ is generally small in the range 0 to :t21kJ/mol [27]. Hbt. is a binary interaction term and is given by: AHw, , - -

An.

(4)

where/~r/m is the heat of mixing of the binary solution and ~ is a geometric constant which depends on the crystallographic orientation of the surface. AHb~, is generally larger than AHx and typically falls in the range 0 to +65kJ/mol [27]. AH~ is the solute strain energy contribution which is given by: aa,

- -24,

- 0 2

(4Sr, + 310")

(5)

189 where K is the bulk modulus of the solute and S is the shear modulus of the solvent and r i and q are the ionic radii of the solute and solvent respectively. AH~ is zero when the solute ion is smaller than the solvent ion since no strain is generated in the lattice, but ranges from 0 to -150kJ/mol where the solute ion size is larger than the solvent ion. The solute strain term provides a driving force for yttria segregation in Y203-ZrO2 since the ionic radius of the Ya+ ion (0.893~) is larger than the Zr4+ ion (0.79~) resulting in AHseg of-5.3kJ/mol. According to Wynblatt and McCune the model described above does have some limited application to oxides [29]. They have suggested for oxides with no anion impurities and only homovalent cation impurities, that the nearest neighbour interactions may be replaced by the next nearest neighbour interactions i.e., by cation-cation interactions. Very few oxides however, have only one species of homovalent cation impurity hence the use of equation (1), on this basis alone, is in most cases not justified. Three major problems which exist with the determination of the free energies of segregation in oxides are: the co-incident segregation of a number of solute species, interaction between the species and the development of large space charge regions associated with defects and aliovalent cation impurities. Segregation of multiple species and the interactions between them have been dealt with in the theories of Guttmann [30] and Fowler [31,32]. The self-interaction of segregating solute species has been examined by Fowler who introduced an interaction term into the free energy in equation (1). Guttmann extended Fowler's theory by allowing for interactions between different solute species. Thus for a N component system the interface atomic fraction X/of species i:

x~expCAa, IR~3

x'-

N-1 '

(6)

E x:to (aa, /Rr - 1

1"1

where X~ is the bulk fraction of species i and AGi is given by:

A G i .. AGOg + ];[l'i(ll4l :~Xl :'X2 '"" XN)

(7)

and Wi(0qj,XI,X2,..,XN) is the solute interaction term with: (8) a0 " ZNo(~O - 0.SCe. + ~ ) ) Z is the coordination number, NO is Avogadro's number and eij is the bond energy. It is immediately apparent from equation (6) that a plot of the logarithm of the concentration against the reciprocal temperature will not directly yield the free energy of segregation. Hence the application of equation (1) to data obtained from multiple component systems is likely to lead to misleading results. Space charge regions in oxides develop as a result of Schottky and Frenkel disorder where there is an inequilavence in the formation energies between cation and anion vacancies [33,34]. In yttria-zirconia, anion vacancies are the predominant point defects, and their high mobility compared to the cations ya+ or Zr4*, ensures an uneven space charge distribution near

190 the source of anions such as the surface [35, 36]. In most practical situations aliovalent cations exist in greater number than point defects and thus dominate the nature of the space charge distribution near the surface. Hwang and Chen [13], for example, have calculated that temperatures greater than 2400~ would be required for the number of oxygen vacancies to dominate over a 1% cation dopant concentration.

tO

...,.

L_

Surface double layer

c 0) o to

region of negative charge

o

,, bulk

0 =

VO~ Yzr Distance from grain surface

Figure 3. Vacancy (V"o), Oxygen anion (O-) and Solute Y'z~distribution near the surface. The surface is positively charged due to the high mobility of anion vacancies and segregation of aloivalent cations. There is a region of subsurface enrichment of anions with a total charge counterbalancing the surface positive charge.

To account for the electrostatic component of the imbalanced electrostatic interaction of aliovalent cations, it has been proposed by several workers [27, 30, 37] that equation (2) should include an additional term, Hei. A more rigorous approach, however, involves the minimization of the free energy of the surface subject to the conditions of charge neutrality. The system is closed so there is also conservation of mass. The free energy G is defined as: L

G-

f~(~,n~) 0

+ 1/2p(x)#(x) - ]~c

(9)

191 where L is a distance into the surface characteristic of the bulk, ni and Fi are the number density and formation energy of the ith defect, p(x) is the charge density, ~(x) is the potential and Sc is the configuration.,d entropy. ~(x) is determined via Poisson's equation using p(x):

d2r dx 2

4u p(x)/e

(10)

p(x) is evaluated from the number densities hi(X), hence equations (9) and (10) must be solved self-consistently. The cation, anion and vacancy distribution near the surface of a grain is presented schematically in Fig. 3. A space charge region develops within the surface of Y203-ZrO2 due to the high mobility of anion vacancies (Vo) compared to cations. The high mobility of V"o means that they will concentrate near the surface producing a positively charged aniondepleted region [35]. To maintain charge neutrality, the subsurface region will be anion-rich, but this excess concentration of anions will decay to the bulk concentration at greater depths. Furthermore, segregation of yttria and other aliovalent cations with a charge less than four will increase the vacancy concentration via: (11) Most treatments of the development of surface space charge regions have concentrated on alkali halides [33, 34, 38], although Desu and Payne have studied the perovskite BaTiO3 [39, 40]. For Y203-ZrO2 the evaluation of the charge distribution near a surface is complicated by the lack of information on the formation energies of defects or the binding energy between dopants and point defects [36].

2.3. lnterfacial development during sintering Sintering is a process where a packed powder fuses to form a single dense body. There are two types of sintering processes of interest for sintering of Y203-ZrO2 powders; these are liquid and solid state sintering [41-44]. As stated in sect. 2.1, most ZrO2-based ceramics contain silicate impurities which readily form liquid phases in the grain boundary network at typical sintering temperatures (>900~ The process of sintering is driven by the requirement for the packed powder to reduce its total free energy, which, unless there are phase transformations, is dominated by surface free energy. This is achieved through densification and grain coarsening. Densification is defined as the removal of gas/solid interfaces and replacement with lower energy liquid/solid or solid/solid interfaces and ideally results in a sintered body without porosity (Fig. 4). Coarsening proceeds by grain growth where large grains consume smaller ones, thereby reducing the surface area of the system and hence the total surface energy, porosity may still exist after the coarsening process (Fig. 4). Segregation, which occurs at all stages of sintering, is a third process which minimizes the free energy of the system. During segregation the bulk and the surface of grains strive to reach a new equilibrium in response to changes at the grain boundary. The accumulation of

192

Sintering Mechanisms Densification

9169

OL_)

Coarsening Packed Powder Both

Figure 4. Sintering proceeds via densification and coarsening. During densification there is loss of porosity and during coarsening there is grain growth.

these impurities and solute at the grain interface can significantly inhibit grain growth during sintering by reducing grain boundary mobility. In many applications, e.g., superplasticity or moisture sensitive applications, small grain size is an advantage. Hence the correlation of grain growth kinetics with surface solute and impurity concentrations is of significant practical importance since it could lead to new methods of engineering ceramics. The free energy of a sintering system can be partitioned in a number of different ways. Consider the following partition:

A F - AFsu~

+

AF~

+

AFju~

+

AFeo~s

(12)

193 AFsve~ is the free energy related to the surface of the grains where the surface is treated as

a separate phase:

aFs0~

-

M

//

1

1

| I~o,~j+Z~jdn~

(13)

The first term in equation 13 describes the change of surface energy with incremental changes in the surface area dA. Each term in the sum represents the different contribution of different types (total M) of interfaces e.g. solid/gas, solid/solid or solid/liquid, to the surface energy. The second term is a summation of the product of the surface chemical potential Pr by the number n~ of moles of species i (total N) in the surface (~). la~'i may take on the following form in an ionic system [39]: I~ir

"

~

~

+RT/n

(a~) + Z e U |

(14)

where pim is the standard value, ai* is the activity of species i and U~ is the potential in the surface phase. The second term in equation 12, AFz~e, describes the free energy changes associated with the impurity phases present in the grain boundary network. The impurity phases in Y2Os-ZrOz ceramics are generally silicate in nature and may become liquid at low temperatures promoting liquid phase sintering. The major contributions to the free energy of the glass phases will be from their surface energy and chemical potential terms, hence: P

AGao.

I~oaa'dA. + ~l~gaa'n,

(15)

where all the symbols have the same meaning as before but in this instance apply to the impurity phase. There are P types of impurity phase/impurity phase or impurity phase/ceramic interfaces and Q types of impurity phases. It has been assumed that the glassy phase is incompressible hence no PdV term has been included in equation (15). The third term in equation (12) for the free energy of the bulk would have the same format as equation (15) except the superscripts "IMP" are replaced by "BULK". The surface energy term is included for the bulk to account for the interface between the surface phase and the bulk phase. Finally the contribution of pores varies considerably during sintering from around 40% of the total volume in the pressed powder to less than 2% in the fully dense state. The free energy of the pores can be defined as: s

A GeoaE - P d V ~

+ ~ l~ ~ni ii

(16)

194

) Adhesion (a) Highly Dispersed forms of Silicon, e.g. Silicates, Silica

(b)

Silicates move to grain boundary at newly formed neck. Other impurities also migrate to neck (see Table 2) densification may occur if centre to centre distance is reduced

C T') J

r

f (c) Grain boundary migration, reduction of surface area and accumulation of impurity phase (coarsening & densification)

(d)

Wetting of grain boundary surface at melting point of silicate. Silicate film may be continuous or discontinuous depending on wettability and quantity

(e)

Stationary or nearly stationary network of grain boundaries

A

9Mass transport of material along grain boundaries

B

9 At higher temperature lattice diffusion is also significant. Material may migrate away from the grain boundaries or to the grain boundaries

Figure 5. The development and distribution of silicate impurity phases during sintering and post-sinter annealing.

195 The rate of approach towards equilibrium will be determined by the kinetics of the system. The diffusion rates of impurities and impurity phases, vacancies, the annihilation rate of surface vacancies with the atoms of the sintering atmosphere and the grain boundary mobility will all have a role to play. The primary concern here is a qualitative description of the redistribution of matter within the sintering system. The first stage of the sintering process is neck formation between microcrystallites within the pressed powder. On the atomic scale, the contact points between microcrystallites within a pressed powder are lower energy sites than other parts of the surface. In the unsintered powder these contact points may arise from adhesion due to the presence of residual salts, hydroxides, impurities such as silicates [ 15] or van der Waals attraction between grains (Fig. 5(a)). Presumably, during the initial stages of sintering surface hydroxides are eliminated being replaced by M-O-M bonds, where the M's may be surface impurities, solute or solvent cations. At higher temperatures neck formation can be promoted by the presence of a liquid phase wetting of the grain surface since there is an adhesion pressure exerted by the liquid on the grains which draws them together. The origin of this adhesion force is the lower pressure within the liquid phase compared to adjacent regions of the interface containing only vapour [44]. Neck formation may result in densification if the centre to centre distance of necked grains is reduced (Fig. 5(b)). Matter transport to the neck region can occur via a number of different mechanisms as listed in Table 1 and depicted in Fig. 5(b). Curvature differences on the surface of microcrystallites result in chemical potential and vacancy gradients which drive matter transport [41]. For example, the Gibbs-Thomson equation relates the vacancy concentration under a curved surface, C(r), to the surface energy "t and the radius of curvature r : C(r) - C..exp{2u } rkT

(17)

TABLE 1 Mechanisms for Matter Redistribution During Sintering. Mechanism Liquid Phase Transport Surface Diffusion Vapour Diffusion Bulk Diffusion

where k and T have their usual meaning, 6"_ is the defect concentration in the bulk and f~ is the volume occupied by the vacancy. During sintering there will be a net flow of vacancies, described by Fick's law, from the neck to areas of large curvature which will be balanced by a flow of atoms in the opposite direction as depicted in Fig. 5(c). An alternative picture is that the surface tension at regions of the surface with large curvature is greater than at small

196 curvature thus providing a driving force for the redistribution of material. In Y203-ZrO2,the ever-present impurity silicate phases provide the dominant mechanism of matter redistribution. Once the neck is larger than the cross-sectional area of the smallest grain then it is free to migrate through the grain. It is this stage of sintering that is the most complex to describe since both grain coarsening and densification occur concurrently by the mechanisms described in Table 1 [45]. Pore networks are gradually eliminated until high densities are obtained. Furthermore, the intermediate stage of sintering is characterized by clustering between interconnected networks of micro-crystallites and the development of true grain boundary networks within these clusters. At the end of this stage of sintering (not easily defined) the clusters and networks have coalesced to a dense body which is approaching its maximum density. During the later stages of sintering grain growth behaviour can be described by equations of the general form [ 13, 46, 47]:

dn _ ~ n .

K ( t - to)

(18)

where d and do are the initial and actual grain size after time to and t respectively. K is a temperature dependent rate constant given by: 0

g-

re -~

and Ko is proportional to the grain boundary mobility [13, 46].

(19)

n is a constant which

TABLE 2 Grain Growth Control Mechanisms 1. Mechanism Anomalous Grain Growth Normal Grain Growth Impurity Drag (High Solubility) Impurity Drag (Low Solubility) Diffusion Through Continuous Second Phase i In he absence of pores or a second phase.

designates the grain growth mechanism (Table 2). It can be seen in Table 2 that the value of n is not unique for a particular type of grain growth mechanism. For example in the absence of pores or a second phase, there are two types of mechanisms which give rise to values of two for n. Surface analysis of fracture surfaces provides a method of distinguishing between

197 these two mechanisms by determining whether there is any surface segregation of high solubility impurities or solute. Similarly, for n=3 surface analysis can distinguish between impurity silicate phases and grain enrichment by depth profiling through the fracture surface. As will be seen in section 4.3.1., values for n of three are frequently observed for Y203-ZrO2 and surface analysis indicates that both impurity phases and surface segregation contribute to retarded grain growth in this system, although the impurity phase may only have a significant role at temperatures where it wets the grain boundary network. Grain boundary mobility (or velocity) can be modified by a number of mechanisms including pinning due to a second phase, pinning or drag due to the presence of pores or drag due to the accumulation of impurities, components and/or defects at the boundary. In Y-TZP there is sufficient evidence to suggest that grain growth oft proceeds via a solute drag mechanism where the accumulation of yttria and impurities impedes the boundary mobility [ 13, 47-50]. In Y-FSZ, however, normal grain growth is observed since the relative difference between the surface and bulk yttria content is much smaller than is the case for Y-TZP. The retardation of grain growth by diffusion control through grain boundary silicate phases would appear to be minimal since parabolic grain growth is normal in Y-FSZ, where impurity phases are known to form [17, 19, 23]. This suggests that either the diffusion coefficients of cations through the silicate phases are generally high or that the grain boundary phases rarely form continuous films. The latter is more likely, since the impurity phases change in morphology with temperature and there may be only narrow temperature regimes where they form a continuous film in the grain boundary network. Pinning of grains boundaries is also known in Y203-ZrO2, where A1203 additions result in smaller average grain size. It seems likely that pinning may also occur due to the presence of insoluble or high melting point impurity phases at triple points. The role of pore drag is much less than solute drag in Y203-ZrO2. Pores are rarely observed and in Y-'IT_~ and while pore density is much higher in Y-FSZ than Y-TZP, there is no evidence to suggest that pores impede grain growth since normal grain growth is observed in the cubic phase field. Under conditions of impurity drag control the rate of grain growth dG/dt is defined by: dt where M b is the grain boundary mobility, AP is the pressure gradient across a curved grain boundary and pOe,at is the dragging force on the boundary. AP is defined by [51]"

AP-

2--7 P

(21)

198 where y is the interface energy and p is the radius of curvature, pORao has been defined by Cahn as [52]" -t-w

e

-

ae'ax

(22)

&

where Nv is the number of atoms/unit volume, C(x) and E(x) are the impurity or solute concentration and interaction energy respectively at a distance x from the grain boundary. CB~, is the bulk concentration of the impurity or solute component. The actual form of the impurity drag term evaluated from equation 22 will depend on the velocity of the grain boundary with respect to the diffusion coefficient of the species that may reside at the grain boundary [53]. Equation 22 is usually evaluated in the limiting cases of high and low grain boundary velocities [52, 54]. Generally, the fastest diffusing species will exert a drag on a high velocity grain boundary because they are the only species able to keep pace with the boundary. For low velocity boundaries the slow diffusing species exert drag because the faster moving species are can easily redistribute [42].

C(x)

Co

~ 2 - z ~

_ _. C Bulk

x

Bulk Grain Boundary Figure 6. Hypothetical enrichment profile, C(x), of yttrium near the surface of a grain. The area under the curve can be approximated by a rectangle of with ~5and height Co-Cam where CBun, is the bulk concentration.

2.4. Grain boundaries in the fully dense state Grain boundary drag would appear to be the most significant cause of grain growth retardation in Y203-ZrO2 ceramics of commercial importance. Studies of yttria tetragonal

199 polycrystal (Y-TZP) and doped CeO2-Y203-ZrO2 indicate that a major role is played by impurity and yttria accumulation at the migrating grain interface. From equation 22 it can be seen that the dragging force will increase with either an increase in the concentration at the grain interface or an increase in the interaction energy of impurities or yttrium with the grain boundary. In the former case Theunissen proposed a model of grain boundary enrichment of yttria which could be invoked to qualitatively describe the slower grain growth kinetics in the two phase field regime. Grain boundary enrichment as depicted in Fig. 6 is proportional to the area under depth profile which can be calculated if C(x) is known exactly or approximated by a rectangle of width ~5and height Co. Using this approximation, Theunissen found, using AES and XPS, that Co-Cbu,, was larger for the tetragonal phase hence the dragging force was larger and consequently, the grain size smaller. This same conclusion could be inferred from earlier work of Theunissen et al. where Y203 reached the same equilibrium level at the external surface regardless of the bulk concentration suggesting that the enrichment level in the tetragonal phase was larger and the drag greater [55]. Hwang and Chen clearly defined the role of impurities in promoting grain boundary drag in Y203-ZrO2 and CeO2-Y203-ZrO2 [13]. They found that the grain size increased from small to largest with the following sequence of solutes: Ca~+ < Mg2+
3. SOME COMMENTS ON SURFACE ANALYTICAL TECHNIQUES.

3.1. Core level binding energy shifts A number of excellent review articles and books cover the topic of surface analysis in general [56-58], and its application to ZrO2-based ceramics [59]. Generally, XPS has only been used to give quantitative analysis of interfacial segregation; very little attention has been paid to the spectroscopy which, through binding energy shifts in the photoelectron peaks, can provide information on the chemical environment and insight into surface defect structures. There are, however, some notable exceptions. For example, XPS and X-ray Auger electron spectroscopy (XAES) have used extensively by Hughes et al. [ 17, 23, 60] and Ingo et al. [25, 61] to give information about impurity silicate phases in ZrO2-based ceramics through the silicon spectroscopy. Hughes and Badwal have used the Y 3d photoelectron peak to examine glass phases at the external surface of Y-TZP. Ingo et al. have determined Ce 3+ content of various CeO2-Y203-ZrO2's and Mercendalli et al. [5] have used the Zr 3d photoelectron peaks to determine the degree of interaction between Si and Y-FSZ. Ideally, it would be desirable to use the Zr 3d or Y 3d binding energy shifts to identify different polymorphs of ZrO2 or to establish the yttria content of a surface. As can be seen in Fig. 7, however, the Zr 3d photoelectron line displays very little sensitivity to changes in solid solution concentration. Studies by Leonov et al. [62] and Strekalovskii et al. [63] indicated that there was no significant change in the Zr 3d 5/2 binding energy with increasing

200

183.0

o3 i(1) IE LU (::: 9

"ro

(

182.5

.,,.m

rn

"o L_

N 182.0

181.5

I I I I I 1 11 0 5

I I I I 11 10

11 I I I I 1 15 20

mol % Y203

Figure 7. Zr 3d5/2 binding energy versus mol%Y203. (ll)-Leonov et al. [62], (A)-Strekalovskii et al. [63], (0) Morant et al. [64], (A) De Gonzalez et al. [65] and (r'l)-Hughes unpublished data.

Y203 content, although reliminary results by Hughes suggest that the dependence may be stronger than previously thought. (Differences in the absolute values between the results of Leonov and Strekalovskii, probably arise due to spectrometer calibration and the choice of internal binding energy reference.) From Fig. 7 it also appears that the Zr 3d binding energy is lower in solid solution with Y203 than the pure oxide grown on Zr foil [64, 65] although it may be similar to m-ZrO2 [66]. These differences in the Zr 3d binding energies of either the pure oxides or the solid solutions may be due to the presence of different polymorphs in the powders. To date no systematic study has been made of the variation of the photoelectron binding energies with yttria content in conjunction with XRD measurements to determine the polymorphs present in the powders. Insensitivity on the yttria content is also displayed by other spectroscopic parameters such as binding energy splittings between photoelectron lines, e.g. the O ls - Zr 3d binding energy splitting is 347 + 0.1 eV over a broad range of Y203 contents. Furthermore, loss structure on

201

Zr 3d loss structure A

m

I

0

I

15.0

I

300

1

45.0

i

60.0

AE

Figure 8. Loss structure to the high binding energy side of the Zr 3d photoelectron peaks. A 2p ~ 3s ligand excitation, B - Bulk Plasmon and C - Zr 4d ~ Zr 4p transition (conduction band).

the photoelectron lines where the escaping photoelectron interacts with the valence band also appears to be insensitive to structure. Marietta and Pignataro [67] have found that the loss line structure labelled A, B and C are due to the excitation of ligand np electrons in the valence band to (n+l)s unoccupied state in the conduction band, bulk plasmon and excitation of a Zr 4d --~ Zr 4p level, respectively. However, despite the contribution of the valence and conduction band electrons to these process there appears to be little difference between the loss structure of cubic Y-FSZ (Fig. 8(a)) compared to the m-ZrO 2 (Fig. 8(b)) except for a small shift in the oxygen 2p ~ 3s transition. While the dependence of the photoelectron and Auger peaks on Y203 content is, at best, weak, changes in peak position with the different polymorphs are more definitive. The combination of the Zr 3d photoelectron and the Zr MNV Auger line as depicted in Fig. 9, forms a 2-dimensional chemical state plot, which can provide an even better method of polymorph identification. (The Zr MNV Auger transition involves valence levels which are sensitive to bonding configurations and hence display a shift with changes in environment.) For example, the monoclinic phase of ZrO2, either pure or in solid solution with Y203 can be distinguished from the tetragonal and cubic phases on the basis of its modified t~' parameter (diagonal lines in Fig. 9).

3.2. Defect structures The stabilization of the high temperature phases of ZrO2 by alloying with other oxides where the cations are aliovalent with respect to Zr4+ results hi the production of oxygen

202

329.1 147.1 v

j J

L,.

L_

(D

11)

rW ~ r-

s

146.6

328.6

E

t~ n t..

i..

,a:

<

328.1 Z ~

(D

146.1

m-ZrO 2

-o r . 0

"1o o

N

327.6 145.6 182.4

182.2

182.0

Zr3d Binding Energy (eV)

Figure 9. Zr 2-dimensional chemical state plot. Tetragonal zirconia (t-ZrO 2) can, in principle, be distinguished from cubic (Y-FSZ) and monoclinic zirconia (m-ZrO2) [66].

vacancies. The distribution of both the oxygen anion vacancies (Vo) and the y3+ cations has been the subject of considerable speculation and investigation. Early neutron [68] and X-ray diffraction [69] studies suggested that the stabilizing cations (y3+ or Ca 2+) were nearest neighbours to the anion vacancies. Recent extended X-ray absorption f'me edge (EXAFS) [70] and electron spin resonance (ESR) [71, 72] studies have revealed that a significant population of anion vacancies are more likely associated with zirconium cation sites rather than the yttrium cations. In EXAFS studies of 18wt% Y203-ZrO2, Catlow et al. [70] found that the coordination number was less than 8 for Zr but equal to eight for Y. They concluded that up to 48% of the Zr cations may be heptacoordinated in their material. Furthermore, Orera et a1.[72] have shown using ESR, that while holes are trapped at oxygen anions adjacent to Y cations, the most dominant intrinsic defects are associated with anion vacancies adjacent to Zr cations sites. These T (VoZr'Vo) and C defects (Zr'Vo) with hexa- and hepta-coordinated Zr cations are depicted in Fig. 10. XRD measurements suggest relaxation of the oxygen ions adjacent towards the anion vacancies [69]. There was no evidence of adjacent Y cations and anion vacancies. Both the T and C defects are effective electron traps and are easily detected in photoemission experiments. Indeed, it is impossible not to detect them since low energy photoelectrons become trapped at these defect sites giving rise to low binding energy broadening on the Zr 3d photoelectron lines typical of Zr ~ (Fig. 11). No systematic XPS

203

(a)

(b)

~ ~ ~

9

o-Jto/I (3"

(3" C-defect

o

o

(3"

I~ T-defect

9

- Zr 4 +

9 -O

Figure 10. The two types of defect structure in yttria-zirconia which may give rise to a Zfl § signal in the Zr 3d spectrum. Zrz~-Vo or ZrzcVo-Zrz~ are efficient electron traps which effectively trap low energy photoelectrons.

studies have yet correlated the Zr spectral features associated with the T and C defects with the ESR observations. This correlation could be of significant importance, particularly in studying the tetragonal to monoclinic phase transition since the T and C defects are sites with lower coordination Zr cations and hence could be nucleation sites for phase transformation. For example, relaxtion around the oxygen anion vacancy produces a trigonal symmetry similar to the hepta-co-ordinated Zr cation structure in monoclinic zirconia. Indeed, Hughes et al. [73] have found considerable changes in the %Zr"~ as a function of treatment time of Y-TZP in moist air at 200~ there appeared to be some correlation with the development of the monoclinic phase (see sect. 8).

4. Y203-ZrO 2

4.1. Phase diagram In order to effectively describe segregation in Y203-ZrO2 systems, it is necessary to examine the Y203-ZrO2 phase diagram, which delineates a number of different phases as a function of composition and temperature (Fig. 12) [74] ZrO2, which at ambient is in the monoclinic form, is transformed to tetragonal structure at 1170~ The tetragonal phase transforms to the cubic structure at 2370~ All these phase transformations are reversible on cooling. On addition of about 8.5mo1% Y203 the cubic phase can be stabilized at room temperature [75] Compositions with about 2-2.5mo1% Y203have tetragonal symmetry. This

204

-

Zr 3d

6 r O

o

5

O O I--

>,

4

t-

a}

3

t-Zr 3+

.....

.. ~I

9

,

..... 9 ,

..

9

-....

::.:

,:.

-..~.._ . : . . . . . . . - : . . . ~ . ...............

172.0

9

1

180.0

176.0

i 84.0

I

188.0

,

I

192.0

Binding Energy (eV)

Figure 11. Zr 3d photoelectron spectrum displaying low binding energy broadening, indicating the presence of Zr3§ associated with photoelectrons trapped at the defect structures depicted in Fig. 10. 2500

vA P2000

- T

C

1500

~1000 E 500 0

0

i 2

4

i 6

Y203 Concentration

I 8

10

(mole%)

Figure 12. Yttria-Zirconia phase diagram after Scott [74]. T = tetragonal, M = monoclinic and C = Cubic phases.

205 phase transforms under stress to monoclinic zirconia providing the toughening mechanism utilized in yttria tetragonal zirconia polycrystal (Y-TZP) ceramics. All compositions between 2 and 8.5mo1% Y203 are usually a mixture of cubic and tetragonal (t) phases above 600~ and t' and t phases below 600~ Extended thermal treatment leads to solute partitioning into an Y203-poor tetragonal phase and an Y203-rich cubic phase (t' at room temperature) [76]. The redistribution of material is an extremely complex process during this phase of sintering. The presence of glassy impurity phases is generally considered to enhance solute partitioning by assisting mass transport of matter from one phase to another. Yttria fully-stabilized zirconia (Y-FSZ) is formed at higher yttria additions (>8.5mo1% Y~O3) where the cubic phase is fully stabilized at room temperature. The study of segregation phenomena is, in principal, much simpler in Y-FSZ since there is no solute partitioning.

0,4--

(a)

1 100~

.o al

n.o 0.31 E 0

0.2

co 0.1

I (b)

N

0.2

0

100

150

200

Annealing Time (minutes)

Figure 13. (a) S i / ~ and (b) Y/Zr XPS atomic ratios versus annealing time at 1100~ @) - Ar or (v, O) - air atmospheres. (v,v); 0 = 70 ~ and (@,o); 0 = 0 ~

(v,

206

4.2. Single crystal cubic-stabilized Y203-Zr02 Segregation studies of single crystals have an impact on three main areas:- (i) the fundamental science of segregation, (ii) technologies which utilize single crystals and (iii) their implication for grain growth mechanisms in polycrystalline systems.

0.3

-

a

"'

( )

.~""""--"~-~-

1100~

o 0.2 I~ 0.1 Z

/,e,

.

~ ~ ...e

I,,

I

i

50

lOO

15o

o12.0

04

0

a: 8.0 0

4.0

o

I

Annealing Time (minutes)

2200

Figure 14. (a) Na/Zr and (b) Fe/Zr XPS atomic ratios versus annealing time at 1100~ 0 ) - Ar or (v, O) - air atmospheres. (-,v); 0 = 70 ~ and (O,o); 0 = 0 ~

(-,

In single crystals segregation ideally occurs by lattice diffusion. Consequently studies of segregation in single crystals provide an opportunity to investigate this particular segregation mechanism in isolation from the grain boundary network. These studies provide an important base for developing theoretical models of the mechanisms of segregation. There are two aspects of segregation which need to be considered in real oxide systems. First, there is competitive segregation of species which should be considered in the framework of sophisticated models such as Guttman's theory for multicomponent systems [30]. Second, in ionic systems such as oxides, segregation of cations with a charge different to the valency of

207 cations in the matrix, i.e. aliovalent cations, generally results in the development of a considerable space charge region in the vicinity of the surface [77]. Hughes studied the segregation of impurities and yttrium in single crystals of yttria fullystabilized zirconia (Y-FSZ) annealed under oxygen and argon atmospheres [9]. Segregation of yttrium and impurities such as Si, Na and Fe, were observed in both atmospheres after annealing in the temperature range 900 to 1500~ For example, during annealing in air at 1100~ the S ~ ratio increased for the fh-st 150 minutes then remained constant for longer times (Fig. 13). The angle dependence indicated that the Si was concentrated in a surface layer since the Si/Zr ratios were much larger for 0=70 ~ The dependence of the S ~ ratio on annealing time at 1100~ in Ar was similar to that of specimens annealed in air. During annealing in air the Y/Zr ratios also increased to equilibrium values of 0.28 (0 = 0~ and 0.32 (0=70 ~ which was attained at much shorter times than for Si segregation. Furthermore, the Y/Zr ratios were larger at 0=70 ~ than 0= 0 ~ indicating the presence of a distinct surface layer. During annealing in Ar the Y/Zr ratios displayed an initial increase, peaked between 12 and 70 minutes annealing, then decreased to equilibrium values for annealing times longer than 100 minutes. For short annealing times the Y/Zr ratios were similar at 0=70 ~ and 0= 0 ~ indicating a homogenous distribution of Y throughout the volume analysed by XPS. However, at longer annealing times the Y/Zr ratios were less at 0=70 ~ than at 0---0~ suggesting that the external surface was yttrium-poor. It is worth noting that the degree of yttrium surface enrichment observed for the single crystal studies is similar that observed by Theunissen et al. [24] in polycrystalline systems suggesting that in polycrystalline systems enrichment by lattice diffusion is an important process. The Na/Zr and Fe/Zr ratios for the crystals annealed at 1100~ are displayed in Fig. 14. For air annealing the N ~ and Fe/Zr ratios increased significantly within the first twelve minutes of annealing, but at times longer than 30 minutes both ratios appeared to have reached equilibrium. The large difference in magnitude between 0 = 0 ~ and 70 ~ is evidence that the iron and sodium have segregated onto the external surface of the crystal forming a mixed layer with silicon and yttrium. Under Ar annealing the Na/Zr ratios displayed peaks at around 25 minutes annealing then decreased to equilibrium levels which were approximately the same for both angles suggesting a homogenous mixture of Na throughout the surface volume examined by XPS. No Fe was observed to segregate to the external surface. It is clear that even for single crystals of Y-FSZ, there is concurrent segregation of a number of species. Furthermore, the interaction between these species and the external environment, such as the annealing atmosphere, may enhance or suppress segregation of any of the species. The interaction between species is manifested, for example, in the complicated interplay between the Si, Y, Na and Fe segregation behaviour. A qualitative summary of the depth distribution is given in Fig. 15. In both atmospheres species appeared to segregate strongly at short annealing times which for air-annealing led to large equilibrium values. For Ar-annealing, however, further annealing led to desegregation of all species except for silicon. These changes suggest that compositional modifications occurring at the surface as a function of time changed the free energy conditions for segregation. It is evident for both sintering atmospheres that as the annealing time increases the depth of the surface modified region also increases to the point where, at equilibrium, the depth is at least 10nm. Furthermore, the segregation of Y, Na and Fe appears to be dependent on the sintering atmosphere, whereas silicon is not. This indicates

208

TIME

150 mins

45 mins

Air

o.31~ -

y

O.3

y

0.2

0.2~Na -

0.1 0

:,= 0.3 n" .o_ 0.2

E o0.1

7(1

Fe

i

|

=

7O

d(nm)

d(nm)

>I f

<

0.3

J

si

0.2

o.11

0"2~~Si

0.1 7O

d(nm)

Y

Na

70

d(nm)

Figure 15. Segregation in air and Ar annealed single crystals of 9.5mol%Y203-ZrO2. Yttrium/Zr and impurity/Zr atomic ratios as depth profiles after annealing at 1100~ for the stated times.

that silicon may well segregate via a different mechanism, possibly along defect sites, such as low-angle grain boundaries given that the solubility of silicon in YzO3-ZxOz is generally accepted as being low [78]. These striking differences in segregation behaviour under the different annealing atmospheres are most likely related to the availability of oxygen during annealing and how the surface charge and space charge region responds to changes in the oxygen concentration. The strong segregation behaviour of Na at short annealing times is probably related to its large ion size (0.97) and effective charge of minus three. However, at longer annealing times in air it may combine with surface oxygen and volatilise from the surface. The loss of Na appeared to initiate desegregation of Y since the Y/Zr levels decreased after 100 minutes annealing. (Yttria is not volatile under these annealing conditions hence decreases in the Y/Zr ratio represent desegregation from the surface rather than volatilisation.) The case for Fe is not so clear. Fe3+(0.64A) and FeE+(0.74A) both have smaller ionic radii than Zr4+ (0.79~) and might therefore be expected to be attracted to the grain surface only on the basis of their effective charge. However, no iron segregation was observed during Ar-annealing, where Fe 2§ with the largest effective charge, might be expected to be

209 the dominant iron cation. These results have important ramifications for electrode response and sintering behaviour in different atmospheres. For example, Nowomy et al. [ 12] observed that the work function in Zr0.s2Y0.1802_s took up to 150h to equilibrate after a change of annealing atmosphere from reducing (low 02 partial pressure) to oxidizing (high O2 partial pressure) conditions. AES measurements indicated that this slow response was due to the surface segregation of AI and Ca, cations with slow diffusion rates compared to oxygen anions. From the perspective of sintering, impurity drag during grain growth, which limits grain size through the surface segregation of aliovalent cations, may in principle be controlled by a judicious choice of impurities and sintering atmosphere (see sect. 4.4.1). From a technological viewpoint, single crystal Y-FSZ has been proposed as a substrate for silicon on insulator (SOI) devices and also for superconductors. In the fabrication of SOI devices one major problem is the density of defects at the Si/Y-FSZ interface. One method to remove this defective region is by oxidation of the interfacial Si so that ultimately the Si/Y-FSZ interface can be replaced with the Si/SiO 2 interface [5]. Loebs et al. found that oxidation occurred readily, but in an uncontrolled fashion when submicron films of Si were deposited by vacuum sublimation onto Y-FSZ heated to 900~ [6]. Mercandalli et al. used electrochemical techniques to provide a higher concentration of oxygen at the Si/Y-FSZ interface and thus eliminate problem associated with the reduction of the Y-FSZ via a redox reaction [5]. They concluded that the rate limiting step was the transport of oxygen across the Y-FSZ/SiO 2 interface which could be significantly effected by segregation phenomena. While the kinetics of impurity and yttrium segregation at 900~ are slow, Mercandalli has described post deposition annealing treatment of 1200~ for times up to 2 hours where segregation is rapid. Hence the accumulation of impurities and yttrium is likely to have a significant effect on the transport kinetics of oxygen in the surface of the Y-FSZ single crystal. Pellet et al. [79, 80] inverted the experiment described above by ion beam sputtering Y-FSZ onto single crystal Si<100> substrates. This method eliminates the problem of oxidizing the Si film at the Si/Y-FSZ interface since a SiO 2 layer can be grown on the silicon before deposition. In both the work of Pellet et al. and Mercandalli et al. no infusion of Si into YFSZ was observed which may provide some evidence for the insolubility of silicon in yttriazirconia. Segregation of impurities such as Fe and Si are also a problem when Y-FSZ single crystals are used as substrates for high Tc superconductors. Silicon has been reported to destroy the superconducting state [10]. In the deposition of superconductors, however, different preparation techniques may have a marked influence on the degree of segregation. Kawasaki et al., deposited superconductor material by sputtering from sintered discs onto Y-FSZ held at temperatures between 580 to 700~ Segregation phenomena are not expected to be a problem in this temperature range.

4.3. Polycrystalline fully-stabilized Y203-ZF02 Polycrystalline yttria fully-stabilized Y203-ZrO2 (Y-FSZ) finds application mostly in the area of solid electrolytes for oxygen sensors and fuel cells, although partially stabilized Y203ZrO2 also competes in the fuel cell area because of its greater mechanical stability [11]. Oxygen ion conductivity in these solid electrolytes basically consists of two contributions, namely that due the lattice conductivity, which depends on the bulk properties and that due

210 to interfaces, primarily the grain boundarynetwork, which depends on interfacial properties. The presence of the grain boundary network itself, and impurity phases which reside in the network both effect the grain boundary conductivity. A capacitance is developed across a clean grain boundary due simply to lattice mismatch (intrinsic effect) [19]. In addition, in most Y203-ZrO2 systems distinct intergranular impurity phases develop between adjacent grains which provide an impedance to oxygen ion transport (extrinsic effect). In high temperature applications the contribution of the grain boundary resistivity to total resistivity is small compared to the lattice resistivity. In low temperature applications, where oxygen sensor development has been heading, the contribution of the grain boundary resistivity to total resistivity becomes increasingly important. Verkerk et al. [81] demonstrated that the grain boundary conductivity depended on the grain size and rate of cooling from the sintering temperature. These results were proven to be caused by compositional differences between the surface and the bulk of grains. Hence segregation of Yttrium and impurities in polycrystaUine Y-FSZ is of interest because it can modify grain boundary conductivity. For this reason many studies of Y-FSZ have sought to relate the conducting properties to the composition of the grain boundary [17, 19, 82-85]. Of equal importance is the influence of segregating species on grain growth kinetics which will be addressed in the next section. In Y203-ZrO2 Burggraaf et al. have found that the grain boundary conductivity of Y-FSZ to be considerably lower than the bulk [81, 86, 87]. Auger electron spectroscopic (AES) studies of the intergranular fracture surfaces found that the grain boundary experienced considerable enrichment of yttrium. For example a sample 17mol%YO1.5, designated ZY17, had a grain boundary enriched in yttrium to ZY24.5 after 3 hours sintering at 1456K. Ar+ depth profiles indicated that the yttrium levels decreased to the bulk level with depth from the surface although, for experimental reasons no quantitative assessment of the depth could be made. Based on studies of the conductivity of bulk ZY24.5 [87-89] they concluded that the lower surface conductivity in ZY 17 was consistent with the conductivity of the yttriumenriched surface (ZY24.5) rather than the bulk composition. Further studies by Burggraaf, Winnubst and others indicated that the segregation observed in ZY17 was common to Y203-Za'O2 materials with a wide range of compositions from 4 up to 32 mol% Y203 [55]. They found on the external and fracture surface of yttria-zirconia, that the segregated layer was approximately 2-4nm thick and only developed after extended periods of annealing at temperatures as low as 600~ [55, 90]. It is evident from these studies that surface enrichment of yttrium is to be expected in Y203-ZrO2. As seen in the previous section it is probably driven by the electrostatic attraction of yttrium to the grain boundary as well as ion-size mismatch. In detailed studies of the influence of sintering atmosphere on the distribution of impurity phases in Y-FSZ, Hughes and Badwal found a strong correlation between XPS atomic ratios on the fracture surface and the grain boundary impedance which were monitored as a function of sintering temperature. Under Ar-sintering, the XPS atomic ratios for Y/Zr and Si/Zr 2displayed a single peak at 1300~ as a function of temperature (Fig. 16(a-b)). This peak corresponded to an inflection point in the grain boundary resistivity (Fig. 16(c)) and was attributed to a grain boundary impurity silicate phase spread throughout the grain boundary network at 1300~ At lower temperatures the silicate phase was too viscous to wet the grain boundary network and at temperatures above 1300~ it began to dewet the grain boundary network and migrated to the external surface as evidenced by increases in the XPS atomic

211 ratios on the external surface (Fig. 17) and decreases on the fracture surface (Fig. 16(a-b)).

I I

o

0.32 (-

".G E Z o "E o

0.30

..=

r fit.

X

a) Y/Zr

Ar

0.36 _

--

0.34

0.2 0 . ~ :~7 b) Si/Zr 0.15~ /'NN 0.1 I--

/

~e

o ~- 0.32 E 7 0.3

__11

b) Si/Zr E o ~ 0.3

oo 5 ~ " 2 00.~~,~c) E

--rr"

o,oooF

ot

a) YIZr

~ =

~

x

0.25

~.~A~

10,00

1300

1400 Temperature

1500

1600

1700

(~

Figure 16. (a) Y/Zr, (b) Si/Zr atomic ratios for fracture surfaces and (c) grain boundary resistivity versus sintering temperature. Specimens were sintered for 4h in Argon.

1400

1500

1600

Temperature ~

Figure 17. (a) Y/Zr and (b) Si/Zr atomic ratios from the external surface versus sintering temperature. Specimens were sintered for 4h in Argon.

Sintering in air produced two peaks in the XPS atomic ratios for Y/Zr and Si/Zr on the fracture surface (Fig. 18(a-b)) which again correlated with the grain boundary impedance (Fig. 18(c)). In this case, the ceramic attained maximum density between 1200 and 1300~ which correlated with maxima in the XPS ratios on the fractured surface suggesting that the grain boundary impurity phase wetted the grain boundary network in this temperature range and furthermore promoted densification. Between 1350 and 1450~ both the XPS atomic ratios and the grain boundary impedance decreased compared to 1300~ This decrease is contrary to the trend that might be expected given grain growth between 1300 and 1450~ It should be noted that as the grain size increases the grain boundary surface area diminishes, hence for a given amount of impurity phase evenly distributed throughout the grain boundary network, the XPS atomic ratios on that surface would be expected to increase. Consequently the decreases in the atomic ratios represent a difference in morphology of the impurity phase within the grain boundary network between 1350 and 1450~ compared to 1300~ The key to understanding this behaviour is provided from XPS of the external surface of

212 the air-sintered discs and NMR studies of the decomposition of Na-Y-silicate glasses annealed

AIR

a) Y/Zr o n" $ .s Z r 13.

0.30

0.34 0.32 0.30

022 o

[

r .~

E zo 2

<

0.06 p-

g_..r

S,,Zr

~

0.14

/t /

--

~

.A ~ /

Ol

X Na/Zr

"~

20,000[-

tr-

10,000 r l 1200

1:3

13100 14100

I 1500

I 1600

1 lu(

Temperature (~

Figure 18. (a) Y/Zr, (b) Si/Zr atomic ratios obtained from fracture surfaces and (c) grain boundary resistivity versus sintering temperature. Specimens were sintered for 4h in air.

1200

1300

1400

1500

1600

1700

Tomperature ~

Figure 19. (a) Si/Zr and (b) Na/Zr atomic ratios from the external surface versus sintering temperature. Specimens were sintered for 4h in air.

in air and Ar environments. Increases between 1350 and 145&C of the Si/Zr and Na/Zr ratios on the external surface suggest the migration of a Na-rich silicate from the grain boundary network (Fig. 19). Moreover, it was evident from the XPS of the fracture surface that some silicate remained in the grain boundary network. Clearly, the silicate resident in the grain boundary network between 1350 and 1450~ is different to that present at 1200 or 130&C and also to the Na-rich silicate that migrates to the external surface between 1350 and 1450~ Thus for air-sintered specimens it appears that the silicate which is formed in the grain boundary network between 1200 and 1300~ either does not form at higher temperatures or otherwise decomposes into a number of different silicates. Some evidence for the latter was provided by NMR and XRD studies of the decomposition of some Na-Y-Silicates after annealing in air and Ar environments. For example, the Si NMR spectrum of Na3YSi207 glass after heating in Ar for 10h at 1500~ indicated that only glass is present since the broad resonance at 87ppm suggests a range of silicon sites (Fig. 20(a)).

213

a) Ar

Na 3YSi 207

b) Air

I

-60

I

-70

I

-80

I

-90

I

- 1oo

I

- 110

29Si N M R Shift ( p p m )

Figure 20. 29Si MAS NMR spectra of the glass Na3YSi207 after annealing in (a) Ar and (b) air for 10h. The fine structure in (b) indicates a large degree of crystallization. (Courtesy of Dr. M.E. Smith.)

After heating the same glass in air for 10h, however, a number of distinct peaks were present on a broad base centred at the same position as the glass after heating in Ar (Fig. 20(b)). The narrow resonances in Fig. 20(b) indicated a high degree of crystallinity and bear some similarity to the peak positions (although not peak intensity) for Y2Si207. XRD of the material also indicated reflections due to Y2Si207. In addition, the small broad base in Fig. 20(b) is evidence that some glass was still present. At temperatures above 145&C the XPS ratios on the fractured surface and the grain boundary impedance again displayed increases suggesting that the silicate phase formed at lower temperatures has wetted the grain boundary network. Apart from yttria redistribution through the movement of the grain boundary impurity phases through the grain boundary network, there was also yttrium enrichment of the surface of grains. This is particularly evident in Fig. 16, where the Y/Zr ratio is well above the bulk ratio at 1600 and 1700~ where no impurity phase is present in the grain boundary.

4.4. Polycrystalline tetragonai Y203-Zr02 In the Y203-Zr02phase diagram, materials with yttria contents of around 2 to 8.5mo1% Y203fall in the tetragonal and cubic two phase field. The tetragonal phase is in a metastable state at room temperature, and is only stabilized by its small grain size. (In a sintered body it is further stabilized by the strain energy of the surrounding matrix.) Normal processing at elevated temperatures brings these compositions into the tetragonal and cubic phase field where the tetragonal phase is fully stable. However, because the system is not at

214 thermodynamic equilibrium solute redistribution results in a partitioning of the ceramic into an yttria-rich and -poor microstructure; a process called solute partitioning. The high strength and toughness of Y-TZP makes it the most interesting polymorph of the Y203-ZrO2 system. Its strength is derived from the small grain size whereas its toughness is a result of the transformation toughening mechanism. The toughness is developed when cracks propagating through the matrix cause transformation of the metastable tetragonal (t) phase to the monoclinic (m) phase. A three percent volume expansion associated with this transformation closes the crack tip thus preventing further propagation. Apart from its strength and toughness, superplasticity has also been observed in Y-TZP which could lead to forming and extrusion of these types of ceramics. The higher strength and toughness of Y-TZP coupled with its high ionic conductivity at elevated temperatures has led to its use in fuel cells. 16

1700~ mol%Y203 (cubic)

12L/

I,~

c '~ 614

~ v ~

~

1.s mo,~Y=O=

/ ~

(cubic-tmragonal)

2, 0I 0

1

10

I

20

1

30

I

40

I

50

l

60

1

70

1

80

1

90

100

Time Hrs.

Figure 21. Grain size (pm) versus sintering time at 1700~ for (~) 8mo1%, (11) 4mo1% and (x) 1.5mo1% Y203-ZrO2. Grain growth for 4mol%Y203-ZrO2 (two-phase region) is much slower compared to the other compositions. After Sakuma and Yoshizawa [93] (reproduced by permission).

4.4.1. Solute Partitioning and Grain Growth It has been recognized for some time that Y-TZP with Y203 contents between 2 and 6mo1%, Y203-ZrO2 exhibits retarded grain growth [91]. Indeed, the grain growth rate has been reported to be some 30 to 250 times faster in the cubic phase compared to the tetragonal phase [92]. Grain growth has also been demonstrated to be faster for low Y203 contents in the tetragonal single phase field [93]. The sluggish grain growth for 4 mol% Y203-ZrO2 is

215 clearly demonstrated in Fig. 21 by comparision to 1.5 and 8mol%Y203-ZrO2. Retarded grain growth in Y~O3-ZrO2 coincides with the compositions in the two phase field and is closely related to the process of solute partitioning [76]. Materials sintered in the two phase field exist in a metastable state with a reasonably homogenous yttria distribution. However, during post-sinter annealing the yttria redistributes creating Y~O3 poor grains of 1.5mol%Y203 representing the solubility limit of Y203 in the tetragonal phase [94]. The Y203 also redistributes into an Y~O3 rich phase of nominally 6mo1% which .is cubic at sintering and annealing temperatures but transforms to the t' (tetragonal) phase at room temperature [93]. Redistribution into these phases can take as little as 10h at 1700~ but the equilibration time increases as the temperature decreases. In a number of studies, analysis of the grain growth kinetics in the two phase field yielded cubic behaviour as defined by equation 18, i.e. n=3 [50, 76, 93, 95], whereas grain growth kinetics for compositions either side of this region are quadratic suggesting normal grain growth [92]. grain boundary

40-

(~ i AZY17

30-I

bulk

O >O

g

20-

',

o

.

Co (ZY17) .

.

.

.

.

.

,

10-

~~~~~~~_

_Co_(_zY_5_8_)

Figure 22. Grain boundary enrichment of yttria after Theunissen [47]. The surfaces of both tetragonal and cubic YzO3-ZrOz enrich to the same level as determined by Auger electron spectroscopy and XPS, i.e. 30 - 34mo1% YOLs. The enrichment factor (total area under the curve) compared to the bulk is much larger for the tetragonal phase (ZY5.8), hence the effects of solute drag are larger. (Reproduced by permission.)

A number of different models have been proposed to explain the retarded grain growth kinetics in the two phase field regime. Lange et al. [96] proposed that strain energy within grains acted to retard grain growth. Lattice strain was proposed to develop from compositional differences as grain boundaries migrated from YzO3-rich grains through YzO3-poor grains leaving the yttria concentration difference behind them. On the other hand Sakuma and

216 Yoshizawa proposed that grain growth was restricted in the two phase regime simply by Zener pinning of one phase (major) by the other (minor) phase [93]. The cubic nature of grain growth kinetics suggests, however, that grain growth is restricted by a solute drag mechanism on the grain boundary mobility as proposed by Theunissen et al. [47-49] and Lee and Chen [92]. Burggraaf et al. have observed, using AES and XPS, that the surface Y203 composition of Y203-ZrO2 is always around 30 at% regardless of the bulk composition in a variety of YzO3-ZrO2 with Y203 compositions spanning the two phase field as well as the cubic phase [55]. Hence the enrichment ratio is much larger for the tetragonal phase than the cubic phase as depicted in Fig. 22. Consequently, according to equation 22, the amount of grain boundary drag can be expected to be considerably more in the tetragonal phase than the cubic phase.

(a) as-sintered

r

.o

0

I I e) t~

O~ 0 0 0

distance, d Gram Boundary (GB)

(c) at equilibrium

(b) during solute partitioning

0 C 0 0

0 C

I

0

I

distance, d Grain

Boundary (GB)

0

distance, d Grain Boundary

(GB)

Figure 23. Schematic representation of solute partitioning. The as-sintered material has a homogenous distribution of yttria (a). During post-sinter annealing (b) the loss of yttria from one grain and the gain of yttria in the adjacent grain will result in surface enrichment for both grains. After equilibrium is reached (c) the surface enrichment will be less for the grains with high bulk yttria concentration hence normal grain growth can proceed (assuming that both types of grains experience the same surface enrichment).

The reason for this enrichment lies in the solute partitioning. During the process of yttrium redistribution from yttria-poor to yttria-rich grains the surface of both types of grains will be enriched until equilibrium is reached (Fig. 23). The yttria-poor grains develop surface enrichment because yttria is being withdrawn from the bulk through the surface and the yttriarich have surface enrichment because the bulk yttria content must be built up via diffusion from the surface. It is only after equilibrium is reached that normal grain growth can proceed [97].

217

1450~

I~ / 9TS3/1

~ 2 ~1.8

1.6

I

1.40

I 100

2.4

I

I 200

1450~

2.2

I 3OO

b

l9

]

I

/

9TS2/2

1.8

1.6

2't

1500~

2.2

E1

1.4

.

8

--

9TS2/1 9TS3/2

~

100

200 time (hrs)

3oo

Figure 24. Y203 content in the tetragonal phase during annealing at the indicated temperatures. Y203 contents have been determined using EDS microanalysis in transmission [76]. TS3/1 and TS2/1 contain high levels of A1203 (650ppm) and silica (910ppm) respectively whereas TS3/2 and TS2/2 are relatively pure.

Solute partitioning can be enhanced by the presence of impurities as pointed out by Nauer and Carry in XRD studies of the yttria composition of the yttria poor tetragonal phase (Fig. 24). They found that the time taken to reach the equilibrium Y203 concentration was faster for impure materials. For example, in Fig. 24, the %Y203 for yttria-poor tetragonal grains has been determined as a function of time at various temperatures. In Fig. 24(a) TS3/1 (650ppm A1203) has 13 times the impurity alumina content but only a third the impurity silica content

218 (20ppm) of TS3/2. In Fig. 24(b) TS2/1 has over 46 times the impurity silica content and thirteen times the A1203 content of TS2/2. Both TS3/1 and TS2/1 reach the equilibrium Y203 content more rapidly than the pure specimens. Indeed, the pure materials fail to reach the equilibrium content after extensive annealing at 1500~ To date, no detailed surface analysis studies have been undertaken of the change in the composition of the fracture surface of these types of specimens in their evolution towards equilibrium. Hughes and Badwal, however, have examined the changing composition of glassy phases as well as the surface composition of grains on the external surface of two YTZP's of vastly different impurity contents [98]. They found a number of interesting results. First, the level of impurity segregation to the external surface of polished discs was more dependent on the composition of the glassy phase rather than the absolute level of impurity within the material. Second, the nature of the impurity phase appeared to play an important role in the degree of yttrium redistribution. The ramifications of these points are obvious. The temperature and rate of phase partitioning could potentially be adjusted by an appropriate choice of impurities.

4.4.2. Superplasticity In 1986 Wakai and coworkers reported on the potentially important discovery of tensile elongations in Y-TZP of over 100% [99]. Since that report many other studies have been undertaken of superplasticity in Y-TZP, both in compression and tension which have been comprehensively reviewed by Maehara and Langdon [ 100]. Indeed, forming applications have already been demonstrated for Y-TZP [101]. Superplasticity has also been demonstrated in A120~{-TZP [102, 103]. While there is no debate that Y-TZP undergoes superplastic deformation at low strain rates, there is considerable debate as to the mechanism. For deformation in superplastically deformed materials, the strain rate is defined as: "

A{I}Pd o'oxp(-~

(23)

where A, p and n are constants, Q is the activation energy, R T has it usual meaning, d is the grain size and ~ is the flow stress. Controversy has centred around the value of n (or its reciprocal m, the strain rate sensitivity). Studies of superplasticity in Y-TZP have focused on the determination of n and Q. Unlike metals where n is essentially constant, values of n for Y-TZP typically fall between 1.9 and 3.2 suggesting different mechanisms of deformation. Hermansson et al. [ 104], for example, found n to vary from 1.9 to 2.4 as the level of intentionally added impurity increased. They maintained, through TEM observations, that the additional glass had no effect on the thickness of the grain boundary film, but increased the amount of glass at triple junctions resulting in rounded grains which, presumably, could move past each other more freely during superplastic deformation. On the other hand, Nieh et al. [105] and Maehara and Langdon [100] determined values of n closer to three, although Nieh et al. later concluded that may be closer to two when the effects of dynamic grain growth were taken into account.

219 I I

~

. - . . . .

grain " boundary thickness

\

.-~

upper grain - ~ ' - - _/

\, ~ ~\ ~i I

I I

,

~

~

elongation

.. 1t

Figure 25. In the model proposed by Langdon for superplastic deformation n can vary from one for high coverage of the grain boundary with impurity phase to three for clean grain boundaries [ 106].

Figure 26. For the lower grain to slide past the two upper grains, its apices must sweep out a path indicated by the hatched areas suggesting the rapid redistribution of material. The presence of glassy phases can assist matter transport.

Values of n close to two (or lower) are consistent with a grain boundary sliding mechanism which is supported by the observation of equiaxed, rather than elongated, grains after superplastic deformation. The mechanism of grain boundary sliding is unclear, however, Langdon has put forward a plausible model where the value of n varies according to the fraction of surface area covered by impurity phase (Fig. 25) [106]. Specifically, n has low values for a high level of grain boundary coverage (approaching 1) whereas it is closer to three for pure materials. The large variation in n arises because of the method of matter redistribution during superplastic deformation. It is clear in Fig. 26 that for the bottom grain to slide past the two upper grains that it must sweep out a path indicated by the hatched areas. To achieve this, matter from the apices of all three grains must be redistributed, which for the case of clean grain boundaries is achieved via the flow of vacancies and surface and bulk diffusion. Redistribution is much faster, however, in the presence of a grain boundary impurity phase due to liquid phase assisted transport. One further point to note is that the value of n decreases during grain growth which occurs during superplastic deformation [103, 106, 107]. Langon has suggested that this may be due to an increase in the relative coverage of the grain boundary by the grain boundary impurity phase with a reduction in the total grain boundary surface area upon grain growth. With the degree of coverage of the grain boundary by impurity phases being of such critical importance to the mechanism of superplastic deformation it is surprising that there

220 are no extensive studies of the grain interface using surface spectroscopies. To date XPS and AES have only been used in one study by Nieh et al. [108], who claimed to have clean grain boundaries. This claim should be viewed with some caution since (i) if there is epitaxial growth of impurity phases onto the surface of grains, as claimed by Clarke, then their detection may be difficult by TEM and (ii) small submonolayer coverage or island formation may be difficult to detect using XPS. 1,2oo I

t.

~E j::-

ill

11)

0 0

I I I I I I I I 100 200 300 400 500 600 700 800 Aging Temperature ~

Figure 27. Degradation in flexural strength as a function of aging time in moist environments. (O) after Matsui et al. [115] and (x) after Watanabe et al. [114].

4.5 Low Temperature Degradation The mechanical degradation of sintered fully-dense Y-TZP in moist environments has been the subject of intense research for a number of years. Microcracking [109-111] and a diminution of mechanical properties [112-115] are characteristics of the degradation which has been shown to be due to the transformation of surface tetragonal (t) grains to the monoclinic phase (m). Fig. 27 displays typical moisture sensitive degradation in this case where there is a considerable loss of bend strength between 100 and 400~ after treatment in moist environments. The transformationproceeds from the surface inwards with the grain boundaries probably providing an important pathway for subsurface attack. X-ray diffraction (XRD) has been used monitor the increase in time of monoclinic phase content of the surface at temperatures between 100 to 400~ [110, 111, 113, 116-118]. The transformation has been shown to be dependent on the grain size, and the yttria content of the Y-TZP [114, 119]. As can be seen from the result of Chen and Lu (Fig. 28(a) as the grain size decreases the %m-ZrO 2 developed also decreases. Indeed, there appears to be a critical grain size below which the transformation does not proceed. Watanabe et al. [114],

221

,~176 / 8~F

1.15p, m

O o "E

20

0

0

1

2

3

4

5

10

A g e i n g T i m e (hrs)

O o "E

~ 4o ~;

0.49jff,m 2O

00

~

~j

200

,

300

j~,j,

400

500

Ageing Temperature (~

Figure 28. Effects of grain size on the development of the monoclinic phase in Y-TZP (a) with aging time in a moist air stream at 18&C and (b) aging temperature for 5h after treatment in a moist air stream [119].

for example, found that the critical grain size increased from 0.2 to 0.6pm with an increase of the yttria content from 2 to 5 mol% Y203. Certainly as the yttria content is increased the tetragonal phase becomes more stable and consequently there is less transformation to the monoclinic phase [109, 113]. Sato and Shimada [120], and Chen and Lu [119, 121] have shown that the production of the monoclinic phase is, however, completely reversible and all the monoclinic phase can be removed by heating above Ms, the martensitic start temperature. Ms, which defines the t ~ m transformation temperature in a constrained matrix, decreases with decreasing grain size (Fig. 28(b)) and increasing yttria content. To date, there is no definitive theory of why the tetragonal phase transforms to the monoclinic phase upon attack by water, although there are a number of contending models. It is clear that the attack proceeds from the surface inwards, hence surface sensitive techniques have an important role to play in identifying surface species present after exposure to moist environments. The study of moisture sensitivity is made more complex by a variety of specimen preparation and treatment conditions. Since the amount of transformation depends

222 on the grain size and yttria content of the Y-TZP then specimen preparation, which determines the microstructure, becomes an important variable. The microstructure will be determined by the impurity levels of the starting powders, sintering temperatures and times, cooling rates, sintering atmospheres and the role of the impurities in promoting redistribution of yttria and inhomogenous grain growth [98, 122].

4.5.1. Models Under free energy considerations the tetragonal phase of yttria-zirconia only exists if the free energy change of the tetragonal--+monoclinic (t---~m) is > 0 [123, 124], i.e.:

AtTt.m - (t7c" - Gc t) + (Us," - U,, ~) + (Us" - Us t) z0

(24)

Gc is the chemical free energy and Us, is the strain energy associated with the transformed particle. There are a number of contributions to Us, including strain associated with a volume expansion in x, y and z axial directions as well as a shear strain which contributes the largest to the volume expansion [ 116]. Us is the change in energy/unit volume associated with the surface and is given by:

AU,- --~(V. - g,V,)

(25)

where d is the diameter of the transformed inclusion, g, is the quotient of tetragonal to monoclinic interfacial surface areas and Yi is the specific interfacial energy of phase i. At room temperature both G'c and U',, are larger than G~c and U=s, [124]. As the grain size decreases the relative contribution of the surface energy term increases, stabilizing the tetragonal phase. It is within this free energy framework that a number of models have been developed which endeavour to explain the t--+m phase transformation in moist environments. Lange et al. [ 118] proposed that in moist environments, precipitates of Y(OH) s formed within the surface grains and were accompanied by a depletion of yttrium in that part of grain surrounding the precipitate. This depleted zone subsequently acted as a nucleation site for the transformation to monoclinic phase (Fig 29(a)). Yoshimura [75] proposed that water adsorbs onto the surface of the ceramic by hydrogen bonding to a bridging Zr-O-Zr bond. Subsequently, there is scission of the bridging oxygen bond, resulting the formation of two Zr-OH bonds. The scission of the surface bridging Zr-OZr bonds permits subsurface access to 1-120 and further attack by water vapour, resulting in microcrack development. The presence of microcracks at the surface of Y-TZP, results in the accumulation of strain and hence the transformation to monoclinic (Fig. 29(b)). Sato and Shimada proposed another model where the internal stress of the tetragonal phase is relieved by the attack of water vapour along the grain boundaries resulting in the dissolution of cations [ 125]. Sato and Shimada [120] and Winnubst and Burggraaf [124] have also suggested that the adsorption of hydroxyl ions onto the surface of Y-TZP lowers the surface energy difference between the tetragonal and monoclinic phases, thereby spontaneously promoting the transformation.

223 Y(OH) 3

(a)

Y(OH)3 (b)

H

\

Yttriumdepletedzone

oJ H I

/ \Zr

/ \Zr"

0 O~Zr/

/ Zr

[] 0 ~Z/

o\ / OH

Zr HO ;//O/ ~ [],,, ~Zr

"~~Zr/H HO/~0 ~Zr?/Y~O "-~ O~ ~OH ~Zr \~ ~\ o/ ~,~- S.ain 0

'2, O ' ~ J ' k ' ~J _ r-,.l A c c u m u l a t o Zr

Figure 29. Models of low temperature degradation in moist environments. (a) Model attributed to Lange et al. [118] and (b) Yoshimura et al. [128] (see text for explanation).

There are also additional effects which may alter the stability of the tetragonal phase during sintering. For example, yttria concentration gradients may occur in the surface of grain due to segregation and/or the dissolution of glassy phases within the grain boundary network [55, 62, 98, 108, 122, 126]. While inhomogenous yttria distributions within grains will probably not effect AGc in equation 24, they may effect the strain energy terms through changes in the elastic moduli and the surface energy terms. Furthermore, the presence of glassy phases within the grain boundary network will modify the interfacial energies. A glassy phase which wets the surface of the grains will lower 7' and will further promote the t-->m transition since /f, is less than U", in the absence of glassy phases. To give some idea of the amount of m-ZrO2 developed in moist environments the %m-ZrO2 phase developed in pure and impure Y-TZP after a range of treatments is presented in Table 3 [73]. It can be seen that no monoclinic phase was detected on the surface of the discs after they were polished and heated to 600~ Autoclave treatment of both ceramics resulted in the most transformation to m-ZrO2. Treatment in moist air for 48 hours at 200, 250 and 300~

224 resulted in significant transformation for the impure Y-TZP at all temperatures with a maximum at 250~ whereas the pure Y-TZP only displayed any significant transformation at 200~ After treatment for 70 hours at 120~ and 170~ there was 73 %m-ZrO 2 and 77%mZrO2 respectively. Indeed, after treatment at 170~ the pure Y-TZP had disintegrated leaving only flakes and powder from the former disc whereas the impure Y-TZP disc was still intact with no visible macroscopic degradation other than some surface rumpling. Chen and Lu [119, 121] have found that as the grain size decreased the maximum temperature at which the t--->m transition occurred also decreased which was attributed to a lowering of the martensitic start temperature (Ms) in the constrained matrix of a sintered material. On this basis alone, given the larger grain size of the pure Y-TZP, its M s should have been higher than the impure Y-'IT_~. However, ~ for the pure material appears to be around 250~ whereas it is above 300~ for the impure material. TABLE 3 %m-ZrO2 Generated Under Various Treatment Conditions for Pure and Impure Y-TZP Treatment Grain Size (lain) Polished + 600~ (5h Dry

Pure

Impure

0.8

0.9

0

0

Air) Polished + 600~ (5h dry air) + 200~ (48h moist air)

62, 61

44, 43

250~ (48h moist air)

11, 0

58, 57

300~ (48h moist air)

0

55

120~ (70h autoclave)

73

73

170~ (70.5h autoclave)

80

77

Clearly other factors are coming into play which may include the Y203 content of surface grains which will modify all the free energy terms in equation (24), the presence of a glassy phase in the grain boundary network near the surface which modifies the surface energy of the tetragonal phase and the rate of ingress of water, differences in grain size distributions, differences in the defect distributions or a combination of all these factors. In the study of moisture sensitivity surface sensitive techniques such as XPS and FTIR have generally been employed to examine surfaces after severe treatments. Hernandez et al. [111, 127] observed YO(OH) on the surface of Y-TZP after hydrothermal treatment, but none on Ce,Y-TZP after similar treatment. Lepist6 and M/intyl/~ [ 117] observed significant changes in the hydroxide component of the O l s signal and changes to the Zr 3d signal after extended exposure of Y-TZP to moist air. Yoshimura et al. [ 128] and Shigematsu et al. [ 129] observed significant increases in the OH stretch after treatment in moist air or hydrothermally. These studies however only characterize the corrosion products but reveal nothing of the change to

225

a) Y/Zr

0.077 -0.069

'~o

,o

/9= 0 ~

o

0.061 I'-&&'& o

~9

o~ ft. X

= 70~

0"053 1 7 ~ ~ -

1.5

__&

o

b) O/Zr 0.5

50

c) % Zr 3+

40 30

o

20

1~t 0

0

1

10

I

20

I

30

I

40

9

I

50

I

60

% Monoclinic Zr02

Figure 30. XPS atomic ratios vs %m-ZrO2 generated in a moist air stream at 200~ [73]. (a) Y/Zr, 0= 0 ~ (O,@) and 70 ~ (A, -). (b) O-/Zr and OH/Zr ratios (@) - pure Y-TZP and (O) impure Y-TZP at 0 = 0 ~ (c) % Zr~ vs %m-ZrO2 at 0=0". (Q) - pure and (O) - impure Y-TZP.

the surface during transformation. Hughes et al. have studied the kinetics of transformation using XPS for the treatment of YTZP in moist air and hydrothermally. For example, in Fig. 30(a) the Y/Zr ratio measured by XPS at 0---0 and 70 ~ is plotted against the %m-ZrO2 of the surface. At both 0=0 and 70 ~ there were considerable changes up to 15%m-ZzO2; initial differences in composition were removed at around 8%m-ZrO 2 where the two surfaces approached a similar composition. For the surface (0---700, 2.4nm) the Y/Zr ratios decreased monotonically for both Y-TZP's to an equilibrium value of-0.052 at 15%m-ZrO2. By contrast changes in the Y/Zr ratios for the top 7.0nm (0---if') were complex between 0 and 15%m-ZrO2. They f'wst increased between 0 and -6%m-ZrO2, but between 6 and 15% %m-ZrO2 the behaviour was unclear and above 15%mZrO 2 both displayed similar values. These results suggest that there is a surface depletion and subsurface enrichment of yttrium. Large increases in O- and Zr s+ were also observed in the 0-15%m-ZrO2 region. There was

226 no accompanying increase in the amount of OH except at high %m-ZrO2 i.e. long treatment times. The percentage of the Zr~ present in the surface is plotted against the %m-ZrO2 content of the surface in Fig. 30(b). The percentage of Zr~ peaked at around 35% for the pure and 50% for the impure Y-TZP's respectively at 10%m-ZrO2. For longer exposures the percentage Zr~§ decreased to around 20%. As discussed in section 3.2, Zr ~§ arises from electrons trapped at vacancy sites adjacent to Zr cations. The changes in Fig. 30(c) therefore represent significant changes in the vacancy concentration in the surface during the process of transformation from the tetragonal to monoclinic phases. Zr~ is also evident in the Zr 3d spectra of Hemandez et al. [111] and Lepist6 and M~tyl~i [117], after moisture sensitive degradation of Y-TZP, although it was not present in Ce, Y-TZP after similar treatment. From these results it is clear that the vacancy concentration and annealing of vacancies by O- have an important role to play in the transformation. Vacancy annealing results in a strain field within the surface of the grains which induces the t ~ m transformation. The likelihood of the transformation is improved by the migration of Y to the subsurface region leaving an yttria-poor surface. The exact mechanism for the production of O- from H20 and the role of hydrogen remains elusive. The results of Narita et al. [130] and Shubert et al. [131] indicate that hydrogen is taken up by the Y-TZP, although it is unclear as to whether it occurs as interstitial hydrogen ions, OH or some other form. Hydrogen in the form of H § may penetrate the surface, residing at interstitial sites, and provide charge balance for the excess O-. The adsorption of oxygen and hydrogen result in the well documented weight gain of the ceramic [130]. A similar mechanism was invoked by Badwal and Nardella [132] to explain the presence of a considerable amount of m-ZrO2 at the anodic side of Y-TZP during current flow. Hydroxides of Zr only developed after long exposure to a moist air stream at 200~ or autoclave environment at 120 or 170~ Hydroxides of yttrium were only observed after autoclave treatment at 170~ For example, the Y 3d spectra after treatment with moist air at 200~ were adequetley fitted with a single component (Y3d5/2-3/2 doublet) representing yttrium cations in the lattice (Fig. 31a). However, after autoclave treatment at 170~ there was clear evidence of yttrium hydroxide formation and a second component was fitted to the Y 3d spectrum (Fig 31b). Hence, hydroxide formation should be viewed as symptomatic of the degradation in moist environments, i.e. corrosion product, rather than the initiator of the t ~ m transformation. These phases presumably result from the dissolution of Zr4§ and Y~ from the matrix with the subsequent deposition of hydroxide onto the external surface of the ceramic. They may however, have a role to play in the macroscopic and microscopic disintegration of the ceramic via a stress corrosion cracking mechanism or corrosion attack along grain boundaries resulting in hydroxide formation. Impurities have a dual role to play in the transformation; first grain boundary impurity phases modify the characteristics of the intergranular region and second soluble impurities can presumably modify the characteristics of the grain surface. The extent and nature of modification is difficult to estimate in either case. It is significant that after treatment in moist air, the Si/Zr ratios increased in the same fashion with the development of the %m-ZrO 2 for both the pure and impure Y-TZP's discussed above. The increase was much slower for the impure Y-TZP as a function of time suggesting that the presence and/or nature of the impurity phase may have some controlling role on the kinetics of transformation. The slower kinetics could be related to the different composition of the grain boundary impurity phase [98]. In

227

oxide

a) Moist Air

..

b) Autoclave

153.0

157.0

161.0

165.0

Binding Energy (eV)

Figure 31. Y 3d spectra for pure Y-TZP after treatment in (a) moist air at 200~ (48h) and (b) autoclave at 170~ (170h). Each doublet represents a single yttrium species. Only yttrium in an oxide environment was present in (a) but the hatched component in (b) indicates that some hydroxide is present.

the autoclave experiments Si was detected on the surface of both Y-TZP's at short autoclave times but none was detected after 70 hours treatment at either 120~ or 170~ Chemical analysis of the autoclave solution revealed small amounts of Si (up to 3ppm) which could be the result of dissolution of impurity phases from the ceramic. From free energy considerations, the presence of some form of silicon on the external surface of the ceramic could change the surface energy. Holmes et al. [133] concluded, on the basis of experiments where they measured the heat of immersion in water, that the surface energy was much lower for the tetragonal phase (500---)600 erg/cm2) than the monoclinic phase (-1100 erg/cm2). Earlier work of Harkins and Boyd [134] gave the heats of immersion of SiO2 and ZrSiO4 as 600 and 850 ergs/cm2 respectively suggesting that the presence of some oxidized form of silicon on the external surface of the ceramic may increase the surface energy of the tetragonal phase thereby reducing the free energy barrier to transformation. Lattice soluble impurities are commonly present in Y-TZP, which form glassy phases with SiO2. Hughes and Badwal [ 122] have found that it is possible to produce significant impurity and yttrium concentration gradients within the surface of grains. These concentration gradients will significantly alter the vacancy concentration and the presence of the glassy phase in which the t-ZrO2 grains are imbedded should modify the surface energy of the phase.

228

Ce Oxidation States a) CeO 2

iv,,,

~u'"

/'L~

1 885.0

._

.~

I 905.0

__,,,.,,.,___

I 925.0

Binding Energy (eV)

Figure 32. XPS Ce 3d spectra. (a) Ce4+ spectrum from CeO2 and (b) reduced CeO 2 with a mixture of Ce 3+ and Ce 4+. The Ce 3d5/2 (3d3/2) features labelled v (u) and v " (u") correspond to Ce4+ mixed final states 4f 1 (Sd 6s) ~ 3d 9 and 4ta (5d 6s) ~ 3d9 whereas v " ' (u'") can be attributed to a final state Ce4+ 4f ~ (5d 6s) ~ 3d9. The Ce 3+ features labelled v ~ (u~ and v' (u') are attributed to Ce 3+ mixed final states 4f 2 (Sd 6s) ~ 3d 9 and 4f 1 (5d 6s) ~ 3d 9 respectively [141]. Hence the u " ' is the only peak which can be attributed to Ce4+.

5 CeO2-Y203-Zr02. The ternary systems CeO2-YzO3-ZrO2 and Ce203-Y203-ZrO2 are closely allied since Ce203Y203-ZrO2 is very sensitive to sintering atmosphere because the degree of reduction of Ce4+ to Ce 3+ can depend markedly on the partial pressure of 02 [135]. For example, XPS and thermogravimetric analysis of 25.5CeO2-2.SY203-72ZrO2 (by weight) treated under reducing conditions (10%HJN2) indicated a considerable weight loss between 560~ and 780~ accompanied by the generation of Ce 3+ [61]. The temperature range 560 to 780~ coincides well with the observed temperature programmed reduction behaviour of pure ceria from CeO2 to CeOl.s2 [136]. Reduction of the ceria results in the production of 02 which at temperatures in excess of 1300~ can lead to bloating and a loss of density [137]. Furthermore, upon generation of Ce 3§ the phase relations change resulting in the detabilization of the tetragonal phase and generation of the monoclinic phase [138]. Again the weight loss at higher temperature corresponds well the reduction CeO2 to Ce203 [139]. Weight loss under nitrogen for the binary system CeO2-ZrO2 (9-12mo1%) displays similar behaviour beginning at 1150~ It is clear that the amount of Ce 3+ can produce a significant alteration in the properties of the CeO2-Y203-ZrO2 system since it can modify grain growth [13], sintered density [ 137, 140] and the phase relationships [ 138]. While Ce spectroscopy in XPS is complicated due to the

229 presence of a satellite structure arising from hybridization with O 2p orbitals and partial occupancy of the 4f levels in CeO2 [ 141, 142], the amount of Ce3+ is readily estimated from the peak labelled u'" in the Ce 3d spectrum in Fig. 32. This feature arises solely from Ce4+ and represents a fixed percentage of the Ce 3d spectrum of a material with cerium in the 4+ oxidation state. Hence the amount of Ce3§ can be determined from the ratio of the integral u ' " peak area to total Ce 3d area [143, 144].

SILICON C O M P O U N D S

/

1611.0

/ ..o.L/

V

1712.0

t3rj L (D t"

uJ o 1610.0

1711.0

.i-, r"

.-I .-I

v

1609.0

/2 Z/

1608.0 104.0

103.0

>/ / / // 102.0

1710.0

1709.0 101.0

Si2p Binding Energy

Figure 33. Si 2-dimensional chemical state plot for segregated impurity phases for air (@) and Ar (O) sintered Y-CSZ, on the fracture surface (11) and external (r'l) surfaces of 25.5wt%CeO22.5Y203-72ZrO2 [23] and the fracture surface of (v) of 30wt%A1203-Y-TZP after sintering at various temperatures in Air. I = sheet silicates, II = chain silicates, A = Aerosil and Q = quartz.

The CeO2-Y203-ZrO2 system has been investigated for a range of applications from thermal barrier coatings (TBF) to a less moisture sensitive TZP. A common requirement for both these applications is small grain size. Smaller grain sizes are obtained for the ternary system compared to the binary CeO2-ZrO2 system, [49, 145] but the smallest grain sizes are observed in the binary Y203-ZrO2 system. The grain growth kinetics in CeO2-Y203-ZrO2and CeO2-ZrO2 are complex [140, 145]. At temperatures below 1200~ it is likely that the grain growth kinetics in CeOE-Y203-ZrO2 are retarded by solute drag (n=3) due to the segregation of y3+, as in the binary YEO3-ZrO2system. At higher temperatures Ce3+is generated which segregates to the grain surface competing with y3+. The dragging effect of Ce~ is presumably less than Y~ since grain growth becomes more rapid. In the binary CeO2-ZrO2 system generally normal (n=2) grain growth is observed, although after long sintering times anomalous (n=l) grain

230

growth is observed which is accompanied by a loss of density [ 140, 145]. CeO2-Y203-ZrO2 thermal barrier coatings are used in aircraft engines principally for thermal insulation of metallic surfaces resulting in lower metal temperatures and reducing the impact of thermal gradients [8]. Metallic components are generally coated by plasma spraying a partially stabilized ZrO2-based topcoat onto an oxidation resistant subcoat. Ceria additions have been found to produce a high fracture toughness (although this is lowered with small additions of Y~O3) in the tetragonal phase [139] and also display high deformation plasticity [146]. An important feature of these types of coatings is their mechanical integrity because flaws in the coating can produce local hotspots leading to component failure.

o) t_ ~) r" UJ .o_

1388

1461

1387

1460

1386

1459

G) r ._ v ..J _J v


/ 75.0

74.0

73.0

AI2p Binding Energy

Figure 34. A1 2-dimensional chemical state plot for impurity phases on the fracture surface of 5 and 30wt% (a) and the external surface of 5 and 30wt% (11) A12OjY-TZP and the fracture surface (@) of 25.5wt%CeO2-2.5Y203-72ZrO2 [126]. I = aluminium hydroxides and oxides, II = Aluminosilicates with part octahedral aluminium and III = aluminosilicates with fully tetrahedral aluminium.

The likelihood of failure is increased when there is segregation and accumulation of impurity phases in the grain boundary network, where their presence can adversely modify the fracture mechanics leading to intergranular fracture. In a number of studies, Ingo et al. [61, 126, 147] have examined segregation and other effects on the external and fracture surface of thermal barrier coatings. They found that silicon and yttrium segregate to both surfaces at sintering temperatures above 900~ with no apparent difference in the level of segregation to either surface. Aluminium also segregated at higher temperatures. Furthermore, cerium desegregated from both surfaces, suggesting that Ce was not incorporated in the

231 impurity phase. 2-D chemical state plots for the silicon and aluminium indicated that the impurity phase was silicate in nature and at least some of the aluminium was in tetrahedral coordination [147]. Using silicon and aluminium 2-D chemical state plots, Ingo et al. demonstrated how XPS can be used to identify segregated silicon as silicate and furthermore that aluminium was incorporated into the silicate. The identification of aluminosilicate was evident from the Si 2-D chemical state plot in Fig. 33, since the data for the segregated phase is intermediate between silicates not containing aluminium and surface aluminosilicates in A1203-Y-TZP [60]. The A1 2-D chemical state plot conf'mns the similarity between the aluminium environment for the segregated impurity phases and aluminosilicates as well as indicating that there was a mixture of tetrahedral and octahedral aluminium sites thus providing some structural information about the glassy phase (Fig. 34).

6. AIzO3-YzO3-ZrO2

A1203-Y203-ZrO2 composites form an important class of structural and ionic conducting ceramics. Tsukuma et al. [ 148] first reported bend strengths of around 2400MPa for around 20wt% A1203 addition. It was evident that HIP'ing was an important step in achieving the high strengths of these materials since it reduced the high porosity obtained under normal sintedng conditions [148, 149]. Superplastic elongations of around 625% have also been reported for A1EO3/Y-TZP [150]. Additions of A1203 have also been reported to reduce the grain boundary resistivity of yttria-zirconia [151, 152]. Rajendran et al. [153] reported that the addition of 10wt% A1203 to Y-TZP produced a minimum in the grain boundary resistivity for coprecipitated powders. Clearly for superplastic deformation and ionic conducting applications the nature and purity of the grain boundary is extremely important. Considerable attention has been given to understanding the mechanism whereby A1203 additions modify the grain boundary conductivity in yttria-zirconia. Earlier work concentrated on physical mixtures of alumina with stabilized-zirconia. Improvement to grain boundary conductivity and sinterability was attributed by Buffer and Drennan to the scavenging ability of alumina additions for impurity silicates [ 151 ]. In their model, developed for mechanically mixed powders, migrating, silicate-rich grain boundaries come into contact with alumina particles, where reaction of the impurity silicate immobilizes the silicate, furthermore rapid impurity diffusion along the grain boundary towards the alumina particle effectively cleans the adjacent boundary region of impurity silicate. Later work has concentrated on coprecipitated powders, where the mechanism of grain boundary cleaning will be slightly modified. Rajendran et al. demonstrated that the grain boundary resistivity had a minimum at 10wt% alumina addition for coprecipitated A1203Y203-ZrO2 corresponding to approximately one alumina grain per grain boundary [153]. In a fashion similar to physically mixed powders, glassy grain boundary phases were observed by TEM at the A1203/Y-TZP interfaces. XPS studies of the fracture surfaces revealed, however, that at least some amorphous alumina was still in excess in the grain boundary along with silicate after crystallization of most of the alumina to tx-A1203. The XPS spectroscopic data in the form of Si and A1-2D chemical state plots (Figs. 33 and 34) revealed that the impurity phase bore some similarity to the aluminosilicates detected in 30wt%A1203/Y-TZP composite [60], suggesting that the grain boundary phase was

232 aluminosilicate in nature. Furthermore, depth profiles into the fracture surface revealed only a thin layer of aluminosilicate at low sintering temperatures (-2.4nm) but much deeper profiles at 1700~ (>7.5nm). Based on these results, it was proposed that silicate scavenging by alumina in coprecipitated powders was at least two-step. At low temperatures (<1300~ impurity silicate reacts with the amorphous grain boundary alumina. At these temperatures the Y-TZP is effectively embedding in an amorphous alumina matrix. Between 1200~ and 1300~ the amorphous alumina crystallizes to o~-A1203,however, some aluminosilicate remains in the grain boundary. This aluminosilicate was only removed from the grain boundary at 1500~ as revealed by the XPS AI/~ atomic ratio which indicated that there was a homogenous mixture of the alumina and Y-TZP phases. The Si/Zr depth profile presented in Fig. 35 can therefore be considered as a profile through the alumina grains in the fracture surface. The profile therefore suggests that there is a zone in the surface of the alumina grains rich in aluminosilicate, conf'maaing that the alumina has successfully scavenged the impurity silicate. 0.24

Depth Profile 30wt% AIzO3-Y-TZP

0.20 t~ nL

0.16

E Z E o

0.12

ll) .0

O~

0.08

1700~

0.04 0.0

0

30 Sputtering Time

60

90

(sec)

Figure 35. Depth profile into the fracture surface of 30wt%A1203-Y-TZP after sintering for 2h [601.

7. CONCLUSIONS The interfacial chemistry in yttria-zirconia ceramics can significantly modify many properties. Intergranular impurity phases can restrict conduction paths for oxygen anions and yttria enrichment within the surface of grains can increase anion impedance. Yttria and impurity cation enrichment at the surface of grains can inhibit grain growth through solute drag. The presence of impurity silicate phases within the grain boundary network may assist in plastic deformation. Furthermore, sintering conditions such as atmosphere and time may

233 result changes in the grain surface and impurity phase compositions further modifying other properties. It has been demonstrated that surface spectroscopies, in particular X-ray photoelectron spectroscopy can provide much information on the chemistry within the grain boundary network with relative ease. As well as providing quantitative analysis of grain boundary composition, XPS can be used to determine the oxidation state of species present and even provide information on vacancy concentrations within the surface. The information gained from XPS and other surface techniques currently only assists our understanding of macroscopic properties which depend surface properties. However, ceramic processing is moving towards smaller grain-sized ceramics in which the surface modified region becomes an increasingly larger fraction of the total volume of the ceramic. For example, current grain sizes are frequently 300nm and the surface modified region 3nm, hence the surface modified region constitutes only 6% of the total volume. With order of magnitude decrease in the grain size the surface would constitute nearly 50% of the total volume of the ceramic. In such systems the need for surface spectroscopies as well as electron microscopy will become increasingly acute.

8. REFERENCES

8 9 10 11

12 13 14 15 16 17

M.J. Murray and S.P.S Badwal, Mat. Sci. Forum, 34-36 (1988) 213. T. Yazawa, Present State of Solid Oxide Fuel Cells in Japan in: Proceedings of the First International Symposium on Solid Oxide Fuel Cells, S.C. Singhal (ed.), The Electrochemical Society Inc., Pennington, 1989, pp. 387-393. K. Tanabe, Mater. Chem. Phys., 13 (1985) 347. H.C. Yao, H.K. Stepien and H.S. Gandhi, J. Catal., 61 (1980) 547. L.M. Mercandalli, D. Dieumegard and J. Siejka Epitaxial Growth of Silicon on YttriaStabilized Zirconia (YSZ) and Subsequent Oxidation of the Semiconductor-YSZ Interface, E-MRS, Syrup. 12, Dielectric Layers on Semiconductors: Novel Technilogical Developments, Strassbourg, (1986) 153. V.A. Loebs, T.W. Haas and J.S. Solomon, J. Vac. Sci. Technol., AI (1983) 596. M. Kawasaki, S. Nagata, Y. Sato, M. Funabashi, T. Hasegawa, K. Kishio, K. Kitazawa, K. Fueki and H. Koinuma, Jap. J. Appl. Phys., 26 (1987) L738. R.A. Miller, Surf. Coat. Tech., 30 (1987) 1. A.E. Hughes, to be submitted. Q.Y. Ma, E.S. Yang and Chin-An Chang, J. Appl. Phys., 66 (1989) 1866. S.P.S. Badwal, M.J. Bannister and W.G. Garrett, J. Phys. E: Sci. Instrum., 20 (1987) 531. J. Nowotny, M. Sloma and W. Weppner, Solid State Ionics, 28-30, (1988) 1445. Shyh-Lung Hwang and I-Wei Chen, J. Am. Ceram. Soc., 73 (1990) 3269. D.R. Clarke, J. Am. Ceram. Soc., 70 (1987) 15. F.F. Lange, J. Am. Ceram. Soc., 72 (1989) 3. Phase Diagrams for Ceramists, E.M. Levin et al. (eds), American Ceramic Society, Columbus Ohio, 1964- 1985. S.P.S. Badwal and A.E. Hughes, J. Eur. Ceram. Soc., 10 (1992) 115.

234 18 19 20 21 22

23 24 25 26 27

28 29 30 31 32 33 34 35 36 37

38

39 40 41 42 43 44 45 46

47

48

E.P. Buffer and A.H. Heuer, J. Am. Ceram. Soc., 68, (1985) 197. S.P.S. Badwal and J. Drennan, J. Mater. Sci., 22 (1987) 3231. B.V. Narashima Rao and T.P. Schreiber, J. Am. Ceram. Soc., 65 (1982) C44. R. Chaim, M. Riihle and A.H. Heuer, J. Am. Ceram. Soc. 68 (1985) 427. M.L. Mecartney and Y.J. Lin, Grain Boundary Chemistry and the Properties of Zirconia Ceramics in: Microbeam Analysis - 1991, D.G. Howitt (ed.), San Francisco Press Inc., San Francisco 1991, p. 163. A.E. Hughes and B.A. Sexton, J. Mater. Sci., 24 (1988) 1057. S.P.S. Badwal and A.E. Hughes, Proc. 2nd Int Syrup. on Solid Oxide Fuel Cells, 1991, EUR1991, EUR13564, pp.445-54. M. ArfeUi, G.M. Ingo and G. Mattogno, Surf. Int. Anal., 16 (1990) 452. J.F. Shackelford, P.S. Nicholson and W.W. Smeltzer, Ceram. Bull., 53 (1974) 865. P. Wynblatt and R.C. Mc Cune, Surfaces and Interfaces in Ceramic and CeramicMetal Systems, Materials Science Research Vol. 14, J. Pask and A. Evans (eds.), Plenum Press, New York, 1981, p. 83. R.C. McCune, W.T. Donlon and R.C. Ku, J. Am. Ceram. Soc., 69 (1986) C-196. R.C. McCune and P. Wynblatt, J. Am. Ceram. Soc., 66 (1983) 111. M. Guttmann, Surf. Sci., 53 (1975) 213. M.S. Seah, J. Phys F: Metal Phys., 10 (1980) 1043. S. Hofmann, J. Chim. Phys., 84 (1987) 141. K.L. Kliewer and J.S. Kohler, Phys. Rev., 1411 (1965) A1226. R.B. Poeppel and J.M. Blakely, Surf. Sci., 15 (1969) 507. B.C.H. Steele and E.P. Buffer, Brit. Cer. Soc. Proc., 36 (1985) 45. L. Heyne, Interfacial Effects in Mass Transport in Ionic Solids in: Mass Transport in Solids, F. Bernie and C.R.A. Catlow (eds.), Plenum Press, 1983, pp. 425-456. J. Nowomy, Surfaces and Grain Boundary Segregation in Metal Oxides in: Proc. NATO ASI Ser., Ser E, Surface and Interface of Ceramic Materials, Kleewer Acad., 1989, pp. 205-239. M.F. Yan, R.M. Cannon and H.K. Bowen, Space Charge Contribution to Solute Segregation Near Grain Boundaries, in: Grain Boundaries in Semiconductors, H.J Pike, G.E. Seamy and C.H. Leamy (eds.), Mat. Res. Symp. Ser., 5 (1982) 57. S.B. Desu and D.A. Payne, J. Am. Ceram. Soc., 73 (1990) 3391. S.B. Desu and D.A. Payne, J. Am. Ceram. Soc., 73 (1990) 3398. N.J. Shaw, Pow. Metall. Int., 21(3) (1989) 16. N.J. Shaw, Pow. Metall. Int., 21(5) (1989) 31. N.J. Shaw, Pow. Metall. Int., 21(6) (1989) 25. R.M. German, Liquid Phase Sintering, Plenum Press, New York, 1985. M.F. Ashby, Acta metall., 22 (1974) 275. R.J. Brook, Controlled Grain Growth in: Ceramic Fabrication Processes, Treatise on Materials Science and Technology, Vol. 9, F.F.Y. Wang (ed.), Academic Press Inc., New York 1976, p.331. G.S.A.M. Theunissen, Microstructure, Fracture Toughness and Strength of (Ultra)FineGrained Tetragonal Zirconia Ceramics, Ph. D Thesis (1991), University of Twente, Twente, The Netherlands. A.J.A. Winnubst, G.S.A.M. Theunissen, W.F.M. Groot Zevert and A.J. Burggraaf, Sci. 53

235

49 50 51 52 53 54 55 56 57 58 59

60 61 62 63 64 65 66 67 68 69 70 71 72 73

74 75 76 77 78

79

Ceram., 14 (1988) 309. G.S.A.M. Theunissen, A.J.A. Winnubst and A.J. Burggraaf, J. Eur. Ceram. Soc., 9 (1992) 251. T.G. Nieh and J. Wadsworth, Mat. Res. Soc. Symp. Proc., 196 (1990) 331. P.W. Atldns, Physical Chemistry, Oxford University Press, 1978, p. 193. J.W. Cahn, Acta metall., 10 (1962) 789. J.W. Cahn and R.W. Balluffi, Scripta metall., 13 (1979) 499. R.J. Brook, Scripta metall., 2 (1968) 375. G.S.A.M. Theunissen, A.J.A. Winnubst and A.J. Burggraaf, J. Mater. Sci. Letts., 8 (1989) 55. J.T. Grant, Surf. Int. Anal., 14 (1989) 271. M.P. Seah, Surf. Int. Anal., 9 (1986) 85. M.P. Seah, Electron Micros., II (1983) 521. S.P.S. Badwal, J. Drennan and A.E. Hughes, Segregation in Oxygen-Ion Conducting Solid Electrolytes and its Influence on Electrical Properties, in: Science of Ceramic Interfaces, J. Nowomy (ed.), Elsevier Science Publishers (Amsterdam) 1991, pp. 227285. A.E. Hughes and S. Rajendran, Mater. Forum., 13 (1989) 303. G.M. Ingo, G. Righini and L. Scoppio, App. Surf. Sci., 55 (1992) 257. A.I. Leonov, Yu. P. Kostikov, I.K. Ivanov, N.S. Andreeva, E.M. Trusova, Izv. Ak. Naul SSSR, Neorg. Mat., 16 (1980) 1076. V.N. Strekalovskii, Yu. N. Makurin, E.G. Vovkotrub, M.G. Chudinov and V.P. Gorelov, Izv. Ak. Nauk SSSR, Neorg. Mat., 21 (1985) 589. C. Morant, J.M. Sanz, L. Gal~, L. Soriano and F. Rueda, Surf. Sci., 218 (1989) 331. C.O. De Gonz(dez and E. A. Garcfa, Surf. Sci., 193 (1988) 305. D. Majumdar and D. Chatterjee, J. App. Phys., 70 (1991) 988. G. Marletta and S. Pignataro, Chem. Phys. Leas., 124 (1986) 414. D. Steele and B.E.F. Fender, J. Phys. C: Solid State Phys., 7 (1974) 1. J.B. Cohen, M. Morinaga and J. Faber Jr., Solid State Ionics, 3/4 (1981) 61. C.R.A. Catlow, A.V. Chadwick, G.N. Greaves and L.M. Moroney, J. Am. Ceram. Soc., 69 (1986) 272. C.B. Azzoni and A. Paleari, Phys. Rev. B., 40 (1989) 6518. V.M. Orera, R.I. Merino, Y. Chen, R. Cases and P.J. Alonso, Phys. Rev. B, 42 (1990) 9782. A.E. Hughes, F. Ciachi and S.P.S. Badwal, Surface Transformation of Y-TZP in Moist Environments in: Science and Technology of Zirconia V, Technomic Publishing Co., Lancaster, USA, to be published. H.G. Scott, J. Mater. Sci., 10 (1975) 1527. M. Yoshimura, Ceram. Bull., 67 (1988) 1950. M. Nauer and C. Carry, Mat. Sci. Forum., 94-96 (1992) 871. R.C. McCune, Mat. Res. Soc. Sym. Proc., 48 (1985) 27. The author could find no data to substantiate this claim. However, solid solution of SiO2 in ZrO2 does not exceed 0.1%. (see ref. 16, Fig. 2400) and similar solubility may well be expected for Y203-ZrO2. C. Pellet, C. Schwebel and P. Hesto, Thin Solid Films, 175 (1989) 23.

236 80

81 82 83 84 85 86 87 88 89 90 91 92

93 94 95 96

97 98 99 100 101

102 103 104

C. Pellet, C. Schwebel, P. Legagneux and J. Siejka, Yttria Stabilized Zirconia Heteroepitaxy on Silicon by Ion Beam Sputtering, in: Heterostructures on Silicon: One Step Further with Silicon, Y.I. Nissim and E. Rosencher (ed.), Kluwer Academic Publishers, 1989, p.335. M.J. Verkerk, B.J. Middlelhuis and A.J. Burggraaf, Solid State Ionics, 6 (1982) 159. K. Schindler, D. Schmeisser, U. Vohrer, H.D. Weimhtifer and W. G6pel, Sens. Actuat., 17 (1989) 555. A.J. B urggraaf, M. van Hemert, D. Scholten and A.J.A. Winnubst, International Symposium on the Reactivity of Solids, 10~ Proceedings, 1985, p.797. A.J.A. Winnubst, P.M. Kroot and A.J. Burggraaf, J. Phys. Chem. Solids, 44 (1983) 955. S.P.S. Badwal, J. Drennan, A.E. Hughes and B.A. Sexton, Mat. Sci. Forum., 34-36 (1988) 195. T. van Dijk and A.J. Burggraaf, Phys. State Solidi (A)., 63 (1981) 229. M.J. Verkerk, A.J.A. Winnubst and A.J. Burggraaf, J. Mater. Sci., 17 (1982) 3113. M.V. Inozemtsev and M.V. Perfil'ev and V.P. Gorelov, Electrokhimiya, 12 (1976) 1231. Y. Suzuki, T. Takahashi and N. Nagae, Solid State Ionics, 3-4, (1981) 483. G.S.A.M. Theunissen, A.J.A. Winnubst and A.J. Burggra~, J. Mat. Sci., 27 (1992) 5057. F.F. Lange, J. Am. Ceram. Soc., 69 (1986) 240. I.G. Lee and I-Wei Chen, Sintering and Grain Growth in Tetragonal and Cubic Zirconia, in: Sintering '87, Proc. 4~ International Symposium on Science and Technology of Sintering, S. S6miya, M. Shimada, M. Yoshimtwa and R. Watanabe (eds.), Elsevier Applied Science, Tokyo, 1987, pp. 340-345. T. Sakuma and Y. Yoshizawa, Mat. Sci. Forum., 94-96 (1992) 865. R. Ruh, K.S. Mazdiyasni, P.G. Valentine and H.O. Bielstein, J. Am. Ceram. Soc., 67 (1984) C190. T. Chaffer, T. Gervais, L. Chermant, J.L. Chermant and M. Coster, J. Eur. Ceram. Soc., 10 (1992) 299. F.F. Lange, D.B. Marshall and J.R. Porter: Ultrastructure Processing of Adavnced Ceramics, J.D. Mackenzie and D.R. Ulrich (eds.), John Wiley & Sons, New York, 1988, p. 159. T. Stoto, M. Nauer and C. Carry, J. Am. Ceram. Soc., 74 (1991) 2615. A.E. Hughes and S.P.S. Badwal, Mat. Forum, 15 (1991) 261. F. Wakai, S. Sakaguchi, K. Kanayama and H. Kato, Adv. Ceram. Mater., 1 (1986) 259. Y. Maehara and T.G. Langdon, J. Mater. Sci., 25 (1990) 2275. F. Wakai, S. Sakaguchi, K. Kanayama, H. Kato and H. Onishi, Ceramic Materials and Components for Engineers, W. Bunk and H Hausner (eds.), Deutsche Keramische Gesellschaft, Bad Honnef, 1986. F. Wakai and H. Kato, Adv. Ceram. Mater., 3 (1988) 71. T.G. Nieh, C.M. McNally and J. Wadsworth, Scripta metall., 23 (1989) 457. T. Hermansson, H. Swan and G.L. Dunlop, Euro-Ceramics, G. de Eith, R.A. Terpstra and R. Metselaar (eds.), Elsevier Applied Science, London, 1989, vol. 3, 3.329-3.333.

237 105 106 107 108 109 110 111 112 113

114

115

116 117 118 119 120 121 122 123 124 125 126 127

128 129

130

T.G. Nieh, C.M. Tomasello and J. Wadsworth, Mat. Res. Soc. Symp. Proc., 196 (1990) 343. T.G. Langdon, Superplasticity in Aerospace II, T.R. McNelley and H.C. Heikkenen (Eds.) The Minerals, Metals & Materials Society, 1990, pp.3-18. Y. Yoshizawa and T. Sakuma, J. Am. Ceram. Soc., 73 (1990) 3069. T.G. Nieh, D.L. Yaney and J. Wadsworth, Scripta metall., 23 (1989) 2007. J. Wang and R. Stevens, Br. Ceram. Proc., 42 (1989) 167. T. Sato and M. Shimada, J. Am. Ceram. Soc., 67 (1984) C212. M.T. Hernandez, J.R. Jurado, P. Duran and J.L.G. Fierro, J. Am. Ceram. Soc., 74 (1991) 1254. P.J. Whalen, F. Reidinger and R.F. Antrim, J. Am. Ceram. Soc., 72 (1989) 319. K. Tsukuma, Y. Kubota and T.Tsukidate in: Science and Technology of Zirconia II, N. Claussen, M. Rhiile and A.H. Heuer (eds.), American Ceramics Society, Columbus OH., Adv. Ceram., 12 (1984), 382. M Watanabe, S.Iio and I. Fukuura in: Science and Technology of Zirconia II, N. Claussen, M. Rhiile and A.H. Heuer (eds.), American Ceramics Society, Columbus OH., Adv. Ceram., 12 (1984) 391. M. Matsui, T. Soma and I. Oda in: Science and Technology of Zirconia II, N. Claussen, M. Rhiile and A.H. Heuer (eds.), American Ceramics Society, Columbus OH., Adv. Ceram., 12 (1984) 371. J. Wang and R. Stevens, J. Mater. Sci. Letts., 8 (1989) 1195. T.T. Lepist6 and T.A. M/i.ntyl/i, Ceram. Eng. Sci. Proc., 1O (1989) 658. F.F. Lange, G.L. Dunlop and B.I. Davis, J. Am. Ceram. Soc., 69 (1986) 237. S-Y. Chen and H-Y. Lu, J. Mater. Sci., 24 (1989) 453. T. Sato and M. Shimada, Br. Ceram. Proc., 38 (1986) 151. H-Y. Lu and S-Y. Chen, J. Am. Ceram. Soc., 70 (1987) 537. A.E. Hughes and S.P.S. Badwal, Solid State Ionics, 46 (1991) 265. F.F. Lange, J. Mater. Sci., 17 (1982) 225. A.J.A. Winnubst and A.J. Burggraaf in: Science and Technology of Zirconia III, Adv. Ceram. 24A (1988) 39. T. Sato and M. Shimada, J. Am. Ceram. Soc., 68 (1985) 356. G.M. Ingo, G. Mattogno, N. Zacchetti, P. Scardi and R. Dal Maschio, J. Mat. Sci. Letts., 1O (1991) 320. M. T. Hemandez, J.R. Jurado and P. Duran: XPS Studies to Elucidate the Aging Behaviour of Y203-TZP-CeO2 Ceramics, in: Ceramics Today -Tomorrow's Ceramics, Materials Science Monographs 66B, P. Vincenzini (ed.), Elsevier Science Publishers B.V., Amsterdam, 1991, p.649-657. M. Yoshimura, T. Noma, K. Kawabata and S. S6miya, J. Mater. Sci. Letts., 6 (1987) 465. T. Shigematsu, Y. Nakao and N. Nakanishi: Effect of Water on the Tetragonal to Monoclinic Phase Transformation of ZrO2 with Y~O3, in Shape Memory Materials, vol. 9, M. Doyama, S. Sfmiya and R. Chang (eds.), Materials Research Society, Pittsburgh, 1989, pp. 549-554. N. Narita, S. Leng, T. Inada and K. Higashida: Enviromental Effect on Phase Stability in Y-TZP Ceramics, in: Sintering '87, Proc. 4 ~ International Symposium on Science

238

131

132 133 134 135 136

137

138 139 140 141 142 143 144 145 146 147 148 149 150 151 152 153

and Technology of Sintering, S. SSmiya, M. Shimada, M. Yoshimura and R. Watanabe (eds.), Elsevier Applied Science, Tokyo, 1987, pp. 1130-1135. P. Kountouros, H. Shubert and G..Petzow, Low Temperature Degradation and Mechanical Properties of TiO2-Y203-Stabilized Tetragonal Zirconia Polycrystals in: Science and Technology of Zirconia V, Technomic Publishing Co., Lancaster, USA, to be published. S.P.S. Badwal and N. Nardella, Appl. Phys. A, 49 (1989) 13. H.F. Holmes, E.L. Fuller, Jr. and R.B. Gammage, J. Phys. Chem., 76 (1972) 1497. W.D. Harkins and G.E. Boyd, J. Amer. Chem. Soc., 64 (1942) 1195. S. Meriani, Mat. Sci. Eng., AI09 (1989) 121. A. Laachir, V. Perrichon, A. Badri, J. Lamotte, E. Catherine, J.C. Lavalley, J. E1 Fallah, L. Hilaire, F. le Normand, E. QuemEr6, G.N. Sauvion and O. Touret, J. Chem. Soc. Far. Trans., 87 (1991) 1601. M. T. Hernandez, J.R. Jurado and P. Duran: Processing, Characterization and Sintering Behaviour of Y203-TZP and Y203-TZP-CeO2 Ceramics, in: Ceramics Today Tomorrow's Ceramics, Materials Science Monographs 66B, P. Vincenzini (ed.), Elsevier Science Publishers B.V., Amsterdam, 1991, pp.1395-1404. K. Heussner and N. Claussen, J. Am. Ceram. Soc., 72 (1989) 1044. D.J. Bevan, J. Inorg. Nucl. Chem., 4 (1955) 49. J-G Duh, H-T Dai and B-S Chiou, J. Am. Ceram. Soc., 71 (1988) 813. F. Le Normand, J. E1 FaUah, L. Hillaire, P. L6gar6, A. Kotani and J.C. Parlebas, Sol. State Comm., 71 (1989) 885. S. Imada and T. Jo, J. Phys. Soc. Jap, 58 (1989) 2665. J.Z. Shyu, K. Otto, W.L.H. Watldns, G.W. Graham, R.K. Belitz and H.S. Gandhi, J. Catal., 114 (1988) 23. M. Hoang, A.E. Hughes and T.W. Turney, submitted to Appl. Surf. Sci. P. Duran, M. Gonzalez, J.R. Jurado and C. Moure, Ceram. Trans., 12 (1990) 945. A.V. Virkar and R.L.K. Matsumoto, J. Am. Ceram. Soc., 69 (1986) C224. G.M. Ingo, J. Am. Ceram. Soc., 74 (1991) 381. T.Tsukuma, K. Ueda and M. Shimada, J. Am. Ceram. Soc., 68 (1985) C4. S. Rajendran, M.V. Swain and H.J. Rossell, J. Mater. Sci., 23 (1988) 1805. T.G. Nieh and J. Wadsworth, Acta metall, mater., 39 (1991) 3037. E.P. Buffer and J. Drennan, J. Am. Ceram. Soc., 65 (1982) 474. J. Drennan and S.P.S. Badwal, Science and Technology of Zirconia III, Adv. Ceram. 24B (1988) 807. S. Rajendran, J. Drennan and S.P.S. Badwal, Mat. Sci. Letts., 6 (1987) 1431.

Science of Ceramic Interfaces II J. Nowomy (Editor) 9 1994 Elsevier Science B.V. All rights reserved.

239

Important role of the interfaces in the high temperature superconductors S.X. Dou and H.K. Uu School of Materials Science and Engineering, University of New South Wales, Kensington, NSW, 2033 Abstract The composition, structure, orientation and properties of the interfaces play a crucial role in determining the properties of the HTSC materials. The interfaces act as Josephson weak links, and in effect control the Jc in random oriented polycrystalline materials. Chemical examination on the grain boundaries of the HTSC shows that the composition even in the clean boundaries deviates from the interior of the grains. Some high angle boundaries show a flux pinning characteristic. Studies on the grain boundary structures reveal that the lattice matches can be described by a coincident site lattice (CSL) model or a constrained CSL model for YBa2C3OT_~(YBCO). Most of the grain boundaries in the textured Bi-Pb-Sr-Ca-Cu-O(BPSCCO) are parallel to the basal plane, while they are perpendicular to the basal plane in YBCO. The impressive high critical current densities Jc obtained for the Ag-clad Bibased superconducting tapes are of great interest for high current and high field applications, and hence has attracted more attention to the grain boundary studies in these materials. A strong correlation between the high Jc and textured platelike grain morphology has suggested the "brickwall" model. Most of the grain boundaries in the textured BSCCO are twist boundaries in (001) with c-axis as rotation axis. Due to the mica-like layered structure, most of the boundaries are in low-angle boundary region. Grain boundary dislocations are commonly observed in these twist boundaries with high density. The presence of such a high density of the dislocations is probably due to the deviation from misorientation between grains in order for the interface to retain the low energy configuration over most of its area. Recently, studies on the transport and magnetisation Jc, and the irreversibility behaviour of the textured Ag-clad Bi-based tapes show no evidence of grain boundaries for being the weak links and the overall Jc is controlled by the intragrain pinning at high temperatures. Weak link behaviour becomes evident at low temperatures. The interfaces between the HTSC oxides and the silver sheath exhibit unusual mechanical properties that are responsible for the improved strength and flexibility of the Ag/HTSC composites.

240

1. INTRODUCTION

Interface and grain boundary problems have been the central research topic in the high temperature superconductors(HTSC). It is well established that grain boundaries act as Josephson weak links resulting in low critical current densities in HTSC, which are the major barrier for realising practical applications of these materials. Despite enormous advances in various aspects of HTSC through intensive research efforts, the grain boundary problems remain as a great challenge to the research community from the viewpoint of both the electromagnetic properties and microstructures of the grain boundaries. These problems are characteristic of ceramics and, in superconductors, exist as a result of their granular nature and high critical transition temperature itself. Thus, to understand why the interface problem is particularly important and why the study of the interface problems is so challenging we must first address the special properties of the HTSC. Extensive research work has revealed that the known copper-oxide superconductors all have the following features in common: (1) intergrowth layer structures, which lead to two dimensionality and strongly anisotropic properties [1,2]; (2) mixed valencies, which are the origin of the charge carriers, which have a very low density compared to conventional superconductors [3]; and (3) extremely short coherence lengths and large penetration depths. One of the consequences of these inherent properties is weak links at grain boundaries, which lead to very low intergrain current densities since the boundaries act as easy paths for Josephson vortices. So the critical current derlsity(J~ will be determined by the pinning of these vortices. The superconducting coherence length for the HTSC is 2 nm in the a-b direction and 0.3 nm along c direction. This extremely short coherence length defines the maximum grain boundary thickness through which the supercurrent can pass without showing Josephson weak link characteristics. This also defines the resolution for microstructural and electromagnetic characterisation. The large anisotropy imposes even more severe requirements on the spatial resolution. Another problem is that, in high-T c superconductors, the magnetic flux line creep is unusually rapid. As the "giant flux creep" proceeds, the critical current density declines. The flux creep in effect controls the size of the critical current density [4]. Especially, the giant flux creep dominates the magnetic behaviour of the Bi-Sr-Ca-Cu-O (BSCCO) even more strongly than YBa2Cu3Oz_` (YBCO) [5] since the BSCCO systems are more anisotropic than YBCO. BSCCO exhibited a wide reversible regime [6] and low melting temperature of the Abrikosov Lattice [7]. In the conventional alloy low temperature superconductors, grain boundaries and interfaces act as effective pinning centres since the coherence length for these materials is one or two orders of magnitude greater than that for HTSC. To pin the flux lines in place remains an unsolved question since superconductivity of oxide high-T c materials is much more serlsitive to small-scale structural and chemical imperfections than in conventional superconductors. To give rise to an appreciable pinning strength, it is necessary to introduce pinning centres the size of the dimension of a coherence length. Otherwise, these defects such as grain boundaries and interfaces may well act in decoupling and weak link roles. Thus, it

241

is a great challenge to design interfaces and grain boundaries as effective pinning sites rather than weak links. 2. WEAK UNKS AT GRAIN BOUNDARIES 2.1 IMPURITY IN THE BOUNDARIES The typical thickness of the grain boundaries varies from subnanometer to several ten nanometres depending on the processing procedure. Grain boundaries with impurities or glass phase were often observed in the bulk HTSC materials with a thickness over 10 nm for a random oriented polycrystalline YBa2Cu3OT_x(YBCO) sample. It is obvious that this type of grain boundaries will act as weak links rather than pinning centres. As the materials processing technique is improved, the bulk materials with clean boundaries have been fabricated. But the critical current densities of the bulk materials with clean grain boundaries are still two orders of magnitude smaller than that for single crystals. 2.2 GRAIN MISORIENTATION The grain misorientation is largely responsible for the drastic degradation in the Jc. In an elegant bicrystal experiment, Dimos et al have clearly shown that the transport Jc cross the boundary degrades as the misorientation increases, suggesting that high-angle grain boundaries are generally deleterious to the overall Jc [8,9]. Figure 1 shows the ratio of the Jc in the grain boundaries to that in grains versus the misorientation angle e in the basal plane. Dimos et aJ. identified two regimes of the behaviour: (1) boundaries with very small misoirentations(_< 5~ that had little effect on the Jc value and (2) boundaries with misorientation angles >_100 that had strongly reduced Jc values and showed a Josephson weak link behaviour. The degradation of Jc for the misoriented samples is significantly enhanced in magnetic field as reported by Larbalestier et al.[10]. It is noted that Jc across a 22 ~ [001] Josephson junction-coupled bicrystal declines nearly one order of magnitude in weak fields of less than 50 Gauss as shown in Figure 2. Such Jc characteristics are clearly serious problem for high-current and high-field applications. It is evident that grain misorientation is one of the major obstacles for large scale applications of the HTSC. 2.3 CHARGE DISTRIBUTION IN THE BOUNDARIES

In addition to the grain boundary misorientation, some other mechanisms that may cause weak link boundaries have also been proposed. For example, fluctuation in the charge distribution along the grain boundaries as a result of low charge carrier density in the HTSC may form barrier for the transport current[3]. The low carrier density leads directly to a decreased capacity to electrically screen the net charge in the materials. Accumulation of net charge tends to occur at grain boundaries resulting from the deviation from the bulk composition or the segregation of the impurity ions in the vicinity of the interface. Therefore, the potential barrier at the boundaries could prevent the free transmission of current,

242

especially in low carrier density materials. This phenomenon has been used for making ceramic varistors and positive temperature coefficient thermistors. In the HTSC materials, the significance of the boundary potential barrier was first pointed out by Enomoto et al. [11] in Ba(Pb,Bi)O3(BPBO) based on the observation of Josephson junction characteristics across grain boundaries in the V-I relation. The same situation was found in the HTSC[12]. Since the charge carrier density in HTSC may not be increased significantly, more studies directed toward reducing the net charge at grain boundaries are needed.

12

L

l= 0 l - -a

=

09

.,o L

06

~

03

--'

0.1 0 0 0

qb V

E

0.01

z

O0

-200

0

Applied Field, 0.001

200

Gauss

|,,,,i||iI|,,,iiiiii||,1111111|,|,I,iIiiI|,,

0

10

20

0

30

40

Fig. 1. The ratio of J= in grain boundaries to that in grains vs the misorietation angle O in the basal plane( courtesy of D. Dimos et al.19D

Fig. 2. Normalised J=(H)characteristics of a 22" [001] Josephson junction-coupled bicrystal( courtesy of Larbalestier et al.[lO D

2.4 COMPOSITION CHANGE IN THE BOUNDARIES

Composition variation in the grain boundaries has been observed in HTSC and is another important factor responsible for the weak link behaviour of the grain boundaries. In our early work we showed that the YBCO single crystal was rich in Cu towards grain boundary(Figure 3)[13]. The compositional analyses on a YBCO single crystal were performed with a TEM using EDS. An electron transparent thin film was ion-milled. A typical crystal of approximately 0.2 mm 2 was systematically probed in the horizontal and vertical directions in order to observe possible change in the cation ratios of Cu/Y and Cu/Ba. The central area was found to consist of the cations in ratios reasonably consistent with stoichiometry of Y:Ba:Cu = 1:2:3, whereas changes in stoichiometry appeared away from the centre. The composition change towards grain boundaries was further studied in more details by a number of groups. Dimos et al. associated the weak link behaviour of the epitaxial thin film grain boundaries with density of grain boundary dislocations, because both show the same dependence on the misorientation angle[9]. Since the dislocation cores are nonsuperconducting, the overlap of these dislocation cores at the critical angle value of 10~ results in the Josephson weak link

400

243

behaviour. Gao et a1.[14] have proposed that copper oxide grain boundary cores are nonsuperconducting, which is in good agreement with the dislocation model and the change in cation composition near the boundaries. Recently, Babcock et al. have confirmed that a 26 ~ [001] Josephson-coupled boundary appears to be rich in Cu relative to the grain interiors as deduced from the scanning TEM equipped with EDS[15]. However, they argued that the measured excess Cu concentration corresponds to an approximately 0.08 nm thick layer of copper oxide, so it is unlikely that a thin, wetting layer of this oxide is fully responsible for the grain boundary's Josephson weak link behaviour.

Elemental Ratio Horizontal Line

Elemental Ratio

YBa=CusOx

[ 4t

V, Vertical Line

i

-zOO -,oo o .,oo .zoo Distance from Centre (nm) J

l

-zoo - , ~ ' o .,~ .zoo Distance from Centre (nm) L

&

i

9

9

Fig. 3. Cation compositional distribution in the YBCO crystal It is known that oxygen in HTSC is highly mobile and superconductivity is sensitive to the change of oxygen content. The change of oxygen concentration in the grain boundaries may be an important origin of the weak link behaviour. Zhu et al. have studied the misorientation of over 200 pairs of adjacent grains separated by grain boundaries in melt-textured YBCO[16]. They found that some boundaries may adjust the length of the c-axis by controlling oxygen content to produce a favourable condition to meet the lattice matching requirements. Thus, some grain boundaries are "naturally oxygen deficient" in a fully oxygenated sample and such boundaries will likely act as weak links. 2.5 PHASE CHANGE NEAR THE BOUNDARIES

Phase change near the interfaces and boundaries has been observed in the HTSC materials. This is particularly important for the systems with multiphase, for example, the Bi-Sr-Ca-Cu-O HTSC compounds have three phases: 2201 (7-22K), 2212(85K-92K) and 2223(110K)[17-19]. Substitution of Bi by Pb stabilises the

244

2223 phase[20-22]. Ramesh et al. have examined the grain boundaries of the leaded and unleaded Bi-based samples[23]. They found that for the unleaded samples, the structure inside the grains consisted mainly of 2223 polytypoid, whereas upon approaching the grain boundaries the structure changed to the 2212 polytypoid and finally adjacent to the boundary the 2201 polytypoid was observed. In the leaded samples, the grain boundary microstructure is very different. In the samples with the optimum composition, the 2223 polytypoid is continuous up to the grain boundary interface. The difference in the grain boundary microstructure is reflected in the resistive transition. The resistive transition for the leaded sample is in general sharp, free of step, while the step in the resistivity plot is always observed for the unleaded sample. A similar situation has been found in the silver/superconductor interface where the low T c phase Bi-2212 or Bi-2201 was detected adjacent to the interface in the Ag-sheathed Bi-based HTSC composites. For YBCO, it was found that the orthorhombicity of the interface layer differed greatly from the bulk of the grain. The surface layer is much less orthorhombic than the bulk as shown in Figure 4[24]. The reduced orthorhombicity of the surface layers can be regarded as a universal characteristic of crystalline YBCO. The reduce orthorhombicity of the surface layers may indicate a lower oxygen content at the interface than the bulk. 3. GEOMETRICAL MODELS FOR BOUNDARIES 3.1 FLUX PINNING BOUNDARIES As described in the previous section, grain boundaries, in particular, the high angle boundaries in polycrystalline HTSC act as Josephson weak links due largely to the short coherence length and high anisotropy. However, the impressive high Jc values in high fields has been achieved in the Ag-sheathed Bi-based HTSC tapes[25-29]. These results suggest that Josephson weak links at the grain boundaries can be minimised. In particular, Babcock et al. directly demonstrated that some large-angle boundaries in bicrystals of YBCO transport high current as shown in Figure 5[30]. It is interesting to note that the Jc - H characteristic of all three bicrystals with a grain boundary angle of 3~ 140 and 90 ~ were relatively flat for applied fields of up to 5 T, which approaches the effective upper critical field for H parallel to the c-axis at 77 K. It is also noted that this Jc - H behaviour is similar to that of single crystal. Although the grain boundary angles 14~ and 90 ~ should belong to high angle boundaries, they do not show a sign of weak link behaviour. These boundaries are considered as flux pinning characteristic rather than Josephson weak link behaviour. These flux pinning boundaries were also found in highly-textured polycrystalline thin film YBCO[31]. The weak link free and flux pinning boundaries are of interest from the viewpoints of both fundamental research and technical applications. These results also point out the importance of the structural characterisation of the grain boundary in HTSC. 3.2 GEOMETRICAL MODEL FOR LATTICE MATCH

245

It is well established that the Coincident Site Lattice(CSL) model has been successfully used for characterisation of the boundary structure of the cubic materials. According to this model, when lattices of two crystals are interpenetrated each other, a relatively large fraction of the lattice sites of one crystal would be coincident with those of the other crystal. The ratio of lattice sites to the coincident sites is denoted by Z. Recently several groups have shown that the lattice match in YBCO can be described by the CSL model or "near" and "constrained" CSL model[32-34].

1

16.8

200 9 9SAME CRYSTAL 9DIFFERENT CRYSTALS

t3.9

190

tle=~ L~ mgl= =,/It=

ol

30[001]

/,,,

I-

=o~

.J

,

g

t0.8

180

7~

1to

c

=

l

-

I.-

g

r a.4

1-1t6o ,,f

O.0 - t.1

__t

~.

/

-6.2 L 0

, t

, 2

i 3

I~ 1-1 BULK -T-=I( x - KaY) H 150

't

, , , , , i t i i ~ 140 4 5 6 7 8 9 10 1t 12 13 DEPTH ETCHED ( ~ m )

Fig. 4. Lattice parameters a and b (A=b-a) from the known X-ray value Ao = 0.007 nm as a function of depth from the YBCO crystal surface (courtesy of D.J. Werder et al. [24])

\

e-

:3

o o~

a

1000 I~176176 ......

L_

o

II . . . . . . . q

~

P". . . . .

li" ......

t ....

"i~-~176176

6

8

Applied Field, Tesla

Fig. 5. Jo(H) curves for flux-grown YBCO bicrystals that show flux-pinning characteristics (courtesy of S.E. Babcock et a1.[301)

The first implication of the geometrical lattice match model is that the high degree of lattice match is always associated with a low grain boundary free energy configuration as evidenced by Smith et al.'s histogram (Figure 6) of the angular distribution of flux-grown bicrystals for a set of [001] grain boundaries[35]. It is noted that the peaks which reflect low energy misorientation at the various labelled angles all correspond to low-Z, highly lattice-matched misorientation. Zhu et al. found that more than 70% of the grain boundaries have small misorientation angles in the melt-textured YBCO. Large angle boundaries with preferred orientation observed in the melt-textured samples indicate the existence of low energy boundaries. The misorientation of these boundaries can be described with the CSL model. Another important feature of these models is the presence of grain boundary dislocation network associated with small deviations from the special, low energy

246

misorientation. Figure 7 is a TEM image of a grain boundary dislocation network observed in the melt-textured YBCO, suggesting the presence of strain field at the high angle boundaries. Although the correlation between the electromagnetic properties and these geometrical models is unclear, it is believed that the grain boundary structure models will help us understand some aspects of physical properties of the HTSC. For example, using electron energy loss spectroscopy Zhu et al. found that the grain boundary hole concentration correlates to the potential for achieving a highly-lattice-matched constrained CSL interface through tuning of the c-axis parameter. The necessary adjustment can be accomplished by local changes in the oxygen concentration. The lattice match at the 900 [010] boundary should promote full oxygen stoichiometry. Thus, the constrained CSL-based model explains the high J= at special high angle boundary [36]. Various mechanisms for explaining the flux pinning grain boundaries have been proposed but agreement has not been reached.

Fig. 6. Histogram of measured misorietation angles for 0[001] flux-grown YBCO bicrystals( courtesy of D.A. Smith [35])

Fig. 7. TEM image of a grain boundary dislocation network observed in the melttextured YBCO

4. GRAIN BOUNDARIES IN TEXTURED Bi-Pb-Sr-Ca-Cu-O 4.1 HIGH J= IN Ag/Bi-BASED HTSC TAPES Both Bi-based superconductors and YBa2CU3OT_x have layered structures, which lead to strong anisotropy of their properties. The anisotropy parameter, y, = (m=/ m,) ~, where m= and ma are the effective superconducting masses for pair motion along the c direction and the a-b plane, respectively, is reported to be in the range 25-50 for Bi-based materials [37,38], while it is estimated to be 5 for YBa2Cu3Oz_x [39]. This difference leads Bi-based materials to be considered two

247

dimensional (2D), whereas YBa2Cu3OT_x is considered three dimensional (3D). Thus, it is expected that the grain boundary weak link effect in the Bi-based materials would be more pronounced than that in YBCO. However, A significant improvement in the critical current density (Jc) and flexibility has been achieved in the Ag-clad Bi2Sr2CaCu2Os+x (2212)[40-42] and Ag-clad (Bi,Pb)2Sr2Ca2CU3Olo+x (2223)[43-46] using a powder-in-tube technique. In particular, the resistance to magnetic field in Bi-based materials has been greatly improved in the last year. 4.2 MAGNETIC FIELD DEPENDENCE OF Jc FOR THESE TAPES Critical current densities of lx104 A/cm 2 for Ag-clad (Bi,Pb)2Sr2Ca2CU3Olo+y tape [47] at 77 K and 1 T have been achieved. It is more encouraging that the Jc of Ag/BiSrCaCuO wires at 4.2 K is already comparable to conventional metallic superconductors and even superior at fields above 20 T [27]. Recently, Jc over 10,000 A/cm 2 at 77 K has been achieved for long length tapes up to 120 meters by a number of research groups. Recently, Dou et al [48-52] have reported Jc values up to 4 x 104 A/cm 2 at 77 K and zero field and 9 x 103 A/cm 2 at 77 K and 1 T for Ag-sheathed 2223 tapes prepared by using a high T c phase formation-decomposition-reformation (PFDR) process as shown in Figure 8. The PFDR processed tapes (A to D) exhibit a noticeable increase in the Jc at 77 K and 1 T over the normal processed tapes (E,F). The Jc-H behaviour of the PFDR processed tape C is better than the best results reported by Sato et al [53]. At 77 K and 1 T the Jc of the former's still holds 33% of its zero field value while the Jc of the latter's holds 23% of its zero field value. The improved Jc and Jc-B behaviour has been attributed to the improvement of grain alignment and grain boundary weak links, decrease of nonsuperconducting phases and uniform dispersion of the impurities derived from the high T c phase formation-decomposition-reformation process[52]. The microstructure of the PFDR tape C (19000 A/cm 2) shows well connected, aligned and elongated grains, with a high apparent density and no large secondary phase particles (Figure 9). A double-step characteristic in the Jc-H curves for various high T c superconductors has been observed and interpreted by IEkin et al [54]. At low magnetic field (H < 300-500 Oe ), a rapid drop in the Jc is attributed to the Josephson weak links in the grain boundaries. It is noticed, however, that the Jc did not show a large drop in low magnetic field for tape A-D in Figure 8, indicating a significant improvement in weak-link structure in these tapes. At fields higher than 500 Oe, a plateau region observed is due to the flux pinning within grains. Above 1 T the Jc at 77 K for all Ag/2223 tapes reported so far dropped rapidly due to either thermally activated flux creep or approaching upper critical field. At lower temperature thermally activated flux creep was suppressed and the Jc values greater 10s Ncm 2 at 4.2 K and 23 T have been reported for Ag-sheathed tape with thickness less than 30 Hm [55]. Osamura et al showed that at 4.2 K the Jc for a 20/Jm thick tape is one order magnitude higher than that for a 45/Jm thick tape [56]. Figure 10 shows the Jc magnetic field dependence at 4.2 K for a Agsheathed 2223 tape with a thickness of 45 pm, a Jc of 2.26 x 10s N c m 2 in 0 T and 7.0 x 104 Ncm 2 in 15 T has been achieved. This thickness dependence may be

248

attributable to the effect of self-generated fields.

Fig.

8.

J,-H

curve

(Bi,Pb)2Sr2Ca2Cu30~o,7 tapes

for Ag-clad at 77 K[49]

Fig. 9. SEM image of the melt-processed Ag-clad 2223 tape with Jc=19000 A/cm 2 [52]

At 4.2 K Ag-sheathed 2212 wires show a better Jc-B behaviour than Agsheathed 2223 wires. Heine et aJ [40] were among the first to report that Jc up to 5.5 x 104 Ncm 2 at the zero field and up to 1.5 x 104 A/cm 2 at 26 T and at 4.2 K was achieved, which is superior to the commercial NbTi, NbSn and (NbTa)3Sn superconducting wires. This significant improvement in Jc-B behaviour is attributed to the melt processing which has the following advantages over the sintering process: (i) better grain connectivity, cleaner grain boundaries and higher degree of grain alignment reducing the weak links; (ii) finely dispersed second phases acting as pinning centres, increasing the pinning strength and reducing flux creep. The Jc-B dependence of Ag-sheathed 2223 tapes is highly sensitive to the field orientation as shown in Figure 11. The Jc drops drastically when the applied field is perpendicular to the tape surface. This is essentially attributed to the anisotropic properties. For example, the upper critical field in the c-axis H% (.L) is only about 2 T at 76 K in the Bi-2212 superconductors [57]. These high current densities and excellent Jc- H behaviour strongly suggest that the weak links are diminished in the textured Bi-based HTSC tapes. On the other hand, the Jc values of highly textured YBCO prepared by the melt-texture growth technique show a much stronger field dependence[58]. This indicates that the grain boundary structure in the melt-textured YBCO could be different from that in Ag-clad Bi-based tapes. The grain boundary planes in the melt-textured YBCO are closely perpendicular to the a-b plane, while most grain boundary planes in

249

Ag-clad Bi-based tapes are parallel, or closely parallel to tha a-b plane[52,59]. This fundamental difference in the grain boundary orientation in the two type textured materials suggests that the effect of grain boundary on the transport properties could follow different models.

1.oo'"0

****

o:eo

J._32oo

II

A/cm =

N=3k

0.O0

10 s.

J,=73oo A/era .I =

t~.lrl}

0.80

0

0.70

H , tape surface, 4.2 K

0.60 0.50

10"

o

~ .........

H (T)

1o

15

Fig. 10. Transport Jc vs. H at 4.2 K for Ag/2223 tape

4.3

"BRICKWALL"

MODEL

0.40

~'

so

=oo

ANGLE (3)

15o

~oo

Fig. 11. The normalised J= versus 0, the angle between the tape surface and H, for the tapes treated at 832~ for different times

It has been realised that the impressive transport properties of the Ag-clad Bi-based superconducting wires are attributable to a desirable combination of the plate-like morphology with the thermomechanical deformation process, which results in an excellent c-axis alignment and grain connectivity between a-b planes in these wires. The large contact area between the plate-like grains substantially reduces the resistance for current flow in the c-axis direction. A "brickwall" model has been proposed to illustrate the current transport mechanism in the c-axis aligned tapes[60]. Each brick represents a crystal grain, and all the crystals are assumed oriented along a common c-axis normal to the tape plane, while the orientation of the a and b axes of the grains is random in the tape plane. The thickness of the grains is small compared with the length along the principal tape axis, while their width is comparable to the length. This model closely resembles the experimentally observed microstructure as shown in Figure 9. According to this model, the net horizontal supercurrent passes from brick to brick chiefly through the c-axis twist grain boundaries. At higher temperatures, however, the Jc of these tapes shows a rapid decline with increasing magnetic field and a pronounced anisotropy, due to the thermally activated flux creep. This suggests that at higher temperatures the transport current

250

density is not limited by the grain boundary weak links but by the intragrain currents. Thus, it is evident that the flux pinning becomes important for maintaining the high Jc of the Ag-clad Bi-based superconducting wires in magnetic field at high temperatures. 4.4 NO EVIDENCE OF WEAK LINKS OF Ag/Bi-BASED TAPES AT 77 K To study the effect of grain boundaries, the transport Jc measured by a four probe method are compared with the magnetisation Jc at 4.2 K and 77 K as shown in Figure 12. Transport Jc measures the intergranular critical currents, while the magnetisation Jc has contributions from both inter- and intra-granular critical currents. Because the magnetisation measurements usually use much severer voltage standard(10 -~~ V/cm, according to ref.[61]) than that for transport measurements, magnetisation measurements are more sensitive to the flux motion and give low values of Jc under magnetic fields. It is, therefore, necessary to compare the behaviour of these two types of Jc - B at different temperatures.

1

~ ' - " 1 0 "~

~10

~

TRANSPORT MAGNETISATION

J, J,

77K, BIIC

4.

0.! 10 4 .

o

CCCCO Jo TRANSPORT ~"~'~o~ J, MAGNETISATION

zl~ I 0.00 10

'

'

,

,

,

,

, ,

,

!

.

.

.

.

.

.

,

,

,

!

,

,

,

, ,

0.40 0.80 MAGNETIC FIELD (T)

,

, , T

1.20

0 . 0 |

,

0.00

,

,

,

,

,

,

,

,

i

,

2.00

,

,

,

,

,

MAGNETIC

,

,

,

I

,

,

,

4.00

,

,

,

FIELD

,

,

,

I

,

(T)

,

T

,

6.00

Fig. 12. Normalised Jc versus fields measured by both the transport and magnetisation procedures at 77 K(a) and 4.2 K(b) All the measurements were done with field parallel to the c-axis. At 77 K the normalised Jc determined by magnetisation drops drastically with increasing magnetic field and Jc can not be measured above 0.1 T, while the transport Jc still retains a finite value up to 1 T, which is in agreement with results reported by Maley et a1161]. At 4.2 K, the measured irreversible magnetisation AM decreases only rather slowly with field, and when translated to a critical current yields values considerably larger than the measured transport current. This is evidence that at low temperatures, AM contains a large intragranular component, arising from strong

251

flux pinning within the grains. The different behaviour at low and high temperatures suggests that the critical current is controlled by different mechanisms. At high temperatures thermally activated flux motion become pronounced, flux pinning is weak and hence intragrain current controls the J=. On the other hand, at low temperatures the flux pinning is strong due to the intrinsic pinning[62], the intragrain current is high, so the grain boundary weak links become relatively important, limiting the transport current at low fields. The effect of the weak links depends on the quality of the grain boundaries which, in turn, depends on the materials processing. On the basis of magnetic moment(m) measurements, Angadi et a1163] proposed a length scale on which the intergranular supercurrents circulate to evaluate the relative contribution of the inter- and intragranular currents to the critical currents. The measured length scale for a normal tape at 77 K is consistent with the tape dimensions(Figure 13)[64]. The length scale decreases only slightly from 3.2 mm at 5 mT to 2.8 mm at 40 mT; at this temperature the irreversibility field H,, is of the order of 100 mT. This demonstrates directly that the tapes are fully connected at high temperatures and in fields up to the irreversibility lines and the grain boundaries display no evidence of being weak-links. This situation is very different from that in polycrystalline YBCO, where just a few mT is enough to decouple the grains and so cause a drastic reduction in length scale. To investigate the grains, we cut open a tape, split the silver sheet, scraped out the powder from the tape and measured the magnetic moment hysteresis loops for the tape powder. Figure 14 compares the loops for the tape and the ex-tape powder. At all fields and temperatures, the former is nearly an order of magnitude larger than the latter, which shows immediately that the dominant contribution to the tape loop(Am) is from intergranular current. The grain boundaries show no sign of weak link behaviour. The weak link behaviour of the tapes is also studied using the irreversible J c - B loop as shown in Figure 15 for a tape at various temperatures. The loop can be ascribed to the existence of grain boundary weak links. When the applied field increases the flux density in the grain boundaries is higher than that in the grains, leading to a higher field in the grain boundaries than the average applied field. On the other hand, when the applied field decreases the flux density in the grain boundaries declines more rapidly than that in the grains since the grains have much stronger pinning than the grain boundaries, resulting in a lower field in the boundaries than the average applied field. Thus, the J= on the increasing fields is lower than that on the decreasing field if the grain boundaries act as weak links. It was found that this loop diminished for the good Ag-clad Bi-2223 tapes at 77 K, whereas it became evident at 4.2 K, suggesting that the grain boundaries do not act as weak links at 77 K but weak links are relatively pronounced at 4.2 K. Recently, we have shown that the Jc at 77 K does not drop with increasing length of the tapes up to 12 meters(Figure 16) and it is only limited by "the bottle neck" along the length of the tapes. The grain boundaries show no accumulative resistance at least up to 12 meters, suggesting that the grain boundaries are not controlling the Jr A Jr over 8,000 A/cm 2 at 77 K was achieved for a 49 filament tape. This tape was rewound into a solenoid of a diameter of 14 mm without

252

degradation of the Jc. It is more encouraging that the reproducibility of the production of long tapes is better than that for the short tapes because of the high thermal conductivity of the Ag sheath that reduces the thermal fluctuation during the heat treatment.

4-

0.03" 9

9

9

9

5OK

9

A

E

oi

T/K r-

o

o 5 lb ~'s 2o 2's ~

~

field (roT)

4b 4'5-so

Fig. 13. Variation of the length scale with field at 77 K for a normal tape[64]

:X)O

10000

0

10000

20000

H (Oe)

Fig. 14. m-H loops at 50 K for ex-tape powder and Ag/Bi-2223 tape with H perpendicular to tape plane[64]

8.0 -

10'

6.0

(',4 ~10 0

S4492152 Bllab ..... 4.2k 20k ..... 40k 60k ..... 77k 84k .... 98 6 k 96k

4.

< 0 10 ~

~o' ~ " 6 " ' ~ ' " ~ "

4.0

with

length

2.0

' ~ ' " ~ " ','o'' ','i" '~','' ','6'"

8(T)

Fig. 15. Jr history effect at various temperatures with H parallel to the tape plane

length (M)

Fig. 16. Variation of the Jc along the length of a tape

253

Flux creep effect has been recognised to be much more important in the Bibased materials than in YBCO. The temperature dependence of the electrical resistivity at low current density for 2212 and YBCO single crystals has been measured [65]. It was found that at H > 1 T the resistivity of Bi-single crystal was higher than Cu in the temperature range above 30 K. However, the achievement of high Jc in Ag/2223 wires at 77 K and 1 T demonstrates that the flux creep can be pinned by improving the flux pinning [50]. Hampshire et al have studied the Jc-B and Jc-T behaviour of Ag/2223 wires at 10 K intervals from 4.2 K to 77 K in fields 1 to 20 T [66]. They found that there was no evidence of any catastrophic change in Jc (H), in contrast to the prediction by flux line lattice (FLL) melting experiments [67], suggesting that FLL melting is not controlling the Jc of these tapes and FLL melting can be delayed to higher fields and/or temperatures by improving flux pinning. 4.5 GRAIN B O U N D A R Y STRUCTURE IN TEXTURED BSCCO

Although studies on the grain boundary weak link behaviour in the textured BSCCO has been carried out, the properties of the grain boundary structures in these materials are poorly understood compared with those for YBCO, due mainly to the difficulties for obtaining high quality bicrystals of these materials. Recently, several groups have studied the twist boundaries, grain boundary dislocations, high angle and low energy boundaries in the textured BSCCO[23,28,48,52,59,68,69]. Some important points are summarised as follows. TEM studies reveal that the controlled melt-textured tapes contain clean boundaries, leading to a strong bonding between grains and responsible for the slow drop of J= with magnetic field. The low critical current densities in the bulk Bibased materials are mainly attributed to the weak links in the grain boundaries where the impurities, pores and glass phases are often observed. Most of the grain boundaries in the textured BSCCO are twist boundaries, where successive layers have the a- and b-axes interchanged as shown in Figure 17. The twist boundary plane is on (001) with c-axis as rotation axis. Due to the mica-like layered structure, most of the boundaries are in low-angle boundary region with a rotation angle less than 10~ The grain boundary planes appear to be atomically fiat. Grain boundary dislocations are commonly observed in these twist boundaries as shown in Figure 18. The presence of grain boundary dislocations is due to the weak Van der Weals bonding between adjacent BiO layers. Cleavage between BiO layers and sliding of the BiO layers over each other occur easily. These displaced layers form dislocations in the interface regions. Figure 19 shows a set of nearly parallel dislocations on one side of the grain boundary. The density of dislocations is still high around 1 x 101~ 2 although the sample was annealed for a long period at 835~ The background "tweed" texture between dislocation lines may be a result of precipitation of the excess components from decomposition of the 2223 during the short period of melt process. Some dislocations are tangled documenting a certain degree of recovery due to annealing. The presence of such a high density of the dislocations is probably due to the deviation from misorientation between grains. The dislocations may also be induced

254

by the strain between the grains. The dislocations can be part of the equilibrium boundary structure in order for the interface to retain the low energy configuration over most of its area.

Fig. 17. TEM image showing a twist boundary with c-axis as a common axis

Fig. 18. Dislocation network at twist boundary observed in Ag~i-2223 tape

High-angle boundaries are also found in the melt-textured BSCCO. High resolution TEM image shows that the boundaries are strongly faceted into a number of segments associated with steps of one (002) lattice spacing with boundary planes parallel to the basal plane[59]. The low interracial energy of the (001) twist boundaries is the result of the highly 2D structure of the BSCCO. 4.6 GRAIN AUGNMENT Owing to the large separation between the two Bi-O layers (0.3 nm), the Bibased materials exhibit a micaceous morphology and are easily cleaved. Mechanical deformation, in particular, rolling and pressing process takes the advantage of the plate-like morphology for grain alignment in the a-b plane direction. The degree of alignment has been considered as a major contributing factor for the high J= in the Ag-sheathed Bi-based tapes. It has been realised that repetitive deformation and annealing is necessary to achieve clean grain boundaries and high degree of grain alignment, and hence high J=. Yamada et al [46], Osamura et al [56] and Dou et al [70] have correlated the Jc with the number of intermediate pressings. The degree of texturing is dependent on the oxide core thickness. It was found that degree of texturing decreased from the Ag/superconductor interface towards the centre of the oxide layer [71]. Thus, the high Jc values observed in the

255

thinner tapes are attributed to the improvement of grain alignment. Wilhelm et al studied the J= of samples with different core thickness [44]. They found that the J= dropped an order of magnitude when the tape core thickness increases from 20 pm to 70 pm. The J= dependence on the thickness may be attributable to the degree of texturing or/and the effect of self-generated fields. The relative importance of the two factors depends on processing procedure used. Recently, we have shown that the degree of texturing does not change within the entire oxide core thickness for the tapes prepared using a controlled melting process[70]. Thus, in this case, the self-generated fields are mainly responsible for thickness dependence. The degree of grain alignment of the tapes is usually demonstrated by X-ray diffraction patterns. However, XRD patterns only give the degree of alignment in a very thin layer on the surface and do not reflect the alignment situation in the entire thickness of sample since the texturing degree may vary from surface to the centre of the tapes. By rotating a tape sample in magnetic field a peak was observed in the plot of J= versus 8 which is the angle between the rotating tape and the applied field. The Jc shows a maximum when the applied field is parallel to the tape surface and a minimum when the H is perpendicular to the tape surface. The ratio of a = Jc(H.Lc)/Jc(HIc) gives a good indication of the degree of grain alignment in the entire thickness of the tape since the magnetic field (for example, 0.08 T) penetrates the tape. Thus, this can be used to compare the degree of texturing in various samples. Figure 11 shows the normalised J= versus 8 for the tapes treated at 832~ for different periods of time. Figure 20 shows Jc(H.Lc)/J=(HIc) measured at

Fig. 19. Bright-field TEM image showing a high density of dislocations in a Ag--clad 2223 tape

Fig. 20. The J, and degree of texture, ct, versus annealing time for the tapes presented in Fig. 111701

256

77 K and a field of 0.08 T versus annealing time for the samples given in Figure 3. The J= values at 77 K and zero field are also shown in the figure. It is seen that both the J= and the degree of texture (a) increase with the annealing time up to 230 h, and then start to drop. Possible mechanism for the texture formation in the silver-sheathed Bi-Pb-Sr-Ca-Cu-O wires may be the combined effect of mechanical deformation and the liquid phase sintering. Jin et al [72] proposed the following mechanisms for the formation of the texture in the silver-sheathed Bi-PbSr-Ca-Cu-O wires: (i) deformation induced alignment of fractured grains, (ii) annealing texture, and (iii) Ag/oxide interface induced texture. In addition, the liquid phase formed during prolonged partial melt sintering probably plays an important role for the grain alignment. Recently, Hu et al. found that Jc was not significantly affected by a small angle misorientation of the grains(<10 ~ in the Ag-clad BPSCCO tapes[73]. They found that the critical current anisotropy I=(B,e) where e is the angle between field and a-b plane is influenced only by the c-component of the magnetic field and can thus be expressed as Ic(B-Sine). This is considered as evidence for two dimensional behaviour of the Bi-based HTSC. The data do not follow this relationship, if the angle e is smaller than about 10~ Comparison of different tapes reveal that within misorientation angle 10~ the degree of texture is not a controlling factor for the Jc other factors such as pinning ability may play important role. This is consistent with the results on YBCO thin film bicrystals reported by Dimos et al.[9]. Melting process has been successfully used for processing YBCO and Bi2212 in order to achieve a good grain connection and alignment. However, the melting process can not be simply used for processing Bi-2223 since it will cause 2223 phase degradation and segregation[29]. Two processes have been developed to achieve the liquid phase formation during the heat treatment of Ag-clad Bibased 2223 tapes. One of the procedures is to anneal the tape at a temperature as close the melt temperature as possible in order to achieve a partial melting without decomposition of the 2223 phase. The tapes produced by this procedure show a good grain connectivity, crystalline alignment and high physical densities with an average 2212 fraction around 6% to 10%. A Jc of 32,000 A/cm 2 at 77 K and zero field has been achieved [70]. Another procedure is the use of a short period of melt at higher temperatures as discussed in previous sections[48-51]. 5. FLUX PINNING MECHANISM IN BSCCO

It is known that grain boundaries act as effective pinning centres in the conventional alloy superconductors. However, the grain boundaries in the HTSC are normally the cause for weak links rather than flux pinning sites. Some high angle boundaries have been found and identified to have flux pinning effect in YBCO as mentioned in previous section. The effect of grain boundaries, interfaces, and defects on flux pinning in BSCCO is unclear. 5.1 INTRINSIC PINNING Tachiki and Takahashi proposed an intrinsic pinning mechanism to interpret

257

the high J= in the oxide superconductors[62]. In the high-Tr oxide superconductors, the CuO2 layers are strongly superconductive, while the other layers are weakly superconductive or non-superconductive. Since the coherence length along c-axis is short, the radius of the vortex core along c-axis is small. The layer structure itself works as strong pinning centres of the vortices. This intrinsic pinning is responsible for the superior J= - B behaviour of the Ag-clad Bi-based superconductor tapes at low temperatures. For effective pinning, the grain alignment is essential for the oxide superconductors. 5.2 DEFECT PINNING Some pinning mechanisms such as Y2BaCuO4 precipitate pinning[74], and twin plane pinning[75] have been demonstrated in YBa2Cu3Oz_x(Y-123). The pinning mechanism in Ag-clad Bi-based superconducting wires, however, has not been well studied. It is not clear that the defects such as dislocations, stacking faults and interfaces can act as pinning centre. In fact, it is even not certain whether the defects such as dislocations are present in the samples after prolonged annealing. The flux pinning properties have been widely investigated through the measurements of irreversibility line (IL). The IL, defined as the disappearance of pinning at a certain temperature and certain field, can be determined from the DC magnetisation measurements under field cooling and zero field cooling, and the AC susceptibility measurements under varying DC fields. The irreversibility lines (IL) for melt-processed and normal processed samples were determined from magnetisation measurements as shown in Figure 21. For comparison, the IL for caxis aligned powder samples YBCO, 2223[76] and single crystalline TIBa2Ca~Cu30~('i'I-1223) [77] are also included. All the measurements were made with the field parallel to c-axis. It is generally believed that the IL is a function of the coupling strength between the superconducting CuO2 planes. The distance between the inter-CuO2 planes is taken as a measure of the coupling strength. The position of the IL for Y123(0.84 nm), T1-1223(0.88 nm) and Bi-2223(1.23 nm) appears to obey well this rule. However, the significant shift of the IL for the same Bi-based materials observed in Figure 21 can not be explained by this rule. It is noted that the IL for the normal tape(J= - 9900 A/cm 2) is slightly shifted to higher temperature compared to the IL for the powder Bi-2223 sample, while for the melt-processed tape (Jc - 19,000 Ncm2), the IL is positioned at remarkably higher temperature than that for the normal processed tape and the powder Bi2223. Furthermore, the IL for the melt-processed tape does not show an abrupt change which is commonly observed in the Bi-based powders, single crystals and thin films. This has been considered as a result of the transition of the three-dimensional vortex glass state to 2D pancake-like vortex lines[78]. The major difference between the c-axis aligned powders and c-axis aligned Ag/Bi-2223 tapes is whether they have interface. In addition, the grains in the tape samples may contain more defects than those for the powders due to the thermomechanical process and melting in the case of tapes. Thus, the significant enhancement of the IL can be attributed to a strong interracial pinning and defect pinning.

258

To investigate the grains, we cut open a tape, split the silver sheet, scraped out the powder from the tape and measured the irreversibility line for the tape powder. Figure 22 compares the IL for the powder from the tape and normal powder. It is clear that the IL for the tape powder is still positioned at higher temperatures than the normal powder. The grain boundaries in the tape powder have been at least partially destroyed, but the significant shift of the IL for the tape powder remains nearly the same as that for the tape. This suggests that the pinning within the grains themselves is improved by the harsh thermomechanical process. 8.00

8.00

Calcined powder 2223: rm powder from 2223 tape: 6.00

6.00

\ ~

p. v

t: 4.00

=

2.00

0.00 1 7 i l l 0.00

4.00

2.00

w i i,

ii

iw

i'!'1

0.20

iI

ii

ill

0.40

w IIll

l l l l

0.80

T,,,/%

l|

| 1 1 I I 1 , 1 1 !

0.80

, ~ '

1.00

Fig. 21. Irrevers~ility fines for the meltand normal tape in comparison with c-axis aligned powder Y123 and Bi-2223176] and single crystal T11223177]

processed tape,

o.o0 .... ; 0.00

......

. . . . . . . . . . . . . . . . .

0.20

0.40

o o ~ i ....

0.60

,, , r,

, ,~,r

0.80

,

1.00

Fig. 22. Irrevers~ility lines for the powder scraped from a Ag--clad Bi-2223 tape and the Bi-2223 normal powder

Figure 23 shows the lattice fringe of the melt-processed tape with electron beam perpendicular to c-axis. The tape has high concentration of stacking faults along c-axis. Figure 19 is a TEM image showing a high density of dislocation lines (10~~ 2) in a melt-processed Ag-clad tape. We have also shown a high density of interfacial dislocations in this tape in our previous paper[52]. These planar and linear defects, which have a comparable line spacing to the flux lines in a 2 T field, may act as desirable pinning sites, responsible for the enhancement of the IL of the Ag-sheathed Bi-based wires. If the defects act as pinning centres it would be expected that an increase of the density of the defects will result in an increase in the Jc. To increase the density of the defects an additional uniaxial pressing was applied to the annealed tape samples without further heat treatment and the effect of uniaxial compressive strain on the flux pinning was investigated. The present approach allows to isolate and

259

identify the sole effect of the room temperature uniaxial pressing on the flux pinning in the Ag-clad Bi-based superconducting tapes. The results of magnetisation measurements clearly demonstrate that the norrnalised Jc for the uniaxial pressed tape has an improved performance in magnetic fields compared with that without the final pressing as shown in Figure 24. This enhanced flux pinning effect becomes more pronounced at high temperatures. Since the major effect of the uniaxial pressing is the creation of the defects, the present results indicate that the defects such as dislocations act as pinning centres in the tape.

Fig. 23. TEM image showing heavy stacking faults in the melt-processed tape

Fig. 24. NormaJised magnetisation J , : - H curves at 50 K for the tape samples AP(after pressing) and BPCoefore pressing) with H normal to the tape surface

6. SPECIAL ROLE OF AG/BSCCO INTERFACE ON MECHANICAL PROPERTIES

High-T 9materials, as with all ceramics, are very brittle and difficult to shape and handle. For large-scale application, superconducting wires have to be reasonably flexible so that they can be wound into coils for handling, storage, and installation. Thus, it is important to overcome the brittleness problem of high-T c materials before practical applications can be realised. Silver has been successfully employed to fabricate metal/superconductor composites, such as metal-clad wires, tapes, and multifilaments, to compensate for the ceramic's brittleness. Furthermore, the metal provides a means of thermal dissipation, thus stabilising the superconductor environment. The stabilising effect is of fundamental importance for type II superconductors, in which undesirably large local rises in temperature can develop through flux jumping in the mixed state. Figure 25 shows the normalised Jc versus bend strains for single tapes[79,80]. It can be seen that the tape with a superconductor thickness of 25/Jm

260

and a total thickness of 90 pm could be bent to a strain as large as 0.98%(a radius of 4 mm) without damage to the grain linkages, as indicated by the unchanged critical current. It is of interest that the same tape could be flexed at progressively smaller radii of curvature prior to reaching the critical bend strain. With increasing thickness of the superconducting tape, the critical bend strain increased rapidly. The bending ability became poor when the thickness of the superconductor was greater than 120 pm. Since the bending property of the silver is far better than the ceramic oxide the bend strain for the silver/oxide composites will strongly depend on the ratio of area of the silver/oxide interface layer to the volume of the composite. Figure 26 shows the normalised J= versus surface bend strains for both the single tape and multifilamentary tape with the same thickness of 0.2 mm. The critical strain for the single tape is 0.2%, while the critical bend strain for the 19 multifilamentary tape is 0.8%. The critical strain for the multifilamentary tape is slightly higher than those reported previously(0.7%)[81]. This may be attributable to the higher ratio of Ag to the oxides in the former (8 vol% oxide) than in the later(9 to 13% oxides). Sato et a1182] reported that the critical tensile strain increased from 0.1% to 0.7% when the numbers of filaments in the tape increased from 1 to 1296.

1.20 1.20

AE-CLAD

Bi(2223)

SINGLE TAPE

~

1.00

1.00

~SINGLE TAPE (SB) c-'c,.,~MULTIFILAMENT(SB) ,:,,:,,:,,:~ MULTIFILAMENT(DB)

0.80

0.80 *o u

0.60

0.40

0.40 "

o.o 0.00

0.60

~ ....

o.oo

, ....

o.so

0.20

0.00

,.~o' 'i.~o' ":/.;d ' ~o' "36o

SURFACE

BEND

STRAIN

(~),E.

Fig. 25. Dependence of the critical surface bend strains on the thickness of the oxide core of Ag-clad Bi-2223 single tapes

......... , ......... , ......... ~. . . . . . . . . , ....... 0.00 l.O0 2.00 3.00 4.00 E, BEND STRAIN (~)

~', 5.00

Fig. 26. Comparison of the bend strains for single tape and multifilamentary tape

For the tensile tests, the normalised Jc versus tensile strain for the single and multifilamentary tapes, both having the same thickness of 0.2 mm is shown in Figure 27. It is seen that the critical current density remained unchanged up to a strain of 0.65% for the 19 filamentary tape while, the Jc started to drop at a strain rate of 0.15% for the single tape. The Jr of both single and multifilamentary tapes

261

declined rapidly to about 50% of its zero strain value and then showed a gradual decrease. It is argued that the difference in critical strain reported by various group may be attributable to the different volume ratio of the superconducting oxide to the silver rather than the different type of the precursors. This suggests that the interracial layer has better tensile strength than the bulk oxides since the higher volume ratio of silver versus oxide means higher volume ratio of the interfacial layer to the bulk oxide, in regards of the oxide-filament configuration in the metal matrix studied here. It is proposed that an interfacial composite layer was formed through a strong bonding between the oxide and Ag due to the eutectic reaction during the heat treatment[83,84]. Figure 28(a) is the SEM image showing the cracks on the top surface of the oxide core after splitting the Ag sheath from the tape of 0.16 mm thick in bend test at a bend strain of 1%. At this stage, the normalised Jc is about 10% of the zero strain Jc. Figure 28(b) shows the fracture surface of the Ag sheath that was split

1.20

-

oc.:~c.~MULTIFILAMENT WITH STRESS : - - ' - : - : . - ~ MULTIFILAMENT WITHOUT STRESS SINGLE TAPE WITH STRESS

1.00

l

0.80

o "~o 0.80

0.40

0.20

0.00

{

.........

0.00

, .........

!.00

, .........

2.00

TENSILE

STRAIN

~ .........

3.00

(2)

,,

4.00

Fig. 27. Normalised Jc versus tensile strain for single and mutifilamentary tapes with total thickness 0.2 mm from the bend test tape. No crack was observed on the surface. This Ag sheath with a thin layer of oxides less than 5 pm on the Ag surface has experienced much more severe bend stress with a strain greater than 1.5%. The four probe technique performed on the single side Ag sheath shows a finite amount of critical current under 1 pV/cm criterion. This demonstrates that the interfacial layer between the oxide and Ag has higher flexibility and higher strength than the bulk oxide layer. The Ag sheath not only assists in compensating for the brittleness of the

262

ceramic superconductor, it also provides protection for the material from atmospheric degradation and corrosion. The Ag-clad wires maintained nearly constant critical current densities for 2 years, up to the preparation of this paper. Most of the wires showed no change after cycling more than ten times between liquid nitrogen temperature and room temperature.

Fig. 28. SEM image of the top fracture surface of the oxide core for a tape of 0.2 mm thick after bending to a strain of 1.0%(a) and SEM image of the fracture surface on the Ag sheath with a bend strain greater than 1.5%(b) 7. SUMMARY

The composition, structure, orientation and properties of the interfaces play a crucial role in determining the properties of the HTSC materials. The interfaces act as Josephson weak links, and in effect control the Jc in random oriented polycrystalline materials. Chemical examination on the grain boundaries of the HTSC shows that the composition even in the clean boundaries deviates from the interior of the grains. Some high angle boundaries show a flux pinning characteristic. Studies on the grain boundary structures revea that the lattice matches can be described by a coincident site lattice (CSL) model or a constrained CSL model for YBazC3OT_x(YBCO). The impressive high critical current densities Jc obtained for Ag-clad Bibased superconducting tapes is of great interest for high current and high field applications, and hence has attracted more attention to the grain boundary studies in these materials. A strong correlation between the high Jc and textured platelike grain morphology has suggested the "brickwall' model. Most of the grain

263

boundaries in the textured BSCCO are twist boundaries in (001) with c-axis as rotation axis. Grain boundary dislocations are commonly observed in these twist boundaries with high density. The presence of such a high density of the dislocations is probably due to the deviation from misorientation between grains in order for the interface to retain the low energy configuration over most of its area. Recently, studies on the transport and magnetisation Jc, and the irreversibility behaviour of the textured Ag-clad Bi-based tapes show no evidence of grain boundaries for being the weak links and the overall Jc is controlled by the intragrain pinning at high temperatures but weak link behaviour becomes evident at low temperatures. The interfaces between the HTSC oxides and the silver sheath exhibit unusual mechanical properties that are responsible for the improved strength and flexibility of the Ag/HTSC composites. ACKNOWLEDGMENTS

The authors gratefully acknowledge Y.C. Guo, J. Yau, J.X. Jin, Q.Y. Hu Associate Professors Y. Zhao, C.C. Sorrell, and all members of Superconductivity Group for their dedication and contribution to this project. Sincere thanks are given to our collaborators: Drs. A.D. Caplin, H.W. Weber, E.D. Collings, N. Sawides, P. Munroe, D.L. Shi, E.R. Vance and Mr. W.M. Bian for their help and support and to Metal Manufactures Ltd. and commonwealth Department of Industry, Technology and Commerce for the financial support. REFERENCES

[1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11]

Goodenough, J.B. and Manthiram, A.: J. Solid State Chem., 1990, 88, 115 Raveau, B., Michel, C., and Hervrieu, M.: J. Solid State Chem., 1990, 88, 140 Kitazawa, K.: Amer. Ceram. Soc. Bull., 1989, 68, 880 Campbell, A.M.: Cryogenics (UK)., 1990, 30 Suppl., 809 Kes, P.H.: Physica C., 1988, 153-155, 1121 Yeshurum, Y., Malozemoff, A.P., Worthington, T.K., Yandroski, R.M., Krysinelbaum, L., Holtzberg, F.H., Dinger, T.R. and Chandrashekhar, G.V.: Cryogenics, 1989, 29, 258 Gammel, P.L., Schneemeyer, L.F., Waszczak, J.V. and Bishop, D.J.: Phys. Rev. Lett., 1988, 61, 1666 Dimos, D., Chaudhari, P., Mannhart, P. and LeGoues, F.K., Phys. Rev. Lett., 1988 61,219 Dimos, D., Chaudhari, P., Mannhart, P. and LeGoues, F.K., J. Phys. Rev., 1990, B41, 4038 Larbelestier, D.C., Babcock, S.E., Cai, X.Y., Field, M.B., Gao, Y., Heing, N.F., Kaiser, D.L., Merkle, K.L., Williams, L.K., and Zhang, N., Physica C 1991, 185-189, 315 Enomoto, Y., Suzuki, M., Murakami, T., Inukai, T. and Inamura, I., Jpn. J.

264

[12] [13]

[14] [15] [16] [17] [18] [19] [2Ol

[21] [22]

[23] [24]

[25] [26] [271

[28] [29] [30] [31] [32] [33] [34]

Appl. Phys. 1981, 20, L661 Koch, R.H., Umbach, C.P., Clark, G.L., Chaudhari, P., and Laibowitz, R.B., Appl. Phys. Lett., 1987, 51,200 Dou, S.X., Ringer, S., Bourdillon, A.B., Sorrell, C.C. and Easterling, K.E., The Australian Phys. 1987, 24, 165 Gao, Y., Merkle, K.L., Bai, G., Chang, H.L.M., and Lam, D.J., Physica C, 1991, 174, 1 Babcock, S.E., Zhang, N., Gao, Y., Cai, X.Y., Kaiser, D.L., Larbalestier, D.C., and Merkle, K.L., J. Adv. Sci. (Japan), 1992, 4, 119 Zhu, Y., Zhang, H., Wang, H., and Suenaga, M., J.Mater. Res., 1991, 6,2507 H. Maeda, Y. Tanaka, M. Fukutomi, and T. Asano: Jpn.J. Appl. Phys. Lett., 1988, 27, L209 R.M. Hazen, C.T. Prewitt, R.J. Angel, N.L. Ross, L.W. Finger, C.G. Hadidiacos, D.R. Vablen, P.J. Heaney, P.H. Hor, R.L. Meng, Y.Y. Sun, Y.Q. Wang,Y.Y. Xue, Z.J. Huang, L. Gao, J. Bechtold, and C.W. Chu: Phys. Rev. Lett., 1988, 60, 1174 J.M. Tarascon, Y. Le Page, and P. Barboux, B.G. Bagley, L.H. Greene, W.R. Mckinnon, G.W. Hull, M. Giroud, and D.M. Hwang: Phys. Rev. B., 1988, 37, 9382 S.A. Sunshine, T. Siegrist, L.F. Schneemeyer, D.W. Murphy, R.J. Cava, B. Batlogg, R.B. van Dover, R.M. Fleming, S.H. Glarum, S. Nakahara, R. Farrow, J.J. Krajewski, S.M. Zahurak, J.V. Waszczak, J.H. Marshall, L.W. Rupp Jr, and W.F. Peck: Phys. Rev. B., 1988, 38, 893 H.K. Uu, S.X. Dou, N. Sawides, J.P. Zhou, N.X. Tan, A.J. Bourdillon, and C.C. Sorrell: Physica C., 1989, 157, 93 S.X. Dou, H.K. Uu, A.J. Bourdillon, M. Kviz, N.X. Tan, and C.C. Sorrell: Phys. Rev. B., 1989, 40, 5266 R. Ramesh, S.M. Green, and G. Thomas: in Studies of High Temperature Superconductors ed by A. Narlikar, p361. Nova Science Publishers, New York, 1990 D.J. Werder, C.H. Chen, M. Gurvitch, B. Miller, L.F. Scheemeyer, and J.V. Waszczak: Physica C, 1989, 160, 441 K. Sato, T. Hikata, and Y. Iwasa: Appl. Phys. Lett., 1990, 57, 1928 Y. Yamada, B. Oberst, and R. FI0kiger: Supercond. Sci. Technol., 1991, 4, 165 J. Tenbrink, M. Wilhelm, K. Heine, and H. Krauth: presented at the Applied Superconductivity Conf., Snowmass Village, CO (September 24-28, 1990). S.X. Dou, H.K. Liu, and Y.C. Guo: Physica C., 1992, 194, 343--350 Y.C. Guo, H.K. Liu, and S.X. Dou: Physica C, 1992, 200, 147 S.E. Babcock, X.Y. Cai, D.C. Larbalestier, and D.L. Kaiser: Nature 1990, 347, 167 C.B. Eom, A.F. Marshall, Y. Suzuki, B. Boyer, R.F. Pease, and T.H. Geballe, Nature 353, 544 (1991) A.Singh, Chandrasekhar, and A.H. King, Acta Cryst. B45 (1990) 117 Y. Zhu, H. Zhang, H. Wang, and M. Suenaga, J. Mater. Res., 6 (1990) 2507 S.E. Babcock and D.C. Larbalestier, J. Mater. Res. 5 (1990) 919

265

[35]

[36] [37]

[38] [39] [40l [41] [42]

[43] [441 [45] [461 [47]

[481 [49] [501 [51] [52] [53] [54]

[55] [56] [57] [58] [59] [6Ol [611 [62]

D.A. Smith, M.F. Chisholm and J. Clabes, Appl. Phys. Lett., 53 (1988) 2344 Y. Zhu, Philos. Mag. submitted J.J. Un, E.L. Benitez, S.J. Poon, M.A. Subramanian, J.Gopalakrishnan, and A.W. Sleight: Phys. Rev. B., 1988, 38, 5095 D.E. Farrell, S. Bonham, J. Foster, Y.C. Chang, P.Z. Jiang, K.G. Vanvervoot, and D.J. tam: Phys. Rev. Lett., 1989, 63, 782 Y. lye, T. Tamegai, T. Sakakibara, T. Goto, N. Miura, H. Takeya, and H. Takei: Physica C., 1988, 153-155, 26 K. Heine, J. Tenbrink, and M. Thoner: Appl. Phys. Lett., 1989, 55, 2441 N. Enomoto, H. Kikuchi, N. Uno, H. Kunakura, K. Togano, and K. Watanabe: Jpn.J. Appl. Phys. 1990, 29, L447 S.X. Dou, H.K. Uu, J. Wang, M.H. Apperley, C.C. Sorrell, S.J. Guo, B. Loberg, and K.E. Easterling: Physica C., 1990, 172, 63 S.X. Dou, H.K. Uu, M.H. Apperley, K.H. Song, and C.C. Sorrell: Supercond. Sci. Technol., 1990, 3, 138 M. Wilhelm, H.W. Heuw011er, and G. Ries: Physica C, 1991, 185-189. T. Hikata, M. Ueyama, H. Mukai, and K. Sato: Cryogenics., 1990, 30, 924 Y. Yamada, K. Jikihara, T. Hasebe, T. Yanagiya, S. Yasuhara, M. Ishihara, T. Asano, and Y. Tanaka: Jpn. J. Appl. Phys., 1990, 29, L456 M. Ueyama, T. Hikata, T. Kato, and K. Sato: Jpn. J. Appl. Phys. 30 (1991) L1384 S.X. Dou, H.K. Uu, and Y. C. Guo: Appl. Phys. Lett., 1992, 60, 2929 Y.C. Guo, H.K. Liu and S.X. Dou, Appl. Supercond. 1, 25 (1993) S.X. Dou, Y.C. Guo, H.K. Uu, and D.L. Shi, Cryogenics, (in press) S.X. Dou, H.K. Uu, Y.C. Guo, D.L. Shi, M.D. Sumption, and E.W. Collings, Proceedings of ISS'92, 16-19 Nov. 1992, Kobe, Japan H.K. Uu, Y.C. Guo, and S.X. Dou, Supercond. Sci. Technol., 1992, 5, 59198 K. Sato, T. Hikata, H. Mukai, M. Ueyama, N. Shibuta, T. Kato, T. Masuda, M. Nagata, K. Iwata, and T. Mitsui: Transactions on Magnetics., 1991, 27, 1231 J.W. Ekin, T.M. Larson, A.M. Hermann, Z.Z. Sheng, K. Togano, and H. Kumukura: Physica C, 1989, 160, 489 K. Sato, T. Hikata, M. Ueyama, H. Mukai, N. Shibuta, T. Kato, and T. Masuda: Cryogenics., 1991, 31,687 K. Osamura, S.S. Oh, and S. Ochiai: Supercond. Sci. Technol., 1990, 3, 143 T.T.M. Palstra, B. Batlogg, L.F. Schnneemeyer, R.B. van Dover, and J.V. Waszczak: Phys. Rev. B., 1988, 38, 5102 S. Jin, T.H. Teifel, R.C. Sherwood, R.B. van Dover, M.E. Davis, G.W. Kammlott, and R.A. Fastnacht: Phys. Rev. B., 1988, 37, 7850 Y. Gao, C.T. Wu, Y. Shi, K.L. Merkle and K.C. Goretta, Appl. Supercond. 1 (1993) 131 L.N. Bulaevskii, J.R. Clem, L.I. Glazman, and A.P. Malozemoff, Phys. Rev. B45 (1992) 2545 M.P. Maley, P.J. Kung, J.Y. Coulter, W.L. Carter, G.N. Riley, and M.E. McHenry, Phys. Rev. B 45, 7566(1992) M. Tachiki and S. Takahashi, Solid State Comm. 70, 291(1989)

266

[631 [641 [651 [66] [67]

[68] [69] [70] [71] [72] [73] [74] [75] [76] [77] [78] [79]

[80] [81]

[82] [83] [84]

M.A. Angadi, et al, Physica C, 1991, 177, 479 S.M. Cassidy, L.F. Cohen, M.N. Cuthbert, J.R. Laverty, G.K. Perkins, A.D. Caplin, S.X. Dou, Y.C. Guo, H.K. Uu, F. Uu, and E.L. Wolf: J. Alloys and Compounds(in press) T.T.M. Palstra, B. Batlogg, R.B. van Dover, L.F. Schneemeyer, and J.V. Waszczak: Phys. Rev. B., 1990, 41, 6621 D.P. Hampshire, S.S. Oh, K. Osamura, and D.C. Larbalestier: Supercond. Sci. Technol., 1990, 3, 560 R.N. Kleiman, P.L. Gammel, L.F. Scheemayer, J.V. Waszczak, and D.J. Bishop: Phys. Rev. Lett., 1989, 62, 2331 S.X. Dou, H.K. Liu, J. Wang, and W.M. Bian: Supercond. Sci. Technol. 1991, 4, 21 S.X. Dou, H.K. Uu, J. Wang, K.H. Song, and G.J. Bowden: Physica C, 1990, 171,293 S.X. Dou, H.K. Uu, and Y.C. Guo: Proceedings of BHTSC-92, 25 - 29 May, 1992, Bejing, PRC Y. Yamada, B. Oberst, and R. Flufiger: Supercond. Sci. Technol. 4 (1991) 165 S. Jin, and J. Graebner: Materials Sci. and Eng., 1991, B7, 243 Q.Y. Hu, H.W. Weber, S.X. Dou, H.K. Liu, and H.W. Neumuller, J. Alloys and Compounds(in press) M. Murakami, M. Morita, K. Doi, and M. Miyamoto: Jpn. Appl. Phys. 1989, 28, 1189 L.J. Swartzendruber, D.L. Kaiser, F.W. Gayle, L.H. Bennett, and A. Roytburd: Appl. Phys. Lett., 58, 1566(1991) M. Suenaga, D.O. Welch and R. Budhani, Supercond. Sci. Technol 5, $1 (1992) T. Nabatame, Y. Saito, T. Doi, K. Aihara, T. Kamo, and S.Matsuda, S., Physica C 1991, 190, 114 P.H. Kes, Supercond. Sci. Technol. 5, $41 (1992) S.X. Dou, H.K. Liu, Y.C. Guo, and C.C. Sorrell: Physica C, 185-189, 2494, 1991 S.X. Dou, Y.C. Guo, J. Yau, H.K. Uu, and S.X. Dou: Appl. Supercond. (in press) A. Otto, L.J. Masur, J. Gannon, E. Podtburg, D. Daly, G. J. Yurek and A.P.Malozemoff, Proceedings of Applied Superconductivity Conference, Aug. 25-28, 1992, Chicago, IL, USA K. Sato, T. Hikata, H. Mukai, M. Ueyama, N. Shibuta, T. Sato, T. Masuda, M. Nagata, K. Iwata and T. Mitsui, IEEE Transactions on Magnetics, Vol. 27, No. 1231(1991) S.X. Dou, H.K. Liu, K.H. Song, M.H. Apperley, C.C. Sorrell, A.J. Gouch, and N. Sawides: Physica C., 160, 533 S.X. Dou, K.H. Song, H.K. Uu, C.C. Sorrell, M.H. Apperley, and N. Sawides: Appl. Phys. Lett., 1990, 56, 493

Science of Ceramic Interfaces II J. Nowotny (Editor) 9 1994 Elsevier Science B.V. All rights reserved.

267

THE ROLE OF INTERFACES NUCLEAR TECHNOLOGY

IN

MICHIO YAMAWAKI

Nuclear Engineering Research Laboratory, Faculty of Engineering, University of Tokyo, Hongo, Bunkyo-ku, Tokyo 113, Japan

Abstract For fission reactor fuels, U O ,

or (U,

Pu)O,,

grain boundaries play

important roles. During sintering, motion of grain boundaries causes grain

growth and microstru~tural development. Grain boundaries are a

path for the

fission gas release. As methods to enhance grain growth, oxidative and doping sinterings are useful. Increase in grain size is benefical for depressing

fission gas release, while adverse from the view point of the pellet cladding interaction due to increment in the creep strength. As for fusion reactor

plasma-facing materials, surface impurities play an important role to control

the transport of hydrogen isostope particles in the material. A new concept of

superpermeability may supply an effective method to separate hydrogen isostopes

from the helium exhaust. In the case of fusion reactor breeder materials, surface and grain boundaries play vital roles in effective recovering of

tritium from the ceramic breeder materials.

1. Introduction Interfaces of materials play important roles in various aspects of the

nuclear technology. In this review, recent topics are focused on regarding the areas of the fission reactor oxide fuels, fusion reactor plasma-facing

materials and fusion reactor breeding materials. Interfaces exert beneficial as well as adverse effects in various ways so that it is vitally important to understand and control their roles in the nuclear technology.

268 2 . Fission reactor fuels Microstructures such as grain boundaries and pore distribution are recognized

to have a strong influence on the property of U 0 2

or ( U ,

Pu)

02,

typical

fission reactor fuels. The property of grain boundaries is an important factor f

or the formation of such microstructures during sintering processes.

During sintering, after bonding is established between initial powder

particles, motion of grain boundaries takes place to cause grain growth and

microstructural development with the interface energy operating as the driving force.

Grain boundaries of U 0 2

pellets are considered to act as an important path

for the fission gas release ( F G R ) .

boundaries of U 0 2 1 1 ]

A typical model of such a role of grain

is shown in Figure

1. The gas atoms, being in dynamic

solution, diffuse to grain boundaries. The gas atoms reaching the grain

boundaries form gas bubbles on grain faces as well as at grain edges. The

development of grain edge porosity eventually forms an interlinked network of

tunnels and provides thereby an easy path for gas atoms produced deep in the body of the fuel.

9 -..

single

0

9

.

.

"

"

:-'.~

9

atoms

.

". 0

9

"

"

".

"

'

....

.

".o.

.

"

"

f!

open pore

:

"

"

9 .---..~

9

.

.

.

.

....

- ~

. ' . . . ,

9 -

o

."

""-'.'" "

?

""

o 9.

.

".

"

.

.

.

o .

"

"

. 0

I "

channel

"

D

.

O"

.

.

. . . . . 9

. 9

.

.,..... ,

9 .

0 "

~

9

.

.

g r a i n edge

~

9

.0

9

.

. . . . .

0 0

-..

'

I

grain

oundary

bubbles

small g r a i n bubbles

Fig.1 Schematic of mechanism of f i s s i o n gas r e l e a s e [ l ] .

269 Centre temp : 1500 dog C Fue[ radius : .53 cm Groin radius : 2.5, S ,15 microns Surface letup : 500 dog ,~ Oiff C o e f f . A : 7E-06 cm'/s Oiff Coeff 0 = 69503 Ftux depr fact = 1 Boundary conc = 5E+15 otoms/cm2

The above model i m p l i e s t h a t i n c r e a s e of U O 2

grain size

w i l l i n c r e a s e the f i s s i o n retension,

gas

thereby decreasing

the f i s s i o n

gas r e l e a s e .

prediction

This

has been confirmed in Rating

the Over Ramp Test a t Halden as result

showing the e f f e c t

g r a i n s i z e of U O ~

374 W/cm

5O

of

on the

FGR

[2] is p r e s e n t e d in

Figure

2.

-~

---,--

4o

Q.

i n c r e a s e the

d

g r a i n growth. One is to s i n t e r

2.5 micron grcins

--o--

5 micron groins

--+--

15 micron groins

0 t_

There are mainly two methods to e f f e c t i v e l y

:

A typical

well as at S t u d s v i k .

3O

0

U O2 under o x i d a t i v e c o n d i t i o n s (oxidative sintering),

and the

o

o t h e r is to use a d d i t i v e s

CO

admixed to pure U O ~

C . _0

(doping

20

ol

sintering).

In the former

method, the k i n e t i c s densification lattice

diffusion

proportionally

coefficient

the r e l a t i v e

density

f o r the l a t t e r

increased 3 [5].

or

"~

Effect

of grain

30

40-

~0

size

predicted

by the

p r e s e n t model f o r c o n s t a n t power and temperature o ~ati tfuel centre t e m p e r a t u r e i pe 5oo*c on ) 2].

t h e r e b y enhancing t h e g r a i n [6,7].While, or S i O 2

can cause f o r m a t i o n o f i n t e r f a c e phase of low m e l t i n g p o i n t , in an e x t r a o r d i n a r y

g r a i n growth due to l i q i u d phase

98

96

r a t e at g r a i n b o u n d a r i e s ,

a d d i t i o n of A 1 2 0 3

20

100

i n c r e a s e s the d i f f u s i o n

resulting

io

As

doping s i n t e r i n g

method, a d d i t i o n of N b 2 0 5

growth r a t e

0

Burnup. biWd/kgU02 Fig.2

up to the v a l u e of 0.25

as shown in F i g u r e

TiO2

of

increases

to x 2 1 3 , 4 ] .

of s i n t e r e d p e l l e t with x

10

of

is governed by the

U which in U O 2 + ~ Actually,

~

u_

~

94

~

92

.,..,

q,)

"m

90

~

88

0.0

o:,

"

0:2

"

0'.3

"

o.,

Stoichiometric deviation x Fig. 3. Sintered relative density versus stoichiometric deviation, under CO:, at 1100*C (rq) and 1200~ ( , ) at the heating rate of 500 K h.-1 for 45 rain [ 5 ].

270

sintering. affect

Such a d d i t i v e s

1 x lO-S:

not only the g r a i n s i z e

(1400 *C,30MPa)_

but a l s o the pore d i s t r i b u t i o n , especially

inner-grain

pore

(,-.

\

distribution.

0

.~

On t h e o t h e r hand,

from the

0

-bJ v

view p o i n t of p e l l e t - c l a d d i n g interaction

(PC

in g r a i n s i z e

I),

increase

account of a r e s u l t i n g

The

c r e e p r a t e d e c r e a s e s with the

l x 10-"

g r a i n d i a m e t e r as shown in 4.

In t h i s

respect,

an

I

10

optimization

of g r a i n s i z e

distribution

can be r e a l i z e d

with a r e l a t i v e

Particle

ease by means

of t h e o x i d a t i v e

sintering.

Fig.4

coarse grains

structure'

1

appropriately oxygen p a r t i a l

!! ~

by

controlling

r~

~:,oo I I <~

pressure during

by t a k i n g

sintering

path of |

Figure

5 [8].

grain structure

the

~'~#?7 7 ~ @ ~ ~

I

o

I I

Icli~

i

p r o d u c t sweep-out in the of a c o a r s e g r a i n

structure.

'400

1

o

~ ~ ~

ll:

~

jo

/

Oo/ o

t

.;

. ....

"i,

....

l:

-:"~2 " "

p

".,

~ io 2i

o

o

, 7.'-\.~ . e

~

02o

of the

volume of the c o a r s e g r a i n s as typical

as

of a f i n e g r a i n s t r u c t u r e

as w e l l as r e d u c t i o n fission

6001 I

the

a d v a n t a g e s such as p l a s t i c i t y typical

B~176

Such a bimodal offers

Pellet

I

'-~.,

the c o u r s e of the s i n t e r i n g , actually,

9 /zm)

UO.

I I~ ~:'(Ti:-.-'zo',,~,

Such a

can be a t t a i n e d

("D5o

18oo7 - 7 :I'L:"" ":;

network of f i n e g r a i n s was structure

diameter

Creep Rate of S i n t e r e d

Sinterin9 Temperature

with

in a c o h e r e n t

proposed for U O = .

30

The

bimodal g r a i n s t r u c t u t e ' c a l l e d 'rocks-in-sand

~ 0

w

enhance-

ment of t h e c r e e p s t r e n g t h .

Figure

0

i.,

is a d v e r s e on

200 0

2.00 2.05 2.10 2.15 2.20 2..25 2.30 2.35 - O/U-Ratio

Fig. 5. Grain structures and method of control, shown by superimposing sintering paths on U/O phase diagram[ 8 ].

271

3. Fusion reactor plasma-facing materials Thermal or radiation-induced

6

segregation of non-metal

5

impurities or certain alloying

elements on a surface drastically changes the properties of the

"~

surface of metals or alloys. Figure

6

~

I

J

i

a

!_

4 3

shows an observed

correlation between the ion-induced

2

nickel membrane of the thickness

1

hydrogen permeation rate of a of 2 0 ~ m

''

D2*----.-Ni (20.urn)

and the surface

0

concentration of sulfur as

I0

"

?

,

,

20

I

30

,

,J

z,O

Cs / a t %

analyzed by Auger electron

spectroscopy [9,10]. This result clearly shows that the surface

barrier due to a sulfur monolayer

Fig. 6. Relationship between the deuterium permeation rate and surface concentration ofsulfur on nickel at 773K[9,10].

on the membrane i n l e t side gives

H2,He,CH 4

r i s e to an enhancement of hydrogen t r a n s p o r t through the back side surface. This p r i n c i p l e was u t i l i z e d by A. I. L i v s h i t s et a l .

[11-13] to

design a new pump to be used to s e p a r a t e D/T-mixture from helium at the exhaust of a fusion r e a t o r . Their model experiment was performed by u t i l i z i n g

a s e t - u p as shown

s c h e m a t i c a l l y in Figure 7 . high vacuum ( U H V )

vessel,

A ultrain the

form of 50cm-dia t u b e ( F i g u r e models a s t r e t c h of the (International

7),

I TER

Thermonuclear

Experimental Reactor) pumping d u c t . A cylidrical

membrane may be

mounted j u s t along the v e s s e l w a l l s , but a c t u a l l y a 10cm-dia, 20cm-long membrane of 1 0 0 ~ m utilized.

t h i c k Nb was

Four ribbons of N b ,

Ta

X

pump. ~"~~ms] I

pump Fig.7

UHV apparatus for model experiments.

(1) Nb ~n~mb~n~, (2)~tomiz~to~, (3) ~h~mocouple, (4) quartz shields (to lessen the recombination losses) [ 11 ].

272

and W

serve as atomizators

where thermalized hydrogen isotope molecules (D2,

DT

o n

~,I 5 . 0 o

and T2) will be atomized.

62.0

The effects of switching

on of the atomizator are shown in Figure 8.

O9 60

The

,,1.0

k_

inlet pressure comes down,

and the outlet pressure

Q_ .t=_

reach a steady-state level,

_~0.0 ~-

through the membrane.

o_

presssure does not rise back

6 =10

correspondingly rises to

01,

2

. . . . 4~ 6

Time,

Ator~'zator "Compression"

due to permeation of hydrogen

on

Remarkably, the inlet

after the initial drop,

1

hrs.

on

f

i

:

-' ',

which testifies of the back reemission flux being

_o10 -3

virtually zero. That. means

Time,

hrs.

that almost the whole flux

~10 o

the membrane. This is why

F i g . 8 E f f e c t of s w i t c h i n g on of the a t o m i z a t i o n ~n the hydrogen pressures in upstream and ownstream chamb rs. The compression o arrow indicates a moment when the downstream chamber valve (V at the plot) is partly closed[ 12].

absorbed penetrates through

this phenomenon was named

'superpermeation' by A.I.Livshits.

4.

-4,

.

0

2 4 6

Fusion r e a c t o r b l a n k e t b r e e d e r m a t e r i a l s E s t i m a t i o n of the p r e c i s e t r i t i u m

i n v e n t o r y in the b l a n k e t is e s s e n t i a l

the design of a f u s i o n r e a c t o r f u e l c y c l e . tritium

In s i n t e r e d

r e l e a s e c o n s i s t s of m i g r a t i o n s t e p s such as d i f f u s i o n

in g r a i n ,

d i f f u s i o n along i n t e r - c o n n e c t e d pores and r e a c t i o n on the s u r f a c e . unit processes, tritium

r e l e a s e e x p e r i m e n t s [14-19],

d e s o r p t i o n and exchange r e a c t i o n s in-situ

tritium

Among t h e s e

the mechanism of s u r f a c e r e a c t i o n s remains to be c l a r i f i e d .

T r i t i u m m i g r a t i o n phenomena in s o l i d b r e e d e r s have been s t u d i e s in-situ

for

solid breeding materials,

[20,21],

out-of-pile

and modeling s t u d i e s

r e l e a s e e x p e r i m e n t s named T T T E x

by combining

e x p e r i m e n t s such as [22-23].

The

performed by members

i n c l u d i n g the p r e s e n t author involve the use of the r e s e a r c h r e a c t o r 'YAYOI' at the U n i v e r s i t y of Tokyo, whose neutron f l u x at t h e core c e n t e r

is

273

8 x 1 011 n/cm 2 at its maximum power of 2kWt,. In T T T E x ,

tritium

diffusivities within grains and along interconnected pores were evaluated by

using L i s O ,

LiAl02

particles of different sizes and H e +

3%H2

sweep

gas [14,15]. The chemical form and the release kinetics were found to depend on the composition of the sweep gas[14,15].

This implies that surface

phenomena such as adsorption and exchange reactions affect the tritium release behavior.

Polycrystalline particles of L i , O

were used in the experiments. In in-situ

tritium release experiments, about 5 g of the sample was loaded in a type-316 stainless steel (S S) sample cell ( 1 0 m m

diameter) in a type of packed column,

and 316S S wool was used to fix the sample in the ceil.

In in-situ tritium release experiments using sweep gases of H ~ + H 2 0

+D,O,

or H e

it was found that tritium residence time in the sample decreased with

the increase of water vapor pressure. After correction of the diffusion-relevant residence time which was estimated by the experimental result based on runs

using sweep gas of H e + 3 % H 3 ,

surface-relevant tritium residence time was

found to be inversely proportional to the square root of the water vapor

pressure in the sweep gas.

For the sweep gas of H e + D 2 0 ,

surface residence

time was about 1.4 times larger than that of H e + H 2 0 ( F i g u r e

9),

implying

1.4 times smaller in the surface reaction rate. This isotope effect shows that

movement of H H(D)-O-] Li,O

or D

atoms is important in the activated complex [ - O - T -

to form water molecules.

and L i 2 A I O 2

The amount of water vapor adsorbed on

was reported to be corelated with the water vapor pressure

by about half power dependence [20]. This dependence shows that water vapor is

dissociatively adsorbed as surface hydroxyl group.

From the results of this

adsorption isotherm and the results of in-situ tritium release experiments including H 2 0 - D 2 0

surface-OT

isotope effect, it was concluded that recombination of

and - O H ( o r

-OD)

103

'

is the rate-determining step.

I

'

'

I

9

'

'

,

_

-0.5 lO s HeO

..... <:1

LizO 600 ~ IOII 0o

I0'

I0 z

03

PH,o , Po,o / P c F J g. 9

H20 and D20 vapor pressure dependencies of surface tritium residence time for Li20 by

hz-situ tritium release experiments [ 20 ] .

274

Tritium release

c u r v e s o b s e r v e d in t h e o u t - o f - p i l e

are shown in F i g u r e is p l o t t e d

was o b s e r v e d .

constant

(K)

curve d e v i a t e s

was c a l c u l a t e d

in t h e l a t t e r

rate

In the i n i t i a l

From the s l o p e of t h i s

from the i n i t i a l

Apparent r e a c t i o n

(Figure

straight

1 1).

105i

....

from the a d s o r p t i o n

I ....

site

500~

reaction

with a l a r g e r

'

'

1 1 as being straight

adsorption

| , ' , , I

'

line

mechanism such as

,

,

energy.

, , , , , !

,

from second slope ~ ' ~ from first slope ix 9 Sweep 15Po o 9Chorge 15Po

-,

the

some e l a p s e d t i m e .

from t h e i n i t i a l

'

reaction 1 O,

from t h e s l o p e of the l i n e

surface

,,-.,.

I ....

after

experiment amount

a linear

K is a l s o shown in F i g u r e

The d e v i a t i o n

was c o n s i d e r e d to be caused by a d i f f e r e n t release

portion,

line the o v e r a l l

line portion

This c a l c u l a t e d

l a b e l e d ' f r o m the second s l o p e ' . tritium

release tritium

As shown in F i g u r e

c o n s t a n t was a l s o c a l c u l a t e d

portion.

tritium

where the l o g a r i t h m of the r e s i d u a l

against the elapsed time.

relationship rate

1 0

,

,

1 , 1 ,

,

i

, , l , ,

r

o to

E

o

r

.o_

Li20

_!

,

,

O.

10 3

0

,

o

o

g:0.~ 500~ ,

, 1 1 1 | I

I

,

l

,

, , , l , l

PH,0 (Po)

I0

,

00

Fig.11 Water vapor pressure dependence of o v e r a l l r e a c t i o n rate constant ( K ) f o r t r i t i u m release by H e + H 2 0 purge gas. L i , - O , o u t - o f - p i l e experiments [ 2 0 ] . ,

,

,

I

5

,

,

,

,

I

,

I0 Time (min)

,

,

,

5 R e l e a s e r a t e was found not t o be a f f e c t e d by the w a t e r vapor p r e s s u r e

Fig. 10. Changeof the residual tritium amount with time when desorbed by He + H20 purge gas. Li20, out-of-pile experiments [ 20 ].

c h a r g i n g gas but by t h e vapor p r e s s s u r e t h e p u r g i n g gas as c l e a r l y Figure

vapor p r e s s u r e

1 1.

Using the r e a c t i o n

and - O H

Dependence of K on t h e water

is the r a t e - c o n t r o l l i n g

rate constant

surface-related

tritium

(K)

= I/K.

again t h a t

step.

o b t a i n e d from the o u t - o f - p i l e

r e s i d e n c e time ( I / S )

corresponding

that of the in-situ tritium release experiment can be calculated equation of I / S

in

shown in

in t h e p u r g i n g gas was about a h a l f power, s u p p o r t i n g

recombination of-OT experiment,

in the

to

from the

This comparison is shown in Figure 1 2 .

Surface-

related tritium residence time estimated from the i n i t i a l straight portion of

275

w

w

103 (J

w

s

w iwj

,

w

'

I

!

w

w w ~|

J

9

|

I

w

~ ,ri

0 from oul-of-pile in silu exneri experimenl (Isl slope) / " v ment - 9from oul-of-pile // o - experimenl (2nd s l o p e P ' - ~ o

Li20 ~

G.) r

,,...,.

ca,)

-

~" 102

I

""'""b',~

600

_

9

|

1

,

*

i

_

,1,1

"-._

9

1

1

,

J

!

| , , | 1

,

i

i

i

| 1 , ~

I00

I0 PH,O (Po)

Fig.12 Tritium surface residence time estimated from the out-of-pile experiment. Comparison with the in-situ experiment[2o]. the o u t - o f - p i l e

experiment was by about a f a c t o r of 5

observed in the i n - s i t u

tritium

release experiment.

s h o r t e r than t h a t

However, the r e s i d e n c e time

e s t i m a t e d from the second p o r t i o n of the o u t - o f p i l e r e l e a s e experiment was found to be comparable to t h a t of the i n - s i t u experiment t r i t i a t e d

heterogeneous surface. release rate,

In the o u t - o f - p i l e

A s m a l l e r a d s o r p t i o n energy c o r r e s p o n d s to a l a r g e r

while a l a r g e r energy to a s m a l l e r r e l e a s e r a t e ,

the l a t e r p o r t i o n of the r e s i d u a l t r i t i u m in-situ

experiment.

water vapor is c o n s i d e r e d to be adsorbed on the

tritium

change in F i g u r e

as observed in

1 0.

However, in the

r e l e a s e e x p e r i m e n t , t r i t i u m atom is produced in the g r a i n by

n u c l e a r r e a c t i o n and then d i f f u s e s to the s u r f a c e where i t r e a s o n a b l e to c o n s i d e r t h a t on the s u r f a c e ,

s i t e of a l a r g e r a d s o r p t i o n energy during the i n - s i t u Summarizing the above, s u r f a c e r e a c t i o n s

is trapped.

experiments.

is

experiment.

in the t r i t i u m

release processes

from s o l i d b r e e d i n g m a t e r i a l were i n v e s t i g a t e d by performing i n - s i t u of-pile

It

t r i t i u m tends to be adsorbed a t the

and o u t -

The s u r f a c e r e a c t i o n was found to be much a f f e c t e d by the

chemical composition of sweep gas. When water vapor was added to the sweep gas, tritium

r e l e a s e r e a c t i o n on the s u r f a c e was enhanced by a c o n s i d e r a b l y l a r g e

recombination r a t e between - O T was a l s o enhanced and H T

and - O H .

fraction

increased.

By H2 a d d i t i o n ,

the r e l e a s e

rate

276

References

1. W. Hering, J. Nucl. Mater., 114(1983)41. 2. I.D.Palmer, C.Vitanza and R.J.White, OECD Halden Reactor Project HWR-155(1986). 3. H.Assmann, W.Dorr and M.Peehs, J.Nucl.Mater.,

140(1986)1.

4. H.-J.Matzke, in:Plutonium 1975 and other Actinides, Eds.H.Blank and R.Linner (North-Holland,Amsterdam,1976)pp.801-803. 5. H.Cherrel, P.Dehaudt, B.Fancois and J.F.Banmard,J.Nucl.Mater.,

189(1992)175.

6. H.Assmann, W.Dorr, G.Grandle, G.Maier, M.Peehs, J . N u c l . M a t e r . ,

98(1981)216.

7. K.Une, J . N u c l . M a t e r . ,

158(1988)210.

8. H.Assmann, W.Dorr and M.Phees, J.Am.Cer. Soc., 67(1984)631. 9. M.Yamawaki, T.Tanabe, K.Yamaguchi and K.Kiyoshi, Nucl. I n s t . Meth., B23(1987) 498. lO.K.Yamaguchi and M.Yamawaki, ISIJ I n t e r n a t i o n a l , 29(1989)560. ll.A.Livshits,

M.E.Notkin, A.A.Samartsev, A.O.Busnyuk, and V . I . P i t u n i o v ,

17th Symposium on Fusion Technology, Rome, 14-18/9/92. 12.A.Livshits, M.E.Notkin, A.A.Samartsev and I . P . G r i g o r i a d i ,

J.Nucl.Mater.,

178

(1991)1. 13 .A . L i v s h i t s , M.E.Notkin, A.A.Samartsev , J . N u c l . M a t e r . , 14.T .Terai, et a l . , Fusion Eng. and Desi gn, 7(1989)345.

170(1990)79.

15.S .Tanaka et a l . , J.Nucl.Mater.,155-15 7(1988)533. 16.R .A.Verall et a l . , J . N u c l . M a t e r . , 155 -157(1988)533. 17.M .Briec, et a l . , J.Nucl.Mater., 155-1 57(1989)549. 18.H .Kwast, et a l . , Fusion Eng. and Design, 8(1989)3459. 19.T 9Kurasawa, et a l . , J . N u c l . M a t e r . , 155-157(1988)544. 20.S 9Tanaka et a l . , Fusion Technol,, 14(1988)816. 21.W .Breitung e t a l 9 Fusion Eng. and Design, 8(1989)323. 22.T .Terai, et a l 9 Fusion Eng. and Design, 8(1989)349. 23. W.Breitung, and H,Werle, Proc. 2nd I n t e r n a t i o n a l Symposium on the Fabrication and P r o p e r t i es of Lithium Ceramics, Apr. 23-27, 1989, I n d i a n a p o l i s .

Science of Ceramic Interfaces II J. Nowomy(Editor) 9 1994Elsevier Science B.V. All rights reserved.

277

Some aspects of grain boundary diffusion in oxides E. G. Moya Laboratoire de M~tallurgie, URA CNRS 443, Facult~ des Sciences et Techniques de St J~rSme, 13397 Marseille Cedex 20, France Abstract Some f u n d a m e n t a l and e x p e r i m e n t a l aspects of g r a i n b o u n d a r y diffusion in oxides are reviewed. Examples are given of solid state reactions associated with interface diffusion. The segregation aspect is examined. The different conditions required to m e a s u r e grain b o u n d a r y diffusion are discussed. Some experimental works are presented. 1. INTRODUCTION The history of a solid material, from its preparation to its ageing and deterioration, includes a series of steps, a large n u m b e r of them ruled by diffusion processes. Several types of diffusion are generally distinguished: bulk, dislocation, grain boundary and surface diffusion. In metals, extensive studies have shown that the last three kinetic processes are much faster than the first one. They are named short-circuit diffusion paths. Grain boundary and surface diffusion, which present some similarities, are often joined under the name of interface diffusion. In oxides, experimental diffusion data are not as numerous and clear as in metals. This fact has appeared in recent review papers on interface diffusion in oxides [1-5]. This scarcity is to be opposed to the n u m b e r of phenomena, including t r ans f or m at i ons , chemical reactions and physical properties which appear to be related to the presence of interfaces and to their diffusional properties. These phenomena play an i m p o r t a n t p a r t in the processing and the behavior of ceramic oxides. To u n d e r s t a n d them, the investigator must dispose of a whole set of data on kinetic and equilibrium parameters. In this paper, we have a look at the state of research in this field. Taking into account what has been done in previous reviews [1-5], we deal with some selected f u n d a m e n t a l and experimental aspects, with a view to exemplifying interests, difficulties and progresses related to this subject.

278 2. INTERFACE DIFFUSION AND SOLID STATE REACTIONS A number of experiments which are not direct diffusion ones involve some transport of matter in the solid state as a slow step. Comparison of the kinetics with available diffusion coefficient data is used to give information about the kind of diffusion which controls the process. We have selected two examples in which interface diffusion is thought to play a predominant part. 2.1. Oxidation of metals According to Wagner's theory [6], the growth of an oxide on a metal surface is controlled by the transport of reacting species through the already formed scale. The rate constant kp for film growth is expressed as a function of diffusion coefficients of metal and oxygen ions. Equation I Oxide MOa having predominant electronic conductivity D*: tracer self-diffusion coefficients of the oxide - ao2: molecular oxygen activity - f: correlation factor for diffusion - the limits of integration are the two interfaces of the film -

kp=

(

+

d(Ln ao2)

-

Atkinson [7] has reported several examples which confirm the validity of this theory: oxidation of Co in CoO, of Cu in Cu20 at 900 and 1000~ of Fe in FeO, of Ni in NiO above about 1000~ In these cases the oxidation rate is described by bulk diffusion of the cation, and the oxide growth occurs at the oxide-oxygen interface. In other cases where bulk diffusion is slow, the growth rates are faster than those predicted by bulk diffusion mechanisms: NiO below 1000~ Fe304 below 600~ Fe203, Cr203 and Al203 at all temperatures [7]. Oxides grown at the surface of metals have a polycrystalline structure. Pores are present, mainly at the triple points. During the growth it is not excluded that microfissures form under stress and are subsequently filled with new oxide. Hence these short-circuit diffusion paths may play a predominant role. Conclusive evidence can be obtained if the appropriate interface diffusion coefficient is put in Wagner's equation, which yields the oxidation constant. Atkinson used this procedure in the case of the oxidation of nickel [8]. The "effective diffusion coefficient" of Ni in the oxide film below 1000~ was written:

279

Eq~aation 2

5 D eft = (1-~) D + ~ D' = D + 2D'g

- ~ 9fraction of sites associated with the short circuit path - D and D', bulk and interface diffusion coefficients - 5 : boundary width g" grain size 2 5 / g ,, 1 -

-

The comparison shows t h a t kp measured and kp calculated using tracer diffusion data in polycrystalline scales are in agreement within experimental error. The correlation between oxidation and short-circuit diffusion has been extended to other systems, as documented in recent reviews [9]. The formation of protective oxide scales is of high interest with regard to corrosion. Alloys able to form A1203 or Cr203 layers are protected by these slow growing oxides. However, the growth rate is higher t h a n w hat would be expected from a l u m i n i u m and oxygen, or chromium and oxygen diffusion kinetics in corresponding single crystals. Short-circuit p a t h s (dislocations, grai n boundaries and microcracks) are likely the predominant means of transport of m a t t e r during growth. As a result, any p a r a m e t e r which would block or decrease diffusion along these paths would play an i m p o r t a n t role in the improvement of resistance to corrosion of such alloys. A body of work has been done to understand the beneficial effect of yttrium in corrosion resistant alloys. Among possible explanations, it was suggested t h a t this effect is due to a slowing of aluminium or oxygen diffusion in alumina grain boundaries [1011]. 2. 2 Metal.ceramic b o n d Metal-ceramic couples offer a wide range of applications and alumina is often used as the ceramic par t of the couples. Recent studies on copperalumina and silver alumina bond [12-15] have drawn attention on the diffusion processes which occur in alumina during the formation of the junction or s u b s e q u e n t t r e a t m e n t s . For example, in the so-called diffusion-bonding technique, the metal is put in contact with the ceramic and the system is heated at temperatures from 0.5 to 0.9 Tm (Tin= melting temperature of the metal). A minimal temperature is necessary to obtain the bonding. In this way, alumina binds to aluminium at 450~ silver at 600~ copper at 700~ nickel at 900~ iron at 1300~ platinum at 1400~ Whereas the temperatures involved are relatively high for metals, they are low for alumina, which melts at 2030 ~ (Tin/2 = 880~ Using Brown and Ashby correlations for diffusion constants [16] we can see that, in alumina, the diffusion rate is reduced by about 10 orders of magnitude from the melting t e m p e r a t u r e to half the melting temperature. At Tm/2, the bulk diffusion coefficient can be estimated by D = 10 -24 m2s 1. In metals, at 0.9 Tm, D =10 -13 m2s-1. Hence the mean penetration " f ~

has the following values"

280 Metals at 0,9Tin 2.10-5m

~D~ (m2s -1) (t= 10h)

Alumina at Tin/2 2.10-10m

This simple calculation shows that only some lattice planes near the surface "see" a diffusion effect in alumina in such conditions. The time required to reach a penetration of some micrometers would be several thousand years. Thus, in metal ceramic couples, alumina can be considered as opposing a "barrier" to diffusion penetration. Is this observed experimentally? Determination of diffusion profile by SIMS on a single crystal of alumina, after deposition of silver and a 10 hour diffusion annealing at 900~ yields a mean penetration depth ~]Dt - 10 nm [17]. However, mechanical tests on silver-alumina ceramic couples show that the tenacity of alumina in the zone close to the junction is changed on relatively large distances (= 50 ~tm at least) [15]. This can be explained by the fact that the ceramic is polycrystalline (grain size = 10 ~tm). Referring to radiotracer analysis of silver diffusion in polycrystalline alumina [18], it appears that a considerable amount of silver is transported far inside the material: over tens of micrometers (Fig 1). Then, when bulk diffusion is almost negligible, grain boundary diffusion becomes a preponderant factor; this effect must be kept in mind when one wants to know what happens to a metal ceramic bond after high temperature treatments.

A

~e.-.....o~e.,...~t. ~ ~ ~ , . , . . e~

f~o

oe ~

J~eo',,,,,.o.o...__..

rk" .

. ,

~

~O

O

9

O~,

a) b) c)

d) e)

f) g) h) i) j) k)

i

200

400

600

T = 827~ T = 916~ T = 916~

T -

= T = T

T = T = T = T = T =

950~

1202~ 13050C

1202"C !000"C 1102~ 1360~ 1402~

~ =-241h10ms t: ,, 48h t: - 240h10md.n

t = 337h x: - 2 h t: =, 1 h l l m i n

t: ~ ~ t ~

= = ,, =

19h18min 95h42nd.n 47h12nd.n 13h45m'in 4hl2nd.n

yuJmi

Fig 1: silver grain boundary diffusion in alumina, observed by radiotracer [18]

281 3. GRAIN BOUNDARY DIFFUSION AND DEFECTS IN OXIDES From the point of view of diffusion, it is necessary to consider the complexity of the defect population in oxides as compared to metals. One has to take into account: - intrinsic defects, formed according to the Schottky or the Frenkel model - extrinsic defects, brought by impurities or dopants - non- stoichiometric defects, in relation with the partial oxygen pressure - association of defects, likely to occur at a high level of concentration. The formation of these defects is treated like equilibrium reactions, ruled by equilibrium constants, so that the concentrations of all types of defects are highly interdependent. Interfaces bring an additional complexity to this picture. We have to consider the possible influences of three modifications of the crystals close to or at interfaces on interphase diffusion. &l. Space

charge

and intrinsic

segregation

In ionic compounds the distribution of the constituents may be different at an interface and in the bulk at sufficient distance from this interface. This difference is compensated by the space charge. The space charge region extends to about 10 to 100 interatomic spacings [19-20]. The contribution of the space charge to the boundary diffusivity in KC1 has been theoretically calculated by Yan et al.[21] These authors concluded that the space-charge diffusivity enhancement is very small. 3.2. Extrinsic segregation In ceramics, the problems associated with the presence of impurities deserve special attention. Impurities often have a low solubility in crystals, so that they tend to segregate at interfaces. Interface segregation plays a double role in interface diffusion. a) The segregation of impurities may change the defects concentrations in the boundary and introduces supplementary interactions compared to the bulk. b) When the diffusing element is an impurity, the segregation factor a must be included in the parameter P = D'.5.a expressing the boundary diffusivity in sectioning methods. A n u m b e r of theoretical and experimental studies on interface segregation have been carried out. Wynblatt and MC Cune [22], Egdell and P a r k e r [23] have reviewed the current state of knowledge for surface segregation in metal oxides, whereas Kingery [24] dealt with the chemistry of ceramic grain boundaries. Nowotny [25-26] has reviewed both the subjects of grain boundary and surface segregation. In the latter papers, the influence of non-stoichiometry is pointed out. Whereas the profile of solute enrichment as a function of the distance from the interface is sharp in stoichiometric compounds such as A1203, in largely non-stoichiometric compounds such as CoO a smoother dependence is observed. Thus the thickness of the interface

282 region may vary from one-two lattice layers to several tens. This observation could also have implications in interface diffusion. In non-stoichiometric compounds, the slab model generally used to describe interface diffusion would no longer be valid. As a basis for discussions in the rest of this paper, we report here some results of interface segregation in oxide materials. Combining Auger analysis, electron microscopy and electron probe analysis, Johnson and Stein [27] studied the chemistry of grain boundaries in sintered alumina. They found that calcium was present at the boundaries as a segregant in an equilibrium state, in the temperature range 1700-1900~ Their results were fitted to the McLean isotherm (derived from Langmuir type adsorption) (Eq.3). EQuation

3

_

c. exp(-AGs/RT) Cgb=

1+c. exp(-AGs/RT)

- Cgb: concentration at grain boundaries - c: concentration in the bulk AGs: binding free energy

They found: AGs (Ca/A1203)=-120 kJ mo1-1 Jupp, Stein and Smith [28] found calcium enrichment ratios in excess of 103 in alumina. Scanning transmission electron microscopy was applied by Li and Kingery [29] to study grain boundary segregation in doped alumina. Their results show that grain boundary enrichment varies widely with the nature of the solute. The authors analysed these results in terms of solute misfit in the lattice. The strain energy is function of the difference of radii of the solute and the solvent ions and can be estimated by: W = f ( rsolute- rsolvent)2 = f e2 This energy put in the McLean equation leads to the following dependence at a given temperature" l o g i cb ~ e2 The data plotted in this way fall on a straight line, which shows that the dominant driving force for segregation is the relief of misfit strain energy. Cook and Schrott [30] drew similar conclusion from SEM and AES study of Ca segregation. Their results along with those of Li and Kingery are reported on Fig 2.

283 _: 10 =~9 ss S

r

J

p

p

,/

/s

s/ sS

10 ~

,y,,,"/'~

Ca

p

Z

1

s/ s*"

Z O

y

>.

I0 z

o=

Z (~ uu >. k<( _J I.tJ

m

-

~o

s~s"

Sc

/ s SAp

s~ ~Z,n,,"" "n.~,,""'" i!~~, AI203 ..

Gram Dotalclanes

i STEM restat. Li and K i n g l y o AES result, this s ~ l y .

1

10

-1

'

0.0

s,, SS

0.2

.,

I .

0.4

I,

0.6

_

I

I

0.8

1.0

1.2

MISRT STRAIN PARAMETER. E

2

Fig.2: Grain boundary enrichment as a function of the misfit strain energy

[30]. Duffy and Tasker[31] have performed computer simulations of grain boundaries in ionic crystals. For NiO they calculated the segregation energies reported in Table 1. The effect of ion size on the interaction energy is exomined by comparing isovalent impurities with a range of ionic radii. The values of interaction energies and boundary impurity concentrations are reported in Table 2. It appears clearly that the magnitude of the interaction energy increases with the size mismatch of the defect. Table 1 Impurity segregation energie s near tilt bounda.ries in NiO[31] Boundary

Co2+

A13+

Ce4+

-8.7

" ' -17.4

-111

(320)/[001]

-21.2

-37.6

-167

(211)/[011]

- 11.6

-23.5

-39.5

( 122)/[011 ]

-5.8

-51.1

-71.4

(310)/[001]

"

284 Table 2 Interaction energies of isovalent impurities at the (211)/[011] tilt boundary in NiO, with the corresponding boundary impurity concentrations for Cbulk = 10-4 and T = 1000K [31] Interaction Cgrain boundary Impurity Ionic radii (nm) energy (kJ/mol) (Equation 3) Mg2+

0,0722

-9,6

3,2. 10 -4

Co 2+

0,0745

-11,4

3,9. 10-4

Fe 2+

0,0771

-23,8

1,7.10 -3

Mn2+

0,0836

-47,1

2,8. 10-2

Ca 2+

0,1021

-114,3

0,99

Sr 2+

0,1196

-207

1

Ba 2+

0,1377

-326

1

3.3. Precipitation along grain-boundaries When the impurity level is too high with regard to the solubility, a precipitation may occur along grain boundaries. I n t e r g r a n u l a r phases sometimes exhibit a glassy structure. The chemical diffusion coefficient t e m p e r a t u r e dependence in single crystal lmol% yttria-doped alumina presents an anomaly around 1550~ [32] This temperature is close to that at which intergranular precipitates dissolve in polycrystalline samples (T>1600~ The two effects are correlated and it is assumed that some short-circuit diffusion effect caused by the precipitates at lower temperatures results in increased diffusion. 4. GRAIN BOUNDARY DIFFUSION: DIRECT MEASURFAVIENT METHODS The f u n d a m e n t a l s of grain and boundary i n t e r p h a s e diffusion are presented and discussed in the book of Kaur and Gust [33]. In this paragraph, we recall briefly the general aspects necessary for basic understanding, and we discuss more extensively some special points, related mainly to oxide materials, to which our attention has been drawn in the course of works we have done on the suject. The sectioning method has been the most widely used to investigate grain boundary diffusion. Generally, the penetration of the diffusing element is followed by using radiotracers. Recently, some works have been done by SIMS. Analysis of the data is based on Fisher's model ( Fig 3).

285

X

g

X

.~

I I I i

D

'

Y

D

Y

Fig 3: Fisher's model for grain boundary diffusion

Fig.4: Diffusion in a polycrystalline solid according to the B-regime

The tracer diffuses into the material in two ways: -from the initial surface into the bulk, with a diffusion coefficient D -along the free surfaces, grain boundaries or dislocations with a diffusion coefficient D' >> D. Simultaneously, the tracer penetrates the grains by lateral bulk diffusion. The concentration in the slices parallel to the surface are determined by appropriate sectionings and activity countings, yielding the penetration profile. In adequate conditions, this profile has the typical shape shown on Fig.5.

Bulk diffusion ~ t,l

Intermediary zone I

Bulk diffusion &

Intermediary zone

0 m

Grain boundary diffusion ion diffusion

y or T1 y or Fig.5 Typical diffusion profiles of (a)grain boundary- (b) dislocation - diffusion experiments Part I: bulk diffusion Part II: intermediary zone Part III: grain boundary or dislocation diffusion zone The equations relative to these diffusion problems are functions of the p a r a m e t e r s listed in Table 3. The conditions necessary to obtain typical concentration profiles are listed in Table 4, the solutions in Table 5.

286 Table 3 Main diffusion parameters D D'

t 5 a

g 0~

= y/4-

y/4D'fa [~= 2 D ' ~ P--D'Sa Dd. a 2. o~

Bulk diffusion coefficient Boundary diffusion coefficient. It may be written also Dgb for grain boundaries, Dsb for dislocations arranged in sub -boundaries, Ddfor isolated dislocations diffusion time boundary width dislocation radii grain size segregation factor mean penetration reduced depth reduced distance from the boundary parameter which characterises the enhancement of diffusivity experimental parameter deduced from grain boundary sectioning analysis experimental parameter deduced from dislocation sectioning analysis

Table 4 Conditions required to obtain a typical penetration profile in a sectioning experiment Conditions to be fulfilled 1) First part 9can be missed if the thickness Ay of the successive removed slices is large compared with the mean penetration

Ay < < ' ~

2) B r e . m e : the profiles due to adjacent boundaries do not overlap. However, bulk diffusion occurs out of the boundary slab.

100 5 < ~

3) Separation of bulk and ~rain boundary diffusion profiles: on the first part,resulting from bulk diffusion from the surface, the decrease in concentration is abrupt. The contribution of grain boundaries is preponderant on the last part, which is flat.

D'5a ~= 2D.~_> 10

4) .Last part: this part can be missed when the concentrations are too low

C2 > background

< d/20

287 Table 5 Solutions of grain boundary diffusion ec uations Equation 4 Cl: concentration resulting from bulk diffusion from the surface C=Cl +C2 c2: average concentration resulting from grain boundary diffusion Equation 5

instantaneous source

rl 2 c 1 = co. exp (--~-) .Equation 6 Cl = co erfc (T1/ 2)

constant source

Equation 7

Suzuoka [34,35] instantaneous source solution

D' 5 a =

- dLn c2~-5/3 ( ~75 ) (~_)1/2

1~0,0~3

Equation 8 D'~a=

- d L n c2~-5/3 4D 1/2 0,661 ( d-~5 " (t)

Whipple[36]-Le Claire[37] constant source solution

REMARKS

a) concerning condition n~ Equations 7 and 8 show t h a t determination of diffusion p a r a m e t e r P implies that the bulk diffusion parameter D is known. As a matter of fact, the D'Sa sectioning method leads to the ratio ~ . The ideal conditions for grain boundary diffusion determinations would include preliminary m e a s u r e m e n t of the bulk diffusion coefficients on the same material and in the same range of temperatures. This case is rare. When the mean penetration ~ is small, it is difficult to obtain a precise value of D on the first part of the penetration profile. Generally D values are obtained using single crystals, and the investigator has to be sure that the possible differences in impurity level in the two kinds of material do not influence bulk diffusion. Another procedure which may result in errors is extrapolating to low t e m p e r a t u r e s data obtained in noticeably higher temperatures. Because of uncertainty on activation energy, D values may be erroneous by more than an order of magnitude. Ceramic oxides with very low diffusion coefficients constitute a particular problem for the classical radiotracer sectioning technique. When the mean penetration is lower than the thickness of removed layers, the very first part of the curve may be missed.

288 On Fig. 6 we have reported the available data on bulk cationic selfdiffusion coefficients in some oxides, as a function of the reduced temperature T/TIn, along with the line which correlates diffusion in fcc metals [16]. Data for A1203 have been updated from recent determinations on single crystals [38]

10-10

lo"11.] 10-12

Co/CoO

~o" 1 4.i!

T"

!

10-15~ u~ 10-16~ Q

I0-19~ i0-20.~ AI/AI203

~o-21.;

10-22 10-23 10-24

fcc metals

Cr/Cr203 !

1

2

Tm/T

-

3

Fig 6: Arrhenius plots of bulk self- diffusion coefficients of cationic species in some oxides, using a reduced reciprocal t e m p e r a t u r e scale. The correlated values for fcc metals are also plotted in view of comparison All A1203 [38], Co/CoO [39], Ni/NiO [40], Cr/Cr203 [41], fcc metals [16] From Fig. 6, it appears that the diffusion into A1203 et Cr203 is very slow. In these materials the problem of bulk diffusion determinations is particularly striking. This can be illustrated by the case of silver diffusion into alumina. First, bulk diffusion coefficients were determined in single crystals of aalumina with the radiotracer technique associated to mechanical sectioning [42]. An example of a curve is given on Fig.7. These curves reveal the influence of dislocations. Because of the decrease in the slope of the curve resulting from such short-circuits, it did not appear obviously that the very first p a r t of the penetration profile was lacking. As m a t t e r of fact the experimentals points belonged to part II of the curve drawn on Fig. 5b. This was demonstrated later t h a n k s to the SIMS technique [17]. An example of a curve is given on Fig.8. It can be seen t h a t p a r t I, corresponding to Fig. 5b, extents less than 0.1 ~m. Consequently, diffusion coefficients are by about 2 orders of m a g n i t u d e s m a l l e r t h a n those found previously by radiotracers. The Arrhenius equation for this system must be written: D (m2s "1) = 1 10-7 exp (-303 (kJ/mo1-1) / R T ) instead of D (m2s -1) = 2,.4 10 -4 exp (-331 (kJ/mo1-1) / RT)

289

Ar [arbitrary units t

~~..~48

~,,,15 h

,

.

48 h(vacuum)

..

9.. l

;

..

5 _

,

10

l

15

y[.m]

Fig 7: Tracer diffusion profiles of silver diffusion in monocrystalline a l u m i n a T = 916~ - Diffusion times are indicated on the plots. The activity reporte d here is the residual activity (at the surface of the sample)

105 ~. c

104

= 10

"I4..i -

Ag /single crystal alumina 4-

4-

900~ -I-

3

§ §

0

t" §

10 2

§ §

= 101

§247 §

c

§

§

§

10 0

§247247 §

§247 §

§

§

§

4-

i0 -I

9

"

"

"

"

"

9

"

"

I

5

"

"

~'

"

"

"

"

"

"

I

"

"

"

"

"

10 y (10

"

"

-8

"

"

I

15

"

"

"

"

"'"

"

"

"

"

20

m)

Fig 8: Diffusion of silver in monocrystalline alumina. T = 900~ - t = 10 hours This penetration profile was obtained by the SIMS technique.

290 This does not qualitatively change the results on grain boundary diffusion, as shown on Fig. 1. Quantitatively, there is a change in absolute values of grain boundary diffusion parameters (which are lowered 10 fold) and on the activation energy (which is modified by half the difference between real and wrong bulk activation energy). This gives: D' ~ a (m3s "1) = 1 . 9 10 -7 e x p (-307 (kJ/mo1-1) / R T ) instead of D' 5 a (m3s -1) = 9 10 -6 exp (-321 (kJ/mol-1) / RT) When the aim of a work is to compare grain boundary diffusivities in a material submitted to intentional modifications, it is of a fundamental importance to follow the bulk diffusion at the same time. In other cases, the results should be given in the form D' 5 a/~f-D rather than absolute D' ~ a values. A more serious mistake can be encountered when low bulk penetrations are associated with too large sectionings: in this case, the first and the second part of the curve (Fig. 5a) are lacking and the observable part, corresponding to grain boundary diffusion, is misinterpreted as due to bulk penetration. Although the theoretical equations 5 or 6 are not followed, sometimes a part of the curve may be assimilated to one of these equations. In this case, however, the influence of the diffusion time on the profile is lessened: instead of moving as a function of tl/2, a point at a given c/co moves as a function of tl/4. To conclude this discussion, the importance of fitting experimental sectionings to the mean penetration distance should be emphazised. Indeed, when the mean penetration distance is small compared to the thickness of slices obtained by abrasion, the near surface counts are neglected or undetermined. Confusion may result, and the curve may be erroneously attributed to a fast bulk diffusion. One must be particularly careful when studying diffusion at different temperatures. The time must be increased when decreasing temperature in order to keep ~ almost constant. d) concerning the sectioning process The sectioning method, which is based on the i n t e g r a t i o n of concentrations in slices parallel to the surface, is optimized when the analyzed area includes many grains. This is generally done with ceramics having grains of some tens of micrometers, the radiotracer technique allowing zones of some millimeters to be analyzed. When dealing with the sectionings, it should be kept in mind that the useful part of the penetration profile must be taken at sufficient depth to make the bulk contribution thoroughly negligible. It should also be verified that the plot logc2 = f(y) does not exhibit a gradual curvature. As a matter of fact, an intermediary part (labeled II on Fig. 5a) is sometimes observed. For example, the diffusion of Ca in polycrystalline NiO grown by oxidation of Ni at 976~ during 24 hours has lead to the profile reported on Fig.9 [43]. The plot exhibits three parts, labeled I, II and III. Part I is attributed to bulk diffusion.

291

10 4

~' ' Activiiy

(ms/s)

10 3

Ca/polycrystalline 976~ - 24 hours II

10 2

III

101

10 0

NiO

I

'

0

'

'

I

"

9

I.

10

.

.

.

-;

20

!

30

Y(l~m) Fig 9: Diffusion of Ca in polycrystalline NiO. T= 976~ - t = 24 hours Profile determined by the radiotracer technique

10 3 -Activity 10 2 ~

(cts/s)

Ca/single crystal NiO 976~ - 38 days

i

101

II

10 0 10 -1t, 0

10

20

30 y (.u.m)

Fig. 10: Diffusion of Ca in single crystal NiO T = 976~ - t = 38 days Profile determined by the radiotracer technique.

292 It was checked that the diffusion coefficient D = 1.4 10 -18 m2s -1 is close to t h a t m e a s u r e d on a single crystal (part I, Fig. 10) after diffusion at 976~ during 38 days (D = 1.3 10 -18 m2s-1). For the polycrystalline sample, part II was tentatively interpreted as the result of diffusion along isolated dislocations. The mathematical analysis given by Le Claire and Rabinovitch [44] indicates t h a t curves log c = f(y) are s t r a i g h t lines w h e n there is a diffusion along dislocations. Their slope allows one to calculate Dd a 2 a (Table 3). Applied to part II of curves obtained on polycrystalline oxide (Fig.9) and on single crystals (Fig.10), the analysis leads to very close values: Dd a 2 a = 2.3 10 -30 m4s -1 on polycrystalline oxide Dd a 2 a = 9 10-30 m4s -1 on single crystal oxide These values are in the same order of m a g n i t u d e and they m a y both represent diffusion along dislocations. Accordingly, the third p a r t (III) of the curve (Fig. 9) was t h o u g h t to represent grain b o u n d a r y diffusion, and was analyzed using Equation7. It leads to: Dgb 5 a = 9.4 10 -22 m 3 s -1 Assuming as usual t h a t the values for a and 5 are about 10 -9 m, one obtains values which are in the same order of magnitude for Dd a and Dgb a. Dislocations and grain boundaries would similarly increase the diffusion compared to the bulk. However, comparison of curves 9 and 10 shows t h a t the depths affected by the diffusional process may be different when dislocations only, or dislocations and grain boundaries act as short circuit diffusion paths. This example emphazises the advantage of the radiotracer-sectioning technique, which allows one to follow the penetration of the tracer on several tens of micrometers or more. c) c o n c e r n i n g ~

parameter

A proper choice of values of ~ D t is important when referring to condition 1 of Table 4, as detailed in remark 1. A better look at this table shows that~]Dt appears in all other conditions. Hence, the range of values given to this parameter has to be chosen carefully. When ~ is higher than the grain size, the penetrations due to adjacent boundaries overlap. The diffusion is no longer carried out in the B-regime (condition2) and the penetration profile does not allow bulk and grain boundary contributions to be separated properly. The inequality at the right of condition 2, and condition 3, work in the same sense, and lead to favoring small values of ~/Dt. However, a lower limit is imposed not only by conditionl as already seen, but also by condition 4. It must be kept in mind that, in the B regime, because of the very small thickness of the slab (one or few atomic spacings), the tracer detected in slices parallel to the surface is m a i n l y t h a t b r o u g h t by bulk diffusion process from grain boundaries. What are the factors which determine the average concentration?

293 These factors may be deduced from the detailed expressions of c. EQuation 9

constant source -co: constant surface concentration - cII: depends only on T1 ~-1/2 as shown by Suzuoka. Computed values of this parameter are reported on Fig.10

-

_

2 ~1

c= co [erfc~-+ ClI ~

EQuation 10

instantaneous source M: a m o u n t of tracer deposited at the surface cII: depends only on TI ~-1/2 as shown by Suzuoka. Computed values of this parameter are reported on F i g . l l

-

_

-

+

c= ~]~D t

-

To characterize the range of c values in the grain b o u n d a r y diffusion zone, on Fig.10 and 11 we have extrapolated the cII values computed by Suzuoka to TI ~-1/2 = 0 (Fig. 11). Although the Ce values thus obtained by extrapolation have no physical meaning, they can be considered as giving an indication of the upper limit of c2 values. We have thus obtained: Constant source

I n s t a n t a n e o u s source

Ce

5Nf-~

co

g

Ce cO

8~

or

Equation 11

Equation 12 Equation 13

8 M . , ] 2D (Dt)l/4 Ce = g-~' ~ D,5 a From equations 11 and 13, it appears clearly that very small values of would m a k e the m e a n concentration in the grain b o u n d a r y diffusion zone undetectable. The above discussion may be summarized by stressing the importance of p a r a m e t e r ~]Dt, the selected values of which m u s t be r a t h e r high to satisfy conditions 1 and 4, and r a t h e r small to satisfy conditions 2 (right part) and 3. When the typical curve schematized on Fig.5a is not obtained, the possibility of a wrong choice of the value of this parameter m u s t be considered. d) c o n c e r n | n g 13p a r a m e t e r A clear separation of bulk and grain boundary contribution implies t h a t value is high. Assuming a mean p e n e t r a t i o n ~ of about 10 -6 m, and 8 = 10 -9 m, condition 3 gives D ' a / D > 2 10 4 When a = 1 (self diffusion) D'/D > 2 10 4

294

c[[ i

m

c n P ~2

=

CONSTANT

, - ~ 2.5

n

%

a)

%

-.

SOURCE

INSTANTANEOUS

SOURCE

10"-

10"2

10"" -

10-3

10": o (3=Q= 9 t3=10 i,,=

0

I0" -

t

2

4

6

~p-i/2

e ~ - ~

9 13 = 1 0 _!

8

2

....

i .....

4

~p-~/2

!

6

w

8

Fig.ll: Plot of cII or cII ~-~ vs 1] ~-1/2 according to Suzuoka [35]

a) Constant source b) Instantaneous source

~.

1,500

.

1,400

. . . . . . . . +

1,300 '=

1,200

"5" ~,1oo ~,~ 1,000 ,

0,900

0

2

4

6

' 8

10

Fig.12: Theoretical slope- b(Ln c) / b( rl ~-1/2) vs 1] ~-1/2 computed by Suzuoka [35]

[

295 In the case of metals, from correlations on fcc metals [16,44], D ' / D ratio as high as 107 are calculated at half the melting point, using 8 = 10-9 m. According to data on NiO published by Atkinson [46], ratio D'/D would be equal only at 5 104 at T = Tm/2 = 844~ Obviously, in the case of self-diffusion, the grain-boundary e n h a n c e m e n t of diffusion requires special conditions (low temperatures, small m e a n penetrations, high activity of the deposit) to be observed. This is not the case when diffusion experiments are carried out with hetero-elements. The segregation factor a may increase [3 values significantly, so t h a t experimental conditions are less drastic t h a n for self-diffusion. This was observed in the case of Ca diffusion in NiO [43]. e) concerning the solutions

given by Eq. 7 and 8

The computed Suzuoka values reported on Fig. 11 show t h a t curves log c = f(vl) or fly) are not straight lines. The slope increases slightly when 11 or y increases. Le Claire has shown that a better linear fitting is obtained when plotting log c= f(y6/5). Hence, equations 7 and 8, which allow the determination of D' 8 a from b log c = b (y6/5) are generally used. We have used an alternative equation. Considering that the curvature of logc = fly) plot is undetectable in the limited range of values of the parameter Vl[~-1/2 investigated, we have expressed the slope in this range as a function of o t h e r p a r a m e t e r s . This has been done u s i n g the t a b u l a t e d values ~(Ln c) / ~( 11 ~-1/2) given by Suzuoka [35], reported on Fig. 12. We thus obtain the following equation EQuation 14 ~=

[0.96 + ( 0. 9216 + 0.454 p Ym )1/2] 4.6 p ~ - ~

Ym: middle of the penetration range in the second part of the curve. p = -(blog ClOJ by)

This equation leads to D' 8 a values which differ by less than 2% from those obtained with equations 7 and 8. The difference is well below experimental error. The a d v a n t a g e s of plots log c= f (y) compared to log c = f(y6/5) is that they are easier to read and to compare. 5. REVIEWS ON A1203, Cr203, NiO a n d CoO For complete reviews, the reader may refer to [46 to 50]. Data obtained by direct measurements on a-A1203, Cr203 and NiO are reported on Tables 6, 7 and 8 and in Figs.13 14 and 15. We have added the most recent work from our laboratory on a-A1203 and NiO, along with work done on CoO (Table 9 and Fig 16). The ratios D'a/D are reported on Table 10.

296 Table 6 Diffusion in a - a l u m i n a (D' values are calculated from Dgb 5 a or Dd a 2 a experimental values assuming 5 = a = 10 -9 m)-* SIMS technique Bulk Diffusion D (m 2 s -1) Element Authors Temperature Reference ~ A1 Paladino and 1670-1905 2.8 10 -4 exp(-477 (kJ tool-l)/RT) Kingery [51] A1 Lesage, leGall, 8 10-2o Philibert, 1610 Bernardini [38] 1650 3 10-19 Cr Lesage, Huntz 1200-1700 6.9 10 -11 exp(-265(kJ mol-1)/RT) Petot-Ervas [52] Cr* Moya et al. [53] 1000-1500 1 10 -10 exp(-288.6(kJ mol-1)/RT) Fe

1200-1700

Ag*

Lesage, Huntz Petot-Ervas [52] Lesage, Huntz Petot-Ervas [52] Moya et al. [17]

800-1150

1 10 -7 exp(-303(kJ mol-1)/RT)

Cu*

Moya et al. [53]

800-1100

0.11 exp(-411(kJ mol-1)/RT)

Ni

O* Element A1

Cr

3.6 10 -9 exp(-300(kJ mol-1)/RT) 2.5 10 -10 exp(-280(kJ tool-l)/RT)

Oishi and 1500-1700 5.6 10 -2 exp(-665(kJ mol-1)/RT) Kingery. [54] P r o t , Miloche 1521-1630 9.9 10 -3 exp(-626(kJ mol-1)/RT) and Monty [55] Dislocation, sub-boundary and grain-boundary diffusion Authors D' (m 2 s -1) Temperature oC Reference Lesage, leGall, Philibert, 1610 Dislocations 4 10 -13 Bernardini [38] 1650 9 10-13 Lagrange[56] 1200-1500 5 10 -3 exp(-341(kJ mol-1)/RT)

Fe

Lagrange[56]

Ni

Loudjani][57]

Ag

Moya et al. [17]

800-1150

Ag

Moya et al. [ 17]

800-1150

Dislocations 8.8 exp(-307(kJ mol-1)/RT) 1.9 102 exp(-307(kJ mol-1)/RT)

O

Oishi and Kingery. [54] P r o t , Miloche and Monty [55]

1450-1780

1.6 exp(-460(kJ mol-1)/RT)

1521-1630

Sub-boundaries 3 109 exp(-877(kJ mol-1)/RT)

O*

1200-1500

4 10-6 exp(-212(kJ mol-1)/RT) 1.1 10 -2 exp(-283(kJ mol-1)/RT)

297

10-10J . . . . . . !

:/3

E

E~ =.

t=.=

o E3

. 15

.......

10 111 ' "--" A l ( b ) 10-12.,.

__ _~ ......

O~

Ni

i

i

10-13..,.

.

.

.

.

,

_ _ _ j

__

10-14• 10

1 5 - -_~ 9

_

I..

,! ....

,

";\i ii

ii ............

10-18..--A~(b-).-X

"k~----

10-16_ 10-17__ 10-19-_

~o-2O.. 10 - 2 1

_

10-22

_,

4

|~

! \ O(a,b)

.

I!

..... !

""

.. _ : : , ~ x ---

Ni,

Cr(a).. .

.

.

'._~ . . . . .

"/lCJ

.

.

,

i

5

J

6

7

8

.

_ ~

.

%

9

10

104/T Fig. 13: Diffusion in A1203 Bold characters refer to bulk diffusion, italic characters to dislocation, subboundary or grain-boundary diffusion. A1 (a) [51], m Al(b) [38], Cr(b) [53], Fe [52],Ag [17] Cr(a) [52], Ni [52]: the lines representative of these two sets of data cannot be distinguished at the scale of the graph. O(a), [54], O(b) [55]: the lines representative of these two sets of data cannot be distinguished at the scale of the graph. a Al(b) [38]: these data were obtained on single crystals and refer to dislocation diffusion. The work is in progress. Cr [56], Fe [56], Ni [57] Ag [17]: Earlier published data [42] have been corrected using bulk diffusion coefficients determined by SIMS [17]. O(a) [54]

298 Table 7 Diffusion in Cr203 (D' values are calculated from Dgb 5 a experimental values assuming 5 = 10 -9 m) - * SIMS technique Bulk Diffusion D (m 2 S-I) Temperature Element Authors Reference ~ 1.2 10-2o flOO Atkinson and Cr Taylor[58] Cr*

Sabioni et al [59]

Cr*

King and Park

O*

King and Park

1200 1300 1100

10-22 4.8 10 .22 3.7 10-22

1100

8.3 10 -18

950-1100

1.4 10 -12 exp(-115(kJ mol-1)! RT)

[60] [60] Benlyamani et al

[61] .,

Cr*

Grain-boundary diffusion Authors Temperature Reference ~ Atkinson and 1100 Taylor[58] Sabioni et a1159] 1200 1300 King and Park 1100

O*

King and Park

Element Cr Cr*

D' (m 2 s -1) 3.8 10 -17 10-19 4.7 10 -18 8.6 10-17

[6O] 1100

2.1 10 -14

950-1100

1.3 10 -5 exp(-163(kJ mol-1)/RT)

[60] Benlyamani et al

[61]

299 _

10.101 11

i r,~

r

E

J

n

10-14 o

l

''i''-'

I

;

- - -

10

a

-....... ~-I 1500~

-'

.,,-13 / U

|

_!

-~

i

~---=-----=

.... "-

-

s

J

/,

'

'

'

~

:-" - .......

- ~-

-- -

;

=--i

!. . . . . .

L

l

-~-

mO

-

_! -

-

1n-20.;10-21

.

-

.......

:~.,--=

.

......

....

:_---

-.---:

:

'

-:

..

Cr

m

...... : . . . . . .

t,r-,l~

.....

~-

. . . . . .

!

(b)

A

.

.

.

.

--

" o

.

.

.

.

.

t

.

,

~

_

.

.

.

i: . . . .

_--_

:

.

.

.

!i

.

10 "22

.....

I0 - 24

.,

4

F .

5

,

:! ;

6

i .... ,

:

t 7

.

]

....... .

8

''

"

r

10-23

F

'

. . . . . . . . . . . . . . . . . .

. . . . . . . . . . ~--~- -~-: . . . . _. . . . . . . . . . . . . . . . . . . . . . j . . . . . . . . ' = 4Cr (b) ~--~-- --~- ....... ' ............C r ( a ) -.~ ..........

,

. . . .

i

!

_....... _~

e

...... i

i

i

I

9

_

-.......... 1000~

:

i

,

_

'"'=1

.................

15

10 "16

9

!

,

. . . .

_ ............

i........................" ___~ .....................

.

'i

..... : ! _

9

....

10

104/T

Fig. 14: Diffusion in Cr203 Bold characters refer to bulk diffusion, italic characters to dislocation, subboundary or grain-boundary diffusion m Cr(a) [59] &

Cr(b),

[60]

9 O [60], S [61], El Cr(a) [59], A Cr(b) [60] = 0[6O] -S [61]

300 Table 8 Diffusion in NiO (D' values are calculated from Dgb 5 a or Dd a 2 a experimental values assuming 5 = a = 10 -9 m) Bulk Diffusion Element Authors Temperature D (m 2 S-1) Reference ~ Ni Atkinson and 755-1400 2.2 10-6exp(-247 (kJ mol-1)/RT) Taylor [40] Cr Chen,Peterson 1192-1700642 8.6 10-7exp(-282 (kJ mol- 1)/RT) Robinson [62] Co Chen and 1179-1649 9.1 10-7 exp(-226.7(kJ mol-1)/RT) Peterson [63] Ca Tabet, Dolin and 1300-1600 6.8 10-7 exp(-328(kJ mol-1)/RT) Monty [64] O* Dubois, Monty 1190-1600 5 10 -3 exp(-540.3(kJ mol-1)/RT) and Philibert[65] Dislocation, sub-boundary and grain-boundary diffusion Element Authors D' (m 2 s -1) Temperature Reference ~ Ni Atkinson and 500-800 3 10-5 exp(-171.7(kJ mol-1)/RT) Taylor[46] Cr Chen and 906 to 1.7 10-11 to Peterson[67] 1263 1.6 10-9

(a) Cr Co

Atkinson and Taylor [66] Chen and Peterson[67]

600-1100 802 to 1043

Regime C (b) 6.5 10 .7 exp(-193(kJ tool-!)/RT) 5.1 10-12 to 1.4 10-9

Atki'nson and Taylor [68] Atkinson and Taylor [66] Amalhay et al [43]

5oo4oo

3.8 10 -4 exp(-172(kJ tool-l)/RT)

700-1100

Regime C (b) 6.3 10 -8 exp(-193(kJ mol-1)/RT) dislocations 9.4 10-12 2.3 10-11 8.4 10-13 1.4 10-11

(a)

Co Ce Ca

Ca

Amalhay et al

[43] O*

Atkinson, Pummery and Monty [69]

976 1000

976 1000 1100 to 1600

2.2 10-16 to 1.3 10-13

(c)

a) The activation energy is not given in the paper but is reported to be similar to that of bulk diffusion b) Regime C corresponds to the condition ~[Dt <<6 c) The activation energy is not given in the paper

301

10" 10-8 10-9 !

E

o tm t,,..

i

i

i

i

i

i

I

,,,

P

I

"----- Cr(a)

10- 10,,_ 10-1 1.~

i

10- 12.,_ 10-13-.(

1000~ ..........

I 500oc Ill

. .... j . . . .

Co(a)

Ca(b)!

1

I

.

[

.

.

.

.

. ,

,

i

'

I ~~i--

",~ i

. . . . .

:

10- 14.,_ 10 15._ ._Co,Cr(a)&..LJ 10 16._ 10- 17..~ ~ N

i----"XCrCb) ---!.... ......:;: ........ x. 1

10 18.,_ 10 19..,_

~o-2O._ 10-21 10-22.1-4

i

~------1 6

o

,

i='~L 8

10

~, 12

14

10

4

/T

Fig. 15: Diffusion in NiO Bold characters refer to bulk diffusion, italic characters to dislocation, subboundary or grain-boundary diffusion Ni [40] Cr(a) [62], Co [63]: the lines representative of these two sets of data cannot be distinguished at the scale of the graph. Ca(a) [64], 9 Ca(b) [43] 0[65]

Ni [46] Cr(a) [67], Cr(b) [66] Co (a) [67], Co(b) [68] 0 Ca(b) [43] O [69]

302 Table 9 Diffusion inCoO (D' values are calculated from Dgb ~ a or Dd a 2 a e x p e r i m e n t a l values a s s u m i n g 5 = a = 10 -9 m) Bulk Diffusion Element D (m 2 s -1) Authors Temperature oC Reference Chen, Peterson, Co 963-1683 5 10 -7 exp(-160.6 (kJ mol-1)/RT) Reeves [70] Kowalski et al 4 10 -14 Co 953 [71] Ni Hoshino a n d 1000-1450 3.3 10-6exp(-197.5 (kJ mol-1)/RT) Peterson [72] Ni Kowalski et al 1.5 10 -14 953 [71] Cr Chen, quoted in 1300-1500 vacuum [72] 8.2 10 -7 exp(-252(kJ mol-1)/RT) Cr Ca

Element Co Ni Cr

Ca

Kowalski et al [71] Kowalski et al [71]

953 1109 953 953 G r a i n - b o u n d a r y diffusion Authors Temperature oC Reference Kowalski et al 953 [71] Kowalski et al 953 [71] Kowalski et al 953 [71] 1109

Kowalski et al [71]

953 953

1.2 10 -17 7 10 -17 2.9 10 -14 1.6 10 -14 D' (m 2 s -1) 6.3 10 -9 2.6 10 -9 1.5 10 -11 2 10 -lo 4.4 1043 2.4 10 -8

303 !

I

~

]r

=1 ~ ...... l

1000 10-8

E ,=,

E3

!

i

.

~0- 9 . . . . . .

i

! ............... i ...............

i0.I0

j_

'-

9

........

= =-

~

'~

.....

s

,o

-

, &

--:-

".

i=

t..__

o

. "~..

--- .

i,i

i

.

.

Cr

.

.

0

Co(b)

"

........ n

Ni(b)

....

.

.

.

.

.

10 10

13

"~~-~_,._ I i i ~.~--~.~_ 'i Ni(a) I = ............. "'~.~_ ._.-~.. ~

10

1 4

[ ....... ~

10

15

10-1 10-17

6

,

-;-

.

_

10"18/

.

.

.

.

.

.

.

.

.

i

i

_ . . . .

., . . . . . . . .

i

5

9

~ ' e

. . . . . . . . . .

.'

! ......... F

9 ' i' 4

,

.

.

_=_

i

. . . . .

...... __.

.

.

' .

.

,,

.

1"

.

.

Cr(b)

,~

i

J~ .......... 1 L. . . . .

Co(b)

. . . . . . .

~

.

I

. . . . . . ~i ....................

|

. . . . . . . .

,

9 . . . .

E'

T

.~....... : . -

-i-c-r(,i)

.

! __.._ ~_!A

,,-~ 12

,.I

.

(b)

~!.C0
_--

E]~,l

if

....

~

!

c.-

~! ,,, ,.~

.

.

.

.

.

.

.

i

.

i

.

!!

L!A

.

'

.

_

9

:...

;

~

t" 6

7

8 10

4

9

10

/T

Fig. 16: Diffusion in CoO Bold characters refer to bulk diffusion, italic characters to g r a i n - b o u n d a r y diffusion Co (a) [70], 9 Co(b) [71] Ni (a) [72] = Ni(b) [71] Cr(a) [72] & Cr(b) [71] m Ca [71] 0

Co(b) [71]

m N i (b) [71] A Cr(b) [71] 1:3 Ca [71]

304 Table 10 Ratio D' a / D for self and impurity diffusion ("pseudo-self diffusion" refers to an impurity cation close to the matrix cation) Self or I Tm/T D' ~/D Impurity I Tm/T D' ~ / D "useudo-self' q diffusion i t diffusion 1.2 5 104 1.3 A1/A1203 5 106 " Fe/AI203 dislocations ,,,

O/A1203 sub-bound.

1.1

1 106

Ni/A1203

1.3

3 107

Cr/AI203

1.3

4 105

Ag/A1203 dislocations

1.3

1.3 109

Ag/A1203

1.3

1.4109

S]Cr203

1.9

1.5 105

Cr/Cr203

1.8 1.9

1 103 2 105

O/Cr203

1.9

2.5 105

Ni/NiO ...

Co/COO" Ni/CoO

4 104

1

Cr/NiO

1 104

Co/NiO

1.3 105

Ca/NiO

1.8

9 106

1.7

1.6105

Ca/CoO

1.7

1.5 106

1.7

1.7 105

Cr/CoO

1.7

1.2 106

COMMENTS AND DISCUSSION Compared with that on metals, the amount of work done on grain boundary diffusion in oxides is very limited. In this part of the paper we intend to examine only direct diffusion measurements, which would support indirect determinations made by sintering, creep .... With this restriction it appears that the state of research in this field is still at its beginning. Among "common" oxides, only three have been the subjects of systematic investigations" A1203, MgO and NiO. To a lesser extent, data have been given on Cr203, Cu20 and stabilized ZrO2. When concerned with nuclear material, data exist on ThO2 and U02. A number of review papers have been devoted to particular oxides: MgO and NiO [46,47], (z-A1203 [48,49]. On a more general level, the available data on

305 bulk and grain boundary diffusion in oxides have been collected and examined in the reviews of Monty and Atkinson [3] and of Ddchamps and Barbier [4]. They also appear in the Handbook of Grain and Interphase Boundary Diffusion Data [50]. From discussions and comments given in the review papers, it appears that even for the three most investigated oxides, the situation is far from being satisfactory. Works from different sources often lead to contradictory results. The main reasons behind the discrepancies are the presence of impurities and of second phases along grain boundaries, and the imperfect s t r u c t u r e of polycrystalline oxide samples, which could be responsible for artefacts [4]. a) Recent set of data for A1203 and Cr203~ The new values recently determined for A1 diffusion in single crystals of A1203 [38] are lower by 2 or 3 orders of magnitude than those found in 1962 by Paladino and Kingery for polycrystalline alumina. The diffusion coefficients of A1 are close to those of oxygen. Short-circuit diffusion paths appear to play a role even in single crystals. These new data call for a re-examination of the mechanism of oxidation of aluminium and of alloys forming protective layers of alumina. The SIMS technique, which makes it possible to measure very small diffusion coefficients, has opened a field of investigations for A1203 and Cr203. By this technique data have been obtained for the systems O/A1203, Cu/Al203, O/Cr203, Cr/Cr203. In the case of Ag/A1203, it was found that values obtained by the radiotracer technique were too large. The dislocation and grain boundary diffusivities have been re-evaluated according to SIMS data. The "corrected" values are reported on Table 6. b) comparison between dislocation and grain-boundary diffusion data When bulk diffusion coefficients are low (diffusion in alumina, Ca diffusion in NiO), curves typical of short-circuit diffusion are found both with single crystal and polycrystals. The dislocation related curves may have a higher slope than the grain-boundary curves, because of different geometrical conditions. However, it should be noted that the diffusivities obtained for the two kinds of defects are very close. c) first values about grain boundmT diffusion in Up to now, there was no data on this process in CoO. The diffusion of different elements in the polycrystalline scales of the oxide grown on highpurity metal has been studied recently in our laboratory. Diffusions in NiO and CoO present similarities but the bulk diffusion coefficients are higher in CoO because of its degree of non-stoichiometry. As a consequence, grain boundary diffusion can be investigated only in a limited range of temperatures close to 950~ [71]. Chromium is an exception since it diffuses very slowly in CoO, making investigations at higher temperatures possible. It is not possible to work at T<900~ where the stable form of cobalt oxide is Co304.

306

d) the segregation effect The influence of segregation is obvious when considering the ratios D'a/D reported on Table 10. For example, Ca, which because of its size, tends to segregate at interfaces [Table 2], is characterised by D' a / D values higher than those of the cation of the matrix, both in NiO and in CoO. [43,71].

e) activation energies When activation energies are available, they do not follow the rule observed in metals: Q' < Q. In a number of cases, the equality or even the inverse inequality is observed: Q' > Q. The values of activation energy are affected by large uncertainties and may differ by several kJ when varying the number of experimental points in Arrhenius plots. However, the number of cases where the relation Q'> Q. was observed calls for discussion. As a matter of fact, it is not clear how an increase in diffusion rate in grain-boundaries compared to the bulk can be obtained with a higher activation energy. The activation energies may be written as: Q = AHf + A Hm + C AHf, AH'f, A Hm and A H'm: formation and migration enthalpies of the defect responsible Q'= AH'f + A H'm +AHs + C' for diffusion C and C': Arrhenius dependences of the correlation factors AHs: segregation.enthalpy In metals, the increase in diffusion coefficients is thus explained by a decrease in the formation and migration enthalpies of vacancies along the grain boundaries. With impurities, the apparent activation energy Q'= -~ (LogD' 5 a) / ~ (l/T) also includes the segregation enthalpy AHs which is a negative term. All these terms contribute to the decrease in activation energy which is observed experimentally. In oxides, we have to imagine a more complex process to explain the observed grain boundary diffusion activation energy. As a first possibility, the unavoidable impurities may affect the grain boundary diffusion. It can be assumed t h a t high concentrations of impurities or precipitates in the boundaries slow the atomic motion. In this case, the effect would be more pronounced at lower temperatures, since the impurity segregation and precipitation equilibria are favored by low temperatures. The result would clearly be an artificial increase in the apparent activation energy. A second hypothesis is based on the possibility of a grain boundary migration during the diffusion annealing. It has been shown by Glaeser and Evans [73] that this process leads to an apparent grain boundary diffusion coeefficient higher t h a n the real one. Thus, as the migration occurs preferentially at high temperatures, the apparent activation energy is too high. As a third possibility, we suggest that, because of the rearrangement of the crystal structure at the interfaces, diffusion along these interfaces cannot be interpreted in the same way as in the bulk. For example, the concentration

307 of interstitials could be higher than in the bulk. The simple analysis given above for metals would no longer be valid, and one would have to take a complex diffusion mechanism into account. Theoretical and experimental data are insufficient at this date to discuss this question further, but it would be of high interest to confirm and explain the increase in diffusivity along the grain boundaries which is associated with equivalent or higher activation energies in comparison with the bulk. 6. CONCLUSIONS From this survey, it can be concluded that research on grain boundary diffusion in oxides must advance simultaneously on different grounds. The practical aspect is important when using industrial ceramics; diffusion studies have to be carried out on these materials to elucidate their behavior. It should be noted that the amount of matter transported via interfaces in fine grained ceramics may be significantly higher than that transported through the bulk. Fundamental studies provide a guide for improving the qualities of technological ceramics. Theoretical studies need to be supported by good quality experiments. In the oxides, progress in the understanding of grain boundary diffusion is linked to the development of refined analysis techniques, and to disposing of pure and well-defined materials. 7. REFERENCES

10 11 12 13 14 15

A. Atkinson, Solid State Ionics, 28-30, 1377-1387, (1988) Hj Matzke, Surfaces and Interfaces of Ceramic Materials, L.C. Dufour et al eds., 241-272 (1989) C Monty and A. Atkinson, Cryst. Latt.Def. and Amorph. Mat., 18, 97120 (1989) M. D~champs and F. Barbier, Science of Ceramic Interfaces,, J. Nowotny ed., Elsevier Science Publishers, 323-369 (1991) Hj. Matzke, Phil. Mag. 64, 1181-1200 (1981) C. Wagner, Z. Phys. Chem. B, 21, 25 (1933) A. Atkinson, Mat. Sci. and Technology, 4, 1046-1051 (1988) A. Atkinson, R.I. Taylor and A.E. Hughes Phil. Mag. A, 45, 823-833 (1982) W.W. Smeltzer, External and Internal Surfaces in Metal Oxides, L.C. Dufour and J. Nowotny eds., Trans Tech Publications, 29, 151-172 (1988) A. M. Huntz, Mat. Sci. and Techn. 4, 1079-1088 (1988) M. LeGall, Thesis, Paris, 1992 C. B~raud, Thesis, Lyon, 1986 M. Courbi~re, Thesis, Lyon, 1986 C. B~raud, M. Courbi~re, C. Esnouf, D. Juve, D. Treheux, J. of Mat. Sci. 24, 4545-4554 (1989) B. S~rier, Thesis, Lyon, 1991

308 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47

A.M. Brown and M.F. Ashby, Acta Metall.,28, 1085-1001 (1980) F. Moya, E. Moya, D. Juvd, D. Treheux, C. Grattepain, to be published L. Badrour, E.G. Moya, J. Bernardini and F. Moya, J. Phys. Chem. Solids, 50, 551-561 (1989) W.D. Kingery, J. Am. Ceram. Soc., 57, 1-8, 1974 and 57, 74-83 (1974) S. M. Mukkhopadhyay and J. M. Blakely, Science of Ceramics Interfaces, J. Nowotny ed., Elsevier Science Publishers, 205-225, 1991 M.F. Yan, R.M. Cannon, H.K. Bowen, and R.L. Coble, J. Am. Ceram. Soc., 60, 120-127 (1974) P. Wynblatt and R.C. Mc Cune, Surface and Near-Surface Chemistry of Oxide Materials, Nowotny and Dufour eds., Elsevier Science Publishers, 247-279, 1988 R. G. Egdell and S.C. Parker, Science of Ceramics Interfaces, J. Nowotny ed., Elsevier Science Publishers, 41-78, 1991 W.D. Kingery, Pure Appl. Chem. 56,1703-14 (1984) J. Nowotny, Surfaces and Interfaces of Ceramic Materials, Dufour et al eds., Kluwer Academic Publishers, 205-239, 1989 J. Nowotny,.Science of Ceramics Interfaces, J. Nowotny ed., Elsevier Science Publishers, 79-204, 1991 W.C. Johnson and D.F. Stein, J. Am. Ceram. Soc., 58, 485-488 (1975) R. S. Jupp, D.F. Stein, D.W. Smith, J. Mat. Sci., 15, 96-102 (1980) C.W. Li and W.D. Kingery, Adv. Ceram., 10, 379-393 (1984) R. F. Cook and A.G. Schrott, J. Am. Ceram. Soc., 71, 50-58 (1988) R..M. Duffy and P.W. Tasker, Phil.Mag., 50, 143-154 (1984) G. Petot-Ervas, D. Deweirder, M. Loudjani, B. Lesage and A.M. Huntz Advances in Ceramics, 23,125-135 (1987) I. Kaur and W. Gust, Fundamentals of Grain and Interphase Boundary Diffusion, Ziegler Press, Stuttgart, 1988 T. Suzuoka, Trans. Jpn. Inst. Met., 2, 25 (1961) T. Suzuoka, J. Phys. Soc. Jpn, 19, 839 (1964) R.T.P. Whipple, Phil. Mag.,45,1225 (1954) A.D. LeClaire, Br. J. Appl. Phys., 14,351 (1963) B. Lesage, M. LeGall, J. Philibert, J. Bernardini, to be published W.K. Chen, N.L. Peterson, W.Y. Reeves, Physical Review, 186, 887-891, (1969) A. Atkinson and R.I. Taylor, Phil. Mag.A, 39, 581 (1979) A. C. S. Sabioni, B. Lesage, M. Huntz, J. Besson, C. Dolin and C. Monty, Colloque de Physique, 51, C1-611-C1-616 (1990) L. Badrour, E.G. Moya, J. Bernardini and F. Moya, Scripta Metall., 20, 1217 (1986) M. Amalhay, E. Moya, F. Moya, to be published E.G. Moya, F. Moya and J. Nowotny, Interface Segregation and Related Process in Materials,, Trans. Tech. Pub., Zurich, 239-283, 1991 A.D. LeClaire and A. Rabinovitch, J. Phys. C: Solid State Phys.., 14, 38633879 (1981); 15, 3455-3471 (1982); 16, 2087-2104 (1983); 17, 991-1000 (1984) W. Gust, S. Mayer, A. B6gel and B. Predel, J. Physique C4, 46, 537 (1985) A. Atkinson and R. Taylor, Phil. Mag. A, 43, 979-988 (1981) M. Ddchamps and F. Barbier, Non-Stoichiometric compounds, J. Nowotny and W. Weppner eds., Kluwer Acad. Publishers, 221-236, 1989

309 48 49

50 51 52 53 54 55 56 57 58 5~

60 61 62 63 64 65 66 67 68 69

70 71 72 73

E.G. Moya and F. Moya, External and Internal Surfaces in metal oxides, L.C. Dufour and J. Nowotny eds., Trans Tech Publications, 237248, 1988 E. Moya and F. Moya, Non-Stoichiometric compounds, J. Nowotny and W. Weppner eds., Kluwer Acad. Publishers, 363-387, 1989 I. Kaur, W. Gust and L. Kozma, Handbook of Grain and Interphase Boundary Diffusion Data, Volumes 1 and 2, Ziegler Press, Stuttgart,1989 A.E. Paladino and W. D. Kingery,J. of Chem. Physics, 37, 957-962 (1962) B. Lesage, A.M. Huntz and G. Petot-Ervas, Rad. Effects, 75, 283-299 (1983) F. Moya, E. Moya, D. Juv~, D. Treheux, C. Grattepain and M. Aucouturier, Scripta Met. and Materiala, 28, 343-348 (1993) Y. Oishi and W. D. Kingery,J. Chem. Physics, 33,480 (1960) D. Prot, M. Miloche and C. Monty, Colloque de Physique, 51, C1-10271033, (1990) M. H. Lagrange, Thesis 3 ~cycle, Universit~ Paris Sud, 1982 M.K. Loudjani, Thesis, Universit~ Paris Sud, 1992 A. Atkinson and R.I. Taylor, Transport in non-stoichiometric compounds, Plenum Publishing Corporation, 285-295,1985 A.C.S. Sabioni, B. Lesage, A.M. Huntz, J. Besson, C. Dolin, and C. Monty, CoUoque de Physique, 51, C1-611-616 (1990) W.E. King and J.H. Park, Colloque de Physique, 51, C1-551-556 (1990) M. Benlyamani, F. Ajersch and G. Kennedy, J. Electrochem. Soc., 136, 843-846 (1989) W. K. Chen, N.L. Peterson and L.C. Robinson, J. Phys. Chem. Solids, 34, 705-709 (1973) W.K. Chen and N. L. Peterson,J. Phys. Chem. Solids, 33, 881-892 (1972) N. Tabet, C. Dolin and C. Monty, Rev. Int. hautes Temp. Refract. Fr., 19, 413-416 (1982) C. Dubois, C. Monty and J. Philibert, Phil. Mag. 46, 419-433 (1982) A. Atkinson and R.I. Taylor, J. Phys. Chem. Solids, 47, 315-323 (1986) W.K. Chen and N. L. Peterson, J. Am. Ceram. Soc, 63, 566-570 (1980) A. Atkinson and R.I. Taylor, Phil. Mag. 45, 583-592 (1982) A. Atkinson, F.C.W. Pummery and C. Monty, Transports in nonstoichiometric compounds, G. Simkovitch and S. Stubican eds., Plenum Press, New York, 359-370, 1985 W.K. Chen, N.L. Peterson and W.T. Reeves, Phys. Rev., 18{;, 887-891 (1969) K. Kowalski, F. Moya, E. Moya and J. Nowotny, to be published K. Hoshino and N. L. Peterson, J. Phys. Chem. Solids,, 45, 963-972 (1984) A.M. Glaeser and J. W. Evans, Acta Metall. 34, 1545-1552 (1986)

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Science of Ceramic Interfaces II J. Nowomy(Editor) © 1994Elsevier Science B.V. All rights reserved.

311

C h e m i c a l a n d S t r u c t u r a l A l t e r a t i o n in t h e S u r f a c e L a y e r s of O x i d e s and Sulphides Roger St.C. Smart with Pawittar Arora, Robert Hayes, Byung-Sub Kim, Clive Prestidge and John Ralston Particle and Surface Technology Research Group, University of South Australia, The Levels, South Australia 5095

Al~tract Surface analytical and high resolution microscopic techniques have been used to study the surface modification of oxide and sulphide materials resulting in major alteration of their surface reactivity. These surface modifications include restructuring of the first few unit cells of the surface layers, new surface "phases", and the effect of impurity atoms on oxidation site initiation. Ionic oxides (e.g. MgO) have been found to restructure on exposure to water vapour or solution to produce {100}-based pits and protusions with dimensions of a few unit cells. Some semiconducting oxides (e.g. CoO) also r e s t r u c t u r e in AFM images but the scale (10-20nm) is larger and the restructured regions are rounded and not obviously {100}-based. These oxides show initial dissolution kinetics in which the dissolution rate increases with increasing pH during the restructuring process. Low temperature plasma reactions have been used to produce silicate structures in the first few oxide layers on nickel metal. The silicate layers are strongly passivating against hydrolysis and acid attack. Non-stoichiometric iron sulphide surfaces (i.e. Fel. xS) oxidise in air and solution by the loss of iron ions from the sulphide lattice to form hydroxide overlayers. The underlying sulphide lattice has been shown to form a crystalline, defective tetragonal Fe2S3 surface phase in which linear change of Sn atoms have a S-S distance similar to elemental sulphur. Oxidation and reaction of PbS surfaces in air and in solution has been extensively studied using the STM technique. Impurity sites in the cleaved (100) surfaces initiate oxidation in preference to low-coordination sites at edges and corners of steps and ledges. Synthetic, pure PbS shows much slower oxidation and preference for the low-coordination sites. The implications of these surface modifications will be discussed in relation to dissolution and leaching kinetics, flotation separation of sulphide minerals and weathering of oxide and sulphide surfaces.

312 1.

INTRODUCTION

The unifying theme in the study of oxide surfaces described here has been the physical and chemical alterations respo.nsible for changes in surface reactivity and dissolution kinetics. Some of our contributions to the theory of oxide reactivity and dissolution have been summarised in major reviews [1-3]. In particular, this work has included: derivation of a systematic approach to prediction of oxide dissolution rates in acidic solution based on the solid state structure of the oxide and solution conditions; the identification of rate-determining mechanisms for each class of oxide in different pH conditions dependent on reactions such as charge-transfer control, oxide protonation, reactant and product diffusion, surface area alteration during dissolution and base-catalysed hydrolysis; -

the observation and explanation of major restructuring of ionic oxide surfaces immediately after immersion in aqueous solution (i.e. before dissolution); experimental evaluation of different possible oxide protonation models based on log (rate per unit surface area) versus pH dependence; dependence of semi-conducting oxide dissolution kinetics on defect properties of the oxide phase and correlation with semiconductor theory, and; correlation of dissolution kinetics of ZrO2, TiO2, Al203, SiO2 in alkaline solution with relative rates of base-catalysed hydrolysis reactions. Dissolution rates per unit surface area for binary oxides, in ionic, semiconducting (i.e. p-type or n-type) or covalent insulating categories, under similar conditions (i.e. pH, solution temperature, extent of dissolution) vary by more than ten orders of magnitude. Different mechanisms obviously govern the dissolution kinetics of oxides in different categories and oxides in the same category with markedly different dissolution rates. Consequently, different theories are required to model the dependence of the rates on solution or solid, particularly solid surface, properties. Even for the same oxide, the kinetics of dissolution in different stages of initial, advanced and n e a r - s a t u r a t i o n conditions require different theoretical approaches, e.g. electrical double layer stabilisation; reaction (e.g. protonation) rate control; and diffusion rate control. A previous review [1] has summarised the characteristics of five different kinetic regimes, namely:- near-saturation (equilibrium) Nernstian conditions; diffusion rate control (e.g. stirring); surface reaction rate control (acidic); solid state charge-transfer rate control; and base catalysed hydrolysis reaction rate control (alkaline). The processes or mechanisms contributing to dissolution for each oxide category are summarised in Table 1 reproduced from reference [ 1].

313 Table 1 OXIDES IONIC (ions at surface)

$

/ p-type conductivity / \ high low

insulating

$

COVALENT (lattice network of covalent atoms)

SEMICONDUCTING (partial ionicity at surface)

$

$

$

\ n-type conductivity / \ high low

$

$

insulating

$

OXIDE DI$SOLUTI0~q RATES (Examples)

$ Very fast MgO CaO BeO

$

$

$

Fast CoO MnO TiO

Slow NiO CuO UO2

Fast ZnO CdO

$

Slow a-Fe203 MnO2 SnO2

$

Very slow a-A1203 TiO2 SiO2

DISSOLUTION PROCESSES Surface reconstruction

$

Charge transfer to surface

Base-catalysed hydrolysis*

$

$

Alteration of

Ion formation

Ion formation

surface charge

(M n+ and 02-)

(ie M(OH) y+)

$

$

$

M n+ transfer

M n+ transfer

M(OH) y+

to solution

to solution

transfer to solution

$

$

H + transfer to 02- ion

H + transfer to 02- ion

$

$

OH- (or H20) transfer to solution / *eg>Al

O

\ Al< +OH-

$ H + transfer to oxide species

$

OH- (or H20 transfer) to solution

OH- transfer to surface

(H20) >A1-OH + - O - A l <

-~

2(>A1-OH) + OH-

314 A list of factors influencing oxide reactivity is also summarised in Table 2 from reference [ 1]. Table 2 ROLE OF SOLID OXIDE

Factor~

Relationship to Reactivity

Degree of ionic or covalent character in bonding

Nature of surface reaction; ion formation in the surface; charge transfer; protonation

Nature and concentration of semiconducting lattice defects

Charge transfer;

Nature and concentration of charge alloy dopants Morphology and surface faceting

Nature and concentration of 9 active (defect) surface sites (eg valence states, coordination) Specific features of surface states

Charge transfer; surface

Surface reconstruction; concentration of surface sites; protonation; surface exchange Ion formation; protonation; effect of redox potential 9 Redox effects; adsorption; surface conductivity

ROLE OF SOL(~rION Concentration of reactants, : Kinetic regime (initial, advanced, products, ionic additives near-saturation); potential control pH

Protonation, base-catalysed hydrolysis; potential control; diffusion

Redox potential

Ion transfer; concentration of active sites

Complexing agents

Ion transfer; adsorption; surface charge

Surface active agents

Surface blocking; ion transfer; surface charge

Temperature

Reaction-mechanism category; surface reconstruction (& phase changes)

315 The mechanisms of surface reconstruction and alteration of surface charge together with factors relating to morphology and surface faceting will be particularly relevant to studies described in this chapter. The inhibition of rate-determining mechanisms for dissolution of semiconducting oxides will also be considered. Incorporation of oxides into ceramics, glasses, thin films and bonding layer structures has comprised a major extension of the fundamental studies of these binary oxides into areas of application [4-11]. The alteration of sulphide surfaces in ambient conditions and in solution generally produces oxide, hydroxide and oxyhydroxide products on their surfaces which, in some cases, can control the subsequent reactivity and dissolution of the surface layers. These oxidation reactions have provided the link between the understanding of oxide surfaces and that of sulphides. In particular, a combination of surface analytical techniques (i.e. XPS, scanning Auger spectroscopy, SIMS, STM, AFM) with analytical high resolution SEM (i.e. field emission) has been used to study the chemical and structural alteration of sulphide surfaces. Oxidised products have been identified in three distinct forms, namely: as colloidal flocs of Fe(OH)3 precipitated from its solution; as oxidised fine (i.e.
316 can silicate structures be formed in thin oxide layers on metal surfaces at low temperatures in the same form as those previously found in the surface layers of bulk NiO [ 15]? what are the chemical and physical forms of the reaction products and altered sulphide surfaces after non-stoichiometric Fel.xS is reacted in air or solution [ 16]? what are the growth kinetics and distribution of oxidation products on PbS surfaces in air and solution [17]? How are these parameters affected by defect and impurity sites in the PbS surfaces? -

These examples will illustrate a range of different mechanisms which can result in major alteration of the surface properties of both oxide and sulphide materials. 2.

SOLU'I~ON INDUCED RESTRUCTURING OF OXIDE SURFACES

We have previously shown [1, 2] that ionic oxide surfaces almost totally restructure on exposure to water vapour or solution. The surfaces of MgO formed by burning magnesium in air, as nearly perfect {100} interfaces (Fig. la) alter immediately after immersion in solution to produce {100}-based rectangular pits and protusions with dimensions of a few unit cells (see Fig. lb). Less developed restructuring of similar form has also been observed with MgO smoke surfaces cooled to <10~ and transferred rapidly to the electron microscope vacuum without immersion in solution (Fig. lc). The resulting condensation of water on the surfaces was sufficient to cause restructuring of the initially almost-perfect {100} surfaces without dissolution. It was confirmed that the dramatic restructuring immediately after immersion in solution took place before any significant (i.e. <1 monolayer) dissolution could be measured using dissolved Mg 2+ as the indicator of such dissolution. Other experiments [1, 2] have also shown that this restructuring occurs for all MgO surfaces irrespective of their initial form, e.g. as smoke or decomposed hydroxide or carbonate materials. Hence, the extensive changes to the ionic oxide surfaces take place as a result of contact with water in vapour or solution form before any significant dissolution into bulk solution has occurred. The restructuring occurs immediately after immersion in solution but much more slowly in contact with water vapour in air. Parallel experiments using the computer controlled dissolution system, in which dissolution rates as log rate (mol m - 2 s - 1 ) versus pH plots were calculated from changes in the [H +] concentration with time using a fast pH electrode, gave reproducible plots in which the initial regime, i.e. immediately after immersion in solution, showed a significant increase in log rate with increasing pH [1, 2] (see Fig. 2). Hence, during and immediately after the period in which the dramatic surface restructuring takes place, there is a

317

Figure 1. Phase contrast TEM micrographs of MgO particles a) fresh MgO smoke (top left); b) after condensation of water vapour on to smoke surfaces at -10~ for <10s (lower centre) and c) after immersion for <5s in pH3 nitric acid solution (top right). The bar in each micrograph is 20nm. Refer references [1,2].

318 significant reduction in [H+] concentration represented both by the rise in pH and the rise in log (rate) calculated from the overall reaction MgO + 2H 2+ -~ Mg 2+ + H20 CaO, a second ionic oxide has the same characteristics, namely restructuring of the surface immediately on immersion in aqueous solution and initial dissolution kinetics in which the log rate per unit surface area increases with increasing pH [2]. The rate in the second case is again calculated from the measured change in pH r a t h e r than from m e a s u r e m e n t of dissolved Ca 2+ suggesting that the surface restructuring occurs before any substantial bulk dissolution has taken place. The process of surface r e s t r u c t u r i n g described above has two major consequences. The magnitude of the surface electric field is massively reduced by the increase in the dimensions of the distance across which the potential decays from the solid surface to the bulk solution. On the initial flat {100} surfaces the 8 term in the AV/8 field expression is relatively small and well defined, depending on the solution conditions in the electrical double layer. After restructuring, producing projections on the scale of 1-5nm, the electric field, and consequently the surface free energy, is dramatically lowered. Accompanying this change is the production of a v a s t l y i n c r e a s e d concentration of surface defect sites in the form of low-coordination edge and corner positions. These sites have been shown in both theoretical [18-22] and experimental [23-26] studies to be responsible for the reactivity of ionic oxide surfaces in the adsorption of simple molecules (e.g. H2, CO, H20) from the gas phase. For instance, a defect-free {100} MgO surface is unable to dissociate H2 or CO molecules [20-22] and does not produce surface roughening due to H20 adsorption. Defect sites are propagated across the surface by the adsorption of H20 or H+ but the initial defect nucleation sites are required for this reaction [23-26]. Hence, the observed restructuring represents rapid propagation of reaction across the surface simultaneously producing a highly defective set of low-coordination surface sites. The ionic oxides generally dissolve relatively rapidly (i.e. >10 -4 mol m -2 s-l). A second group of semiconducting oxides, still retaining quite high ionic character, dissolve more slowly than MgO and CaO (i.e. >10 -6 mol m -2 s -1) but still show the same initial dissolution kinetics as the ionic oxides. In particular, both CoO and MnO have an initial solution phase in which the log rate per unit surface area increases with increasing pH [1] (see Fig. 2). In contrast, the relatively very slowly dissolving semiconducting oxides, like NiO (i.e. 10 -10 mol m -2 s -1) do not show an initial dissolution regime of the same form but rather a monotonic decrease in log rate with increasing pH. There is therefore a suggestion that the former group of semiconducting oxides may exhibit surface restructuring of the same form as that found for MgO. There has been some difficulty in testing this hypothesis previously because two experimental conditions are required:- an experimental technique capable of imaging surface structure at the sub-nm level and; ideally, initially flat

319 surfaces of the oxide before solution or water contact. These conditions were uniquely met by the combination of TEM and the electron-transparent MgO smoke particles but similar preparations are not available for the semiconducting oxides. The relatively new technique of atomic force microscopy (AFM) has however provided a technique capable of imaging at the atomic scale so that the preparation conditions of the oxide surfaces are not as demanding.

[Mg++] 2 !

-5.4 (D t._

v

0

O

4 !

5x 10 -4 I

n

-5.6 -5.8 o

-6.0

-6.2

I 1

5

I 2x10 ~

I 2,0

15x10 ~

5

[ M g ++1

[Mg++]

-6.4

10

I 2.5

' 10

' 10x10 6

[ M g ++1

I 3.0

l 3.5

l 4.0

pH Dissolution rate (mol m 2 s 1 ) of MgO smoke in HCIO 4 01%

~

3%

[] 5%

Figure 2. Log dissolution rate (mol m'2s -1) vs pH for MgO smoke in HC104 at different starting pH values of 2.02, 2.47, 2.89 and 3.46 for 1%, 3% and 5% dissolution as marked. The subsidary axes show Mg 2+ concentrations. Note the initial increases in dissolution rate with increasing pH at each starting pH value. Refer references [1,2]. We have been attempting to obtain similar evidence of the restructuring of CoO surfaces on exposure to water vapour or liquid using the AFM technique. The CoO sample was synthetically prepared in the form of relatively large single crystals (i.e. a few mm dimensions) from the melt by Dr. J. Nowotny. The Nanoscope II AFM instrument had been previously calibrated giving atomic imaging of mica, graphite and other surfaces with observations of restructuring due to reaction at the sub-nm level.

320 In the first set of experiments, a fresh fracture face of the CoO crystal was mounted on a sample stub. Representative images of the initial surface were examined from a number of different regions of the fracture face. The fracture face was then dipped in distilled deionised water, rapidly removed and dried using a tissue to extract water without contact with the exposed fracture face. A hot air blower was then applied and it was re-examined at a number of different regions of the surface. Figs 3a and b present large area images before and after solution immersion respectively. Even at this low magnification, it is immediately a p p a r e n t t h a t some form of surface alteration has occurred, evidenced by the raised regions in Fig. 3b and by the change in the z distance parameter from 30 to 100nm between the two images. At higher magnification, using areas selected for their relative flatness in the previous two images, Figs 4a and b show similar effects with an increase of the z parameter from 5.0nm to 20nm. Surface features in both images, however, appear relatively rounded despite the NaC1 structure of CoO.{100}-based facets similar to those found with MgO are not evident. Images obtained at higher magnification t h a n those in Figs 4a and b were very blurred with features extremely difficult to determine. The reason for the obscurity of these images is not clear given the ability of the AFM system to image other surfaces at atomic level in air. It may be an instrumental difficulty associated with the morphology of the fracture surface seen in Fig. 3 or, alternatively, a feature of this type of semiconducting oxide. On all areas examined from both the initial surface and t h a t after solution immersion, the same general features were found. They can be further exemplified in line scan diagrams of Figs 4a and b where the difference in surface roughness between the initial dry surface (i.e. r = 5.4 + 0.5nm) and that after solution immersion (r = 17+ 3nm) is clearly apparent. The AFM images described here were obtained in air whereas the TEM images for MgO restructuring in Fig. 1 were obtained in -10 -5 Torr vacuum. Experiments to test the effect of evacuation on the AFM images are currently being conducted but it m us t be noted t h a t differences in surface structure arising from this factor are actually relevant to the dissolution mechanisms. Preliminary experiments in which the surface changes are imaged in-situ after water is added to the CoO fracture faces have confirmed the larger scale and rounded surface structures. The first image is obtained ~2min. after solution addition and this roughened structure does not change significantly over periods >60min. The initial surface alteration apparently occurs in the first 2min. of immersion. Hence, it is clear that CoO surfaces do substantially restructure after solution immersion during the period in which the solution kinetics show rapid uptake of H + in similar form to that found for MgO dissolution kinetics. However, the surface restructuring is different in form and magnitude. The restructured regions are considerably broader and higher (i.e. on the scale 10-20nm compared with 1-5nm for MgO) and the distinct {100} faceting is not found.

321

Figure 3. AFM images from a freshly cleaved CoO single crystal face before (upper) and after (lower) immersion in distilled water (pH 5.5) for 60s (recorded after drying in air). Different regions (1000xl000nm) of the surface are shown but care was taken to select areas representative of surface features in each case.

322

Figure 4. AFM images from the same CoO samples as Figure 3 at higher magnification (i.e. 250x250nm) showing rounded surface features after immersion (lower image).

323 Rather, the surface appears to produce rounded features with less evidence of distinct low-coordination sites. Correspondingly, the dissolution rate of CoO is two orders of magnitude lower than that of MgO. This evidence suggests that surface restructuring of many oxides of both ionic and relatively conductive semiconducting type may occur immediately after immersion in bulk solution. The simplistic model of flat {100} faces for these oxides in solution is therefore manifestly incorrect even on single crystal surfaces and the corresponding reactivity to be anticipated on these altered oxide surfaces will be much higher than that derived from measurements of the dry crystal surfaces before immersion. 3.

F O R M A T I O N OF S U R F A C E SILICATE S T R U C T U R E S IN OXIDE FILMS

Bulk nickel oxide can be modified by reaction with vapour transported silica at temperatures as low as 550~ [15]. The silica, transported in the form of SiO:H20 complexes, is incorporated into the surface layers of the NiO to form silicate structures evidenced by an infrared absorption at ~1045cm -1 attributable to orthosilicate, diorthosilicate or chain silicate structures. The surface morphology of the NiO is d r a m a t i c a l l y a l t e r e d to become predominantly {111} (Fig. 5). The {111} faces are atomically smooth but show separated pyramidal pits still based on {111} faceting (Fig. 5b). The silicate structures in the surface layers produce a highly unreactive surface for which the dissolution rates per unit surface area in dilute acid are decreased more t h a n 100 times, i.e. from -10 -9 mol m -2 s -1 to <10 -11 mol m -2 s -1. Bulk concentrations of silicon <10 -3 mol % were measured after reaction of the nickel oxide in evacuated quartz tubes. The reduction in dissolution rate, in addition to the observed changes in surface morphology, can be attributed to changes in the electronic properties of the nickel oxide since the dissolution kinetics for this oxide are controlled by charge transfer in the p-type semiconductor [27,28]. Sensitive electron diffraction and phase contrast microscopy studies have not indicated any separate phase formation or phase boundaries in the nickel oxide surfaces after vapour phase reaction with the silica (Fig. 5c). XPS measurements of the surface layers gave silicon contents of the order of 9 mol % confirming that the majority of the silicon is segregated into the surface layers. This previous work with bulk NiO suggested t h a t silica might also be incorporated into the oxide films on nickel metal surfaces conferring similar surface reaction and passivation. Experiments with vapour phase deposition in sealed quartz tubes, sol/gel deposition of silica and plasma reactions have all been trialled on oxidised nickel surfaces. The sol/gel deposition is difficult to control generally resulting in silica layers which obscure the interface with the oxides. Information on the nature and extent of the reaction is therefore difficult to obtain non-destructively. Both the vapour phase deposition and the plasma reaction have, however, verified the formation of silicate structures in the oxide surface layers. The plasma reaction is more practical in application and will be briefly described here.

324

Figure 5. a) SEM micrograph of NiO single crystals annealed in air at 1450~ (top left); b) SEM image of NiO particles after annealing in sealed evacuated (10 -5 Torr) quartz tubes for 4h at ll00~ (top right); c) TEM micrograph of typical smoke particle from Fig. 5b) showing very smooth {111} faces without phase differentiation (lower centre). The surface of the pure nickel is cleaned using the method of Tegart, i.e. a mixture of nitric, sulphuric, orthophosphoric and glacial acetic acids [29]. The surface of the clean nickel is oxidised in a H20 plasma (100W, 5 mins) with the oxide surface layer depth verified using XPS. Subsequent reaction of the

325 oxidised nickel surface is carried out in a plasma containing both H20 and tetraethoxysilane for varying periods of time. The H20:TEOS volume ratio can be varied between ~10:1 and 20:1 depending on the required rate of deposition of reactive silica species on to the oxide surface. The process of fragmentation of the TEOS in the plasma produces hydrocarbon, oxyhydrocarbon and a variety of SiOn radicals. The SiO radical is one of the major species and this species may combine with H20 or with CH2 species in the plasma. The limited extent of oxidation of the nickel and subsequent reaction with silica has made observation of silicate structures using FTIR difficult and equivocal. The reaction has been confirmed, however, using XPS and, particularly, the modified Auger parameter. This parameter is derived from the difference in kinetic energies of the Auger and photoelectrons plus the energy of the exciting radiation as in: a" = KE(Auger) - KE(photoelectron) + hv = KE(Auger) + BE(photoelectron) determined for the same spectrum on the same sample. The modified Auger p a r a m e t e r removes difficulties with binding energy determinations due to surface charging and chemical shifts in this p a r a m e t e r directly reflect changes in screening energy [30]. Compilations of the type in Fig. 6 have been d e t e r m i n e d for silicon in a variety of different compounds as oxides, hydroxides, silicates, aluminosilicates, silicones etc.[31]. In our experiments, the KE (KLL) Auger peak and Si2p BE have been determined for a variety of plasma treatments as a function of time. Some of the Si2p spectra are illustrated in Fig.7 for silicon surface concentrations varying from 1.3% to 21% as the time of reaction is increased from 2mins to 30mins. The combined parameters contributing to the Auger parameter are plotted in Fig. 6 to illustrate the changes in the chemical nature of the silicate film with time. In the first stage of plasma treatment, the combination of Auger and photoelectron p a r a m e t e r s places the silicate structure close to that corresponding to hemimorphite, a hydroxysilicate of Si2075- type. With continuing reaction, the structures become more like chain and layered silicates as in spodumene. Finally, after 30mins of reaction, the surface is covered with a layer corresponding to silica species, as in silica gel, oxidised silicon quartz etc., and there is no remaining exposure of nickel at the surface. Preliminary experiments with dissolution kinetics have these surfaces are highly passivated even after 2mins Quantitative results from these experiments are currently will be available in later publications. Some preliminary Table 3.

also confirmed that of plasma reaction. being processed and results are given in

326

Table 3 Dissolution Rates (mol m "2 s -1) of N i c k e l Plate and N i c k e l Plasma-Treated with H20:TEOS in pH2 Nitric Acid at 22~ T i m e (min) N i c k e l Plate Nickel:TEOS

300 2.0x10 -5 < l x l 0 -s

500 2.6x10 -5 < l x l 0 -s

1611

1000 3.2x10 -5 6.9x10 -s

1713

1712

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1609

,

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103

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Si2p Binding Energy, eV Figure 6. S o m e examples of compounds h - hemimorphite p - pyrophyllite HS - hydroxysodalite g - silica gel pw - p s e u d o w o l l a s t o n i t e q - quartz w - wollastonite o x - oxidised Si sp - s p o d u m e n e v - Vycor glass The three points m a r k e d f~ represent m e a s u r e d Si2p, Si(KLL) and Auger p a r a m e t e r s from the oxidised nickel surface for different t i m e s of reaction in the H 2 0 : T E O S plasma.

327

10

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BINDING ENERGY, eV Figure 7. XPS spectra from the Si2p region for the oxidised nickel surface after a) 2min; b) 5 min and c) 30 min exposure to the H 2 0 : T E O S plasma. The atomic concentrations of Si were determined to be 1.3%, 3.6% and 21.3% respectively. Hence, oxide surface layers on metal may be able to be selectively reacted with silica to produce highly passivated silicate structures a few monolayers thick. This selective reaction m a y be useful in practical applications where corrosion of the metal surface can be inhibited by prior plasma treatment.

328 4.

SOLU~ON-INDUCED RES'I~UCTURING OF SULPHIDE ~ A C E S

Like oxide surfaces, sulphide mineral surfaces undergo major changes in the surface layers due to structural and chemical r e a r r a n g e m e n t s after immersion in solution. The iron-containing sulphide minerals, e.g. pyrrhotite (Fel.xS), pyrite (FeS2), chalcopyrite (CuFeS2), are known to react in solution by the loss of iron ions from the sulphide lattice to form hydroxide overlayers [3236]. XPS studies, in addition to the increasing surface concentration of O ls signals (in the form of hydroxide) with time of immersion in the solution, also show S2p spectra exhibiting a second doublet shifted to higher binding energy attributed to the sulphur signal of a highly defective, sulphur-rich surface of the original sulphide. In particular, Buckley, Hamilton and Woods [32-34] have reviewed previous studies of pyrrhotite surfaces subjected to reaction in air and in solution using XPS, linear potential sweep v o l t a m m e t r y and chemical analyses. Fresh fracture faces of this sulphide mineral oxidise immediately on exposure to ambient air or to solution forming an overlayer of iron (III) hydroxide or hydrated oxide covering an iron-deficient sulphide lattice, the metal content of which decreases with increasing exposure or oxidation time. Sulphur-oxy anions, e.g. sulphate, are only formed after air exposure for relatively long times (i.e. > 10 days) and these species are generally soluble in bulk solution. Pyrrhotite fracture surfaces reacted in acid solution [34] had severely iron-depleted surfaces layers of the sulphide with the majority of the S2p intensity in the higher binding energy doublet shifted 2.0eV from the sulphide peak. This binding energy is close to but still 0.2eV less than that of elemental sulphur. Hence, it is inferred that soluble Fe(II) species are formed leaving the iron-deficient sulphide surface exposed to further solution attack. In a recent study within our group [16], we have examined the structure and chemical n a t u r e of the altered sulphide surface after acid solution attack in more detail. Of particular interest was the possible restructuring of the irondepleted surface layers to new, more stable structures as the loss of iron continues during the reaction. For instance, continuing depletion of the Fel.xS lattice might have led to the formation of a pyrite-like (FeS2) by surface restructuring. Some evidence for this possibility was obtained by studying the dissolution rates of the pyrrhotite surface as a function of the duration of acid attack. At 60~ in pH 1.7 perchloric acid, the initial dissolution rate was found to be 8.5x 10 -6 mol m-2s -1 in air-equilibrated solution but the initial rate reduces rapidly leading to a dissolution rate approaching zero after ~30mins with total dissolution of the pyrrhotite sample comprising <20%. Comparison with the dissolution kinetics of pyrite under the same conditions of acid solution and temperature shows that the initial dissolution rate is two orders of magnitude slower, i.e. 3x10 -8 mol m -2 s -1 but this very slow rate falls to zero after <30mins of solution attack. Hence, the continuing dissolution of pyrrhotite in acid solution exhibits dissolution kinetics similar to t h a t of a pyrite-like surface u n d e r the same conditions. The activation energy, e s t i m a t e d from temperature studies, is ~75kJ mo1-1 consistent with surface reaction control rather than diffusion in solution or in the surface layers.

329 Fig. 8 illustrates the components of Fe2p3/2, O ls and S2p curve fitted signals from XPS spectra of pyrrhotite surfaces ground in 0.05M perchloric acid at 25~ under argon for 10mins before transfer to the spectrometer. The Fe2p3/2 spectra have contributions from ferric oxide/hydroxide (near 712.5eV), ferrous oxide/hydroxide (near 709eV) and FeS (near 707.4eV). The O ls spectra have three components corresponding to the oxide (near 530.1eV), the hydroxide Fe(2p 3/2)

714.9 713.8 712.7 711.6 710.5 709.4 708.3 707.2 706.1

O(Is) LU Z

536

535

534

533

532

531

!

530

529

i

528

i

F .. F~.xS ~-x~

...............

~l i

172.0 170.6 169.2 167.8 166.4 165.0 163.6 162.2 160.8 159.4 158.0

BINDING ENERGY, eV Figure 8. Fe2p, O ls and S2p spectra from XPS of pyrrhotite (Fel-xS) surfaces (before ion etching) ground in 0.05M HC]O4 at 25~ under argon for 10rains before transfer to the spectrometer as a slurry. Curve-fitted components are shown. The original spectrum is the broken line with the fitted components and calculated envelope as full lines. Refer reference [16].

330 (near 532.0eV) and condensed water (probably ice near 534.0eV). The S2p signal has two doublets with the predominant high binding energy doublet shifted 1.8eV from the doublet due to the unaltered sulphide (i.e. $2p3/2 near 161.1eV). The percentage of the reacted surface layers in the form of each of these species, together with carbonate and sulphate species, has been estimated [ 16]. It is important to note that the same binding energy near 707eV in Fe2p spectra is found for all of the iron sulphides including pyrrhotite and pyrite so that this peak does not distinguish between different structural forms of the sulphide. We have established that, for heavily oxidised pyrrhotite surfaces showing no remaining sulphide peak n e a r 707eV, this peak is restored in Fe2p spectra by acid attack in periods from 10-75 mins in solution. The form of this sulphide layer after removal of the oxidation products is therefore of interest. X-ray diffraction (XRD) was used in conjunction with XPS to obtain information on the reacted sulphide surface layers after acid attack. XRD traces from the reacted surface were compared with a 50:50wt.% mixture of pyrrhotite and pyrite minerals. The acid-treated pyrrhotite showed four diffraction peaks corresponding to pyrite at 20 angles of 33.35 ~, 38.7 ~, 43.50 ~ and 55.82 ~ but the major pyrite peak at 47.80 ~ was missing and the pyrite peak at 38.70 ~ was extremely weak whereas it is normally the strongest pyrite peak. Calculated powder diffraction patterns for a number of alternative structures were compared with the acid-treated XRD patterns. This revealed t h a t a tetragonal intermediate Fe2S3 structure, based on a disordered form of pyrite

Figure 9. Calculated lattice structures for a) pyrite (FeS2) (left) and b) the tetragonal intermediate Fe2S2 structure (FeSI-I.5) (right) viewed down the [001] zone. The statistical occupancy of S sites for the Fe2S3 structure have been shown. Refer reference [16}.

331 (i.e. NaCl-type) r a t h e r t h a n pyrrhotite (i.e. NiAs-type) was successful in predicting most features of the altered phase. This structure is illustrated in Fig. 9b where the statistical occupancy of sulphur sites creating a sulphur sublattice with shortest S-S bonds (1.91A) similar to $8 (2.08A) is found. In particular, linear chains of Sn polysulphide like structure as in Fig. 9b are likely to be responsible for the electronic rearrangement giving rise to the high binding energy shifted doublet in the S2p signal. The Fe2p3/2 signals, however, are still likely to a p p e a r n e a r 707eV because the n e a r e s t neighbour coordination in all three structures is essentially the same. As iron is extracted from the sulphide lattice during continuing reaction in solution, the structural rearrangement may then proceed as: pyrrhotite (hexagonal P3 21) FeSI-I.5 (tetragonal P42/nnm) -~ FeS2 (cubic Pa-3). The continuing depletion of iron may be responsible for structural rearrangement in the surface layers reducing the reactivity of the surface towards oxidation and dissolution as found for pyrite surfaces. Pyrrhotite minerals encountered in base metal sulphide ores exhibit a variety of responses to flotation separation but, at acidic pH values and short conditioning times in solution, good flotation can be achieved without the addition of collector (i.e. hydrophobic adsorbate) molecules [37,38]. Longer conditioning times, however, result in extensive surface oxidation and the flotation response is severely diminished. Some of the reasons for this flotation behaviour can now be explained. After acid reaction in deoxygenated solution, the S-rich phase is formed on a significant proportion of the surface in the early stages of reaction. This phase provides hydrophobic character for bubbleparticle attachment and flotation. With time in air-equilibrated conditions, the surface continues to react and dissolve, iron oxide/hydroxide species are formed both in-situ and by precipitation, and the exposure of both the pyrrhotite and the altered hydrophobic sulphide surface is lost. The surface becomes hydrophilic and bubble-particle attachment is greatly reduced. F u r t h e r detail on the surface reactions and restructuring can be found in reference [16]. 5.

OXIDATION OF S U L P H I D E SURFACES: E F F E C T S OF I M P U R I T Y SITES

Like the iron sulphides, surface oxidation and dissolution of galena PbS is of practical importance since it is one of the most common sulphide minerals concentrated by the flotation method. An u n d e r s t a n d i n g of the surface concentration and distribution of different oxidation products on the mineral surface before the addition of surface-active reagents, is central to control of the subsequent bubble-particle attachment and flotation. As with other sulphides, the corrosion process in air and in solution is also important in applications of t h e s e m a t e r i a l s in s e m i c o n d u c t i n g and l e a c h i n g / w e a t h e r i n g / t o x i c i t y applications. In a first study of this system, scanning tunnelling microscopy (STM) was used to examine the oxidation of fresh fracture faces of stepped {100} PbS surfaces in air over periods up to 270min. [17]. The initial surfaces revealed

332 large areas of atomically flat (100) faces with one and two unit cell steps as shown in Fig. 10a. There are some few atomic scale (i.e. <0.3nm) defects or reaction sites seen in the images as vertical protusions even in the first image after cleavage. In constant current mode, they appear as sharp, local changes in the tip distance d within a unit cell parameter. Natural galena crystals supplied by Wards Natural Science Establishment (New York) were used in these measurements. Electron probe microanalysis show that they contained, on average, minor impurities as 720ppm Fe, l l 0 p p m Zn and 25ppm Cu. In STM, the tunnelling current density (j) is approximated by: j = (e 2 ko/4~hd) V exp(-kod) where ko is related to the work function. Hence, changes in the work function at the local impurity sites is apparently responsible for the initial vertical projection. Remarkable changes can be seen in the images as a function of time as illustrated in Fig. 10. Systematic growth of oxidation products as roughly conical structures on the same sites are seen on the galena surface. In the early images, the lateral dimensions of these surface features are <0.6nm but, after 270min. (Fig. 10d), the dimensions of the largest reaction structures were ~9nm diameter and ~4nm height with overlap between adjacent oxidised regions becoming apparent. No clear preference for growth at step edges (i.e. low coordination sites) was found but careful matching of images revealed that the defect sites in the {100} surfaces provide the initial points for reaction. Parallel XPS monitoring of the surface after successive reaction times suggests that limited adsorption of peroxide (O-) and hydroxide (OH-) species occurs initially but that the predominant products in the early oxidation phase are carbonate species from CO2 reaction with these adsorbed oxygen species. After 120min., hydroxycarbonate species are also formed. The S2p spectra show only sulphide species at all times with no evidence for sulphide-oxygen species during the reaction in air. A schematic illustration of the possible mechanism of carbonate formation, together with adsorbed oxide and hydroxide species, is shown in Fig. 11. The reaction of the galena surfaces in air to produce carbonate and hydroxycarbonate species clearly distinguishes this sulphide mineral from the iron-containing sulphides discussed in the previous section where the predominant oxidation product formed on the surface is ferric hydroxide. More recent studies have compared the oxidation and solution of galena derived from synthetic (>99.999% pure) and natural (100-200ppm impurities) sources. These studies have been carried out in-situ using a modified version of the CSIRO STM, a system allowing environmental control of pH, Eh and purging gas with the solution in contact with the surface. Commercial glasscoated Pt tips were used in these studies and images from both constant current and constant height mode were examined.

333

Figure 10. STM 3-dimensional rotated images from a 70x70nm area of a) freshly cleaved PbS surface (gal000); b) the same surface after 150min exposure to air (gall50); c) after 210min in air (gal210) and d) after 250min in air (ga1250). Note the increases in the vertical scales, i.e. 1.8, 3.1, 3.8, 5.8nm respectively. There are slight shifts between the images b u t the individual growing oxidation produces reposition at the same place in successive images when these shifts are corrected. Refer reference [ 17]. A significantly different mechanism, compared with reaction in air, is found in aqueous solution as illustrated in the sequence of images in Fig. 12 where galena was reacted under air - equilibrated water at pH6 for periods up to 90min. The main features of the reaction are the appearance of dissolution pits growing in x, y and z directions with dimensions corresponding to multiples of unit cells, i.e. 0.6nm, 1.2nm etc. With natural galena, it is noted t h a t reaction

334

Figure 12. STM images from PbS recorded in-situ in pH6 (air-equilibrated) solution. The images were recorded initially (i.e. <5min) and after 30, 70, 90 minutes as marked. Note the dissolution pitting, i.e. no increase in the z parameter, in contrast to the reaction in air.

335 takes place relatively indiscriminantly on both flat faces and step edges. With synthetic galena, both in air and in solution, oxidation is significantly slower (i.e. approximately halved) and occurs primarily on the step edges in the initial phases of attack. The presence of impurities in the natural galena appears to control the oxidation initiation and subsequent propagation. The mechanism for oxidative dissolution of galena, as proposed by Buckley, Hamilton and Woods [32,39,40], involves two stages, namely: PbS k_~ Pb 2+ + S 0 + 2e4S 0 + 3 02 + 4e- ~2 2S2032or S O+ 2 02 + 2e-

~2

SO42~

The dissolution of PbS is apparently congruent since XPS does not detect the build-up of oxide and hydroxide products on the surface and the S2p spectra remain unaltered from that of the original sulphide. This suggests that the reactions with rate constants k2 are much faster than the initial oxidation kl so that the meta-stable sulphur phase does not exist for any significant time on the galena surface. As with the iron sulphides this sulphur phase may be in the form of rapidly reacting polysulphide species or structures. 100II A

80-

oxygen air 9 nitrogen

E , 0m 0 ej} {/)

u O cO

E

60-

40-

20-

0

99

0

w

w

w

w

]

lOO

200

Time / min

Figure 13. Extent of PbS monolayer dissolved (i.e. topographically altered in STM images) as a function of time in pH6 solution for different purging gases (i.e. oxygen, air, nitrogen).

336 Similar forms of reaction and altered surface layers are found at both more acidic pH (i.e. 3) and more alkaline pH (i.e. 9.5). The effect of the presence of oxygen in the solution during this reaction is also dramatic. From the successive STM images, it is possible to estimate the extent of the monolayer that has been dissolved or reacted to form pits as a function of time in the solution. Fig. 13 illustrates the kinetics of galena dissolution at pH6 for different purging gases in the solution, i.e. nitrogen, air, oxygen. It is clear that the absence of oxygen in solution, using nitrogen purging, limits the surface reaction almost to zero during the first 50min. whereas air or oxygen increase this dissolution rate by many orders of magnitude. 6. SUMMARY AND CONCLUSIONS These examples illustrate some of the chemical and structural alterations that can occur in the surface layers of oxides and sulphides during reaction in air or solution. The combination of evidence from a variety of different surface analytical and high resolution microscopy techniques can add a great deal to our understanding of the nature of these changes and the mechanisms causing them. The results described in this paper can be summarised as follows: 6.1 Ionic oxide surfaces almost totally restructure on exposure to water vapour or solution producing {100}-based rectangular pits and protusions with dimensions of a few unit cells. This r e s t r u c t u r i n g occurs before any significant dissolution takes place whilst there is a significant reduction in [H +] concentration represented by a rise in pH. Semiconducting oxides (e.g. COO), with similar initial dissolution kinetics, also show substantial surface restructuring after solution immersion. Restructured regions are, however, considerably broader and higher (i.e. 10-20nm compared with 1-5nm for MgO) and distinct {100} faceting is not found. Rounded features with less evidence of distinct low-coordination sites result from the restructuring and the dissolution rate of CoO is correspondingly two orders of magnitude lower than that of MgO. Surface restructuring of many oxides of both ionic and semiconducting type may occur immediately after immersion in bulk solution resulting in manifestly different surface structures and reaction sites from those derived from observation of the crystal surfaces before immersion. 6.2 Silicate structures can be formed in the surface layers of both bulk NiO and oxidised nickel as a result of low temperature reactions using both vapour phase deposition and plasma reactions. These silicate structures have been verified using both FTIR (for bulk NiO) and the modified Auger parameter in XPS spectra (for oxidised nickel). The layers are strongly passivating in corrosion and acid attack on these surfaces. The plasma reaction using water:tetraethoxysilane mixtures provides better control of the surface products. 6.3 Iron sulphide mineral surfaces react in solution by the loss of iron ions from the sulphide lattice to form hydroxide overlayers. Removal of these hydroxide products during acid attack on the surface reveals a sulphur-rich

337 layer in which restructuring of Fel.xS to a tetragonal Fe2S3 intermediate based on a defective pyrite lattice, is found. In this structure, the linear chains of Sn polysulphides have a S-S distance closely similar to that of elemental sulphur apparently conferring hydrophobicity on these surface layers. The initial rapid dissolution of the p y r r h o t i t e mineral is d r a m a t i c a l l y slowed by this restructuring to dissolution rates similar to that of pyrite. 6.4 Galena PbS surfaces oxidise in air at initiation sites imaged as both impurity defect sites on flat {100} surfaces and low coordination sites at step edges. In air, carbonate and hydroxycarbonate species are formed at these sites. After immersion in solution, a m e c h a n i s m in which congruent dissolution occurs creating pits of unit cell dimensions has been found in STM images. The oxidation and dissolution kinetics are dramatically affected by the presence of impurity defect sites and the presence of oxygen in solution. The mechanism of corrosion in both cases can be clearly demonstrated by combined STM/XPS studies.

7. R E F ~ C F ~ [1]

R.L. Segall, R.St.C. Smart and P.S. Turner, Oxide Surfaces in Solution in Surface and Near-Surface Chemistry of Oxide Materials. L-C Dufour and J. Nowotny (eds.), Elsevier (1988) 527-576.

[2]

P.S. Turner, C.F. Jones, S. Myhra, F.B. Neall, D.K. Pham and R.St.C. Smart, Dissolution Mechanisms of Oxides and T i t a n a t e Ceramics Electron Microscope and Surface Analytical Studies, in Surfaces and Interfaces of Ceramic Materials. L.C. Dufour et al (eds.) Kluwer Academic Publishers (1989) 663-690.

[3]

S. Myhra, R.St.C. Smart and P.S. Turner, Scanning Microscopy 2 (1988) 715-734.

[4]

D.R. Cousens, R.A. Lewis, S. Myhra, R.L. Segall, R.St.C. Smart, P.S. Turner and T.J. White. J. Amer. Ceram. Soc. 68 (1985) 64-70.

[5]

S. Myhra, R.L. Segall, R.St.C.Smart, P.S. T u r n e r and T.J. White. Intergranular Films and Pore Surfaces in Synroc C. Materials Research Society Vol. 50 (1985) 291-298.

[6]

J.L. Woolfrey, D.M. Levins, R.St.C. S m a r t and M. Stephenson. Amer. Ceramic Soc., 12 (1987) 1739-1746.

[7]

R.A. Lewis, S. Myhra, R.L. Segall, R.St.C. Smart and P.S. Turner. J. Non Cryst. Solids 53 (1982) 299-313.

[8]

R.A. Lewis and R.St.C. Smart. Phys. Chem. Glasses 24 (1983) 93-97.

Bull.

338 [9]

R.St.C. Smart. Appl. Surface Sci. 22/23 (1985) 90-99.

[10] D.K. Pham, S. Myhra, F.B. Neall, R.St.C. Smart and P.S. Turner. Amer. Ceram. Soc. (1993) in press..

J.

[11] R.St.C. Smart, Materials Technology in Surface Analysis Methods in Materials Science, J. O'Connor, B. Sexton and R.St.C. Smart (eds.) Springer-Verlag (Berlin) (1992) 319-336. [12] R.St.C.Smart. Min. Eng. 4 (1991) 891-910. [13] L. LaVanier, S. Grano and R.St.C. Smart. press.

Int. J. Min. Proc. (1993) in

[14] S. Grano, J. Ralston and R.St.C. Smart. Int. J. Min. Proc. 30 (1990) 69-97. [15] W.R. Pease, R.L. Segall, R.St.C. Smart and P.S. Turner. J. Chem. Soc. Faraday I 76 (1980) 1510-1519. [16] C.F. Jones, S. LeCount, R. St.C. Smart and T.J. White. Appl. Surf. Sci. 55 (1992) 65-85. [17] K. Laajalehto, R. St.C. Smart, J. Ralston and E. Suoninen. Sci. 64 (1993) 29-39.

Appl.Surf.

[18] V.M. Bermudez. Prog. Surf. Sci., 11 (1981) 1-64. [19] A.B. Kunz and M.P. Guse. Chem. Phys. Letts. 45 (1977) 18-21. [20] E.A. Colbourn and W.C. Mackrodt. Surf. Sci. 117 (1982) 571-580. [21] E.A. Colbourn and W.C. Mackrodt. Solid State Ionics 8 (1983) 221-231. [22] E.A. Colbourn, J. Kendrick and W.C. Mackrodt. Surf. Sci. 126 (1983) 550551. [23] C.F. Jones, R.A. Reeve, R. Rigg, R.L. Segall, R.St.C.Smart and P.S. Turner. J. Chem. Soc. Faraday 1 80 (1984) 2609-2617. [24] S. Coluccia, A.J. Tench and R.L. Segall. J Chem. Soc. Faraday I 77 (1981) 2203-2207. [25] C.F. Jones, R.L. Segall, R.St.C. Smart and P.S. Turner. A, A374 (1981) 141-153.

Proc. Roy. Soc.

[26] C.F. Jones, R.L. Segall, R.St.C. Smart and P.S. Turner. J. Materials Sci. Letters 3 (1984) 810-812. [27] C.F. Jones, R.L. Segall, R.St.C. Smart and P.S. Turner. Faraday 1 74 (1978) 1615-1623.

J. Chem. Soc.

339 [28] C.F. Jones, R.L. Segall, R.St.C. Smart and P.S. Turner. Faraday I 73 (1977) 1710-1720.

J. Chem. Soc.

[29] W.J. Tegart, The Electrolytic and Chemical Polishing of Metals. Pergamon (1956) Chapter 10, 96. [30] D. Briggs and M.P. Seah (eds) Practical Surface Analysis by Auger and X-ray Photoelectron Spectroscopy. Wiley, N.Y. (1987) 477-509. [31] C.D. Wagner, D.E. Passoja, H.F. Hillery, T.G. Kinisky, H.A. Six, W.J. Jansen and J.A. Taylor. J. Vac. SCI. Tech. 21 (1982) 933-944. [32] A.I~. Buckley, I.C. Hamilton and R. Woods, in: Proc. Int. Syrup. on Electrochemistry in Mineral and Metal Processing II, Georgia, Atlanta, 15-19 May 1988. Proc. Vol. 88-21, P.E. Richardson and R. Woods (eds.) 234246. [33] A.N. Buckley and R. Woods, Appl. Surf. SCI. 22/23 (1985) 280. [34] A.N. Buckley and R. Woods, Appl. Surf. Sci. 20 (1985) 472. [35] A.N. Buckley and R. Woods, Appl. Surf. Sci. 27 (1987) 437. [36] A.N. Buckley and R. Woods, Aust. J. Chem. 37 (1984) 2403 [37] G.W. Heyes and W.J. Trahar, in: Proc. Int. Symp. on Electrochemistry in Mineral and Metal Processing. Vol. 84-10, P.E. Richardson, S. Srinivasan and R. Woods. (electrochemical Society, Pennington, N.J. (1984) 219. [38] D. Fornasiero, M. Montalti and J. Ralston, unpublished results. [39] I.C. Hamilton and R. Woods, J. Electroanal. Chem. 118 (1981) 327. [40] A.N. Buckley, I.C. Hamilton and R. Woods. Developments in Mineral Processing Elsevier, Amsterdam (1985) 41-60.

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Science of Ceramic Interfaces II J. Nowotny (Editor) 9 1994 Elsevier Science B.V. All rights reserved.

341

Interfaces in ceramic substrates Koichi Niwa Fujitsu Laboratories LTD. 10-1 Morinosato-Wakamiya Astugi 243-01 Japan.

Abstract

Three interfaces in ceramic substrates for electronic applications are reviewed. The substrates discussed here are aluminum nitraide A1N, glass/ceramic composite and magnesia. Calcium added A1N shows very tight and highly purified grain boundaries after 2000~ firing. The lattice constant of A1N changed by diffusion of oxygen in A1N lattice. This decreases the thermal conductivity. Alumina dispersed in B203-SIO2 glass react with silica chain structure which prevent crystallization of glass/ceramic composite. This cause a steep change in thermal expansion curve. B i-Sr-Ca-Cu-O compound has been deposited on MgO single crystal. The Bi superconducting film grows with c-axis perpendicular and parallel to MgO surface.

1. INTRODUCTION Ceramic substrate is an important key material for packaging LSIs. The properties of the ceramic substrates are mainly affected by the composition of the substrate itself. For example, thermal conductivity of alumina strongly depends on the purity of A1203, the flexural strength is also affected by the purity 1). However the macroscopical interface between grains has really strong influence to these properties. The interfaces in ceramic substrates will affect enormously to its properties such as thermal conductivity or thermal expansion. In this paper, interface of the important ceramic substrates in electronic applications are introduced. They are aluminum nitride, glass/ceramics composite and magnesia substrate for superconductive ceramics.

342

2. ALUMINUM NITRIDE In LSI packaging, a high thermal conductivity ceramic material is important for dissipating heat from high power chip2). Aluminum Nitride A1N is one of candidate which has a capability to show a high thermal conductivity. Thermal conductivity of A1N is influenced by the impurities in ceramic body such as A1203 and pores. Useful sintering additives for A1N are CaO and Y203 which react with the oxide of aluminum during firing3,4). Additive CaO forms liquid phase of calcium aluminate above 1360~ So the liquid phase evaporate above about 1400~ and amount of oxygen and calcium decrease as firing temperature increases (Fig.I)5). The liquid phase accelerate densification and grain growth. 2.0 1.5 v

rG) tO 0

1.0

,.m,,

(a ~3 ". , mo. u)

0.5

rr

0

1400

Calcium 1600

1800

Oxygen 2000

2200

Sintering temperature (~

Figure 1. Change in residual oxygen and calcium in A1N as a function of firing temperature. Calcium was added as CaCO3 (3.57 wt%). Used A1N powder contains 1 wt% of oxygen impurity. After three hours firing at 1800~ the microstructure of CaO added A1N exhibits curved grain boundaries, but after 15 hours firing specimen become dense and grain boundaries show straight line. At three hours firing, Ca is observed at the grain boundaries (Fig. 2). At 15 hour firing, boundaries

343

become very thin. But secondary phase can be observed clearly. X-ray diffraction analysis suggests that the secondary phase is 5C3A(5CaO 3A1203). The specimen fired at 2000~ for 15 hours shows clean and tight boundaries (Fig. 3). Thermal conductivity, however, does not increase even though very small amount of oxygen and Ca at grain boundaries. Makihara et al.5) reported concem to this phenomena that c-axis lattice constant decreases as firing temperature increases.

Figure 2. TEM-EDX analysis for Ca added A1N fired at 1800~ 3 hours: A: Ca-Kt~ line scan for triple point. B: Ca-Kt~ composition image for grains and grain boundaries.

344

l~tm

I

I

Figure 3. TEM photograph of Ca added AIN fired at 2000~ for 15 hours.

4.981

i

!

!

!

0Raw powder

l'

i

i

Additive-CaCO3

c 9 -"

4.980

e,, o o

0

o

~

4.979-

x

6 4.978 1400

I

I

I

1600

1800

2000

__

2200

Sintering temperature (~ Figure 4. C-axis lattice constant of Ca added A1N.

345

250

2.0

200 .-> o :3

1O

c:

0 o

o3

E !-_

I-"

Jctivity

150-

er

e-

1.0 o o

100

o3 :3

0.5 "~

tD n"

5O

O" 1400

1600

1800

2000

0 2200

Sintering temperature (~

Figure 5. Change in thermal conductivity of Ca added as a function of firing temperature. As shown in Fig. 4, an average c-axis lattice constant of 1500~ firing specimen is 4.9798 A, but lattice constant decreases to 4.9786/~ at 2000~ firing. G.A. Slack6) reported that c-axis lattice constant decreases due to the oxygen concentration. The oxygen enters the A1N lattice and substitutes for the nitrogen up to very high oxygen concentrations. Oxygen concentration of 2000~ firing specimen is suggested as the level of 30x1020 cm-3 from G.A. Slack. Figure 5 shows the thermal conductivies as a function of the firing temperatures. The maximum point of thermal conductivity in this figure is obtained at 1700~ firing. So the interface between A1N grain and grain boundary will have some gradient of oxygen concentration due to the evaporation of calcium aluminate liquid phase and diffusion of oxygen to the A1N lattice.

3. LOW TEMPERATURE SINTERING CERAMICS Glass/ceramic composite shows low sintering temperature as low as about 1000~ The densification of glass/ceramic composite is promoted by the liquid phase of softened glass. Microstructure of the glass/ceramic composite is a simple mixture of glass matrix and ceramic powders. The ceramic powder disperses in the glass matrix. The roll of glass in the composite is lowering both sintering temperature and dielectric constant. The interface

346

between glass and ceramics also play an important roll that prevent the crystallization of glass matrix and cause of cracks. When borosilicate glass is applied to the composite, if the ceramic powder is cordierite, cristobalite crystallization will occur as shown in Fig. 6. Cristobalite formed in glass matrix makes the thermal expansion curve steep in the temperature range between 100-200~ due to the volume change of cristobalite caused by ~ to 13 phase transition7). However, If alumina is applied as a ceramic powder, cristobalite does not appear in glass matrix. And the borosilicate (B203-SIO2) glass-cordierite system with 5 wt% of alumina shows not steep TEC curve to compare that of the system without alumina (Fig. 7). This suggest that the addition of alumina to B 2 0 3 - S I O 2 glass can prevent the cristobalite crystallization8).

Figure 6. Cristobalite formed in glass matrix.

06 c 4b

9O. U3 C0 Q. X

i

Gloss-50%cord~ertte ~ ~ n o

I

-602 E :, I

r-

ot 0

~

w

t

%

I00 200 Temperoture (~

300

Figure 7. TEC of glass-cordierite system.

347

A strong interest is in the interface between alumina and glass matrix. Figures 8-10 show TEM-EDX results on glass rich composition (G) which was abserved cristobalite crystals in glass matrix, and alumina rich composition (A) in which glass matrix cristobalite was not observed. As shown in Fig. 8, :omposition (A) shows a small amount of alumina diffusion on every observed spot of glass matrix. Composition (G), however, does not show the alumina diffusion on every spot. From this result, it is clearly understood that diffusion of aluminum ion into B203-SIO2 glass matrix can prevent the cristobalite crystallization. 3000 Si

A =., C

=

2000

~2 m

>.

=

1000

II

0

1

1

I

I 2 3 Energy (keY) JI

i I It.L

4

Fig.8 - A 3000

3000 Si

A

Si

r

= 2000-

C

= 2000

~2

,4 L_

>,

c

1000

~" 1000

-

0 e~

II

II

0

I-.J

0

1

d

2 3 Energy (keY) Fig.8 - B2

4

~ /I

0 0

1

I

I I L

2

Energy (keY)

3

Fig.8- B1

Figure 8. TEM-EDS analysis results for two glass-alumina systems. A: Composition (A) B203-SIO2 glass- 30 wt% A1203 system, B" Composition (G) B203-SIO2 glass - 10 wt% A1203 system.

4

348

Mull~le

A

M

/M

M M

121

r" I1) E

20

0

30 20

40

(deg)

Figure 9. X-ray diffraction pattern of glass-mulite system.

i | Crlslobohle

I

cryslolhzollon

i

L)

~6

%

~t

35%gloss

i 1

x

k)

w4

- X%mulhte

-

sllico

I

t

2

0

I0 Mulhle

20 contenl

30

X (wl%)

Figure 10. Change in TEC as a function of mulite content (glassmulite + silica- glass).

In the B203-SIO2 glass-mullite system, same effect on crystallization is obtained. The minimum addition amount of mullite is about 5 wt% to prevent the crystobalite crystallization. X-ray diffraction results show the specimen of 2 wt% mullite does not prevent crystobalite crystallization. The TEC of the system 35 wt% B203-SIO2 glass-x wt% mullite-silica drastically changes at x = 2 wt%, however, at 5 wt% the TEC is as small as 2 x 1 0 - 6 / ~ which is smaller than 1/3 that of x = 2 wt%.

349

4. MAGNESIA SUBSTRATE Magnesia (MgO) single crystal is often applied as the substrate for the superconducting ceramic thin film deposition by PVD because MgO has a lattice constant of 4.2 A (a-, b-axis) which is close to that of Bi-Sr-Ca-Cu-O compound(3.8/~) and MgO does not react with Bi-compound during depositing. Bi-compound has been deposited on MgO (100) plane. Deposition has been done by RF sputtering in an atmosphere of Ar/O2 = 1. An X-ray study confirmed that deposited film was the high-Tc phase of B i superconductor. The 2000 A thick film was obtained after 9 hours deposition. The surface morphology was observed and shown in Fig. 11. Rectangular shape surface textures and granular CuO extracts are observed. The observed rectangular texture has two orientations perpendicular each other. The pattern of this texture seems to be a replica of MgO(100) lattice pattern, and the c-axis of Bi superconducting crystal is parallel to substrate surface. However it is reported that the crystal growth on MgO(100) is not epitaxial in the case of Bi- compound 9) because lattice mismatch is about 10% between MgO and Bi superconducting crystal. Usually the first monolayer is observed as Bi-oxide and on this layer Sr-oxide monolayer will be formed. Then the Bi-system crystals are formed, the c-axis is perpendicular to the MgO(100) plane(Fig.12(a)). Observed crystals of different orientation are on the same surface, which c-axis parallel to MgO(100) plane (Fig.12(b)). The rectangular pattern shown in Fig.11 suggests that the c-axis is parallel to MgO(100) plane. This result suggests that in B i-compound grows on MgO(100) substrate which has c-axis perpendicular and parallel to MgO(100) surface and orientations of parallel crystal are determined by orientation of MgO (Fig.12(c)). 5. CONCLUSIONS Interfaces in ceramic substrates have been reviewed. The thermal conductivity of A1N affected strongly by the residual oxygen at grain boundaries. However even though grain boundaries highly purified, if oxygen diffuses and enters into A1N lattice, the thermal conductivity will become low. B203-SIO2 glass/ceramic composite shows a low thermal expansion. But the thermal expansion increases if crystobalite crystallization occurred. Diffusion of aluminum ion to the glass matrix can prevent the crystallization. Bi-compound superconductive thin ceramics, deposited on MgO single crystals, show a rectangular shape surface. And Bi-compound grows in both orientation of perpendicular and parallel to MgO (100). The lattice parameter of the B i-compound is about 10% of mismatch with MgO (100).

350

Figure 12. Crystal growth on MgO substrate. (a) shows c-axis orientation is perpendicular to substrate surface. (b) shows c-axis orientation is parallel to substrate surface. (c) is a schematic texture of observed surface.

351

6. R E F E R E N C E S

1 2 3 4 5 6 7 8 9

C.A. Bruch, Amer. Ceram. Soc. Bull. 41 (1962) 799. W. Werdecker and F. Aldinger, Proc. 34th ECC (1984) 402. N. Kumamoto and H. Taniguchi, J. Mat. Sci. Let. 3 (1984) 471. K. Koyama, A. Tsuge, H. Inoue and H. Ohota, J. Mat. Sci. Let. 1 (1982) 325. H. Makihara and N. Kamehara, Proc. Mat. Res. Symp. 245 (1992) 437. G.A. Slack, J. Phys. Chem. Solids 34 (1973) 321. K. Niwa, Y. Imanaka, N. Kamehara and S. Aoki, Advanced in Ceramics 26 (1989) 323. Y. Imanaka, S. Aoki, N. Kamehara and K. Niwa, Yogeyo-Kyokai-shi 95 (1987)1119. M. Kawai, Appl. Phys. in Japan, 60 (1991) 470.

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Science of Ceramic Interfaces II J. Nowotny (Editor) 9 1994 Elsevier Science B.V. All rights reserved.

353

Diffusion-induced grain boundary phenomena in metals and oxide ceramics E. Rabkin, C.Y. Ma and W. Gust Max-Planck-Institut ftir Metallforschung and Institut ftir Metallkunde Seestr. 92, D-70174 Stuttgart, Germany

Abstract In the first part of the paper recent experimental results on diffusion-induced grain boundary migration in ceramics are reviewed. It is demonstrated that all types of diffusiondriven phenomena known from the study of metallic systems occur also in oxide ceramics. Amorphous films on grain boundaries in ceramics play an important role in migration processes. It is shown that none of the existing theories of diffusion-induced grain boundary migration can describe satisfactorily the whole set of experimental data. In the second part of the paper, a new approach to the phenomenon is proposed, based on the gradient term in the expression for the free energy of binary alloys. The developed theory describes satisfactorily the stationary grain boundary migration and initial stages of diffusion-induced grain boundary migration. A new vacancy theory of the phenomenon is also reviewed.

1. INTRODUCTION About 20 years have gone after publishing the pioneering work of den Broeder [ 1] where the diffusion-induced grain boundary migration (DIGM) has been revealed. The phenomenon can be described as the motion of a grain boundary (GB) perpendicular to its plane occurring when solute atoms diffuse along it (see Fig. l a). The peak of the interest to the phenomenon was in the end of 70ths and the beginning of 80ths. At that time DIGM was experimentally observed in more than 30 metallic systems and theoretical explanations were suggested, which are widely accepted till now. The main results of investigations of DIGM at that time are collected in excellent review papers of King [2] and Handwerker [3]. From the end of 80ths the interest of research to DIGM and related phenomena has been decreased. We think that it is due to the lack of new ideas concerning the atomic mechanism of the phenomenon. Any new experimental result till now has been discussed in terms of advantages or disadvantages of the coherence strain and dislocation climb models. We will consider in detail in the present paper some recent new ideas: Fournelle's vacancy diffusion model of DIGM [4] and a model based on the gradient term in the expression for the free energy of an alloy [5]. In the fh:st part of the paper, a brief review of DIGM and related phenomena in ceramics will be presented, because review papers [2,3] are mainly referred to the metallic systems.

354 2. GENERAL DEFINITIONS Let us introduce some definitions. In the microstructure of multiphase materials (as most ceramics are) layers of the phase with the lower melting point often separate the grains with the higher melting point. At the temperatures above the melting point of the intergranular phase these layers are liquid. In many cases the diffusion of solute atoms along such liquid layers leads to its migration and produces microstructures resembling those at DIGM (see Fig. l b). Such phenomenon is known as liquid-film migration (LFM). LFM in metallic systems is often observed during liquid-phase assisted sintering of refractory metals [6,7]. LFM occurs also if the intensive thermodynamic parameters of the system are rapidly changed. For example, if the temperature of the equilibrated two-phase system with liquid films is increased rapidly, then liquid films will migrate leaving behind them the solid solution of the new equilibrium concentration. Another phenomenon closely connected with DIGM is diffusioninduced recrystalization (DIR). During alloying experiments, when solute atoms are supplied from the surface of the sample, fine grains are often nucleated having high solute concentrations near the surface [8]. Below that DIR zone the microstructure is often formed by the DIGM process (Fig. 1c). Figure 2 shows the typical morphologies of DIGM (a), LFM (b), and DIR (c,d) formed in the Ni(Cu) system.

Figure 1. Scheme showing possible types of diffusion-induced phenomena in solids. (a) Diffusion of solute atoms from the surface of the sample leads to the GB migration. Migrating GB leaves behind it the alloy enriched by the solute (DIGM). (b) Changing of the temperature of the two-phase system with liquid films leads to their migration. Migrating films leave behind them the alloy with the new equilibrium concentration of the solute (LFM). (c) The layer of new fine grains is nucleating at the surface of the sample, from which solute atoms are supplied. -The original positions of interfaces are shown by the broken lines and the directions of the GB migration by the thin arrows.

355

Figure 2. Light microscopy (LM) and scanning electron microscopy (SEM) images of microstructures developing in the Ni(Cu) system after the diffusional anneal. (a,c) Crosssections; (b,d) parallel sections.

356 3. DIFFUSION-DRIVEN INTERFACE PHENOMENA IN NONMETALS 3.1. DIGM in ZrO 2 -based ceramics The microstructure evolution in ZrO 2 - based ceramics partially stabilized by 3.2 wt.% CaO + 1.1 wt.% MgO has been studied after annealing at 1400 ~ [9]. Such a composition is in the two-phase field [cubic phase (c) + tetragonal phase (t)] of the ternary phase diagram. Transmission electron microscopy (TEM) of samples showed that most GBs are covered by the 5 to 30 nm thick glassy phase. Such intergranular GB phases are often observed in ceramic systems [10]. After quite a long anneal the composition of that intergranular phase was 27 wt.% CaO + 17 wt.% MgO + 56 wt.% SiO2 . It forms at the manufacturing process due to the interaction of SiO2 impurities with the main components. TEM observations showed also fine (about 10 nm) intragranular t-ZrO 2 precipitates in the c-ZrO 2 matrix grains. Near the GBs wetted by the glassy phase precipitate-free zones of maximal 10 Bm width were observed. The concentration of CaO in these zones was about 0.7 wt.% higher than in the two-phase region (c+t). The following mechanism of the formation of such zones was proposed: The migrating

Figure 3. Evolution of the microstructure of Z r O 2 - based ceramics during the DIGM process. (a) The GB wetted by the liquid phase (marked by a solid line) dissolves during the migration process the fine t-ZrO 2 precipitates. The excess of CaO in the swept zone is provided from the liquid layer on the GB [9]. (b) The GB migrates simultaneously with the growth of the t-ZrO 2 lens. The excess of Y203 in the swept zone is provided from the growing lens [ 12]. The GB migration direction is shown by the arrow.

357 GB wetted by the liquid film dissolves the two-phase grain decreasing the coherence strain and interface energy connected with the t-precipitates. Single phase CaO enriched c-solid solution is then forming at the other grain. The liquid film along the GB provides CaO needed for this process and, hence, decreases its width during the migration process (see Fig. 3a). Such a process has common features with DIGM, LFM and discontinuous dissolution [11]. Analogously to the discontinuous dissolution process, the two-phase structure disappears by the GB migration process, but usually the two-phase structure of such a phenomenon is lamellar with an interlamellar spacing of about 10 ~m, which is much higher than the characteristic size of rounded t-ZrO 2 precipitates observed [9]. The similarity with the LFM process is obvious, because the SiO2 -rich liquid phase on the GB is involved in the process. However, such a liquid layer only few nm wide is under strong influence of the adjacent crystalline grains and cannot be treated as an equilibrium liquid (the term quasi-liquid is often used for such situations [10]). In [12] another zirconia-stabilized ceramic of ZrO 2 - 8 wt.% Y203 composition was investigated. In the as-prepared ceramic the c-ZrO 2 grains were found to have the tweed microstructure arising from fine-scale homogeneous precipitation of t-ZrO 2. Again, along the GBs wetted by the amorphous phase precipitation-free zones were observed after annealing in the temperature range from 700 to 1400~ Lenses of t-ZrO 2 were also observed at c/c GBs and c/t interphase boundaries, and one t-ZrO 2 lens together with one precipitation-free zone occurred in approximately all c-ZrO 2 grains. The composition analysis showed that the content of Y203 in the t-ZrO 2 lens is lower than in the surrounding tweed c-ZrO 2 grains, and the content of Y203 in the precipitation-free zones is higher than in the surrounding tweed c-ZrO 2 grains. So Y203 necessary for the formation of precipitation-free zones during GB migration is supplying not from the GB amorphous film as in the case of Ca-stabilized ZrO 2 [9], but from the simultaneously growing t-ZrO 2 lens (see Fig. 3b). TEM investigations of the GB behaviour were performed on fully tetragonal ceramic of composition 90 mol% ZrO 2 +10 mol% CeO 2 [13]. In the as-prepared material an amorphous GB phase about 1 nm thick was found to wet many GBs and to form pockets at the triple junctions. In situ observations of samples exposed to extensive electron irradiation during prolonged TEM experiments showed the GB migration. Areas where GBs migrated were free from the intergranular phase, and sometimes the dislocation structures were visible on wavy GBs in their final positions. Composition measurements showed that the Ce content in areas swept by the migrating GBs is lower than in the surrounding material. The following mechanism of DIGM was proposed: Electron irradiation partially reduces Ce 4+ ions to Ce 3+ which is accompanied by an increase of about 10% of the corresponding ionic radius. High elastic strains in the lattice provide the driving force for the GB migration. Ce 3§ with the higher ionic radius diffuses along the quasi-liquid GB layer to the surface of the sample, and the growing grain with the lower Ce content is unstrained. Simultaneous diffusion of Si, A1 and other impurities from the quasi-liquid GB layer decrease the amount of the liquid phase, and when it disappears completely, the GB stops and the lattice mismatch between the growing and shrinking grains is accommodated by the GB dislocations. The GB can also stop if the driving force is equilibrated with the GB curvature. Using the expression for the elastic coherence strain energy Eel = Y(nl )rl 2(Ac) 2,

(1)

358 where Y(nl) is an appropriate, orientation dependent elastic modulus of the shrinking grain, rl is the atomic misfit, and Ac is the difference in composition between the shrinking and growing grains, and using the condition of the balance between elastic and surface energies Eel =2y/re,

(2)

where y is the GB surface tension and r c the critical GB curvature at which it stops, one can obtain the relationship between r c and Ac"

(Ac)2- 1/r~.

(3)

Indeed, such a relationship was observed experimentally for curvatures higher than 25 nm. DIGM and DIR phenomena were also observed in (Pbo.91Lao.09)(Zr0.91Tio.09)0.9803 ceramic during the anneal in PbO atmosphere and in air [ 14]. In the former case, GBs migrated due to the PbO diffusion into the specimen along GBs and, in the latter case, due to the PbO evaporation (dealloying).

3.2. Observations in other ceramic materials DIGM was observed in calcite (CaCO3) in contact with the SrCO 3 melt [15]. 40 different GBs in bicrystals with different crystallographic parameters were studied. The main features of the process can be summarized as follows: The DIGM process has a reversible character; that means that some GBs change the direction of migration. The Sr concentration in doubly swept zones is higher than that in singly swept zones. The portion of migrating GBs increases with the increase of the misorientation angle 0, and only a little portion of the smallangle GBs migrated. That can be connected with the fact that DIGM is a two-stage process of nucleation and growth, and the probability of the formation of the nuclei of critical size is directly proportional to the GB energy. DIGM and DIR were reported also in AI20 3 at either alloying and dealloyng with Fe20 3 [ 16]. GBs which have migrated often were of faceted shape. It is generally accepted that such faceting proves the coherence strain mechanism of DIGM, because the elastic modulus Y(n 1) determining according Eq.(1) the coherence strain energy is anisotropic and GB facets should minimize that energy. Problems connected with the different valency of matrix and diffusing atoms were avoided during the study of DIGM in SrTiO 3. Ba or Ca ions diffused into the SrTiO 3 matrix during annealing in BaTiO 3 and CaO powders and, as far as valency state of Sr, Ba and Ca are identical, the electric potential gradient in the diffusion zone was absent [17]. All classical features of DIGM microstructures were observed after anneals at 1200 and 1400 ~ There is some controversy about DIGM and DIR in the NiO(O) system. It was shown by Parthasarathy and Shewmon [ 18] that a layer of fine grains appears on the surface of the NiO foil annealed together with the decarburized Ni foil. The phenomenon was identified as DIR caused by the oxygen diffusion and partial NiO reduction within its homogeneity range. However, later it was shown that under the same experimental conditions fine new grains nucleating at the NiO surface are oxygen free and the phenomenon should be characterized as a reduction of oxide, but not as DIR [19]. LFM and DIGM were observed in the two-phase TiN-Ni material with additions of TiC [20]. It was found that the migration rate of liquid films increases parabolically with the

359 observed lattice parameter difference between the TiN and Ti(NC) solid, which is consistent with the coherence strain model.

4. W H A T IS STILL UNCLEAR IN DIGM? For the interpretation of the experimental results two models are widely employed: the coherence strain model and the dislocation climb model. According to the coherence strain model the solubility of the solute atoms in the GB phase is different on the sides of the shrinking and growing grains because of the elastic energy Eel [see Eq.(1)] accumulated in the volume diffusion zone ahead of the migrating GB. The growing grain is assumed to be unstrained at such a consideration. The difference of solubilities leads to the concentration gradient and atom flow across the GB and its migration. That model allows one to understand the following features of DIGM: 1. There is a correlation between the atomic misfit parameter rl and the rate of DIGM. 2. The faceted shape of GBs sometimes observed at their final positions is connected with the anisotropy of the elastic modulus Y(n 1). 3. The change in the direction of DIGM. If the GB is pinned at two points, it will increase its curvature during the migration process. That decreases the velocity of the GB and increases the size of the diffusion zone ahead of the GB. At some moment the coherency in that diffusion zone is breaking, which decreases the coherence strain energy. After that the GB begins to migrate under the influence of its curvature in the opposite direction.

Figure 4. Parallel sectional view 10-15 ~tm below the surface of a specimen annealed at 1073 K for 2 days showing DIGM and DIR during Cu diffusion into Ni-19.4 at.% Pd [23].

360 However, there are some serious limitations which do not allow one to accept that model as applicable to all cases when DIGM occurs: 1. There are observations of DIGM in systems where the misfit parameter is negligible, such as in the Au(Ag) [21] and Pt(Pd) [22] systems. Recently, a special test for the validity of the coherence strain model has been performed [23]. DIGM and DIR were observed during Cu diffusion in a polycrystalline Ni-19.4 at.% Pd alloy in the temperature interval where considerable volume diffusion ahead of the GB was expected (Fig. 4). Ni-19.4 at.% Pd has the same lattice parameter as pure Cu, so q--0 and no coherence strain should develop in the diffusion zone. Consequently, at least in this case some other mechanism should be involved in the DIGM process. 2. The coherence strain model imphes the complete incoherence across the migrating GB. It is certainly not true in the case of small-angle GBs and special GBs with a low value of the inverse coincidence cite density E. But only such GBs should exist in two-phase systems, where LFM and DIGM were observed simultaneously [6,7,20], because all random GBs with a high energy in such systems are wetted by the hquid phase. It should be noted that in ZrO 2 -based ceramics DIGM was observed only at the GBs wetted by the layer of the amorphous phase[9,12,13], so the coherence strain model can be apphcable to those systems. If we imply the complete coherence of two grains at the GB, some driving force, however, still will exist due to the anisotropy of the Y(n 1) modulus. It vanishes only in the case of symmetrical GBs. Therefore, the numerous DIGM observations on symmetrical GBs are speaking against the coherence strain model. 3. In the case of symmetrical GBs the model does not explain the nucleation of DIGM, because the coherently strained layers on both sides of the GB make it more stable against small perturbations of its shape. In the dislocation climb model there are also some unclear moments: the model does not predict the migration of a single GB in two different directions simultaneously, the migration direction reversal and the occurrence of DIGM in interstitial soLd solutions. The model is certainly not true in the case of LFM, because hquid films have not any internal dislocation structure.

5. MODEL BASED ON THE GRADIENT TERM IN THE EXPRESSION FOR THE FREE ENERGY OF AN ALLOY We will start from the following expression for the free energy of one mol of an inhomogeneous alloy in the framework of the theory of regular solutions: G = Go +RT[clnc +(1-c)ln(1-c)] +Dx:(1-c)

+~(~C/c)X)2,

(4)

where c is the concentration of the solute and the solute distribution is inhomogeneous along the x direction, f~ is an interaction parameter and Go depends linearly on the concentration. The last term in Eq.(4) arises, like f~, due to the interaction between atoms. That term determines, for example, the kinetics of spinodal decomposition at the early stages, when the concentration gradients are small [24]. Its nature can be understood using a simple graphic construction for the free energy (Fig.5). On the c axis the concentration change Ac on the distance ~ is shown, where ~ is a characteristic distance of atomic interaction. In that

361

G

| |

|

c0 Figure 5. Schematic dependence of the free energy G of a binary alloy on the concentration c, showing the nature of the gradient term. Ac is the concentration change on the distance ~ of atomic interaction. The solid segment shows the upper limit of the gradient term contribution to the free energy.

concentration interval atoms of the alloy are "feeling" each other. The upper limit of the energy increase due to the inhomogeneity can be estimated under the assumption that the alloy of concentration c o is composed from the fine mixture of alloys of compositions bounding the interval Ac. The increase in the concentration gradient leads to the increase of Ac and corresponding increase in the inhomogeneous part of the free energy. The coefficient ~ can be estimated as RTc~ 2 [24], or can be considered as a phenomenological parameter of the theory. Using Eq.(4) one can obtain for the part of the chemical potential that depends on concentration RT[ln(1-c)+2x 4 c2+x-1 ~2{ (~C/~X)2 _2C(~2C/~X2)}]

(5a)

Ps = RT[lnc+21:-1 (1-c)2+I: -1 11/2{(OC/~X)2 +2( 1-c)(/)2c//)x2) }],

(Sb)

~1,m -~

and

where Its and ~1.m are the chemical potentials of the solute and matrix atoms, respectively; x = T/Tc, where Tc is the critical temperature of solid solution decomposition.

5.1. Stationary GB motion We assume the GB to be a plane layer of a homogeneous phase of thickness ~i. The GB moves with a constant velocity v and leaves behind it an alloying zone with the constant solute concentration Co (Fig. 6). We further assume stationary solute distribution ahead of the GB, with the solute concentration Cl in the bulk atomic layers neighbouring the GB. This means that

362

Figure 6. A GB moving with the velocity v. The stationary distribution of the solute ahead of the GB is given by Eq. (7). The driving force for the migration process arises due to the gradient term in Eq. (4), in spite of the small amplitude of the difference co - Cl.

all atoms entering the GB due to its movement should be transmitted to the growing grain by the cross flows, and stationary solute distribution is then reconstructed by the atomic fluxes along the GB. Fluxes along and across the GB are assumed to be independent. Using Einstein's formula for the atom mobility, that condition can be written as follows:

Dbm (1 --C b ) ( ~ m l -- ~l'mO)rib

RT8

= v(1-cl)nv

(6a)

and Db, C b (~-I'sl -- l~l.sO)rib = vclnv, RT8

where

Dbmand

(6b)

D~ denote the coefficients of matrix and solute atom self-diffusion across the

GB, respectively; c b is the solute concentration in the GB; ~l.mi and ktm0 are the chemical potentials of the matrix atoms on the right and left surfaces of the GB layer, respectively; the same quantities for the solute atoms differ by the "s" subscript only; n b and nv are the numbers of atoms in the unit volume of the GB and bulk material, respectively. Because the 8 value is small, the derivatives of chemical potentials have been replaced by the differences. For further consideration we will assume that the solute distribution in the alloying zone is homogeneous, and the solute distribution ahead of the migrating GB is given by the stationary solution of the diffusion equation for a plane moving with the constant velocity v and having the constant concentration Cl:

363 V

C(X) = c l e x p ( - ~ x ) ,

(7)

where D is the bulk interdiffusion coefficient. Strictly speaking, D should differ from those values determined by standard methods of bulk interdiffusion studies due to the importance of the gradient term [the last term on the right side of Eq.(4)] at short distances. It is clear that the contribution of this term to the diffusion equation should be of the order of magnitude ~2(~4c/~x4), or, using Eq.(7), ~2(v/D)2(O2c//)x2). The coefficient of the second order derivative can be written as (~/d) 2, where d is the characteristic width of the bulk diffusion zone ahead of the GB. We will see later that in our case (~/d)<
Dbs [c~ - 4'~(1- c~ - C0)Cb / D28nv'l; 1 "] )-1 ](I_Cb)C ~ V Dbsllt2Cbnb Dbm[(l_c 0 --4XC0 =

(8)

For estimations we simplify the expression assuming Co= Cb= 0.5, D~= Dbm= D b and ~ = nb. Then we will get from Eq.(8)

4D2~Sx v = DbXi/2 .

(9)

This equation connects the velocity of the stationary motion of a GB with the bulk interdiffusion coefficient D, the boundary width ~i, the coefficient Db of diffusion across the GB, the relative temperature x and the average distance ~ of atomic interaction. It is interesting to note that according to Eq.(9) the GB velocity is proportional to the boundary width 8, while in theories where simply the chemical driving force is involved that dependence is inverse (see, for example, [25]). That means that the theory discussed can be useful in the understanding of the DIGM phenomena in ceramics, where GBs are wetted by the glassy phase. The theory is also completely compatible with the phenomenological description of DIGM proposed by Brechet and Purdy [26]. They introduced the phenomenological parameter 1 which can be treated as a thickness of the layer of material on both sides of the GB influencing the motion of the GB. In our consideration, the parameter l has a clear physical meaning: it denotes the average distance of atomic interaction and is analogous to our parameter ~. The free energy graphic construction allowing one to calculate the driving force for the DIGM process has been proposed [26]. The driving force AG' can be presented as the shaded on Fig.7 area divided by the difference c 1 - Cl', where c 1' is the concentration of the solute on the distance l ahead of the moving GB. Expanding the free energy at the concentration c o to the second order one can derive the following expression for the driving force:

364

Figure 7. The thermodynamic driving force for DIGM can be represented as the averaged over the concentration interval c I - c 1' difference between the free energy of the alloy G(c) and the values of the tangent to the G(c) function at c o [26].

Ao'=

7t

0c'

i[c x _c0l x

)C- o o

(10)

Let us perform some numerical estimations on the basis of Eq.(9) for the case of DIGM in a Cu polycrystal during Ni diffusion [27]. It has been shown that at a temperature of 900~ the displacement of the GB is approximately 10 I.tm after an 1 hour anneal and is not followed by a noticeable concentration discontinuity across the boundary. For an estimation on the basis of Eq.(9) we assume B to be of the same order of magnitude as ~ and approximately equal to 1 nm (about 3 interatomic distances being a normal value for the GB width [11]). The interdiffusion coefficient at 900~ for Cu-50 at.% Ni is about 10-15 m2/s [28], and z=2 [29]. Then, from Eq.(8), Db/D=103. This seems to be a reasonable value, because for this homological temperature (0.75Tin, where Tm is the melting temperature) the characteristic value of the ratio of the coefficients for GB and bulk diffusion is about 104 [ 11 ]. For Db/D= 103 the average width, d, of the bulk diffusion zone ahead of the boundary should be d=D/v=0.5 gm. This justifies the approximations made in obtaining Eq.(8), because the ratio ~s/d is small.

365

5.2. Initial stages of DIGM We consider the stability of the GB plane against small perturbations of spherical shape. The behaviour of such a perturbation is dependent on the difference of the chemical potentials of the solute atoms on both sides of the GB. If the chemical potential of the solute atoms is lower on the internal surface of the sphere (see Fig.8), solute atoms diffusing along the GB will be deposited on that surface and the perturbation will grow. We will show below that such an instability really develops due to the asymmetry of the solution of the diffusion equation for a sphere. Due to such an asymmetry, even in the case when the volume concentration of solute on both sides of the spherical segment is the same (Cl), the gradient term in Eq.(4) gives rise to the chemical potential difference which overcomes the capillary force. In further consideration we will assume the size of bulk diffusion zone to be much smaller than the radius of curvature of half-spherical nuclei a, so standard solutions of the diffusion equation for a sphere [30] are applicable:

. . . . . .

a

Growing s/Shrinking grai3n [__grain o

I

cD r

I

t

o

3 .... Distance

Figure 8. (a) Spherical bulge of radius a on the GB. (b) The distribution of the solute diffusing from the GB to the bulk. At the straight parts of the GB (1,3) distributions are symmetrical, while at the bulge (2) the distribution is asymmetric. The gradient of solute distribution is higher on the side of the shrinking grain. Therefore, solute atoms diffusing along the GB should be deposited on the growing grain, because the contribution of the inhomogeneity to the chemical potential is lower there.

366

C1

2~-t

a/

(11)

for the diffusion outside the sphere, and

2a ~ (-1)" nxr = n sin

---~-exp[( n Z x 2a2D t )

&Cl = 1 + ~

(12)

for the diffusion inside the sphere, where t is the annealing time and r is the distance from the center of the curvature. Combining Eqs.(ll) and (12) with Eq. (5b), where x should be substituted on r (due to the spherical symmetry gradc---Oc/Or) one can get

~t~l = RTcq/2Cl

(1 1)E/1 1) 4,1Cl,] +~

Cl

+~

+ ~ a

(13)

for the arising from the gradient term part of the chemical potential ~l.sl of the solute atoms on the external side of the GB, and

[52Cl - 2(1 - c 1)S], gt~0 = 4RTcv2cl a2

(14)

where

: exo(

(15)

for the corresponding part of the chemical potential gts0 of the solute atoms on the internal surface of a half-sphere nuclei. As in the case of stationary GB motion we will assume Cl= 0.5 and the average size of bulk diffusion zone x/-~=0.01 a. Then, for the difference A~t = ~Lsl -~l.s0 we will get: A~t - 170

RTc~ 2 a2 .

(16)

Comparing the chemical potential difference given by Eq.(16) with the corresponding difference caused by the GB curvature, one can conclude that for the nuclei sizes

367 a < 170 RTc~t2 , 2Ttos

(17)

where y is the GB surface tension and o)s is the molar volume of solute atoms in the GB, the spherical perturbation of the GB shape will grow. For further numerical estimations we assume Tc=1000 K, y --- 0.5 J/m 2, O)s=10-5 m3/mol, and V --1 nm. Then, from Eq.(16), a < 0.14 ktm. That is characteristic size of a bulge from which the DIGM process develops, known from the experiment [31]. It should be noted, however, that the characteristic size of the bulk diffusion zone at such an estimation (0.01a) appears to be of the same order of magnitude as ~. That means that the gradient term should be taken into account in the bulk diffusion equations as well, and solutions (11) and (12) are, strictly speaking, not valid. So at chosen values of the parameters, the results obtained are rather qualitative than quantitative. Consider the physical sense of the obtained instability. Solute atoms diffusing along the GB have two possibilities to be accepted in the lattice structure of the adjacent grains. If the binary system exhibits an inmiscibility gap, then there should be an effective attraction between the solute atoms. But from the internal side of the bulge the total amount of solute atoms in the layer of thickness V is more than from the external side of the bulge. So there should be a net attractive force acting on the diffusing atoms and directed towards the grain from the internal side of the bulge. The deposition of solute atoms on that grain leads to its growth and development of the bulge. Bulges can also nucleate on moving GB. It is clear that the probability of the formation of a bulge opposite to the GB velocity direction is higher for slower boundaries. So, for GBs being slow enough the migration direction reversal can occur. It is consistent with the in situ observations of DIGM in calcite [15]: it was found that the migration direction reversal is sometimes preceded by the GB migration slowing.

6. FOURNELLE'S VACANCY MECHANISM OF DIGM Foumelle has proposed recently the new mechanism for DIGM based on the dependence of the equilibrium vacancy concentration in the alloy on its composition [4]. If the composition discontinuity exists across the GB, then the corresponding discontinuity in the vacancy concentration should lead to the flow of vacancies across the GB and its migration. If one assumes that the GB segment has a constant thickness 8 and a constant composition, then the segment should always contain a fixed number of atomic sites. Thus, whenever an atom jumps into a vacancy from an adjacent grain, another atom should jump from the boundary into an adjacent grain. Thus, the net flux of vacancies across the GB determines the corresponding net flux of atoms and the GB migration. The GB velocity can then be written as follows: b v = Jvo) ,

where

(18)

jb is the flux of vacancies across the boundary and co is the atomic volume. According

to Fick's first law jb is given by

368 0 m

Jbv = - D b cv

1

cv ~5

(19) '

where D bv is the diffusivity of vacancies in the boundary, and c~ and c 1v are the equilibrium vacancy concentrations in the shrinking and growing grain, respectively. Under some further simplifying assumptions one can get the following expression for the velocity:

v :-g- expt-TJ expt--ffj[exp{-

RT

'

(20)

where AS~ and AH~ are the activation entropy and enthalpy, respectively, for the formation of vacancies in the shrinking grain, AG~ is the difference between the free energies of formation of vacancies in the shrinking and growing grain. It was found that Eq. (20) describes adequately the experimental results for DIGM in the Cu(Zn) system.

7. CONCLUSIONS The following conclusions can be drawn. 1. The principal features of DIGM, LFM and DIR in ceramic systems are the same as in metals. 2. In ZrO 2 - based ceramics precipitation-free zones are forming around GBs migrating during DIGM. Migrating GBs dissolve small tetragonal particles formed during homogeneous nucleation process. Quasi-liquid films at the GBs play an important role for DIGM. During the migration process the thickness of the GB film decreases, and sometimes a GB dislocation structure is forming at the final position of the GB. 3. A new model of DIGM is proposed. The model is based on the gradient term in the expression for the free energy of a binary alloy and requires some volume diffusion from the GB into the bulk. The predictions of the model are in good agreement with the experiment. 4. It is shown that the vacancy model of DIGM allows one to express the migration rate of the GB through the thermodynamic properties of the binary system. The model is applicable at low temperatures, when the concentration discontinuity exists across the GB. 5. There is no universal mechanism of DIGM, which describes all possible experimental situations. Mechanism maps for DIGM and DIR proposed [3] can be very useful in the analysis of particular experimental situations.

8. ACKNOWLEDGEMENTS E.I.R. wish to thank the Alexander-von-Humboldt Foundation for the support of his work at the Max-Planck Institute in Stuttgart. Prof. L.S. Shvindlerman, Prof. R.A. Fournelle, Dr B.B. Straumal and Dr F. Bartels are greatly acknowledged for helpful discussions.

369 9. REFERENCES

1. J.A. den Broeder, Acta metall., 20 (1972) 319. 2. A.H. King, Int. Mater. Rev., 32 (1987) 173. 3. C.A. Handwerker, in: Diffusion Phenomena in Thin Films and Microelectronic Materials, ed. D.Gupta and P. Ho, Noyes, Park Ridge, New Jersey (1988) 245. 4. R.A. Foumelle, Mater. Sci. Forum, in press. 5. E.I. Rabkin, L.S. Shvindlerman and W.Gust, submitted to Interface Science. 6. Y.D. Song, S.T. Ahn and D.N. Yoon, Acta metall., 33 (1985) 1907. 7. H.K. Kang, S. Hackney and D.N. Yoon, Acta metall., 36 (1988) 695. 8. F.J.A. den Broeder, Thin Solid Films, 124 (1985) 135. 9. E.P. Butler and A.H. Heuer, J. Am. Ceram. Soc., 68 (1985) 197. 10. E.I. Rabkin and W. Gust, this volume. 11. I. Kaur and W. Gust, Fundamentals of Grain and Interphase Boundary Diffusion, Ziegler Press, Stuttgart (1989). 12. R. Chaim, A.H. Heuer and D.G. Brandon, J. Am. Ceram. Soc., 69 (1986) 243. 13. H.K. Schmid, J. Am. Ceram. Soc., 74 (1991) 387. 14. J.-J. Kim, B.-M. Song, D.-Y. Kim and D.N. Yoon, Ceram. Bull., 65 (1986) 1390. 15. R.S. Hay and B. Evans, Acta metall., 35 (1987) 2049. 16. H.Y. Lee and S.-J. L. Kang, Acta metall, mater., 38 (1990) 1307. 17. K.J. Yoon, D.N. Yoon and S.-J. L. Kang, Ceramics Int., 16 (1990) 151. 18. T.A. Parthasarathy and P.G. Shewmon, Acta metall., 32 (1984) 29. 19. B. Bergman, N. Lange and J./~gren, Acta metall., 36 (1988) 883. 20. Y.-J. Baik and K.Y. Eun, J. Am. Ceram. Soc., 74 (1991) 1397. 21. D.B. Butrymowicz, D.E. Newbury, D. TurnbuU and J.W. Cahn, Scripta metall., 18 (1982) 277. 22. F.J.A. den Broeder, M. Klerk, J.M. Vandenberg and R.A. Hamm, Acta metall., 31 (1983) 285. 23. C.Y. Ma, W. Laub, R.A. Foumelle, W. Gust and B. Predel, Z. Metallk., 83 (1992) 633. 24. J.W. Cahn, Trans. AIME, 242 (1968) 166. 25. J.E. Krzanowski and S.M. Alien, Acta metall., 31 (1983) 213. 26. Y.J.M. Brechet and G.R. Purdy, Acta metall., 37 (1989) 2253. 27. F.J.A. den Broeder and S. Nakahara, Scripta metall., 17 (1983) 399. 28. H. Mehrer (ed.), Diffusion in Solid Metals and Alloys, Landolt-B0rnstein New Series, Vol. 11]/26, Springer-Verlag, Berlin (1990) 303. 29. T.B. Massalski et al. (eds.), Binary Alloy Phase Diagrams, ASM International, Materials Park, Ohio (1990) 1444. 30. J. Crank, The Mathematics of Diffusion, Clarendon Press, Oxford (1964). 31. R.W. Balluffi and J.W. Cahn, Acta metall., 29 (1981) 493.

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Science of Ceramic Interfaces II J. Nowotny (Editor) 9 1994 Elsevier Science B.V. All rights reserved.

371

Thin films on grain boundaries in metals and ceramics and their importance for the properties of the materials E. Rabkin Max-Planck-Institut for Metallforschung and Institut far Metallkunde Seestr. 92, D-70174, Stuttgart, Germany

Abstract Thermodynamic and kinetic conditions, under which grain boundaries can be described as thick, are considered. It is shown that in many cases such thick grain boundaries consist of thin films of a phase different from that in the adjacent grains. Such films can be liquid, amorphous or can have a crystalline structure different from that in adjacent grains. A variety of materials are considered, from pure metals to high-temperature superconducting ceramics. Phase diagrams for grain boundaries are constructed. The lines on such diagrams are constructed from the points of phase transitions on grain boundaries, when the layer of another phase appears on them. In complete analogy with the bulk phase diagrams such lines can be treated as grain boundary solidus or grain boundary solvus. The influence of the phase state of grain boundaries on the properties of polycrystalline materials (such as mechanical properties and grain growth behaviour) is considered.

1. INTRODUCTION Both computer simulation and direct observation of the grain boundaries (GBs) in solids show that generally the GB width does not exceed 2-3 interatomic distances. The dependence of crystal physical properties as a function of the distance across the GB is smooth, exhibiting an extreme at the GB location, so the GB cannot be treated as a layer of special homogeneous GB phase, as is often supposed by the different simplified models of GB phenomena, for example, in Fisher's model for GB diffusion [1]. However, it was shown recently that in some circumstances GBs may be considered as quite wide (about 10 interatomic distances or more) layers of a phase, which is different from the bulk phase (for example, liquid or amorphous) and is bounded by two sharp, well-defined "crystal phase-GB phase" interfaces. Experimental indications of the existence of such GB layers were obtained, for example, in the investigation of Fe-Si-Zn ternary alloys, in a variety of ceramics, including high-temperature superconductors, in Au-Cu and Co-Pt alloys near the ordering temperatures and in incoherent twin boundaries in Cu and Ag. There is also indirect data which can be interpreted in terms of the thick grain boundaries. We will review below these recent results in more detail.

372 2. GRAIN BOUNDARY WETTING 2.1. Bicrystal in contact with the melt If a bicrystal is in contact with the liquid phase, then a grain boundary groove is formed at the surface (Fig. la). Such a groove is characterized by the contact angle 0. The case of 0 = 0 corresponds to complete wetting of the GB by the melt (Fig.l b, c). The shape of the GB groove is dependent upon the mechanism of atom transport and the thermodynamic state of the system. If the bicrystal far from the GB is in thermodynamic equilibrium with the surrounding melt and the shape changes occur due to the dissolution-condensation of the bicrystal material, then the GB groove profile may be calculated analytically [2] and will look as is shown schematically in Figs. 1a and b for incomplete and complete wetting, respectively. It should be noted that in the case of complete wetting by the equilibrium melt (Fig. l b) the GB groove has the finite size and differs from the GB groove in the case of incomplete wetting (Fig. l a) only quantitatively. If the melt is undersaturated with the bicrystal material and the dissolution occurs, then the GB groove will look like a long narrow channel, filled with the melt (Fig. lc), as was first shown in [3]. It was found recently [4-6] that if the solute diffuses from the melt to the bicrystal, the GB directly under the wetting groove (as marked by arrows in Figs. l b and c) may be considered as a quasi-liquid layer about 10 interatomic distances wide. An analogy with a known phenomenon may be found: when a drop of liquid is placed on a solid surface and the liquid wets the surface completely, then the spreading of the drop is accompanied by the presence of a microscopic precursor film ahead of the drop [7]. In the above-mentioned works [4-6] Zn diffusion along tilt GBs in Fe-Si alloys in the presence of a Zn-rich liquid phase has been studied by electron probe microanalysis (Fig.2). By such measurements the dependence of the diffusant concentration on the distance perpendicular to the GB plane is determined firstly. That dependence has a maximum at the GB position. The concentration of the maximum, cb , is accepted as the concentration of the diffusant in the bulk layers adjacent to the GB and in local equilibrium with it. It was found that in the case of complete wetting of the GB by the melt (type Fig. 1c) the penetration profiles of Zn along GBs in Fisher's coordinates (log cb vs y, where y is the distance along the GB) consist of two sections, one with a small slope at high Zn concentrations and one with a large slope at

Figure 1. Bicrystal in contact with the liquid phase (L): (a) incomplete wetting of the GB by the melt, 0>0; (b) and (c) complete wetting, 0--0. (b) The melt is in equilibrium with the crystal; (c) the crystal dissolves in the melt. Arrows denote the region of existence of a thin quasi-liquid film on the GB.

373

Figure 2. Schematic diagram illustrating the electron microprobe measurements of the zinc concentration on the GB.

i ii1.8

7.6 857

975 ~q

1o

21 809

63 790

quasi-L ,!

102h 745"C Sv

I I

' f I I

r"= 1.0

bB I I

I

o,I

t

I

Y Figure 3. The dependencies of the zinc concentration Cb on 430 <001> tilt GBs in Fe-5 at.% Si on the depth y for different temperatures. At the concentrations Cbs and Cbsv the value of GB diffusivity changes abruptly [4]. Cbs and Cbsv correspond to the GB solidus and solvus, respectively (see Fig.4).

374

low Zn concentrations (Fig.3). The transition from one type of behaviour to the other was found to occur at a definite Zn concentration at the GB, which is an equilibrium characteristic of a GB and can be considered as the concentration of the GB phase transition. That concentration depends on the temperature. The ratio of the GB diffusivities in the two regions was approximately 102 which is an indication of the presence of a highly disordered layer in the GBs at high Zn concentrations. The GB region with an extraordinary high diffusivity corresponds to the area marked by arrows in Fig. l c and may be considered as an analogue of a precursor film. The dependence of the concentration of GB phase transition upon temperature (GB phase diagrams) has been determined for alloys with different Si contents. It follows from the Gibbs phase rule that there should not be two-phase regions on such diagrams. Such a GB phase transition line can be considered as the GB solidus line at temperatures above the pefitecfic temperature when the solid is in equilibrium with the liquid (Fig.4). At the temperatures below the pefitectic temperature the line of GB phase transitions can be treated as GB solvus line, because at that temperatures the ternary Fe-Si-Zn solid solution is in equilibrium with the Zn-fich F-phase. So we will use the designations CbS (GB solidus) and CbSv (GB solvus) for the concentrations of the GB phase transition above and below the pefitectic temperature, respectively. Below we will use also the term "premelting" for both situations because of the common manifestations and in order to underline the fact that the phase nucleating on GBs is more disordered than the surrounding bulk. In complete analogy with the bulk solidus line a disordered phase appears on the GBs when crossing that line in direction of high Zn concentrations. The only difference with the bulk solidus is the structure of that disordered layer. For instance, a long-range order can be presented in the quasi-liquid GB layer due to the disturbing influence of the adjacent crystalline grains. Such long-range order disappear when approaching the bulk solidus line, where the complete wetting of the GB by the equilibrium melt occurs. In spite of the fact that the GB phase diagram (Fig.4) is determined for the 430 <001> tilt GB only, the principal features of the GB phase diagrams for other GBs should be the same. Factors decreasing the GB surface tension should shift the CbS and CbSv lines in the direction of higher Zn concentrations. Such a shift can be observed for special large-angle GBs with a low energy. For low-angle GBs the critical misofientation should exist, below which the GB phase transition is unobservable (CbS = CS and CbSv= CSv ). Let us summarize the main features of GB phase diagrams obtained in [4-6]. 1. For all Si contents sharp peaks directed towards low Zn concentrations were observed on GB phase diagrams at the pefitecfic temperature of the Fe-Zn system (782~ (Fig.4). For reference, the Fe-fich end of the binary Fe-Si phase diagram (Fig. 5a) and the Fe rich end of the binary Fe-Zn phase diagram (Fig. 5b) are also shown. 2. In an alloy containing 5 at.% Si a peak on the GB phase diagram directed towards low Zn content was also observed in the temperature vicinity of the Curie point (Fig.4). Such effects are often observed at the intersection of the line of a second-order phase transition with the line of a first-order phase transition and can be explained within the framework of a simple thermodynamical theory [9]. In alloys containing 10 and 12 at.% Si the Curie temperature is below the temperature interval studied and such peaks have not been observed. All of the above may be illustrated by a three-dimensional GB phase diagram having the coordinates temperature, Zn concentration and Si concentration (Fig.6). Such a diagram looks like a twodimensional surface from the left side of which (at low Zn concentrations) GBs with low zinc adsorption are stable, and from the fight side of which GBs are replaced by the layer of quasiliquid (above pefitectic line) or by the layer of F-like phase (below the pefitecfic line). Two

375

T {~ Fe-5at.% S i - Z n 9 0 0 _ 43~

S

CbS//c

a+L 800 er

%

700

Cb I

I

0

I

(zwF I

I

I

I0

I

I

I

I

20

c {at.% Zn}

Figure 4. Fe-5at.% Si-Zn phase diagram [6]. CbS and CbSv are the GB solidus line and GB solvus line, respectively, for a 430<001> tilt GB. cs is the bulk solidus line and Csv is the bulk solvus line. Tr~r is the peritectic temperature and Tc is the Curie temperature. Both at cbS and CbSvthe GB diffusivity changes abruptly.

T(~

[ L

1500

a+L 1000 "-

500 ...... 0

10

"'""l'"'"" 20

c {at.%SI}

Tper

'-" Tc

~+F -

I

I

I

I

0

10

20

30

40

c{at.%Zn}

Figure 5. (a) Fe-rich end of the Fe-Si phase diagram. The binary Fe-Si alloys used for our investigations (5, 6, 10, 12 and 14 at.% Si) are marked on the top. (b) Fe-rich end of the Fe-Zn phase diagram [8].

376

Figure 6. Two-dimensional GB phase transition surface in "Zn concentration-Si concentrationtemperature" space. channels are clearly seen on that surface. The first at constant temperature is associated with the peritectic temperature, and the second, which is bent in direction to low temperatures, is associated with the Curie temperature. 3. Below the Curie point for the Fe-5 at.% Si alloy the premelting line is very close to the bulk solvus line, and below the atomic ordering A2-B2 temperature in an alloy containing 12 at.% Si the complete wetting of the GB by the zinc-rich melt disappears simultaneously with the GB premelting phase transition. So it may be concluded that atomic and spin ordering shortens the region of stability of a GB in a premelted state. We will give detailed analysis of disordered GB film stability later, and here mention only that in the case of the ferromagnetic transition a long-range attractive contribution to the free energy of the premelting layer arises, which is known to prevent wetting. The origin of such a contribution is the attraction of two ferromagnetic fragments divided by the paramagnetic disordered layer along the GB. In the case of atomic ordering A2-B2 in the bulk an additional contribution to the free energy of both the GB and the disordered premelting layer arises due to the disorder in the interface core.

377 Since the premelting layer is more disordered than the GB, such a contribution should be greater for a premelting layer and it looses its stability in bulk ordered state. We should emphasize, howeCer, that above results have been obtained during the study of zinc diffusion from the liquid environment to the solid bicrystal free from zinc. Even if the thermodynamic equilibrium is present locally in the diffusion zone the vacancy flux can strongly influence the observed behaviour. In this respect experiments on intemal wetting, when liquid is present inside the solid, are more pure in a physical sense, and are also more relevant in understanding the behaviour of multiphase systems.

2.2. Wetting in multiphase systems Let us consider a two-component material at such a temperature that the solid solution can coexist with the liquid phase. We will also assume that the condition 27SL
378

e

T

(y+L CbS \ \

L

~Cs

EL

a

T1

-

-

-

o

12 .

.

.

.

.

.

.

o

.

.

.

.

.

.

.

.

.

Q

.

.

.

.

.

.

.

.

.

.

.

.

.

.

Q -

-

(~

A

CB

Figure 7. Different possible microstructures of the two-component material at the temperature T1, at which coexistence of liquid (L) and solid (S) is possible and the liquid phase wets the GBs completely, and the corresponding phase diagram. (a) Polycrystal without GB layers, (b) GBs in premelted state: thick lines correspond to GBs with quasi-liquid layers, (c) a small amount of liquid phase, (d) excess of liquid phase, and (e) corresponding part of the binary phase diagram (schematically). cs and cL are the bulk solidus and liquidus lines, respectively, and cbs is the GB solidus (premelfing) line.

379 in ceramic materials is connected with their manufacturing method. For example, the liquid phase sintering process is often used to densify them, or ceramic materials can be produced by controlled, but incomplete crystallization of a glass (glass ceramics). As an example, the high resolution electron microscopy (HREM) image of a GB in a high-purity Si3N4 ceramics, containing an SiO 2 intergranular phase and doped with 220 ppm Ca, which demonstrates an 11 A wide equilibrium layer of an amorphous phase, is shown in Fig.8. There is also some indication of the presence of an amorphous phase at GBs in YBa2Cu307 superconducting ceramics [17]. It was found by HREM studies that the core region of a small-angle 7.5o <001> tilt GB in this material consists of a continuous string of triangle regions with amorphous appearance and a spacing of about 9 nm. Such amorphous regions have not been found in two other 10 <001> and 5o <001> tilt GBs studied. This is quite easy to understand: the surface tension of a small-angle GB increases rapidly with increasing misorientation angle, while the dependence of crystal - amorphous phase interface surface tension on the orientation is not so strong, and the presence of an amorphous layer on the GB is thermodynamically favourable above some critical misorientation. It was also found that amorphous regions at a GB may grow when they are exposured to the 300 keV electron beam and then shrink after subsequent annealing, but never disappear. This is an indication of the equilibrium character of the intergranular phase. Intergranular glassy films in various superconducting ceramics on the large - angle GBs were observed in [ 18]. It was also found that special GBs (Z3, Z5 and Z13 GBs) are often free

Figure 8. The equilibrium layer of an amorphous SiOg phase in a high-purity Si3N4 doped with 220 ppm Ca. The equilibrium film thickness is 11+ 1 A [ 16].

380 from the intergranular phase, which is consistent with their low surface tension. Presence of 412 nm wide intergranular films on GBs was also reported [19]. The presence of anaorphous or liquid films on GBs should strongly affect the deformation behaviour of the material (we will discuss that question in more details later). It was found during the high- temperature deformation of a Bi-Pb-Sr-Ca-Cu-O superconducting ceramic [20] that at a temperature of 800"0C the specimen could be deformed to as much as 90% without any apparent work-hardening. At a temperature only slightly below (780 0C) specimens fractured after deformation to about 40%. Though HREM studies of GBs in the samples after deformation at 800 0C have not revealed any GB amorphous phases, partial melting of GBs may occur during the deformation process. Such melted regions at GBs should lead to the rapid increase of GB diffusion and sliding rates, and, hence, to the high deformation values.

3. STABILITY OF THIN INTERGRANULAR FILMS 3.1. Two-phase region According to Clarke [21 ], the balance of normal forces, acting on the thin film between two grains, is the following. p+pc+p'=0,

(1)

where p is the applied pressure, Pc is the capillary pressure and p' is the disjoining pressure. The disjoining pressure is a net pressure due to all the interracial forces that might be acting across the intergranular liquid phase, or, in other words, an effective pressure which arises due to the interaction of two solid-liquid interfaces bounding the liquid film. For the most general situation the disjoining pressure consists of the contributions such as those due to the van der Waals or dispersive forces, P'disp , electrical double layer interactions, P'~dl , interaction resulting from the solute adsorption on the two grain surfaces, P'aas, hydrogen-bonding forces, P'hb, anY structural or steric forces, P'st, interaction resulting from the discreteness of the phonon spectra of the liquid film due to it finite width, P'ph, and, in metals, interaction resulting from the discreteness of the conducting electrons spectra of the liquid film due to the it finite width, P'ce. Thus, the disjoining pressure can be formally expressed as '-) p,~) ~-) (• P'= P'~p +P'~a, + P ' ~ + +P' +P'~) +P'o~ ,

(2)

where (+), (-) and (+_) correspond to the attraction, repulsion and variable interaction of two crystal-liquid interfaces, respectively. For the calculation of the net disjoining pressure only the van der Waals and steric forces were taken into account [21]. It has been shown [22], however, that the phonon contribution is very important at high temperatures, and in some circumstances it can exceed the van der Waals force by a factor of 1~, where h is the film thickness and b is the interatomic distance. Analogously to the fact that van der Waals force in the liquid layer results from the fluctuating electromagnetic field energy dependence on h, the phonon contribution to the net force arises from the fluctuating phonon field energy dependence on h. But a good agreement between the calculated and experimental values of h can, in principle, justify the approximation made in [21]. So now we will follow Clarke's analysis [21].

381 For the parallel slab geometry the dispersion force acting on the liquid layer surrounded by the two grains of phase ~ can be expressed as follows [23]:

p,

[

h i e~ (i~)- e~ (i~) disp= 8~2h 3 e~ (i~) + e~ (i~)

]2

d~ = H~r / 6 ~ h

3

(3)

where e~ (iq) and e~ (iq) are the dielectric permittivity of solid and liquid phases, respectively, as a function of an imaginary frequency ig, h is Planck's constant divided by 21t and H ~ is the so-called nonretarded Hamaker constant for two solid slabs separated by the liquid f'flm. Using the procedure proposed in [24], Clarke calculated the Hamaker constant for two alumina slabs separated by fused silica, which turned out to be 5.895kT = 1.63x10 -19 J at 2000 K (k is the Boltzmann constant and T is the absolute temperature). The less accurate estimations using Tabor-Winterton approximation [25] were also made for other ceramic systems; for example, for silicon nitride grains divided by the amorphous SiO2 layer H ~ = 7.6x10 o20 J. It is important that the van der Waals force given by Eq.(3) always corresponds to the attraction of the two crystal-liquid interfaces. So in the absence of an external pressure p and any capillary pressure Pc, the thin intergranular film cannot be stable. In order to stabilize it a repulsive structural or steric force of sufficient strength is necessary. It was recently shown that if impurity cations are present in the intergranular film in ceramics, the electrical doublelayer repulsive interaction is of the same order of magnitude as the structural one and should be taken into account [ 16]. For example, the equilibrium thickness of the intergranular film in Fig.8 was determined by the balance of dispersive, structural and electrical double-layer interaction. The structural force can be calculated using a Landau-Ginsburg expression for the free energy density g of a liquid phase: g=go + arl 2(x) + K[arl(x) / 3x] 2,

(4)

where go, a and K are positive constants. Expression (4) is widely used in solid state physics for describing magnetic phenomena (here rl may be expressed in terms of spontaneous magnetization), spinodal decomposition (11 is connected with the alloy composition), ordering reactions (q is connected with the sublattice occupation probabilities), superconductivity (11 is the density of superconducting electrons), and so on. In our case q may be considered as a crystalline order parameter in a given plane, which should be equal to q0 in one grain and -q0 in the other grain, taking a zero value in the middle of the liquid layer. It means that the two adjacent grains have an opposite orientation (this gives an upper estimation for the structural force). Using standard methods for the free energy minimization one can obtain for the structural disjoining pressure: P'st =

~)o , sinh (h / 2~)

(5)

382

where 4)0 =aq02 and ~ = K/a is a correlation length. 4i0 represents here the free energy of ordering and may be roughly estimated as the heat of melting of the intergranular film material; in the case of silica it gives the value of 9.6 J/mol. Using equations (2) and (5), the equilibrium thickness of the film may be found by the equation

H ~ / 6~h 3 =

aq~ sinh (h / 2~) "

(6)

In the above example (two alumina slabs separated by fused silica) assuming that the correlation length ~ , which physically represents the effective distance of atom interaction [26], is equal to the molecular size of the SiC4 unit, i.e. 0.3 nm; the stable thickness of the film is then 2 nm. An estimation made for other systems where amorphous films have been observed showed that the equilibrium film thicknesses are approximately of the same value. From Eq.(6) the compressive stress may be calculated, which should be applied to the sample in order to remove the equilibrium intergranular phase. In the above example this varies from 70 MPa to 6.5 GPa, depending on the 11o value. This is a very large pressure. Application of smaller compressive stresses can alter the film thickness, but not to a great extent. For instance, to decrease the stable thickness by a distance of about 0.3 nm, a compressive stress of 15 MPa is necessary.

3.2. Single-phaseregion In a single-phase region of the phase diagram, a new term in the disjoining pressure due to the fact that the liquid phase is no longer thermodynamically stable. In a one-component system this term may be easily estimated (see, for example, [27]):

P'therm arises

L AT P'therm . . . . ,

(7)

VmTm

where L is the heat of melting, V m is the molar volume of the material, Tm is the melting temperature and AT is the undercooling. For silica we will have P'therm = 4.2x 108 (AT/Tm) Pa. This means that for supercoolings up to the AT-0.15T m the quasi-liquid film will be stable on the grain boundary. This corresponds to the situation shown in Fig. 7c (GBs are in the premelted state). Let us now estimate the thermodynamic disjoining pressure in the case of twocomponent systems. If in the solid solution of concentration co with cubic grains of average size d a layer of liquid phase enriched by solute appears, then the concentration of the solute in the solid solution decreases by the value Ac. The difference between the free energy of such a heterogeneous system and the free energy of a homogeneous solid solution of concentration co gives the excess free energy AG of a nonequilibrium liquid:

AG=-

- RT

Vm

Ir

/)G~

+

o-oo}

It O~C7~o c,.-Co

Ac=-

E

GL-Go ]Ac,

~ R T (g~-lX 2 ) + ~

v~.

C L ~C 0

(8)

383

G Solid

A

,.-Ac--~ i | o i

A

Co

Cs

eL

B

Concentration Figure 9. Dependence of the free energy G of liquid and solid phases on the second component concentration of a binary A-B system. AG is the excess free energy of the nonequilibrium liquid phase, co is the composition of the solid, and cs and CL are the concentrations of bulk solidus and liquidus lines, respectively.

gl

4 1.5

10

50 2.0

140 1.1

350 580 740 890 1170 1400 MPa 1.6 1.0 1.3 3.9 3.0 6.0 h o%

-

-

.

-

o

t

-

@

I

1.0

0.1

-

IT M

"'-I

I

I

I

I

I

Figure 10. The dependence of the zinc concentration cb on 430 <100> tilt GBs in an Fe-6 at.% Si alloy on the depth y for different values of hydrostatic pressure at a constant temperature of 905 oC. The annealing times are given below the values for the pressure used. The concentration Cbs corresponds to the premelting phase transition on the GB [28].

384 where R is the gas constant, ~ts2and ~t~ are the chemical potentials of solute and solvent atoms in the solid phase, G O and G e are the free energies of the solid and nonequilibrium liquid, respectively. Equation (8) allows a very simple graphic interpretation: The free energy of a heterogeneous system of the same average composition co as a homogeneous one is given by the point A on the corresponding tangent (Fig. 9), and the AG value corresponds to the difference between the energies of the heterogeneous and homogeneous systems. Ac value may be estimated from the substance balance: Ac= (CL- co)h/d,

(9)

Fe-6 at.% Si-Zn 43~

tilt GB

PW

T (~ Tper

800~----~ i

1.5 ~ ,

"~

700 0.5 ~ ~ ~ 0

/

c~ ~'~_...~J

Tc

.// 10

15

5

Fe-6 at.% Si Figure 11. Three-dimensional Fe-6 at.% Si-Zn phase diagram on which the locations of the bulk solidus (Cs) and solvus (Csv), and the solidus (CbS) and solvus (cbSv) for 430 <001> tilt GBs are shown [28]. Pw is the pressure at which the transition from complete to incomplete wetting occurs.

385 where CL is the concentration of the solute in the liquid. Assuming an ideal solid solution behaviour, the difference A=cs-c o , where cs is the solubility limit, to be small, substituting G O by G s and G~. by its equilibrium value one can obtain from Eqs. (8) and (9) the expression for the thermodynamic disjoining pressure: P'therm =

RT(cr. -Cs) VmCs(1-Cs)

(Cs - Co).

(10)

Let us start some numerical estimations for the Fe-Si-Zn system, where the premelted state of GBs was detected in a single-phase region of the phase diagram, corresponding to a value of A -0.1. In the Fe-Zn binary system the peritectic temperature Tper=1055 K, cs = 0.42 and cL=0.93 [8]. Then, from Eq.(10) we have p'therm=300 MPa for the thermodynamical disjoining pressure which suppresses the quasi-liquid film on the GB. Application of the same external pressure should suppress the quasi-liquid film on the GB, too. Indeed, it was demonstrated [28] that at hydrostatic pressures higher than 500 MPa the region of accelerated GB diffusivity, which may be considered as a manifestation of GB premelting, disappears (Fig. 10).

Figure 12. Free energies of the solid and liquid phases of a binary A-B system in the presence of a ferromagnetic-paramagnetic transformation. Cc is the concentration at which the magnetic transition occurs. Shaded area corresponds to the decrease in free energy of an alloy due to magnetic ordering. Solid segments denote the excess free energy of the liquid phase.

386 The transition also manifested itself by the change in dihedral angle 0 at the tip of GB groove from zero (at pressures p<500 MPa) to nonzero values, increasing with p (p>500 MPa). In Fig.11 a schematic three-dimensional GB phase diagram in the temperature, Zn concentration and pressure coordinates is shown in two sections, at p=0 (isobaric) and at T=905 ~ for different pressures (isothermal). It should be noted that in spite of its simplicity the graphic free energy construction (Fig.9) is very useful in the qualitative understanding of various phenomena. Consider, for example, the change of thermodynamic disjoining pressure in the temperature vicinity of the Curie line. The corresponding free energy construction is given in Fig.12. Cc is the concentration of the ferromagnetic-paramagnetic phase transition, so at concentrations C
4.

9R PHASE IN THE STRUCTURE OF TWINS IN Cu AND Ag

Recently, it was found that a few interatomic distance wide layers of another ordered phase exist on GBs in pure metals with low stacking fault energy [29-31]. In these papers computer simulation and HREM observations of incoherent twin boundaries (ICB) in Ag and Cu were performed, and their relative surface tension was measured. Besides the E3 value, another parameter determining the inclination of the twin plane is important. The inclination angle 9 can vary from 9 =0 0 for coherent twin boundary (CTB) to 9 =90 0 for symmetrical incoherent twin boundary (SlTB). In the CTB lattice matching is very good, and the surface tension of such a boundary is small (0.05 of the surface tension of the free surface {011 } in Cu). At the deviation from the coherent orientation the surface energy of a twin increases. Surprisingly, another symmetrical orientation of the twin plane (SITB, 9 =90 0) does not correspond to the energy minimum. The second local GB surface tension mini"mum occurs at 9 = 82 0, and the SITB has a relatively high energy. Such unusual behaviour was connected with the existence of a layer of an ordered 9R phase on ITBs [30], which was proved by HREM observations in Cu and Ag. The 9R structure is relatively close-packed and does not disturb close-packed planes. Denoting by the arrow, relative shift of one close-packed atomic plane with respect to another, the face-centred cubic (fee) structure can be represented by the [,], where square brackets denote one repeat unit. The hexagonal closed-packed (hop) structure in these notations corresponds to [~], and the 9R structure to [***]. One can see that the 9R structure is the simplest after fcc and hcp lattices way to construct the crystalline lattice from the close-packed planes, and its occurrence is not surprising. Detailed consideration shows that the structure of the ITB can be described as a sandwich of the 9R phase, separated

387

9R

9R

jJ Jl J

J

f

a

GB

GB

Figure 13. {111 } planes configuration in the sandwich of the 9R phase on the GB and the adjacent grains. (a) Schematic representation for SITB" (b) schematic representation for ITB with the inclination ~ =90 o-ct/2= 82 o. In the case of SITB interplanar spacing in the 9R phase is reduced by the factor of cost~, which increases the energy. Note that the 9R phase is bounded by asymmetrical small-angle GBs in the case of SITB and by the symmetrical ones in the case of ITB with inclination ~ = 82 o. (c) HREM image of SITB in Cu (courtesy of Dr Ernst).

388 from the adjacent grains by small-angle GBs. In this configuration, however, { 111 } planes of the 9R phase are tilted with respect to those in the grains at an angle o~= 13.50. The coherence conjugation of the grain with the 9R phase in SITB requires the decrease of interplanar spacing of {111 } planes in the sandwich with a corresponding increase in energy (Fig. 13a). If the SITB is rotated by the angle ct/2-- 70, then { 111 } planes in the 9R phase regain their normal interplanar spacing with the corresponding decrease in energy (Fig. 13b). HREM image of SITB in Cu bicrystal is shown on Fig. 13c.

5. GRAIN BOUNDARY AMORPHIZATION DURING PLASTIC DEFORMATION The idea of possible excitation of GB structure during plastic deformation is closely connected with the well-known effect of the spreading of dislocations, trapped in the GB [32]. Transmission electron microscopy (TEM) observations show that the core of the dislocation trapped in the large-angle GB becomes more and more wide, and at some characteristic time td the image of the core disappears. The image disappears only above some critical temperature Td. The parameters Td and td are very sensitive to the type of the GB: It was found that trapped dislocations are more stable at low-energy special GBs, for example in twins. The process of spreading can be formally considered as an incorporation of the "quant" of disorder (dislocation core) into the GB, which produces the disorder-like excitation of the GB structure with the characteristic size rd (the effective radius of spreading) and time of life td . Above some critical value of the dislocation flux Jv towards the GB its structure will remain permanently disordered. Now we consider this process in more details. The density of excited GB sections is given by the equation [33] Pb = Jvtd = ~Evtd ] b,

(11)

where ev is the strain rate averaged over two adjacent grains, ~ is the coefficient depending on the geometry of dislocations falling on GB and b is the interatomic distance. If the excited sections of the GB begin to overlap, the entire GB can be considered as an excited one. That is equivalent to the condition Pb >rd . Taking into account Eq.(11) one can obtain the critical value of the strain rate ~vl, above which the entire GB is excited: evl = b / ~rdtd.

(12)

At strain rates higher than ~vl the GB represents a continuous layer of spreading cores of lattice dislocations. This means that in the GB thermally activated displacements and rearrangements of atoms take place, and the GB can be considered as a quasi-liquid: In this state the diffusivity, mobility and sliding rate of the GB increase considerably. The increase in GB diffusivity during superplastic deformation in Zn-A1 [34], Ni-Cr and Ni-Mo alloys [35] can be understood using such a concept.

389

6. GRAIN BOUNDARIES AT THE BULK ORDER-DISORDER TRANSFORMATIONS Above considerations were limited to the case, when the sandwich of a new phase replaces the GB and makes it wider. It may be an amorphous or quasi-liquid phase (premelting) as in the case of Fe-Si-Zn alloys and various ceramics or 9R crystalline phase as in the case of twin boundaries in Cu and Ag. Another situation is possible when the layer of another phase is formed on the GB and does not destroy the GB itself. Obviously, the atomic structure of the new phase should not differ considerably from the structure of the grain phase for such a phenomenon, and the surface tension of two crystal-GB phase interfaces should be much smaller than the GB surface tension. Then the increase in energy of the system at the creation of two new interfaces can be compensated by the same decrease in energy of the original GB hidden in the layer of the new phase. It is quite natural to expect such GB behaviour near the bulk ordering transition. Because the GB is a plane defect breaking the translation symmetry of the crystal, one can expect that the GB surface tension is somewhat lower in the disordered phase than in the coexisting ordered phase in the two-phase region of the phase diagram. The condition for complete wetting of a GB by a disordered phase can be written as follows: ~/ord

dis

ss > Yss + 2Yod,

(13)

where 7ss - ord dis and Yss are the surface tensions of the GB in the ordered and disordered phase, respectively, and Yoa is the surface tension of the order-disorder interface. Analogously to GB wetting by the melt, if there is only a small quantity of the disordered phase or in the singlephase region of the ordered phase stability a microscopic layer of disordered phase may exist on the GB. Again the width of such a layer is governed by the competition of the attractive van der Waals and thermodynamic forces and the repulsive structural components of the disjoining pressure. It should be noted that the disordered phase can wet (replace) the antiphase boundaries (APB) in ordered alloys. This phenomenon is well investigated both theoretically [36] and experimentally. For instance, existence of a prewetting layer on APBs was found in Co3Pt7 at temperatures of 40 K below the critical temperature Tcr of the bulk order-disorder transition [37]. The behaviour of CTB in the Cu3Au alloys in the temperature vicinity of orderdisorder transition was studied by in situ TEM heating experiments [38]. There are two possible types of CTB atomic structure in Cu3Au: symmetrical (S) and asymmetrical (A) which can be obtained from the S structure by rigid-body displacement in the direction <011 > of the two crystals relative to each other. For the S structure the first and second-nearest neighbour atoms and the distances between them are the same as in a perfect crystal, while in the A structure, like in APB, some nearest and next-nearest neighbours are wrong. It was found [38] that an indication of the disordered layer on CTB appears at 4-10 K below bulk order-disorder transition temperature in S-type CTB and at 7-25 K below for A-type CTB. The prewetting behaviour was shown to be reversible, which is an indication of the equilibrium character of GB films. The only exception was connected with the grain boundary dislocation (GBD): during the cooling of the sample the contrast from the disordered phase near the GBD remained while in the surrounding CTB it had disappeared. It was concluded that the longrange order parameter might be zero at dislocations at a lower temperature than at the interfaces.

390

The same conclusion was made during the study of GB interdiffusion in the Sn-In system [39, 40]. Cahn has pointed out [41] that the GB diffusion may be an indicator of the prewetting f'flrn on the GB. Now we consider this question in more details. The occurrence of the prewetting film on the GB means a change in the original GB structure and a corresponding abrupt change in the GB diffusivity, ff the prewetting film does not participate in the accelerated diffusion transport across the GB, at the bulk transition temperature the GB diffusivity should not exhibit any singularities (Fig.14). Such behaviour was actually observed during the study of In diffusion along large-angle twist GBs in tin and tin-germanium interphase boundaries (IPBs) [39,40]. It was observed that the abrupt change in the GB diffusivity, which is an indication of the appearance of the prewetting film, is 17-25 K in the case of GBs and 1-15 K in the case of IPB below the bulk transition point. It is not by chance that these values are in good agreement with those obtained for the Cu3Au alloy [38]. It is

Figure 14. Schematic diagram showing the variation of the bulk and GB diffusion coefficients in the vicinity of an order-disorder transformation. In the temperature dependence of the bulk diffusion coefficient there is a break at the transition temperature. On the upper part of the diagram the successive states of a GB at different temperatures T 1 ,T 3 and T 4 (see Fig.15) are shown (shaded areas denote ordered phase). The temperature dependence of the GB diffusion coefficient exhibits a step at the temperature, at which a prewetting layer of disordered phase appears on the GB (prewetting phase transition). The change in the GB diffusivity is due to the transformation of its structure, which remains unchanged at the bulk order-disorder transition. At this temperature there is not any singularity in the temperature dependence of the GB diffusion coefficient.

391

T

!% B

~

%

A

...... ~ / ~ /

o

f

Complete GB wetting

iT1 Prewetting Incompletewetting ~ of GB

_~ \

\ B

Figure 15. Schematic binary phase diagram of a system with an order-disorder transformation. o~ and c~' are the disordered and ordered solid solutions, respectively. At some compositions GBs become disordered at temperatures lower than the corresponding temperatures for the bulk (GB prewetting). Above the wetting transition temperature Tw condition (13) is satisfied and complete wetting of the GB by the disordered phase occurs in the two-phase region of the bulk phase diagram.

Figure 16. Schematic diagram showing the subsequent stages of disorder nucleation at the GB. The temperatures T1, T2 and T3 are below the bulk order-disorder transition temperature (see Fig.15). Shaded areas denote the ordered phase. The symbol ..]_ is representing a GB dislocation.

392 known that in ordering systems the most rapid change of long-range order parameter occurs in the temperature interval AT--- 0.1 Tod below the temperature Tod for the order-disorder transition in the bulk (see Fig. 15). Because the surface tensions change during ordering depend on this parameter, if condition (13) is not true far below the transition, then it can be fulfilled in this temperature interval. It gives an upper estimation of the difference between the bulk transition temperature and the temperature of the prewetting film occurrence at the GB of approximately 50 K. In [39] the dependence of the GB phase transition temperature upon the misorientation angle of twist <001> GBs in tin in the angular vicinity of special Z17 misorientation was also studied. It was found that the GB phase transition temperature decreases with increasing deviation from the special misorientation. Based on the fact that the main feature in the structure of near - special GBs is the presence of secondary GB dislocations (SGBDs), it was concluded that at the core of SGBD the disordered phase nucleates at temperatures lower than those for the surrounding GB. The following macroscopic picture of the nucleation of disordered prewetting film on a GB can be proposed, based on [38, 39] (Fig.16): The disordered phase nucleates firstly at the cores of the SGBDs, then, at higher temperatures, the clouds of disordered phase around the SGBD begin to overlap and at the GB phase transition temperature a continuous film of disordered phase forms on the GB. Then, at the bulk order disorder temperature, this film spreads over a macroscopic distance into the adjacent grains.

7. GRAIN BOUNDARY ROUGHENING There is another phenomenon which may be classified as nucleating on the GB disorder. That is GB roughening, when GB becomes unstable with respect to the nucleation of step-like defects. High concentrations of steps with different signs makes the disordered layer on the GB wide enough for TEM observations. Consider this phenomenon in more detail. It is known that special GBs of arbitrary orientation may be faceted, that is, they may be divided into planar sections with inclinations corresponding to the energy minima [31]. According to Cahn [41 ], each such section may be treated as a different GB phase with different values of the intensive thermodynamic parameter ~ (inclination). A similar phenomenon is also observed on the surface of solids. Many single crystals are of a faceted shape with flat sections corresponding to low index crystalline planes which possess particularly low surface tension. For a small deviation of a plane orientation from a low index orientation, steps occur on the surface and the energetically favourable orientation is partially conserved. In this case, however, the surface tension increases by a value proportional to the free energy of steps and their density (if the interaction between the steps is neglected). With increasing temperature the contribution of the steps to the entropy increases and at some temperature formation of steps may become energetically favourable. Their spontaneous formation begins, and every surface atom can be considered as being in a step. All the characteristic sizes in this process are in the order of interatomic distances. Transformation of an atomically flat surface into an atomically rough one occurs: that is, a roughening phase transition takes place. Theoretical calculations show that the width of such a disordered region on the surface below the transition temperature TR diverges as (TR -T)-l/4 and above TR, formally speaking, it should be equal to infinity [42]. It is obvious from the above that the surface tension of the rough surface will vary smoothly with orientation and a crystal surface, corresponding to such orientations, must be of a rounded (non-faceted) shape.

393 In principle, the same phenomenon may also be assumed to occur on GBs. With increasing temperature the boundary may become unstable with respect to the introduction of steps and transform into an atomically rough one. The surface tension of such a boundary will vary smoothly with inclination, and macroscopically the transition is manifested in defaceting. The boundary becomes flat but somewhat diffuse (of course, not of the infinite width, as follows from theory). For large-angle GBs such an experiment has recently been performed by Hsieh and Balluffi [43]. In this work asymmetrical Z3 <111> tilt boundaries in gold and aluminum and an asymmetric Z11 boundary in aluminum were obtained. At room temperature all the boundaries possessed a clearly defined facet structure. A direct observation of the boundaries on heating inside of an electron microscope demonstrated that with increasing temperature the facets gradually become diffuse and at a high temperature the boundaries became flat and thick. As temperature decreases the boundaries resumed their faceted structure again. Some observations of amorphous layers at GBs in ceramics show that the crystalamorphous phase interface can be faceted along low-index planes [44]. So, the above considerations show that instead of the single characteristic temperature for such ceramics, when the intergranular phase becomes soft or melts, there can be another one, corresponding to the roughening of the two crystal-amorphous phase interfaces.

8. HOW GRAIN BOUNDARY FILMS AFFECT ON THE PROPERTIES OF MATERIALS 8.1. Mechanical properties For the situation of complete wetting of the GB by the melt (as shown in Fig. l c) an embritflement of the polycrystal occurs [45]. This phenomenon is well-studied experimentally, however, no attention was paid to the possible role of the GB premelting layer ahead of the spreading macroscopic liquid film. The fact that under applied tensile stress liquid metal embritflement proceeds much more rapidly than without stress, could be an indication of the presence of such films. Indeed, for the GBs oriented perpendicular to the strain axis the tensile stress applied can be considered as a "negative" pressure, which, according to Eq.(1), can make a film on the GB more stable or create it if Eq.(1) was not originally fulfilled without applied stress [46]. It is well-known that glassy phases at GBs in ceramics limit the high-temperature mechanical strength and creep behaviour. Above some temperature the intergranular phase becomes soft or melts, and at this temperature a drastic decrease in strength is observed. In order to improve the high-temperature properties of silicon nitride and other ceramics, exhibiting intergranular phases, GB amorphous phases should be modified in some manner. This can be achieved with the aid of Y203 and A1203 additives, which crystallize partially amorphous layers on the GB [44]. It was also shown that amorphous films at GBs in nuclear waste ceramics can be dissolved by deionized water. Therefore, the amorphous phase in these ceramics should de avoided [ 14]. Consider, according to [47-49], the main features of deformation behaviour of the polycrystalline material containing a liquid or amorphous intergranular phase. It should be noted that the mechanical behaviour has never been investigated in composition ranges shown in Figs.5b and c where the width of the GB layers is only a few interatomic distances. In the studies mentioned the liquid phase was present in such a quantity that it formed macroscopic layers at the GBs, observable by light microscope or TEM at a low magnification. However,

394 one should expect the same principal features of mechanical behaviour of such materials, or, some intermediate situation between the behaviour of the material with macroscopic films on the GBs and single-phase polycrystals. The main features of the high-temperature deformation of polycrystalline alumina with the intergranular glass phase can be described as follows: Cavitation occurs within the glass phase during creep, flexural deformation and fracture. Cavities nucleate at the three- and four-grain junctions and propagate by viscous flow into the two-grain junction channels. The macroscopic distribution of the cavities is inhomogeneous: they are localized into bands. A glass phase extrusion on the free surfaces of the sample was also observed. In [48], where model Cu-10 wt. % Bi alloy was studied, the liquid phase redistribution was shown to be important during the first stages of deformation. At this stage, the liquid is squeezed out of the highly stressed boundaries, which are perpendicular to the compression axis, and flows into less highly stressed boundaries. As the creep develops, the GB sliding becomes important; sliding is accommodated by cavitation in the liquid. The stress strain dependencies for different strain rates and temperatures were obtained for an aluminium alloys containing layers of the quaternary (A1-Mg-Cu-Zn) eutectic with the melting temperature of 477 0C between the Al-based solid solution grains with the melting temperature 610 0C [49]. It was found that at temperatures above the intergranular phase melting temperature the behaviour of the material changes considerably. At high strain rates damage occurs at about 0.1 strain. At low strain rates, a superplastic-type behaviour due to the

A1-Mg-Zn-Cu "E"

~ = 1 0 - 4 S -1

e=10% tl)

4 2

Liquid phase on GBs

Solid phase on GBs 477 I

440

460

480

500

520

Temperature (~ Figure 17. Axial flow stress as a function of the test temperature, measured for 0.1 strain at a low strain rate of 10-4 s-1 for an Al-based AI-Cu-Mg-Zn alloy with layers of the quaternary eutectic on the GBs. An abrupt decrease in the stress is connected with the intergranular phase melting at 477 0C [49].

395 intergranular liquid phase was found. The dependence of the axial flow stress measured for 0.1 strain on the temperature at the deformation rate of 10-4 s-1 is shown in Fig. 17, exhibiting a step in the temperature of the intergranular phase melting.

8.2. Grain boundary migration Migration of the liquid films formed due to the complete wetting of the GBs by the liquid phase is a well-studied phenomenon [50]. The driving force for the migration of such liquid films is, in complete analogy with GBs, the decrease in area of the two crystal-wetting phase interfaces. Considering the problem of migration of the APB wetted by disordered phase, Krzanowski and Allen [51 ] obtained the expression for the APB velocity, which can be also applied to the analysis of the liquid film or premelted GB migration:

v=

2DT(k 1 + k 2) G~ (x~)(x~ -xV)2h

(14)

where D is the interdiffusion coefficient in the wetting phase, y is the surface tension of crystalwetting interphase, k 1 and k 2 are the curvatures of the wetting film, G~ (x ~) is the second derivative of the free energy of the wetting phase (denoted as a) per unit volume, with respect to concentration, evaluated at x~; x ~ and x v are atomic fractions of the matrix component in the wetting phase and the bulk, respectively, and h is the wetting film thickness. In the mentioned work [51 ] the kinetics of the coarsening of the APB structure in Fe-23 at.% AI and Fe-24 at.% A1 was studied in the temperature region, in which there is a small amount of the disordered phase in the alloy, which wets the GBs. It was shown that the layer of the disordered o~ phase greatly (by a factor of 8) slows the domain coarsening kinetics. The same slowing of the migration rate can also be expected for GBs with the wetting (premelting) layers. For example, the stability of the grain size during superplastic deformation [52] can be explained by such slowing down. Another important phenomenon which involves GB migration is the diffusion - induced grain boundary migration (DIGM). During DIGM the GB moves away from its original position when solute diffuses along it. For wetting liquid films on GBs the corresponding phenomenon (liquid film migration, LFM) is also well-established [53]. The driving force for LFM is provided by the coherence strains, arising as the solute diffuses out of the liquid film to the bulk of the crystal. The expression for the velocity may be obtained from Eq.(12) by introducing the additional elastic driving force Eel to the numerator: Eel = Y(nl)~2(Cs _%)2,

(15)

where Y(n 1) is an appropriate, orientation dependent elastic modulus of the shrinking grain, is the atomic misfit, and c s - c o is the difference in composition between the solid far away from the liquid film and the solid in contact with the liquid. Equation (13) was obtained with the assumption that the growing grain is stress-free. As far as we know, there are not systematic observations of GB migration in the vicinity of order-disorder transformations. Activationless migration of GBs in Zn [54] can be connected with the roughening of GBs. Indeed, if the elementary act of GB migration consists

396 of creation of a step-like defect and its propagation along the boundary, then the activation energy of migration consists of two terms: the free energy of creation of the step-like defect and the free energy of its migration. Above the roughening temperature the first term vanishes and the activation energy for migration decreases. So, the Arrhenius plot of the GB mobility should exhibit the upward curvature.

8.3. Diffusion-activated sintering The sintering rate for refractory metal powders can be vastly improved through the addition of the appropriate transition metal, such as Ni, Pd, Pt, Co and Fe [55]. The rate of activated sintering depends on the transition metal concentration, and a barrier concentration c* exists, above which enhanced sintering proceeds. It was established [56] that the sintering acceleration is due to the fast Mo and W atom transport along GBs, which can be covered by a 4-5 interatomic distances thick layer of the phase enriched by the transition metal. At the concentration c* such a layer can appear due to the premelting phase transition on the GB. It was also found [57] that diffusion of the transition metal activates the secondary recrystallisation of a W wire, and Auger electron spectroscopy measurements showed that the activators Ni and Pd are segregated to the GBs up to the recrystallization front in the form of thin layers of a separate phase about 2 nm in thickness.

9. CONCLUDING REMARKS As shown above thin layers of a separate phase can exist on grain boundaries under a wide range of thermodynamical parameters. Sometimes, the term "interface science and engineering" is used to underline the importance of knowledge about grain boundary structure and properties for the understanding of the properties of polycrystals [44]. Thin films interpenetrating the polycrystalline material provide unique possibilities for design. Though decreasing the strength, such films induce superplastic behaviour during deformation. The high rate of diffusion along such films can accelerate processes, where the GB diffusion is involved, for example, grain growth or GB sliding. On the other hand, it follows from Eq.(14) that the GB velocity is proportional to the ratio D/h. So, if the increase in h exceeds the increase in D for the prewetting film formation on the GB, such a film can brake the GB migration, as in the case of APB wetted by a disordered phase. This can be used, if the grain size should be stabilized, for example, during superplastic deformation. It was also shown recently that the presence of "thick" GBs can accelerate the alloying process during DIGM [58]. Conceming the electrical properties of polycrystals, one can speak about 3-dimensional electronic devices. For example, if the thickness of an intergranular phase in superconducting ceramic has the same order of magnitude as the coherence length of superconducting electrons, then such boundaries constitute a network of Josephson junctions. It was also shown that at the temperature of the wetting of APB in CoPt3 alloy by the disordered phase a singularity in the temperature dependence of dp/dT is observed, where 9 is the electrical conductivity of the sample [59].

10. ACKNOWLEDGEMENTS E.R. wish to thank the Alexander-von-Humboldt Foundation for the support of his work at the Max-Planck Institute in Stuttgart. Prof. L.S. Shvindlerman, Prof. D.R. Clarke,

397 Prof. V. Vitek, Dr B.B. Straumal, Dr T. Muschik and Ms L. Bray are greatly acknowledged for helpful discussions.

11. R E F E R E N C E S

.

.

8. .

10. 11. 12. 13. 14. 15. 16.

17. 18. 19. 20. 21. 22. 23. 24. 25. 26. 27. 28. 29.

I. Kaur and W. Gust, Fundamentals of Grain and Interphase Boundary Diffusion, Ziegler Press, Stuttgart, 1989. W.W. Mullins, J. Appl. Phys., 30 (1959) 77. V.E. Fradkov and A.V. Galina, Penetration of Liquid Metal along Grain Boundaries, Preprint ISSP, Chernogolovka, 1986. E.I. Rabkin, V.N. Semenov, L.S. Shvindlerman and B.B. Straumal, Acta metall, mater., 39 (1991) 627. O.I. Noskovic, E.I. Rabkin, V.N. Semenov, B.B. Straumal and L.S. Shvindlerman, Acta metall, mater., 39 (1991) 3091. B.B. Straumal, O.I. Noskovich, V.N. Semenov, L.S. Shvindlerman, W. Gust and B. Predel, Acta metall, mater., 40 (1992) 795. P.G. De Gennes, Rev. Mod. Phys., 57 (1985) 827. T.B. Massalski et al. (eds.), Binary Alloy Phase Diagrams, ASM International, Materials Park, OH, 1990, p. 1444. T. Nishizawa, M. Hasebe and M. Ko, Acta metall., 27 (1979) 817. N. Eustadopolos, Int. Met. Rev., 28 (1983) 189. P.J. Wray, Acta metall., 24 (1976) 125. D.R. Clarke and G. Thomas, J. Am. Ceram. Soc., 60 (1977) 491. D.R. Clarke, J. Appl. Phys. 49 (1978) 2407. D.R. Clarke, J. Am. Ceram. Soc., 64 (1981) C89. D.R. Clarke, J. Am. Ceram. Soc., 72 (1989) 1604. H.-J. Kleebe, M.K. Cinibulk, I. Tanaka, J. Bmley, R.M. Cannon, D.R. Clarke, M.J. Hoffmann, and M. Rtihle, Mat. Res. Soc. Symp. Proc. Vol. 287, Silicon Nitride Ceramics, I-Wei Chen et al. (eds.), Pittsburgh, PA, 1993, p.65. M.F. Chisholm and D.A. Smith, Phil. Mag. A, 59 (1989) 181. J.Y. Laval, C. Delamarre, M.H. Berger and C. Cabanel, J. Physique, Coll.C1, 51 (1990) 991. J. Alarco and G.L. Dunlop, J. Physique, Coll.C1, 51 (1990) 959. T. Murase, K. Kuroda, K. Suzuki and H. Saka, Phil. Mag. A, 62 (1990) 583. D.R. Clarke, J. Am. Ceram. Soc., 70 (1987) 15. A.A. Chernov and UV. Miheev, Poverhnost', 6 (1986) 14. I.E. Dzyaloshinskii, E.M. Lifshitz and L.P. Pitaevskii, Adv. Phys., 10 (1961) 165. D.B. Hough and L.R. White, Adv. Colloid Interface Sci., 14 (1980) 3. D. Tabor and R. Winterton, Proc. R. Soc. London, Ser. A, 312 (1969) 435. J.W. Cahn, Acta metall., 9 (1961) 795. H. Gleiter and B. Chalmers, High-Angle Grain Boundaries, Pergamon Press, Oxford, 1972, p. 2. L.S. Shvindlerman, W. Lojkowski, E.I. Rabkin and B.B. Straumal, J. Physique, Coll. C1, 51 (1990) 629. F. Ernst, M.W. FinNs, D. Hoffman, T. Muschik, U. Sch/3nberger, U. Wolf and M. Methfessel, Phys. Rev. Lett., 69 (1992) 620.

398 30. 31. 32. 33. 34. 35. 36. 37. 38. 39. 40. 41. 42. 43. 44. 45. 46. 47. 48. 49. 50. 51. 52. 53. 54. 55. 56. 57.

58. 59.

U. Wolf, F. Emst, T. Muschik, M.W. Finnis and H.F. Fischmeister, Phil. Mag. A, 66 (1992) 991. T. Muschik, W. Laub, U. Wolf, M.W. Finnis and W. Gust, Acta metall, mater., 41 (1993) 2163. P.H. Pumphrey and H. Gleiter, Phil. Mag. A, 30 (1974) 593. V.N. Perevezentsev, V.V. Rybin and V.N. Chuvil'deev, Acta metall, mater., 40 (1992) 887. S.V. Zemskii, P.L. Gruzin and A.S. Tikhonov, Fizika Chim. Obr. Mater., 5 (1971) 83. S.V. Zemskii, N.E. Fomin, G.K. Mal'tseva and V.A. Karpel'ev, Fizika Chim. Obr. Mater., 4 (1978) 91 B. Widom, Phys. Rev. Lett., 34 (1975) 999. Ch. Leroux, A. Loiseau, M.C. Cadeville, D. Broddin and G. Van Tendello, J. Phys.: Condens. Matter, 2 (1990) 3479. F.D. Tichelaar, F.W. Schapink and Xiaofeng Li, Phil. Mag. A, 65 (1992) 913. E.I. Rabkin, L.S. Shvindlerman and B.B. Straumal, J. Less-Common Met., 158 (1990) 23. E.I. Rabkin, L.S. Shvindlerman and B.B. Straumal, J. Less-Common Met., 159 (1990) 43. J.W. Cahn, J. Physique, Co11.C6, 43 (1982) 199. T. Ohta and K. Kawasaki, Prog. Theor. Phys.,60 (1978) 365. T.E. Hsieh and R.W. Balluffi, Acta metall., 37 (1989) 2133. M.G. Norton and C.B. Carter, MRS Bulletin, 10 (1990) 51. H. Gleiter and B. Chalmers, High-Angle Grain Boundaries, Pergamon Press, Oxford, 1972, p. 89. E.E. Glickman, private communication (1992). D.R. Clarke, J. Mater. Sci., 20 (1985) 1321. B.L. Vaandrager and G.M. Pharr, Acta metall., 37 (1989) 1057. B. Baudelet, M.C. Dang and F. Bordeaux, Scripta metall, mater., 26 (1992) 573. Y.-D. Song, S.-T. Ahn and D.N. Yoon, Acta metall., 33 (1985) 1907. J.E. Krzanowski and S.M. Allen, Acta metall., 31 (1983) 213. K.A. Padmanabhan and G.T. Davies, Superplasticity, Springer, Berlin, 1980. A.H. King, Int. Mater. Rev., 32 (1987) 173. Ch.V. Kopezky, L.S. Shvindlerman and V.G. Sursaeva, Scripta metall., 12 (1978) 953. P.E. Zovas, R.M. German, K.S. Kwang and C.J. Li, J. Metals, 35 (1983) 28. V.V. Skorokhod, V.V. Panichkina, M. Pavicevic, D. Uskokovic and M.M. Ristic, Science of Sint., 16 (1984) 7. L. Kozma, R. Warren and E.-Th. Henig, in: Proc. 7th Riso Int. Symp. Met. Mater. Sci., N. Hansen, D. Juul Jensen, T. Lefters and B. Ralph (eds.), Riso Natl. Lab., Roskilde, 1986. E.I. Rabkin, L.S. Shvindlerman and W. Gust, submitted to Interface Science, 1993. Ch. Leroux, A. Loiseau, M.C. Cadeville and F. Ducastelle, Europhys. Lett., 12 (1990) 155.

Science of Ceramic Interfaces II J. Nowomy (Editor) 9 1994 Elsevier Science B.V. All rights reserved.

Systematic Understanding interfacial phenomena Keizo Uematsu

of

399

ceramic

processing

and

related

and Yao Zhang

Department of Chemistry, Nagaoka University of Technology Kamitomioka 1603-1, Nagaoka, Niigata, Japan 940-21 Abstract T h i s p a p e r p r e s e n t s a new a p p r o a c h for the d e v e l o p m e n t of b e t t e r ceramics. Important roles of interfacial phenomena are emphasized. The following subjects will be discussed with the particular emphasis on the problems and their origin which are encountered in the ceramic production process of powder compaction route.

I. Liquid immersion method - Opening a new approach This is a novel and very powerful method for characterizing processing defects in ceramic green body. Green bodies can be made transparent by an immersion l i q u i d w h i c h s u p p r e s s e s the r e f l e c t i o n at the i n t e r f a c e of particle and s u r r o u n d i n g s . Subsequent examination w i t h an o p t i c a l microscope provides us detailed informations which have never been obtained before. New informations on the micro- and macro structure of green body open a new methodology in the research of ceramic processing. 2. Systematic understanding on features from granules to ceramics Against the implicit expectation, a few very large macro-defects of processing origin were found above in even the best ceramic green bodies f o r m e d w i t h the g r a n u l e s by the p o w d e r c o m p a c t i o n p r o c e s s . They grow during densification with the interfacial energy as the driving force and can not be removed by sintering. Importance in the achievement of defectfree green body is emphasized for the ceramics with better properties. 3. Binder segregation during processing of ceramics Binder segregation occurs during the drying process by the flow of binder solution to the drying surface. The segregation of binder is a major source of defect in green body. A detailed model is proposed to analyze the segregation process. It considers the flow of binder to the surface and back diffusion which is driven by the concentration gradient established by the segregation. Adsorption of binder plays an important role. The model can explain the result of experiment accurately. I. Introduction The elimination of processing defect is crucial for the development of better ceramics which have improved properties[l,2]. Both experiment and t h e o r y show that l a r g e d e f e c t s in g r e e n b o d y are n e v e r r e m o v e d d u r i n g sintering and form fracture origin in ceramics[3-5]. These results are also supported by the long experience in industries that processing of ceramics, particularly the forming process is the critical factor which governs the properties of c e r a m i c s . V a r i o u s s c i e n c e s are i n v o l v e d in f o r m i n g of ceramics. Especially, interfacial chemistry plays a critical role in pre-

400 forming process. Systematic understanding of p r o c e s s i n g m u s t be v e r y important for producing ceramics with greatly improved properties. P r o c e s s i n g , h o w e v e r , has b e e n a d i f f i c u l t s u b j e c t for f u n d a m e n t a l study. One of the m a j o r o b s t a c l e has b e e n the a b s e n c e of a d e q u a t e analytical means which can be used to characterize intermediate products involved in processing. Compared to the rapid progress in other area of ceramic science, scientific understanding of processing is slow and has remained in the state of "Black box". The m a i n r e q u i r e m e n t in the a n a l y t i c a l tool for g r e e n b o d y is the capability of detecting features of low concentration. High resolution of SEM is virtually useless for most cases in processing. This can be easily understood if one recall the very low concentration of fracture origin. We must find and eliminate processing defect of the concentration, say one in whole test peace. In mercury porosimetry, which is another popular tool in processing study, the p o r e size d i s t r i b u t i o n determined essentially represents the a v e r a g e o v e r the w h o l e s p e c i m e n v o l u m e s u b j e c t e d for examination. Information on features of low concentration is wiped out. No information on large processing defect can be obtained with this method[6]. It can be used to understanding the pore size-densification relation which is governed by the average feature of micro-pores. We recently removed the obstacle, at least partially, by developing a new characterization method of green body. This is the liquid immersion technique, which involves an immersion liquid to make green body optically transparent and a subsequent transmission optical microscopic examination[711]. F e a t u r e s of m i c r o n l e v e l s are e a s i l y found. Furthermore, the technique is best suited to detect large processing defect of extremely low concentration. The technique can be applied to major ceramic systems such as alumina [7-9 ], zirconia [ i0, II ], silicon nitride, sil icon carbide and barium titanate, etc, and e n a b l e us to s t a r t the s y s t e m a t i c s t u d y of processing. Namely, this technique opens a new methodology in the research of ceramic processing. This paper presents following subjects. In e a c h subject, the interfacial phenomena play an important role. i. The novel characterization method - liquid immersion method 2. Systematic understanding on features from granules to ceramics 3. Binder segregation during ceramic processing- A possible source of defect in green body. 2. The liquid immersion technique 2.1. The principle of the method The o u t l i n e of an u n i q u e o p t i c a l m i c r o s c o p i c technique w i l l be presented in this section with alumina granules as an example. The specimen used in this section is of commercial product made by spray-drying process. According to the manufacture, the product is of very high purity and can be densified to over 99% of theoretical density at the sintering temperature around 1300~ Fig. 1 illustrates the principle to achieve transparency of specimen. In Fig. l(a), the g r e e n b o d y is o p a q u e . T h i s is d u e to the s e v e r e s c a t t e r i n g of l i g h t at the p a r t i c l e - a i r i n t e r f a c e , but not due to the absorption of light by alumina. The fraction of light reflected at the interface R is given by the equation 1 for light incidence in the direction

401 normal

to the interface[12], R =(n'-l)2/(n'+l)2

(i)

where n' shows the relative refractive index. In the present system, at each interface of alumina (n = 1.77) and air (n = 1.00), approximately 10% of light is reflected. The green body containing numerous interface should be opaque, for all light entering the green body is scattered by reflection in random direction, and no direct transmission of light is allowed. The scattering of light is eliminated by filling the empty space between alumina particles w i t h an i m m e r s i o n liquid which has the refractive index approximately e q u a l to t h a t of a l u m i n a . In t h i s c a s e the r e l a t i v e refractive index becomes approximately unity, and the equation 1 shows that no reflection of light happens at the interface as shown in Fig. l(b).

Figure 2. Effect of refractive index of the immersion liquid on the transparency of alumina green body (n=1.77). From left to right, n=l.0, 1.69, 1.74, 1.76, 1.79

402

Fig. 2 shows the effect of refractive index of immersion liquid on the t r a n s p a r e n c y of a l u m i n a green b o d i e s [ 7 ] . E x c e p t b r o m o n a p h t h a l e n e (n = 1.64), all immersion liquid made the green body transparent and the letter on the other side of the s p e c i m e n s are v i s i b l e . The t r a n s p a r e n c y was improved by close matching of refractive index. Although the transparency was apparently poor with the bromonaphthalene, closer examination showed that good transparency can be achieved with a thin specimen. The optimum refractive index of liquid depends on two factors; i. the size of feature subjected for examination, and 2. the thickness of specimen. For thick specimen, good matching of refractive index is needed to achieve good transmittance. Good matching, however, makes the observation of small features difficult due to little light reflection at the interface of the f e a t u r e and the m a t r i x . For thin specimen, it is p o s s i b l e to a p p l y an immersion liquid which has large mismatching of refractive index. The high light scattering at the interface enable us to detect and characterize small features in green bodies. T y p i c a l p r o c e d u r e for p r e p a r i n g thin t r a n s p a r e n t s p e c i m e n is as follows. First, the green body is fractured with a knife. A small piece is hold with fingers and is hand polished with a grinding paper. No precaution is r e q u i r e d for s u r f a c e r o u g h n e s s , for it is c o v e r e d w i t h the liquid. Thinned specimens (0.2 mm) is immersed in an adequate liquid and, if needed, is evacuated for 5 min. For the observation of powder granules, no special pre-treatment is needed. They can be examined as-prepared. C l e a r l y , the e x p e r i m e n t a l p r o c e d u r e is simple, quick, and does not r e q u i r e any special e q u i p m e n t . T h e s e factors, t o g e t h e r with the h i g h resolution power of optical microscopy make the present method ideal for studying processing of ceramics. The direct image will provide detailed information which cannot be obtained by other techniques. T a b l e 1 s h o w s s y s t e m w h i c h can be e x a m i n e d with this unique characterization method. They include important systems such as alumina, zirconia, silicon nitride, silicon carbide, barium titanate. If immersion liquid with adequate refractive index is available, there is no difficulty in examination.

Table 1 Systems which can be examined with the liquid immersion technique System alumina zirconia magnesia silica spinel silicon nitride PLZT silicon carbide

Refractive Index 1.77 2.20 1.74 1.55 1.72 2.05 2.5 2.38

403 2.2. Structure of granules prepared by spray-drying Fig. 3 shows a SEM micrograph of molded and fractured granules of a typical high grade alumina. The granules have a nearly spherical shape and do not contain large pores. Fig. 4 shows photomicrographs of immersed granules. Some granules have balloon-like structure with a thin wall (lower middle of micrographs). In this structure, many small spherical granules are found inside the thin

Figure 4. Internal structure of alumina granules optical microscope

examined by the

404 spherical wall. When a small spherical granule overlaps a large one, a r o u n d f e a t u r e a p p e a r s in a l a r g e g r a n u l e . A s e r i e s of o b s e r v a t i o n at various focus position is needed to distinguish whether the feature is a pore or an overlapped granule; When the position of focus is varied from the top to the bottom of a granule, the small feature is brought into focus near the middle if it is a pore. If it is an overlapped granule, it can be b r o u g h t into f o c u s w h i l e no p o r t i o n of the l a r g e g r a n u l e is in focus. (Trapped gas bubbles can be readily distinguished from structural pores by their very sharp contrast). In these observation, bromonaphthalene gave good transparency, good contrast and easy handling, and is best suited for this specimen.

Figure 5. Detailed structure of alumina granules examined by the optical microscope of transmission mode with the immersion liquid technique. Fig. 5 shows a more detailed micrograph taken at higher magnification. This micrograph shows the presence of 3-dimensional features in all of the large granules. R o d - or p l a t e - l i k e f e a t u r e s are f o u n d in m a n y of the granules. Less distinct dark-light contrast was also found, showing the p r e s e n c e of some o t h e r k i n d of n o n - u n i f o r m i t y . Apart from these nonuniformities, however, most of granules appear solid, and do not contain observable pores. Fig. 6 shows SEM micrographs of the rod- or plate-like feature. The size and shape of this feature are similar to those found by the optical micrography shown in Fig. 5. In this structure, the particles are much more densely packed than the surrounding matrix. The concentration of this aggregate particle was very low; over 500 ground surfaces of granules have to be carefully examined to find one. Without knowing the presence of the strange feature, one would not have tried t h i s k i n d of l a b o r i o u s examination. This result exemplifies the weakness of SEM in the study of processing. Equation 1 can be used to estimate the maximum granule thickness that can be made transparent. The relative refractive index of alumina against bromonaphthalene is 1.07, causing only 0.1% reflection or loss of incident

405

Figure 6. SEM micrographs for the cross section of alumina granules. (a) low magnification, (b) high magnification l i g h t at e a c h i n t e r f a c e . A l o n g t h e c e n t e r a x i s of e a c h g r a n u l e , approximately 500 alumina particles are present in this specimen (particle size of this alumina: 0.1um). Recalling that two interfaces are present in each particle, only 37% of incident light is transmitted directly through the granule. The portion of directly transmitted light may be even larger than the calculated value. Scattering of light waves depends strongly on the size of the particle, and becomes very small for particles with sizes much smaller than the wavelength used for observation (~0.5 um). Scattering is significant for particles of a size comparable to the wavelength : Better m a t c h i n g of r e f r a c t i v e index is n e e d e d for s p e c i m e n w i t h p a r t i c l e size around 0.5 um. There are at least two possible mechanisms which produce contrast and make the internal structure visible in the optical micrographs. One is the difference of light absorption between the immersion liquid and alumina, and the other is the light lost by scattering at the alumina-liquid interface. To obtain high contrast in the former mechanism, relative difference of

406

absorption coefficient should be high for the two relevant phases. High c o n t r a s t by the l a t t e r m e c h a n i s m r e q u i r e s l a r g e r e l a t i v e d i f f e r e n c e of refractive index for the two phases. This is, however, not compatible with high transparency of the specimen required for microscopic examination. To obtain the best result, experimental conditions must be optimized. 3. SMstematic

understanding

on features

from granules

to ceramics

To understand how the ceramic structure be developed in processing, structures are examined and compared for green and fully sintered bodies. The m o r p h o l o g i c a l c h a n g e of l a r g e p o r e s d u r i n g d e n s i f i c a t i o n is a l s o examined. 3.1. Structure of green body Materials subjected in this section are prepared with the same granules examined above. Granules were uniaxially pressed into bar or pellet at 10MPa and subsequently isostatically pressed at various pressures (20600MPa) to prepare green bodies. Densities of specimens were determined f r o m t h e i r d i m e n s i o n s and w e i g h t s a f t e r b i n d e r r e m o v a l . For i m m e r s i o n liquids, 2-bromonaphthalene(n=l.64) was used. The internal structures of green bodies were also examined on a fractured and/or ground surface with SEM for c o m p a r i s o n . The p o r e size d i s t r i b u t i o n was d e t e r m i n e d by the mercury porosimetry. Fig. 7 shows a series of observations along the vertical direction for an immersed thinned specimen which was isostatically pressed at 20 MPa. These micrographs show a characteristic feature of round shape which is distributed 3-dimensionally in green bodies. The round shape suggests that the feature corresponds to unfractured granules which survived compaction. The sharp boundary line suggests a very thin boundary between unfractured granule and surrounding matrix. The boundary is visible when viewed only f r o m the t a n g e n t i a l d i r e c t i o n ; The r o u n d line r e p r e s e n t s the m a x i m u m diameter of the granule. Most of the spaces between granules are less than 5um. Fig. 8 s h o w s the s i m i l a r o b s e r v a t i o n s for s p e c i m e n s i s o s t a t i c a l l y pressed at various pressures. With increasing pressure, the granules are deformed and packed more tightly. The void spaces between granules are narrowed and their concentration decreases with increasing pressure. The detection of void becomes increasingly difficult with increasing pressure. However, presence of defect was noted even in the specimen isostatically p r e s s e d at 300 MPa. B e s i d e s t h e s e f e a t u r e s , a s m a l l rod- or p l a t e - l i k e feature was also found in the matrix of all specimens. As explained above, they corresponds to the agglomerates of particles. Clearly, agglomerates are not destroyed during compaction process. Fig. 9 shows SEM micrographs of fractured surfaces of specimens which were isostatically pressed at various pressures. At a low forming pressure, many round features were found on the fractured surface, showing that most of granules were not fractured and were preserved in the green body. At pressures exceeding i00 MPa, no round feature was found on the fractured surface. However, the roughness of the surface tended to decrease with increasing pressure, s u g g e s t i n g the p r e s e n c e of a c e r t a i n s t r u c t u r e . According to the supplier, the granules are hard and require high forming pressure which is consistent to the present result.

407

Figure 7. Optical microscopic examination of internal structure along the depth direction.

408

Figure 8. Effect of compaction pressure on the znternal structure of green bodies examined with the liquid immersion technique. (a) 20MPa, (b) 100MPa, (c) 300MPa

409

Figure 9. Effect of compaction pressure on the internal of green bodies examined with SEM. (a) 20MPa, (b) 100MPa, (c) 300MPa

structure

Fig. i0 shows the cumulative pore size distribution determined by the mercury porosimetry for g r e e n b o d i e s p r e p a r e d at v a r i o u s p r e s s u r e s . Clearly, the majority of pores have the size under 1 um. With increasing p r e s s u r e of f o r m i n g , the total v o l u m e of p o r e d e c r e a s e s , and the m o s t

410

frequent pore size which corresponds to the sharp rise in the cumulative pore size curve also decreases. For example, the curve rises sharply at 70 n m a n d 50 n m in t h e g r e e n bodies prepared at 20 M P a a n d 6 0 0 M P a , respectively. The most frequent pore size determined from these curves decreases with increasing pressure and tends to become constant at high pressure; 64 nm at 20MPa, 53 nm at i00 MPa, and 47 nm at 300 MPa and 600 MPa. T h e r e a l s o is a small but f i n i t e a m o u n t of p o r e in the r e g i o n exceeding 1 um for all specimens. However, detailed characterization of t h e s e l a r g e p o r e s is d i f f i c u l t due to t h e i r s m a l l t o t a l v o l u m e s . No c o r r e l a t i o n was f o u n d b e t w e e n the v o l u m e of l a r g e p o r e and the f o r m i n g pressure. The relative densities determined by mercury porosimetry were 53.7, 55.6, 56.2, 59.4 and 59.7% of theoretical for specimens formed at 20, 40, i00, 300 and 600 MPa, r e s p e c t i v e l y . T h e s e v a l u e s a g r e e d to t h o s e determined from weight and size of specimens within 0.5%, except for the specimen formed at i00 MPa for which the mercury porosimetry gives 1.4% lower density. The source of crack-like voids is the unfractured granules in green bodies formed at low pressures. Formation mechanism of this type of void in green body will be discussed in the later section. In contrast to the apparent high concentration of defect, the volume fractions are not significant for large pores. The apparently conflicting result is due to very large examination volume of the present technique. These pores are actually distributed over large specimen volume, and its v o l u m e f r a c t i o n is low. The c h a r a c t e r i s t i c s f o u n d in this s t u d y are believed to be very universal and provide guide-line for the study of more advanced processing. With conventional characterization methods, however, nothing special was found for the present specimen. The pore size distributions determined by the mercury porosimetry were similar to those found by Zheng et ai.[13] and Roosen et al[14]. The small matrix pores represent the majority of total pore volume fractions. The p o r e s i z e s are a p p r o x i m a t e l y a half of the primary particle size and are understandable. The total volume of pore and

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Figure I0. Effect of compaction pressure distribution in alumina green body.

on the pore

size

411

the most frequent pore size decrease with increasing forming pressure. This behavior clearly reflects the r e d u c e d inter-particle distance with increasing pressure. 3.2. Structure

of partially

sintered body[15]

In this section, the morphological change will be studied for large pores in ceramic green body during early stage of densification at ll00~ for 1-64h. The results will show that large processing pores are preserved after densification up to the relative density 83.9% with their shape and size unchanged significantly. The pore size distribution determined with mercury porosimetry will show growth of matrix pore of small size. The specimens are prepared with the same granules used above. The pressure of isostatic pressing is 300 MPa. After the binder removal, the specimen is sintered at ll00~ for 1-64 h in air. Fig. ii shows the observation by the liquid immersion technique on the internal/pore s t r u c t u r e of s p e c i m e n s i n t e r e d for v a r i o u s times. The concentration of pore is low, and the round feature of complete shape is not found in the specimen sintered for lh. The structure of this specimen is almost the same to that of the green body. Many crack-like features are present at junctions of three granules in this specimen. In the specimens sintered for 16 and 64h, although the observed optical image became less sharp due probably to the increased light scattering at particle boundaries, crack-like p o r e s are v i s i b l e . T h e i r size and c o n c e n t r a t i o n are not changed markedly from those of the green body. Fig. 12 shows the cumulative pore volume for the specimen formed at 300MPa and sintered for various periods. Pores with size over 60nm were not found in this specimen. Majority of pores has the size between 40 and 50nm in green body, and 50 and 60nm in all sintered specimens. The pore size increased slightly in the initial sintering period, but remained almost c o n s t a n t t h e r e a f t e r at this f o r m i n g p r e s s u r e . B e h a v i o r of m i c r o - p o r e s during densification is interesting, and will be discussed later in this paper. Mercury porosimetry is, h o w e v e r , powerless in the s t u d y of processing defects. Clearly, large processing pores are preserved in the early stage of densification in sintering. That is, once introduced in the green body, large processing p o r e s can not be r e m o v e d at l e a s t in this stage of densification. If t h e y p e r s i s t t h r o u g h the s u b s e q u e n t d e n s i f i c a t i o n process, they become flaws governing the fracture strength of the whole sintered body. With the exceptionally large size and crack-like shape, they are believed to behave as very detrimental pores which reduce the fracture strength of the sintered body significantly. The presence of large pore in a highly densified body will indeed be shown in the subsequent section in which thinned specimen will be examined directly by an optical microscope, and the r e l a t i o n b e t w e e n p o r e and f r a c t u r e s t r e n g t h w i l l be d i s c u s s e d . These direct observations will strongly show that these pores indeed have s i g n i f i c a n t e f f e c t on the s t r e n g t h of c e r a m i c s and a b e t t e r p r o c e s s i n g before firing is needed for their elimination in the sintered bodies. The stability of large pore in sintering is consistent to the past theoretical arguments of Kingery[3] and Lange et al[4]. According to their discussion, thermodynamic stability of pore is governed by the curvature of pore surface and varies with the relative size of pore and grains. Large pores are characteristically coordinated by many grains, and have negatively

412

Figure ii. Change of internal structure with sintering examined with the l i q u i d i m m e r s i o n t e c h n i q u e . The s p e c i m e n is a l u m i n a body w h i c h is f o r m e d at 300MPa and s i n t e r e d for various time at l l00~ (a) lh, (b) 16h, (c)64h.

413

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Figure 12. Change of pore size distribution with sintering time in alumina body which was formed at 300MPa and sintered at ii00 ~ curved surface. The s u r f a c e e n e r g y of t h i s n e g a t i v e l y curved surface provides the driving force for pore growth. Contrary, small pores have positively curved surface and shrink with time. Although the growth of large pore with sintering is consistent to the thermodynamic argument, an alternative kinetic explanation also seems to be possible. This explanation is b a s e d on the d i f f e r e n t i a l s h r i n k a g e in agglomerated particle and the resultant formation of cracks as presented by Lange et al[16]. In the present study, granules are considered to be a large agglomerate which have a higher particle packing density than the boundary region of granules. More rapid densification among particles in granules can leave cracks at the boundaries of granules. The i m a g e for the b o u n d a r y regions became more noticeable with increasing sintering time and/or density at least in the initial sintering stage. This result shows the enhancement of non-uniformity through the difference of local sintering among particles in highly packed and loosely packed regions as discussed above. Other microstructural features became less noticeable with densification. Small pores tended to disappear with densification, which is c o n s i s t e n t to the p r e d i c t i o n of s i n t e r i n g theory[ i, 2 ]. 3.3. Morphological

change of processing

defect during densification

The morphological change of natural pore with densification was more directly examined in the alumina green body prepared above. Pore inside the body was examined before and after sintering. Growth of large pore with densification was observed. The specimen was similar to w h a t w a s u s e d in t h e s e c t i o n 3. 2. Photomicrographs were taken for various large pores in the green bodies. After the examination, the immersion liquid is removed by evaporation and the s p e c i m e n was s i n t e r e d at 1150 ~ for a n o t h e r g i v e n p e r i o d in air for densification. The same pores for which the photomicrographs were taken

414

above were examined again to understand the morphological change of pores associated with densification. This densification-examination procedure was repeated for 4 times with the total sintering period up to 64 h. Fig. 13 shows the m o r p h o l o g i c a l c h a n g e of m i c r o s t r u c t u r e d u r i n g densification. The relative density of the green body is 59.3%. The green body (a) shows a large crack-like pore in the center of the photomicrograph. Many dark dots and light p l a t e - l i k e s t r u c t u r e s are also noted. After densification to 67.6% of theoretical in Fig. 13(b), the crack-like pore grows; particularly its width increased significantly. Dark dots which were present in the green body became hardly noticeable. The size and shape of plate-like s t r u c t u r e did not c h a n g e significantly. After further densification to 83.9% of theoretical in Fig. 13(c), the crack-like defect has grown further, and its edges has become dull. Dark dots are almost absent. H o w e v e r , the size and shape of p l a t e - l i k e s t r u c t u r e have been hardly changed. In a d d i t i o n to the f e a t u r e s e x p l a i n e d above, all m i c r o g r a p h s show less d i s t i n c t dark p a t t e r n s . They c l e a r l y show the presence of another type of non-uniformity in the specimen, likely a long scale fluctuation of powder packing density. However, their origin will not be examined further in this study. Morphological change of large natural pore can be clearly understood and is consistent to the discussion of the section 3. 2. Some of the large pores are believed to shrink in the later stage of densification where grains have grown large enough to make the sizes of pore and matrix grains comparable[17]. Removal of large pore as presented in Fig. 13 may be difficult even in the final stage of densification. Results of next section will show that large pores over 100um of size are present in normally sintered alumina ceramics. 3.4.

Structure of fully sintered body[18]

As large processing pores are preserved in the intermediate stage of densification, they are likely left in the fully sintered ceramics. In this section, this expectation is examined. Green bodies are prepared in the same method at various compaction pressures as explained above and sintered at 1350~ for lh in air. Thin specimens (0.5 mm) for microstructural examination were cut from the central r e g i o n of s i n t e r e d b o d i e s and their both faces w e r e p o l i s h e d . The translucent thin specimens thus prepared were examined with transmission optical microscope. Specimen for strength measurement were also prepared and subjected for 3-point flexural test. Fig. 14 shows the transmission optical microscopic image of specimens. All s p e c i m e n s c o n t a i n s pores w h i c h w e r e c h a r a c t e r i s t i c a l l y a r r a y e d in approximately circular shape. The concentration of pores decreased with increasing compaction pressure in forming. The origin of these pores are clearly the circular structure of pores which were found in the green and the partially sintered bodies. As expected, large processing pores are preserved in the fully sintered bodies. Fig. 15 shows the strength of specimens prepared at various compaction pressures. The strength increased with increasing compaction pressure. The pores found above clearly behave as fracture origin in sintered ceramics. The slope of the figure decreased with increasing compaction pressure: The fluctuation of strength increased with increasing compaction pressure or/and average strength. This behavior is understandable. Large processing pores

415

Figure 13. Morphological change of processing pore in alumina during sintering. (a) green body, (b) 16h, (c) 64h.

416

Figure 14. Internal structure of slntered body which are formed at various pressures. (a) 40MPa, (b) lOOMpa, (c) 300MPa

417

99.9 99~

95 90~-

.~ o .a o

a.

'

'

20MPa

I 00MPa .

5o- i

I

---

9 --

300MPa

3o 20

,,..,.

I

IO

e

oL _

a

'

40MPa

5

$.

' 2--

~0o

, i

,

I

I

/

400 ' 5 C)O strength (MPo)

,,I

600

Figure 15. Flexural strength of specimens formed at various pressure. of high concentration and of similar size exist in the specimen formed at low compaction pressure. One of them is always present in the stressed region of specimen in flexural test, limiting the strength within a narrow region. In specimen formed at high compaction pressure, the concentration of large p r o c e s s i n g p o r e s is low, and the p r o b a b i l i t y of f i n d i n g large processing pore in the stressed region in flexural specimen is low. In this case, the strength is governed by, in addition to large processing pore, uncontrolled defects such as surface cracks. The strength distribution clearly becomes wide in this case. Fig. 16 shows the fracture surfaces of specimens. Wavy patterns of t y p i c a l size ca. 50um are n o t e d in s p e c i m e n s f o r m e d at low c o m p a c t i o n pressure. The fracture surface became smoother with increasing compaction pressure. The w a v y p a t t e r n c l e a r l y c o r r e s p o n d s to the weak r e g i o n s in specimens. During fracture of specimen, cracks must have propagated through these weak regions selectively, leaving wavy patterns. The origin of these w e a k r e g i o n s is c l e a r l y the p r o c e s s i n g p o r e s left at the b o u n d a r i e s of granules in green bodies.

4. Binder segregation during defect in green body

ceramic

processing - A possible source of

Results in the previous section clearly shows that processing pores are left in green bodies at the boundaries of granules and that they survive the following densification process, forming fracture origin in the sintered c e r a m i c s and r e d u c i n g the s t r e n g t h of c e r a m i c s . In the d e v e l o p m e n t of better ceramics, it is crucial to clarify first the mechanism involved in the f o r m a t i o n of p r o c e s s i n g p o r e s and then to e l i m i n a t e them. In this section, the binder segregation is experimentally confirmed with a simple experiment, and the underlying mechanism is discussed with a mathematical model.

418

Figure 16. Fracture surfaces of specimens which were formed various pressures and sintered to near full density. (a) 40MPa, (b) 100MPa, (c) 300MPa

at

419

4.1. Experimental

confirmation

of binder segregation[19]

Binder segregated in the surface region of granules is one of the very likely sources of pore formation as discussed by Lukaciewitz[20]. This surface layer of high binder content must be formed by the flow of binder s o l u t i o n to the s u r f a c e of s l u r r y d r o p l e t d u r i n g s p r a y d r y i n g p r o c e s s . Since the flow of binder solution is driven by the capillary force of the d r y i n g s u r f a c e and is v e r y u n i v e r s a l in any d r y i n g p r o c e s s , binder segregation is very likely present in any granules prepared by spray drying process. The segregation phenomena involved in the drying of slurry must be understood for the system of powder-binder solution. In this section, the segregation of binder on the drying surface of slurry is examined. The system examined is alumina-PVA, which is one of the most popular system in the processing of ceramics and is the best suited for the model. Alumina powder and PVA used in experiments were both of commercial grades. Weighed amounts of PVA were dissolved in distilled water to form binder solutions of various concentrations (1.6-3.85wt%). The solution (200g) and alumina (300g) were placed in a ball mill and mixed for 24h. The slurry was placed in a small TEFRON beaker and was suspended in a vertical electrical furnace with a thin stainless-steel wire, which was connected to an electric balance with a digital output. The output of the balance was fed to a personal computer and recorded. The drying temperatures were 40 ~ 90~ The actual temperature of the slurry were determined by placing a thin alumel-cromel thermocouple in the middle of slurry. The dried slurry was approximately 4 mm thick. A small piece (10mm x 10mm x 4mm) was cut from the center region of the dried slurry, and was then divided into 5 sections of approximately equal weight from the surface to bottom region. The PVA content in each section was determined with the thermogravimetry. Fig. 17 shows the effect of binder concentration on the distribution of PVA in the consolidated slurry dried at 90~ At all binder concentrations, the I00

50

"

C - 1.60 wt %

90"C

"~ 8 0 o

9 C-3.85

wt%

m C=2.24

wt% wt%

A C=1.60

60

"~ 4 0

E

E

e

--~--

30

40oC

--~---

60oC

"m--

90oC

-.---

80oC

~ ao

40

C 0 0

~

.

0

zo

a.

0 0

a

,

,

,

I

2

3

4

Depth

(mm)

Figure 17. Distribution of PVA in the dried samples of various concentrations of PVA

R

0

5

-

9

-

l

I

I

I

I

0

I

2

3

4

Depth ( m m )

Figure 18. Distribution of PVA in the sample with 1.6wt% PVA dried at various temperatures.

5

420

concentration of PVA was the highest at the top region of the specimen and was evenly low in the inner region. The binder concentration in the top region increased with increasing binder concentration in the solution. Clearly the binder moves to the top region of the specimen in the drying process. F i g . 18 s h o w s t h e e f f e c t of d r y i n g t e m p e r a t u r e on the b i n d e r distribution. The binder distribution became increasingly uniform with decreasing drying temperature. At the drying temperature 90~ the PVA concentration in the top layer is approximately 4 times higher than in the inner regions. At 40~ it was only 1.5 times higher than in the inner regions.

60

I00 IQ

0 ,<

80

=

E

60

E

Clb

5O

40 e

30

O

40

Ob

=

r

< > n

90"C

C - 3 . 8 5 wt %

C-3.85

wt~

20

o

20

I0 I

4~0

I

50

I

60

I

70

'

80

'

90

Temperolure (~

Figure 19. Relation between PVA content in the surface layer of dried sample and drying temperature

40

I

Woter

content 1%)

Figure 20. Effect of PVA concentration on drying rate

F i g . 19 s h o w s t h e e f f e c t of d r y i n g t e m p e r a t u r e on the b i n d e r concentration present in t h e t o p l a y e r of s p e c i m e n s . The binder concentration in the top layer clearly increased with increasing drying temperature. In the s y s t e m w i t h low b i n d e r c o n c e n t r a t i o n (l.6wt%), it increased gradually with increasing temperature. Whereas, it increased r a p i d l y at 4 0 - 6 0 ~ and b e c a m e c o n s t a n t in the s y s t e m with high b i n d e r concentration (3.85wt%).

Fig. 20 show the effect of PVA concentration on the drying rate of the slurry. At the b e g i n n i n g of the d r y i n g p r o c e s s , the d r y i n g rate was constant (constant rate drying). At a critical water content, the drying rate started to decrease with decreasing water content. The drying rate was slightly lower in the specimen of high PVA content than that of low PVA content. The critical water content was lower in the specimen of low PVA content than that of high PVA content. Fig. 21 shows the effect of drying temperature on the drying rate. With i n c r e a s i n g d r y i n g t e m p e r a t u r e , the d r y i n g rate i n c r e a s e d and the

421

600

.~l 50

A

E

e

C - 3 . 8 5 wt % "" 0

40-

E e,,, 0

90~

i

".

30

s,.

tl. :E

400

m m Q e-

~' 2 0

=0

-1-

200

L

o I0

0

40

30

20

Water

I0

0

0

content (%)

Figure 21. Effect of drying temperature on drying rate

I

I

I

I

20

40

60

80

I00

PVA content ( m g / g N = 0 3 )

Figure 22. Relation between PVA content in the surface layer of dried sample and surface hardness.

critical water content increased. Note, however, that this critical water content does not correspond to the state where the shrinkage with drying stops. With the density of dried body, it is possible to determine the apparent water content where the shrinkage stops in drying. This water content is presented by the dotted line in the figure. This value is much h i g h e r t h a n the a p p a r e n t w a t e r c o n t e n t at the e n d of the c o n s t a n t r a t e period. Table 2 summarizes the critical water contents and constant drying rate periods for all specimens. The constant rate period decreased drastically with increasing drying temperature. Fig. 22 shows the relation between the PVA content and Vickers microhardness in the surface region of specimen. The hardness clearly increased Table 2 Drying behaviors of the slurries at various conditions PVA(wt%) 3.85 3.85 3.85 3.85 2.24 1.60 1.60 1.60 1.60

Drying temp.(~ 90 80 60 40 90 90 80 60 40

Critical water cont.(%) 30.2 30.2 27 22 29.8 28.1 27.4 25.2 19.6

Drying time(h) 5.7 7.8 16 83 5.3 4.9 6.3 13 67

Const. rate period(h) 0.9 1.2 3.7 33.2 I.i 1.2 1.7 3.9 34.4

Av. Slurry temp.(~ 60 55 40 30 60 60 55 40 30

422

with increasing PVA content in the surface region. The hardness was 26MPa for i n n e r r e g i o n s of all s p e c i m e n s and was m u c h l o w e r t h a n t h a t of the surface. This study confirmed marked segregation of PVA binder on the drying surface. After the drying of gel[21], the surface segregation of binder during drying process is assumed to occur through the following mechanism: Binder molecules dissolved in a solvent are carried to the surface by the f l o w of s o l v e n t and are a c c u m u l a t e d t h e r e a f t e r the s o l v e n t v a p o r i z e s . Although back diffusion is driven by this concentration gradient and is favorable for reducing the surface segregation, a large portion of binder molecules contained in the solution eventually moves to the surface and segregates. 4. 2. M e c h a n i s m of binder segregation

in ceramic processing

The surface segregation of PVA results from its migration to surface caused by a capillary force and is reduced by the back-diffusion driven by the concentration gradient of PVA. After the drying of a gel, drying of a ceramic slurry consists of two stages, a constant rate period(CRP) and a falling rate period(FRP)[22-24]. During the CRP, the surface of slurry is kept wet, and the water evaporates at the exterior surface. To be evaporated, water inside the slurry must be carried to the surface by the capillary force. At this stage, the volume

layer: 1

i-1 i i+1 qr.~-, qm, qm.~

C!m~

._ ,0rat,

n Cran-1

t ,orat,on -

O. s.onf col

+O.,us.on q~-i qD, qoul

CDn-1

C1 (1) tO

o

X

n

surface

~

C

n interior

Figure 23. Schematic diagram of the mathematical movement of PVA

model

for

423 of the slurry decreases as the liquid content in the slurry decreases with drying. The distance between particles decreases with drying in this stage. At a certain water content, the particles touches and further shrinkage is not allowed. The supply of water to the surface stops and the drying front (the water/vapor interface) retreats into the slurry, The drying stage enters the falling rate period; the rate of evaporation decreases[21] with drying. If the slurry contains binder soluble in water, the binder molecules are carried to the drying surface by the flow of solution in the constant rate p e r i o d . N a m e l y , the d r i v i n g f o r c e of b i n d e r s e g r e g a t i o n is the capillary force. Once the concentration of binder in surface region is r a i s e d r e l a t i v e to the i n t e r i o r r e g i o n , d i f f u s i o n p r o c e s s s t a r t s to be involved. With the concentration gradient of binder established, binder molecules in s u r f a c e region diffuse back to the interior region. Based on the behaviors of binder as stated above, a mathematical model can be e s t a b l i s h e d as s h o w n in Fig. 23. The d r y i n g s l u r r y is d i v i d e d to n layers. The binder molecules can move among layers as shown in the model. In the m o d e l , qe is e v a p o r a t i n g rate p e r u n i t area, qmi the a m o u n t of b i n d e r w h i c h m i g r a t e s f r o m the ( i + l ) - t h l a y e r i n t o i-th l a y e r p e r u n i t interface area in a time interval t, and qDi the amount of binder which diffuses from the i-th layer into the (i+l)-th layer per unit interface area in the time. In the model, some assumptions were made as follows: (i) The binder is uniformly distributed in each layer. (2) The rate for the liquid to transfer to the surface is equal to the evaporating rate. (3) Drying shrinkage occurs uniformly all over the slurry, and only in the direction of depth. (4) The diffusion coefficient of binder is affected only by temperature. (5) The binder molecules can transfer until a upper limit concentration of binder 45wt%. (6) The migration stops as drying shrinkage stops. The amount of binder in each layer after a certain time interval t can be calculated. By the numerical solution of Fick's first law for onedimensional diffusion, the amount of binder diffusing from the i-th layer into the (i+l)-th layer qDi is qDi=-D*t*(Ci+l, j-l-Ci, j-l)/hj-I

(3)

w h e r e D is t h e d i f f u s i o n coefficient, t the time interval, C the concentration of the binder, subscript i and (i+l) represent the i-th and the (i+l)-th layers respectively, j-I represents the time (j-l)*t, and h the thickness of the layer. The amount of binder diffusing into i-th layer qDi-I can be calculated in the same way. Then the change of the amount of binder in the i-th layer which is caused by diffusion is given by AqDi=qDi_l-qD i

(4)

=D*t*(Ci_l,j_l-2Ci,j_l+Ci+l,J_l)/hj_l water layer

The migration of binder can be calculated as follows. The amount of evaporated in the time t is qe*t, that is, the volume of the first (surface layer) should decrease by qe*t (specific gravity of water is

424

taken as i). As the drying shrinkage is uniform throughout the drying slurry during the CRP, and the volume reduction of each layer is equal and qe*t/n, the first layer will be supplied with solution of the volume (qe*tqe*t/n) by the second layer, and the amount of binder carried into the first layer by this flow of solution is qml=qe*t*(l-i/n)*C2

(5)

Similarly, the change of the amount of binder is caused by the flow of solution is given by Aqmi=qmi-qmi_ 1

in the i-th layer which

(6)

=qe*t* {(l-i/n)*Ci, j-l- (I- (i-l)/n)*Ci_ 1 } The amount Wi, j and the concentration C . of binder in the i-th layer 1,3 at the time j*t are given by the following equations by considering the mass transport by diffusion and the flow of solution. Wi,j=Wi,j_l+AqDi+Aqmi

(7)

Ci,j=Wi,j/V j

(8)

(i=1,2, ....... n) In this equation, V~ is the volume of the solution in each layer at the time 3 j*t. The initial condition and the boundary condition are respectively at j=0,

Ci, j=C 0

(i=1,2 ....... ,n)

at i=n,

qDi=0; qmi=0

(9) (i0)

where C O represents the initial concentration. If the evaporation rate of solvent qe' the diffusion coefficient binder D, and the initial conditions of slurry are known, the movement free binder can be simulated by the above equations. The distribution binder during and after the drying process can be calculated.

of of of

4. 3. Results

Distribution of PVA in the slurry of PVA-water-alumina system after the drying process was calculated by the mathematical model stated above. Fig. 24 shows an example of drying rate and the slurry temperature at the surrounding temperature 90~ After increased rapidly at the beginning of drying process, the drying temperature increases gradually with time and again i n c r e a s e d r a p i d l y at the end of the p r o c e s s . The d r y i n g rate i n c r e a s e d s h a r p l y at the b e g i n n i n g of the process, and a f t e r r e m a i n e d constant for short time, it decreased gradually toward the end of process. Fig. 25 shows the diffusion coefficient of PVA in porous specimens. The diffusion coefficient increased with increasing temperature as was found in aqueous solution. However, the absolute value of diffusion coefficient is o r d e r s m a g n i t u d e lower in the p o r o u s m a t e r i a l than in the a q u e o u s solution. The small diffusion coefficient shows that the diffusion of PVA

425

4

90 80

E~

60

ID

50 40

"~

30 0

20 0

80

160

240

320

Dr~ng time (rain) Figure 24. Relation between and drying time.

7

I

drying

i

i

I

I

rate,

I

I

I

I

temperature

of slurry

i

65

o

2

i5

0

g

I

290 300

310 320

330 340 350 360

Temperature Figure 25. Diffusion

(K)

of coefficient

of PVA in slurry

involves adsorption-desorption process. The diffusion of PVA in slurry is believed to be similar to that in the porous material, and is represented by this figure. Fig. 26 shows the distribution of PVA in the specimen of various PVA contents dried at 90~ The points in the figure are the experimental results, and the lines are the results of simulation. The results show that

426 150 A

O r < ~100

9 3.85wt% I 2.24wt% 9 1.60wt%

,,,,,,

E ,lw

e~

c o

0

50

0

1

2

3

4

5

Depth (ram) Figure 26. Comparison of theoretical and experimental distributions of PVA. Solid lines show theoretical results.

100 A

r

O

80

O~A--3.85wt%

~ 6o E e

40

.

c 0

o

a

20

a.

0

9

30

I

40

.

II

50

I

I

60

9

I

70

9

I

80

9

I

90

9

100

Temperature (~ Figure 27. Comparison of theoretical and experimental PVA content on the top layer of dried bodies. Solid lines show theoretical results. the simulation agrees well with experimental results of PVA distribution, and that the mathematical model represents the segregation of PVA during drying process. Fig. 27 shows the PVA at the top layer of specimen dried at various temperatures. The points show the experimental results, and the lines shows the result of simulation. The agreement is good in the high temperature

427

region, but was only fair at the low temperature region. The fundamental agreement shows that the mathematical model is basically correct. However, the fair agreement in the low temperature region also suggests the presence of a certain unknown parameter involved in the segregation process. Segregation of binder in drying process can be expressed by the present mathematical model. Effects of drying temperature and concentration on the segregation of PVA could be understood with the present model as shown in Figs. 26 and 27. The s e g r e g a t i o n became increasingly significant with increasing temperature. This was due to the increased drying rate relative to the d i f f u s i o n rate w i t h i n c r e a s i n g t e m p e r a t u r e . The t e m p e r a t u r e c h a n g e of drying rate is much larger than that of diffusion rate. In the case of high concentration of PVA, the constant rate period of drying was shortened at high drying temperature. The constant rate period was f i n i s h e d w h e n the d r y i n g s h r i n k a g e c o u l d still take p l a c e . This behavior can be explained by the formation of thick layer of gelated PVA (approximately 100um) in the surface region[7,9]. 4. 4. Effect of adsorption

on binder segregation

in spray-dried granules

The same mechanism discussed above must be involved in the segregation of binder on the surface of granules made by the spray drying process. This segregated binder can explain the formation of non-uniform structure in green body formed through powder compaction process. The mechanical property should be different for the surface and the i n t e r i o r r e g i o n s of g r a n u l e s ; s u r f a c e r e g i o n can be c o n s i d e r e d as a composite with high organic content and only deformation without volume change is allowed during the compaction process. No change is expected in the average distance between particles in the compaction process. Whereas, the compaction can reduce the particle distance for the interior region of granules where the binder content is low. After the binder is removed by heating, regions of low powder packing density are left in the boundary regions of granules. Adsorption of binder on solid surface is often found. It may affects the segregation of binder in drying process, for it immobilize the binder. In addition to the analysis of the former section, the effect of adsorption should be taken into account in a more detailed model. The amount of PVA segregated should be governed mainly by the following three factors; i. the f l o w of w a t e r w h i c h c a r r i e s b i n d e r m o l e c u l e s t o w a r d s u r f a c e . 2. the diffusion which sends the binder molecules backwards against the water flow. 3. Adsorption of binder on alumina surface. Adsorption reduces the concentration of binder in the liquid phase, and is expected to reduce the surface segregation of binder. Our preliminary experiment showed that it can explain the disagreement between the theory and the experiment in Fig. 27. The adsorption of PVA on alumina increased with decreasing temperature. At low temperature the large fraction of PVA added was immobilized by adsorption. The above simple theory overestimates the amount of mobile binder and the surface segregation. Adsorption is a complicated subject and is affected by various factors, which include temperature, time, surface conditions, etc. Importance of time can not be overemphasized. Adsorption is often slow process in the system involving polymer. A certain system needs over one year period for equilibration. This may explain the importance of aging often practiced in

428

the production of traditional ceramic. Understanding the role of adsorption will be needed. The relevance of adsorption on ceramic processing has been only briefly examined, and must be an important subject for future study. 5. Conclusions This paper is the first which shows the systematic understanding for the formation mechanism of fracture origin in high performance ceramics. B i n d e r is s e g r e g a t e d at the s u r f a c e of g r a n u l e s by the flow of b i n d e r solution to the its surface in drying stage. The segregated binder forms low-density regions at the boundaries of granules in the green body. These regions are not eliminated during the densification process due to the low densification rate in these regions or/and thermodynamic stability of large pores, forming defects in final ceramics. These defects behave as fracture origin, reducing strength of ceramics. Various interfacial phenomena are invloved in the whole process, some of which were discussed in this paper. These detailed systematic understanding owes almost entirely to the novel characterization method presented in this paper. With the mathematical model also presented in this paper to explain the segregation behaivor of binder, this novel characterization method opens up a new methodology for the research in ceramic processing. 6. References

i. W. D. Kingery, "Ceramic Processing Before Firing", Ed. by G. Y. Onoda and L. L. Hench, pp. 291, Wiley, New York (1978) 2. F. F. Lange, "Fracture Mechanics of Ceramics", Vol. i, Ed. by R. C. Bradt, D. P. H. Hasselman and F. F. Lange, pp. 3, Plenum, New York (1974) 3. W. D. Kingery and B. Francois, pp. 471 in Sintering and Related Phenomena, Ed. by G. C. Kuczynski, N. A. Hooton, and C. F. Gibbon, Gordon and Breach, New York, 1976 4. B. J. Kellett and F. F. Lange, J. Am. Ceram. Soc., 72(1989)725 5. K. Uematsu, M. Miyashita, J.-Y. Kim, and N. Uchida, J. Mater. Sci., 75 (1992)1016 6. J. Zheng, J. S. Reed, J. Am. Ceram. Soc,. 75(1992)3498 7. K. Uematsu, J-Y. Kim, Z. Kato, N. Uchida and K. Saito, Nippon Seramikkusu Kyokai Gakujuturonnbunnsi, 98151515-6(1990) 8. K. Uematsu, J.-Y. Kim, M. Miyashita, N. Uchida and K. Saito, J. Am. Ceram. Soc., 73(1990)2555 9. K. Uematsu, M. Miyashita, J.-Y. Kim, Z. Kato, and N. Uchida, J. Am. Ceram. Soc., 741912170-74(1991) I0. J. Y. Kim, M. Inoue, Z. Kato, N. Uchida, K. Saito and K. Uematsu, J. Mater. Sci., 26(1991)2215 ii. J.-Y. Kim, M. Miyashita, M. Inoue, N. Uchida, K. Saito and K. Uematsu, J. Mater. Sci., 27(1992)587 12. W. D. Kingery, H. K. Bowen and D. R. Uhlmann, in Introduction to Ceramics, 2nd Ed., Chap. 13, pp.646-703, Wiley, New York(1976) 13. J. Zheng and J. S. Reed, J. Am. Ceram. Soc., 72(1989)810 14. A. Roosen and H. K. Bowen, J. Am. Ceram. Soc., 71(1988)970 15. J.-Y. Kim, M. Miyashita, N. Uchida and K. Uematsu, J. Mater. Sci., in press

429 16. F. F. Lange, J. Am. Ceram. Soc., 67(1984)83 17. F. F. Lange and Metcalf, J. Am. Ceram. Soc., 66(1983)398 18. M. Miyashia, J.-Y. Kim, Z. Kato, N. Uchida and K. Uematsu, J. Ceram. Soc. Japan, 100(1992)1357 19. Y. Zhang, X.-X. Tang, Z. Kato, N. Uchida and K. Uematsu, J. Ceram. Soc. Japan, 100(1992) 20. S. J. Lukasiewicz, J. Am. Ceram. Soc., 72141617-24(1989) 21. G. W. Scherer, J. Am. Ceram. Soc., 73(1990)3 22. R. K. Dwriredi, J. Mater. Sci. Lett., 5(1986)373 23. E. J. Crosby and W. R. Marshall, Jr., Chem. Eng. Progr., 54(1958)56 24. H. H. Macey, Trans. Br. Ceram. Soc., 41(1942)73

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Science of CeramicInterfacesII J. Nowotny(Editor) 9 1994ElsevierScienceB.V. All rightsreserved.

431

Interface Phenomena in Synroc, a Titanate-based Nuclear Waste Ceramic E.R. Vance, C.J. Ball, M.G. Blackford, R.A. Day, G.R. Lumpkin, K.L. Smith, K.P. Hart, P. McGlirm and G.J. Thorogood, Advanced Materials Program, Ansto, Private Mail Bag 1, Menai 2234, New South Wales, Australia

Abstract Several aspects of Synroc which fall into the broad class of interface phenomena are discussed. These are radiation damage processes which give rise to interfaces between damage tracks and neighbouring unirradiated material, intergranular films which have deleterious effects on chemical durability, and aqueous leaching of Synroc which takes place prim_arily at the interface between the solid and groundwater. 1. INTRODUCTION Synroc was formulated in 1978 by Prof. A.E. Ringwood and colleagues[I] at the Australian National University. It is a dense ceramic designed to incorporate and immobilise waste fission products and actinides (see Table 1) arising from reprocessing of spent nuclear fuel. Table 1 Approximate composition of High-level Waste (HLW) Oxide

(wt%)

Oxide

Cr20 s Cs20(+Rb2 O) Y203 Ru02 Ag20 Ce2Os(+La203)

0.8 8.3 1.6 7.5 0.2 12.2

Fe20 a 3.8 SrO 2.7 ZrO 2 12.5 Rh203 1.2 CdO 0.2 Nd203 15.5 (+Pr6011+PIn20 z)

Gd2Oa 3.7 (+Sm2Oa+Eu2Oa+Am2Os+Cm2Oa) UO 2 5.2 (+NpO2+Pu02)

(wt%)

Oxide

(wt%)

NiO BaO MoOs PdO TeO 2 P205

0.3 3.9 13.3 3.7 1.9 1.7

432 Minor elements at >0.1% level not included. In non-active simulations of waste-loaded Synroc, Ce is substituted for the actinide elements and the bracketed elements are substituted by the chemically-similar main elements. Synroc has been under development at Ansto since 1979 in collaboration with the Australian National University. It is composed mainly of thermodynamically compatible titanate phases which have very resistate mineral analogues in nature. It is formed from a very reactive calcined mixture of waste and additional cations (see Tables 1 and 2) by hot-pressing at 1150-1200~ MPa for a few hours in sealed metal containers[2]. The overall composition, in terms of oxides is as follows (wt%): A1203(4.3); BaO(4.4); CAO(8.9); TIO2(57.0); ZRO2(5.4); HLW(20). The phases, their approximate (waste-free) compositions and abundances, and the waste elements which they contain in dilute solid solution, are shown in Table 2 for Synroc-C, the ceramic designed for Purex reprocessing waste. Table 2 Composition and mineralogy of Synroc-C Phase Hollandite, Ba1.14(A1,Ti)2.28Ti5.71016 Zirconolite, CaZrWi20 7 Perovskite, CaWiO3 Ti oxides Ca-A1 titanates and aluminas Alloy phases

(a)RE =

rare earths,

An

=

(wt%)

Radionuclides

30 30 20 10 5 5

Cs,Rb RE,An(a) RE,An(a)

Mo,Tc,Pd ....

actinides

Thousands of aqueous leaching tests covering a wide range of temperature, pH, and groundwater simulations have been carried out over the years and confirm excellent chemical durability of Synroc. The leach rates of the most soluble elements, i.e. the alkalis and alkaline earths, are in the range of 0.010.1g/m2/day at temperatures of order 100~ for short times (1-10 days), and they decrease to 103-104g/m2/day at times of order l0 s days[2]. The temperature of 100~ corresponds to that expected in the repository soon ailer emplacement, although there are abundant data which show that Synroc is very resistant to aqueous attack at temperatures as high as 400 or 500~ Recent work on the Synroc project at Ansto has centred on technology and process development, Synroc design for reprocessing wastes of different compositions, radiation damage studies, and increasing relevant chemical

433 knowledge of the behaviour of Synroc in a variety of aqueous media at different temperatures to allow credible prediction of repository behaviour over times of 103-106 years. A key feature of the project is the existence of numerous collaborative agreements with other countries. Interface phenomena play an important part in the design and performance of all ceramics. Unfortunately, relatively little is known of interface behaviour in Synroc. In this paper however we discuss three aspects of Synroc which lie with the general class of interface phenomena, namely:

(a)

atomic-scale phenomena at the interfaces of radiation damage "tracks" and the surrounding undisturbed crystalline matrix,

(b)

intergranular glassy phase formation, and

(c)

solid-liquid reactions occurring when Synroc is exposed to aqueous solutions.

Differential stresses due to radiation damage in the different phases in Synroc also constitute an interface effect. 2. RADIATION DAMAGE EFFECTS The approximate numbers of disturbed atoms produced by different nuclear particles in solids in general and the particles' trajectory lengths are given in Table 3. Table 3 Approximate numbers of disturbed atoms and trajectory lengths for different energetic nuclear particles in solids Damage Agency Alpha particle Alpha recoil Beta particle GAmma ray

Energy(MeV) 5 0.1 0.1-1 1

Range (pro) Atomic Displacements 20 0.02 103 108

102 10z 1 10-2

Many more atoms may be displaced from their initial sites during the damage event than are lei~ on structurally "incorrect" sites alter the "thermal spike" has decayed, and recovery of some of the relatively permanently disturbed atoms can take place subsequently. The recovery process is obviously enhanced at higher temperatures.

434 Many studies of radiation damage effects on the structure and leachability in aqueous media of Synroc and its constituent minerals have been made, using (a) natural metamict perovskite and zirconolite minerals which have been damaged over timespans of millions of years via the alpha-decay of U and Th (and subsequent daughter products) present in dilute solid solution when the minerals were formed; (b) fast neutrons from nuclear reactors; (c) heavy ions in accelerators; and (d) doping with short-lived alpha-emitters such as 238pu or 2~Cm (respective half-lives = 87 and 18 years). In Synroc containing our reference level of 20 wt% of high-level waste from Purex reprocessing of power reactor fuel, the radiation damage after densification is virtually all due to alpha-decay events from waste actinides. There is no doubt that nuclear particles in susceptible materials produce localised regions of permanent structural dsmage which delineate their tracks. For example, alpha-recoil tracks have been seen in mica[3]. In met~mict minerals, the alpha-recoil tracks occur in conjunction with an equal number of alpha-particles, which do not produce visible tracks, together with fission tracks which arise from the spontaneous fission of 238U. The fission tracks are much less abundant than alpha-tracks (the half-life for spontaneous fission of 238U is of the order of 1016 years, compared to 109-10~~years for alpha decay of U and Th). However they are much easier to observe as the energy per fission is roughly 200 MeV, compared to 0.1 MeV for the alpha-recoil (see above) and many more atoms are disturbed than in an alpha event. Certainly fission tracks in minerals are subject to preferential erosion with aqueous etchants such as HF, boiling NaOH or H3PO4. However, it is not clear whether the tracks can be etched out by groundwater in the same way, although the weakening of the bonding between the atoms in the d~maged region would suggest enhanced leachability. Similar considerations apply to alpha-recoil tracks. The localised stress between the damaged region and the matrix arising from the increase in the molar vol,_,me of the d_~_m_agedregion (see below) might also suggest enhanced leachability by a stress corrosion mechanism. However to make theoretical calculations using realistic models of these effects would not be easy. As the density of tracks increases, there is an increasing amount of track overlap, and the whole crystal may eventually become completely disordered. Under these conditions the material almost always registers a nett volume expansion, which can be as large as several percent, and which can be readily detected by macroscopic density measurements. In the early stages of the damage process (low radiation fluences), both X-ray diffraction and density measurements indicate a lattice expansion. From simple calculations based on particle fluences and displaced atoms per event, it can be deduced that the disturbed region in a single alpha-recoil track may well be semi-amorphous, with an associated expansion of several percent. These phenomena can occur

435 in radioactive Synroc after long periods. Some recent results [4] showing the density decrease for Synroc-C doped with ~ C m are shown in Figure 1.

0

o

|

|

|

|

2

4

6

8

|

|

1o l~

J

|

i4

l~

is

Alpha Decay Dose (I0" ~ g " ) l

.

o 1o2

A

lo 3

I

lo 4

!

lo s

ASe C~m~)

Figure 1. Changes of density after alpha doses corresponding to indicated Synroc geological disposal times for (a) Synroc stored at 200~ (b) Synroc stored at 30~ (c) Synroc containing additional Na processing impurities stored at 30~ The change of slope in curve (c) is attributed to microcracking which however can be eliminated by adjustment of the Synroc composition to allow for the Na content[4]. Studies of natural zirconolites are also in progress to try to correlate the structural radiation damage with the content of the uranium and thorium series isotopes, and to evaluate the effect of the damage on the release of actinides in Synroc. Figure 2 shows a high-resolution micrograph obtained from a zirconolite which has sustained a small ~mount of damage. A further interface problem arising from radiation damage derives from the fact that the different phases in Synroc contain different concentrations of alpha-emitting ions, have intrinsically different responses to alpha decays, and are generally anisotropic. Even a non-cubic polycrystalline single-phase material such as BeO displays anisotropic expansions and this produces severe intergranular stresses[5]. Abundant studies show that rare earths and uranium in Synroc are strongly partitioned into perovskite and zirconolite[6]. On the basis of ionic sizes and general chemistry, rare earths are good simulants for trivalent actinides and (tetravalent) uranium is a good simulant for tetravalent actinides. Therefore perovskite and zirconolite should be the principal hosts for actinides in Synroc. Experimentally, in ~Cm-doped Synroc, zirconolite and perovskite display much more structural damage than the other phases[7]. The propensity of the irradiated polycrystalline ceramic to display

436 intergranular radiation d~mage will depend on such factors as grain size, overall radiation fluence, grain boundary adhesion etc.

Figure 2. High resolution transmission electron micrograph of zirconolite from Kaiserstuhl, Germany. The 16 Myr old specimen contains 4 wt% ThO2 and 1.8 wt% of UO2 and has accumulated a fluence of 1.5 x 10TM alphaa/mg, equivalent to 0.1 displacement per atom. Mottled diffvaction contrast and local areas of damage (arrows) are characteristic of the early stages of the radiation-induced transition from the crystalline to the ~morphous state. 3. INTERGRANULAR FII~MS Synroc which contains small quantities of Si is observed to contain glassy, Sirich grain boundary phases[8]. Although Si is not a component of the major titanate phases comprising Synroc, the appearance of Si-rich phases is

437 surprising since perovskite can react with SiO2 to form sphene (CaTiSiOs). Perhaps the answer lies in SiO~ reacting with a little alkali (Cs fission products or Na contomination) early in the processing cycle to form metastable glassy phases which wet the grain boundaries of the later-forming phases, but are slow to react with them. Longer hot-pressing times may well remove the glassy phases. These glassy phases give rise to potential difficulties in several ways: (a) Sirich glass is considerably more leachable than the crystalline titanates; (b) when the glass is preferentially leached, the titanate grains originally separated by the glass are decoupled, so the total surface area of material exposed to the water is increased; (c) effect (b) may be greatly amplified if the lattice stresses due to differential radiation expansion of the different phases cause microcracking (see above). 4. LEACHING AT THE WATER-SYNROC INTERFACE The leaching of Synroc is a very slow process at ordinary temperatures, say 70-200~ and corresponds to the passage of the order of an atomic monolayer per day into solution, with a tendency for back-deposition of ions onto the Synroc surface, depending on the solubility of a given species and its potential for being sorbed onto the altered surface. Thus it is a difficult process to study experimentally. To perform experiments where the concentration of the leached species builds up in solution to easily-measurable values, high specific surface areas, small amounts of aqueous solution and long dissolution times are necessary. Even then, it is difficult to be sure that the species concentration is not due to a very small-region of poorly-mixed material which may not have the general or "average" microstructure. Other possibilities of error include the presence of very small fragments, adhering weakly to polished surfaces, as well as glassy phases (see above). Thus in addition to solution analyses we are currently leaning towards the use of surface techniques, such as secondary ion mass spectroscopy, nuclear techniques utilising energetic recoil analyses, and transmission electron microscopy to obtain results which are representative of the solid as a whole, rather than its most-leachable regions. Aqueous durability tests in the 70-150~ temperature range show by transmission electron microscopy that some limited dissolution of surficial perovskite occurs and that there is formation of thin amorphous Ti-O films at the lower temperatures and thin crusts of small (<0.1 pro) anatase crystals at the higher temperatures[9-11]). These Ti-rich secondary phases form relatively quickly, i.e. in 28 days at 70~ Ti-O films are seen and anatase is observed after only 1 day at 150~ Scanning electron microscopy of transverse sections of leached Synroc showed that after 532 days at 150~ surficial perovskite was leached to only a depth of 0.2 urn, and zirconolite showed no signs of attack.

438 Myhra et al[12] stated that "the main reaction with the lattice of the oxide phases is base-catalysed hydrolysis breaking the M-O-M bonds to give a hydroxylated-MOH structure" and that "cations ... released in this reaction must diffuse through either the hydrolysed surface layer, recrystallised products or the intergranular films". Pham et al[13] conducted dissolution tests on single-phase perovskite and found that even a 10 nm-thick hydroxylated layer could inhibit diffusion sufficiently to slow down Ca 2+ release to immeasurably small rates. Ph~m[14] also showed that concentrations of 5 and 20 ppm of Ca 2§ in the water were sufficient to prevent any nett disssolution of perovskite at 70 and 150~ respectively. At least under near-neutral conditions, it appears that the long-term dissolution rate of Synroe is controlled by the dissolution rate of its least durable phase (perovskite) which is present in sufficient quantity to form connected paths throughout the bulk of the material. Some work has been done at Ansto on trying to reduce the concentration of perovskite below the percolation limit, estimated to be in the 10-15% range. This can be done[15], but of course short-term (real-time) leach rates are insensitive to percolation, when only fractions of surface perovskite grains are leached, and the leach rate decrease due to the reduction in perovskite would only become apparent at very long leaching times. The preparation of individual Synroc phases is difficult. It is an experimental fact that hot-pressing a stoichiometric alkoxide-derived precursor under Synroc processing conditions (1200~ not yield entirely single-phase or theoretically-dense material, so it is necessary to use higher temperatures and pressures to try to obtain the ideal dense single-phases. High-temperature sintering(1500~ for some days) gives essentially singlephase material, but such material is not fully dense. To date, no experimental efforts have been 100% successful. However the use of alumina dies will allow hot-pressing at higher temperatures and enable redox control and it is possible to buffer the compositions to prevent the co-formation of unwanted relatively leachable phases. This will enable a comprehensive effort to study the effects of surface area-to-volume ratio, water composition, pH, Eh, and temperature on Synroe and its constituent phases. 5. CONCLUSIONS AND FINAL REMARKS Synroc-C is a well-characterised ceromic which on the basis of 15 years of work, mainly at the laboratory scale, but also at a scale of contained (inactive) monoliths weighing tens of kg., appears eminently suitable for disposal of highlevel nuclear waste. Here we have described several phenomena - radiation damage, intergranular films, and aqueous leaching - which can be discussed as interface effects. These are by no means the only possible such features; for example, some future studies of Synroc will address complex defect structures formed from partly ordered solute distributions within nominally single phase

439 grains. The essential features of radiation damage are a lattice expansion in the affected regions and possible enhanced leachability due to the weakening of the atomic bonding in the damaged region and the stress near the boundary of the damaged region. Intergranular films, which are probably metastable, are undesirable for Synroc because they act as less leach-resistant regions. Leaching at the solid-groundwater is the key property governing the acceptability of Synroc as a vehicle for the immobilisation of HLW; it occurs very slowly, and is a very complex interaction as it depends on pH, Eh, temperature, ionic strength etc. 6. R E F E R E N C E S

1 A.E. Ringwood, S.E. Kesson, N.G. Ware, W.D. Hibberson and A. Major, Nature, 278, 219 (1979). 2 A.E. Ringwood, S.E. Kesson, K.D. Reeve, D.M. Levins and E.J. Ramm, in Radioactive Waste Forms for the Future, Ed. W. Lutze, Elsevier, pp. 233-334 (1988). 3 W.H. Huang and R.M. Walker, Science 155, 1103 (1967). 4 tLD. Reeve, E.R. Vance, I~P. Hart, K.L. Smith, G.R. Lumpkin and D.J. Mercer, Proc. Austceram '92, Ed. M.J. Bannister, p. 1014 (1992). 5 B.S. Hickman, T.M. Sabine and R.A. Coyle, J. Nucl. Mater., 6, 190 (1962). 6 I~L. Smith, G.R. Lumpkin, M.G. Blackford and R.A. Day, Proc. Materials Research Society Meeting, 1992, Boston, U.S.A. 7 H. Mitamura, S. Matsumoto, I~P. Hart, T. Miyazaki, E.R. Vance, Y. Tamura, Y. Togashi and T.J. White, J. Amer. Ceram. Soc., 75, 392 (1992). 8 J.A. Cooper, D.R. Cousens, J.A. Hanna, R.A. Lewis, S. Myhra, R.L. Segall, R.St.C. Smart, P.S. Turner and T. White, J. Amer. Ceram. Soc., 69, 347 (1986). 9 G.R. Lumpkin, K.L. Smith and M.G. Blackford, J. Mater. Res. 6, 2218 (1991). 10 ILL. Smith, G.R. Lumpkin, M.G. Blackford, R.A., Day and K.P. Hart, J. Nucl. Mater., 190, 287 (1992). 11 ILL. Smith, K.P. Hart, G.R. Lumpkin, P. McGlinn, J. Bartlett, P. Lain and M.G. Blackford, Materials Research Society Symposium Proceedings, Pittsburgh, PA, U.S.A., 212, 167 (1991). 12 S. Myhra, D.K. Pham, R.StC. Smart and P.S. Turner, in Science of Ceramic Interfaces, Ed. J. Nowotny, Elsevier, p. 569 (1991). 13 D.K. Pham, F.B. Neall, S. Myhra, R.St.C. Smart and P.S. Turner, in Scientific Basis for Nuclear Waste Management XII, Ed. W. Lutze and R.C. Ewing, Materials Research Society, Pittsburgh, PA, U.S.A., p. 231 (1989). 14 D.K. Pham, Ph.D. Thesis, Griffith University, (1992). 15 E.R. Vance, K.L. Smith, G.J. Thorogood, B.D. Begg, S.S. Moricca, P.J. Angel, M.W.A. Stewart, M.G. Blackford and C.J. Ball, in Scientific Basis for Nuclear Waste Management XV, Ed. C.G. Sombret, Materials Research Society, Pittsburgh, PA, U.S.A., p.235 (1992).

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Science of Ceramic Interfaces II J. Nowotny (Editor) 9 1994 Elsevier Science B.V. All rights reserved.

441

Chemical and polar nanodomain fluctuations in relaxor-type lead scandium tantalate L.A. Bursill, J.L. Peng and H. Qian School of Physics, The University of Melbourne Parkville, 3052, Vic., Australia

ABSTRACT Polar and chemical nanocrystalline doma:in textures have been observed directly, using high-resolution transmission electron microscopy (HRTEM), for lead scandium tantalate (PST). This is a relaxor-type electronic or functional ceramic, derived from the perovskite structure type. Atomic structure images having 0.17nm resolution revealed directly the Pb, Ta and Sc atom positions for <101> zone axis projections. Careful comparison of experimental and computer-simulated images allowed the Pb atom shifts to be measured with respect to those of Sc and Ta atoms, yielding 0.035nm, directed along one of the eight <111> pseudocubic directions. The local structure is trigonal, space group R3. Local variations of polarization vector were mapped directly off the HRTEM images in favourable cases, where the crystal tapered to a very thin edge. High-resolution (0.4am) dark-field images were used to locate polar domain fluctuations occurring with length scales of 1-5nm and time scales of 0.04-1s. The images were recorded using video techniques. Thus it was possible to establish that spontaneous polar domain fluctuations occurred on the same nanometer scales as the smallest of the chemical domains for disordered PST, which has short-range order due to Sc and Ta segregation onto alternate (111) planes. The chemical domain textures have been modelled using Monte Carlo (MCS) and next-nearest-neighbour Ising (NNNI) models. These results provide a statistical physics framework for developing analytical theories for the dielectric response of relaxor-type ceramics. Thus the frequency and temperature variation of the permittivity are due essentially to dipolar-type fluctuations on nanometre scales. These simulations also help us understand and quantify the domain textures and domain wall configurations which may occur in complex perovskite-type oxides, as they appear in HRTEM images.

442 1. I N T R O D U C T I O N Polar and chemical domain textures on nanocrystalline length scales are believed responsible for the diffuse phase transitions and unusually high dielectric response exhibited by so-called relaxor-type ceramic materials (Setter and Cross, 1980a; Cross 1987). Thus the permittivity may reach values greater than 30,000 at 1Khz over a useful range of temperature for lead scandium tantalate (PST), depending on appropriate ceramic processing routes (see for example Setter and Cross, 1980a,b; Chu, Reaney and Setter 1992). The peak permittivity may vary from about 7500 to above 30000, depending on the specimen preparation. The unusual dielectric properties have been linked theoretically to chemical inhomogeneities presumed to exist due to chemical disorder of e.g. Sc and Ta on the perovskite (ABO3) B site. This was the classical theory due to Smolenskii and Isupov (1954); Smolenskii (1970)). Most recently there have been intensive efforts to apply spin glass theory, originally developed for magnetic glasses, to explain relaxor-like behaviour; Cross (1987) and Viehland et al 1990; 1991a,b,c,d) being the principal exponents. Note, however, that whatever theoretical approach is adopted, a point is reached where it is necessary to make assertions concerning the atomic structures underlying the extremely high density of polar and chemical domains and, of equal importance, the domain walls existing in this material. Following Cross (1987) we restrict use of the word relaxor to dielectrics which exhibit the following characterisitics: a significant frequency-dependence (dispersion) of the peak permittivity; a ferroelectric response under high fields at lower temperatures; but no sharp normal ferroelectric phase transition or Curie temperature. Relaxors are characterized by nanometre scale chemical heterogeneity and large values of the rms polarization (Prms) at temperatures well above that of the dielectric maximum. There may also be spectroscopic evidence that the fluctuations are dynamical, with a diffuse transition from nano- through micro- to macro- polar domains at low temperatures and high applied electric fields. In the zero-field-cooled state, a relaxational process typical of a classic dielectric relaxor, cf Debye theory of dipolar relaxation (Hench and West, 1991) is observed for temperatures well above the peak permittivity. There is also no readily observable macroscopic polarization or departure from Pm3m simple cubic symmetry. PST is referred to as relaxor-like since it may exhibit all of the above characteristics except that for higher degrees of chemical ordering a distinct ferroelectric transition may be detected. High-resolution transmission electron microscopic (HRTEM) studies of PST i.e. Pb(Sco.s, Tao.50)3 (Bursill, Peng, Chu, Setter and Reaney 1992) showed significant new results: including the observation of a hierarchy of nanodomain textures, observed by conventional and high-resolution dark-field imaging techniques. Furthermore the first direct measurement of the magnitude of the Pb atom shifts, measured with respect to those of Ta and Sc, was obtained, yielding magnitude 0.3)I along < 111 > directions of the

443 pseudocubic unit cell. In the present paper we review some recent results of our HRTEM imaging as well as our NNNI simulations. These lead to a better understanding of the dielectric response of PST, in terms of the extremely high density of chemical and polar domain walls which characterize this class of electronic material. The HRTEM results are first described in sections 2,3. The Pb, Ta and Sc atomic columns are next identified and the local direction and magnitude of the Pb atom shifts are measured, relative to the Ta and Sc atom positions, using image analysis techniques (section 4). It is then possible to discuss some of the local atomic structures of both chemical and polar domain walls, using information from the experimental images. Polar domain fluctuations are then described (section 5) and the NNNI computer simulations described in section 6. The results are then discussed with reference to recent theoretical approaches to the prediction of relaxor-type dielectric response (section 7). 2. E X P E R I M E N T A L Very well-characterized, reproducible ceramic preparations of PST were provided by Prof. Nava Setter, Labo de Ceramique, EPFL, Lausanne, Suisse. These were made by Chu Fan, who also measured the dielectric response. Ion- thinned sections suitable for HRTEM were provided by Dr. Ian Reaney. Some ceramic specimens were also thinned by fracture (Bursill and McLaren, 1965). These were most valuable for HRTEM studies since there was no ion-beam induced damage. Some electron-induced damage nevertheless occurred, which had to be taken care of, to avoid artifactual conclusions, as mentioned in the results section. Two specimens are discussed below, referred to as PST-o and PST-d. Both were classified as containing sufficient PbO for the perovskite phase to be fully stoichiometric (see Chu, Reaney and Setter 1992). The long-range chemical order parameter is defined as

,2__ (II11/I2oo)/(II11/12OO)s=1

(1)

where I111 and I200 refer to powder X-ray diffraction intensities for the pseudocubic 111 and 200 Bragg beams. The measured values were s = 0.93 and 0.10 for pst-o and pst-d respectively. These two specimens represent the most ordered and the most disordered specimens available to us. Thin edges of these specimens were examined at the University of Melbourne, using a JEOL-4000EX electron microscope operating at 400keV. The ultrahigh resolution pole-pieces had Cs = 0.94mm. The three principal zone axes <100>, <110> and <111> of PST were accessible using a + / - 20~c double-tilt top-entry goniometer. All of the results reported below refer to <110> zone axis projections. The specimen height was carefully adjusted to an optimum focussing current, when the

444 objective lens astigmatism, as well as the optical alignment parameters could be set precisely against calibrated values. The adjustment of the desired crystal orientation was achieved using a Gatan TV camera and image-intensifier. A video cassette recorder was used to record dynamic events, usually for dark-field imaging conditions. Special attention was paid to minimizing the effects of mechanical vibrations and acoustic interference, enabling point-to-point resolutions of better than 0.17nm to be achieved at the Scherzer defocus image condition. Computer-simulated images were obtained using Melbourne University Multislice (MUM) software. 3. H R T E M R E S U L T S 3.1 Introduction to the technique

A parallel beam of high energy electrons is passed through an extremely thin section less than say 50nm thick. The transmitted diffracted beams are collected by the objective lens and finally used to form a magnified image of the atomic structure. Provided the specimen is thin enough and certain electron optical conditions are met then it may be possible to directly image the atomic columns of atoms lying parallel to the incident beam. Thus a two-dimensional projection of the atomic structure along a principal zone axis of a crystal is the primary result. Complications arise if the crystal becomes too thick (greater than 10nm say), if the specimen and/or instrumental alignments are not precise (better than 0.1mRad say), if the specimen moves more than 0.1nm during the recording or becomes damaged in the electron beam. Definitive image analysis should be made using image-matching techniques based on computer simulations. Reference may be made to Goodman and Moodie (1974) for the theory and Cowley (1989) for reviews of both experiment and computing. 3.2 H R T E M results Figures la,b show HRTEM images of pst-o and pst-d viewed along the [li0] projection. These two images were obtained for essentially identical imaging conditions. The ordered specimen pst-o gives apparently a perfect structure image across the field of view. For pst-d however there is already a high degree of nanoscale heterogeneity, indicated by sharp changes in contrast and fine detail of the images . It is shown below, by use of computer-simulated image analysis, that these should correspond to nanocrystalline regions associated with chemically ordered and also polar domains. These features will emerge as the analysis proceeds below.

445

Fig.la,b HRTEM images of pst-o and pst-d for the [liO] projection. Note the heterogeneous nature of b compared to a.

446 4. A N A L Y S I S O F H I G H R E S O L U T I O N I M A G E S 4.1 Introduction Principles for analysis of the HRTEM images and the corresponding nanodomain textures were first established by careful comparison of computer simulated images with selected areas of the experimental images. We start with the known structure for the chemically-ordered phase, based on space group Fm3m. Then we determine by imagematching the values of the relevant electron optical parameters, e.g. crystal thickness and objective lens defocus values, for the ordered structure image. According to the weak phase object approximation for a very thin crystal (Cowley, 1989) the images most resemble the crystal structure when the objective lens current is set to the Scherzer defocus. Then the regions of darker intensity in the images correspond to the projection of columns of atoms in the structure. In the case of PST we then establish that the variation of intensity corresponding to columns of different types of atoms may carry a one-to-one correspondence with the different species of atom in the ordered structure. Furthermore it then becomes possible to measure directly the Pb atom shifts with respect to the Ta and Sc atoms. Finally the atomic structures of some polar domain walls are discussed.

4.2 Computing Facility Selected images were digitized from high-quality prints of electron micrographs or directly from negatives using a Wild-Leitz macroscope fitted with a television camera. A 64 level grey scale with 512x512 pixels was used to input images to an IBM RISC 6000 computer via a PCVISION Frame Grabber. It was found convenient to obtain the Fourier transforms (FT) and power spectra (PS), as well as output the digitized image, using SEMPER image-processing programs provided by SYNTRONICS (Cambridge, UK) Images were output using a QMS 600 dpi laser-printer with 64 grey levels.

4.3 Computer Simulation Techniques The calculations were based on the physical optics approach to electron diffraction and imaging (Goodman and Moodie, 1974). All the nonlinear N-beam dynamic scattering, as well as lens aberrations and Fresnel propagation effects of the objective lens were included. Periodic continuation methods, as developed for atomic resolution images (MacLagen, Bursill and Spargo, 1977) were obtained using MUM (Melbourne University Multislice) software.

447 4.4 S t r u c t u r a l b a c k g r o u n d PST has the perovskite ABO3 structure shown in Fig.2a, where the corner-shared BO6 octahedra enclose the larger A cation. The chemically-disordered simple cubic Pm3m structure (a= 4A) is usually assumed to have random occupancy of the B sites by Ta and Sc atoms. This is the ideal pst-d structure, as shown in Fig.2b for the <110> projection. If the Ta and Sc atoms order on alternate {111} planes the unit cell is doubled (a'=8/~) when the space group becomes Fm3m. This is the ideal pst-o structure, as shown in Fig.2c. At liquid nitrogen temperatures the ordered structure was reported to be trigonal (Groves, 1985). However there have not been any definitive structural studies reported for ordered or disordered PST. Note that the complex polar and chemical domain textures are not readily analyzed using Rietveld powder diffraction techniques. Details of the crystallographic data used in the present analysis are given in Peng and Bursill (1993). The space group was taken as R3, with hexagonal cell parameters a = 0.575 nm, c = 1.410 nm. Initially ideal atom coordinates transformed from the Fm3m ordered PST structure were obtained. Later a series of Pb, Ta and Sc atom shifts, consistent with R3m space group symmetry were deduced for the image simulation, as detailed by Peng and Bursill (1993). The effects of oxygen shifts were also investigated by applying shifts antiparallel to the best-fit cation shifts and having up to one-third the magnitude of the Pb shifts. There appeared to be no detectable effects of oxygen shifts on the set of images under discussion. 4.5 C h e m i c a l m i c r o d o m a i n s t r u c t u r e 4.5.1 S t r u c t u r e image of o r d e r e d m i c r o d o m a i n s An enlarged HRTEM image selected from the very thin edge of Fig.la is shown in Fig.3. Inset (a) is the power spectrum of the selected image; (b) is a computersimulated image match and (c) is the corresponding power spectrum (equivalent to the Fraunhofer diffraction pattern) of inset (b). Under the Scherzer defocus condition, the image contains three distinct levels of black spot intensities (inset b) which have one-toone correspondences to Pb (z=82), Ta (z=73) and Sc (z=21) atoms respectively. It should be possible therefore to see directly whether Ta and Sc atoms order on alternate {111} planes. In order to investigate this further the image intensity profile along adjacent [001] atomic rows was scanned for both the digitized experimental image and the corresponding computer-simulated image. The results are given in Fig.4a,b for the simulated and experimental images respectively. In Fig.4a the Pb atom row showed relatively uniform intensity whereas the (Ta,Sc) atomic rows show two levels of intensity, with the darker

448

Fig.3 Enlargement from thin edge of Fig.la. Inset (a) is the power spectrum. (b) is a computer-simulated image match (c) is the corresponding power spectrum of inset (b).

449

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451 level corresponding to Ta. For Fig.4b we note first of all that there are shifts of the Pb atom positions which were not evident for the simulated image Fig.4a. This is because the simulation was made for an ideal perovskite unit cell. We return to discuss shifts below. Note that there is of course significant noise for the experimental images. Background intensity variations also occur. Nevertheless the experimental atom row profiles do show the expected variation of intensity levels, with approximately equal density for the Pb atoms but alternating darker and lighter density for the Ta,Sc sites, in good agreement with the simulated profiles (Fig.4a). This result clearly demonstrates that Fig.3 represents a structure image of an ordered chemical microdomain. This is the first direct evidence for the chemical ordering of Ta and Sc in B sites of a perovskite-type structure. Some variations of intensity at Pb atom positions are visible in Figs.3,4b. These are presumeably due to some deficit of Pb, which may occur during preparation of this material (Chu, Reaney and Setter, 1992). There may also be Pb-loss during the HRTEM observations. Further study of Pb deficit related defects including Pb or PbO vacancies, platelet defects (loops), stacking faults, helical dislocations, and voids facetted on (111) and (110) planes, for a specially-prepared Pb-deficient sample, will be discussed elsewhere (Peng, Bursill, Reaney, Chu and Setter, 1993). It remains to determine the magnitudes and directions of the Pb atom shifts evident in Fig.3. Fig.5a shows a further enlargement of Fig.3, where the dark spots may be identified with the Pb, Ta and Sc atom positions. The lower diagram (Fig.5b) identifies the atom columns, as verified by computer simulated images. Simulated images were obtained corresponding to Pb shifts of 0.0, 0.2, 0.4, and 0.6/~, along a <111> direction (see Peng and Bursill, 1993 for full details of the image-matching procedure and results). Note that in this case the Pb atom shifts lie parallel to a [111] direction i.e. perpendicular to the projection axis [10-1]. Note also that shifts of 0.015nm were assumed for both Ta and Sc atoms. It was not possible to match the experimental images for Pb shifts of less than 0.05nm. Note that there may be additional antiparallel shifts of the oxygen atoms, which are not imaged here. The result was that the Pb shift appears to be about 0.035nm, measured with respect to the Ta,Sc sublattice.

4.5.2 Determination of local polarization characteristics If we assume R3 space group symmetry for polar nanodomains, as resulted from the analysis of Fig.5 above (again we refer to Peng and Bursill,1993 for full details of the space group analysis) then there are in principle eight such domains expected to form within the high-temperature Pm3m average structure. Fig.6 shows the corresponding atomic projections along a [10-1]pseudocubic direction. Note that four cases arise with polarization vector P8 lying in the plane of the figure; the remaining four cases, where P projects along [001] or [00-1] give only two distinct projections. The directions of P~ are

452

Fig.5 (a) Enlargement of Fig.4b showing Pb,Ta and Sc atom positions. Note Pb shifts along [111], magnitude approx. 0.03nm wrt Ta and Sc.(b) shows atom labelling for (a).

453 indicated on Fig.6. Some consideration of this figure reveals that there are essentially two types of atomic projection which should appear in the HRTEM images. The first case was analysed in Fig.5; it showed zig-zag lines parallel to [001], corresponding to the sequence Pb-Sc-Pb-Ta-Pb-, where the Pb shifts lie in the plane of the drawing along one of [111], [-1-1-1], [-1-111 or [1-1-11. The second case should give a long-short-long-short sequence for Pb-Sc-Pb-Ta-Pb- along [001], where the Pb shifts are along [11-11, [-1-11], [-111] or [1-1-1]. A search of the HRTEM structure images quickly revealed the two distinct classes of image. An example of the second case is given as Fig.7; the Pb shifts are indicated. Note that Fig.7 was obtained from the same negative as Fig.5; Thus both polarization vectors coexsist within a single chemically-ordered domain, implying the existence of polar domain walls. Fig. 6 also shows schematically some of the many possible polar domain walls (PDW) which may exist for [10-1] zone axis projections of pst-o. A systematic structural and diperiodic symmetry group analysis reveals that the PDW may be neutral, positively or negatively charged, depending on the relative directions of the polarization vectors connected across the different domain walls. Of course Fig.6 assumes that the PDW are exactly parallel to either (001) or (110) planes, which is likely to occur to a significant extent, but by no means exclusively. A search for HRTEM images showing atomic structure images of P DW is continuing.

4.5.3 Summary The local magnitude and direction of P, may be assigned directly by careful analysis of the HRTEM images in favorable cases; for example where we have a single chemicallyordered domain containing two or more non-overlapping polar domains. Such was the case for Figs. 5,7, which refer to pst-o. We may conclude by asserting that polar domains have been observed directly and the polarization vectors characterized for pst-o. At least a significant fraction of such polar domains are relatively static (frozen). A statistical analysis of polar domain size and direction has not been achieved as yet. For the case of pst-d exactly similar polar nanodomins are expected to occur, although the statictical distribution of sizes and directions will differ quantitatively from that for pst-o. Of course the dynamics of polarization fluctuations will also be quite different, depending critically on the local chemical disorder. Identification of P, for polar domains in chemically-disordered specimens such as pst-d is also possible, although the probability of finding non- overlapping polar domains in these cases is very restrictive. Dynamical fluctuations for the latter further complicate the analysis, which requires painstaking scrutiny of video recordings of the HRTEM images. We are presently developing on-line image processing techniques to meet the challenge of this problem.

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Fig.7 Enlargment of area from Fig.3, showing the second type of Pb shift.

455 5 DARK-FIELD IMAGE RESULTS 5.1 Hierarchy of domain textures

PST showed a variety of textures when imaged using different reflections. For example Fig.8, obtained using a < 111> Fm3m superlattice reflection, shows chemical antiphase domain walls, whose spacing varies from several tens of nm down to 1-5 nm. This was unusual for ordered PST, where the antiphase boundaries had, for the most part, spacings greater than 50 nm (see e.g. top of Fig. 8). Note also that the extent of chemical order may vary markedly in the vicinity of grain boundaries, similar to the variation shown towards the bottom of Fig. 8.. However, for pst-d the chemical domain size was homogeneously 1-5 nm, giving exactly the same contrast as shown at the bottom of Fig. 8 (see also Fig. 2b of Bursill, Peng, Chu et al 1992). One may conclude that disordered PST differs from ordered PST essentially in the chemical domain size; for pst-d the chemical domain walls are only a few nm apart. Perhaps a random distribution of Ta and Sc does not occur in real specimens? (This is to some extent a semantic argument; random does not mean zero short-range order, it does imply zero long-range order). Use of subcell reflections for dark-field imaging allows the polar domain textures to be studied. Macroscopic ferroelectric domains, on the scale of 0.1 - 1 micron, may be found for pst-o if the specimen is cooled well below the Curie temperature (28 C). Reference may be made to Randall, Barber and Whatmore (1987) for examples. The case of pst-d follows in the next section. 5.2 Polar domain fluctuations in d i s o r d e r e d P S T

When we tried to image polar domains in disordered PST, using for example <111> Pm3m subcell dark-field reflecting conditions, we were surprised to see bright and dark patches of contast, which fluctuated continuously in intensity and position with frequency of 0.1-1 second. Fig.9 shows part of a typical sequence. The contrast mechanism may be explained as follows. The dark-field image condition (Fig.9a) is set initially for only a fraction of the eight possible polar domains having polarization vector along one of the < 111 > directions. If a set of polar domains changes polarization to another < 111> direction, or simply reverses sense, then the high intensity dark-field condition will certainly change locally and those domains will go out of contast; others may appear of course, provided they satisfy the high intensity dark-field condition. In the case of polarization reversal the contrast will change due to failure of Friedel's law F(hkl) = F(h,-k,-1) for non-centosymmetric crystals. It is most important to realize that such polar

456

Fig.8 Shows dark-field image for 1/2(111) superlattice reflection showing a distribution of chemical antiphase walls in pst-o; the spacings at the top are typical of pst-o, whereas the fine spacings at the bottom are typical of pst-d.

457

Fig.9 Sequence from a video recording, showing time sequence of polar nanodomain fluctuations in pst-d.

458 fluctuations occurred on the same scale as the smallest of the chemically ordered domains of pst-d, at least at about room temperature. Thus we conclude that polar fluctuations in disordered PST, which presumeably underlie the extremely high dielectric response of this specimen, are strongly correlated with the smallest of the chemically ordered domains. We assert that the peak of the dielectric response should correspond to flipping of polar domains having, statistically, the same dimensions as the chemical domains. Further discussion of this point must be left for later in this paper. Note that the dark-field images also showed some interesting artifactual contrast fluctuations, due to electron beam induced decomposition of the specimen. Thus after prolonged illumination of an area for say 10m or longer nanocrystals of about 5-20nm diameter were seen decorating the original crystals. These continuously changed shape and orientation, on time scales of about ls. After recognizing this effect, which is seen most clearly in HRTEM bright-field images, it was nevertheless clear that the initial fluctuations just described were a fundamental effect. Similar decomposition phenomena occurred for PZT (see Bursill and Peng, 1992; Fig.18) when the fluctuating crystallites were identified as ZrO2. The nanocrystalline particles have not been chemically analyzed for PST. Exactly similar phenomena occurred for PMN; in each case there is probably loss of PbO by decomposition and/or sublimation, leaving eg ScTaO4 in the case of PST. Further analysis of this decomposition phenomema will be published separately (Bursill and Peng, 1993). 6. M O N T E C A R L O A N D N E X T - N E A R E S T - N E I G H B O U R

OF C H E M I C A L M I C R O D O M A I N

SIMULATIONS

TEXTURES

6.1 I n t r o d u c t i o n

Relaxor-type behaviour was assumed by Smolenskii, 1954 to be due to the presence of microheterogenous compositional fluctuations within the crystal so that different volumes of crystal transform to the ferroelectric state at different Curie temperatures. Certainly PST crystals exhibit chemical disorder with respect to Sc and Ta cations on the B site of PST; which may be achieved by appropriate thermal annealing. The phase transition is certainly diffuse, occurring over a range of at least 150 degrees C and it is significantly sharpened by increasing the B-cation ordering (Setter and Cross, 1980a,b). Although these facts support the Smolenskii model the latter does not predict the frequency dependence or dispersion characteristic of relaxors. Cross (1987) has attempted to apply the spin-glass theory (also known as superparaelectric or polar-glass model), originally developed for magnetic glasses, to explain relaxor-like behavior (see Viehland et al 1990,1991a,b,c,d). It is believed that a partitioning on the nanometer scale into short range chemically ordered clusters (nanodomains)

459 sets the scale of the inhomogeneity which underlies the relaxor behavior; these nanometer scale domains are assumed to be dynamical in nature with the dipole moment thermally fluctuating between equivalent directions. However, as the temperature is reduced freezing of the nonapolar regions is required; this allows a frequency dependent response to be predicted. So far there have been no strictly microscopic theories of relaxors. Both Smolenskii and Cross et al use phenomonological Landau or mean-field methods. In this section, by means of Monte Carlo Simulation (MCS), we attempt to model atomistically the distribution of Ta and Sc over the B-site of perovskite. This simple approach is followed in the next section by a more sophisticated statistical treatment using the next-nearest-neighbour Ising model (NNNI). The MCS approach allows an atomistic model to be set up using an L x L x L array. Different degrees of order (specified by s, the long-range order parameter) are then simulated assuming a random distribution of Ta and Sc atoms over the B sites. No interatomic potentials or energetics are invoked and there is no relaxational diffusive process which could lead to an equilibrium state. It assumes essentially the the specimen was quenched instantaneously from a perfectly disordered (random) structure at a temperature close to the melting point. The NNNI approach starts with the same perfectly disordered state, but then allows a relaxational diffusive annealing process to occur, by introducing NN and NNN interactions and minimizing the energy in a series of steps, to be described in detail below. Such simulations have a dual purpose; firstly, to clarify the possible chemically-ordered configurations of Ta and Sc atoms and the inevitable chemical anti-phase boundary or microdomain textures and domain wall configurations. Secondly they provide an appropriate basis for the development of statistical microscopic theories of relaxor-like behaviour. The calculations were made using a three-dimensional simple cubic lattice. For convenient comparison with the HRTEM results the output from the simulations are represented as [110] projections. To simplify even further only slices two unit cells thick are presented in most cases. Note however that the statistical analysis and the underlying computer simulations were all made in three-dimensions. In some cases Fourier Transforms (FT) and/or Power Spectrum (PS) were also obtained. In this report we restrict the discussion to chemical domain structures. As a further step in the analysis we have superimposed polar nanodomain textures onto the chemically-ordered state (s = 0). These results will be published elsewhere (Qian et al, in preparation).

460 6.2 Structural background The structure of PST was shown in Fig.2a. The corner-shared BOe octahedra enclose the larger A cation. The chemically-disordered simple cubic Pm3m structure (a= 4/~) is usually assumed to have random occupancy of the B sites by Ta 5+ and Sc3+ cations (Fig.2b). If the Ta and Sc atoms order on alternate {111} planes the PST unit cell is doubled (a'=8/~) when the space group becomes Fm3m (refer Fig.2c). It is not necessary to go to trigonal space groups to model the Ta,Sc distribution. The MCS and NNNI simulations were made using the B sites of the Pm3m simple cubic 0.4rim cell. Oxygens were not included in the simulations. However Pb atoms were replaced for display of the simulated structures. 6.3 The order parameters Usually the long-range order parameter is determined by calculation from the ratio of the intergrated intensities of 111 (basis) and 200 (superlattice) reflections, see Eq.1 above. The average domain size may be estimated by measuring the broadening of the superlattice reflections with respect to basis peaks (Setter and Cross, 1980b). Of course these measurements may represent average results. The degree of order, however, may be characterized more fully using both long- and short-range order parameters (see e.g.A.A.Bokov and I.P.Rayevsky, 1989). If two types of atoms are present in 1:1 ratio the long-range order parameter is defined as s=2p-1, where p is an occupation probability of atoms on one of the ordered sublattice sites. The shortrange order parameter describes the configuration of atoms immediately surrounding a given atom type. It may be defined as a=2r/-1, where 7/is the probability of finding say a Ta atom as NN for a Sc site. Both parameters were calculated or, where necessary, set for all of the computer simulated structures. 6.4 Monte Carlo Simulation Results Figure 10 shows a statistically disordered structure of PST (s = 0) projected along the [1i0] direction for slice thickness of 2 • dli0. T a 5+ and Sc3+ occupy at random the B-sites of the perovskite structure. Hatched, large and small discs represent Pb, Ta, and Sc respectively. Close examination of this figure revealed seven distinct classes of clusters labelled A,B,C,D,E,F and G in Fig.10. Enlarged representations are shown as Fig. 11. It is important to notice that the translational symmetry is very strongly broken in such a structure. The probability of finding any one class of chemical defect is equal to the maximum cluster size < 40A. Most clusters are 10-20/~; the smallest cluster is

461

Fig.10 Monte Carlo simulation of pst-d (s = 0); note the different cluster types labelled A,B,C,D,E,F and G.

4/~. Long-range ordered structures evolve from this statistically random state, as will be shown in the next section (NNNI model). Three of the five clusters may be identified as follows. A has space group Fm3m with unit cell (a'=8/~). It is simply a cluster of the perfectly ordered Fm3m superlattice structure, which gives 1/2(111) superlattice reflections. B and C would have 1/2(110) and 1/2(001) superlattice reflections respectively with orthorhombic unit cells. Such structures have not been observed for PST. Note that pst-d, our most disordered preparation was consistent with space group Pm3m; it showed near zero intensity for all possible superlattice reflections. As may be expected only the A class clusters survive after application of the NNNI model, as shown in the next section. The above three clusters are uncharged, although Ta 5+ and Sc3+ ions have different charges, since the superlattice structures each contain Ta and Sc atoms in the stoichiometric ratio 1"1. The remaining clusters are Pb-Ta rich (D)i.e.(P2+Tah+O~-) 1+ and Pb-Sc rich (E)i.e. (Pb2+Sc3+O~-)~- Both are charged, positively for Pb-rich and negatively for Pb-rich structures respectively.

462

Fig.ll Detail of defects A,B,C,D,E,F and G, from Fig.10; see text for detailed description.

463 Careful examination of the boundaries between clusters revealed only two wall types differing from the clusters A,B,C,D and E discussed above. These are marked F and G in Fig.ll. Note that F and G are less than two unit cells thick. The non-stoichiometric Pbrich walls F, stoichiometry (p2+T.5+.~ ~3/4'-'*"1/4v3~'rf)2-)1/4+, '2+ are positively charged and the Pb-Sc rich walls G, stoichiometry/p2+Ta5+ ~,~,2+ 1"-}2-)1/4- , are negatively charged Overall the \-- _~..1/4,......3/4v3 net charge averaged over all of these walls must equal zero. Note that the volume occupied by walls appears surprisingly large. This will of course be much less in three dimensions; nevertheless it may still be very significant for understanding the relaxor response in such grossly disordered structures. 6.5 N N N I R e s u l t s 6.5.1 I n t r o d u c t i o n NNNI models are ubiquitous throughout the literature on the theory of magnetism (see e.g. Landau, 1971; Binder and Landau, 1980; Binder and Stauffer, 1985). These methods may be applied to electric dipolar systems by suitably formulating such problems so that most of the analysis carries over from magnetism. In a magnetic system, three types of long-range order can be reached at equilibrium; i.e., ferromagnetic, antiferromagnetic and two equivalent superantiferromagnetic orderings. These are modelled by choosing different combinations of the interatomic potentials. The critical temperature for the high to low temperature phase transition has been studied by several authors with different theoretical models. Results of computerized Monte Carlo calculations have been reported in the literature by Landau, 1971. The long-range "ordered" structure of PST has antiferromagnetic type structure since the Ta and Sc atoms order on alternate {111} planes over the B sites (ref. Fig.2c) so that the critical temperature chosen for the NNNI simulation has to be within an antiferromagnetic-type phase region. As for magnetic systems, the Hamiltonian of the next-nearest-neighbor Ising (NNNI) model is of the form

nn pairs

nnn pairs

i

where aiaj, aiak, o'i -- -t-1 according to the site occupied by an A or B atoms in binary solid solution (spin up or down in the magnet system). In this case we follow the description of atomic interactions for a binary solid solution. Let c~ = J,,,~/Jnn, i.e., the ratio of nearest-neighbour to next-nearest-neighbour interaction strengths. The types of long-range order and the critical temperature Tc(c~), both depend upon the values of Jn~ and c~.

464 For the 2-dimensional Ising model: if J,,,,,, = 0, i.e., for nearest-neighbour interactions the critical temperature To(0) is J,,,,/kTc(O) = -0.4407 and J,,,,n -fi 0 (Onsager, 1944). In the case of next-nearest-neighbour interactions, Tr is a function of a, see e.g. Binder and Shauffer, 1985. 6.5.2 A l g o r i t h m s and procedure The most commonly used algorithms are called single spin-flip and spin exchange; since the former may change the concentration of the A and B atoms of the solution (or the magnetization of a magnetic system), whereas the spin-exchange algorithm does not. Only the latter was acceptable for the present calculation. The Monte Carlo Simulation was calculated on finite L • 2 1 5 (L=40) simple cubic lattices with periodic boundary conditions. The simulations start with an arbitrary spin configuration, i.e., a random distribution between A and B atoms as follows: one lattice site i was occupied by atom A or atoms B and its interaction energy determined with atom in its neighbor j. A positive energy implied an unfavorable state and therefore the atom A and B was exchanged (A --. B, B--. A). Otherwise the exchange probability was calculated and compared with a random number selected between 0 and 1. If the exchange probability was greater that that number, the atom was exchanged. In any case the calculation continued to the next atom. The procedure was repeated until one pass had been completed in a regular (typewriter) or a random fashion through the entire lattice; i.e., Monte Carlo Steps (MCS). Although there was finally no difference between these two procedures the random sequence of visiting sites is more realistic although it is much slower than the regular procedure. Typical atomic slices were output after an appropriate number of Monte Carlo steps; then the domain structures were analyzed, the degrees of order were calculated and the power spectrum were determined at each step. It is possible to show the domain textures clearly in three-dimension space by making use of color graphics. For a two-dimensional representation a multislice technique was used to cut MCS structures normal to the [li0] direction and obtain a two-dimensional slice. Microdomain textures in two (one unit cell thick) or three (two unit cells) slices could then be presented and compared with HRTEM micrographs. The Pb, Ta and Sc cations were finally represented using discs of diameter chosen according to their atomic numbers (Pb=82 > Ta=73 > Sc=21), so as to simulate very approximately the HRTEM structure images.

465 6.5.3 The evolution of chemical m i c r o d o m a i n t e x t u r e s Figs.12(a-1) are NNNI model simulations as follows. The simulation started with a completely random initial configuration (a). The remaining sequence shows the time sequence representing the evolution of increasingly ordered microdomain textures, simulating the annealing of PST as it approaches equilibrium. Finally the Fm3m state (A type in Fig.11) predominates. The order parameters were calculated for both (s) and (cr) for each configuration. Insets show some typical power spectra; note that the intensity of the 1/2(111) superlattice reflections increases from zero for the fully disordered structure towards a maximum value for the fully ordered structure. In Fig.12 type A chemically ordered microdomain textures are emphasized; ordered regions are shown darker whereas lighter regions represent the remaining chemical disorder. The evolution of the local structures A to G may be summarized as follows. Chemically ordered domains nucleate at A type clusters, which grow quickly for the first 4 MC steps after which ordered areas continue to expand. This result is consistent with the observation that widely-spaced chemical domain walls appeared for the pst-o specimen (top of Fig.8), as expected for s = 0.93, whereas for pst-d (s = 0.10) the domain wall density remained extremely high (cf. bottom of Fig.8). The domain size rapidly increases from nano- to micro- to macro- as the number of MCS increases, matching the experimental results of Setter and Cross (1980b; Table 1). There was no growth of the other types of structures. The sequence of disappearance was Pb-Ta (or Pb-Sc, B or C type ) and then F or G type walls; the former two defect types have higher electrostatic energy than the latter. However, some higher energy defects may continue to persist for non-equilibrium states; e.g. Pb-Ta and Pb-Sc rich regions can be found in Fig.12d (labeled D and E, inset). Power spectra from thin edges of the pst-d specimen (Fig.lb) indicate non zero short-range ordering even for the most disordered specimens, which is consistent with the MCS calculation. It is clear that fast-quenching does retain the high temperature state to some significant extent. As the equilibrium state is approached (see the enlargement of Fig.12k in Fig.13a) the structure recovers translational symmetry with widely-spaced antiphase boundaries or chemical domain walls (CDW); these are nothing more than the original F or G type defect structures. Note that kinks (arrowed) occur along the chemical domain walls (CDW). They have D and E type Ta or Sc rich structures as mentioned above. When the short range order parameter reached 0.948 (s = 0.88) the antiphase boundaries almost disappear but those charged small defects remain (see enlargement of Fig.121 shown in Fig.13b). Note that these may be described as topological defects which remain after a closed loop of CDW shrinks to a point or line. These persist indefinitely in the NNNI simulations unless the algorithms are modified to specifically rule them out. Although these charged chemical defects would appear to have prohibitively high electrostatic energies we are intrigued by

466

Fig.12(a-1) Shows the evolution of chemical ordering in pst after NNNI model simulation Note rapid loss of B and C type clusters and development of chemic al antiphase walls. Note also persistence of D and E type small charged defects.

467

Fig.13 (a) Enlargement of Fig.12k, showing detail of chemical antiphase walls. (b) Enlargement of Fig.121 showing small charged defects.

468 the possibility that they may nevertheless persist in the most rapidly quenched specimens. If so, then they would certainly be expected to contribute to the dielectric response of relaxor-like crystals, e.g. by acting as pinning centres for polar walls. Further studies of these small charged defects are being developed; this phenonema is even more interesting for PMN (lead magnesium niobate), the arch-type relaxor, where the Mg/Nb ratio is 1/2. Although CDW were readily located in dark-field images using diffraction contrast imaging conditions, the corresponding atomic structures have not yet been resolved directly images using HRTEM. A basic problem here is to find segments of CDW's at the thin edges of the specimens and which are still parallel to the incident electron beam. We expect to overcome this essentially statistical sampling problem. Computer simulation and image matching using models generated by NNNI and MCS may be necessary to finally identify the disordered structures in pst-d. Note that the antiphase boundaries lead to lower s relative to (a) values. In an extreme case where two perfectly ordered domains are separated by a single anntiphase boundary the long range order parameter is reduced to zero. 7. D I S C U S S I O N

7.1 Microdomain behavior at the ferroelectric/paraelectric phase transition of BaTiO3 Direct evidence for the occurrence of microdomains was first achieved by the authors for BaTiO3 by in situ electron microscopy of the cubic to tetragonal phase transition (Bursill and Peng 1984). It was concluded that these microdomains contained a ferroelastically-distorted (tetragonal) transition structure. The microdomain size was about 40-60/~ diameter, there was no systematic variation of domain size during the phase transition and the microdomains probably disappeared suddenly rather than gradually. The overall periodicity and texture of the microdomains were observed, they clearly possessed a dynamic character on a nanoscale and were detectable using video techniques at 25 frames s -1. In this case the microdomains did not interact readily with the electronbeam. When a ferroelectric crystal is cooled below its Curie temperature, in the absence of external electrical and mechanical stresses, it generally breaks up into domains. The directions of polar axes in the neighboring domains are distributed over all the symmetrically equivalent vectors with equal probability, so retaining the original point symmetry overall. Those are referred as ferroelectric domains or polar domains. The size of polar domains depends on various factors of individual crystals such as defects, composition fluctuations and external and internal electric fields. It may range from macro- (macrodomain) to micro- (microdomain) scales. It was evidenced in the study of BaTiO3 that the ferroelec-

469 tric domain size tends to become very small at the crystal surfaces or edges, and that these small domains (approx. 50/~ dimensions) often tend to fluctuate and oscillate rapidly. 7.2 Polar D o m a i n Fluctuations in P S T

Chemical nanodomains have been observed earlier in partially ordered PST (Randall, Barber and Whatmore 1987). Similarly the same authors observed ferroelectric domains at the F E / P E transition of partially-ordered PST using dark-field TEM techniques. However the present study is the first to record and report thermal fluctuations of polar nanodomains for pst-d at about room temperature. All simple perovskite type compounds have sharp phase transitions (e.g., BaTiO3). Diffuse phase transitions are observed only in solid solutions such as (Srl/3TiO3-Bi2/3TiO3, (Ba,Sr)(Ta,Nb)206, and Ba(Wi,Sn)O3). They are ubiquitous in complex perovskite compounds such as the relaxor types PST and P MN. Consider an arbitrary small region (several ten unit cells, say _< 100A in size) in the MCS of PST as shown in Fig.10 and enlarged as Fig.ll; a Pb-rich cluster (E) situated at the center of the region may be surrounded by all the different types of structure clusters (A, B, C, D, E, F and G) which clearly leads to structural and chemical inhomogeneities. This may be referred to as a frozen structural glassy state (FSGS) in the sense that broken translational symmetry exists due to the gross chemical disorder. The polar domains associated with phase transitions may have different onset temperatures in association with different types of structure cluster. Diffuse phase transitions may then be expected due to the presence of those frozen structural fluctuations within different regions. Note that macroscopic inhomogeneities of structure are not responsible for diffuse phase transitions. This was realized for ordered PST; thus as the size of ordered microdomains increases the phase transition sharpens. Internal electric field inhomogeneities may be referred to as a frozen polar glassy state (FPGS) in the sense that there are electric fields which fluctuate at random over the symmetrically equivalent polarization directions; we refer here to spatial fluctuations, see Viehland et al (1990; 1991a,b,c,d) for detailed discussions on this point. Any theory of relaxor response must account for the existence of the different types of chemical defects elaborated above; both charged and neutral clusters, etc which may be expected to interact strongly with polar domain configurations and dynamics, both in the relaxor state characteristic of pst-d as well as the paraelectric/ferroelectric phase transition ofpst-o. Our in situ experiments reported above imply that in zero-field-cooled experiments on relaxors such as pst-d macroscopic ferroelectric domains do not occur; they appear to be replaced by small polar clusters; giving a glass-like structure with spontaneous polarization vectors statistically distributed over the eight available polarization states. (We assume there is rhombohedral local symmetry for PST in the relaxor state.) This is, at least to

470 some significant extent, a dynamical fluctuating system. The size distribution of polar nanodomains is at least similar to that of the frozen structural glassy state (FSGS). It remains to establish the correspondence, if any, between the FSGS and a FPGS. Whereas these two obviously overlap in space and are interdependent via electrostatic interactions, it is by no means certain there is a one-to-one spatial correspondence. Such nanocrystalline glass-like chemical and polar structures are expected to be isotropic to x-rays and other macroscopic experiments. In the field-cooling experiments, the spontaneous polarization vectors of polar glassy clusters become aligned and nanodomains grow into larger microdomains then macrodomains, imposing a field-constrained ferroelectric state, see e.g. the elegant results obtained for PMN by Ye and Schmid, 1993. Although polar nanodomain fluctuations have now been observed for electron beam experiments using HRTEM and HRDF images it remains an open question as to what extent the dynamical nature of the polar clusters involves dipole moments thermally fluctuating between equivalent polarization directions (polar-polar switching) and/or polar to non-polar fluctuations (polar/non-polar switching). It should be clear from this report that HRTEM bright- and dark-field imaging techniques have a vital and interesting role to play in exposing nanodomain structures and textures as well as spatial and to a limited extent temporal fluctuations. These techniques will become even more powerful in future, as they are combined with MCS and NNNI modelling of the disordered structures. Thus, as demonstrated in the present work, an atomistic statistical physical theoretical approach to the understanding of the structure-property relationships becomes directly accessible. ACKNOWLEDGMENTS This work was supported by the Australian Research Council. We are grateful for the use of the JEOL-4000EX ultrahigh resolution instrument provided at the University of Melbourne, known as the National Advanced Materials Analytical Centre or NAMAC. We are grateful for the support of Prof. Nava Setter for this work, both for her enthusiastic discussions and encouragement, as well as for providing the two specimens through Chu Fan and Ian Reaney. We also appreciate the editor of this volume for his enthusiastic support. REFERENCES

Binder, K. and Landau, D.P. Phys.Rev.B 21 (1980),1941-1962. Binder, K. and Stauffer, D."Applications of the Monte Carlo Method in Statistical

471 Physics" (K. Binder, ed.) 2"dEd. Springer-Verlag (1985)ppl-36. Bokov, A.A. and Rayevesky, I.P. Ferroelectrics 90 (1989) 125-133. Bursill, L.A. and McLaren, A.C, J.Appl.Phys. 36 (1965) 2084-2085. Bursill. L.A. and Peng Ju Lin, Nature (Lond) 311 (1984) 510-514. Bursill, L.A. and Peng Ju Lin, Key Engineering Materials 66/67 (1992) 421-460. Bursill, L.A., Peng JuLin, Fan,C., Setter, N. and Reaney, I., Ferroelectrics, (1992) in press. Bursill, L.A. and Peng, J.L. "Polar Fluctuations in Disordered PST", in press, 1993. Chu Fan, Reaney, I and Setter, N. private commun./manuscript in prep. (1992) Cowley, J.M. "High-Resolution Transmission Electron Microscopy" Oxford Univ. Press, 1989. Cross, L.E., Ferroelectrics, 76 (1987) 241-267. Goodman, P. and Moodie, A.F., Acta Crystallog., A30 (1974) 280-291. Groves, P.J. Phys.C. Solid State Phys. 18 (1985) L1073-1078. Hench, L.L. and West, J.K. "Principles of Electronic Ceramics", J.Wiley, NY (1990) Ch.6. Landau, D.P., J.Appl. Phys. 42 (1971) 1284-1285. MacLagen, D.S., Bursill, L.A. and Spargo, A.E.S. 35 (1977) 757-768. Onsager, L. Phys. Rev. 65 (1944) 117-121. Peng, JuLin and Bursill,L.A. " Electron Optical Analysis of Polar and Chemical Nanodomain Structures in PST", in prep. (1993). Peng, JuLin, Bursill, L.A., Fan, C., Reaney, I. and Setter, N. "HREM study of Pb-deficient PST", in prep. (1993). lxandall, C.A., Barber, D.J. and Whatmore, R.W.J.Micros. 145 (1987) 275-291. Setter, N. and Cross, L.E., J.Appl.Phys 51 (1980a) 4356-4360. Setter, N. and Cross, L.E., J.Mat.Sci. 15 (1980b) 2478-2482. Smolenskii, G.A. and Isupov, V.A., Soviet Journ. Tech. Physics 24 (1954) 1375-1379. Smolenskii, G.A, J.Phys.Soc.Jap. 28 (1970) 26-37 (supplement). Viehland, D., Wuttig, M. and Cross, L.E., Ferroelectrics 120 (1991) 71-77. Viehland, D., Jang, S.J. and Cross, L.E., Philos. Mag. B 64 (1991) 335-344. Viehland, D., Jang, S.J., Cross, L.E., and Wuttig, M., J.Appl.Phys. 6___99 (1991) 414-419. Viehland, D., Li, J.F., Jang, S.J., Cross, L.E., and Wuttig, M., Phys.Rev. B 43, 8316-8320. Viehland, D., Jang, S.J., Cross, L.E., and Wuttig, M., J.Appl.Phys., 68 (1990) 2916-2921. Ye, Z-U and Schmid, H. (1993) Ferroelectrics, in press.

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Science of Ceramic Interfaces II J. Nowomy (Editor) 9 1994 Elsevier Science B.V. All rights reserved.

473

Copper and nickel ultrathin films on metal-oxide crystal surfaces Preben J. Moller* Department of Chemistry, University of Copenhagen, DK-2100 Copenhagen 13, Denmark

Universitetsparken 5,

Abstract The synthesis and in situ characterization of ultrathin films of copper and nickel on single-crystalline surfaces of metal oxides, and in particular the development in the epitaxial and electronic structures of the metal particles and films during atomic layer epitaxial growth, are reviewed. We consider recent results for lowindex surfaces ofa-AleO s, CaO, a-Fe2Os, LaA1Os, MgO, NiO, SrTiOs, TiO2, ZnO and yttria-stabilized ZrO2, using low-energy electron diffraction and electron spectroscopies, and discuss the problem of electron-impact induced surface charging. Analysis by combined electron spectroscopy and photodesorption methods of the reactivity of the metal deposits is briefly discussed for the cases of CO exposure to Ni-deposited TiOe and to Cu-deposited ZnO crystal surfaces.

1. I N T R O D U C ~ O N Studies of single-crystal surfaces of the metal oxides are of fundamental importance in the understanding of a wide range of both basic and applied research, and the interfaces of these surfaces with metal particles and thin films are of decisive importance in many areas of chemistry and materials science: such as in heterogeneous catalysis, e.g. the synthesis of methanol or the oxidation of a range of hydrocarbons; in ceramics, e.g. joining, and the synthesis and reactivity of high-Tc superconductors and of electrical contacts on these, or the synthesis special layered sandwich constructions and composits; in micro-electronics, e.g. electronic-packaging systems and the formation of metal contacts to the oxide layers whose properties may be of semiconductor, insulator or even metal nature, or in synthesis and application of selective gas sensors; in metallurgy, e.g. metal corrosion resistance. In order to u n d e r s t a n d well and, in particular, to control and in some cases to improve the conditions of many of these processes it is fruitful to carry out studies on well-defined single crystal surfaces and their interfaces with the metal deposits. With well-defined surfaces we mean surfaces that reproducible in atomic scale is determined with regard to both geometric, i.e. two-dimension-ally crystallographic, and electronic structure. Not least due to its importance in many large-scale

474 industrial applications this type of research has gained a strong m o m e n t u m over the last decade. The structure and properties of the metal oxide surfaces have been discussed previously in a review [1], metals on oxides were reviewed in 1987 [2], and recently the bulk properties of the transition metal oxides have been reviewed [3]. In the present review we will discuss recent experimental results on geometric and electronic structures of clean MO, M02, M203 and MaMbO 3 metal-oxide crystal surfaces and of the adsorption of copper and nickel on a range of metal-oxide single crystal surfaces, and also give some examples from studies on the reactivity of these surfaces toward carbon monoxide. After a brief section on experimental methods (Sect.2) and a discussion of e-beam induced surfaces charging (Sect. 3) we consider the following six crystal classes: (i) rocksalt (Sect.4), (ii) rutile (Sect. 5), (iii) corundum (Sect. 6), (iv) perovskite (Sect. 7), (v) wurtzite (Sect. 8) and (vi) fluorite (Sect. 9).

2. EXPERIMENTAL METHODS The experimental methods that were used in experiments that we discuss here were all ultra-high-vacuum (UHV) based, i.e. at base pressures below 10 .7 Pa. The base pressure was usually 6x10 9 Pa. Low-(and very-low)-energy electron diffraction (LEED and VLEED, respectively) were used for the characterization of the surface geometrical structure. Very recently the first results were obtained by scanning tunnel microscopy, for the rutile surface, though. That method will give much wanted and new light on the surface geometry. There are severe difficulties, though, due to the fact that most metal oxides are hardly good conductors. The cleanness, stoichiometry [4] and the electronic structures in detail were elucidated by Auger (AES), electron and highresolution electron energy-loss (EELS and HREELS, respectively), laboratory (ultraviolet (UPS) and X-ray (XPS)) and synchrotron-radiation-based photoemission (PES) [5] and target current spectroscopy (TCS) [6], and finally, in some gas-reactivity measurements, temperature-programmed desorption (TPD) and (photo- or laser-induced) desorption mass spectrometry. The metal was deposited in submonolayer-steps at a rate of 1/~Jmin either from an electron beam evaporator or from a Knudsen cell. In all cases the thickness d of the deposited metal layers was monitored by an oscillating quartz crystal microbalance (QCM) and calibrated by AES. For all the above methods there is often much difficulty, as compared to metal crystals, in carrying out experiments on the metal oxides and on submonolayermetal-covered metal oxide surfaces, due to several reasons, the most serious being surface charging problems, particularly for the large-band-gab semiconductors and insulators. Quite often the threshold incident electron energies are low. The range of usable electron energies is quite restricted, therefore. Secondly, the concentration of defects at most of these surfaces is considerably larger t h a n for metals. These experimental difficulties have probably been major reasons for the later development of the surface chemistry for the oxide crystal surfaces. We have succeeded in obtaining AE and EEL spectra and LEED patterns for many large-

475 band-gap metal oxides by creating, by flash, a small (not AES-surfacestoichiometry-detectable) amount of oxygen vacancies in the surface, giving effective positively charged holes t h a t can receive the electrons and thus give sufficient conductivity to allow for the AES, EELS and LEED experiments. For most metal oxide surfaces it is necessary to heat the samples to high temperatures in order to obtain clean and ordered surfaces. The heat t r e a t m e n t may be carried out in air if one makes precautions not to expose the surface to moisture while mounting in the UHV (using flow of nitrogen). When under UHV the surfaces are flashed in s i t u to ~ ~ ~ T--rcleanness and surface crystallographic order using a 1 sharply focussed beam of light from an external Xe lamp. 3. E I , E ~ N - B E A M

,11,,, 0 (5031

015031 Cu(920) Cu (60)

~

201)

u

4

5

1

0

~

200

I

400

I

600

I

800

1

I

._W

INDUCED CHARGING

Let us start the discussion of the experimental inve stigatio ns of me tal-o n- me tal-oxide syste ms with an investigation of the troublesome charging-up behavior of many of the oxide surfaces upon incidence of an electron or light beam, the former yielding a negatively charged surface upon scattering, the latter a positively charged surface. Positive charges are fairly easily removed by electron flooding with a mild cloud of electrons from a 'flood gun' while the negative charges may not similarly be removed as easily (by positive ions) without damage. In some cases [7] negative charged may be removed by continuously using a second electron gun (operating at 1-2 keV) at glancing incidence and thereby changing the "secondary emission crossover", the point where the ratio of the secondary-electron-emission current to the incidentelectron-beam current passes one. We have found it possible to carry out electron diffraction and spectroscopy if the surfaces initially is flashed to high temperatures to create a few (oxygen) vacancies whereby it becomes possible for the electrons to tunnel away to a ground connector. In Figure 1 we demonstrate a familiar behavior of an AE spectrum upon irradiation with an electron beam [8]. Here the electron beam is incident perpendicularly onto the MgO(100) target surface at room temperature with an energy E, of 3.1 keV. The sample had been cleaved in air and then immediately, under

1000 1200 1400

Electron energy (eV)

Figure 1. AES from Cu/MgO(100). See text.

476 1500

/

A

. _ e~

1000

o/ !

. _ e--

/

/

0

/

!

/

=" 500

/

I

L_

/

t.,,o

. . . . .

A-----0/I

1

/ /"

/

.~o/. /

n

-AO/--'-OO/" -A I

2

Incident energy

,,d

I

3

(keY)

.A I

4

Z~b

5

Figure 2. Au~ger surface potential shift with Ep. A: clean MgO(100); [3:30 A Cu; o" 60 A Cu; o: isolated holder. nitrogen, mounted under UHV where it was flashed with the beam of light. The spectra were recorded when sufficient time (a few minutes) necessary to reach an equilibrium (saturated) potential shift had elapsed. First we note t h a t there is no charging problems for the clean surface (curve 1), but upon deposition of small amounts of Cu onto the surface we note t h a t the spectra are shifting towards higher energies (curves 2-5). We can thus follow a shift of the Cu(849eV) peak. At an average Cu thickness dcu of about 15 A (1/k - 0.1 nm), corresponding to about 4 monolayers (ML) of Cu, the Mg and 0 peaks disappeared. This observation agrees reasonably w e l l with the known value of the electron escape depth in Cu which is about 10 A. We have also found t h a t a threshold value of Ep exists above which the charges build up, measured as the potential shift of the spectra, and that no shift was found for Ep < .1 keV (Figure 2). On the other hand, it is necessary t h a t Ep is sufficiently large to produce enough emitted electrons from the surface. We have observed similar behavior on other metal-on-metal- oxide systems as well, such as Cu and Y on LaAI03 (100) [9]. An exponential relationship was found for the shift with de,. When the primary electron beam is aimed at a fixed particular point in the film, the shift is not very stable with beam exposure time; a dynamical behavior of the shift in potential is revealed although that change in the shift is small in comparison to the uncertainties of the data in Figure 2. As to the threshold in E, (Figure 2) we know t h a t electrons will penetrate deeper into the solid the higher Ep . We suggest t h a t the surface will be loaded to the critical threshold value, about 1.6 keV for MgO (and 1.4 keV for LaA10s [9]), consistent with results for many insulating materials [10]. Since the insulator MgO has an energy gap E~ of 7.8 eV, a work function of 3.14.4 eV and the vacuum level located in the gap, there are no states available to hold the arriving incident electrons which then are scattered from the surface,

477 thus explaining the lack of charge buildup on the clean MgO surface even for an Ep of 5 keV. We have explained charging-up results for metal deposition on metal oxides by a three step model [8]. (Step 1): Strong surface charging during ultrathin (a few tenths of ML coverage) deposition of Cu on MgO(100) and MgO(111) surfaces due to electron trapping centers in the interface (as explained by Cu 3d impurity levels located in the MgO band gap). (Step 2): Building up continuous Cu films which can accommodate local conduction bands. (Step 3): When reaching thicker layers (in the order of hundreds of/k and above), charging-up as for an isolated metallic target exposed to an incident electron beam.

4. COPPER AND NICKEL ON ROCKSALT METAL-OXIDE STRUCTURES

5

(Cu/0) ~4

~3

~R(Cu/ig) oo

-

L_

j

I---

o

1

0

0

I

5 10 Cu thickness (A)

Figure 3. Changes in Cu(LMM)/ O(KLL) and Cu(LMM)/Mg(KLL) AES intensity ratios with dcu.

i

15

The high-symmetry cubic rocksalt crystals have the advantage that they cleave very easily and often give cleavages of high quality, i.e. having only relatively few defects. The ideal surface is non-polar (charge neutral) and atomically fiat. Particularly the basal (100) plane of these cubic crystals are often close to ideal for 8surfaces studies. It is terminated at surface in a simple bulk truncation. It is preferable if cleavages can be carried out in the UHV but for the rocksalt crystals, particularly MgO and NiO, there is only little difference compared to surfaces cleaved in air. One should take care, however, to minimize exposure to air (interaction with water vapor ) and therefore make the cleavages in a N2 box. We have carried out synthesis and combined AES-EELS-XPSLEED characterization of ultrathin Cu or Ni films on four rocksalt-type crystals: The alkaline-earth d o oxide surfaces MgO(100) [8,11-13], M g O ( l l l ) [8] and CaO(100) [14,15] were exposed to Cu, and the transition-metal oxide surface NiO(100) to Ni [16]. The system

478 Cu/MgO(100) has also been studied previously by electron microscopy and diffraction [17-18], and recently by AES and HREELS [19]. 4.1. Cu r MgO(100) a n d MgO(111) Magnesium oxide is an insulator with a wide bandgap of 7.8 eV (bulk) [1] and 7.0 eV (surface) [20], and it has a lattice constant of 4.21 A. In section 3 above the preparation of the surfaces of this crystal is described. During deposition of submonolayer amounts of Cu at room t e m p e r a t u r e (RT) onto well-defined (1• (i.e. simple bulk-truncated) MgO(100) and MgO(111) surfaces [11], the originally sharp substrate LEED-pattern becomes gradually blurred until the substrate p a t t e r n is eliminated at a dcu of about 8/k. When dcu has reached about 15/k, a week Cu(100) superstructure p a t t e r n appears for the case of freshly air-cleaved MgO(100) and a, also weak, Cu(111) superstructure for the case of the fresh MgO(111) surface. The p a t t e r n s sharpened with increasing amounts of Cu. In Figure 3 is shown the gradual change in chemical composition, i. e. of the AES signal intensity ratios divided by their respective AES sensitivity factors, of the surface layer with the Cu layer thickness dcu. The growth behavior indicates a Stranski-Krastanov mechanism: two stages, a monolayer followed by 3dimensional islands. This growth mode was later confirmed by Conrad et al. [ 19]. Large spectral changes are observed by EELS during deposition of Cu particles. Strong resonance peaks at 2.6 eV for MgO(100) and at 2.2 eV for M g O ( l l l ) were both eliminated during deposition of submonolayer deposits of Cu, and two new peaks appeared 2.2 eV and 4.5 eV. We have shown [21] t h a t this most probably is due to binding of the first small Cu deposited particles to the oxygen ligands as ions, sticking to Mg 2§ vacancy sites causing the surface defect peak to disappear and two new resonance peaks which have character of cuprous oxide to appear due to Cu impurity-level related electronic transitions. The MgO(111) spectrum only shows a small difference in the 6.1 eV surface-state related peak in comparison to the spectrum of the clean MgO(100) surface. The intensity of the surface-defect peak is a little weaker and the surface-state related peak a little stronger. We also note t h a t the initial deposition has the same influence on the EEL structure as does the change from the (100) to the polar and Mg-richer (111) surface, supporting our claim t h a t the initial Cu deposits have a high possibility of interaction with Mg § sites, i. e. the Vs color centers (holes trapped at a metal ion vacancy, with 5 oxygen ligands in the case of the MgO surfaces). We should point out, though, t h a t the finding of the Vs being the physical origin of the electronic loss structures by no means implies t h a t Fs§ centers are non-existing on oxide surfaces. Already at average coverage near 5 / k the spectra show some characteristics of bulk copper energy-loss features, and at higher coverage, above 15 /k, corresponding to a few monolayers, the spectrum is a bulk copper one. Furthermore we found t h a t the epitaxial copper film after oxidation by CO agrees very well with t h a t from oxidized copper crystals. We have also investigated the effect of heat t r e a t m e n t and reoxydation. Figure 4 [12] thus demonstrate for the Cu/MgO(100) system the effect of oxidation of a previously flashed surface (curve 2), deposition of a very small a m o u n t (0.5/k ) of Cu (curve 3) and the result of heating the latter surface, with the Cu on (curve 4). We see t h a t the loss structure at 2.6 eV still remains after heat t r e a t m e n t in 10 .5 Torr oxygen for 2 h at 430 ~ discarding the loss as an origin of a Fs§ center.

479 Deposition of 0 . 5 / k on MgO(100) produces loss spectra at 2.2 and 4.2 eV, and removes the peak at 2.6 eV. To give further evidence for our above claim of oxygen-bonded copper deposits the latter surface was heated to 300 ~ (curve 4). The 2.2 eV peak was enhanced ( further heat treatment to 410~ for 15 rain did not change the spectrum further) and consequently leaves out the possibility that the peak has originated from Cu clusters which have electronic characteristics different from bulk Cu. We believe that the peak is the well-known loss structure of the semiconductor Cu20, corresponding to a threshold transition from the Cu 3d valence band to the Cu 3s conduction band, in agreement with the above

22.4

17.4

2.6

I A

e-

L. L_

...i I,,1.1 v

z

4 e-

em

I

4O

I

30

I

20

I

10

I

0

Energy loss (eV)

Figure 4. EELS (at Ep= 97 eV) from Cu/MgO(100). (1) clean, 400~ in UHV, (2) clean, 02 at 430~ (3) 0.5 ,s Cu on (2)-surface, (4) the (3)surface heated to 300~ for 15 rain.

480 observation that the spectrum after oxidation corresponds to oxidized Cu. The 4.2 eV peak may be due to the O 2p band to the Cu 3s levels. Furthermore, the above discussed electron trapping by copper deposits also supports the ionized-copperstate mechanism. These experimental results, supported for the (100) surface by theoretical calculation [22], clearly indicate that magnesium vacancies on the (1• 1)MgO(100) surface are trapping centers for the initial copper deposits, and it is hence reasonable to expect that these trapped copper ions are active sites for a subsequent nucleation, so that the copper atoms bind to the oxygen ions on the surface in a bridge configuration subsequently leading to a Cu(100) epitaxial growth through formation of clusters around the active center. Comparing to previous spectroscopic work on Cu(I) and Cu(II) oxides [23] we attribute the appearance of the 2.2 and 4.5 eV peaks in the 0.7/k Cu/MgO(100) EEL spectrum to Cu(II) states fitting into the Vs centers, and propose that the copper atoms forming the bridge configuration to the oxygen ions are in the Cu(I) state. This conclusion was confirmed by XPS measurements in which we used examination of the Modified Auger parameter d of copper (the sum of the 2ps~ binding energy and the Auger LMM kinetic energy), during the submonolayer growth [11], and also by Conrad et al. [19] using HREELS. For the (1• 1)MgO(111) surface we similarly find that the cation vacancies will be occupied by the copper bounded to three anions and thereby forming nucleation centers for Cu(111) epitaxial growth. The above proposed epitaxial growth was confirmed [11] by (somewhat diffuse) LEED patterns, indicating partial epitaxy. 4.2. N i cm MgO(100)

Ultrathin-fi]m deposition of nickel onto MgO crystal surfaces has been investigated by reflection high-energy diffraction (RHEED) [17] and by electron microscopy imaging (moir~ fringes and dark field images) and selected-area diffraction techniques [18]. A preferred epitaxial relationship was concluded [17] for both Cu and Ni deposits as Cu, Ni(110) [111] II MgO(100)[011]. 4.3.

Cu ~m

CaO(100)

The other alkaline-earth oxide crystal we will look at is calcium oxide. It is also an wide-bandgap insulator. It has a bandgap of 7.0 eV (bulk) [1] and 6.2 eV (surface) [20] (the bandgaps get smaller down through the group MgO, CaO, SrO, BaO) and a lattice constant of 4.81 A (the lattice constant increases in the group with the size of the metal ion). The sample was heated in air to 900~ before mounting in UHV, and after flashing with the light-beam a sharp (1• 1) LEED pattern from an AES clean CaO(100) is obtained whereupon EELS were carried out at Ep > 116 eV [14,15]. LEED was previously obtained for the clean CaO(100) surface, and 1% contraction between the outermost two layers and a vertical displacement of the Ca atoms with regard to the O atoms in the surface of less than 2 % were found [24]. Also EELS was obtained for the clean surface [25] but the spectra were only partly interpreted. After the appearance [26] of a theoretical determination of the density-of-states diagram for CaO, Figure 5, we are now able to assign interpretation [14] to the energy losses obtained. We thus assign losses at 5.2, 8.0, 10.0, 18.5, 26, 28.5, 31.0 and 36.4 eV to interband transitions, and losses at 14.2

481 eV to a surface plasmon (cos) and 28.5 eV to 2 cos The losses at 5.2, 8.0 and 14.2 eV were previously assigned (1,24]. We do not agree, however, to a previous [24] assignment of the loss at 36.4 eV to a volume plasmon (cop) but assign to a Ca 3pto- 3d interband transition[15]. Upon deposition of copper at RT onto CaO(100) the 5.2 eV loss is eliminated already after 2/~ of deposited Cu, and a strong loss at 4.3 eV appears in stead. v-bands

,,.

-,,

c-bands

02p

N(E)

02s

100

Ca 3 d . 4 s

50

m

i

I

-10

-5

I

I

O

5

- ~

10

15

20

25 E, eV

Figure 5. Densities-of-state diagram from CaO. From ref. [26].

-

9

9

e/e

A

"E

m

Q~

0

I

1

1

I

2

4

6

8

dc,,(A)

Figure 6. Growth of Cu on CaO(100). See text.

Cu

The other losses of the substrate become attenuated. The surface plasmons at 14.2 and 28.5 eV are strongly attenuated, supporting the interpretation as surface plasmons. The interpretation much resembles t h a t for the Cu/MgO(100) surface. The mode of growth was studied boy AES for depositions 0 < de, < 10 A. For clean CaO surfaces there are no charging-up when using the s t a n d a r d Ep of 3 keV, but for Cudeposited surfaces it was necessary because of the charging to reduce Ep to less t h a n 2 keV. As we see from Figure 6 [14,15], which shows the change in Cu and Ca AES intensity with clc,, a clear break occurs at 2.1 /k, corresponding to completion of one ML. At higher coverage, the substrate intensity is higher and the

482 Cu intensity lower t h a n what would correspond to a F r a n k - van der Merwe (layer-by-layer) growth, hence we have a Stranski-Krastanov (1 or 2 ML followed by 3D-islands) growth mode for this system. In an alternative plot [14,15] the change in the calcium and oxygen intensities (both normalized with respect to the signal from the clean surface) with the Cu intensity is given. Here the break is also clearly shown. A calculation [15] of the thickness of the Cu in the first layer from these a t t e n u a t i o n curves, based on the known inelastic m e a n free p a t h for electrons in Cu, shows agreement with the values of dcu determined by the quartz crystal microbalance. When the surfaces with the deposited Cu layers is heated to 250 ~ coalescence of the Cu particles occurs. They move together on surface, forming individual islands in the equilibrium stage, hence they now after annealing show a VolmerWeber (pure 'islanding') growth mode, demonstrating the difference in diffusion of the Cu atoms with substrate temperature. 4.4. N i cm NiO(100)

Nickel oxide is also a good insulator, a 3d s transition-metal oxide with a band gap of 3.8 eV, and as with the previous rocksalt alkaline-earth oxides the crystal surfaces charge up upon electron irradiation. And also, like these, it cannot be reduced by chemical reduction. But again, high-temperature annealing, in this case to 600 ~ of the front surface as mounted in UHV, produces a sharp (1• LEED p a t t e r n as shown in Figure 7a [16]. In AES it is necessary, though, to reduce ED to 2 keV like for the CaO case. We have studied the growth of Ni on (1• in the 0 < c~i < 145 deposit range at substrate t e m p e r a t u r e s between 125 and 185 ~ At RT, no LEED could be obtained due to charging (the Ep threshold energy required to produce enough emittance of neutralizing electrons from the surface rose from 120 eV to more t h a n 300 eV with dNi), and at t e m p e r a t u r e s higher t h a n 215 ~ no Ni(100) p a t t e r n could be obtained with increasing the deposit dNi up to 80 A, probably due to decreasing sticking probability of the Ni deposits with increasing temperature. A Ni superstructure appeared at a deposit of about 20 ~,, and it became more clear with increasing c~i. Figure 7(b-d) [16] shows the overlayer structure at 145 A of Ni which we have attributed to a Ni(100) layer whose growth direction is parallel to the Ni[010] and [001] directions along NiO[010] and [001], respectively. A comparable result was earlier obtained by reducing a NiO(100) in H 2 and obtaining Ni(100) islands [27]. The epitaxy of the Ni overlayer can be described by using a site-coincidence preference or axial commensurate growth model [16, and refs. therein] in which every 6 unit cells of the Ni lattice are in registry with 5 substrate unit cells along a preferred direction (Figure 8). The strong modulation of the surface potential of the NiO(100) substrate forces Ni atoms to line up along its symmetrical directions, resulting in misfit dislocation of the Ni layer at the interface. The mode of growth was studied by over the 0 < d~ < 16 A range at substrate t e m p e r a t u r e s in the 20 to 135 ~ range. Figure 9 shows the AES results for 20 and 70 ~ respectively. As seen from the figure, the growth follows the Stranski-

483

Figure 7. Ni epitaxy at 460~ and Ni[001] II NiO[001.

Ni(100) IINiO(100), with Ni[010] IINiO[010]

'1 9 II oT ,Io T, 9I1 9 1 9 1 I

I

I

I

I

I

I T

I I

j r

o . o

C~ t~

[1i0]

Figure 8. Misfit dislocation at a Ni/NiO(100) interface, viewed along [110]. Atoms in the (110) plane: filled symbols. Atoms in next-lower layer: open symbols.

484

120 A

e-

..; 100

=

=..

~

80

w v

i

z

o,N

\

-

=

60

e-

U"O~o 0

%

i

0

,,., e,,. ,_

40

o'1

A ...I

~

20 ~

i

i

2 Ni

i

4

I

I

6

deposition

~

,

=

8

i

10

thickness

=

I

i

I~

12 14 (A)

Figure 9. Change in Auger O(KLL) intensity from Ni/NiO(100)with dNi at 293 and 343 K. (curve b was lifted 20 units for clarity). See text

!

16

Krastanov mechanism with a monolayer point at d~i ~ 2/k [16]. In the range 2
5. COPPER AND N I C K E L ON RU'I'H~ STRUCTURF~ Rutile, one of the (three) naturally occurrmog TiO 2 minerals, is tetragonal with lattice dimensions ao=4.594 A and Co=2.958 A. Stoichiometric TiO2 is a wide-(3.1 eV bulk, 1.9 eV surface [20])-bandgap 3d ~ transition-metal oxide, and is in its stoichiometric form a good insulator. By surface h e a t - t r e a t m e n t or by Ar + sputtering [ 1] oxygen vacancies may be created, causing it to be conducting n-type [51], and the surface t i t a n i u m 3d orbitals be partially populated. Due to applications in photocatalysis and as gas sensor it has been much studied. The (110) surface is - due to the high degree of coordination of the metal ions by oxygen - the most stable of its surfaces, and is ideal for surfaces studies of the crystal metal oxides. Its EELS structure is well known [1,29-31], and so is its low energy part of its H R E E L spectrum [30], and its valence-band structure along the F-Z-M line has recently been studied by angle-resolved photoemission with a comparative bulk-band-structure calculation [32], and most recently by angleresolved inverse-photoemission [33] which suggests t h a t the nearly perfect (110) surface is t e r m i n a t e d by a bridging oxygen layer. Despite the stability, Ti02(110) surfaces may nonetheless be t u r n e d into a (1• reconstruction upon heating in UHV, and we have found [34] t h a t there is a

485

Figure 10. LEED from TIO2(110). (a) 550~ for 3 hr --> lx2 reconstrution; (b) 700~ for 1 hr -~ l x l surface; 550~ for +30 rain --> coexistence of ( l x l ) and (lx2) surface structures; (d) + 80 rain; (e) +30 rain (total 140 rain) (ref. 34).

reversible relation between the (lx l) and (lx2) surfaces (the (lx2) TiO2(110)was first seen Kao et al. [35]). As shown in Figures 10a and 10b, respectively, a low

annealing temperature (550 ~ produces a (lx2) structure at the surface, while further annealing at 700 ~ gives a perfect unreconstructed ( l x l ) surface with sharp and bright integral LEED spots. Charging-up is avoided on this otherwise insulating sample since the slight reduction by heating in UHV increases the

486 conductivity, a procedure t h a t is equivalent to n-type doping [34]. The brightness of the weaker half-order spots are dependent on the annealing time (Figure 10c-e). In the initial period of the annealing, a weak line linking the bright spots is

Figure 11. Ball model for TiO2(ll0)-lx2.Large balls: oxygen, small balls: t i t a n i u m (ref. 34).

observed. After extended annealing over hours the lines have changed via coexistence of spots and lines into the haft-order spots. This behavior can be interpreted as coexistence with (lx2) domains at the surface, leading us to the proposal t h a t the (lx2) can be considered as a missing-row model as illustrated in Figure 11. We explain the process as follows. During the annealing there is competition between diffusion of bulk-lattice oxygen and desorption of surface oxygen, i.e. a t h e r m a l equilibrium. That the desorption of the oxygen anions is favored may be due to the fact that each protruded surface oxygen anion is coordinated with two sublayer titanium cations instead of three t i t a n i u m cations as in bulk TiO2, hence some surface oxygen may escape, forming (lx2) domains. At higher t e m p e r a t u r e s the diffusion is faster and reach equilibrium with desorption, resulting in a (1• 1) surface structure. This surface crystallographic arrangement, postulated on the basis of the electron spectroscopy and the LEED results, may be further tested by atomic-scale scanning tunnel microscopy measurements. This missing-row model is supported by corresponding changes in the Auger intensity ratios, and also by the appearance of a Ti s§ state in the bandgap as shown by UPS at 1 eV binding energy, Figure 12. The more (lx2) domains, the more Ti s§ intensity, which makes sense, since when a row of atoms are missing then the Ti 4§ states with sixfold coordination numbers is reduced to a Ti s§ state

487 with fourfold coordination. Very recent synchrotron-radiation based angle-resolved resonance photo emission (RESPE) results [36] indicate, however, that the reconstructed surface does not exhibit Ti 3d states at the Fermi level which would be expected in the case of a 'simple' vacancy model [37], so even though 'off-resonance' valence-band spectra support the above model the missing-row reconstruction model appears to be of more complex n a t u r e and will need refinement. r ..E =

5.1. Cu on TiO2(110) Due to the open structure of the (110) surface, with u n s a t u r a t e d oxygen and t i t a n i u m atoms, we may expect a less inert surface. We consider the cases of adsorption of III I 1 ! !i I1~'"1!111 copper and nickel, respectively, onto C' C D E A' TiO2(110) surfaces. i i I I l I I i I I I l The growth of copper on the EF 2 4 6 8 10 unreconstructed ( l x 1)TiO2(110) Binding energy (eV) surface at RT was studied over the deposition range 0 < dcu < 80/k [27]. After depositing dcu = 7 A a Figure 12. UPS He(I) from TiO2(l10). (a) hexagonal overlayer gradually (lx2) surface t h a t was heated at 550~ appeared, and by further for a few hrs.; (b) (lx2) surface initially deposition the hexagonal overlayer heated to 700~ for 440 min followed by spots became bright and sharp annealing at 550~ for a few hrs; (c) a while the substrate rectangular (Ix 1) unreconstructed surface (ref 34). p a t t e r n further a t t e n u a t e d until it disappeared completely at dcu = 50 /k (Figures 13a-c). Upon annealing the 50 A Cu-covered surface at 160 ~ for 5 rain the substrate p a t t e r n reappeared [Figure 13d] while UPS showed coexistence of a sharp Cu 3d band characteristic of a bulk Cu metal and O 2p of the substrate. This suggests an agglomeration of the Cu atoms r a t h e r t h a n diffusion of the Cu atoms into the substrate lattice during annealing at 100 ~ The copper atoms w_ere found to be positioned in registry with the substrate lattice along the [110] direction, and with the lattice constant of the Cu superlattice in k space being 3 times larger t h a n t h a t of the substrate lattice, as confirmed by the fact t h a t no splitting or distortion of the diffraction spots at coincident points in the LEED p a t t e r n s are observed. The superstructure lattice constant is 2.498/k, thus only 2% smaller t h a n the 2.556/~ of the Cu(111) lattice.

488 The TiO2(ll0) surface is terminated by oxygen atoms, and the surface oxygen atoms are stretched outward from the layer of highest density, forming a onedimensional row along the [001] direction [1]. There is a row of Ti atoms, with 5fold coordinated oxygen neighbors in the subsurface layer, in the middle between each one-dimensional oxygen row. Based upon our LEED observations and symmetry considerations we have concluded that each Cu hexagon in the first monolayer is located between protruded oxygen rows (Figure 14), and since the second Cu layer is expected [38] to adsorb on each b u l k hollow site of the first adsorbed Cu layer, the Cu superlattice will grow in the ABCABC mode leading to the fcc structure. The copper superlattice is stable to at least 80 A thickness.

Figure 13. LEED from Cu/TiO2(l10) for different dcu with Cu deposited at RT. (a) dcu = 7 A; (b) dcu = 15 A; (c) dcu = 50/k; (d) is the (c)-surface heated to 160~ for 5 rain (ref 34).

489 In the same m a n n e r as for the previous metal-on-metal oxide systems the growth mode was followed by AES. Figure 15 shows a clear break at dcu = 3.5 /k, corresponding to one ML coverage. The substrate LEED p a t t e r n was still visible at dcu = 40 A. The growth thus seems to be island formation on top of the initial monolayer, but since the substrate p a t t e r n is still visible at 40 A [34] the growth mode apparently is a Stranski-Krastanov type, islands on top of fairly large monolayer patches ( due to the sharp monolayer LEED pattern). Very recently [39] the growth for Cu on Ti02(110) has been investigated by lowenergy ion scattering which concludes a Volmer-Weber (three-dimensional islands) type of growth. It is at present difficult to judge between the results since the two methods of surface impact are very different, the former [34] using 3 keV electrons (with carefully optimized beam irradiation time versus the accuracy for the measurements]) and the latter [39] using a 1.5 keV He ion beam (with care to limit damage during the measurements). Let us finally suggest an explanation as to why the superlattice LEED spots are

Figure 14. Hexagonal close-packed Cu superstructure on TiO2(110). Shaded balls: Cu atoms; large white balls: oxygen anions; small black balls: Ti cations.

not visible until dcu ~ 7 ,~ has been reached. This thickness corresponds to two ML, and since the copper grows incommensurately on the surface the protruded one-dimensional oxygen rows may interfere with the diffraction from the first monolayer of Cu atoms located between the surface oxygen rows. We end this more detailed discussion of the growth mode by concluding t h a t the copper grows in a slightly contracted hexagonal superlattice in a combination of a one-dimensional c o m m e n s u r a t e (in registry) growth along one substrate direction and an i n c o m m e n s u r a t e growth in the two-dimensional (1• 1) surface, and t h a t the

490 protruded oxygen rows are i m p o r t a n t in the formation of the superlattice.

[]

140

..~

120

D

_A

[]

C=

[]

A

,4

n Cu (M23VV)

A 0 (KLL)

L_

~0 1 0 0

a

I.l.I

'" z

A

80

AD OA

=

.m..

...m

O

60

O

n

o Ti (L3M1M23)

DO

A~

~X~

A

O

-

O

40

~

o

0

D 20

-

A o

O0 0

0

000

0

r'!

9

O

I

I

2

I

a

I

I

Z3

0

I

&

0

I

4 6 8 Cu deposition thickness

0 I

I

10 (A)

0 I

I

12

Figure 15. Changes in Auger Cu, O and Ti intensities w i t h increasing dcu 293-K deposition on TiO2(110).

When we follow the deposition of Cu on the ( l x l ) surface we note a clear evolution of two p e a k s and a shoulder in the valence b a n d He I s p e c t r u m (Figure 17) which are m a i n l y contributed by O 2p bands. These b a n d s are a t t e n u a t e d gradually w i t h do, while a new shoulder, which subsequently becomes a peak, appears in the oxide b a n d g a p due to Cu 3d emission, and we note t h a t the Cu 3d binding energy shifts continuously until reaching a pure copper s p e c t r u m at the thicker layer (where we have the hexagonal superstructure). We see (Fig~ore 17) t h a t the copper surface state S.S. is not clearly established u n t i l about 10 A of Cu (the figure also shows a He satellite).

491

dcd~,)

=E

Simultaneously we note a gradual attenuation of the substrate electronenergy losses while there appear two new interband-transition losses (A and B) and a plasmon loss (C) which shift with dcu. Besides, the doublet Auger Cu 3p spinorbit interaction-peak merges [40] as it does in Cu20 because of the weakened and broadened Cu 3d band overlapping the Oinduced 2p states. These latter bonds are rather weak due to the lower affinity between the Cu and the protruding O t h a n between Ti and O , but still sizeable enough to give a measurable charge transfer in the interface.

era

5

EF

.

Binding energy (eV)

.

.

10

.

12

Figure 16. Change in UP He(I) spectra with do,, from Cu/TiO2(ll0) at 293 K (ref. 40).

A

""I L l

~

3

~

~

2

~

,

-

0

.

A

02 ,

In Figure 18 is shown the change in charge transfer with the binding energy calculation, using the assumption cop- ~]n of the plasma frequency relation with the electron density and associating the change with a decrease of the free-electron density in the adsorbed layer for the O-Cu interaction in the surface. Comparing to the somewhat similar Cs-on-Cu(111) system [41] we may expect this model to extend to submonolayer coverage where only two-dimensional islands exist, since the density of the Cu on Ti02 is calculated to 1.23 • cm 2 from the StranskiKrastanov growth-mode results [34], the density being considerably higher t h a n for

~

1

,

;i/21 ,

1

tie

40 {

{

3'o

{

2'o

{

,'o

Energy loss (eV)

{

o I

Figure 17. Changes in EEL spectra with dcu from Cu/Ti02(ll0) at RT (ref. 40).

492 the Cs/Cu system. By comparing Figure 18 to results from the Cu-on-Cu~O [42], which indicated a charge transfer of 0.31 e/Cu adatom during formation of Cu20, and to XPS results for the deposition of Cu on MgO(100) [12] in which formation of Cu(I) states were o b s e r v e d , then we conclude t h a t the initial deposited Cu atoms probably exist as Cu(I) states on the TiO2 surface.

,_ 0.3 r-

-

0.2

=,-

0.1

O0

O

2.6

I

2.7

I

I

I

I

2.8 2.9 3.0 3.1 Binding energy (eV)

31

.2

Figure 18. Indication of charge transfer between Cu islands and TiO2(110) substrate v s . EB(CU 3d).

A,~c ~,

AREELS

[II I\ ll! i \

Cu/Ti02 (110)

II ~ i \ III I\

I ,~

\1i ....-.... oo e-

,4 ......

e,,era

\

fVa

= ~1

I

=

I

,tT',/; I

I

\~

\

Eo:'OOeV

o; ooo

\

u

/~/

\l/~

o

/ ~

~

95

~ I

clean

Energy loss, EL (eV)

Figure 19. Changes in AREEL spectra with dcu from T i Q ( l l 0 ) at RT. Absorption edge and loss values are marked (ref. 46).

However, we should be aware t h a t this state is quite different from t h a t of bulk Cu20 where we have Cu 3d levels involved in bonds to O, hence not behaving as isolated core states. In a comparative experiment using the catalyst CuC1 as adsorbent on the (l• found [43,44] a rehybridization between the Cu 3d (downwards oriented in the CuC1induced (4• 1) reconstructed surface) and the O 2p levels. The assignment of the oxidation states of Cu and the bonding relations among the various population levels are of considerable interest in these systems, e. g. in the elucidation of the initial stages in some catalytic reactions. We thus conclude t h a t interactions between the Cu 3d levels are very limited, and t h a t the bonding across the interface mainly is contributed by 4sp electrons, resulting in a charge transfer. The role of defect sites is also of great interest. The (1• surface of

493

/• I/

TiO~(110) contains intrinsic defects located mainly on the top few layers (in contrast to Ar+-sputtered surfaces where Ti § states, assigned to 3 d 1 polaronic states localized at crystallographic shear planes [45], are found deeply in the substrate). We have found from UPS results [40] on the (1• surfaces t h a t more charge transfer occurs from the initial deposited Cu atoms (which may be considered as extrinsic defects, playing the role of Cu (§ surface donors) to the defect sites, and t h a t copper multistates exist on the (lx2) surface.

He(I) UPS Cu/Ti 02(110)

J

....,..

I:=

I

a

f

/

~,,~

4.0

oo

( I I I

J Binding energy (eV)

Figure 20. UPS from CufriO2(ll0) at RT. The surface state of Cu hexagonal superlattice film is m a r k e d (ref 46).

The copper atoms t h a t are placed around defect sites probably correspond to higher binding-energy states, and may - due to the poor screening of the titanium-ion pair at the defect site -yield a higher capability of the defect site to take more electrons from copper adatoms. Hence the defect sites seem to have capability to trap electrons donated from copper atom and thus modifying the interface properties. For the Cu/TiO2(110) system, which is the one system of these t h a t we have most intensively studied, we finally discuss some fundamentally interesting, we think, phononelectron angle-resolved-EELS (AREELS) results in conjunction with UPS m e a s u r e m e n t s [46] from deposition of Cu on the (1• surface, using a HREELS

D 7.25 eV

E0= 40 0 eV

A

e ~ dcu(~,) 51 0

=2

560

39 5

660

. m

51 0 56.0

40

660

1'o 1'2

Energy loss, EL (eV)

1'6

Figure 21. Changes in AREEL spectra with scattering angle from Cu/TiO2(110) monolayer and thick film (ref. 46).

494 i n s t r u m e n t in the specular-reflection geometry (where the m o m e n t u m - t r a n s p o r t parallel to surface (kl) is nearly zero, whence the loss functions are analogous to those derived from optical data [46]) at 22 eV < Ep < 40 eV as a function of scattering angle in the range 490 < 0 < 66 ~ and the analyzer angle at 0 ~ in the UPS measurements. Phonon-assisted interband transitions are known to occur in several semiconductor materials, such as Si and Ge, in which the m i n i m u m in the conduction band and the m a x i m u m in the valence band are not located at the same point in the Brillouin zone (BZ), and in optical absorption experiments these indirect transitions require assistance from phonons to m a i n t a i n conservation of momentum. Upon deposition of Cu over the 0 < dcu < 40 A range at E p - 40.0 eV and scattering angle 0 = 56.00 [46] we observe (Figure 19) the appearance of four copper-related losses (A-D) while the substrate losses are attenuated. The intensity of the clean-surface elastic peak decreased almost one order of magnitude during the initial depositions and started increasing again at about the monolayer point d c ~ - 4/k. Also, it was observed t h a t the copper absorption-edge shifted, first toward lower energies and then, passing through a minimum, toward higher energies. The absorption edge could be determined as the cross-point of the extrapolation of a low-energy-power-law-dependent curve and the linear extrapolation of the absorption-edge curve [46]. The 3 losses (A,B and C) of lowest energies are allocated to interband transitions with Cu 3d levels, as described above in the results of the angle-integrated experiments, although there only 2 losses were resolved because an appreciable contribution from (k I ~ 0) scattering in the angleintegrated data smears out the structure of the (k I - 0) loss function [48]. Below 4 A the shift of the Cu 3d band dominates and is interpreted as previously discussed in terms of charge transfer. The fourth loss (D) is, as previously, assigned to a plasmon excitation.For thick (dcu- 39.5/k) overlayers an absorption of 2.08 eV was measured, in agreement with optical data for Cu. This edge was seen to move toward lower energies at coverage near the monolayer point which was surprising since the Cu 3d levels remain unshifted above the monolayer point as observed by UPS (Figure 20). To exclude the possibility of diffraction effects AREEL spectra were obtained (Figure 21) for different 0 both for the monolayer film and for the thick film. Figure 21 indicates both absorption edges and the other losses, A-D. As discussed in the above paragraph, indirect transitions require assistance from the phonons, hence the absorption edge for those semiconductor materials shifts Eg- Eph r a t h e r t h a n to Eg, where Eph is the phonon energy. The similarity between the present case and the case of indirect transition caused us to the proposal [46] of the existence of a substrate-phonon-assisted Cu-overlayer-interband transition, i.e. across the interface. This type of transition has to our knowledge not been reported earlier in the literature. The proposal is supported by a detailed analysis [ref. 44 and references therein] of the experimentally observed electronic transitions from copper, involving both occupied and unoccupied states (we do not need to consider the substrate electronic structure because the interaction to the substrate is weak as discussed above [34]); in brief the analysis is as the following. Due to the many possibilities of vertical transitions inside the BZ the discussion is complicated. The loss features may be dominated more by transitions near the

495

,ix', .-.

w3 Jl

2

X4

0

r~ i

FERMI LEVEL

,,z,

Wl

LU

-

F25,

F

h

X

Z

W

Q

L

A

F

2:

K

X

k (2n/a)

Figure 22. Bandstructure for Cu. We have marked examples of the three types of transitions; see text (ref 46 and refs. therein).

zone boundary t h a n by those near the center (due to the much larger area in the former region), in the most simple case, but to discuss the present case we have suggested a classification of three different types of transitions upon which a direct comparison with simulated energy bands can be made along the BZ symmetry axes The three types, illustrated in Figure 22, are (i) critical-point transitions from the high-symmetry points at which the m o m e n t u m matrix elements are so high t h a t they can contribute to the structure in the loss function; (ii) the transitions t h a t occur between occupied states and the Fermi surface and (iii) the transitions, which we refer to as a volume effect, t h a t occur in the region near the BZ Fermi surface and t h a t are allowed by the dipole-selection rule. The analysis reveals, with the help of Figure 22, t h a t the 2.08 eV absorption edge corresponds to transitions from the osculating points, while the loss intensities around 2.72 eV are contributed both from the osculating points and by the volume effect from the large 5-+6 and 4-+6 band-to-band transitions near X inside BZ and near the Fermi surface around the L neck, and the broad loss B mainly is caused by the volume effect from the large 3-+6, 4-+6 and 5-+6 transition regions. The transitions 6-+7 between sp states and 1-+6 from the bottom of the 3d band dominate loss C. In the monolayer case the three-dimensional (3-D) BZ becomes a two-dimensional (2-D) BZ, and the possibility for vertical transitions is lowered very much, in agreement with the one-order AREELS intensity decrease from the case of the thick layer to the case of the monolayer. A similar analysis as above is made for the 2-D case. In the UPS results of Figure 20 we note t h a t the submonolayer deposition of Cu

496 causes a shift toward a higher binding energy, as we discussed earlier, in contrast to the usual findings of the Cu 3d levels shifting toward EF in comparison to the bulk d band in isolated monolayer films, hence the blunt absorption edge for the monolayer deposit cannot be allocated to pure electronic transitions from the osculating points, but r a t h e r to a substrate-phonon-assisted Cu-interband transition. The energy difference between the two absorption edges is 0.3 eV, within the region of an oxide-substrate phonon, where the f u n d a m e n t a l modes and their combinations yield significant intensity up till at least 0.4 eV in our observations (a Cu phonon cannot cause such a large absorption-edge shift). The intense substrate optical phonons causes a coupling with electronic fluctuations in the copper overlayer, producing long-range disturbance into the vacuum and hence causing inelastic scattering of the incoming electrons by the dipole field into a small angle around the specular direction. Simultaneously the thermally excited phonons are annihilated, and the edge must shift down (phonon creation would need energy and thus have shifted the edge up). We therefore conclude t h a t we have revealed a substrate-phonon-assisted electronic Cu-interband transition at the interface. 5.2. Ni (m TiO2(ll0) Deposition of the transition metal nickel onto Ti02(110) was already early the subject for investigation due to its importance in m e t h a n a t i o n by heterogeneous catalysis, particularly after it had been found [49] t h a t Ni on a rutile-anatase mixture has a considerably stronger activity t h a n on other (SiO2 and AleO s) oxide supports. Kao et al. [35] thus found for deposition at RT over the 0.23 < 0 < 15.8 coverage range on reduced (i.e. heat-treated to release oxygen) a charge transfer from Ti02 to Ni (the Ni atoms become negatively charged) varying between -0.13 and -0.07 e per adatom at 0 - 0.5, by observing the shift in the Auger Ni(LMM) peak and the 2pa ~ core level of Ni, and t h a t the amount of charge transfer was dependent of the surface p r e t r e a t m e n t (annealing / sputtering). This behavior is often ascribed to a characteristic of the important strong metal-support interaction (SMSI) process in heterogeneous catalysis (the effect is characterized by a suppression or loss of the ability of the metal to chemisorb (hydrogen, carbon monoxide) when positioned on a reproducible oxide support, leading to selectivity properties). Later Onishi et al. [50] found t h a t for Ni deposits on a non-reduced TiO2(110) there was a small, 0.1 e per adatom, and as the overlayer grows more dense, the lateral interaction between the Ni adatoms predominates and inhibits the electron transfer through the interface. No epitaxial relationship was found for the Ni overlayer in these studies. We have found [51], however, t h a t epitaxy did in fact occur, but the epitaxial growth of Ni on the (1• 1)TiO2(110) surface cannot be predicted by any existing theory. Two types of nickel islands are formed, a hexagonal structure oriented parallel to the substrate, and a hexagonal structure whose direction of growth is inclined with reference to the substrate plane. The growth of u l t r a t h i n layers of Ni on the (l• surface over the 0 < dNi < 80 A range at RT shows attenuation of the substrate L_EED p a t t e r n up until 10/k where three weak parallel lines along the substrate [110] direction begin to appear, and upon further deposition the intensity of the lines increases, and for dNi > 30 /k the substrate p a t t e r n has completely disappeared. After slightly annealing, the three-line LEED p a t t e r n s changes into many extra spots lying in the lines. After careful analysis of a series of the p a t t e r n s three sets of diffraction

497 spots were analyzed, and it was observed t h a t among these two of the sets are symmetrical with respect to the substrate {001] line in reciprocal space. These two sets move in complex way with E, such t h a t corresponding symmetrical diffraction spots move away from each other, within the line, with increasing Ep. This usually indicates t h a t the overlayer is tilted away from the substrate plane. After a computer simulation based upon Ewald-sphere analysis [16] we noticed t h a t the observed complex movement of the diffraction spots only occurs along the [110] direction in the substrate. The relation between the spots in the [001] direction is somewhat similar to the case (above described) of Cu on the same substrate, hence t h a t Ni is packed in hcp, so we used a hexagonal lattice in the simulation. We found t h a t the best fit for all the p a t t e r n s in 10 eV steps over the 50 < Ep < 180 eV range was obtained for angle of 270 between the ( l l l ) N i plane and the substrate with the tilt axis in the [001] direction, indicating t h a t two types of Ni islands are formed simultaneously upon the initial monolay_er. The first type grows parallel to the substrate in registry with the [110] direction but incommensurately in the other 2-D directions of the substrate, as in the case for Cu on this substrate [34], leading to normal growth of fcc Ni. We found a lattice constant of 2.50/k, i.e. lightly larger t h a n the bulk 2.49/k. For the 270 tilted plane we found t h a t the (131)plane satisfies the requirement: the angle between the two planes is 29.5 ~ in good agreement with the observed 27 ~ The two fcc structured types of islands can be characterized as: N i ( l l l ) II TiO2(ll0), Ni[101] II TiO2[001] and Ni(131) II TiO2(ll0), Ni[101] II WiO2[001] 60

-\ O

~9

50

,4

_O

\

... 40 30

~ 20

13 Ti(L3M23M23 ) 0

-

..--...

z

o O~KtLI

_ 0

R

\

o,%

O,o

-%\

e-

10

~ 3 u ._~.._ . . E l ~ . ~ ~ . .h. . ~

_

~

0~

_

_

O~D

0

2 4 6 8 10 12 14 Ni deposition thickness (A)

16

Figure 23. Changes in Auger Ti and O intensities with dNi on TiO2(ll0) at RT. (a) quasi-isotropic growth model; (b) anisotropic growth model (ref. 16).

The growth mode for Ni on the substrate at RT was studied over the 0 < dNi < 16 range by AES (Figure 23). We find a sharp monolayer breakpoint at dui = 2.5 A. Since the substrate was still visible at 25 A we exclude a layer-by-layer growth mode and assign it to a S transkiKrastanov growth mode. With the same assumptions as for the growth of Ni on NiO(100), as described above, we have analyzed the data using the quasi-isotropic growth model [28] and found t h a t only the isotropic growth model fits correctly the growth behavior. We expect t h a t after the first hcp Ni monolayer with sixfold symmetry, the anisotropic features of the

498 substrate is weakened so much t h a t the observed isotropic growth can occur, as observed. As for the nickel oxide case we have also here determined the density of the Ni islands before coalescence and find 2.9 • 10 is cm 2. The 3-D growth of Ni (after monolayer coverage) hence is described well in both cases by the quasiisotropic growth model. At RT, CO molecules do not adsorp on a clean TiO2(110) surface, but at a nickelpromoted surface it occurs. In a combined valence-band UPS and vibrational HREELS experiment [52] a distinct peak in He(II) UPS at EB=I 1.0 eV below E F grows, and further CO exposure induces two new CO-derived peaks appear at 8.0 and 11.0 eV (as was, independently, found in an UPS study by Onishi et al. [50]) These two peaks can be attributed to the emissions from (1~ + 5(~) and 4(~ states of CO, respectively. The CO adsorbs with its carbon-end oriented towards the surface. The existence of molecular CO at the Ni/TiO2(110) surface (in agreement earlier studies [53] showing associative CO adsorption on a Ni crystal at RT) was confirmed by HREELS spectra [52] on 1-, 2-, 3- and 4-/k Ni-deposited TiO2(l10) surface at RT. The saturation coverage of CO increases with dNi, indicating that CO molecules bind to the Ni atoms rather t h a n to the substrate atoms. At saturation coverahe, CO molecules adsorb simultaneously on the 2-fold bridge-sites and the terminal sites on the (111)-oriented Ni islands on the TiO2(110) support. The occupation of the edge-sites of the Ni islands gives rise to an exeptionally low C-O stretching vibration of 152 meV. This frequency, indicative of a weakened C-O bond, suggests existence of a precursor to the dissociated state, which is important inthe understanding of the bevior of the catalyst supported on the metal oxide. 5.3. Ni (m TiO~(100) The (100) surface of TiO2 reconstructs upon 500~ annealing to a (1• at 800~ to a (1• and at 1200~ to a (1• LEED p a t t e r n [29], and also this subs~raLc surface shows Ti § interband transition upon Ar § sputtering by removing surface oxygen [55]. The interface with Ni is of particular interest in the process of CO hydrogenation. Kao et al. [54] first used a TiO2(100) as substrate for Ni deposition as a model catalyst and found by XPS t h a t Ni atoms at the interface are negatively charged. The growth of Ni upon TiO2(100) was found [56] by AES and confirmed (using the sequential-layer-sputtering model [57]) by secondary ion mass spectrometry (SIMS) to follow a layer-by-layer growth for the first three Ni layers. Then islands of Ni grow. For a non-stoichiometric support no island growth is found and the Ni is diffusing into the bulk even at RT, depending on the amount of oxygen vacancies present [58].

6. COPPER AND NICKEL ON CORUNDUM STRUCTURRS Of the (trigonal-lattice) corundum structures, we will here describe recent results for adsorption of Cu and Ni on a non-transition-metal oxide, alumina, a-A120 s (saphire), and of Cu on a transition-metal oxide, a-Fe2Os (haematite). A1 and Fe are here surrounded by six oxygen ligands to their (distorted) octahedral sites. Among these, the a-A1203 is by far the one t h a t have been studied most intensively due to its very wide applications in fields as diverse as ceramics, composites, microelectronics, refractories, catalysts, membranes and cements. It is unique also with respect to its electronic properties; it has a very wide band gap

499 of 8.7 eV [59]. After heating to 900~ in air followed by h e a t - t r e a t m e n t in UHV to 700~ a sharp lx 1 surface structure is observed [60] (heating the crystal in UHV to higher temperatures, up to 1400~ yields rotated ~]3x~/3, 3~]3x3~]3 or ~]31x~/31 reconstructed AleOs(0001) surfaces [61]). The electronic structure of a-alumina is now well understood, both from a semiempirical (extended Hiickel) calculation [62] and from an embedded-cluster calculation [63]. The loss structure has been studied experimentally by EELS [63-65]. Work on adsorption of Cu and Ni (and other metals) onto single-crystal surfaces of these oxides are scarce, however, perhaps due to experimental difficulties with charging effects as with the earlier discussed MgO. 6.1. Cu r AI208(0001) The crystallographic structure of the Cu/a-A1203(0001) interface was studied by tra_nsmission_electron microscopy [66], and epitaxial (111)Cu II a-A12Os(0001), [211]Cu II [2110]a-AleOs(0001) relationship was demonstrated by back-reflection Laue X-ray diffraction [67], ion-etch characteristics and scanning electron microscopy [68], and the strong adhesion of copper to this surface obtained as a result of an ion-sputtering t r e a t m e n t was investigated by XPS [69,70]. Theoretically the metal-sapphire shear strength was investigated by serfconsistent-field X-alpha scattered-wave cluster molecular-orbital models, and it was found t h a t a chemical bond is established between the metal d-electrons and the nonbonding 2p-electrons of the oxygen anions on the A12Os surface [71], and that the bond strength between the oxygen anions to the close-packed firsttransition-metals is placed between the empirical values for stronger metal-oxygen bonds and weaker bulk-oxide bonds [72]. We have studied [73] the growth of Cu on aAleO~(0001 ) in the 0 < dou < 130 A range, and have found a Cu(11 1)-R30 ~ s u p e r structure on the (lx 1)-aA120~(0001) surface after having heated a 53-/k film to 650~ for 30 rain (Figure 24). The same LEED p a t t e r n as the one shown in Figure 24 could for the same 53-/k surface be obtained also at lower temperatures, for instance 180~ at 15 rain, but then it was necessary to use rather high Ep because of charging effects. By AES, EELS and UPS (in the latter Figure 24. LEED showing rotationally we compared to an A120s aligned epitaxial Cu(111) on a-A1203(O001) film formed on an AI(111) after 650~ heating, surface and made sure that charging effects did not interfere since the O 2p peak

500 did not shift while the Cu peak at the same I I I I l I time shifts tohigher binding energies with dcu [73]) we have found [60] t h a t already from dCu (/~) the initial low-end submonolayer coverages a chemical 27 bond can be established between copper and the 1• smooth substrate, and our data (Figure 25), which show no Cu(0) signal in the initial submonolayer stage, support an interpretation in terms of a weak charge1.7 transfer process where charge is donated from the copper particles to substrate oxygen, so that initially Cu(I) 0.7 states and t h e n Cu(0) states are seen with increasing dcu (by presputtering an A1203(0001) surface it 0 (clean a-AI203) was found [70] t h a t the interface involve Cu(I) d 1~ configuration). Due to an observed I ! I I' = = more rapid decrease of 20 30 40 50 60 70 the aluminum-toKinetic energy (eV) oxygen Auger intensity ratio for the lower (dcu Figure 25. Changes in Cu Auger fine structure < 2/k) deposits it is with dcu. likely t h a t we see initial copper atoms positioned in a-top position of the aluminum" sites which are in 3-fold hollow positions created by the oxygen atoms (Figure 26). Ideally the outermost atoms of the 1• 1 a-AleOs(0001) surface are aluminum, but here the outermost atoms are expected to be oxygen due to the high-temperature oxidation p r e t r e a t m e n t of the sample. This was confirmed by AES. ,..-,,..

L'u .,,....

cIp

.0,.,, e-

,._ Q.I

501

Figur 26. Surface geometric analysis of LEED pattern (Figure 24) from Cu/a-AleOs(0001),

1.3 A Cu/~-AI203(0001) 0.6 ._o

0.4 .'~_ oo

,_

"-

,,Ik

9

Cu/AI AI/0

02 9 ~"''~--=-0

I

100

i

I

I

i

200 300 400 500 Annealing temperature T('C)

=4

600

Cu/0 i

700

Figure 27. AES on the thermal stability of Cu small islands on a-A12Os(0001).

502 From the changes of the Cu-to-O and Cu-to-A1 Auger intensity ratios it is difficult clearly to evaluate the growth mechanism at RT. The ratios initially change linearly with dcu, followed by a slow exponential change, hence a StranskiKrastanov mode, a copper ML followed by subsequent island (cluster) nucleation [60], is indicated. However, later XPS results [75] conclude a Volmer-Weber type of growth, pure cluster nucleation with no monolayer formation at RT (and also for the high-temperature ~]31• reconstructed surface). As in our case, no superstructure LEED pattern was observed at RT.

/~

dcu(/~) 3.3

T/~ 670

330

3.3

220

3.3

25

A

v

vJ

e-. D

~__

0 (clean)

I

I

I

I

I

25

I

20 30 40 50 60 70 Kinetic energy (eV)

Figure 28. Fine-structure Auger analysis on the t h e r m a l stability of 3.3-A Cu thick deposits on a-AleOs(O001).

503 Very recently, though, an Angle Resolved XPS investigation [76] finds t h a t copper initially grows as uniform layers on l x l A12Os(0001 ) followed by formation of clusters after 2-3 atomic layers of Cu deposited at RT, thus agreeing with our conclusion on growth mode (and with the mechanism of charge transfer). A recent combined HREELS and AES investigation [77] also finds t h a t the first monolayeforms a smooth overlayer at 95 K on a thin A12Os film t h a t was formed on an AI(111) surface, i.e. Stranski-Krastanov-type growth, and t h a t very little change was observed by heating the layer to 400 K. We have investigated also the t h e r m a l stability of the copper film [60, 65]. As seen from Figure 27, the interface composition is quite constant until about 400~ This further indicates a chemical bond at the interface between copper and the clean (~-A12Os(0001) surface. We see some decrease for a 2-/k film in the Cu-to-O and Cuto-A1 ratios with increasing temperature; there might be three possible causes for this: (i) interdiffusion between copper and substrate, (ii) copper agglomeration at the surface, and ('di) desorption or evaporation of Cu from the surface. We consider evaporation to be insignificant at t e m p e r a t u r e s lower t h a n 400~ and the r a t h e r unchanged ratios over the RT to 400~ range for monolayer copper on the substrate indicate t h a t interdiffusion has not happened at the interface. Rather this indicate t h a t small copper islands have increased in size with temperature, resulting in the decrease in these Auger ratios. Figure 28 show the changes with t e m p e r a t u r e of the Cu(M2.3VV) transition for a 3.3-A deposition [65]. The heatt r e a t m e n t (for 15 rain) at 200~ caused a split into two peaks (at 55 and 58 eV) in this Auger line compared to t h a t of the u n h e a t e d sample where the kinetic energy for this transition is 57 eV (A1 and Cu(I). We believe t h a t the split is caused by a mixture of Cu(I) and Cu0) states. Furthermore, we note t h a t the peak position is shifted to a lower kinetic energy, relative to the transition in copper mc~al, indicating t h a t agglomeration has occurred as a result of the heat treatment. F u r t h e r heating to higher t e m p e r a t u r e s results in a "~ r--'--] i"-'-'-I decrease of the Cu(0) intensity, here at 58 eV, when compared to the 55-eV peak which is ascribed to Cu(I), suggesting that the Cu(I) on the surface at these F--] I--~ I---] t e m p e r a t u r e s is more stable t h a n Cu(0). b L . v/~ EELS results for a 0.7-A deposit [65] shows t h a t the new RT-copper-induced loss at 5.5 eV, which suggests t h a t m Cu(I) can be established in a charge-transfer process for 1 thicknesses dcu< 2/k, is stable until around 400~ and is decreased by heat t r e a t m e n t to 600~ in agreement with the AES results. 1-3 N Figure 29 illustrates schematically the different situations on the substrate at RT for 2-D and 3-D clusters and after annealing, respectively. o' - AI 2 0 3 Cu In the case of submonolayer-copper thicknesses a quite Cu (111) stable chemical bond between copper and the substrate is thus established, resulting initially in copper in a Figure 29. Growth Cu(I) state, while for thicker films the copper diffuses in model for Cu/athe direction of the surface lattice to form nucleation A12Os(0001). See centers. text.

504 In the above referred HREELS-AES investigation [77] it was found t h a t heating to 700 K caused an inward diffusion of the copper overlayer. It should be noted, however, as also the authors do, t h a t since the AleOJAI(111) substrate possesses metallic a l u m i n u m beneath, and likely intermixed with the AltOs films, the thermodynamic driving force for the diffusion of copper into the bulk is the formation of a Cu-A1 alloy [77,78] in t h a t case (their films were only -4 and -8 /~ thick). 30-/~ thin A1203 films, made by oxydation of thin A1 foils, were also investigated as to the effect of cluster size on the shifts in core-level binding energy and in the kinetic energy [79]. A charging-up behavior in the m e a s u r e m e n t s at certain energies, as previously discussed (Sect. 3), was naturally found also for this strong-insulator system. a-Al2Os(0001) The energy-loss structure for RT deposits of Ni on a-A12Os(0001) (model for the uses in a range of catalytic processes) was obtained for thicknesses of 0.5, 2 and 3/~ [64], and it was concluded t h a t the stoichiometric surface generate chemical bonding with a l u m i n u m dangling bonds as trapping centers (a similar conclusion was obtained for the (i012) surface [64]). A combined LEED and XPS investigation [80] did not observe detectable reaction at RT, but when Ni was deposited at 800~ in UHV the alumina was seen partially reduced upon Ni deposition, and when the deposition was carried out at that t e m p e r a t u r e in the presence of 02 a NiA1204 spinel phase was formed. Here the XPS showed the shake-up peak characteristic of the oxidation state of nickel, and LEED showed an ordered overlayer structure with sixfold s y m m e t r y with the spots from the NiAleO 4 at about half the distance from t h a t of sapphire (there is an approximate factor of 1.7 between the lattice constants of NiA1204 and a-A12Os, and a lattice mismatch of 3.5% between its (111)plane and the substrate surface). Postannealing (in separate RHEED chamber) to t e m p e r a t u r e s up to l l00~ showed dissociation of NiO until complete metallic Ni was recovered without observation of a NiA1204. Sixfold epitaxy was observed both in the UHV and in 5• 10 .7 Tort 02. LEED was very much hampered due to charging, though. Decrease in the intensity in Kikuchi bands (lines) indicated t h a t the overlayers were strained. An ab initio, cluster, unrestricted Hartree-Fock calculation on the electronic bonding properties of the Ni/Al~Os has found [80] an average bond strength of 5.3 eV and an average charge transfer of 1.5 e per adatom from Ni to antibonding states on the alumina surface and proposes t h a t the catalytic action of aluminasupported nickel is in part connected with the presence of nickel as a positive ion rather t h a n as Ni ~ (and thus seems to contradict earlier experimental findings [80] that the catalytic turnover number measured on a single-crystal Ni, with site density derived from the Ni atoms of a (100) plane, is in good agreement with values reported from high-area (alumina) supported catalysts, with site density derived from chemisorption data). As in the case of Cu/AleO3(0001) we will also for the Ni/AleO s system compare briefly to results from deposition of Ni onto thin films of A1203 formed upon clean AI(111) surfaces. In an AES-HREELS investigation on Ni deposited onto thin A120 s films grown by oxidation of an AI(111) surface it was found [83] t h a t the deposited Ni atoms form 3-D clusters on the A12OJAl(lll) substrate at 200 K, and t h a t thermally 6.2. N i r

505 induced smoothing and diffusion of the Ni overlayer occur at characteristic t e m p e r a t u r e s in the 200-700 K range. From a layer with an A1203 thickness of about 3 A it was tentatively postulated t h a t at t e m p e r a t u r e s below 400 K the results suggest a smoothing of the 3-D Ni cluster layer to a more uniform Ni overlayer occurs, and at t e m p e r a t u r e s above 400 K t h a t inward diffusion of the Ni overlayer p r e d o m i n a n t l y occurs. The Ni/A12Os system has been used as catalyst for m e t h a n a t i o n of CO, and no Ni diffusion had been reported for the 200-700 K t e m p e r a t u r e range. However, it was found now t h a t surface segregation of metallic A1 toward the Ni overlayer occurs during the Ni penetration process, explained by an interdifhasion of Ni and A1~ in macroscopic channels within the A12Os layer, p e r m i t t i n g the formation of a Ni-A1 ahoy. It was thus found t h a t the metallic a l u m i n u m b e n e a t h the AleO 3 film plays a major role in driving the inward diffusion of the Ni overlayer. Recently we have investigated this system by AES, XPS and EELS and have found [81] evidence by correlated AES and XPS m e a s u r e m e n t s for a StranskiKrastanov type of growth (a 2-D layer followed by 3-D clusters) for Ni deposited at RT on a 7-AleO s thin film grown on an Al(111) surface. We have calculated the equivalent thickness of the Ni deposits at the end of the 2-D p a r t of the growth by using the Al(111) plane as basis, and the calculated thickness corresponds to the formation of a layer of Ni on a l u m i n a (with a sticking coefficient near one), as also evidenced [84] by t r a n s m i s s i o n electron microscopy (TEM). XPS showed unaffected A1 core-level features, and the appearance of a new peak at 3.5 eV below the final E F, and the peak shifted with dNi towards a position located 2.0 eV below EF. The Ni 2p peaks shift toward lower binding energies. The curves EB(Ni 2p3~) vs. dNi and EK(L2.sMV) vs. dNi each contains 3 steps which we consider for the two regions, I: dsi < 0.4 A and II: dNi > 0.4 A. In region I, in which the 2-D monolayer is formed, we find large changes. EB decreases 2.8 eV, the modified Auger p a r a m e t e r a' decreases 2.2 eV while E K increases 0.6 eV. The variation in a' at the initial stage we interpret as due to a low-nucleation-rate formation of small clusters on the surface, i.e. a size effect, as observed by TEM [85]. The large shifts in EB we assign to a combined effect of the presence of clusters and a formation of a new compound. F o r m a t i o n of nickel oxide is in agreement with an observed large change in the initial effect, ae = -3.9 eV. The negative sign of ae indicates a positive charge for the nickel clusters, which we explain as a charge transfer from Ni 3d toward ligand states, as we described above for the Cu/TiO2 system. Using the method of Kao et al. [84] we found a positive m e a n charge of 0.39 e / a d a t o m with regard to the final deposit, indicating a selective interaction between Ni and O, and t h u s find positively charged Ni as suggested in the above mentioned ab initio calculation [81], although we find a lower positive charge t h a n the calculated value of 1.5 e per adatom. For region I we have thus found formation of 2-D clusters of oxidized nickel and have elucidated the ionic character of the Ni-A1 bond. In region II, E B and E K still approach the bulk values, and we see 3 oscillations in the evolution curves for 0.4 A < dNi < 8 A. The oscillations in the Auger signal we interpret as instability of the deposit, i.e. some of the 2-D clusters nucleate to 3-D structures. This might explain an observed decrease of the Ni m e a n charge toward a value of 0.21 e per adatom. With reference to the previous discussed work [82] of Ni/AleOs/Al(111) we can deduce t h a t metallic Ni are located inside the aggregates, but the importance of the initial effect indicates t h a t the electronic

506 structure is quite different from t h a t of the metallic state. The m e a n value of a' gradually approaches the Ni ~ value, and the amplitude of the oscillations decreases, indicating establishment of the Ni ~ state. For dNi > 3/k the 3-D clusters are growing in size, and the initial effect decreases. The electronic structure seems to be established before reaching 2 ML. When comparing the chemistry in the interfaces of the Cu and Ni on bulk aA1903(0001) crystal substrates with those of the few-/k-thick A120s films formed by oxidation of an A1(111) surface there is a m a r k e d difference in particular with regard to the diffusion behavior, where the A1 in the latter is a strong driving force. However, much useful information have been obtained which help in understanding the chemistry for both systems. As referred to (Sect. 6.1) for the similar case with Cu it was shown theoretically by Johnson and Pepper [71] from cluster model calculations t h a t a primarily covalent bond can be established between metal and the oxygen anions on the A120 s surface. It was also shown t h a t its strength decreases in the series Fe, Ni, Cu, and Ag. The reduction in strength in this series they attribute to primarily to the increasing occupation of antibonding orbitals established by the metaloxygen interaction. Their calculation was expanded to show [72] a uniform decrease from about 5 eV to about 1 eV for binding energies per interracial 0 when an O-covered (0001) basal plane of a-AleO~ binds oxidatively to close-packed metal surfaces of the first transition series, and that, except perhaps for Cu, adhesive failure of interfaces of the structure modeled will occur between the first and second metal-atom layers. The hypothesis and calculations hence have found experimental support. a-Fe2Os(0001) Iron oxide and iron-based materials are very important catalyst used in many reactions such as the Fischer-Tropsch and the ammonia synthesises, in photoassisted electrolysis of water and in gas-sensor applications. The 3d 5 (magnetic) hematite Fe2Os is an insulator also, and its crystal surfaces [87-89] and their reactivity have been of considerable interest recently, but adsorption of copper onto the oxide surfaces has only recently been in the focus_of interest [90,91]. We have studied the clean and copper-deposited a-Fe203(1012) [90] and (0001) [91] by AES, EELS, LEED and HREELS. Figure 30 shows the HREEL spectrum from a clean a-Fe2Os(0001) surface obtained at room t e m p e r a t u r e at 600 off normal incidence. We observe the surface phonon frequencies 61,62 and 6s together with multiple and combination phonon losses. The crystal structure of aFe2Os is anisotropic, and we may expect [92] a loss function 6.3. C u r

P(r

- r 1 Im {'[sl (co) ~1 (co)+ 1] ~}

where ~l and ~1 are the dielectric functions parallel and perpendicular to the crystal c axis, respectively. Here, where we have 11 meV FWHM resolution, we can still see the 62 feature, and even the fourth multiple loss 463 is obtained. These modes are typical Fuchs-Kliewer (surface optical) phonons whose properties are mainly determined by bulk p a r a m e t e r s such as the transverse optical phonon frequency C0TO[93]. The results from calculations using the above function will compare with the experimental HREELS results.

507

v/ / '~ /

x

3,1os

...-,...

:3

2~''~S

I / //

~//

~

\

~

-4

~'~

4~ 3

.,, A,,

,:o:

I

\

I

0

I

100 200 Energy loss (meV)

I

300

I

400

Figure 30. HREELS from clean a-Fe2Os(0001). See text.

Upon depositing Cu on the (1• 1) a-Fe203(0001) substrate surface [91] the surface phonons are gradually changed (Figure 33) while at the same time a hexagonal rotationally aligned Cu (111) superstructure [91] appears and become gradually brighter and sharper until finally, after depositing for 210 s (~ 35 A) only the C u ( l l l ) structure is present. We note t h a t upon the initial deposition of copper a new vibrational feature was found at 18-20 meV (Figure 33). This vibrational peak was maintained even until 120 s deposition (- 20 A) and indicates an Cu-O interaction as judged from a comparison to calculated and experimental results [94-96], suggesting formation of a Cu(I) state at the interface, similarly as was found at RT also for the copper-alumina system [60,74]. This indicates a chargetransfer process as we saw earlier (Sects. 4.1 and 5.1). m

6.4. Cu r a-Fe203(1012) Also growth of Cu on cz-Fe2Os(1012) surfaces have been carried out [90]. The samples are artificially grown small crystal flakes and are more difficult to investigate experimentally t h a n the quite large (0001) growth faces of n a t u r a l haematite crystals. The LEED structure of the (1012_) surface depends upon the surface t r e a t m e n t [89]. Upon flash-annealing the (1012)surface in situ, Cu grows in a Volmer-Weber-type growth mode and no clear superstructure LEED p a t t e r n is observed for Cu deposited at RT. The substrate 1• 1 LEED p a t t e r n gradually becomes a t t e n u a t e d while the substrate plasmon and inter-band-transition energylosses attenuate, and the familiar low-energy interband-loss due to Cu appears.

508 7. COPPER ON PEROVSKITE STRUCTURF_~

Figure 31. LEED patterns from SrTi03(100) and segregatedcalcium superstructures after heating to 900~

Among the perovskite structures we will discuss another Ti(IV) compound, the ternary d o transition-metal cubic SrTiOs. It is, like other MaMbOs perovskite structures where the cation Mb is Ti(1V) and M a is bivalence cation, closely related in its properties to rutile, TiO2. Electronicallly [97,98], as demonstrated by RESPE [99], but also geometrically, since the atomically flat and nonpolar SrTiOs(100) surface may be terminated either by a Ti02 or by a SrO plane [100] (patches of both terminations will coexist, butpreparations may be carried out so that one of them dominates [99]). SrTiOs may for 2D surfaces experiments be considered as a simple cubic structure at RT since its distortions from the stable hightemperature cubic phase are very small. Of the perovskite clean crystal surfaces, only the (100) and (111) surfaces of SrTiOs and the (100) surface of BaTiOs have been investigated The HREEL spectrum from SrTi03(100) is thus well determined [101,102], but with regard to Cu and Ni adsorption only Cu has been reported [102,103] as deposited on the SrTiOs(100), in both cases stimulated by its qualities as a good matching substrate for the YBa2Cu3OT.x high-Tc oxygen-deficient perovskite superconductor of which high-quality epitaxial films had just been grown on that surface. The processing of synthesis of epitaxial YBa2CusOT.x films involves deposition on a 400-600~ SrTiO3(100) substrate followed by annealing in oxygen at 900-950~ Before we proceed with a discussion on the copper adsorption we will briefly discuss an impurity reconstruction that may take place just around the 900~ annealing temperature. In the synthesis of single-crystal SrTi03 it is difficult

509 completely to remove traces of calcium. Even though the bulk analytical content of Ca is only a few ppm in commercial crystals a detailed AES-LEED investigation has showed [104] t h a t Ca will segregate to the surface and change the originally 700~ sharp 1• 1 surface (Figure 31, upper) to a superstructure (Figure 31, mid) (which we attribute to a multiple scattering across the topmost layers) with about 3% Ca in the surface layer, or with only traces of (( 1% Ca to a p(2• reconstruction (Figure 31, lower) when the air-annealing temperature is 890~ and the humidity below 40%. Since the radius of a Ca 2§ ion is not very different from that of Sr 2§ we suggest t h a t the calcium ions exchange with the ions in the surface region, causing a change in the surface potential which induces the p(2• surface reconstruction (segregation of Ca to the surface of oxides has been studied also in the case of MgO, in which AES and LEISS results agree with a theoretical prediction based upon the surface heat of segregation [105]). A model for the occurrence of the p(2• structure could perhaps consist of (1) a combination of a relaxation of the Sr atom in every second unit cell in the topmost layer and (2) a segregation of ca atoms towards the second- or third-layer region, a model that will agree with our observation of the Ca-to-Sr and Ca-to-Ti ratios during development of the p(2• structure. In every second surface unit cell, the Sr atoms would then be displaced upwards, and downwards in the other half of the unit cell. The validity of the model could be explored in direct experiments on displacements of surface atoms. 7.1. Cu (m SrTiO3(100) In the two (independent) investigations [102,103] it was found t h a t Cu at RT grows in a Volmer-Weber type mode, individual clusters forming at the surface, and at high coverage a coalescence was indicated. For a deposits corresponding to 10 ML, AES results indicate [102] t h a t the copper growth and formation of islands are related to the TiO2 layer termination of the surface, and t h a t the copper growth mainly occurs on the 80% of the surface found to be TiO2 layers. No Cu epitaxial superstructures were found at the RT substrate. Growth of Cu onto a SrTiO3(100) substrate at 600 o has, however, shown [106] by RHEED (in synthesis of DyBa2CusOT.x films by MBE) t h a t copper was depositing primarily with (100) and (110) planes parallel to the interface. Here growth of metallic copper as islands was also indicated, and confirmed by AES during depositing a total of 150 A of Cu. The EEL spectrum from the clean SrTiOs(100) surface at RT shows [102] losses at 6.5, 10.2 and 13.1 eV, which are attributed to interband valence-band transitions between occupied O 2p bands and unoccupied Ti 3d levels [97], and a loss at 22.8 eV which can be assigned either to an O 2p-Ti 3d transition or to a plasmon loss (a bulk plasmon loss is suggested at 26.4 eV [97]), and a loss at 28.8 eV which may also be a plasmon loss. Upon deposition of Cu two new peaks appear at 4.5 and 7.0 eV [102]. These are assigned to an interband transition with Cu 3d as an initial state and to a copper plasmon loss, respectively. The plasmon losses and interband-transition losses are a t t e n u a t e d with increasing de,. The results from UPS (He(I)) show (Figure 32), for the clean surface, energy emission from O 2p levels at 5.2 and 7.3 eV. In the band gap there are some occupied states which are assigned to the He(I) "ghost peak", to carbon states or to Ti 3d states. The O 2p is gradually a t t e n u a t e d with dcu while the copper related to pure Cu 3d states appear as an impurity in the bandgap, close to the valence band maximum.

510

dcu( 10 .,,,-..

.E=

5

2 ..-,,,.

Z

1

0 I

Ef-O

1

1

I

-2 -4 -6 Binding energy (eV)

,.

1

-8

Figure 32. UPS He(I) from Cu/SrTiOs(100). Cu is growing within the bandgap (ref. 102).

The copper emission peak appears at a distance of 2.0 eV and 4.2 eV, respectively, from the main O 2p emission peaks. Within experimental error, no band bending is observed. From the UPS results we also obtain the change in work function A#, found to be 0.4 eV (with a quite large uncertainty of • 0.1 eV) during deposition of Cu until saturation just before dcu = 5 A at the # for pure Cu, 4.48-5.10 eV [107], based upon a measured ~ of 4.2 eV [108] for the SrTiOs(100). The positive A~ and the absence of banding indicate a weak interaction and only little charge transfer, and thus only very weakly bound copper to SrTiOs(100). XPS results [103] found no changes in 8r, Ti or O core levels with doll at RT surfaces, and no evidence of Cu-O compound formation, but instead indication of Cu cluster formation. Analysis of the Cu 2ps ~ binding energy showed a gradual shift of 0.2 eV in binding energy, associated with changing cluster size. By a sputter-profile analysis of a 100-A Cu overlayer the XPS furthermore showed

511 quickly appearance of substrate species and it took longer for the Cu 2p signal to disappear, suggesting an irregular substrate coverage by the coalesced Cu overlayer. No oxygen out-diffusion was observed. Annealing of the 100-/k Cu overlayer at 500~ led to significant changes in atomic distributions with sufficient oxygen diffusion to convert the metal overlayer to an oxide [102]. Also Sr outdiffusion was significant. The annealing caused further clustering in the Cu overlayers. The two investigations [102,103] hence show basically the same conclusion as to the behavior of the Cu deposits during growth at RT and annealing.

8. COPPER AND NICKEL ON WURTZITE STRUCTURF~ The only metal oxide crystallizing in the wurtzite structure is ZnO, but then it has been investigated very intensively for several years due to its high importance particularly in heterogeneous catalysis as substrate for copper, but also in gas sensor, varistor and solar cell applications. In the wurtzite structure the zinc ions are coordinated tetrahedrally with the oxygen ions. It is an insulator, E ~ 3 . 4 eV, Its electronic band structure is quite well investigated [109-111]. The valence band is mainly contributed by O 2p levels and the conduction band by Zn 4s and 4p levels. The Zn 3d band lies below the O 2p band, which is u n u s u a l in comparison to the band structure of other transition metal oxides. Electronic levels within the bandgap of clean ZnO surfaces have not been revealed; in stead, the location of states due to oxygen and zinc dangling bonds have been confirmed within the valence band and conduction band, respectively [ 1 12], corresponding to the 7.4-eV peak in the EEL spectrum [113]. It is conducting as ntype semiconductor in its stoichiometric form. It has equilibrium surface concentrations of intrinsic surface defects, interstitial Zn or oxygen vacancies, which deviate from corresponding bulk values, and extremely small concentrations of surface defects, easily obtained under UHV conditions, Figure 33. LEED from a thermally induce degenerate e t c h e d ZnO(0001) (F_v= 63 eV). accumulation layers

512 leading to "metallic surface conductivity" [ 114]. Investigations on the ZnO crystal surfaces have concentrated on the two hexagonal lo_w-index polar surfaces, Znterminated ZnO(0001) and O-terminated ZnO(0001), that have coor_dinatively u n s a t u r a t e d zinc or oxide ions, respectively, and the non-polar ZnO(1010)surface that has eq_ual numbers of zinc and oxide ions in dimer sites. Most recently, also the ZnO(1120) has been studied. Their electronic and structural properties is well known. Very recently, though, an ordered array of sub-sextets around each of the main hexagonal diffraction spots from a ZnO(0001) surface was observed by LEED after prolonged t h e r m a l t r e a t m e n t at 800 K, Figure 33 [115]. The origin of the observed sub-sextets may be understood in terms of diffracted reflection from a ZnO(0001) surface containing hexagonal pits created by a t h e r m a l etching during the annealing [ 115]. The chemistry, particularly toward CO, of Cu particles or thin films on these catalytically important low-index ZnO surfaces has been object for many investigations over the last decade [116], due to the great importance of the lowpressure synthesis of the basic chemical compound methanol (used for further synthesis of a large group of compounds and for direct fuel as substitute for gasoline) by use of binary Cu/ZnO or, usually, ternary Cu/ZnO/A12Os or Cu/ZnO/Cr20 s catalysts, from carbon monoxide and hydrogen (2H2+CO --> CHsOH), but also in the water-gas shift process (H20 + CO --> H2 + COx) and in methanol steam reforming (CH3OH + H90 -~ CO2 + 3H20). The Cu/ZnO system functions as a gas sensor, as well [117]. We will here concentrate on the geometry and electronic structure during growth of ultrathin layers of copper on the above ZnO crystal surfaces. 8.1. Cu ~m ZnO(0001) Growth of copper on the oxygen-terminated ZnO(0001) surface at RT was studied by XPS, AES, ISS and LEED [118] and also by UPS, XPS, Valence Band PES and REPES [116]. It was found that Cu grows, at least for the first ML (2 A) in a 2-D p(l• 1) overlayer mode (Cu atoms in 3-fold hollow sites created by the outer layer of oxygen atoms), followed by formation of rotational]y-aligned epitaxial 3D Cu(111) microcrystals, randomly distributed over the surface but not observable by LEED before about 3 ML (6-8/k) coverage [116,118]. Increasing temperature favors more agglomeration o[118], producing both sharpened overlayer and substrate patterns for a 8-A surface at 523 K (but only substrate-pattern sharpening for a 0.3-ML surface [116]). The growth can thus be described as following a Stranski-Krastanov-type mode for this surface. This is further supported by the ISS results. Very recently [ll9]_it was found t h a t at 130 K the Cu is cationic at tiny coverages at the ZnO(0001) surface, but becomes nearly neutral at coverages beyond a few percent, forming monolayer-clusters until about 505 surface is covered whereafter the Cu islands grow thicker without filling the gaps between the islands. By annealing to 850 K further clustering takes place. XPS found a mild decrease (about 0.6 eV) of the Cu(2p) binding energy with increasing dcu within the first ML, indicating that the Cu species are not completely metallic [116]. The PES at h v = 120 eV shows an increase in intensity at the top edge of the substrate 0 2p valence band near binding energies of 3-4 eV with increasing dcu, due to Cu 3d states which have a very high PES cross section relative to the O 2p levels at 120 eV, and the Cu 3d levels shift 0.4-0.5 eV to lower binding energy (relative to the ZnO features), similar to a relaxation shift m

513 observed for the core level. Furthermore it is noted that a well-defined Cu 4s peak is not observed, indicating that the low-coverage supported Cu is neither extensively oxidized (4s still present) nor atomic in nature and suggest that the 4s level is somewhat delocalized [116]. The Cu 2p core level binding energies and lack of shakeup satellites indicate that the highly dispersed copper is either Cu ~ or Cu § a point which has been controversial for quite some time, and which is important for elucidating the complex mechanisms in the above synthetic processes. By REPES it was found [116] that the relative energy splitting of the dispersed Cu is not significantly different t h a n in Cu metal, showing that the dispersed copper is Cu ~ and not Cu§ However, detailed analysis of at the satellite peak at the Cu 3p --->4s edge shows that the dispersed copper displays unique REPES features showing t h a t the copper is not purely metallic, atomic or oxidized [116]. The deposition resulted initially in downward, followed by upward, band bending and similarly in the work function for this surface (leading finally to the Cu bulk value). The downward part is indicative of charge donation from the surface to the bulk. However, due to the low number of surface sites expected to be depleted by the Cu, this is expected to have only little effect on the chemical nature of the copper overlayer atoms [116]. The contact potential difference between the Cu and ZnO creates a Schottky barrier of 0.5-0.6 eV as measured by the total change in band bending. 8.2. Gu r ZuO(O001) In the growth of Cu on the_zinc-terminated ZnO(0001) there are many similarities with the Cu/ZnO(0001) case. The similarity of the spectroscopic results (Cu 2psa final-state relaxation shifts and the narrowing of the Cu 2psa peak with dcu, and in the Zn 2pa a normalized intensity ratios) thus point to the same growth mode for at least the first monolayer.However, no Cu overlayer structure was observed. Furthermore, the Cu 3d levels in a He(II) experiment appear at the top edge of the oxide 2p Band, not overlapping the oxide band as strongly as in the Cu/ZnO(0001) case. In REPES it is found again that the copper is not purely metallic, atomic or o_xidized and the data for 0.3-ML Cu is almost indistinguishable from the ZnO(0001) surface. The band bending is only upward and the work function only increasing with dcu, however [116]. Deposition of 0.3 ML of Cu onto the above thermally etched ZnO(0001) surface leads [120] to elimination of the sub-sextet structure and to a broadening and attenuation of the integral diffraction spots, and a new hexagonal diffraction appeared gradually appeared near a coverage of 2 ML and it became predominant at an average of 6 atomic layers almost completely covering the surface, showing an epitaxial Cu(111) ordered island formation, with an observed ratio of the unitcell dimensions in reciprocal space acu/az~o - 1.24 which is slightly smaller t h a n the similar (1.27) p a r a m e t e r for bulk Cu. In a low-energy spectroscopy experiment [120], the target current spectrum [6] (S(E) = dJ(E)/dE vs. incident energy E) at perpendicular incidence from the clean ZnO(0001) surface gives a fine structure, where the maxima may be analyzed in terms of the matching method for determining elastic reflection coefficients [6]. The location of the extrema coincides with the location of band-structure critical points in the Brillouin zone for empty electronic states [ 121,122]. The results demonstrates the expected agreement with calculated density-of-states, DOS between the experimental maxima and the DOS-

514 structure in corresponding critical points (L,M,F,A). A higher-energy m a x i m u m is connected with reflection variation due to diffraction from thermally etched pits [115]. From the TCS we also obtain (peak A, the vacuum-level position of the sample) the work function ~ of the ZnO(0001) surface to be about 3.5 eV, in agreement with the literature value [114]. With Cu deposition of 0.3 ML of Cu we obtain A~ = 0.5 eV and with 6 Cu layers ~= 4.5 ev, i. e. smaller t h a n ~ for bulk C u ( l l l ) , thus showing that the surface is not completely covered by C u ( l l l ) islands. The patch-like coverage is also indicated by the broadening of the primary peak. The dispersed copper on the (0001) surface is readily incorporated into the ZnO lattice as a Cu § site with strong CO chemisorption abilities and is, therefore, a likely possibility for high catalytic activity [116]. 8.3. Cu (m ZnO(1010)

19

1 2

...-,,..

c

t'~

3 v

4 e(D r" m

5 I 40

I 30

I 20 Energ/

1 10 loss

I 0 E L(eV)

Figure 34. EELS from Cu/ZnO (1010). eV. See text (ref. 13).

Ep=97

Deposi_tion of Cu on the ZnO(1010) surface was studied by EELS over the 0 < dc, <- 60 A r a n g e at RT [13]. The most striking results are those from the submonolayer region which is shown in Figure 34. Spectrum 1 is from a clean surface annealed in 10 .5 Tort 02 at 320 ~ and spectra 2 and 3 are obtained after deposition of 0.5 and 1/k of Cu, respectively, on the ZnO(1010) surface of (1). Spectrum 4 shows the clean surface yields loss features at 1.9, 4.0, 7.4, 8.8, 12.5, 18,2 and 24.5 eV. These agrees with previous work [109], except for the 1.9 eV peak which was unfortunately considered as induced by the effect of double derivatives of the loss signal. We note t h a t the 1.9-eV loss structure Vstill remains after the heat t r e a t m e n t in O2 (as did the corresponding 2.6-eV loss structure for the MgO(001) surface [13]). This experimental evidence strongly discards the 1.9-ev loss structure (and the 2.6

515 -eV loss for MgO(001)) as candidates for F, § centers [21], but r a t h e r suggests t h a t a surface V, center, a hole trapped at a metal ion vacancy with 3 oxygen ligands (and 5 oxygen ligands in the MgO(001) case) as the physical origin of the loss structure (although it by no means implies t h a t F, § centers do not exist on oxide surfaces). Upon deposition of 0.5 A of Cu a new structure at 2.0 eV arises, and further deposition of 0.5 A of Cu brings up a weak peak at 4.3 eV (similarly appear peaks at 2.2 and 4.3 eV, respectively). We may ask, w h a t is the nature of these peaks. It is not from the transition Cu 3d ~ 4s since in EELS studies on clean Cu, the loss structure was found. The loss feature disappears from the (1010) system upon deposition of 3.7/k, and it does not appear in the spectrum from a continuous Cu thin film [13]. To elucidate the cause of the peak, the 1-A Cu/ZnO(1010) surface was heated (spectrum 4 in Figure 34) [13], and the peak became enhanced, consequently negating the possibility t h a t the 2-eV peak has originated from small Cu clusters, which have electronic properties different from copper (the peak should have decreased in intensity in case of a very small cluster (it has been reported t h a t bulk metal properties may arise even from clusters consisting only 4 copper atoms [ 123]). In a test experiments of deposition of similar small amounts on a clean Si(111) surface, no 2-eV loss was seen. We believe t h a t the 2-eV loss is the well known loss structure of the semiconductor Cu20, corresponding to a threshold transition from the Cu 3d valence band to the Cu 3s conduction. To support this interpretation of the 2-eV peak, a 1-/k-Cu-deposited surface was exposed to 02 (spectrum 5, Figure 34). This exposure decreased the intensity of the 2-eV peak (and in the MgO case elimination of it completely), confirming t h a t the copper, which corresponds to the 2-eV loss peak, is in the oxidized state. In ZnO where the bandgap is 3.4 eV, the Cu 3d impurity levels overlap the 0 2p band. A Cu 4s level, which has donated one electron, therefore superimposes the conduction band. Hence the 4.3-eV peak of the 1-A Cu/ZnO(1010) probably originates from a copper cluster, as also confirmed by the oxydation experiment (and by the peak-growth at higher coverages [13] (in the MgO case, the 4.3-eV peak may be due to the O 2p --->Cu 3d). We have therefore concluded t h a t the initially deposited copper (at the low end of the sub-monolayer coverages) exists in an ionized state and is bonded to oxygen ions in the surface of the substrate, and t h a t this state probably is the most stable one, Cu(I), corresponding to the induced 2-eV loss peak, and also t h a t the 1.9-eV peak originates from electronic resonance of V, centers acting as copper-deposit trapping-interaction centers. As for the ZnO(0001) no Cu-overlayer superstructure has been observed, but the similar photoelectron spectroscopic evidence as for the (0001) surface was found upon Cu deposition [116]. At 15-20/k coverage a new set of diffraction spots were observed t h a t appeared to be rotationally aligned with the rectangular substrate pattern, but the spots were too diffuse to determine the precise symmetry. The Cu 3d levels develop much the same as on the (0001), with the levels appe_aring at the top edge of the oxide 2p band and not overlapping as strongly as (0001). The band bending and the work function changed with dc~ as for the (0001); thus also for this surface a small charge donation from the surface to the bulk is indicated [116]. Dispersed copper is also on this ZnO(1010) surface, as on the zinc-terminated ZnO(0001) surface, readily oxidized and annealed into the ZnO lattice as a Cu(I) site, a coordinatively u n s a t u r a t e d Cs~ site (that is the only copper center found to

516 adsorb CO w i t h high affinity, 88 kJ/mol, and the dispersed atomic Cu and Cu clusters chemisorb CO with approximately the same affinity as copper m e t a l [116]). UPS He(II) e x p e r i m e n t s carried out in the low 110-130 K t e m p e r a t u r e range [116] have indicated t h a t highly dispersed copper on the ZnO(0001) a n d (1010) surfaces chemisorb CO w i t h a b o u t the same affinity as copper metal, while chemisorption on the Cu/ZnO(0001) was m u c h weaker, a n d high-affinity Co chemisorption, often associated w i t h the catalytic active site, was shown to occur at a coordinatively u n s a t u r a t e d t e d r a h e d r a l Cu § site created on the Cu/ZnO(0001) surface upon a n n e a l i n g in oxygen, and furthermore, chemisorption to the Cu § site p e r t u r b s the CO electronic sstructure m u c h more t h a n chemisorption to either Cu ~ or Zn 2§ due to a stronger a and ~ interactions as indicated from valence-band PES. A Cu-CO active site m u s t provide a lower energy p a t h w a y in the synthesis of m e t h a n o l t h a n a similar Zn2+-CO complex to bring about the k n o w n lower activation b a r r i e r in this synthesis, but u n t i l a definite m e c h a n i s m has been elucidated for this reaction, the n a t u r e of copper activation will be dab_ated [116]. L a t e r [124], in a TPD investigation of CO adsorption on Cu/ZnO(0001) at 130 K, a strong CO TPD-peak at about 160K was found for thicker Cu films (or films t h a t have been a n n e a l e d to high t e m p e r a t u r e s to induce 3D clustering), characteristic for adsorption sites t h a t are Cu(111)-like, and loss in CO adsorption capacity was found to be not as great as rthe loss in Cu surface area. This was i n t e r p r e t e d [124] as caused by a CO-induced redispersion of 3D Cu clusters into 2D_islands. A TPD investigation was recently carried out on the Cu/ZnO(1010) surface as well [125]. At RT, submonolayer a m o u n t s were found to enhance the ability of ZnO to adsorb CO drastically. A small a m o u n t of C02 desorbed from the surface as a result of CO reaction w i t h surface oxygen. A desorption order of 2 a n d a desorption activation energy of a_bout 60 kJ/mole were obtained for CO desorption from the Cu-deposited ZnO(1010) surface, which is in good a g r e e m e n t w i t h the binding energy of CO adsorbed on Cu(100) surfaces. Also the effect of light was investigated on one of these lowest-indexed surfaces .......

e-

=. 1.0 r

= 0.5 m

~PJ

i

0

I

I

i

I

I

500 Time of flight, t

i

I

(ps)

i

I

1000

Figure 35. Time-of-flight distributions of CO2 from CO/Cu/ZnO(0001). (1) clean ZnO(0001); (b) dcu = 0.3 ML (ref. 126).

517 [126]. In the adsorption at RT of CO Cu/ZnO(0001) a double-component time-offlight distribution in laser-induced desorption of CO2 was observed (Figure 35) from CO adsorbed on ZnO(0001), and a strong increase of the a m o u n t and a lowering of the threshold for C09 desorption was caused by a 0.3-ML Cu deposition. m 8.4. Cu cm Z n O ( l l 2 0 ) The growth of Cu on ZNO(1120) and the related c h a n g e s in the electronic structure were investigated for deposits until dcu= 40 A [127]. At RT, the growth shows initially a linear decay with dc,, for dc, < 1.2 A, and then an exponential decay, for 3 < dc, < 40 A (Figure 36), agreeing with the monolayer-simultaneousmultflayer (MSM) model (as classified by Argile and Rhead [128]) for thermodynamic unstable growth with a pre-mono]ayer breakpoint. The model is close to a Volmer-Weber type growth mode, i.e. "islanding". The curve for ideal growth without a pre-monoloayer breakpoint, shown as a thin line in Figure 39, exhibits a break at dc,- 2.0 A (it could perhaps be argued t h a t two breaks may be seen prior to the exponentially decaying curve, but that conclusion would lead to drastic discrepancies in the attenuation lengths of the layers). We consider the 2 A to be the deposition thickness corresponding to monolayer coverage. This growth mode is similar to wh_at very recently has been found for Cu growth on the oxygenterminated ZnO(0001) surface [ 119], and to some other meta]/meta]oxide surfaces as well [128,11,102], and one notices that the pre-monolayer breakpoint obtained by Ernst et al. [119] for the (0001) surface practically coincides with the present value. We did not see any Cu superstructure, only gradual attenuation of the substrate pattern, in agreement with the MSM growth mode. The EELS results for the Cu growth on this surface [127] agrees with our previous results [13]_for Cu growth on the ZnO(1010) surface at RT. Coverage-dependentinterband 6O transitions were allocated at 2.2+0.1 and 4.2+0.1 eV, and a Cu-related plasmon increases in energy from 6.7 to 7.0 eV itq~~ 40 with dcu. The plasmon ""40 ~176 frequency COp did not for this surface show a significant shift with dcu at RT for dcu < 6/k. -20 . . . . . At higher coverages, COp approaches the 7.1-eV value of clean Cu, and this behavior at Oi II I I I I I I RT we believe is primarily 0 10 20 30 40 reflecting the surface-to-bulk dcu (A) development of the copper film. To further elucidate the Figure 36. Change of O(KL2sL2s) intensity with behavior of the Cu deposits, dcu from growth of Cu on ZnO(1120) at 300 K.. synchrotronInsert shows submonolayer range (ref. 127).

t/

1

518

1.2 ~Efi 0.8

D

9

,. 0.4

.._.,.

-..

....

----.-

& I

,

I

.

I

4

I

I

I

8

dcu (~.1

I

12

__

--_.~ I

I

16

Figure 37. Cu/ZnO(1120). Changes with dcu in initial- and final-state contributions to AEB(Cu 2psi) (ref. 127).

6.8 ~

,

,

,

i

,

6.4 6.0 ...-..

56 t,~

~. 4.8 "

Ill

4.4

B

9

2.2

A 9

1.8

3

~

4

dcu (A)

~

9

5

.

6

Figure 38. EELS energy shifts with dcu from a 6-A Cu/ZnO(1120) when heated to 875 K. Ep= 98 eV. (Ref. 127).

radiation based investigations were carried out [127] for depositions over the 0 < dcu < 18 A r a n g e . Core-level spectroscopy showed EB(Cu 2Psi) reaching the bulk Cu value at dcu= 17.8/k., and the Auger Ekin(Cu LsM4,fM4,5) reached similarly the bulk Cu value after a monotonic change of about 4 eV, and we noted a narrowing of this peak with increasing dcu, in agreement with an increasing cluster diameter. This observed shift in Auger kinetic energy is given by the difference between the energy shifts of the one-hole initial state and the two-hole final state. A r a t h e r strong increase in E Bwith dcu may be expected from very small supported metal clusters, however, because of additional finalstate effects AE~ compared to the bulk metal reference state [129]. If we assume t h a t the levels involved in the transition are localized and t h a t the screenings of the first and second hole are similar [129], then the change in the Auger p a r a m e t e r Aa= AF~n(]kl) + AEB0) is related to the shift in the final state by which also holds for metallic clusters. Figure 37 shows the changes AEi, (/) and AE~(]) with dcu, and we see, perhaps with exception of the lowest end of the coverage range, t h a t the observed shift entirely is of final-state nature. For this surface we therefore conclude t h a t the combined stability of the initial-state shift and the plasmon

519 frequency indicate t h a t the shift is due primarily to a size effect, and t h a t the combined use of the two effects is very valuable in elucidating the behavior of ultrathin film electronic properties. If we anneal the Cu-deposited ZnO(ll2) surface, however, a quite different behavior occurs in the electronic structure in the Cu-deposited film with increasing dcu. By annealing a 6-A Cu/ZnO(1120) surface to 875 K for 45 rain at 4• s Pa we obtain [127] a sharp Z n O ( l l 2 0 ) - ( l • LEED p a t t e r n appears, and AES analysis shows a signal corresponding to only 0.2 ML, consistent with a strong agglomeration of the copper particles. In the corresponding EELS we observe [ 127] clear copper features with a strong energy loss at 5.7 eV and absence of both the Cu plasmon at 6.7 eV and the oxygen dangling-bond surface loss at 7.5 eV. Upon deposition of Cu onto this surface, a shift of the 5.7-eV peak (Figure 38, curve C) toward higher energy is observed, and for dcu > 5 ,~ the loss energy is equal to the COpvalue of the non-annealed surface. The free-electron plasmon relation applied to this shift gives a charge transfer of 0.28 e per adatom which is exactly the numerical Cu --~ O charge transfer obtained theoretically for Cu20 [130], and it is therefore reasonable to consider the 5.7-eV loss peak as originating from Cu(I) species at the surface or at interstitial or substitutional positions in the top layer. The considerable agglomeration of Cu and the absence of the O dangling bond cause us to suggest substitutional O-diffusion in the surface at elevated temperature, leading to vacancies in the substrate and oxydation of Cu. This is somewhat in contradiction to the work of Didziulis et al. [116] and Ernst et al. [119] who both reported metallic Cu at elevated temperatures. 8.5. Ni on ZnO(0001) a n d ZnO(0001) Contrary to the Cu/ZnO systems, the Ni/ZnO systems have not been controversial. The deposition of nickel onto ZnO surfaces have been much less studied, perhaps because of the smaller range in applications. An investigation by AES and UPS was carried out already early [1_31] on the deposition over the 0.05 to 5 ML of Ni on the ZnO(0001) and ZnO(0001) surfaces at RT. With the sample at RT, the Ni film is found to grow in layer form on both surfaces. The width of the Ni 3d band as seen in He(I) and He(II) UPS has been developed when d~i has reached 1 ML, indicating that the dispersion of the 3d bands for a 2-D Ni film is as great as for 3-D bulk Ni. Peaks at 6 and 4.3 eV is tentatively interpreted as due to Ni chemisorbed to oxygen from the substrate, and peaks at 2.1 and 3.1 ev are attributed to Ni atoms and 2-D Ni clusters, implying a relaxation shift of 2.5 eV for the Ni atom on the Zn-terminated surface. On the oxygen-terminated ZnO(0001) surface an upward band bending of 0.6 eV with respect to the fiat band case is found after Ni deposition to compensate for the contact potential difference. On the zinc-terminated ZnO(0001) surface, nearly no band bending was found.. This is explained by a dipole layer between the surface Zn atoms and the Ni atoms of the first layer having the Ni atoms negatively charged. On the oxygen-terminated face, and not on the zinc-terminated face, oxygen diffusion into the Ni layer is observed, which could be facilitated and perhaps becomes possible only by this band bending.

520 9. COPPER ON FLUORITE METAI~OXIDE S T R U C r U R K S Last we consider the class of fluorite (CaF2) structures. Among the fluoritestructured metal-oxide crystal surfaces only U02 and yttria-stabilized Zr02 have been investigated with regard to metal adsorption., and so far none with Cu or Ni. We have recently investigated Cu adsorption onto yttria-stabilized Zr02 [132]. A pure Zr02 crystal is monoclinic at RT. At very high t e m p e r a t u r e s (2640-2950 K) a bulk cubic phase exist. This cubic phase can, however, be stabilized to lower t e m p e r a t u r e s by addition of certain oxides, such as Y20s (yttria), CaO and MgO, to the pure ZrO2. Yttria-stabilized cubic Zr02, YSZ, has been found to be a suitable substrate for silicon-on-insulator device technology [133] where it may well replace silicon-onsapphire (SOS) devices. It is a defect solid (Y20s)m(ZrO2)l.m solution which maintains a cubic structure over the 0.08 < m < 0.4 range, and the lattice constant increases almost linearly with m from 5.13 to 5.18 A [134]. The defects are oxygen vacancies created to preserve lattice neutrality when y+s ions are substituted for Zr 4§ ions in the CaF2-type structure. These vacancies give rise to high oxygen ion mobilities, and various oxygen sensors and solid electrolytes made of stabilized ZrO2 are based upon this ion-conducting property. F u r t h e r m o r e it is used as a substrate or as bufferlayer for YBa2CusOT. ~ superconductor thin films and as a support for Rh in heterogeneous catalysis such as in synthesis of ethanol from syngas. There are contradictory assignments of the electronic bandstructure scheme in the literature. There has been studies on polycrystalline surfaces, most recently by Wiemhffer et al. [135], and the reported bandgap varies from 3.7 eV [136] to 7.2 eV [137]. We have studied the clean yttriastabilized ZrO2(100) surface with the common composition m = 0.10, using AES, EELS and LEED, and followed the growth of Cu onto t h a t surface over the 0 _< dcu < 50 /k deposition range [132]. The surface was initially cleaned by heat t r e a t m e n t to 1370 K, causing Zr some loss of oxygen from the surface layers (without heati i I 1 t r e a t m e n t to high t e m p e r a t u r e it 220 200 180 160 140 was not possible to obtain LEED Energy loss (eV) p a t t e r n s due to charging). Due to strong overlap of the Zr and Y Figure 39. Core-level EELS from MNN-transitions, core-level EELS yttria-stabilized cubic ZrO2(100). is preferred (Figure 39) for Ep= 514 eV (ref. 132). identification for YSZ [138], although further refinement is ..-....

e~

u..I

...-... LLI

-...... Z

521 1.o . O-KLL 9 Zr-MNN

0.8

-= 9 0.6

OlD~ ~

~

g

.,-

o

0.4

O

z

0.2

0

"

0

10

"

20

30

-

40

50

Depositionthicknessdcu (,~,) Figure 40. Changes in O(KLL) and Zr(MNN) normalized Auger intensities with dcu for yttria-stabilized ZrO2(100) (ref. 132).

needed to extend it into a quantitative method for surface composition analysis. As seen from the normalized AES Zr(MNN) and O(KLL) exponential attenuations with dcu (Figure 40), the growth at RT follows the MSM mode (due to low mobility of the deposited Cu the mode here basically belongs to the Volmer-Weber mode, i. e. pure "islanding"). The LEED pictures show that the surface is of a p ( l • simple bulk-truncation structure.The EELS from the clean surface exhibited strong loss features at 7, 14.5, 26, 34 and 40 eV, respectively. The assignment of the 14.5-eV loss is controversial in the literature since it has been interpreted either to a volume plasmon [139,140], to an interband of 2p to 3s states of oxygen in the conduction band [141] or to a collective excitation of more ligand np electrons in the valence band towards ligand (n+l)s states in the conduction band [142], coinciding somewhat with ref. 138. We therefore performed a structure calculation for an YSZ(100) cubic crystal, with Y doped uniformly in the lattice, and obtained a plasmon energy cop of 21.0 eV (assuming t h a t the valence electrons include 2p electrons of O and 4d and 5s electrons of Zr and Y). Analyzing the EEL spectra we find t h a t the energy loss at 9.5 eV decreases with increasing Ep, and disappears for E, > 250 eV, and this is believed to be one of the characteristics of a surface plasmon. This point supports an assignment of the 14.5-eV loss to a volume plasmon. Furthermore, an analysis of the dielectric function for Zr02 shows t h a t the 14.5-eV loss has character of a collective excitation of valence electrons, hence the assignment to a single-electron transition is not reasonable. Upon Cu deposition, the 14.5-eV loss shifts downwards in energy, for dcu > 6 A, and the relative intensity of the 7.8- and 10.0-eV losses changes drastically. The losses at 7.0 and 10.0 eV are clearly identified as due to the well-known Cu surface and bulk plasmons, respectively [132].

522 The electronic structure of the clean yttria-stabilized Zr02(100) hence has been more clarified, a LEED pattern has been obtained from this surface, the changes caused by Cu deposition have been elucidated, and the mode for the growth of copper clusters on the surface at room temperature has been determined.

Acknowledgements The cooperation with my coworkers in this work, I. Alstrup, J.E.T. Andersen, B. Ealet, Q. Ge, E. Gillet, L. Gui, Q. Guo, J.-W. He, S.A. Komolov, E.F. Lazneva, F. Matthiesen, J. Nerlov, H.N. Waltenburg and M.-C. Wu, the skillful drawings by H. Vib~k, and the support from the Danish Natural Science Research Council through Center for Surface Reactions and through a synchrotron-radiation research grant, the Carlsberg Foundation, the Thomas B. Thrige Foundation and the Elkraft Corp. are gratefully acknowledged. I thank the Research Institute of Electronics for providing excellent conditions and for fruitful comments by Y. Fukuda during a sabbatical stay at Shizuoka University, Japan.

Referenoe~ *) Present address: Research Institute of Electronics, Shizuoka University, H a m a m a t s u 432, Japan. 1 V.E. Henrich, Rep. Progr. Phys. 48 (1985) 1481. 2 L.C. Dufour and M. Pedereau, in: Surface and Near-Surface Chemistry of Oxide Materials, eds. J. Nowotny and L.-C. Dufour, Elsevier, Amsterdam, 1987. 3 P.A. Cox, Transition Metal Oxides: An Introduction to Their Electronic Structure and Properties, Clarendon Press, Oxford, 1992. 4 J.E.T. Andersen and P.J. Moller, Surf. Sci. 258 (1991) 247. 5 G. Ertl and J. Kiippers, Low Energy Electrons and Surface Chemistry, VCH Publ., Weinheim, 1985. 6 (a) S.A. Komolov, Total Current Spectroscopy of Surfaces, Gordon and Breach, Philadelphia, 1992. (b) P.J. Moller and M.H. Mohamed, Vacuum 35 (1985) 29. 7 T.M. French and G.A. Somorjai, J. Phys. Chem. (1970) 2489. 8 P.J. Moller and J.-W. He, Nucl. Instrum. Meth. Phys. Res. B 17 (1986) 137. 9 J.E.T. Andersen and P.J. Moller, J. Vac. Sci. Technol. A10 (1992) 497. 10 W. Wei, J. Vac. Sci. Technol. A6 (1988) 2576. 11 J.-W. He and P.J. Moller, Surf. Sci. 178 (1986) 934. 12 I. Alstrup and P.J. Moller, Appl. Surf. Sci. 33/34 (1988) 143. 13 J.-W. He and P.J. Moller, Surf. Sci. 180 (1987) 411. 14 H.M. Neergaard, MSc thesis (in Danish), Univ. Copenhagen, 1992. 15 P.J. Moiler and H.N. Waltenburg, to be publ. 16 M.-C. Wu and P.J. Moller, Surf. Sci. 279 (1992) 23. 17 D.G. Lord and M. Prutton, Thin Solid Films 21 (1974) 341. 18 K. Takayanagi, K. Yaki and G. Honjo, Thin Solid Films 48 (1978) 137. 19 T. Conrad, J.M. Vohs, P.A. Thiry and R. Caudino, Surf. Interface Anal. 16 (1990) 446.

523 20 C. Noguera, J. Goniakowski and S. Bouette-Russo, Surf. Sci. 287/288 (1993) 188. 21 J.-W. He and P.J. MoUer, Chem. Phys. Lett. 129 (1986) 13. 22 N.C. Bacalis and A.B. Kunz, Phys. Rev. B 32 (1985) 4857. 23 C. Benndorf, H. Caus, B. Egert, H. Seidel and F. Thieme, J. Electron Spectrosc. Related Phenomena 19 (1980) 77. 24 (a) M. Prutton, J.A. Ramsey, J.A. Walker and M.R. Welton-Cook, J. Phys. C: Solid State Phys. 12 (1979) 5271. (b) M. Prutton, J.A. Walker, M.R. Welton-Cook, R.C. Felton and J.A. Ramsey, Surf. Sci. 89 (1979) 95. 25 A.R. Protheroe, A. Steinbrunn and T.E. Gallon, J. Phys. C: Solid State Phys. 15 (1982) 4951. 26 E.V. Stephanova, V.S. Stepanyuk, M.N. Rogaleva, O.V. Farberovich, A.A. Grigorenko and V.V. Mikhailin, Sov. Phys. Solid State 30 (1988) 1329. 27 R.P. Furstenau, G. McDougall and M.A. Langell, Surf. Sci. 150 (1985) 55. 28 Q.-G. Zhu, A.-D. Zhang, E.D. Williams and R.L. Park, Surf. Sci. 172 (1986) 433. 29 Y.W. Chung, W.J. Lo and G.A. Somorjai, Surf. Sci. 64 (1977) 588. 30 W. GSpel, J.A. Anderson, D. Frankel, M.Jaehnig, K. Phillips, J.A. Schfifer and G. Rocker, Surf. Sci. 139 (1984) 333. 31 M.H. Mohamed, H.R. Sadeghi and V.E. Henrich, Phys. Rev. B 37 (1988) 8417. 32 G.N. Raikar, P.J. Hardman, C.A. Muryn, G. van der Laan, P.L. Wincott, G. Thornton and D.W. Bullett, Solid State Commun. 80 (1991) 423. 33 A.K. See, M. Thayer and R.A. Bartynski, Phys. Rev. B 47 (1993) 13722. 34 P.J. Moller and M.-C. Wu, Surf. Sci. 224 (1989) 265. 35 C.C. Kao, S.C. Tsai, M.K. Bahl, Y.W. Chung and W.J. Lo, Surf. Sci. 95 (1980) I. 36 J. Nerlov, Q. Ge and P.J. Moiler, to be published. 37 Z. Zhang, S.-P. Jeng and V.E. Henrich, Phys. Rev. B 43 (1991) 12004. 38 S.C. Wang and G. Ehrlich, Phys. Rev. Letters 62 (1989) 2297. 39 U. Diebold, J.-M. Pan and T.E. Madey, Surf. Sci. 287/288 (1993) 896. 40 M.-C. Wu and P.J. Moller, Surf. Sci. 224 (1989) 250. 41 S.A. Lindgren and L. Walld~n, Phys. Rev. B 22 (1980) 5967. 42 M.-C. Wu and P.J. Moller, Phys. Rev. B. 40 (1989) 6063. 43 M.-C. Wu, Q.-L. Guo and P.J. Moiler, Vacuum 41 (1990) 1418. 44 M.-C. Wu and P.J. Moller, Chem. Phys. Lett. 171 (1990) 136. 45 R.G. Egdell, S. Eriksen and W.R. Flavell, Solid State Commun. 60 (1986) 835 46 M.-C. Wu and P.J. Moiler, Surf. Sci. 235 (1990) 228. 47 N.R. Avery, Surf. Sci. 111 (1981) 338. 48 K.K. Kleinherbers and A. Goldman, Surf. Sci. 133 (1983) 38. 49 M.A. Vannice, J. Catal. 44 (1976) 152. 50 H. Onishi, T. Aruga, C. Egawa and Y. lwasawa, Surf. Sci. 233 (1990) 261. 51 M.-C. Wu and P.J. Moiler, in The Structures of Surfaces Ill (Springer Ser. Surf. Sci. vol 24), eds. S.Y. Tong, M.A. Van Hove, K. Takayanagi and X.D. Xie, Springer-Verlag Berlin, Heidelberg 1991, p. 652. 52. M.-C. Wu and P.J. Moiler, Surf. Sci. 250 (1991) 179. 53 (a) H.H. Madden, J. Kiippers and G. Ertl, J. Chem. Phys. 58 (1973) 3401.

524 (b) K. Christmann, 0. Schober and G. Ertl, J. Chem. Phys. 60 (1974) 4719. 54 C.-C. Kao, S.-C. Tsai and Y.-W. Chung, J. Catal. 73 (1982) 136. 55 V.E. Henrich, G. Dresselhaus and H.J. Zeiger, Phys. Rev. Lett. 36 (1976) 1335. 56 S. Bourgeois, D. Diakit~, F. Jomard, M. Perdereau and R. Poirault, Surf. Sci. 217 (1989) 78. 57 A. Benninghoven, Z. Phys. 230 (1970) 403. 58 S. Bourgeois, F. Jomard and M. Perdereau, Surf. Sci. 249 (1991) 194. 59 S. Ciraci and I.P. Batra, Phys. Rev. B 28 (1983) 982. 60 P.J. Moller and Q. Guo, Thin Solid Films 201 (1991) 267. 61 C.C. Chang, J. Appl. Phys. 39 (1968) 5570. 62 S. Xia, C. Guo, L. Lin and D.E. Ellis, Phys. Rev. B 35 (1987) 7671. 63 J. Olivier and R. Poirier, Surf. Sci. 105 (1981) 347. 64 E. Gillet, B. Ealet and J.L. Berlioz, Surf. Interface Anal. 16 (1990) 461. 65 Q. Guo and P.J. Moller, Surf. Sci. 244 (1991) 228. 66 C.A.M. Mulder and J.T. Klomp, J. Phys. (Paris) 46 (Suppl. C4) (1985) 111. 67 G. Katz, Appl. Phys. Lett. 12 (1968) 161. 68 G. Katz, J. Mater. Sci. 5 (1970) 736. 69 A.G. Schrott, R.D. Thompson and K.N. Tu, MRS Symposia Proceedings, vol. 60, Materials Research Society, Pittsburgh, 1986, p. 331. 70 J.E.E. Baglin, A.G. Schrott, R.D. Thompson, K.N. Tu and A. Segmiiller, Nucl. Instrum. Methods Phys. Res. B 19/20 (1987) 782. 71 K.H. Johnson and S.V. Pepper, J. Appl. Phys. 53 (1982) 6634. 72 K. Nath and A.B. Anderson, Phys. Rev. B 39 (1989) 1013. 73 Q. Guo, P.J. Meller and L. Gui, Acta Phys. Polon. 81 (1992) 647. 74 Q. Guo and P.J. Moller, Vacuum 41 (1990) 1114. 75 M. Gautier, J.P. Dureaud and L. Pham Van, Surf. Sci. 249 (1991) L327. 76 S. Varma, G.S. Chottiner and M. Arhab, J. Vac. Sci. Tech. A 10 (1992) 2857. 77. J.G. Chen, M.L. Colaianni, W.H. Weinberg and J.T. Yates, Jr., Surf. Sci. 279 (1992) 223. 78 J.T. Klomp, in: Ceramic Microstructures '86, Role of Interfaces, eds. J.A. Pask and A.G. Evans, Plenum, New York, 1988, p.307. 79 Q. Zhong and F.S. Ohuchi, J. Vac. Sci. Technol. A 8 (1990) 2107. 80 V. Vijayakrishan and C.N.R. Rao, Surf. Sci. Lett. 255 (1991) L516. 81 A.D. Zdetsis and A.B. Kunz, Phys. Rev. B. 32 (1985) 6358. 82 D.W. Goodman, R.D. Kelley, T.E. Madey and J.T. Yates, Jr., J. Catal. 63 (1980) 226. 83 J.G. Chen, J.E. Crowell and J.T. Yates, Jr., Surf. Sci. 185 (1987) 373. 84 B. Ealet, E. Gillet, V. Nehasil and P.J. Moller, to be published. 85 M. Ricci, thesis, Univ. d'Aix-Marseille III, 1991. 86 C.C. Kao, S.C. Tsai, M.K. Bahl, Y.W. Chung and W.J. Lo, Surf. Sci. 95 (1980) 1. 87 R.L. Kurtz and V.E. Henrich, Surf. Sci. 129 (1983) 345. 88 C. Sanchez, M. Hendewerk, K.D. Sieber and G.A. Somorjai, J. Solid State Chem. 61 (1986) 47. 89 R.J. Lad and V.E. Henrich, Surf. Sci. 193 (1988) 81. 90 P.J. Moller and F. Matthiesen, to be published. 91 Q. Guo and P.J. Moiler, to be published.

525 92 M. Liehr, P.A. Thiry, J.J. Pireaux and R. Caudano, J. Vac. Sci. Technol. A 2 (1984) 1079. 93 P.A. Thiry, M. Liehr, J.J. Pireaux and R. Caudano, Phys. Rev. B 29 (1984) 4824. 94 L.H. Dubois, Surf. Sci. 119 (1982) 399. 95 T. Conrad, J. Ghijsen, J.M. Vohs, P.A. Thiry, R. Caudano and R.L. Johnson, Surf. Sci. 265 (1992) 31. 96 A.P. Baddorf and J.F. Wendelken, Surf. Sci. 256 (1991) 264. 97 V.E. Henrich, G. Dresselhaus and H.J. Zeiger, Phys. Rev. B 17 (1978) 4908. 98 K.C. Mishra, K.H. Johnson and P.C. Schmidt, J. Phys. Chem. Solids 54 (1993) 237. 99 R. Courths, B. Cord and H. Saalfeld, Solid State Commun. 70 (1989) 1047. 100 T. Hikita, T. Hanada, M. Kudo and M. Kawai, Surf. Sci. 287/288 (1993) 377. 101 A.D. Baden, P.A. Cox, R.G. Egdell, A.F. Orchard and R.J.D. Willmer, J. Phys. C: Solid State Phys. 14 (1981) L1081. 102 J.E.T. Andersen and P.J. Moller, Thin Solid Films 186 (1990) 137. 103 D.M. Hill, H.M. Meyer III and J.H. Weaver, J. Appl. Phys. 65 (1989) 4943. 104 J.E.T. Andersen and P.J. Moller, Appl. Phys. Lett. 56 (1990) 1847. 105 R.C. McCune and P. Wynblatt, J. Am. Ceram. Soc. 66 (1983) 111. 106 C. Webb, S.-L. Weng, J.N. Eckstein, N. Missert, K. Char, D.G. Schlom, E.S. Hellman, M.R. Beasley, A. Kapitulnik and J.S. Harris, Jr., Appl. Phys. Lett. 51 (1987) 1191. 107 J. HSlzl and F.K. Schulte, Solid Surface Physics, Springer, Berlin 1979, p. 87. 108 Y.W. Chung and W.B. Weissbard, Phys. Rev. B 20 (1979) 3456. 109 R. Dorn, H. Liith and M. Biichel, Phys. Rev. B 16 (1977) 4675. 110 W. GSpel, J. Pollmann, I. Ivanov and B. Reihl, Phys. Rev. B 26 (1982) 3144. 111 P.J. Moller and J.-W. He, Surf. Sci. 162 (1985) 209. 112 W. GSpel and U. Lampe, Phys. Rev. B 22 (1980) 6447. 113 G. Heiland and H. Liith, in: The Chemical Physics of Solid Surfaces and Heterogeneous Catalysis, vol. 3, eds. D.A. King and D.P. Woodruff, Elsevier, Amsterdam, 1984. 114 K. Jacobi, G. Zwicker and A. Gutmann, Surf. Sci. 141 (1984) 109. 115 P.J. Moller, S.A. Komolov and E.F. Lazneva, Surf. Sci. in press. 116 S.V. Didziulis, K.D. Butcher, S.L. Cohen and E.I. Solomon, J. Am. Chem. Soc. 111 (1989) 7110. 117 A.R. Raju and C.N.R. Rao, Sensors and Actuators B 3 (1991) 305. 118 C.T. Campbell, K.A. Daube and J.M. White, Surf. Sci. 182 (1987) 458. 119 K.H. Ernst, A. Ludviksson, R. Zhang, J. Yoshihara and C.T. Campbell, Phys. Rev. B 47 (1993) 13782. 120 P.J. Moller, S.A. Komolov and E.F. Lazneva, to be published. 121 S.A. Komolov and V.N. Strocov, Phys. Stat. Sol. B 167 (1991) 605. 122 I. Sch~ifer, M. Schliiter and M. Skibowski, Phys. Rev. B 35 (1987) 7663. 123 S. Bloom and I. Ortenburger, Phys. Stat. Sol. B 58 (1973) 561. 124 A. Ludviksson, K.H. Ernst, R. Zhang and C.T. Campbell, J. Catal. 141 (1993) 380.

526 125 Q. Ge and P.J. Moller, to be published. 126 P.J. Moller, S.A. Komolov and E.F. Lazneva, Surf. Sci. Lett. 290 (1993) L677. 127 P.J. Moller and J. Nerlov, Surf. Sci. in press. 128 C. Argile and G.E. Rhead, Surf. Sci. Rep. 10 (1989) 277. 129 G.K. Wertheim, Phys. Rev. B 36 (1987) 9559. 130 S. Evans, D.E. Parry and J.M. Thomas, Farad. Disc. Chem. Soc. 60 (1975) 102. 131 D. Schmeisser and K. Jacobi, Surf. Sci. 88 (1979) 138. 132 Q. Ge and P.J. Moller, Thin Solid Films, in press. 133 H.M. Mannsevit, I. Golecki, L.A. Moudy, J.J. Yang and J.E. Mee, J. Electrochem. Soc. 130 (1983) 1752. 134 V.S. Stubican, R.C. Hink and S.P. Ray, J. Am. Ceram. Soc. 61 (1978) 17. 135 H.D. WiemhSfer, S. Harke and U. Vohrer, Solid State Ionics 40/41 (1990) 433. 136 Z.M. Bluvshtein, G.P. Nizhnikova and U.V. Farberovich, Sov. Phys. Solid State 32 (1990) 548. 137 M. Morinaga, H. Adachi and M. Tsukada, J. Phys. Chem. Solids 44 (1983) 301. 138 J.S Solomon and J.T. Grant, J. Vac. Sci. Technol. A 3 (1985) 373. 139 K.-O. Axellsson, K.-E. Keck and B. Kasemo, Surf. Sci. 164 (1985) 109. 140 C. Palacio, J.M. Sanz and J.M. Martinez-Duart, Surf. Sci. 191 (1987) 385. 141 G.R. Corallo, D.A. Asbury, R.E. Gilbert and G.B. Hoflund, Phys. Rev. B 35 (1987) 9451. 149 G. Marletta and S. Pignataro, Chem. Phys. Lett. 124 (1986) 414.

Science of Ceramic Interfaces II J. Nowotny (Editor) 9 1994 Elsevier Science B.V. All rights reserved.

527

The surface chemistry of tin(IV) oxide: defects, doping and conductivity. R G. Egdell Inorganic Chemistry Laboratory, South Parks Road, Oxford, OXl 3QR, UK.

Abstract. A review is given of the application of contemporary UHV surface science techniques to the study of the surface chemistry of tin(IV) oxide, SnO2. Topics of particular interest include the nature of electronic states associated with oxygen vacancies at SnO2 surfaces; changes in electronic structure that may be brought about by doping (especially with Sb); and the relationship between bulk doping levels and surface dopant concentration. Interest focusses ultimately on our recent attempts to produce n-type doping of SnO2 thin films in the near surface region by implantation of ions including Sb, Bi and Nb and on the conductivity of single crystal surfaces.

1. INTRODUCTION.

Tin(IV) oxide adopts the tetragonal rutile structure shown in figure 1. The structure belongs to the space group D144h and there are two formula units per cell. The intrinsic material is a semiconductor with a direct bandgap of about 3.6eV. However the transition from the uppermost occupied valence band into the lowest conduction band is dipole forbidden, so that the gap is referred to as directforbidden [1-3]. Tin(IV) oxide may be doped n-type by substitutional replacement of Sn by a group V element such as Sb or As or by substitutional replacement of O by a halogen such as F or CI. Oxygen deficiency is also presumed to produce n-type doping. The donor states initially lie about 0.15eV below the conduction band minimum but move rapidly toward the conduction band with increasing doping and at carrier concentrations above about 5x1018/cm 3 the material becomes a metallic conductor [4]. Carrier concentrations of the order of 1021/cm 3 may be achieved in thin film [5,6] or ceramic material. At these high doping levels, the conduction electron plasmon energy reaches a value of more than 0.5eV, ensuring high

528

0

-Sn

Q-o 0

o-

b - 4.737 c -

3.185

Figure 1. The tetragonal rutile structure of SnO 2.

reflectivity at lower energy in the infrared. This property is combined with good transparency in the visible region and there are thus many important technological applications of degenerately doped tin(IV) oxide films as electrically conducting, optically transparent heat mirrors [5]. The widespread activity in the study of the bulk properties of SnO2 during the period 1950-1975 is reflected in a series of three seminal review articles published in 1976 [7-9] dealing with preparation and defect structure [7], and electrical [8] and optical [9] characteristics. Practical interest in the surface properties of tin oxide derives from two areas of application. Tin oxide functions as a catalyst for oxidation of many hydrocarbons [10,11], including both alkanes such as methane [12] and alkenes including propene and the butenes [13]. At temperatures around 380~ undoped tin(IV) oxide oxidises propene to acrolein with 20% conversion at 11% selectivity. Addition of Sb20 4 allows 20% conversion at the lower temperature of 340"C, but now with 59% selectivity. The catalytically active phase has been postulated to be Sb-doped SnO2, with a mechanism involving a rate determining step in which lattice oxygen is extracted by propene to yield an allylic intermediate [11,14]. Surface spectroscopy obviously has considerable potential in helping to elucidate details of these processes. The second area of application involves use of tin oxide as a sensor [15] for detection of reducing gases such as methane, carbon monoxide [15-17] and even arsine [18]. Adsorption of oxygen on the surface of doped tin oxide leads to trapping of charge in ionic species such as 02", O22 or O 2-. Reaction between the reducing gas and the surface adsorbed oxygen releases carriers into the bulk and there is an associated increase in conductivity which forms the basis for the sensor devices. Optimising sensor performance essentially involves maximising conversion of target reducing species, but without concern now for selectivity in the

529

products. Contemporary UHV surface science techniques have played a major role in developing our general understanding of surface chemistry and surface processes. The present paper aims to review the application of these techniques to the study of tin oxide surface chemistry. Table 1 summarises the techniques that have been applied to tin(IV) oxide surfaces, along with a brief indication of the potential information content of the experiments. Attention is restricted in this paper to welldefined surfaces prepared under carefully controlled UHV conditions. However, this does not mean that we deal only with single crystal surfaces: a major theme of the present paper is a comparison between single crystal, ceramic and thin film surfaces. It is beyond our scope to review comprehensively the applied catalysis and sensor literature but, where possible, the relationship between fundamental surface science investigations and technological problems is highlighted.

2. THE STRUCTURE OF TIN(IV) OXIDE SURFACES. 2.1 Single crystal surfaces. Low energy electron diffraction is the most widely used technique for the study of the structure of surfaces. Analysis of the spot positions in LEED allows identification of the dimensions and symmetry of the surface unit cell, but full information about the positions of the scattering centres within the unit cell demands a full dynamical analysis of the intensity of the LEED beams as a function of incident energy. The work on SnO 2 has not yet extended to detailed experiments of this sort. The growth of tin(IV) oxide crystals required for LEED studies is a non-trivial task and indeed it is remarkable that most surface science investigations have used crystals from the same single source, namely the research group of Professor R. Helbig. These crystals are grown by a technique in which SnO 2 is transported as SnO in a reducing stream of hydrogen and then reoxidised to SnO2 by a permeation of oxygen into the growth zone [19]. Crystals grown in this way have large natural (110) growth faces and for this reason structural experimental work has thus far been largely restricted to (110) surfaces. It was first pointed out by P.A. Cox et al [20] that viewed perpendicular to the [110] direction the structure of SnO 2 is built up from three alternating layers containing (lxO) ion/cell, (2xSn+2xO)ions/cell and (lxO)ion/cell. The nominal ionic charges are 2-, 4+ and 2- respectively. It was earlier noted by Tasker that when an ionic crystal has charged planes, the crystal must not terminate to leave a repeating unit with a net dipole moment perpendicular to the surface because this leads to a divergent surface energy [21]. In the case of SNO2(110) this implies that there must be a termination with a sequence of ionic planes [(lxO) (2xO+2xSn) (lxO)] which has an electric quadrupole [(2-)(4+)(2-)] but no dipole. This termination is shown

530

Table 1 UHV techniques applied to tin(IV) oxide surfaces Acronym

Technique

Comments

LEED

Low energy electron diffraction. [22-33,44]

Probes structure and unit cell dimensions of single crystal surfaces.

AES

Auger electron spectroscopy. [29,31,32,44,46,127]

Elemental composition based on analysis of Auger electrons produced by decay of a core hole. Initial core hole produced by electron bombardment. Surface sensitivity determined by pathlength of incident and emitted electrons

XPS

X-ray photoelectron spectroscopy. [20,22,29,37-39,50-52, 54,55,122,126-129]

Surface elemental composition based on analysis of core photoelectrons. Surface sensitivity determined by photoelectron pathlength.

ISS

Ion scattering spectroscopy. [25,28-30,121]

Surface elemental analysis based on elastic binary collisions of incident ions with atoms in top atomic layer of solid.

UPS

Ultraviolet photoemission spectroscopy. [20,24-29, 36-39,44,48,50 121,122]

Occupied bandstructure and density of states, including surface states.

IPES

Inverse photoemission spectroscopy. [46]

Empty bandstructure and density of states, including surface states.

EELS

Electron energy loss spectroscopy. [54,55,126]

Excitation of surface electronic transitions by thermal electron beam. Resolution usually worse than 400meV.

HREELS

High resolution electron energy loss spectroscopy [20,38,39,111]

Excitation of surface phonons and low energy electronic transitions (including conduction electron plasmons in doped material) by monochromatic electron beam. Resolution may be better than 10meV.

531

schematically in figure 2(a) [22]: note the rows of bridging oxygen ions which complete the 6-fold coordination of the tin ions in the rows beneath them. The second set of surface tin ions are 4-coordinate in the surface plane with a 5th oxygen ion in the next plane down. Removal of the bridging oxygen ions reduces the coordination number of the initially 6-fold coordinate tin ions to 4 and leaves a dipolar sequence of charged planes [(4+)(2-)(2-)][(4+)(2-)(2-)] etc. However, owing to the variable valency of tin it is possible to envisage reduction of half of the tin(IV) cations to tin(ll) so that the sequence of charges becomes [(2+)(2-)][(2-)(4+)(2-)] etc. This second posible termination is shown in figure 2(b). Note that in both structure 2(a) and 2(b) the periodicity of the bulk is retained and we expect to observe (lxl) LEED patterns.

Figure 2. Structural models for SNO2(110) proposed by Shen et al. [22]. (a) ( l x l ) oxygen bridge terminated surface (b) ( l x l ) vacuum annealed surface (c) (lx2) structure formed by annealing at 1100K. Large circles are oxygen ions, small circles are tin ions. Reproduced with permission from reference [22].

532 The most extensive structural work on SNO2(110) surfaces is due to D.F. Cox and coworkers [23-30]. This group established that annealing SNO2(110) in about 1 torr or higher pressure of 02 for several minutes at 700K leads to a surface displaying a sharp ( l x l ) LEED pattern, with an intensity ratio between O and Sn peaks in ion scattering spectroscopy characteristic of a surface terminated by bridging oxygen atoms [28-30]. Heating to progressively higher temperatures up to 1000K in UHV leads to a monotonic decrease in the intensity ratio between the O and Sn peaks, although the LEED structure remained ( l x l ) (figure 3). This is attributed to loss of bridging oxygen ions to leave a high temperature ( l x l ) surface terminated by the (2xSn+2xO) plane. The ability to selectively label the bridge site oxygen positions with 180 was demonstrated in an elegant experiment in which a SnO 2 surface was first annealed in 160 2, then annealed in vacuum at 700K. The bare (2xSn+2xO) surface thus produced was then annealed in 180 2 and a new peak appeared in ion scattering spectra due to occupation of the bridging sites by 180. This isotope is distinguishable from 1 6 0 in ISS by virtue of the different isotopic mass. These 180 species could be largely removed in turn by further annealing in vacuo (figure 4). An alternative method of surface preparation involves argon ion bombardment, followed by annealing in UHV. The structure of these surfaces has proved to be L

Oxidized

Sn02 (I 10)

I.O c 0.8

U3

f

"

o 0.6 -

x, ,

"o

.N 0.4 o

E 0.2

/

I1,,..

o Z

,Sputtered surface

o~~'"-~-.~O-__o_

0 "

_ ~,~

.

L(4Xl) _L jC]xl) I

300

I

I

L (1xl) rI ,

I-

.Lc(2X2)l "T" -I , i

I

-I

,

600 900 Annealing Temperature (K)

,

Figure 3. Normalised intensity ratio between O and Sn peaks in ISS of SNO2(110) as a function of temperature of annealing initially oxidised and sputtered surfaces. The range of stability of different LEED reconstructions is indicated. Adapted from reference [29].

533

more controversial. In their pioneering study of SNO2(110) Darville and coworkers. [31,32] observed an evolution of structures with increasing annealing temperatures through a regime where (4xl) and c(lxl) patterns coexisted to a (4xl) pattern at temperatures above 723K and finally to a (lxl) pattern at 773K. (table 2). The O/Sn Auger intensity ratio showed somewhat erratic behaviour with increasing annealing temperature, but in general increased with increasing annealing. By contrast Erickson and Semancik [33] found only diffuse LEED structure for annealing temperatures below 600K with progressive development of (4xl), ( l x l ) and (2xl) phases as the temperature was raised, the last structure only developing at temperatures above 980K. In later publications, D.F. Cox, Fryberger and Semancik found a somewhat different pattern of behaviour with coexistence of c(2x2) and (4xl) structures at temperatures around 800-900K and no indication of the high temperature (lx2) structure [29]. In ion scattering spectroscopy the O/Sn intensity ratio remained roughly constant for annealing temperatures up to 1000K and then increased sharply to the value of the (2xSn+2xO) (lxl) phase (figure 3). However in differential mode Auger spectroscopy [27, 29], the O/Sn intensity passed through a

16

(c)

/~0

N(E)

0 35

~60

(d)

(b)

1

I

0-43

i

I

0.51

E/Eo

0.35

'

' 0-43

'

' 0.51

'

Figure 4. Oxygen signal in ion scattering spectra from SNO2(110). (a) Bridging oxygen terminated surface after annealing in 1 Torr of 160 2 at 700K. (b) After annealing to 700K in vacuum to remove bridging oxygen ions. (c) Following reoxidation at 700K in 180 2. A peak due to bridging 180 now appears. (d) Following a further anneal in vacuum at 700K to remove the bridging 180. Adapted from reference [30].

534

distinct minimum at 600K. The most recent experiments by Thornton and coworkers [22] have found an evolution of structures (4xl) -> (lxl) -> (lx2) very similar to that of Erickson and Semancik, but within somewhat different temperature regimes. They also found a straightforward increase in the intensity ratio between O ls and Sn3d peaks in XPS (figure 5). In summary then the centred structure c(2x2) [29] and c ( l x l ) [31,32] are controversial and the evolution of Auger intensity ratios is apparently at variance with XPS and ISS results. These ambiguities may simply reflect subtle differences in the conditions used for the ion bombardment and uncertainties in the true surface temperature during annealing, which is always difficult to measure in experiments on oxide crystals. There is however a reasonable consensus that post-annealing bombarded crystals can lead to a (lxl) surface with composition similar to that Sn02(llO) 4 x ~ x t~3 C U3 i/1

o

0.40

0.30 0.20

>,

10 ~.~

I

m~ 0.1 c 0.01

8

-0.0

>-0.2 Ca

v

,~ - 0.4 -

0.6

400 800 i2( )0 Anneal temperature(K)

Figure 5. O ls/Sn3d XPS intensity ratios, sheet conductivity and work function measured at room temperature after sequential 10 minute anneals of ion bombarded SNO2(110) to the temperature indicated. Adapted from reference [22].

535

produced by annealing an oxidised surface: this structure corresponds to figure 2(b). The lower temperature (4xl) surface is clearly more strongly oxygen deficient than the ( l x l ) surface and could therefore involve simple removal of every fourth row of in-plane oxygen ions [21,29,31,32]. However it is difficult on this basis to understand the stability of the (4xl) periodicity. D.F. Cox et aL [29] therefore suggest a structure in which an overlayer of SnO(101) sits on top of the SNO2(110). It requires only small expansions in the SnO lattice parameters for the two lattices to coincide exactly in the [110] direction of the substrate and for the lattice parameter of the overlayer to be 4/3 that of the substrate in the [001] direction (figure 6). This coincidence structure will give the observed (4xl) LEED pattern. Finally the high temperature (lx2) pattern has been observed by two groups [21,33] and has a higher oxygen content than the bare (lxl) structure. As suggested by Thornton and coworkers [22], a plausible structure therefore involves incorporation of bridging oxygen atoms at half the available positions (figure 2c). This structure is similar to that suggested for TiO2(110)(lx2) by Meller and Wu [34].

6-699 a)

--~

b)

i ,0223A .......

I go2

o

Figure 6. Proposed coincidence structure for SNO2(110) 4xl. (a) Bare SNO2(1 10) surface with no terminating layer of bridging oxygen. Large circles are oxygen ions, small circles are tin cations. (b) Structure of tin containing plane associated with SnO(101) surface, assuming unrelaxed bulk lattice parameters. The dashed lines demonstrate the extent of relaxation required to produce a coincident lattice giving the observed (4xl) LEED pattern. Adapted from reference [29].

536

Table 2 Reconstructions observed on ion bombarded and annealed SNO2(110) surfaces. Reference

22 400K 840K 900K 1050K

1 lOOK

(lxl)d

(4xi) (lxl)

(Ix2)

29

31,32

600K

523K

(lxl)d 850K c(2x2) 1000K 1050

(4xi)

(lxl)d

33

(4xl)+ 623K c(lxl)

650K

723K

850K

773

(4xi) (lxl)

IO00K

(4xl) (lxl)

(ix2)

d= diffuse pattern only.

2.2 Polycrystalline surfaces. The surfaces of polycrystalline ceramic or thin film samples are often regarded as somewhat intractable by UHV surface scientists [35]. We believe it is important to distinguish between surface studies on high area, poorly sintered powder material and on dense well-sintered ceramic or thin film material. In the former we may reasonably envisage a preponderance of high energy step and edge sites at the surface, as well as hydroxylation at surface sites. In the latter the high energy sites should be largely eliminated and surface will be dehydroxylated. It remains true that there will in general be a range of different atomic planes exposed at the surface and it is impossible to obtain detailed structural information about each of these in the study of e polycrystalline sample. In some cases the samples may nonetheless show pronounced texture. For instance, in our own work on thin film SnO 2 deposited on quartz substrates by radiofrequency sputtering, the polycrystalline samples shows very strong (110) preferred orientation [36,37] and one may then expect close similarities in surface properties to single crystal samples. In addition it will emerge in the next section that the pattern of electronic levels at SnO 2 surfaces is rather similar for single-crystal, ceramic and thin film material and it is therefore appropriate to comment briefly on techniques for producing atomically clean polycrystalline surfaces. P.A. Cox et al [20,38] showed that annealing ceramic SnO 2 in high oxygen partial pressures (150 torr 02) gives surfaces free of carbon or other contamination and it must be presumed that these are analagous to the oxygen bridge terminated surfaces in having all surface tin cations in the maximal

537

oxidation state. Clean surfaces of ceramic [39] and thin film samples [36,37] are also produced by annealing in UHV, but here bridging oxygen atoms will be removed from (110) surfaces and from related sites on other surfaces, so that reduction of surface plane cations from Sn(IV) to Sn(ll) must be anticipated. Finally ion bombardment of polycrystalline surfaces will also lead to removal of in-plane oxygen ions.

3. ELECTRONIC

STRUCTURE AND DEFECT STATES.

3.1. Photoemission and bulk bandstructure. A schematic system of ionic energy levels for SnO 2 is shown in figure 7. Empty 5s and 5p states associated with Sn4+ lie above occupied 2p states of O2-, and the

Sn 5p

Sn

5s

0 2p ( ~ ) 0

2p(6)

0

2s

Sn Z.d

Figure 7. Schematic of ionic energy levels for Sn02.

538

covalent interaction between these two sets of levels is crucial in determining the bandstructure of SnO2. Four bandstructure calculations for SnO2 have been published [40-43], but we confine our discussion to the tight binding calculation of Robertson which gives the clearest chemical insight into the bonding [42]. Within the SnO 2 structure, the Sn 4+ ions are surrounded by 6 oxygen ions with site symmetry D2h. Nonetheless the space group for D4h 14 includes C4 rotational operations coupled to a non-primitive translation and wavefunctions at the F point may be classified according to the irreducible representations of the point group D4h: the relationship between the [` notation usually used in bandstructure calculations and the more chemical SchSnflies notation is given in table 3. The oxygen ions are surrounded by 3 tin ions in a trigonal planar arrangement. If no O-O interactions are included, the upper valence band consists of four oxygen lone pair bands with symmetries F2+, ]-'3+ and [`5 ( the last being a doubly degenerate) that are degenerate and flat. These states are derived from p orbitals on O orthogonal to

Table 3. Relationship between alternative representations of the point group D4h. ['1 + ['2 + ['3 +

Alg A2g Big

]-'1" ]-'2" ['3"

['5 +

Eg

['5"

['4 +

B2g

['4"

notations for the

irreducible

A2u Alu B2u Blu Eu

the plane of the three surrounding tin ions. Inclusion of oxygen-oxygen interactions lifts the degeneracies in the upper valence levels and the bands now disperse over an energy range of about 2eV (figure 8). Nonethless, the upper valence bands retain significant oxygen lone pair character. The uppermost valence band at the r point is F3+. The lowest conduction band state is a combination of Sn5s states that are in-phase within the unit cell with r 1+ symmetry. The transition r 3 + - > ]-'1+ is dipole forbidden but direct, as found experimentally [1-3]. At the r point the lowest conduction band state has 90% Sn5s atomic character, but disperses strongly upward in energy away from r due to very strong mixing with the O2p states. The upper ['4 + Sn5s derived band involves orbitals out-of-phase within the unit cell and shows less strong dispersion because-mixing with O2p orbitals is allowed even at ['. The 6xSn5p derived states have symmetries 2x['5+ (each of these is twofold

539

12 5+ 1.,-

I- 5§ 4,

0 L,

3" 55-1I+ 2+ 4- 5-

c l.iJ

-8 4,

q

-18 F

A

Z S A

V M ~ FAX

W RUZ

Figure 8. The bandstructure of SnO2 derived from the tight binding calculation of Robertson [42]. Within the energy range depicted are the 12xO2p derived bands, 2xSn5s bands and 6xSn5p bands. Adapted from reference [42].

degenerate), r 4- and FI-. These too all show relatively weak dispersion. In terms of the density of states (figure 9), the valence band has an overall width of about 10eV and shows a three-peak structure. The strongest peak is at the top of the valence band and corresponds to the O2p lone pair states alluded to above. The second peak corresponds to states that have acquired significant Sn5p atomic character due to covalent mixing and the states at the bottom of the valence band to states of similarly mixed Sn5s-O2p character. The density of states at the bottom of the conduction band is very low due to the strong dispersion of the FI+ band. There is a strong peak in the empty density of states extending between about 4eV and 8eV above the bottom of the conduction band associated with the Sn5p levels. The energy distribution curve found in photoemission from oxygen annealed polycrystalline SnO 2 excited at 40.8eV photon energy [20] shows excellent agreement with the calculated density of states due to Robertson (figure 9). Energy distribution curves from single crystal SNO2(110) excited with synchrotron radiation at hv=40eV [44] and from SNO2(001 ) [45] are essentially identical to those from polycrystalline material. With synchrotron radiation it was also possible to reconstruct spectra showing the effect of sweeping the kinetic energy of the

540

128

~

f O

IPES

~

_/.,- e

,~,

'

-8-

,Sn {..-~-..s

p

5"

0o ,ng

.

--- DOS I..... I o n e p a i r ~

Figure 9. Density of states (DOS) for SnO2 derived from the bandstructure calculation of Robertson [42]. Both the total DOS and its decomposition into contributions from Sn and O atomic orbitals is shown. For comparison He(ll) PES data from reference [20] (hv=40.8eV) and IPES data from reference [46] (hv=9.8eV), both measured on polycrystalline SnO2 are given. photoelectrons and the photon energy together so as to keep the initial state constant. On passing through the Sn 5d->Sn5p photoabsorption threshold at about hv=35eV distinct resonances were observed for photoionisation of the valence band states with Sn5s and Sn5p admixture (figure 10). This is attributed to decay of the excited state 4dgvn5p 1 (here vn denotes the fully occupied valence band configuration) to a final state 4d1~176 + e(Ek), where the energy Ek of the emitted photoelectron is the same as would be obtained by direct photoemission. The empty electronic states in polycrystalline SnO2 have been investigated [46] using inverse photoemisson spectroscopy at hv=9.8eV. The intensity just above the Fermi energy was very low, in agreement with the bandstructure calculation. Moreover, a broad peak was found in the IPE spectra peaking about 4eV above the Fermi energy (figure 9). This presumably corresponds to the Sn5p states, although the experimental peak is somewhat lower in energy than expected from the bandstructure calculation. The workfunction ~ of polycrystalline vacuum annealed ceramic SnO2 is about 4.5eV [46] and from work on single crystal surfaces it is clear that in situ surface treatments produce variations in ~ that are generally less than 0.5eV

541

,

1.

r"O

I1-1- I- I.

'.

ffl c-

SnO2('110) CIS

~ Gap States II.

' 02p

I.

11-1.

"l

"ll . . . .

I--I

"ll-- 1 - - 1 "-I ..

,1--1 mm-""l. .... a-..... 9

", Sn51~.'

OQ. ffl

",

i-.'

E

,,'

"i- -1

"-,

r n

1

22

30 38 46 Photon Energy (eV)

Figure 10. Constant initial state spectra in the region of the Sn4d->Sn5p core threshold showing variation in intensity of features in the photoemission spectra of SnO2 as a function of photon energy. The O2p data corresponds to the O2p lone pair valence band peak, the Sn5p data to the valence band feature feature associated with states with strong Sn5p admixture and the Sn5s data to the valence band peak associated with states with significant Sn5s admixture. Adapted from reference [44].

[22,24,27,28,29]. It is known moreover that the Fermi energy lies close to the bottom of the bulk conduction band. It follows then that the vacuum level cuts the middle of the high density of empty states associated with the Sn5p levels. Primary electrons produced in the photoemission process may scatter into these empty states by various inelastic processes and have a significant probability of remaining in them before finally being emitted into the vacuum with a low kinetic energy. For SnO 2 these secondary electrons give rise to strong and highly reproducible structure in photoemission spectra of both single crystal [28] and polycrystalline samples [20,37]. The most commonly used laboratory source for excitation of photoemission spectra is the He(I) source with hv =21.2eV. At this energy the onset of the low kinetic energy secondary electrons begins before the intensity of primary emission from the O2p valence bands has reached its minimum and therefore the apparent

542

(a)

(b)

(r

15

10 5 Kinetic energy leV

Figure 11. He(I) photoemission spectra (hv=21.2eV) of (a) single crystal SNO2(110) and (b) polycrystalline SnO 2. The structure marked S is due to secondary electrons. (c) Shows the secondary electron structure resulting from scattering a 25eV electron beam from the sample surface transferred onto the same kinetic energy scale as the photoemission spectra. Adapted from references [20] and [28].

primary valence band width in He(I) spectra is less than at higher photon energies (figure 11). It is again striking that spectra from single crystals and polycrystalline material are essentially identical. Confirmation that the low kinetic energy structure comes from secondary electrons is provided by the observation of structure at the same kinetic energy in the energy distribution curve resulting from scattering of low energy electrons from the tin oxide surface [48].

543

3.2. Oxygen deficient surfaces: states in the bulk bandgap. Given that the bulk bandgap of SnO2 is 3.6eV and that the bulk Fermi energy lies not more than 0.15eV below the conduction band minimum, the onset of photoemission from bulk oxygen valence band states is expected about 3.5eV below the Fermi energy if flat band behaviour pertains at the surface. Note however that the photoemission onset will be broadened by coupling of phonons to the photohole and a Gaussian bandedge may be expected. P.A. Cox et al [20] first noted that in polycrystalline vacuum annealed SnO2 significant photoemission intensity extends as a shoulder from the valence band edge into the lower part of the conduction band. This observation was confirmed by experiments on SNO2(110) carried out by D.F. Cox and coworkers [27-29] who showed that II Sn02(llO) (o) ion bombarded

QLL

I x I0 16-0

8.0 0" 5-0 Binding energy/eV

0

Figure 12. Photoemission spectra of SNO2(110) following three different surface treatments. (a) 2keV ion bombarded surface (b) As (a) followed by annealing in vacuum at 1000K (c) After annealing at 700K in 1Torr of oxygen. The right hand panels show enlarged views of the bandgap region. Adapted from reference [28].

544

VBM

Annealing Temperatu re 400K

~~.600K

CBM EF I

EF

U.I

I

Z <1

F

3-6 VBM

2.0 Binding Energy (eV)

0 CBM

Figure 13. He(I) photoemission difference curves showing the increase in bandgap defect states produced by heating an initially oxidised surface in vacuum to the temperatures indicated. Adapted from reference [28]. appearance of structure in the lower part of the bandgap (figure 12) could be correlated with removal of bridging oxygen ions to produce the bare (2xSn+2xO) (lxl) surface. Annealing at temperatures above 1000K [28] led to the appearance of structure even higher in the gap, extending toward the Fermi energy (figure 13). It was postulated that this was associated with in-plane oxygen vacancies. All the gap states may be largely quenched by exposure to oxygen [28]. It has also been observed that a strong shoulder on the valence band edge extending into the lower part of the bulk bandgap may also be produced by argon ion bombardment of both single crystal [28,29,44,49] and polycrystalline [50] surfaces of SnO2. It is generally believed that the states in the lower part of the bandgap are Sn 2+like, although this is difficult to establish unambiguously. In XPS studies of bulk SnO 2 and SnO, it was shown many years ago that there is essentially no chemical shift in core XPS between these two oxides [51,52]. Detailed calculations for the related pair of compounds PbO and PbO 2 show that this may arise due to

545

cancellation between the contribution to the chemical shift arising from the differing ionic charges by the difference in site potential between the 6-coordinate environment of the M(IV) cation and the 4-coordinate environment of the M(II) cation [53]. Thus even though SnO2 is reduced to a stoichiometry close to that of SnO under prolonged argon ion bombardment, no chemically shifted components appear in Sn3d XPS [29,54]. Likewise in core XPS the removal of bridging oxygen ions does not alter Sn3d core XPS profiles [29]. However there are changes in energy loss spectra that appear diagnostic of the emergence of SnO [55]. Compelling evidence that the gap states in bombarded SNO2(110) are cationic in nature comes from the observation (figure 10) of distinct resonances in the intensity of the gap emission at photon energies corresponding to the Sn4d core excitation threshold [44]. The reason why the Sn 2+ defect levels lie close to the valence band edge is controversial. Simple charge balance considerations dictate that oxygen vacancies leave behind associated electrons but in a rigid band model these should occupy conduction band states. Polaronic self-trapping may stabilise these electron states. But even where this effect is important, for example in the in the closely related oxide TiO 2, oxygen vacancies produce states in the upper part of the bulk bandgap [35,49], and this behaviour is characteristic of most transition metal oxides. Based on the parameters from the tight binding calculation of Robertson [42], Munnix and Schmeits used the scattering theoretic method to calculate the surface band structure for SNO2(110), assuming the surface termination with no bridging oxygen ions [56-58]. Projecting the surface bands onto the bulk bandstructure, two surface states were found in a so-called "stomach" gap toward the bottom of the main O2p valence band. In addition a number of weak surface resonances were found in the upper part of the valence band. In the conduction band region, three surface states are found, the lowest of which is only 0.1eV below the conduction band minimum at the ]" point. Thus, in the calculations the gap is essentially free of surface states, in clear disagreement with experiment. States in the middle of the gap were also calculated to be absent for SNO2(100 ) and SNO2(001) [53]. A way out of this dilemma is clear from the suggestion of P.A. Cox et al [20] that the surface state for SNO2(110 ) with bridging oxygen ions removed arises because of strong mixing between 5s and 5p atomic orbitals at surface sites where the essentially centrosymmetric coordination of bulk sites is lost. In general, mixing between 5s and 5p levels under the influence of a potential V is described by a matrix element <5slVl5p> and will be non zero if V has a component belonging to the same irreducible representation as one of the 5p orbitals in the point group appropriate to the local site symmetry. This is never possible at a site with a centre of symmetry because the 5s and 5p orbitals are of opposite parity and cannot be mixed by centrosymmetric crystal fields. Thus 5s-5p mixing is inhibited at the F point. The

546

(a)

(b)

/

I

/ / /

Sn 5p

,....- s

f

s

Sn 5s %

%

% %

v

INCREASING

<5slY I 5p>

Figure 14. Schematic ionic energy level diagram for Sn (a) in a centrosymmetric environment (b) in a non-centrosymmmetric environment where <5slVl5p> is non zero and there is strong 5s-5p mixing.

condition for mixing is however clearly satisfied by tin ions in the 4- and 5coordinate surface sites of figure 2 and will generally pertain at any surface cation site where the coordination is reduced below its bulk value of 6. The 5s-5p mixing will stabilise the in-phase 5s-5p combination, as shown schematically in figure 14, pushing the resulting state toward the valence band edge. By contrast the out-ofphase combination will be pushed up above the main Sn5p bands. The notion of 5s-5p mixing has been called into question by Themlin et aL [44], but the recent observation of empty defect-related states above the main Sn5p band in inverse photoemission of ion bombarded SnO 2 (figure 15)appears to provide strong evidence in favour of this picture. Recent cluster calculations simulating the density of states in Sn(ll) compounds reinforce this scenario and show that the Sn 2+ lone pair states in SnO have very substantial Sn5p atomic character [59].

547

IPES hv - 9.8eV Sn

c"

u3 O

.:,.

E

L,

O Z

-

1

I

I

I

4 8 12 Energy reletive to Ef/eV 0

Figure 15. Inverse photoemission spectra of SnO 2 as a function of exposure to 500eV argon ion beam at 5mNcm 2 (a) unexposed (b) 1 hour exposure (c) 2 hour exposure (d) 3 hour exposure. In (d), the original spectrum (a) is shown with dots, emphasising the creation of new states about 6eV-8eV above the Fermi energy due to ion bombardment. These are attributed to empty Sn5s-5p hybrid states. Adapted from reference [46].

4. n-TYPE DOPING OF TIN(IV) OXIDE 4.1 Doping in polycrystalline and thin film material. Previous work on doping in polycrystalline and thin film SnO2 divides into two largely unconnected areas. The chemistry literature is concerned mainly with Sbdoped SnO 2 and the relationship between a doped phase Snl.xSbxO 2 and practical oxidation catalysts mentioned in section 1. By contrast the physics literature focusses on measurement of conductivity, Hall coefficients and optical reflectivity and transmittance in thin film material. Antimony-tin oxide catalysts are now usually prepared by coprecipitation of hydrated oxide species from acidic solution of Sn(IV) and Sb(V) by addition of ammonia, followed by calcination at temperatures ranging from 400~ to 1000"C [10]. After firing at temperatures up to 600~ the material is poorly crystalline and high resolution electron microscopy reveals the presence of significant

548

concentrations of amorphous phases [60]. For this reason the observation of only a rutile phase in X-ray diffraction at antimony concentrations up to 70% cannot be taken as evidence of a 70% substitutional limit for Sb in SnO 2. [10, 14]. In fact electrical measurements show that maximum conductivity is achieved at a much lower doping level of 5% [61] or 3% [62], and these lower levels are taken to represent the true solubility limit. This idea is supported by the observation that when samples containing more than 4% Sb are calcined for extended periods at 1000"C, there is progressive loss of Sb which volatilises as Sb40 6 to give a 4% limiting concentration. This doping level corresponds to a carrier concentration of 1. lx1021/cm 3 if each Sb ion acts as a donor centre. The mode of Sb incorporation into maximally doped SnO 2 is of particular interest. Neutron diffraction studies have been carried out on 10% Sb-doped SnO 2 (which represents an Sb concentration beyond the generally recognised equilibrium limit) and suggest an unusual structure with Sb(lll) in interstitial sites within the rutile lattice, with compensating Sn(IV) vacancies [63]. However, a lone pair ion Sb(lll) will give rise to an electronic state deep below the Fermi level of SnO2 and it is difficult within the framework of this model to understand the essentially metallic transport properties of Sb-doped SnO 2. Unfortunately EXAFS, which could in principle probe the Sb coordination more directly at lower Sb doping levels, give scant information because of the proximity of Sn and Sb core edges. [64]. Atomistic simulation studies suggest that Sb(V) will substitute directly into Sn(IV) sites, with charge compensation by an Sn(IV) vacancy for every four Sb(V) ions [64,65]. Of course this model ignores the expected charge compensation by electrons in the Sn5s conduction band. Perhaps the most informative technique has been 121Sb MSssbauer spectroscopy [66,67]. For 4% Sb-doped SnO 2 calcined at 8000C a dominant Sb(V) peak was found, presumed to be associated with subststitutional Sb in Sn(IV) bulk sites, together with a weak Sb(lll) peak arising from dopant ions at surface or grain boundaries. The Sb(lll) signal was much stronger for material with initially higher doping levels calcined for extended periods at 1000"C [67]. The characteristic pale blue colour of Sb-doped SnO2 has been extensively discussed [10 and references therein] in terms of intervalence charge transfer excitation between Sb(lll) and Sb(V), although as we will see below the colouration is more simply rationalised in terms of the plasma reflectivity characteristics. In summary then it is possible to formulate Sb-doped SnO 2 as Snl_xSbxO 2 but the evidence for this is not unambiguous. The thin film literature is very much more extensive. Already by 1976, Jarzebski and Martin had documented over 30 papers dealing with preparation of SnO2 films. Current techniques for film deposition include the following: spray pyrohydrolysis of aqueous or organic solutions containing Sn(IV), usually SnCI 4, on a heated glass or quartz substrate [6,68-72]; chemical vapour deposition, notably using dibutyl tin diacetate (DBTDA) as the volatile tin compound [73-75]; plasma-assisted chemical

549

vapour deposition using SnCI4+O2 [76]; electron beam evaporation of SnO 2 [77]; r.f. sputtering from SnO 2 targets [36,37]; magnetron sputtering from SnO2 targets [78,79] ; and reactive evaporation of Sn through an oxygen plasma [80]. Doping is easily achieved in techniques involving liquid precursors. Thus in spray pyrolysis, Sb doping can be effected by adding SbCI 3 [71] or SbCI 5 [7] to the spray solution, whilst F doping is possible by adding NH4F [72]. In chemical vapour deposition [75], SbCI 5 or CCI3CF 3 can be added to the DBTDA feedstock to produce Sb or F doping. There is also a possibility that CI in the feedstock may produce adventitious n-type doping, although this question has attracted little attention. In sputtering techniques it is possible to use a mixed oxide or metal alloy target [80]. Even where there is no deliberate doping it is often observed that that there is a high concentration of carriers, there being an inverse correlation between grain size in the deposited film and the doping level [81-83]. Since grain size increases with film thickness, the carrier concentration can be varied simply by varying the deposition time [6]. The nature of the doping in these cases remains obscure, but it is tempting to speculate that surface or near surface oxygen vacancies are important. Doping has two effects on the optical properties of thin films. Firstly, occupation of conduction band states blocks transitions from the F point. The bandgap absorption therefore shifts to higher energy with increasing doping, the so-called Moss-Burstein effect [84,85]. A detailed analysis of the shift reveals that renormalisation of the bandgap to progressively lower values accompanies the occupation of progressively higher conduction band states [6]. As with doped In203 [86] and ZnO [87], the upward dispersion of the conduction band makes a greater contribution to the overall Moss-Burstein shift than downward dispersion of the valence band due to the greater valence band effective mass. Secondly, the conduction electrons are associated with changes in the dielectric properties of the films in the infrared that may be described in terms of the equation:

~(o~) = ~(oo). %2/(o,.2 + i~p)

(1)

where e(co) is the frequency dependent dielectric function, e(oo) is the background dielectric constant, COp is the so-called plasmon frequency and ~,p is the plasmon width. The plasmon frequency is given in terms of the carrier concentration ne, the conduction electron effective mass m e and the permittivity of free space eo COp2 = (nee2/meeo)

(2)

The plasmon frequency is important because below a critical screened plasmon frequency 00p/~/e(oo)the reflectivity is very high and drops down toward zero below

550

1-0 0

0-8

I -

I

E

:~u 0-6 E O'Z,

0"2

-

I

I

I 9

I

9

e 9

,

I

4

,

I

8

,

i

12

ne(1020 cm-3) Figure 16. Compilation of effective masses for SnO2 thin films (closed circles) as a function of doping level adapted from reference [6] compared with effective masses from HREELS data(closed rectangles) of reference [39].

this frequency. The reflectivity edge is typically in the mid infrared for doped SnO2. Across the visible region the reflectivity slowly increases, accounting for the blue colouration of macroscopic samples of doped material. Analysis of the reflectivity edge enables (Op and hence m e to be determined, provided that the carrier concentration is known, for example from Hall effect measurements. A compilation of values from several studies [6] shows a general tendency for me to increase with increasing carrier concentration (figure 16). Similar behaviour is found in conventional III-V semiconductors and is a consequence of energy dispersion with wavevector in the conduction band becoming less pronounced the higher one is in the band [88] 4.2

Surface

studies

of

Sb-doped

SnO2.

Single crystals of Sb-doped SnO 2 can be grown by controlled pyrolysis of SnCI4/SbCI 5 [89], but the crystals obtained in this way have thus far been of insufficient size for electron spectroscopic investigation. Experiments have therefore been restricted to ceramic samples, prepared by coprecipitation and prolonged annealing at 1000~ as described above. P.A. Cox et al. first made comparison between 3% Sb-doped SnO2 and the

551 UNDOP D

1,250j 0.6o/ x 25C ...__-.

I. 5

~/ 25o.U

X 2I ~-0 ~ ~

I~-

- JL

0 5 10 15 Binding EnergyleV

I

I

20

Figure 17. Photoemission spectra excited at hv=21.2eV for vacuum annealed polycrystalline Sb-doped SnO2 at the different doping levels indicated. The weak peak at the Fermi energy (zero binding energy) is due to the conduction electrons introduced by doping. Adapted from reference [39].

undoped compound [20,38]. This work was later extended by Egdell et al [39] to a range of doping levels between 0 and 3%. In He(I) photoemission it was found that doping leads to the appearance of a weak but well defined peak close to the Fermi energy (figure 17): the integrated band area of this peak increases linearly with nominal doping level, as expected from a model where each Sb atom donates one electron to the SnO 2 conduction band. The very low intensity of the conduction band peak is a consequence of both the low concentration of conduction electrons relative to the number of O2p valence electrons and of the very low cross section for ionisation of Sn5s-like states: the ratio between one electron cross sections for atomic Sn5s and O2p states at hv=21.2eV is about 0.012 [90]. In fact the intensity ratio between valence and conduction bands is in tolerable agreement with a

552

..

-.25

~

. ........,_.:...,:~;....:..:~,.~

.25 BINDING

.5 ENERGY

~-:..-;:..;.-.,,. :..: .:..

.75 /

eV

Figure 18. He(I) UPS of 3% Sb-doped SnO2 close to the Fermi energy recorded at an instrumental resolution of 35meV. The solid line is calculated from free electron theory, using an effective mass derived from HREELS data. The cutoff is broadened by the Femi-Dirac distribution function. Instrumental broadening is seen to be negligible. Taken from reference [20].

simple calculation assuming atomic cross sections [20,38]. Moreover the shape of the conduction band in 3% Sb-doped SnO2 conforms very well to a free-electronlike profile (figure 18), in agreement with the essentially parabolic dispersion found in band structure calculations [42]. The electron carriers at the surface of Sb-doped SnO2 can also be probed by HREELS. The loss function P(o)) describing spectra probed by this technique is essentially given by the expression: P(oe) = (e2/llvo)) Im [ (E(~)-1) / (e(o))+1) ]

(3)

This has peaks where: Re[r

= -1

(4)

Assuming a simple dielectric function of the form given in equation (1), the loss peaks are found to occur at a frequency shifted from that of the screened bulk plasmon to a surface plasmon frequency O)sp: O)sp = ~/{nee2/meeo(e(oo)+l)}

(5)

553

UNDOPED x2-5 0.3% xl0 1.5o/.

3-0 % 15 I

I

!

I

0 0.5 ] 1.5 2 Energy Loss/eV

Figure 19. HREEL spectra for vacuum annealed polycrystalline Sb-doped SnO 2 at the different doping levels indicated excited with 25eV electron beam. Energy resolution set at 30meV for these measurement. Note plasmon loss peak that moves to higher energy with increasing doping. Adapted from reference [39].

The loss peaks found in HREELS (figure 19) shift to progressively higher energy with increasing doping, as expected from equation (5). Assuming that the carrier concentration corresponds to the bulk Sb doping levels introduced in the synthesis, the effective masses shown in figure 16 are obtained. These are clearly on the high side of those derived from the thin film studies, although the effect is not dramatic. The data may indicate that in ceramic samples not all of the Sb ions act as donor centres; the high apparent value of m e then results from a true n e lower than assumed. However, the high effective mass in HREELS could be due to band narrowing at the surface of the SnO 2. Turning finally to XPS studies of Sb-doped SnO 2, the very surprising result emerges that the surface Sb concentration inferred from the intensity ratio between

554 Sn: I UNDOPED 5/2 ASn:3d3/2

_L22_

I

I

t

I

IJ

/-,75 500 525 550 575

Binding Energy/eV Figure 20. XPS of Sb-doped SnO 2 ceramics in the Sn3d and Sb3d+Ols region. The O ls peak overlaps the Sb3d5/2 peak, but the ratio between Sb3d3/2 and Sn3d3/2 peaks may be appreciated by visual inspection. In the absence of segregation, the intensity ratio should be roughly equal to the bulk doping level. There is clear evidence for dopant segregation. Adapted from reference [39].

Sb and Sn core peaks is very much greater than that expected from the bulk doping level (figure 20). Sb segregation in mixed Sn-Sb oxides was in fact first observed in 1979, but these studies concentrated on the dynamic effects of Sb evaporation from materials with initially high Sb content [91,92]. The data in figure 20 is from equilibrated material that has been held at 1000"C for extended periods. Egdell and coworkers [20,38,39] have rationalised the high Sb intensity in terms of replacement of Sn by Sb in the topmost ionic layer of the samples with full saturation of surface cation sites at 3% Sb bulk doping. It is clear from photoemission and HREELS that the segregated Sb ions cannot act as donor centre. It was suggested that 5s-5p mixing at the surface Sb sites stabilises Iocalised hybrid states that lie lower than the Sn surface state described in section 3.2. These can trap two electrons and

555

SnO2/ Sb ( 3~

UNDOPED Sn 02

E

Sn 5p Sn 5s

EF

DONOR LEVEL Sn 5s-5p

~ S

CONDUCTION BAND EF

Sn /Sb 5p n/Sb 5s

SURFACE -

J I

OZp

VALENCE BAND g(E:)

CONDUCTION BAND

SURFACE ES

OZp

VALENCE BAND g(E:)

Figure 21. Schematic density of states for vacuum annealed surfaces of undoped and Sb-doped SnO2, showing surface states associated with Sn(ll) and Sb(lll), as well as effects of bulk doping. results in surface Sb(lll) cations. Thus the segregated Sb ions do not introduce electrons into the conduction band. For SNO2(110), charge balance considerations dictate that at the bare non-bridged termination 2xSb(lll) ions should replace Sn(ll) + Sn(IV). This model accounts for the removal of the bandgap intensity of photoemission from vacuum annnealed 3% Sb-doped SnO 2, because it is envisaged that the deeper Sb(lll) states overlap the valence band, as shown schematically in figure 21 [20]. It also accounts for the appearance of an Sb(lll) signal in M0ssbauer spectroscopy. The model of reference 20 has been extended by B~langer et al [75] to deal with trapping of carriers in deep surface states at (211) and (301) surfaces in polycrystalline Sb-doped SnO 2 thin films.

4.3 Doping by ion implantation.

Ion implantation is used in a routine way for introducing dopants into conventional semiconductors such as Si or GaAs. There has been surprisingly little

556 II

O O O 0 O00OO

Co)

+0oo o~

9-

o 9

~ ]

e

""

8

9 (b)

"~ 7"li

rr ~ ~ " 0 a+6 .

,o"

e e

........

t21fl @tI01

(a)

5

i

I

0.002 0.003 i/( Temperature / K)

0001

Figure 22. Loglo(resistance)/ohms versus 1/(temperature/K) plots for: (a) 121Sb implanted thin film. 3x1016ions/cm 2 at 90keV. (b)2O9Bi implanted film. 3x1016ions/cm 2 at 90keV. (c) unimplanted film. Adapted from reference [98].

work on doping oxides, although there are reports of reduction in the sheet resistance of SnO 2 thin films by implantation of P [93], As [93-95] or Sb [96,97]. At low doses, the depth distribution of the implanted ions is characterised by a Gaussian distribution, truncated by the surface. However sputtering away of surface ions competes with implantation and the maximum in the concentration profile shifts toward the surface with increasing implant dose. Above a critical dose the maximum reaches the surface and the profile is then said to be sputter limited: further implantation will not increase the surface implant concentration. In our own work we have investigated implantation into polycrystalline but highly (110) textured thin films of SnO2 deposited on amorphous silica substrates by r.f. sputtering. A range of different dopant species including Sb [35,36], Bi [98,99] and Nb [100] have been investigated, but most work has concentrated on implantation at the fixed energy of 90keV. After implantation the dopant is activated by annealing at temperatures around 800~ in air: this process serves to restore the oxygen stoichiometry and induce recrystallisation in the amorphised surface region. The influence of implantation on sheet conductance is shown in figure 22 for Sb and Bi implants. It will be seen that Sb is a much more effective dopant than Bi. The implantation profiles are easily simulated using the computer programmes

557

E 4 3

-

c

(A)

O ,2 C

U C

oU I O ~ 0

1 ~ - ~ 2.0 40

- ~ 60

i 80

100

depth / nm Figure 23. Simulated profiles for implantation of 90keV 121Sb into polycrystalline SnO2 to different total implantation doses (a) lx1015 ions/cm 2 (b) (a) 3x1015 ions/cm 2 (c) 3x1016 ions/cm 2. Adapted from reference [37].

TRIM [101] and PROFILECODE [102] some 9 typical simulations for 90keV 121Sb into SnO 2 are shown in figure 23. The profile is predicted to be sputter limited for doses above 3x1016ions/cm 2 and the surface Sb concentration at this dose exceeds the substitutional solubility limit of Sb in SnO2. The coarse features of the implantation profiles are confirmed by depth profiling with Auger spectroscopy or secondary ion mass spectroscopy. The surface concentration of Sb can be investigated in greater detail with XPS. This technique shows (figure 24) that postannealing of low dose implanted samples leads to segregation of Sb toward the surface, whilst for the sputter limited dose, the surface concentration decreases, presumably due to Sb406 evaporation [36]. Ultraviolet photoemission studies of the thin films are of particular interest. After cleaning by vacuum annealing in UHV, the energy distribution curve for an undoped thin film shows the expected profile with a prominent bandgap shoulder due to surface Sn(ll) cations (figure 25). Sb implantation leads to the appearance of a weak conduction band feature, seem more clearly in the expansion in figure 26. The intensity of this peak correponds to a carrier concentration of 6.8x1019/cm 3, compared with an effective surface Sb concentration seen in XPS of 2.0x1021/cm 3. Thus not all surface Sb ions act as donor centres. This is confirmed by disappearance of the band gap shoulder in UV photoemission, implying

558

20

l~ Unheated l~ 350~ l~l 1000~ ;e

15

t.J

E o <

~

!0

c u'l §

n v

~

5

(a)

(b)

(r

Figure 24. Effective percentage occupancy of cation sites in Sb implanted SnO 2 thin films obtained from intensity of Sb3d3/2 and Sn3d3/2 core level peaks in normal emission AIKo~ XPS after correction for ionisation cross sections. Implantation doses are (a) 1x1015/cm 2 (b) 3x1015/cm 2 (c) 3x1016/cm 2. Unshaded bars refer to unannealed samples and shaded bars to samples subject to post-annealing at indicated temperatures for 12 hours. The percentage occupancy may be converted to a nominal mean dopant concentration (atoms/cm 3) within the XPS sampling depth by multiplying by a factor 2.79x1020. Adapted from reference [37]. replacement of Sn(ll) by Sb(lll) and trapping of electrons into "deep" states overlapping the valence band. Bi doping is seen also to suppress the bandgap states, but without indication of introduction of electrons into the conduction band. This implies occupation of surface sites by Bi(lll), but in the bulk there cannot be straightforward replacement of Sn(IV) by Bi(V).

559

xs0 Ca)

(o)

:" 9o;." -o -~

% 9

b.'..-.N.%..'-'.~.~....~.~-."~. %'.;

2"

(b) v

, .~11b"j

7

; ~ . . . . . , . , . ...~ .-.,b . . . . . ~-..." -o#.-..': -

(c)

(c)

x 500] .

-4

J

.

0 I

.

I

.

I

I

I

4 8 12 16 Binding Energy/eV

I

2O

0.0 0.5 Binding Energy ! eV

59 -I.0

-0.5

!.0

Figure 25. (Left hand panel) He(I) UPS of (110) textured SnO 2 thin films excited at hv=21.2eV after sample cleaning by annealing at 800~ in UHV. (a) Unimplanted film. Note emission intensity in bandgap region below 3.5eV binding energy and very weak intensity at Fermi energy. (b) 2O9Bi implanted film. 3x1016ions/cm 2 at 90keV. Note pronounced attenuation of bandgap emission but continued weak intensity at Fermi energy. (c) 121Sb implanted film. 3x1016ions/cm 2 at 90keV. Note attenuation of bandgap intensity compared with (a) and appearance of well-defined peak at Fermi energy in x500 expansion. Adapted from reference [98]. Figure 26. (Right hand panel). Expansion of structure close to Fermi energy in He(I) photoemission spectra. Details as in figure 22. Adapted from reference [98].

5.

SURFACE

VIBRATIONAL

SPECTROSCOPY.

With a total of 6 atoms per unit cell, SnO2 has 15 non-acoustic normal modes of vibration. From the character tables of Gay [103] these may be shown [104] to span irreducible representations as follows:

560

Alg + A2g + A2u + Big + B2g + 2Blu + Eg + 3E u

(6)

Since translation modes belong to the irreducible representatations A2u and Eu, only the modes of these symmetries are active in infrared spectroscopy. Moreover, inspection of the basis vectors listed by Katiyar [104] shows that the E u modes are polarised perpendicular to the [001] direction whereas the A2u mode is polarised parallel to [001]. Both Summitt [105] and Katiyar et al. [106] have analysed infrared reflection spectra of undoped SnO 2 single crystals using a simple dielectric model of the reflectivity in which the dielectric function in the vibrational region is written in the form:

e(o)) = ~(oo) + ~ [~.i 2 / (o)i2 - 0)2 + io.,-~,i)]

(7)

where the summation extends over Eu modes when the electric vector is perpendicular to the c-axis and over the single A2u mode when the vector is parallel to the c-axis. Within the summation -Qi, e)i and 7i represent the strength, transverse resonance frequency and damping parameter for the ith vibrational mode. An interesting feature of this analysis is that the percentage ionicity Z, which may be inferred from the strength parameters, is significantly less [106] for SnO2 (Z=49%) than for TiO2 (Z=61%). The surface loss condition in equation (4) which governs HREELS leads to excitation of modes at energies between those of both bulk transverse phonons and longitudinal phonons, for which Re[e(e))]=0. The so-called Fuchs-Kliewer surface modes [107] have energies just below those of the bulk longitudinal optical phonons and the dynamic dipole of the surface modes is along the direction of the surface normal. Under excitation by a beam of velocity v a mode of frequency co has a wavevector parallel to the surface of order ~:=~/v. This wavevector in turn governs the decay length into the solid, which is of order 1/~:. For phonons of energy of the order of 0.1eV and beam energies in the range 5eV-10eV, the decay length will be around 100A. Thus although the modes excited in HREELS are tied to the surface and do not propagate completely into the bulk, they nonetheless relate to atoms extending quite a way down into the solid. The HREELS of SNO2(110) is shown in figure 27, along with a spectrum simulated using dielectric theory [108]. In the simulation it was assumed that only E u modes could be excited in experiments on (110) surfaces because the A2u mode is polarised parallel to the surface plane. The two strong loss peaks then correspond to the two highest energy E u modes of SnO 2. The probability of exciting these modes is very high indeed and there is a significant probability of sequential

561

el)

b)

X?

I

I

I

I

0 0.1 0.2 0-3 Energy loss/eV

Figure 27. (a) HREELS of SNO2(110) excited at 5eV beam energy. (b) HREEL spectrum simulated using dielectric theory as discussed in text. Adapted from reference [108].

inelastic scattering with excitation of a second quantum of the surface mode. This gives rise to multiple scattering peaks extending over a wide energy range. These are easily reproduced in the simulation and there appears to be overall good agreement between observed and simulated spectra. This may however be largely fortuitous because more recent theoretical work on the HREELS of non-cubic material has emphasised the importance of the tensor nature of the dielectric function, which should allow the A2u mode to contribute to the observed loss spectrum [109]. In degenerately doped tin oxide, the intensity of surface phonon losses is strongly attenuated due to screening of the long range electric fields associated with surface modes by the free electron carriers [110,111]. The change in surface phonon intensity is evident in the low resolution loss spectra of figure 19 for polycrystalline Sb-doped SnO2. The high intensity of surface phonon losses in HREELS of undoped SnO2 makes it difficult to study adsorbates. This is unfortunate because the technique has the potential to provide a rich insight into tin oxide surface chemistry by vibrational fingerprinting of adsorbates [112]. If large single crystals of doped SnO 2 were to become available, the lower phonon loss intensity

562

would make adsorbate studies more tractable. Although surface sensitive infrared spectroscopy carried out in the reflectionabsorption mode (RAIRS) is a technique of growing importance in the study of the surface chemistry of metals [113], it has only recently been applied to insulating oxide surfaces [114] and there is no work to date of which we are aware on tin oxide. It is beyond our scope to deal with the very extensive literature on the study of adsorbates on high surface area tin oxide gels and powders by infrared spectroscopy: this topic is reviewed in detail in a recent important chapter by P.G. Harrison [115]. However it is appropriate to discuss briefly infrared reflection spectra

o

b

20

o ~,

20

U C

~

"6 a:

O.~~ 20

A. 1200

800

400

1200 800 Wovenumber / cm -I

400

Figure 28. (a). Experimental infrared specular reflection spectra of Sb-doped SnO2 ceramics as a function of bulk doping level. (b) Simulated spectra assuming dielectric function may be derived by adding a plasmon term to the vibrational function and averaging over E u and A2u modes according to the method of Dresselhaus [118].

563

of Sb-doped SnO 2 ceramics measured in air [116,117]. For undoped polycrystalline SnO 2 ceramics, the reflection spectrum shows a typical Reststrahlen-like profile (figure 28), which may be simulated by a procedure that takes proper account of averaging over Eu and A2u modes [118]. Simple simulations in which a plasmon term is added to the vibrational dielectric function suggest that there should be a progressive increase in reflectivity and loss of features in the Reststrahlen region as the free carriers screen out coupling between the phonon modes and the external electromagnetic field associated with the infrared radiation. At the highest doping level, clear discrepancies emerge between simulated and observed spectra. These were attributed to the formation of a depletion layer at the surface of the highly doped samples due to oxygen adsorption on surface Sb atoms. Within this unscreened layer it is possible to excite longitudinal optical phonons in the reflection experiment, accounting for the anomalous reflection dip around 800cm -1.

6. SURFACE

CONDUCTIVITY.

Sensor activity depends on band bending and associated carrier depletion due to oxygen adsorption on sensor surfaces. Practical Taguchi-type sensors are based on a poorly sintered polycrystalline material [119] and here the sensor performance is most probably related to depletion at the necks linking grains together in the polycrystalline assembly. Nonetheless, an understanding of factors governing conductivity at idealised "planar" surfaces is a prerequiste to understanding the more complex behaviour of polycrystalline material and there is by now a significant body of work in this field. It is generally presumed that oxygen vacancies in SnO 2 act as shallow donor levels [120], but we have seen in section 3.2 that surface oxygen deficiency produced for example by ion bombardment leads to states toward the bottom of the bulk bandgap. For this reason there is not a simple correlation between the degree of surface oxygen deficiency and the surface conductivity. There appears to be an overall consensus that conductivity (measured at room temperature) increases after annealing ion bombarded surfaces, but there is conflict as to the temperature of maximum conductivity. Thus Thornton and coworkers [22] found that conductivity of ion bombarded surfaces increases as the oxygen stoichiometry is restored by annealing ion bombarded SnO 2 to progressively higher temperatures up to 900K (figure 5). Above this temperature the conductivity drops sharply as the (lx2) LEED structure emerges. A similar pattern appears in the experiments of Erickson and Semancik [33]. By contrast De Fr~sart et al [32] observed a very broad conductivity maximum for annealing temperatures in the range 673K-873K. Initially thermally oxidised surfaces behave differently from bombarded surfaces,

564

in that the conductivity increases progressively with increasing annealing temperature [28] and there is a correlation between surface oxygen deficit and conductivity. The most recent work in this area is due to Cavicchi and Semancik [121,122] who have developed a method of preparing SNO2(110) l x l surfaces by oxidation in situ with an oxygen plasma torch. These surfaces had much lower conductivity than thermally oxidised surfaces and a higher O/Sn intensity ratio in ISS (but not XPS). Moreover, valence band features showed upward band bending of 0.5eV relative to the (4xl) surface, as compared to only 0.15eV for thermally oxidised surfaces. If it is assumed that the Fermi level lies 0.5eV below the conduction band minimum for the (4xl) surface, then it will be seen that only plasma oxidised surfaces represent true flat band behaviour. However, the placement of the Fermi level is somewhat controversial in SNO2(110) (4xl) and an alternative is that oxygen adsorbs on plasma oxidised surfaces to complete the 6-fold coordination of the bare tin sites of figure 2(a). This results in upward band bending and a surface depletion layer. Adsorption of submonolayer concentrations of tin onto plasma oxidised surfaces produced large increases in surface conductivity. By contrast Pd produced only small conductance changes [122]. In an experiment more closely related to sensor applications, Kohl and coworkers [123] studied changes in four-probe conductance of SNO2(101) single crystal surfaces exposed to a modulated molecular beam of CH3COOH. From a mass spectroscopic study of desorption products and of analagous interactions with C H3COOD and CH 4 it was possible to suggest a detailed scheme for decomposition of the acetic acid on the SnO2 surface. It was suggested that reversible increases in conductivity could be attributed to hydrogen donors, whereas irreversible increases were attributed to oxygen vacancies. Related experiments involving decomposition of C2H5OH on SNO2(110) have been described in conference abstracts [124,125].

7. CONCLUDING

REMARKS.

The application of UHV surface science techniques has led to considerable development in understanding the surface chemistry of SnO2. Not surprisingly most work has concentrated on relatively tractable model systems, notably single crystal surfaces. Although the interaction between SnO2 and the simple adsorbates H20 [24,27,126 ] and 0 2 [28,30,33] has received some attention, there remains a very wide gulf between the UHV surface science literature and the study of adsorbates on high surface area catalyst mimics [115]. In the sensor field there is perhaps more tangible progress in understanding the complex structural and compositional factors that determine the surface conductivity of tin oxide. There is also a growing awareness that surface spectroscopy provides a powerful means of characterising

565

the effects of addititives such as Pt [127,128] or Pd [129] on tin oxide surfaces. There remains nonetheless a need to attempt to characterise well-defined SnO 2 surfaces under the high pressure and high temperature condition relevant to oxidation catalysis and gas sensing. The technique of reflection-absorption infrared spectroscopy would appear to be of particular promise, both as a means of characterising adsorbates and of probing depletion layer phenomena.

8.

ACKNOWLEDGEMENTS.

I am deeply grateful to the many coworkers who have been involved with me in the study of tin(IV) oxide surface chemistry over the past 12 years. Special thanks must go to C. Harding, P. Tavener, P. A. Cox, W.R. Flavell, S. Eriksen, C.S. Rastomjee, R.J. Dale and R.J. Schaeffer.

9.

1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19.

REFERENCES

V.T. Agekyan, Phys. Stat. Solidi,1977, a4, 11. D. FrShlich, R. Klenklies and R. Helbig, Phys. Rev. Letters, 1978, 41, 1750. M. Nagasawa and S. Shionoya, J. Phys. Soc. Japan,1971, 30, 158. J.A. Marley and R.C. Dockerty, Phys. Rev., 1965, 140, A 304. H. KSstlin, FestkSrpeprobleme, 1982, XXll, 229. G. Sanon, R. Rup and A. Mansingh, Phys. Rev. B, 1991, 44, 5672. Z.M. Jarzebski and J.P. Marton, J. Electrochem. Soc.,1976, 123, 199C. Z.M. Jarzebski and J.P. Marton, J. Electrochem. Soc.,1976, 123, 299C. Z.M. Jarzebski and J.P. Marton, J. Electrochem. Soc.,1976, 123, 333C. F.J. Berry, Advances in Catalysis, 1981, 30, 97. Catalysis and Chemical Processes, R. Pearce and W.R. Patterson (eds), Leonard Hill, Glasgow 1981. I. Brown and W.R. Patterson, J.Chem. Soc. Faraday Trans. I, 1983, 79, 1431. J. McAteer, J.Chem. Soc. Faraday Trans. I, 1979, 75, 2768. D.R. Pyke, R. Reid and R.J.D. Tilley, J.Chem. Soc. Faraday Trans. I, 1980, 76, 1174 and references therein. M.J. Madou and S.R. Morrison, Chemical Sensing With Solid State Devices, Academic Press, Boston, 1989. J.F. McAleer, P.T. Mosely, J.O,W. Norris and D.E. Williams, J. Chem. Soc. Faraday Trans. I, 1987, 8:3, 1323. J.F. McAleer, P.T. Mosely, J.O.W. Norris, D.E. Williams, and B.C. Tofield, J. Chem. Soc. Faraday Trans. I, 1988, 84, 441. W. Mokwa, D. Kohl and G. Heiland, Sensors and Actuators, 1985, 8, 101. B. Thiel and R. Helbig, J. Crystal Growth, 1976, 32, 259.

566

20. P.A. Cox, R.G. Egdell, C. Harding, W.R. Patterson and P. Tavener, Surface Science, 1982, 123, 179. 21. P.W. Tasker, J. Phys. C, 1979, 12, 4977. 22. G.L. Shen, R. Casanova, G. Thornton and I. Colera, J. Phys. Condensed Matter, 1991, 3, $291. 23. S. Semancik and D.F. Cox, J. Vac. Si. Technol. A, 1986, 4, 626. 24. D.F. Cox, S. Semancik and P.D. Szuromi, J. Vac. Si. Technol. A, 1986, 4, 627. 25. D.F. Cox, T.B. Fryberger, J.W. Erickson and S. Semancik, J. Vac. Si. Technol. A, 1987, .5, 1170. 26. D.F. Cox, T.B. Fryberger and S. Semancik, J. Vac. Si. Technol. A, 1988, 6, 828. 27 S. Semancik and D.F. Cox, Sensors and Actuators, 1987, 12, 101. 28. D.F. Cox, T.B. Fryberger and S. Semancik, Phys. Rev. B,1988, 38, 2072. 29. D.F. Cox, T.B. Fryberger and S. Semancik, Surface Science, 1989, 224, 121. 30. D.F. Cox and T.B. Fryberger, Surface Science, 1990, 227, L105. 31. E. de Fr~sart, J. Darville and J.M. Gilles, Solid State Commun. 1981, 37, 13. 32. E. de Fr~sart, J. Darville and J.M. Gilles, Applications of Surface Science, 1982, 11/12, 637. 33. J.W. Erickson and S. Semancik, Surface Science, 1987, 187, L658. 34. P.J. Moiler and M-C. Wu, Surface Science, 1989, 224, 265. 35. V.E. Henrich and P.A. Cox, The Surface Science of Metal Oxides, Cambridge University Press, in press. 36. C.S. Rastomjee, R.G. Egdell, M.J. Lee and T.J. Tate, Surface Science, 1991, 259, L769. 37. C.S. Rastomjee, R.G. Egdell, G.C. Georgiadis, M.J. Lee and T.J. Tate, J. Materials Chemistry, 1992, 2, 511. 38. P.A. Cox, R.G. Egdell, C. Harding, A.F. Orchard, W.R. Patterson and P.J. Tavener, Solid State Communications, 1982, 44, 837. 39. R.G. Egdell, W.R. Flavell and P.J. Tavener, J. Solid State Chemistry,1984, 51,345. 40. J. F. Arlinghaus, J. Phys. Chem. Solids, 1974, 35, 931. 41. J.L. Jacquemin and G. Bordure, J. Phys. Chem. Solids, 1975, 36, 1081. 42. J. Robertson, J. Phys. C.,1979, 12, 4767. 43. A. Svane and E. Antoncik, J. Phys. Chem. Solids, 1987, 48, 171. 44. J.M. Themlin, R. Sporken, J. Darville, R. Caudano, J.M. Gilles and R.L. Johnson, Phys. Rev. B, 1990, 42, 11914. 45. P.L. Gobby and G.J. Lapeyre, Proceedings of the 13th International Conference on Semiconductors, Rome 1976, edietd by F.G. Fumi (North Holland, Amsterdam, 1976), page 150. 46. P.C. Hollamby, P.S. Aldridge, G. Moretti, R.G. Egdell and W.R. Flavell, Surface Science, 1993, 280, 393. 47. P. Tavener, Chemistry Part II Thesis, Oxford, 1982. 48. R.G. Egdell, unpublished. 49. R.G. Egdell, S. Eriksen and W.R. Flavell, Solid State Communications, 1986, 60, 835. 50. C.S. Fang, F.M. Pan, W.S. Tse and S.R. Horng, Surface Science, 1989, 211/212, 279.

567

51. 52. 53. 54. 55. 56. 57. 58. 59. 60. 61. 62. 63. 64. 65. 66. 67. 68. 69. 70. 71. 72. 73. 74. 75. 76. 77. 78. 79. 80. 81. 82. 83. 84. 85. 86.

C. Lau and G.K. Wertheim, J. Vac. Sci. Technol., 1978, 15, 622. T.W. Capehart and S.C. Chang, J. Vac. Sci. Technol., 1981, 18, 393. J.M. Thomas and M.J. Tricker, J. Chem. Soc. Faraday Trans. II, 1975, 71,329. R.G. Egdell, S. Eriksen and W.R. Flavell, Surface Science, 1987, 192, 265. D.F. Cox and G. B. Hoflund, Surface Science, 1985, 151,202. S. Munnix and M. Schmeits, Solid State Communications, 1982, 43, 867. S. Munnix and M. Schmeits, Surface Science, 1983, 126, 20. S. Munnix and M. Schmeits, Phys. Rev. B, 1983, 27, 7624. J. Terra and D. Guenzburger, Phys. Rev. B, 1991, 44, 8584. D.J. Smith, L.A. Bursill and F.J. Berry, J. Solid State Chemistry, 1982, 44, 326. K. Wakabayashi, Y. Kamiya and N. Ohta, Bull. Chem. Soc. Jap., 1968, 41, 2776. G.W. Godin, C.C. McCain and E.A. Porter, Proceedings IVth International Congress on Catalysis, Moscow, 1968. F.J. Berry and C. Greaves, J. Chem. Soc. Dalton Trans., 1981, 2447. T.S. Bush, C.R.A. Catlow, A.V. Chadwick, M. Cole, R.M. Geatches, G.N. Greaves and S.M. Tomlinson, J. Materials Chemistry, 1992, 2, 309. C.M. Freeman and C.R.A. Catlow, J. Solid State Chemistry, 1990, 85, 65. F.J. Berry, P.E. Holbourn and F.W.D. Woodhams, J.Chem. Soc. Dalton Trans., 1980, 2241. F.J. Berry and B.J. Laundy, J.Chem. Soc. Dalton Trans., 1981, 1442. A. Rohatgi, T.R. Viverio and L.H. Slack, J. Amer. Ceram. Soc., 1974, 5,7, 278. H. Kim and H.A. Laitinen, J. Amer. Ceram. Soc., 1975, 58, 23. S. Kulaszewicz, Thin Solid Films, 1980, 74, 211. H. Kaneko and K. Miyake, J. Appl. Phys., 1982, 53, 3629. M. Fujimito, T. Urano, S. Murai and Y. Nishi, Jap. J. Appl. Phys., 1989, 28, 2587. J. Kane, H.P. Schweizer and W. Kern, J. Electrochem. Soc., 1975, 122, 1144. J. Kane, H.P. Schweizer and W. Kern, J. Electrochem. Soc., 1976, 123, 270. D. B~langer, J. P. Dodelot, B.A. Lombos and J.l. Dickson, J. Electrochem. Soc., 1985, 132, 1398. P. Turner, R.P. Howson and C.A. Bishop, Thin Solid Films, 1981, 83, 253. M.H. Madhusudhana Reddy, S.R. Jawalekar and A.N. Chandorkar, Thin Solid Films, 1989, 16g, 117. R.G. Goodchild, J.B. Webb and D.F. Williams, J. Appl. Phys, 1985, 57, 2308. T. Minami, H. Nanto and S. Takata, Jap. J. Appl. Phys., 1988, 27, L287. H.S. Randhawa, M.D. Matthews and R.F. Bunshah, Thin Solid Films, 1981, 83, 267. H. Kaneko and K. Miyate, J. Appl. Phys., 1982, 53, 3629. L. Gupta, A. Mansingh and P.K. Srivastava, Thin Solid Films, 1989, 176, 33. A. Fujisawa, T. Nishino and Y. Hamakawa, Jap. J. Appl. Phys., 1988, 27, 552. E. Burstein, Phys. Rev., 1954, g3, 632. T.S. Moss, Proc. Phys. Soc. London Sect. B, 1954, 67, 775. I. Hamberg, C.G. GranquiSt, K.F. Berggren, B.E. Sernelius and L. EngstrSm, Phys. Rev. B, 1984, 30, 3240.

568

87. Z.C. Jin, I. Hamberg, C.G. Granquist, B.E. Sernelius and K.F. Berggren, Thin Solid Films, 1988, 164, 381. 88. T. Inaoka, D.M. Newns and R.G. Egdell, Surface Science, 1987, 186, 290. 89. D. Pyke, R. Reid and R.J.D. Tilley, J. Solid State Chem., 1978, 25, 230. 90. J.J. Yeh and I. Lindau, Atomic Data and Nuclear Data Tables, 1985, 32, 1. 91. Y.M. Cross and D.R. Pyke, J. Catalysis, 1979, 58, 61. 92. Y. Boudeville, F. Figueras, M. Forissier, J.L. Portefaix and J.C. Vedrine, J. Catalysis, 1979, 58, 52. 93. C.F. Wan, R.D. McGrath, W.F. Keenan, Y.S. Tung and S.N. Frank, J. Electrochem. Soc., 1988, 13.5, 985. 94. F.C. Stedile, C.V. Barros Leite, W.H. Schreiner and I.J.R. Baumvol, Thin Solid Films, 1990, 190, 139. 95. J.C. Lou and M.S. Lin, Thin Solid Films, 1983, 110, 21. 96. S.C. Chang, J. Vac. Sci. Technol. A, 1983, 1,524. 97. O. Tabata, S. Kimura and Y. Sato, Proceedings of 5th Meeting on Ion Implantation into Semiconductors, 1976, eds. F. Chernow, J.A. Borders and D.K. Brice, Plenum, New York, 1977. 98. R.S. Dale, C.S. Rastomjee, F.H. Potter, R.G. Egdell and T.J. Tate, Presented at ISSC-VIII/IVC-XlI, The Hague, October 1992. Thin Solid Films, in press. 99. R.S. Dale, Chemistry Part II Thesis, Oxford 1992. 100. R. Schaeffer and R.G. Egdell, to be published. 101. J.F. Ziegler, J.P. Biersack and U. Littmark, The Stopping Range of Ions in Solids, Pergamon Press, New York, 1985. 102. Implant Sciences Corporation, Danvers, Massachussets 1990. 103. J.G. Gay, W.A. Albers Jr. and F.J. Arlinghaus, J. Phys. Chem. Solids, 1968, 29, 1449. 104. R.S. Katiyar, J. Phys. C, 1970, 3, 1087. 105. R. Summitt, J. Applied Physics, 1968, 39, 3762. 106. R.S. Katiyar, P. Dawson, M.M. Hargreave and G.R. Wilkinson, J. Phys. C, 1971, 4, 2421. 107. R. Fuchs and K.L. Kliewer, Phys. Rev. 1965, 140, A2076. 108. P.A. Cox, R.G. Egdell, W.R. Flavell and R. Helbig, Vacuum, 1983, 33, 835. 109. A.A. Lucas and J.P. Vigneron, Solid State Commun., 1984, 49, 327. 110. P.A. Cox, M.D. Hill, F. Peplinskii and R.G. Egdell, Surface Science, 1984, 141, 13. 111. P.A. Cox and J.P. Kemp, Surface Science, 1989, 210, 225. 112. H. Ibach and D.L. Mills, Electron Energy Loss Spectroscopy and Surface Vibrations, Academic Press, New York, 1982. 113. J.T. Yates and T.E. Madey (eds.), Vibrational Spectroscopy of Molecules on Surfaces, Plenum Press, New York, 1987. 114. J. Evans, B. Hayden, F. Mosselmans and A. Murray, Surface Science, 1992, 279, L159. 115 P.G. Harrison, p. 397 in P.G. Harrison (ed.), Chemistry of Tin, Blackie, Glasgow, 1989.

569

116. P.A. Cox, R.G. Egdell, W.R. Flavell, J.P. Kemp, F.H. Potter and C.S. Rastomjee, J. Electron Spectroscopy and Related Phenomena, 1990, 54155, 1173. 117. C.S. Rastomjee, P.A. Cox, R.G. Egdell, J.P. Kemp and W.R. Flavell, J. Materials Chemistry, 1991, 1,451. 118. G.L. Doll, J. Steinbeck, G. Dresselhaus, M.S. Dresselhaus, A.J. Strauss and H.J. Zeiger, Phys. Rev. B, 1987, 36, 8884. 119. N. Taguchi, UK Patent 1,280,809. 120. J. Robertson, Phys. Rev. B, 1984, 30, 3520. 121. R. Cavicchi, M. Tarlov and S. Semancik, J. Vac. Sci. Technol. A, 1990, 8, 2347. 122. R. Cavicchi and S. Semancik, Surface Science, 1991, 257, 70. 123. W. Thoren, D. Kohl and G. Heiland, Surface Science, 1985, 162, 402. 124. H. Jacobs, W. Mokwa, D. Kohl and G. Heiland, Vacuum, 1983, 33, 869. 125. D. Kohl, H. Jacobs, W. Mokwa and G. Heiland, p. 183 in M. Che and G.C. Bond (eds.), Adsorption and Catalysis on Oxide Surfaces, Elsevier, Amsterdam, 1985. 126. D.F. Cox, G.B. Hoflund and H.A. Laitinen, Applications of Surface Science, 1984, 20, 30. 127. G.B. Hoflund, D.F. Cox and H.A. Laitinen, Thin Solid Films, 1981, 83, 261. 128. D.F. Cox, G.B. Hoflund and H.A. Laitinen, Langmuir, 1985, 1,269. 129. B-S. Chiou, J-J. Li and J-G. Duh, J. Electronic Materials, 1988, 17, 485.

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Science of Ceramic Interfaces II J. Nowotny (Editor) 9 1994 Elsevier Science B.V. All rights reserved.

571

IMPACT OF GRAIN BOUNDARIES ON PROPERTIES OF MULLITE AS A SOLID ELECTROLYTE Kazuo Yamana ~, Masald Miyamoto ~, Kenji Doi Nowotny d

b,

Industrial Research Institute of Ishikawa, Tomizu-machi, Kanazawa, Ishikawa 920-02, Japan

Tadahisa Arahori ~ and Janusz

Ceramic

Section,

Ro-1,

b Sumitomo Metal Industries, Ltd., Advanced Materials Division, l - 3 , Ote-machi, 1-chome, Chiyoda-ku, Tokyo 100, Japan c Sumitomo Metal Industries, Ltd., Advanced Technology Research Laboratories, 16, Sunayama, Hasaki, Ibaraki, 314-02, Japan d Australian Nuclear Science and Technology Research Laboratories, Menai, NSW 2234, Australia

Organisation,

Lucas

Heights

Abstract Basic properties of mullite such as chemical composition, phase stability, crystalline structure and microstructure are briefly reviewed. Effect of impurities, resulting in the formation of grain boundary glassy precipitates on high temperature properties such as mechanical properties and ionic conductivity is considered. Properties of mullite which is free of the glassy phase are analyzed as an ionic conductor for high temperature oxygen sensors. Electromotive force (EMF) of an oxygen concentration cell, based on mullite (which is free of impurities or involving an excess of silica) as an oxygen conductor, indicates that the electronic conductivity component at high temperatures assumes negligible values especially at very low oxygen activities (below 5ppm).

572 1. INTRODUCTION MuUite is a ceramic material with interesting properties and a wide range of applications. Recent studies indicate that muUite can also be applied as a sensor-type material with interesting properties as a solid electrolyte. The purpose of the present paper is to overview basic properties of muUite. Impact of grain boundaries, in particular the grain boundary glassy phase, on mechanical properties and ionic conductivity will be discussed in more detail.

2. GENERAL CONSIDERATIONS Mullite is the only stable phase of Al 2 0 3 - S i O z under atmospheric pressure. Its chemical composition varies between 3AI 2 0 3 9 2SiO 2 and 2A1203" Si02. MuUite is fabricated from clay mineral-based raw materials such as kaolinite and by adding A12 0 3. The resulting product involves a glassy intergranular phase which has negative effect on mechanical properties lowering its strength of the material above 800~ Mullite involving the glassy phase has limited applications. It has been applied as a refractory materials for fabrication of potteries [1-4]. The objective of this paper is to analyze several properties of the synthetic muUite, free of the glassy phase, as a solid electrolyte for oxygen sensors operating at high temperatures. Dokko et al.[5] have reported that the glass free mullite exhibits enhanced mechanical properties. The synthesis is based on so called molecular or/and colloidal method [6]. Thus produced material exhibits very high creep and thermal shock resistance as well as much lower thermal expansion and dielectric constant than those of the traditional material [7]. Several properties of mullite are summarized in Table 1. Schuh et al.[8] reported that silica rich muUite exhibits an ionic conductivity. According to Ficher and Janke [9], the mullite phase becomes a pure ionic conductor above 1000 ~ This property make mullite an excellent candidate for oxygen sensors to special applications such as the determination of oxygen activity in steel melts during steel making processes instead of

573

Table 1. P r o p e r t i e s of commercial m u l l i t e alumina[2].

unit

in comparision with those of

Mullite 3A1203 92Si02

Alumina AI20a

S i02

~

38

O. 1

A1203

%

60

99.5

CaO

%

O. 4

MgO

%

-

O. 15

NauO

%

O. 2

O. 05

Specific gravity

-

2.7

3.9

Water absorpt i on

%

O. 0

O. 0

Compress ive s t r e n g t h

blPa

1,400

3, 000

F I exural s t r e n g t h

MPa

180

300

Refractoriness

~

1,850

2, 000

Chemical compos i t i on

Thermal c o n d u c t i v i t y

kcal/mh~

2, 7

7.2

xlO-6/~

5, 6

8. 1

(400 ~ Coef. of thermal expans i on (20-1000 ~ Serv ice temperature

Electrical resistivity

max i mum

~

1,650

1,930

common

~

l, 600

I, 850 > 10 '3

20 ~

ohm 9cm

> 10 '3

500 ~

ohm 9cm

2xlO 8

109

574 zirconia-based sensors. The oxygen sensor based on mullite will be described in more detail.

3. BASIC PROPERTIES 3.1. Phase diagram in the A1 2 0 3 -SiO 2 system There have been many conflicting data and interpretations on the stability of the mullite phase. One may expect that all these apparent conflicts can be explained by differences in both chemical omposifion of starting materials and the experimental procedure applied to faboficate muUite. A detailed phase diagram for the AI 2 0 3 - S i O 2 system was reported by Bowen and Greig [10]. They claimed that mullite melts incongruently at 1810 ~ Aramaki and Roy [11] reported that mullite melts congruently at 1850 ~ involving a solid solution varying between 71.8 wt% of Al 2 0 3 & 28.2 wt% of Si02 (3:2 ratio) and 74.3 wt% of A 1 2 0 3 & 25.6 wt% of SiO~(Fig.1). A metastable solid solution may reach up to 77.3 wt% of A1 2 0 3 (2:1 ratio). According to the phase diagram of Aksay and Pask [12], determined by using a diffusion couple technique, mullite melts incongruenfly at 1828 ~ (Fig.2). 3.2.

Crystal structure The x - r a y diffraction pattern of muUite is similar to that of sillimanite which is a stable phase under high pressures and at elevated temperatures. Thus mullite structure can be derived from the sillimanite structure by substituting alumina by silicon in the tetrahedral sites and by removing appropriate oxygen atoms as required by the neutrality condition [13]. Basically, the mullite structure is an orthorhombic one. It consists of octahedral AlO 8 chains parallel to the c-axis that are cross-linked by tetrahedral (A1Si)O4 chains (Fig.3)[14]. Variation of its composition from 3 A 1 2 0 3 . 2Si02 to 2 A 1 2 0 3 . SiO2 results in changes of the lattice parameters (increase of a and c and decrease of b ) [15]. Thus the following formula for mullite can be proposed 9 IV AI I(AI IVz,zx Si2-2x) 010-x

(1)

575

A 1,0~ [mo] %] 20 i

40

i

i

2000 -

60

i'

i

i

80 "

i

i

-

100 i

Mullite(ss) ) liquid

Liquid

l==---i

l=_..a

.

,

1800

1800

1400

,

I

I

,

,I

a

:o

i

4o

eo

so

1-oo

Al ~O~ [wt%] Fig. 1. The phase diagram of the hl2Oa-Si02 and Roy [ l l ] .

A1 0

20 9

"

system a c c o r d i n g

Iwt%J

40

60

51

w

80 9

9.

/

~c3

//."

.~/ I [

!-

]400L

I 0

100 .

.

.

.

Liquid

2000-

]6oo

to Aramaki

Hu]lite(ss)+Lzquld

.-7--,-. -"7 ~'~,. / i ~. ", I ! ',:.

Si I icaOlul I i t e ( s s ) ; i i . ',. 20 40

Alumina+Liquid

I1: JJ. /

I~

i

I: .ti

60

i Alu, ina

]

,

]

!

Mullite(ss) i

I

i

~

. 80

I

I

j l O0

A1 ~.03 [mo 1%] Fig. 2. The phase diagram of the A1203-Si02 system according to Aksay and Pask [12].

576

(a)

(b)

Fig. 3. Crystal s t r u c t u r e of m u l l i t e projected on the (001) plane : ( a ) h y p o t h e t i c a l " p e r f e c t " m u l l i t e (dark atoms are located at z=O) and (b) average unit cell of m u l l i t e (heavily outlined s i t e s have p a r t i c a l occupancies and r e p r e s e n t displaced atoms due to the presence of oxygen in 03 c o n f i g u r a t i o n s ) [ 1 6 ] .

9 II

0 II,Si

0

|

Fig. 4. Model for l a t t i c e r e c o n s t r u c t i o n around an oxygen vacancy (dashed c i r c I e).

577 where x is the number of oxygen atom missing in average unit cell, and 1V and VI represent coordination states of the cations [16]. Both x - m y and electron diffraction studies indicate a tendency of oxygen vacancies to order resulting in the formation of ordered structure within the mullite composition range between 3:2 and 2:1 of AI 2 0 a to SiO, ratio. Fig.4 shows the structure model illustrating a reconstruction around oxygen vacant site (dashed circle). The AI and Si cations, which lose their tetrahedral coordination because of oxygen removal, are shifted by about 0.12 nm to new lattice sites (indicated by T * in Fig.3). These displaced atoms become fourtly recoordinated through an associated motion of an adjacent oxygen atom into the position indicated as 0 * in Fig.3. 3.3.

Microstructure Mullite, fabricated traditionally from raw materials, involves a glassy phase in the form of an intergranular precipitate. Chemical compositions of the glassy phase in both sintered and fused mullite phases are shown in Table 2. The amount of the glassy phase in refractory materials depends on the content of its impurity (Fig.5). During growing a mullite single crystal by using both Czochralski method and the floating zone method, the glassy phase is located in spaces between the muUite crystals (Fig.6) [17-19]. If alumina and silica are mixed on a molecular level, then the muUite formation becomes controlled by a nucleation process. Then the glassy phase is not present [20-23]. Therefore, the mullite powders is an important source of a glass-free mullite phase [2]. Fig.7 illustrates the SEM photograph of a polished and then thermal-etched muUite specimen which was sintered from synthetic materials at 1650 *C. 3.4.

Mechanical properties Mullite has commonly been applied as a refractory material in construction of heating tin'races. It exhibits an excellent resistance to creep, thermal shock, abrasion, spalling and corrosion of acid slags. However, mullite has never been regarded as a suitable material for high-strength applications at elevated temperatures. The preparation of muUite, which is free of glassy phases, results in a wide range of application as a structural material of good mechanical properties. Fig.8 illustrates a relationship between the bending strength and temperature for the synthetic mullite in

578

T a b l e 2. Chemical c o m p o s i t i o n s of r e f r a c t o r i e s (made of s i n t e r e d and f u s e d mul I i t e ) and t h e g l a s s y p h a s e s i n v o l v e d in t h e refractory material.

BULK

GLASSY PHASE

Sintered mul I i te

Fused mull i te

Sintered mul 1i te

Fused mul I i te

Si02

27. 98

21.39

56.95

21.39

AlcOa

68.81

74.25

28.32

33.71

Fe~03

1.54

O. 54

5.66

5.13

TiO~

O. 40

2.81

1.75

16.96

CaO

O. 18

O. 31

O. 21

2. 46

MgO

O. 08

O. 10

O. 10

O. 45

Na20

O. 16

O. 02

O. 62

O. 45

K~O

O. 75

O. 02

6.39

O. 45

(unit : wt~;)

Amount of glassy phase

12. 5

5.2

20

10

0

0

I

I

0.4

i, 8

I

J

. .,I

,

2,0

I

2,4

2,8

IHPURITIES[Fe20~+TiO~CaO+~gO+Na20*K~O](wt%)

Fig. 5. R e l a t i o n s h i p refractories.

between g l a s s y

p h a s e s and i m p u r i t i e s

in m u l l i t e

579

Fig. 7.

SEM micrograph of the thermally etched surface of the synthetic mullite.

580

~

400 ~-

99Z Altmina

~a

'

n

'

'

-

~

1

~

|

0

300

100

0

200

400

600

800

I000

1200

1400

Fig. 8. Relationship between bending strengthes and temperature for synthetic mullite, t r a d i t i o n a l mullite, and 99% alumina.

68A

600

500

J,;'

I

400 Z

300

200

~" RT

.-~\ --- -

60A /~" 99%Alumina I000

71.8A\\

~

II00

TEMPERATURE

7'

". 1200

1300

1400

[~3]

Fig. 9. Relationship between the bending strength and the temperature for d i f f e r e n t alumina contents 9 60A, 68A, 71.8A (mullite + s i l i c e o u s glassy phase), 74A (mullite + alumina) and 99A (alumina).

581 comparision with those of alumina and natural mullite [3]. The natural muUite has considerably smaller strength than that of alumina, and nearly the same strength as the synthetic mullite. However, the strength of alumina decreases gradually above 1000 ~ On the other hand the synthetic mullite still exhibits its strength at room temperature. Moreover the strength even increases at 1300 ~ The more detailed data of strengthes on synthetic mullite at different alumina content in the Al, 0 3 - S i O , system were reported by Kanzaki [24]. As shown in Fig.9, only silica-rich specimens results in an increase of strength, while alumina-rich specimens exhibit a degradation effect. It is thought that the observed increase in the bending strength above 1200 ~ is due to the presence of siliceous glass phases [25]. 3.5.

Thermal shock resistance Data on the thermal shock resistance of the commercial muUite, collected in Table 3, correspond to the temperature difference between the specimen and water, if the water quenching procedure is applicable. Table 3 also gives the amount of the glassy phase [2]. Fig.10 shows the relationship between the temperature difference of the quenching procedure ( t l T ) and the bending strength after the water quenching. As seen A T gradually increases with the amount of the glassy phase although this difference is relatively small. One may thus assume that the glassy phase, inbetween the mullite grains, has a tendency to lower the crack propagation which apparently affects the strength. The thermal shock resistance of a mullite tube, from the viewpoint of its application for construction of a high temperature oxygen sensor, was discussed in ref. [26]. The analysis of the thermal stress distribution was performed for a mullite tube immersed into iron melt at 1600 ~ by using a finite element method [26]. As shown in Fig.11 the temperature of the tube increases with the elapsed time. After 0.15 seconds the temperature difference between the inner and the outer part of the tube is 300 ~ The stress distribution on the cross section in tube after 0.15 seconds is shown in Fig.12. The maximum stress results in a center of the inner part of the tube. Its value is 257 MPa. As it was already reported muUite shows the tensile stress about 130-140 MPa [26]. Thus it has been suggested that the thermal shock fracture occurrs in the muUite tube. In order to protect the thermal shock fracture, plastic-coating and porous tubes have been employed [26].

582

Table 3. P r o p e r t i e s of commercial m u l l i t e s employed in thermal shock tests.

Sample

Composition A120a/Si02 (wt%)

Density

Grain structure

Glassy phase (wt%)

I

0

A

47 / 49

2.50

Bullite

B

55 / 41

2.65

Mullite

C

60 / 38

2.70

Mullite

D

67 / 31

2.75

Mullite

E

76 / 23

2.98

Mul I i te+A120a - - - -

I

I

10

20

'G

200 ;=I:2

E--,

~

D

m

o . e .... .c........... 0"0

B

~E}--

...... A

"E~-'I ~

100

z Z

200

2 0 AT [~]

-3i0

Fig. lO. Thermal shock r e s i s t a n c e of the commercial m u l l i t e corresponding to water quenching. Symboles A, B, C and D correspond to Table 3.

I

30

583 --

,

,..

1500

A lO00

~--~

500

._,

,

,

,

,

I 0.5

1

I

I

1.0

1.5

2.0

ELAPSED TIME[SEC, ] Fig. ll. The outer and the inner surface-temperatures of the mullite tube in relation to theelapsed time a f t e r immersing into steel melts.

300

200I '-"2 101

~ N

-~oo[

~ -2oo Fig. 12. Relationship between the calculated s t r e s s (after O. 15 sec.) and the

~ -300 IF

wall thickness on cross section L

J

..,

i

waI l thickness ins ide outside B A

(A-B) in the tube shown in Fig. ll.

584 4. MULLITE AS AN OXYGEN CONDUCTOR Oxygen conduction in solids requres that predominant lattice defects are oxygen vacancies which exhibit high mobility such as the defects in yttria-doped zirconia. As mentioned above muUite is a very good oxygen conductor at high temperature. The oxygen vacancies in mullite may be formed as a result of the following reaction 9 A 120a-[-28 i si+Oo

=

2A I si+V 6 + S

i 02

(2)

where Sis~ indicates Si on Si sites, O o indicates 0 on 0 sites, Alsi indicates A1 on Si sites and V 6 indicates vacancy on 0 sites. Fischer and Janke [9] have reported the electromotive force (EMF) of the oxygen concentration cell of the following compositions 9 air / mullite / H 2 0, H 2

(3)

air / muUite / iron melt (1600 ~ )

(4)

air / mullite / AI melt saturated with A1 2 0 3 (1000-1600 ~ )

(5)

where the mullite phase involved an excess of the silic acid. It has been observed that the mullite phase enriched with Si02 exhibits good ionic conductivity at 1600 ~ at p(O 2 ) value above 10-8 atm (Fig.13). The studies indicate that mullite exhibits much better properties for construction of an oxygen sensor in steelmaking than sensors based on zirconia. Similar studies were also reported by Koller [27]. At high temperatures and at low oxygen activities zirconia exhibits mixed (both electronic and ionic) conductivity. The EMF can be expressed as

E- R T P'(o2)1". PH I/4 F I n P(oz)V4+ p~11/4 ....

(6)

where P(O~.) and P'(O~.) are the partial pressures of oxygen at two electrolyte interfaces, P . is defined as the oxygen partial pressure at which both ionic and electronic conductivity components are equal, R is the gas constant, F is the Faraday's constant, and T is the absolute temperature. The tube-type mullite probe was manufactured from high purity materials

585

2500 ! : ~

I / ~

2000 -

L..__j

\%,\\%\% "

~'

\

~.

\

\

\

\

,%\ ~,<..,.'-~.,, -...:-,~\.~\

1500

1000

500

-30

j

,,

I

--20

i

I

--10

....

I

Log P(O~o) [P(Oz) in Pal

J

\

0

Fig. 13. Relationship between the electromotive force (EMF) and the oxygen partial pressure for the mullite e l e c t r o l y t e .

586 coprecipitated by a sol-gel method [28]. Studies on oxygen sensor by using a mullite tube were also performed by Kawai et a1.[29]. Extensive studies were performed by Suito et al. [30]. Their measurements were carded out by using the equipment shown in Fig.14. A mullite probe was dipped into an iron melt and the EMF values were measured at 1600 *C. The response of the mullite probe remained within 10 seconds. Suito et al.[30] have reported the measurements of oxygen activity in the iron melt at different content of A1 as deoxidiser. They have assumed that (1) the mullite solid electrolyte exhibits pure ionic conduction, and that (2) the thermodynamic equilibrium at the elecrolyte/electrode interface is reached. As shown in Fig.15, oxygen activities ( a o ) obtained from EMF measurements are plotted against the Al content, which corresponds to the dissolved Al in the melt as well as the Al ~ O a inclusions. As seen the a o values for the A1 contents below 0.5 wt% are in a good agreement with the Nemst dependence. Fig.16 shows that the values on the zirconia-plug probe corrected by the electronic conductivity component are in an excellent agreement with those of the mullite probe obtained by using the Nemst equation. Suito et al.[30] have concluded that the electrical conductivity of mullite is lower than that of zirconia (by about three orders of magnitude) mainly due to the electronic leakage of zirconia at elevated temperatures. This may indicate that muUite-type oxygen sensor exhibits superior properties to those of zirconia-type sensor especially at very low oxygen activities. Fig.17 illustrates the oxygen activity values determined by using the stoichiometric mullite probe vs. the EMF obtained by using the mullite probe with an excess of either Al ~ O a or SiO ~. As seen the a o values obtained with the mullite probe containing an excess of A1 ~ O a are slightly higher than those determined by using the probe based on muUite of stoichiometric composition. On the other hand the oxygen activity values of mullite involving an excess of SiO ~ are lower than those corresponding to the stoichiometric mullite phase. In order to evaluate the performance of mullite-based oxygen sensor at very low oxygen activities. Nagatani et al. [ 3 1 ] have determined the parameter P n in equation (6) and illustrates the relationship between the EMF determined for the mullite-based sensor and the oxygen activity for the zirconia sensor. As seen the EMF obtained by using the muUite probe is in a good agreement with the Nemst law at very low oxygen activities. On the other hand the EMF values for the zirconia-based sensors exhibit substantial

587 Deoxidized Ar (CO,Ar-5~H2)

I

Mo rod Alumina tube

L-I

m

V-

Water-coo led cap

Alumina reaction tube

Crucible (alumina, graphite) Z i r c o n i a sensor

r-711

J

. .I.I. I

nN

!

Iron melt

"

Zirconia tube

v.

L~r "

Fig. 14.

ldullite sensor

~

Deoxidized Ar

Thermocouple

S c h e m a t i c d i a g r a m of e x p e r i m e n t a l oxygen a c t i v i t y

in i r o n m e l t .

apparatus

for

measuring

588 t

-2 -

t

I

I

1873K _

9

r-----i

9

9

Q

-3

0

0

-4

_

_

9

AI/A120.~ c~. -4

I

I

i

I

-3

-2

-1

0

1

Log[mass-% A1]obs. Fig. 15. Oxygen a c t i v i t y of iron melt measured by the mul I ite-based sensor vs. Al content (deoxidizer) [30].

-2

I

I

r--"-'-I

c~

[mass-% AI_] _g O. 5

, 9r - - , i

o r

-3 -,---4

--4 o

-5

I -5

-4

-3

-2

Log ao[EMF, M u l l i t e ] Fig. 16. Comparision between the oxygen a c t i v i t i e s measured by the zirconia plug-type probe and the mullite probe.

589 --2

-

I

;:~ - 3 0 o~ t:uO

o

--4

/ -5 . . . . -5

Al20.-rich @ Si02-rich 0 l .... -3

I

-4

L o g a o [EMF,

-2

Stoichio~etric]

Fig. 17. Comparision between oxygen a c t i v i t y measured by using mullite with an excess of Al20a and Si02 on one side and the s t o i c h i o m e t r i c mullite.

600

,

,

t

"!

1873K

,

PH

. . . . 15.29(Zr-7M)[32] ~f. =1 ...... 14.04(Zr-9M)[33] 400 . . ~ ~ o n u I O n x -----13.51(Zr-9M)[34] t_.____s

200

-~--~sH82(~ullite) "----13.15(Zr-9M)[31] -200 0.1

| 1

- - ~ .... ! 10

O0

ao [ppm] Fig. 18. Relationship between electromotive force (EbtF) and oxygen a c t i v i t y at various P. values.

590 deviations from the theoretical dependence at very low oxygen activities (below 5ppm). These experimental data indicate that the muUite-based oxygen sensors exhibit much better characteristics at lower oxygen activities than that of zirconia sensors.

5. CONCLUSIONS

It has been documented that mechanical & electrical properties of the mullite phase are in a large extent determined by its grain boundary composition. It has been shown that mullite fabricated from clay-type raw materials involves an intergranular glassy phase precipitate which has negative effect on mechanical properties. Elimination of this glassy phase by either powder processing or by using alumina and silica free of impurities results in substantial improvement of its mechanical properties. Analysis of electrochemical properties of muUite indicates that muUite-based oxygen sensors exhibits much better characteristics than that of sensors based on zirconia. It has been documented that at 1600 ~ muUite exhibits pure ionic conductivity especially at very low oxygen activities. ACKNOWLEDGEMENTS This work was supported by the Ishikawa Prefecture. gratefully acknowledged.

This support is

6. REFERENCES 1 2 3 4 5 6 7 8

Mullite and muUite matrix composites,S.S0miya, R.F.Davis and J.A.Pask (eds.), Am.Ceram.Soc.,Inc., Westerville, Ohio 1990. Mullite, S.Somiya (ed.), Uchida Rokakuho, Tokyo 1986. Mullite 2, S.Somiya (ed.), Uchida Rokakmho, Tokyo 1988. High temperature oxides 4, R.F.Davis and J.A.Pask (eds.), Academic Press Inc., London 1971. P.C.Dokko, J.A.Pask and K.S.Mazdiyasni, J.Am.Ceram.Soc., 60 (1977) 150. I.A.Aksay and S.M.Wiederhorn, J.Am.Ceram.Soc., 74 (1991) 2341. I.A.Aksay, D.M.Dabbs and M.Sarikaya, J.Am.Ceram.Soc., 74 (1991) 2343. B.Schuh, B.Korousic and B.Marincek, Schweizer Archiv, 12 (1968) 380.

591 9 W.A.Fischer and D.Janke, Arch.Eisenh Uttenwes., 40 (1969) 707. 10 N.L.Bowen and J.W.Greig, J.Am.Ceram.Soc., 7 (1924), 238, 410. 11 S.Aramaki and R.Roy, J.Am.Ceram.Soc., 45 (1962) 229. 12 I.A.Aksay and J.A.Pask, J.Am.Ceram.Soc., 58 (1975) 507. 13 W.H.Taylor, Z.Krist., 68 (1928) 503. 14 R.Sadanaga,M.Tokonami and Y.Takeuchi, Acta CrystaUogr., 15 (1962) 65. 15 S.O.AgreU and J.V.Smith, J.Am.Ceram.Soc., 43 (1960) 69. 16 T.Epicier, J.Am.Ceram., 74 (1991) 2359. 17 W.Guse and D.Mateika, J.Crstal Growth, 22 (1974) 237. 18 W.Guse, J.Crystal Growth, 26 (1974) 151. 19 K.Yamana,M.Tokonami,K.Nobugai,N.Morimoto,I.Shindo and M.Koizumi, J.Am.Ceram.Soc., 61 (1978)63. 20 D.X.Li and WJ.Thomson, J.Am.Ceram.Soc., 74 (1991) 2382. 21 B.L.Metcalfe and J.H.Sant, Trans.J.Brit.Ceram.Soc., (1975) 193. 22 H.Ohnishi,T.Kawanami,A.Nakahira and K.Niihaxa, J.Ceram.Soc.Japan, 98 (1990) 541. 23 K.S.Mazdiyasni and L.M.Brown, J.Am.Ceram.Soc., 55 (1972) 548. 24 S.Kanzald, Ceramics Japan, 23 (1988) 1061. 25 T.Mah and K.S.Mazdiyasni, J.Am.Ceram.Soc., 66 (1983) 699. 26 K.Yamana, M.Miyamoto,K.Doi,T.Arahod,Y.Nakamura and F.Terasaki, J.Ceram.SocJapan 101111] (1993) 1214. 27 A.KoUer, Sklar a keramik 43 (1974). 28 K.Yamana,M.Miyamoto,K.Doi,T.Arahori,Y.Nakamura and F.Terasaki, Report.lnd.Res.Inst.Ishikawa, 41 (1993)49. 29 C.Kawai,T.Arahod,K.Yamana,M.Miyamoto,K.Nakagawa,H.Suito and R.Inoue, Japan Patent, Tokuganhei No.28690 (1991). 30 H.Suito,R.Inoue and A.Nagatani, Steel Research, 63 (1992) No.10. 31 A.Nagatani,R.Inoue and H.Suito, J.Appl.Electrochem., 22 (1992) 859. 32 D.Janke and W.A.Fischer, Arch.Eisenhuttenwes. 46 (1975) 755. 33 M.lwase,E.Ichise,M.Takeuchi and T.Yamasaki, Trans.Jpn.lnst.Met., 25 (1984) 43. 34 MJ.T.van WijngaardenJ.M.A.Geldenhuis and RJ.Dippenaar, J.Appl.Electrochem., 18 (1988) 724.

This Page Intentionally Left Blank

Science of CeramicInterfacesII J. Nowomy(Editor) 1994 ElsevierScienceB.V.

High temperature embrittlement c o m p o s i t e s - interface effects

593

of ceramic

matrix

J. L. Cocking Ship Structures and Materials Division, Materials Research Laboratory, DSTO Australia PO Box 50, Ascot Vale, Victoria, 3032, Australia

Abstract A continuous Tyranno (Si-C-Ti-O) fibre-reinforced barium magnesium aluminosilicate (BMAS) glass was mechanically tested at room temperature and at 1000~ Auger electron spectroscopy, scanning electron microscopy, electron probe microanalysis and transmission electron microscopy were used to examine the fibres, the matrix and the fibre/matrix interface of this composite, both prior to and after testing at 1000~ At room temperature the composite contained a carbon interlayer between the fibres and the matrix and behaved in a tough manner during testing. After exposure to air at 1000~ the interface region degraded, which resulted in the formation of bubbles enclosed by a viscous Sirich oxide, and Ca- and As-containing precipitates. The reaction products bonded the fibres to the matrix and the composite exhibited brittle behaviour.

1. INTRODUCTION Successful design of ceramic matrix composites for use at high temperatures is dependent on knowledge of the thermomechanical and thermochemical properties of the phases present in the composite. It has been established [1,2] that ff strong bonding exists between the fibres and matrix of a continuous fibre reinforced composite, a propagating crack will pass through the matrix and fracture the fibres in an almost planar manner. If, however, the bonding is weak the work of fracture is increased, as the fibres will initially debond from the matrix, then fracture and pull out of the matrix. Pullout of the fibres is affected by mechanical bonding between the fibres and the matrix which, in turn, is affected by the relative coefficients of thermal expansion of the fibres and the matrix [3] as well as by friction due to irregularities on the surface of the fibres

[4]. At high temperatures, bonding between the fibres and the matrix may be affected in a number of ways, including:

594 interdiffusion of elements between the fibres and the matrix; gaseous species from the atmosphere diffusing down the fibre/matrix interface and reacting with either or both the fibres and the matrix; degradation of carbon (or other) interlayers; and differential thermal expansion of the fibres and the matrix. An understanding of the evolution of the structure and composition of the interface region with time, at temperatures and in environmental conditions likely to be experienced during service, is a necessary step in the determination of the behaviour of candidate composite systems. In this paper, the results of characterisation of the interface region between continuous Tyranno (Si-C-O-Ti) fibres in a barium magnesium aluminosilicate (BMAS) matrix, at room temperature and after exposure in air at 1000~ are presented, together with a discussion of how changes in the interface region contributed to the change in behaviour of the composite from tough at room temperature to brittle at high temperature. 2. E X P E R I M E N T A L

2.1 Nominal Composition The composite consisted of a matrix of barium magnesium aluminosilicate (BMAS) glass, with a nominal composition of barium osumilite (BaO.2MgO.3A12Oa.9SiO2), reinforced with Grade F, multii~ament Tyranno (Si-CTi-O) fibres. The diameter of the fibres was approximately 8.5~m. The volume fraction of the fibres in the composite was 36%. The lay-up of the plies in the composite was 00/90 ~, with eight layers arranged symmetrically either side of the central axis. The composite was supplied as tiles approximately 150mm square and 3ram thick.

2.2 Mechanical Testing Beam samples nominally 4ram x 3mm in cross-section were cut from tiles for four point bend testing. The surfaces were cut and fine parallel ground using diamond tooling. Notched samples for fracture in the scanning Auger microprobe were 3ram in cross section and 22ram long. The notch was cut 11.4mm from the top of the sample and was 0.4ram deep. The room temperature 4-point bend strength was measured using a support span of 20ram, an inner span of 10ram and a crosshead speed of 0.254ram/rain. For high temperature testing a 4-point rig with a support span of 30mm, and an inner span of 15ram, was used. The crosshead speed was 0.2ram/rain. The furnace took approximately one hour to reach 1000~ from room temperature. Some samples were inserted into the furnace for a "hot" start from around 200~ rather than room temperature. The furnace took approximately 20 minutes to reach 1000~ from the "hot" start temperature. Samples were given a 15 minute soak time at 1000~ in air prior to applying load.

595

2.3 Instrumental techniques Characterisation of the composite was accomplished by a combination of scanning electron microscopy (SEM), Auger electron spectroscopy (AES), electron microprobe analysis (EPMA), x-ray diffraction (XRD) and transmission electron microscopy (TEM). Where possible, the same section of sample was characterised by using more than one technique. SEM studies were carried out using a Cambridge Stereoscan 250 Mark 2, fitted with a Link energy dispersive analyser. The voltage most commonly used for inspection and analysis was 15keV. TEM was undertaken on a Philips CM30 analytical transmission electron microscope fitted with an EDAX thin window energy dispersive analyser. AES analyses were performed using a Perkin Elmer Model 610 scanning Auger microprobe. This instrument has a cylindrical mirror analyser, which is co-axial with the electron gun to ensure that shadowing does not affect analyses on rough surfaces. Most of the AES analyses were done using a 2keV, 5hA electron beam, which resulted in a lateral resolution of approximately l~m. The energy resolution was 0.6%. Examination of the layers below the surface was accomplished by ion milling with 3.5keV Ar ions generated by a differentially pumped ion gun. The base pressure during analysis was 3 x 10-Spa (2 x 10-1~ Tort). During ion milling this rose to 1 x 10-6 Pa (1 x 10-1~ The ion sputter rate of 20nm/min was calibrated using 100nm thick Ta205-coated Ta. The EPMA analysis was conducted with a Cameca SX50 electron probe microanalyser, using a 15keV or 20keV, 20nA electron beam. The instrument has four wavelength dispersive spectrometers (WDS), one of which is inclined for rough surfaces, and an energy dispersive (EDS) analyser. It is capable of determining all elements of greater atomic number than Be. Fractured samples were mounted and positioned with the fracture surface to be examined as near as normal to the beam as possible. For quantitative EPMA analysis of the matrix, polished cross sections of the composite and the standards were first carbon coated. The samples and standards were subsequently lightly polished to remove the carbon layer. They were then coated with a thin layer of aluminium for analysis of the fibres. The standards used for the quantitative analyses of the matrix were NBS glass STD K309 for Si, Ca, O, Ba and A1. MgO was used for Mg. The standards used for quantitative analyses of the fibres were pure elements for Si, C and Ti and MgO for O. X-ray diffraction (XRD) data was collected using a Siemens D500 x-ray diffractometer. The patterns were recorded using Cu Ka radiation, over a 20 range of 10 to 167 ~ at 0.5~ with an accelerating voltage of 40keV, and a current of 40mA.

596 3. RESULTS

3.1 Room Temperature Mechanical Behaviour The room temperature 4-point bend strength was 312+25MPa. Optical and scanning electron microscopy of the samples after testing revealed the fracture mode was an angled shear band through the sample. All samples tested in four point bend at room temperature showed extensive bending, but no samples completely fractured. In the plies in which the fibres were laid across the sample, as shown in Figure 1, the test produced extensive multiple-branched cracks running around the fibres. The fracture surface produced when the samples were fractured in the vacuum of the scanning Auger microprobe, Figure 2, showed the extensive fibre pull-out characteristic of tough behaviour. 3.2 High Temperature Mechanical Behaviour The high temperature 4-point bend strength was 177+40MPa. Thus, the flexural strength of this composite at 1000~ was reduced to between 42% and 75% of samples tested at room temperature. The greater spread in the results at

Figure 1. The fracture mode after four point bend testing at room temperature showed extensive multiple-branched cracks (arrows) running around the fibres.

597 high temperature was attributed to the different times taken to reach 1000~ as this may have affected the development of the reaction products discussed below: The samples failed in a brittle manner at 1000~ with almost no fibre pullout over the entire fracture surface. In a few limited areas, which tended to be located near the centre of the sample, a limited amount of fibre pullout did occur. An example of the appearance of a failed sample is shown in Figure 3. The fibres were strongly bonded to the matrix and an examination of the fracture surface showed the passage of the crack through, rather than around, the fibres. A sample which had been heated to 1000~ for the same length of time as the four point bend samples, then cooled and fractured in the scanning Auger microprobe showed greater fibre pullout near the centre of the sample than the flexure samples, but brittle behaviour near the edges of the sample. 3.3 C h a r a c t e r i s a t i o n o f t h e r o o m t e m p e r a t u r e c o m p o s i t e XRD of the composite revealed that the major phase in the matrix was barium osumilite (BaO.2MgO.3A1203.9SiO2) , with a major second phase, celsian (BaO.A1203.2SiO 2) and minor phases, mullite (2SIO2.3A1203) and cordierite (5SiO2.2A1203.2MgO). The Tyranno fibres were amorphous.

Figure 2. Fracture surface of the composite, after it was fractured in-situ in the scanning Auger microprobe, demonstrating extensive fibre pullout at room temperature.

598

Figure 3. Four point bend at 1000~ showing initial brittle fracture and subsequent delamination along the neutral axis. Examination of the fracture surface near the edge of the crack, at (b) shows the lack of fibre pullout and the passage of the crack through, rather than around, the fibres.

599 Table 1 Bulk concentration, in weight percent, prior to testing as measured by EPMA Element Bulk Matrix Fibre Centre

Ca 0.37

O 45.53

Ba 11.53

Si 23.43

A1 14.67

Mg 3.50

C ND*

Ti ND*

Total 99.03

ND*

19.42

ND*

48.60

ND*

ND*

29.20

2.07

99.29

*ND = not detected

Representative quantitative EPMA analyses of the matrix and the fibres in a polished cross section of as-fabricated tile are presented in Table 1. EPMA analyses also revealed the presence of calcium in low concentrations, 0.37%, in the bulk matrix as well as in a layer less that 0.5~m wide, at the outer surface of the fibres. The area enriched in Ca, shown in Figure 4, was revealed as a bright

Figure 4. EPMA backscattered electron image of a polished cross-section of the as- fabricated tile. The bright halo at the outer edge of the fibre is slightly enriched in Ca.

600 annulus around the fibres when the sample was examined in backscattered electron mode. The concentration of Ca in this layer could not be measured in a truly quantitative manner, as the band was too narrow, which led to dilution effects in the measurements. The estimated concentration of Ca in this region was of the order of 1%. As can be seen in Figure 4, there is a band, which is dark in the backscattered electron image, surrounding the fibres. EPMA analysis could not definitely resolve whether this band was carbon rich, a crack between the matrix and the fibre, or a combination of both. After samples were fractured in the scanning Auger microprobe, Auger spectra were collected from fibres, from troughs in the matrix exposed by fibre pullout, and from the matrix. Several areas and points were selected and these areas were re-analysed during and after ion-milling. AES analysis showed that a thin layer of carbon was present on the outer surface of the fibres after fracture. The carbon layer was of the order of 75nm thick. Beneath the outer layer of carbon, the elements detected in the fibres, in order of decreasing peakto-peak height, were carbon, silicon and oxygen. At some points on the fibres aluminium was also detected beneath the carbon layer. The analyses could not detect the presence of calcium on the outer surface of the fibres. Carbon was also the main element detected on the surface of the troughs. Generally the carbon was thinner on the surface of the grooves than it was on the surface of the fibres. Analysis of the matrix was dif~cult, due to charging. The only elements commonly detected in the matrix were oxygen and silicon. Aluminium was detected less commonly and signals for other elements could not be discriminated from the noise resulting from charge effects. Transmission electron microscopy (TEM) analysis confirmed the presence of an amorphous carbon interlayer, around 300nm thick between the fibres and the matrix. TEM analysis revealed that, at the outermost edge of the fibres, adjacent to the amporphous carbon layer, Mg, Ba and A1 from the matrix had diffused into the fibres, to a depth of less than 250nm. This band in the fibres was amorphous, as were the fibres themselves. The edges of the fibres were smooth. TEM analyses revealed the presence of a small amount of a glassy phase containing Si, O, A1 and Ba, as well as precipitates of arsenic oxide, which is used as a fining agent, and ZrO2, which is used as a nucleating agent.

3.4 Characterisation of the 1000~ composite Examination of the composite which had been exposed to 1000~ revealed major microstructural changes. As a result of high temperature exposure the interface contained a large number of bubbles. Bubbles were not present in the as-received composite. Figure 5 illustrates that the bubbles formed during exposure at 1000~ were of variable, but always small, diameter (usually less than 2~m). The largest bubbles were found on the outer surface of the samples and the population density of the bubbles decreased from the edge of the sample to the centre.

601

Figure 5. (a) Dark field optical micrograph of the outer surface and (b) SEM micrograph of the fracture surface showing bubbles formed along the fibre/matrix interface. The bubbles were associated with a reaction product which resembled honeycomb. Examination of the honeycomb, Figure 6, showed that it had wet

602 the fibres, effectively bonding the fibres to the matrix. Analysis of the honeycomb was conducted using the WDS spectrometers of the electron microprobe analyser. Quantitative analyses, presented in Table 2, of the areas marked in Figure 6, revealed that these areas were enriched in silicon and depleted in aluminium and magnesium when compared with the bulk composition of the matrix. The relative atomic ratios of (Sil_xAlx):O in areas A and B of Figure 6 were 1:2.20 and 1:1.96 respectively, with x ~ 0.2. Thus it

Table 2 Composition of honeycomb i

Element Area A# Area A**

i

Ca 0.50 0.25

As 0.00 0.00

O 51.65 67.32

Ba Si 8.96 33.80 1.36 25.09

AI 7.22 5.58

Mg 0.46 0.39

Area B# 0.49 0.20 43.42 8.69 30.25 8.35 0.57 Area B** 0.25 0.05 55.72 1.30 22.12 6.35 0.48 # = Weight Percent ** = Atom Percent * = by difference

C 0.00" 0.00 8.03* 13.73"

Figure 6. Reaction product formed between the fibres and the matrix. The compositions of the points marked are listed in Table 2.

603 appears that the honeycomb which encapsulated the bubbles resembled silica modified by the addition of aluminium. Quantitative AES analysis of the honeycomb reaction product was unsuccessful due to charging of the surface, which disrupted the spectra collected.

Figure 7. Fracture surface showing the formation of reaction products has a time dependent factor. Near the edge of the sample (a) reaction products coat the fibres, whereas near the centre of the fracture (b) there were areas where the fibres were smooth, although some bubbles had formed.

604 Degradation of the interface was not uniform from fibre to fibre or across the fracture surface of the samples. As can be seen from an examination of Figure 7, degradation was greater near the edges of the fracture surface than near the centre. In the latter areas the fibres remained smooth, with little reaction evident. TEM examination showed that the thin carbon interlayer still remained, but it was thinner than in the as-fabricated composite. In addition, the edges of the Tyranno fibres were irregular, rather than smooth. When viewed optically, using dark field illumination, the fibres were blue, in contrast to their colourless appearance in the as-received composite. Another feature of the interface region was the presence of crystalline precipitates, such as those shown in Figure 8. EDS analysis performed with SEM and WDS analysis using EPMA indicated the presence of high concentrations of arsenic and calcium in the precipitates. In some instances, Figure 9, the calcium-rich precipitates surrounded the fibres in a manner which reflected the concentration of calcium in the as-received composite, Figure 4. Most commonly the precipitates were discretely distributed, rather than being present as a continuous band around the fibres. As can be seen from an examination of Figure 8, the bubbles and the crystalline precipitates were not attached to each other and the precipitates were also present within the matrix. EPMA of the fracture surface, using WDS analysis, was used to quantify the composition of the precipitates. Their composition was variable, however increases in the concentration of calcium were always associated with increases in the concentration of arsenic. An indication of the range of composition is presented in Table 3. As the fracture surfaces had been carbon coated for initial examination in the SEM, no quantitative analysis of the fibres was undertaken. On the precipitates, where high concentrations of carbon were indicated, corroboration of the measured concentrations was obtained by repeating the analyses without carbon and obtaining the carbon concentration by difference, and good correlation was found in most cases.

Table 3 Composition of precipitates (Weight Percent) Element Ca Precipitate 1 8.94 Precipitate 2 13.42 Precipitate 3 8.66 ....Precipitate 4 12.77 * = by difference

".As 22.83 29.80 22.21 29.46

O i 33.03 27.20 37.55 14.65

Ba 10.21 7.57 13.15 7.58

Si 10.31 11.16 5.18 12.26

AI 6.18 4.94 2.47 5.29

Mg 1.26 1.69 0.41 1.14

C 7.25 4.23* 10.37" 16.85"

605

Figure 8. EPMA secondary electron (a) and backscattered electron (b) images, together with x-ray area maps for calcium (c) and arsenic (d) illustrating presence of high concentrations of arsenic and calcium in the precipitates [the bright areas in (a) and (b)]. In (a) and (b) the bubbles appear as the dark areas unattached to the precipitates.

To determine if any diffusion of elements between the fibres and the matrix occurred during testing at 1000~ AES was used to produce line scans for Ca, C, Si, A1 and O, such as Figure 10, across the ends of fractured fibres which had been milled to below the oxidised surface layer. The line scans corroborated the enrichment of calcium near the fibre/matrix interface. They also revealed a

606 layer concentrated in Si, in the form of silicate rather than carbide, between the Ca-rich layer and the fibre. There was some indication of diffusion of silicon out of the fibres, but this observation must be treated with caution as the counting statistics were poor.

Figure 9. X-ray area map for calcium, performed with SEM, showing that the precipitates surrounding the fibres on the exposed fracture surface are calcium-rich.

607

Ca-rich

!

L_ .*..,

.1o L ~4 .m

~e - 3

. . . . Si a s c a r b i d e

..

t

"-5

. . . . Si as s i l i c a t e

o. o~

ffl

*-.6

oooo C i

.,,

i i

i ,

!

",,/,[i""

e e o ~

** ~

Ca-rich

~4,~, ,,,4'

4t

e

***%

Matrix

,"I' \ "!,, ~

....~

.-:....

'

!:', :" :.!

~ io

.!

":

#

it~

Fibre

9 io

i,

!L

t i

Ca

9

~i~

;

""....

:

;~]' ~:,,........ "~

9

.>_

.

9

"'1~ .,:

~2

.....

.

o :

....

e-

.

.a.

:'"

~'~

"I-. r l

-

:

"

9 9

:"

9176 ""

::....: Matrix

.

! 9 9

i

"

.

.|

"

~!'

%

.:.~

I

J

L

!

3

6

9

12

Distance

r']

9

15

i

18

I

21

I

24

(micrometres)

Figure 10. AES line scans through the matrix and across the ends of a fractured fibre on the fracture surface.

4. DISCUSSION In composites many interfaces exist, including those between the phases within the matrix, those between nucleating and fining agents and the matrix and fibres and those between the fibres and the matrix phases. The mechanical behaviour of this particular composite was controlled by the fibre/matrix interface. Examination of this interface showed that a thin, amorphous carbon interlayer formed during fabrication. Another product of processing was the diffusion, to varying distances, of calcium, aluminium, magnesium and barium into the outermost part of the fibres. The presence of calcium in the fibres, as well as in the matrix (Figure 4) was unexpected, although the presence of calcium has been noted in other nominally Ca-free composites [9]. The diffusion of barium, magnesium or aluminium from the matrix into the fibres in Tyranno fibre/celsian and Tyranno fibre/cordierite composites has been observed previously [3,5]. The formation of carbon layers at fibre matrix interfaces has been widely observed [1,3,5,9-12] in SiC and Si-C-O fibre-reinforced silicate matrix systems. It has been reported that SiC fibres containing high concentrations of oxygen

608 produce thicker carbon interlayers than fibres containing little or no carbon [13], and that additions of multivalent oxides, such as the fining agent As203 also cause thicker carbon interlayers to form [ 13]. The mechanism of formation of the carbon layers is still conjectural. Proposed mechanisms include: the diffusion of excess carbon within the fibres to the fibre surface as a result of the affinity of carbon for the residual oxygen in the fibre/matrix interface [ 11]; the supersaturation of the interface by dissolved CO and CO 2 generated during processing; and the oxidation of the fibres by either oxygen or carbon monoxide during processing if the partial pressure of the gas is high [6,13]. The latter mechanism results in the formation of silica as well as carbon, and there was no evidence of the formation of a silica-rich layer between the carbon layer and the fibres in the as-fabricated Tyranno/BMAS composite. The carbon interlayer aided the creation of a composite which exhibited extensive fibre/matrix debonding and fibre pullout (Figures 1 , 2 ) . AES examination of the fracture surfaces produced at ambient temperatures (Figure 2) indicated that the debonding of the fibres occurred within the carbon interlayer, which is to be expected given that the interracial shear strength of carbon is only of the order of 2 MPa [12]. When the composite was exposed to an oxidising atmosphere at IO00~ the bonding between the fibres and the matrix increased to the extent that the composite behaved as a brittle monolithic material, dependent on the strength of the bonded fibres and the matrix, without any benefit of crack blunting or bridging, or of fibre pullout. It has been reported that mechnical bonding between the fibres and the matrix can be affected by the relative coefficients of thermal expansion of the fibres and the matrix [3] as well as by friction due to irregularities on the surface of the fibres [4]. The coefficients of thermal expansion (CTE) of the fibres, the major matrix phases and the composite are similar, and of the order of 2.5 - 4 x 10"6~ -1 [3,5,6,7,8]. The differences in CTE were considered to have had a relatively minimal effect on the properties of the interface. TEM examination showed that, after processing, the surfaces of the Tyranno fibres were smooth, which would minimise friction effects. After exposure at 1000~ the surfaces of the fibres no longer remained as smooth as previously. This, alone, did not cause the reduction in the toughness of the composite. Rather, the interface itself had degraded, resulting in the formation of reaction products which reduced the toughness of the composite. Although AES and TEM analyses showed that some carbon remained in the interface, the original interface was transformed to a network of bubbles (Figures 5-7) encapsulated in a silica-rich honeycomb. In addition, the Ca-rich annulus around the fibres had disappeared. Calcium and arsenic rich precipitates formed both at the fibre/matrix interface and within the matrix. (Figures 8 and 9).

609 It is considered that the degradation of the interface, and consequently embrittlement of the composite was initiated by the diffusion of gaseous oxygen along the fibre/matrix interface. From an examination of Figure 7, it is apparent that diffusion of oxygen, and the associated increase in the activity of oxygen was critical to the degradation of the interface. When the activity of gaseous oxygen was sufficiently high, the carbon interlayer oxidised to form CO according to the reaction 2C + 0 2 --> 2C0

(I)

Oxidation of the carbon interlayer exposed the surface of the fibres to carbon monoxide as well as oxygen. Both oxygen and carbon monoxide can oxidise the fibres by the reactions (2) or (3) respectively: SiC + 0 2 --+ SiO 2 + C

(2)

SiC + CO ~ SiO 2 + 3C

(3)

Both reactions result in the formation of silica and carbon. The carbon formed could either be subsequently re-oxidised if more oxygen was available, or deposited at the interface. TEM analysis showed the presence of a thinner carbon layer after exposure at 1000~ as well as a reduction in the smoothness of the surface of the Tyranno fibres. At the outer surface of the sample, where the oxygen activity was greatest and where there were fewer physical constraints, the bubbles, which were encapsulated by a silica-rich oxide, were larger and more numerous than those formed within the sample. Near the centre of the samples, in the absence of gaseous oxygen, the activity of oxygen in the fibres was not great enough to initiate degradation of the interface. EPMA analysis showed that the Si- and O-rich layer was doped with aluminium. AES line scans (Figure 10) revealed the presence of a silicate rich layer near the fibres as well as a Ca-rich layer which was nearer the matrix. It is proposed that the presence of calcium in the outer 0.5 ~m of the fibres, as a result of processing of the composite, initially acted as a flux for reactions (2) and (3) by reducing the activity of SiO 2. Silicates form as three dimensional networks of Si-O tetrahedra. When calcium is inserted into silicates it initially causes bond rupture of the Si-O tetrahedra, due to its greater ionic character [16], and then it subsequently disrupts the continuity of the Si-O tetrahedra by forming non-bridging bonds with oxygen. The calcium could subsequently undergo an ion exchange reaction with aluminium from the matrix. The addition of trivalent aluminium ions would improve the cohesiveness of the Si-O network, as they would conform to the Si-O tetrahedral structure and, at the same time, they would decrease the number of non-bonding oxygen atoms. The driving force for the ion exchange reaction could have been the reaction of calcium with arsenic oxide and it is apparent that the kinetics of that reaction

510

were rapid. The crystalline structure of the discrete Ca- and As- rich precipitates is currently unknown. As20 3 is a fining agent added to glasses to eliminate bubbles formed during fabrication [17,18]. Arsenic oxide increases the solubility of oxygen [17,18,19] and perhaps CO 2 [17,18] in glass. The mechanism for the former reaction is: As20 3 + 02 ~-~ As20 5

(4)

while the carbon dioxide dissolves to form an interstitial carbonate group [20]. The most commonly found arsenate is Ca3(AsO4) 2, however none of the measured compositions of precipitates (Table 3) approaches this compound. The structure of the precipitates is under investigation by TEM. The reaction products effectively bond the fibres to the matrix, creating a material whose toughness is reduced by the removal of previously available modes of energy dissipation, as well as by reductions in the inherent strength of the fibres and matrix due to exposure at high temperatures. Tyranno fibres are synthesised from polycarbosilane and are amorphous unless heated for extended periods at temperatures in excess of 1200~ [3,5,14,15]. The titanium is reported [3] to be homogeneously distributed through the non-crystalline network, as is that amount of carbon which is in excess of the stoichiometric requirement. After heat treatment in air at 1000~ for times longer than one hour, the tensile strength of the fibres is reduced to around 65% of its original strength [14,15]. The elastic modulus is also degraded [15], but to a smaller degree. The flexural strength of this Tyranno fibre/BMAS composite exposed for comparatively short times (around 60 minutes) in air at 1000~ was reduced to between 42 and 75% that of composites tested at room temperature.

5. SUMMARY

This Tyranno fibre reinforced BMAS composite had a carbon interlayer between the fibres and the matrix after fabrication. The carbon layer ensured extensive fibre/matrix debonding and fibre pullout during mechanical tests at room temperature. After exposure in air at 1000~ the fibre/matrix interface was decorated by bubbles in a viscous silica-rich oxide, discrete Ca- and As-rich precipitates and a thinner carbon interlayer. It is considered that high activity of oxygen at the fibre/matrix interface initiated the formation of CO bubbles and the associated formation of silica. The reactions which occurred had a time dependent factor, related to the rate of diffusion of oxygen along the fibre/matrix interfaces, as the evidence of degradation was not as well developed near the centre of the fractured samples as it was near the edges. The silica-rich reaction products effectively bonded the fibres to the matrix, preventing fibre pullout and energy dissipation of the crack front. As a consequence of the increased fibre/matrix bonding, the composite failed in a brittle manner during four point bend testing at 1000~

611 It is not yet clear if the rate of formation of the silica-rich phase in this composite was increased by the reaction of arsenic with calcium. Further characterisation and environmental testing of this material is planned to try and resolve some of the outstanding questions.

6. R E F E R E N C E S

1. A.G. Evans and D.B. Marshall, "The Mechanical Behavior of Ceramic Matrix Composites" Rockwell International Science Center Report SC 5432.AR (1989) 2. A. Briggs and R.W. Davidge, in Conf. Proc. Whisker and Fiber Toughened Ceramics, Oak Ridge TN USA, 7-9 June 1988, ASM Publ., (1988) 153-163 3. V.S.R. Murthy, Li Jie and M.H. Lewis, Ceram. Eng. Sci. Proc., 10 [7-8] (t989) 938-951 4. P.D. Jero, R. J. Kerans and T.A. Parthasarathy, J. Am. Ceram. Soc., 74 No.ll (1991) 2793-1801 5. M.H. Lewis and V.S.R. Murthy, Composites Sci. & Techn., 42 (1991) 221249 6. K. Prewo, in High Temp./High Perf. Composites, MRS Symp. Proc., Eds. F.D. Lempkey, S.G. Fishman, A.G. Evans and J.R. Strife, 120 (1988) 145156 7. A.I. Kingon and R.F. Davis, in Engineered Materials Handbook, Ceramics and Glasses, Volume 4, ASM Publ., (1991) 758 - 774 8. D. Steel, ARE Holton Heath, UK, private communication 9. S. Mischief, H.E. Bishop and J.J.R. Davies, Surf. & Interface Anal., 18 (1992) 23-31 10. J.J. Bennan, Mat. Sci. & Eng., A126, (1990) 203-223 11. T.W. Coyle, H.M. Chan and U.V. Deshmukh, Interfaces in Polymer, Ceramic & Metal Matrix Composites, Ed. Hatsuo Ishida, Elsevier Science Publ., (1988) 489-501 12. M.D. Thouless, O. Sbaizero, E. Bischoff and E.Y. Luh, in High Temp./High Perf. Composites, MRS Symp. Proc., Eds. F.D. Lempkey, S.G. Fishman, A.G. Evans and J.R. Strife, 120 (1988) 333-339 13. P.M. Benson, K.E. Spear and C.G. Pantano, Ceram. Eng. Sci. Proc., 9 [7-8] (1988) 663-670 14. E. Fitzer, in Conf. Proc. Whisker and Fiber Toughened Ceramics, Oak Ridge TN USA, 7-9 June 1988, ASM Publ., (1988) 9-52 15. J. Lipowitz, Ceram. Bull., 70 No.12 (1991) 1888-1894 16. F.V. Tooley, The Handbook of Glass Manufacture, Ashlee NY Publ., Section 19 Vol II (1984) 999-1019 17. M. Cable and M.A. Haroon, Glass Techn., 11 No.2 (1970) 48-53 18. M. Cable, A.R. Clarke and M.A. Haroon, Glass Techn., 10 No.1 (1969) 15-21 19. S.W. Lee and R.A. Condrate, J. Am. Ceram. Soc., 73 No. 10 (1990) 28732878 20. E.L. Swarts, Ceram. Eng. and Sci. Proc., 7 [3-4] (1986) 390-403

612 7. ACKNOWLEDGEMENTS The author acknowledges many collaborators on this project, including P. Richards and A. Cogdon, who were responsible for material preparation and testing, G. Cumming who performed the AES analyses, A. Gunner who performed the EPMA analyses, J. Russell and R. Muscat who produced the scanning electron micrographs, L. Hammond who produced the XRD patterns and R. O'DonneU who undertook the transmission electron microscopy.

Science of CeramicInterfacesII J. Nowotny(Editor) 9 1994ElsevierScienceB.V. All rightsreserved.

613

Multilayer ceramic capacitors with nickel electrodes Hiroshi Kishi and Nobutatsu Yamaoka Taiyo Yuden Co., Ltd., Takasaki, Gunma, 370, Japan

Abstract Electrical properties and microstructures of holmium-doped (Ba,Ca,Mg)(Ti,Zr)O3 system were studied. Addition of Ho20 3 had little effect in preventing the dielectrics from reducing at high temperature, but the resistivity at low temperature of Ho-doped samples increased with increasing amount of Ho20 3 when they were treated in oxidizing atmosphere at the cooling stage. The newly developed dielectric materials exhibited high dielectric constant 18000 for Y5V specification, 3200 for X7R specification. Using these dielectric materials, Nielectrode multilayer ceramic capacitors with high specific capacitance (Y5V 4.7~F, X7R I~F in EIA 1206 size) were developed. They showed highly reliable electrical characteristics.

1. I N T R O D U ~ O N Multilayer ceramic capacitors (MLCs) play a important part among surface mounting devices (SMDs) using in electronic equipments such as personal computers, VCRs, radios, communication networks and so on. The expansion of SMDs requires advanced MLCs with large capacitance above l~tF for the wider applications, and the subsequent price reduction. At present time, relaxor based MLCs with Ag-Pd electrodes or BaTiO 3 based MLCs with nickel inner electrodes have been studied respectively to realize the miniature size and the reduction of production cost [ 1-11]. This paper deals with MLCs with nickel electrodes. Conventional MLCs based on BaTiO 3 have been fabricated with noble metals such as Pt or Pd as inner electrodes. However these electrodes are very

614 expensive and share a large part of production cost, especially on a MLC with capacitance above I~F. There are certain ways to solve such problems and we attempted to adopt nickel metal for inner and external electrodes. A cost comparison of a few noble metals and base metals is shown in Table 1. As you can see in Table 1, the material cost of a nickel metal is very cheap. It is about 1/500 to the material price of paradium or 1/150 to that of a Ag/Pd(70/30) alloy, respectively. In this case, the dielectric materials must be fired in a low oxygen partial pressure to prevent nickel metal from the oxidization. However, conventional dielectric materials are easily reduced under such a firing condition and become semiconducting. We previously developed high dielectric constant MLCs with nickel electrodes using dielectric materials based on barium titanate and alkaline earth lithiumsilicate glass[10,11]. Ni electrode MLCs have been increasingly produced to replace tantalum electrolytic capacitors, especially in the capacitance range between 0.1~tF and 10~tF. In this paper, the influence of Ho20 3 on electrical properties of (Ba,Ca,Mg)(Ti,Zr)O3-glass component system are studied. Equilibrium conductivity at high temperature and resistivity at the cooling stage of sintering process are measured. We also present the properties of newly developed Ni-electrode MLCs using these materials.

Table 1. Physical properties and price of metals metal

melting t e m p (~

resistivity ~. cm),

sintering atmosphere

price (Yen/g)

Pd

1552

10~11

air

480

Ag

961

1.5

air

15

Ag70/Pd30

1200

15~20

air

155

Ni

1453

7~8

reducing

0.9

Cu

1081

1.6

reducing

0.3

615 2. DIELECTRIC MATERIALS 2.1. C o m p o s i t i o n and powder preparation The composition studied are shown in Table 2. The dielectric materials were prepared by the following procedure. The raw materials for the basic perovskite compounds were BaCO3 , CaCO 3, MgCO3, TiO2 and ZrO 2 9Also, pre-calcined fine grained BaTiO8 was provided as a starting material to obtain X7R dielectrics. They were weighed and mixed by ball milling, and then calcined. Then, Li2O-SiO2-CaO glass component, fine grained Ho2O 3 and small additives were added to calcined perovskite powders, ball milled and then dried. The average particle size of powders were about 0.8 microns. Li20-SiO 2 -CaO glass component is effective sintering aid. So the dielectric material can be fired at about 1200~

Table 2. Compositions of dielectics specification Y5V

X7R

Base materials

(Bao.948Cao.05M~.o05)01.003 9(Tio.s4Zr0.16)O2 (Ba 1.0IMp.01 )O1.02 9(Ti0.98Zr0.02)O2

additives Ho 203 (0~2atomic%) Ho203 (0~2atomic%)

Li20-SiO2-CaOglass (0.5wt%) Li20-SiO2-CaOglass (0.Swt%)

2.2. E q u i l i b r i u m e l e c t r i c a l conductivity The powder, with an appropriate organic binder, was pressed into bars 18ram by 7ram. After burning out the binder, the bars were placed in a stabilized ZrO 2 boat and fired in a tube furnace at about 1200~ in a low oxygen atmosphere controlled by N2 -He -H20 gasses. The equilibrium conductivities were measured by the four-point dc technique using platinum electrodes as a function of oxygen partial pressures P(O2 )=10 15 to 10 5 Pa. Oxygen partial pressures were regulated by CO-CO 2 and Ar-O 2

616 controlled atmosphere, and were monitored by a stabilized zirconia oxygen activity cell placed near the samples. Figure 1 shows the equilibrium electrical conductivity at 1000~ of the 1.6 atomic% Ho-doped and undoped samples of X7R composition compared with pure BaTiO 3. As can be seen in Figure 1, the position of the conductivity minimum of both samples of X7R composition shifted to lower P(O 2) and they were similar between these samples. It is well known that the acceptor impurities such as MgO and MnO are effective in compensating oxygen vacancies caused by a reduction reaction. M a n y studies have been made on h i g h - t e m p e r a t u r e equilibrium electrical conductivity with acceptor- or donor-doped BaTiO 3 [12-16]. Takada et al. showed that the rare earth ions in BaTiO 3 occupy both Ba and Ti sites [17]. The conductivity minimum will move toward lower P(O 2) if holmium acts as a typical acceptor dopant. However, we can easily see from Figure 1 that P(O 2) at the conductivity minimum of each sample of X7R composition is equal. Similar results have been obtained for the Y5V samples. Therefore, it appears that the holmium in this system does not behave merely as an acceptor dopant.

OX7R H020 3 undoped Q X 7 R H 0 2 0 3 1.6atomic% r-1pure B a T i O 3 ~"-o

E r

-1

"q.

o

9 ,%

~I:1%

v

-2

",q. f'~

O

%%

"-.q..

-3 t e m p e r a t u r e = IO00~ -4

-

I

-13

I

!

I

1

I

I

I

I

-11

-9

-7

-5

-3

-1

1

3

Log P(O 2) Figure 1.

5

(Pa)

Effect ofHo20 3 addition on the equilibrium electrical conductivities at 1000~

(X7R composition)

6t7

2.3. Electrical conductivity at the cooling stage Figure 2 shows that the effects of H020 3 content on the conductivity at the cooling stage in two different atmospheres. The resistivity at low temperature decreased with increasing amounts of H0203 under such conditions when the samples were cooled in reducing atmosphere. It is apparent that the resistivity in a reducing atmosphere become hard to recover with doping Ho. However, in the case of the oxidizing atmosphere, the resistivity increased with increasing amount of H020 3. It is considered that the reduction of dielectrics caused low resistivity by addition of H0203 under reducing conditions. On the other hand, it

is expected that the concentration of oxygen vacancies of Ho-doped samples is decreased more easily than that of undoped one under oxidizing atmosphere.

This indicates that understanding of the kinetic oxidation process at the cooling stage is very important in order to improve reliability of Ni-electrode MLCs.

0 ~ '~ I-

Ho20 3 undoped 0.8atomic% -1.6atomic% R)

-

~.

-2|

"

N,, ,

f

-2

~ ~

E

r

E-4

",X

-4 v

'~(~. 1000"C P(O2)= 30Pa 9

-6 cooling rate -350"c/h~

Ho20 3 undoped _ ~ 0.Satomic% -1.6atomic% (X7R) cooling rate -350"C/hr

-6

X~\

-8 I

0.8 i

1.2 I

1000

I

500 200 T(~ (a)reducing atmosphere

Figure 2.

l

I

l

o18

1.6 2.0 103T-1(1/K)

11

,

.6

103T-1 (l/K) .

I

.

I

i

1000

500 200 T(~ (b)oxidizing atmosphere

Influence of holmium addition on the cooling stage in two different atmospheres. (X7R composition)

618 2.4. Dielectric properties The disk samles were sintered at 1200~ for two hours under a reducing atmosphere and cooled under oxidizing atmosphere P(O2)=30Pa. An In-Ga alloy electrode was attached to the both surfaces of the sintered samples. Dielectric constant and dissipation factor were measured by LCR meter at lkHz with 1Vrms. Insulation resistance was measured with a high-resistance meter in one minute after applying 100V dc. Figure 3 shows the resistivity and dielectric constant of Y5V samples sintered various oxygen parti.al pressures. They showed high resistivity (>10 12 Qcm) at regions higher than 10-8pa. Figures 4(a) and 4(b) show the effect of Ho20 3 content on the dielectric characteristics. In the case of Y5V samples, the Curie point shifted to lower temperature by addition of Ho20 3 and yielded high dielectric constant K=18000 at room temperature. Whereas, in the case of X7R samples, the Curie point of BaTiO a (around 120~ remained slightly and temperature dependence of dielectric constant became more stable with increasing amounts of Ho20 3. So, the dielectric temperature characteristics of Ho-doped samples satisfied X7R specifications and dielectric constant at room temperature was about 3200.

Resistivity

20000

1012 o

Dielectric constant 0

16000 108

0 .,,.,

~

~

9 ,,,,4

Y5V

9 ,",d

Ho20.~ 1.0atomic%

,,,,q

12000 104 cooling stage

IO00*C P(02) = 30Pa

8000 i

I

10 "9

10 "s

I

10-7

I

I

10-6

10-5

i

I0-4

P(O2)(Pa)

Figure 3.

Resistivity and dielectric constant of Y5V composition ceramics sintered under various oxygen partial pressures at 1200~

619

I

!

I

I

!

1

I

l

!

!

!

!

!

w

I

I

1

!

I

25000 (a)YSV Ho=O

20000

Ho =0.5

t~

1200"(2 P(Oz) = 10-6pa

0

o

15000 cooling stage

g.t

eo o -9o ~

--

a

10000

5000

s

I

I

-55

I

I

I

I

I

-25

I

I

I

I

I

25

I

I

I

I

!

85

!

125

Temperature(~ 5000

.

.

.

.

.

i

i

i

4000[

..... (b)X7R

i

j

.

.

.

.

.

j

i

I

i

i

J

Ho=0

= .

~

eO

3000

Ho=l 5

"

I,d

o 2000 o

.p,d

1200"(2 P ( O 2 ) = 1 0 - 6 p a

1000

o

cooling stage P(O2)=30Pa

0 -55

-25

25

85

125

Temperature(*C)

F i g u r e 4.

E f f e c t of H o 2 0 3 a d d i t i o n on t e m p r a t u r e c o n s t a n t . (a)YSV, (b)X7R

d e p e n d e n c e of d i e l e c t r i c

620 2.5. M i c r o s t r u c t u a l observation The microstructures of the ceramics were observed by scanning electron microscope (SEM),transmission electron microscopy (TEM) and investigated by energy dispersive X-ray spectroscopy (EDX). Figures 5(a) and 5(b) show the SEM photographs of as-sintered surfaces of Y5V and X7R samples. It can be seen that the size of grains were quite uniform with average grain diameter of about 3 microns for the Y5V dielectrics and about 1 microns for the X7R dielectrics, respectively. TEM photograph of the Ho-doped X7R ceramics is shown in Figure 6. The formation of the grain core-grain shell structure was confirmed [18]. The core phase showed a typical ferroelectric domain pattern. Figure 7 shows EDX spectra of the core and shell phases. It was found that the shell phase was composed of Ho,Zr,Ca and BaTiO3, whereas the core phase was almost equal to that of pure BaTiO 3. It can be regarded that fiat temperature dependence of dielectric constant of the Ho-doped X7R ceramics is originated from core-shell structure.

(a)YSV Figure 5.

(b)X7R SEM photographs of as-sintered surfaces.

621

Figure 6.

A) grain shell

TEM photographs of holmium-doped X7R ceramics.

Ba,Ti

0

Ba,Ti

(,-

0,000

Range=

10.230 key

10,110 4

Figure 7. EDX spectrum of holmium-doped X7R ceramics.

622 3. MULTILAYER CERAMIC CAPACITORS

3.1. Fabrication of MLC Multilayer ceramic capacitors were prepared by a green sheet method shown in Figure 8. Ni paste for inner electrodes was printed on green sheet that were made by a

Ceramic powder] .. ! ] slurry I

doctor blade method. Some green sheets were stacked and pressed into a bar and cut into pieces. Terminal Ni electrodes were

[Doctor blade castin~ [ ! Green sheet .I

[

i

formed on each end of the green chips. After I Electrod~ printin~ (Ni) burning out the binder, the chips were ! Laminating ] sintered at 1200~ for two hours under a ! I reducing atmosphere, except when under Cutting ] 1000~ at the cooling stage. During cooling [ ! to room temperature from 1000~ the I Terminating ](Ni) atmosphere was controlled to P(O2)=30 Pa. I Firing ,](N2/H2/H20) Because nickel was partially oxidized above I 1100~ under oxidizing atmosphere P(O2)=30 Pa, we changed the atmosphere at 1000~ of the cooling stage.

Figure 8. Flow c h a r t of fabrication process of the multilayer chip capacitors with Ni electrodes.

3.2. Electrical properties of MLCs The dimensions and electrical characteristics of newly developed MLCs are summarized in Table 3. Y5V 4.7~F and X7R 1.0~F chip capacitors with nickel electrodes in EIA 1206 size (3.2mm x 1.6mm) were developed. Figures 9(a) and 9 (b) show SEM photographs of thermal etched cross-section of MLCs. The highly accelerated life test (HALT) were performed on Y5V 1.0~tF and X7R 0.1~tF chip capacitors compared with conventional MLCs with Pd electrodes. Figures 10 shows results of HALT under test condition of 165~ The lifetime under HALT of the developed MLCs were equal to those of conventional MLCs. Figure 11 shows the DC breakdown voltage measured at 25~ The breakdown voltage ofY5V 10~F capacitor was around 350V and was much higher than that

623 Table 3. Dimensions and characteristics of the MLCs. SIZE(EIA) Ceramic layer ~ m )

1210 7.1

1206 6.8

0805 11

1206 10.2

0805 12

Capacitance ~ F )

10.0

4.7

1.0

1.0

0.1

10.1 2.8xi0 3

9.4 3.8xi0 3

Y5V

Y5V

Dissipation factor (%)

I.R.(MQ) TCC

4.9 2.8 4.1 xlO 4 4.3xi0 3 Y5V

X7R

1.7 5xlO 4 X7R

Figure 9. SEM photographs of cross section ofMLCs. (a)Y5V (b)X7R

624

9

I

99-

!

Ni pt~

90-

Ni

,~?,6 ~"

.,..~

,~ 5 0 -

Pd

(D r (D

Y5V 1.0

~//~~J/

10

9

m ~

<

X7R 0.1 ~tF

1-

Temperature 165~ Voltage 350V

-

103

I

I

10 4

105

,, 10 6

Time(sec) Figure 10. Results of accelerated life test of Ni-electrode MLCs and Pd-electrode MLCs.

>

1000-

X7R 1.0ptF

9 >

o

X7R 0.1~tF

Y5V Y5V Y5V 1.0taF 4.71xF 10gF 500 -

%) C~ dD

Ta.electrolytic capacitor 12I

Figure 11. DC b r e a k d o w n voltage of the Ni-MLCs c o m p a r e d with t a n t a l u m electrolytic capacitor.

625 of the tantalum elctrolytic capacitor. Figure 12 shows the capacitance change vs. DC applied voltage of Ni-electrode MLCs (Y5V 10~tF and X7R lpF) compared with relaxor based Y5V 10pF capacitor. The capacitance change of Ni-electrode Y5V capacitor was larger than that of relaxor based MLC. It is disadvantage of BaTiO 3 based dielectric materials. The frequency dependence of the impedance was measured for the Y5V 10pF capacitor, aluminum electrolytic capacitor of 47pF and tantalum electrolytic capacitor of 22~F. These results are shown in Figure 13. The impedance of MLC was smaller than those of electrolytic capacitors over the range about 100kHz. It shows the equivalent series resistance (ESR) at resonant frequency is as small as 5raft.

Further more, from various reliable tests such as humidity load life test, thermal shock test and so on, it was confirmed that the newly developed MLCs had highly reliable electrical characteristics.

20

!

!

!

!

w

i

w

i

,

10 3 t

I

c,

-20

10

= o~ r

c,

.~ ~

-40 E -

-60

!

I

-

1. Al.electrolytic 471xF 2. Ta.electrolytic 22g.F " 3. Ni-MLC Y5V 10gF

0 e~

I

~

1

---...

10-1

-80 (Ni)~Y5V 10gF !

A

4

I

I

8

!

J

12

I

|

16

I

20

DC voltage(V) Figure 12. Capacitance change as a function of DC bias voltage for the Ni-MLCs and relaxor MLC.

10-3 1

1

I

I

10

10 2

10 3

I

10 4

10 5

Frequency (kHz) Figure 13. Frequency dependence of impedance for the Ni-MLC, alminum electrolytic capacitor and tantalum electrolytic capaditor.

626 4. CONCLUSION It was found that the addition of Ho203 had little effect in preventing the dielectrics from being reduced at high temperature, but the resistivities at low temperature of Ho-doped samples increased with increasing Ho20 3 when they were treated in oxidizing atmosphere at the cooling stage. Understanding of the dynamic oxidation process at the cooling stage is very important to improve reliability of Ni-electrode MLCs. With control of the firing process, Ni-electrode MLCs with high specific capacitance were developed. They showed highly reliable electrical characteristics.

REFERENCES

1. M.Yonezawa, K. Utsumi and T.Ohno, Proc. 1st. Meeting on Ferroelectric Materials and their Applications, (1977) 297. 2. K.Furukawa, S.Fujiwara and T.Ogasawara, Proc. of the Japan-US Study seminar on Dielectrics and Piezoelectric Ceramics, (1982) T-4. 3. J.Kato, Y. Yokotani, M.Nishida, S.Kawashima and H.Ouchi, Jpn. J. Appl. Phys., 24 suppl.24-2 (1985) 90. 4. Y. Yamashita, O.Furukawa, H.Kanai, M.Imai and M.Harata, Ceramic Transactions ,eds. H.C.Ling and M.EYan, Am. Ceram. Soc., 8 (1990) 35. 5. K.Tsuzuku and M.Fujimoto, Proc. of IMC 1992, (1992) 588. 6. J.Kato, Y.Yokotani, H.Kagata and H.Niwa, Jpn. J. Appl. Phys., 26 Suppl.26-2 (1987) 90. 7. I.Burn and G.H.Maher, J. Mater. Sci., 10 (1975) 633. 8. Y.Sakabe, I~Minai and I~Wakino, Jpn. J. Appl. Phys., 20 Suppl20-4 (1981) 147. 9. N.Fujikawa, T.Shintome, H.Utaki and N.Yokoe, Proc. of IMC 1986, (1986) 202. 10. H.Kishi, S.Murai, H.Chazono, M.Oshio and N.Yamaoka, Jpn. J. Appl. Phys., 26 Suppl.26-2 (1987) 31. 11. H.Saito, H.Chazono, H.Kishi and N.Yamaoka, Jpn. J. Appl. Phys.,30 No.9B (1991) 2307. 12. S.A.Long and R.N.Blumenthal, J. Am. Ceram. Soc., 54 (1971) 515. 13./~Seuter, Philips Res. Repts., Suppl. 3 (1974)

627 14. J.Daniels and I~H.H~dtl, Philips Res. Repts., 31 (1976) 489. 15. J.Daniels, Philips Res. Repts.,31 (1976) 505. 16. N.G.Eror and D.M.Smyth, J. Solid State Chem., 24 (1978) 235. 17. I~Takada, E.Chang and D.M.Smyth, Advances in Ceramics, eds. J.B.Blum and W.R.Cannon, Am. Ceram. Soc., 19 (1985) 147. 18. D.Hennings and G.Rosenstein, J. Am. Ceram. Soc., 67 (1984) 249.

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Science of Ceramic Interfaces II J. Nowomy (Editor) 9 1994 Elsevier Science B.V. All rights reserved.

629

Corrosion Properties of Ion plated TiN and CrN Coatings Hiroshi Ichimura and Atsuo Kawana Sumitomo Metal Mining Co. Ltd., Central Research Laboratory 18-5, 3 Chome, Nakakokubun Ichikawa, Chiba, 272, Japan Abstract High temperature oxidation of TiN and CrN films was studied in air at temperature from 923 to 1173 K. Oxidation rates were calculated by the mass gain with reaction time. Formed oxide layers were analyzed by XRD, SEM, SAM, and EPMA. The oxidation of TiN films is controlled by the oxygen diffusion through the formed TiO2 layer. The oxidation of CrN, however, is controlled by the outward diffusion of Cr ion through Cr203 layer formed along the grain boundaries. Oxidation resistance of CrN was superior than that of TiN. The pin-hole densities of TiN and CrN films were evaluated as a function of film thickness by means of Ferroxyl test. The observed pin-hole densities of TiN and CrN were decreased with increasing film thickness. The sufficient film thickness for TiN coatings with pin-hole zero was estimated as more than 15/~m for TiN. The pin-hole densities of CrN films were remarkably less than that of TiN films. Corrosion resistance in 10 % HCI solution was measured for the TiN and CrN coatings deposited under the various conditions. CrN coatings showed the high corrosion resistance. I. Introduction

The metallic nitride films such as titanium nitride have high hardness, high strengths, and high wear resistance. Therefore, these films have been applied extensively as the coatings on cutting tools, die molds and so on. The performance of coated parts depends on not only mechanical properties such as hardness and friction coefficient but also corrosion resistance of coated film. TiN coatings deposited by physical vapor deposition (PVD) have been successfully applied in industrial usage for wear resistance, but they suffer from a limited corrosion resistance. For nitride coatings, there are many reports with respect to mechanical properties such as hardness, adhesion, and wear resistance [1-4]. However, there are few reports on the corrosion properties of nitride films. TiN starts to oxidize already at temperatures as low as 823 K [5]. MiJnster and Schlamp [5]

630 concluded that the kinetics of the high temperature oxidation reaction of TiN films deposited by chemical vapor deposition (CVD) follow a linear-parabolic law, and the initial interface reaction control yields rapidly to a diffusioncontrolled mechanism. Tampieri et al. [6] reported that the oxidation of hotpressed TiN ceramics indicates diffusion controlled kinetics where the diffusion of oxygen or the desorption of nitrogen through the oxide scale would be the rate governing step. Kawana and Ichimura [7, 8] reported that the high temperature oxidation resistance of CrN films is superior than that of TiN. Aqueous corrosion behaviors of coatings depend on not only chemical stability of film itself but also the properties such as pin-hole density, film thickness, adhesion between deposited film and substrate, and the substrate materials. Kado et al. [9] reported that the observed pin-hole densities of TiN films that deposited by plasma CVD method were decreased with increasing film thickness. They found that the sufficient film thickness for pin-hole zero coatings was about 15/~m. We have measured the pin-hole densities of ion-plated TiN and CrN films as a function of film thickness by means of Ferroxyl test, and studied the corrosion resistance in 10% HCI solution for TiN and CrN coatings that prepared under the various deposition conditions, This paper presents the summary of corrosion properties of nitride coatings and to discuss the mechanism of high temperature oxidation and aqueous corrosion.

2. Experimental details TiN and CrN were deposited onto austenite stainless steel in a cathodic arc ion plating apparatus as illustrated schematically in Fig.1. The metal ions generated by vacuum arc from cathode react with nitrogen, and deposited onto the substrate as a nitride film. TiN and CrN films were prepared by using a Ti metal cathode and a Cr one, respectively. |

_A

JNeu,

Plasma ~teriatm~ [

~-q....._6~.

Ftux 7 ARC- # ,

~~eactive

Hateriat '

~--J--~Va~u um

~ s

IF-

Figure 1. Schematic diagram of an arc ion plating apparatus.

631 The substrates having 15 mm square and 5 mm thickness in size were mechanically polished to an average surface roughness less than 0.05/~m. Then the substrates were cleaned by ultrasonically in an ethanol bath, and subjected to freon drying. These substrates were mounted in the vacuum chamber, and pumped to a pressure of 10 -5 Torr. The substrates were treated by metal ion bombardment to be heat and to remove the surface oxide layer. Nitrogen was introduced to maintain a chamber pressure at 30 mTorr for TiN and 50 mTorr for CrN during deposition. TiN and CrN were deposited at the negati.ve bias voltage of 300 V and the substrate temperature at 600 K. The film thickness for oxidation kinetic experiments were kept about 5/~m by adjusting the coating time. Oxidation test was conducted as follows; the specimens were kept in a platinum basket and set in the electric furnace with air flow of 500 cm3/min. The samples were heated ranging from 923 to 1173 K for 0.6 to 60 ks. After each oxidation period, samples were taken from the furnace and measured the weight change. The chemical characteristics of the oxidized films were determined by XRD. A Scanning Electron Microscope (SEM) was used to observe the morphology of films. The compositional depth profiles for specimens before and after oxidation were measured by Electron Micro Prove Analysis (EPMA) and Scanning Auger Microprobe (SAM). The pin-hole densities of coated films were measured by means of Ferroxyl Test as shown in Fig.2, which authorized by JIS-H8663. In this test, the pin-holes of films were revealed as blue spots that were formed by the reaction between the ferroxyl reagent and Fe element in the substrate materials. Corrosion behavior of coatings was investigated by HCI immersion test. The specimens were immersed into 10% HCI solution. After each immersion period, samples were taken from the solution and measured the weight change.

carbon plate

/

filter paper / with Ferroxyl solution test piece

' '"

J

I

/

/

..... ~~

'-----

8cm

carbon p l a t e ~

S uz'face/

'-

voltmeter recorder

0"2A/dm2

"

m

,,

9

Figure 2. Schematic diagram of Ferroxyl test.

632

3. Results and discussion 3.1 High temperature oxidation of TiN Ichimura and Kawana [10] reported the mechanism of high temperature oxidation of ion-plated TiN films. Figure 3 presents the XRD patterns of TiN films before oxidation (a), and after oxidation (b) at 973 K for 15 ks, respectively. TiN and TiO2(rutile) have been revealed by X-ray diffraction performed on the gurface of the oxidized samples. Thermodynamically [11], we should expect the oxide scale on titanium nitride to consist of a sequence Of layers of different oxides of titanium (TiO Ti203, Ti305, TiO2), but in this experiment, only rutile phase has been detectecl as a titanium oxide.

_

A C',4 C"4

z

v--

~

O

I----

r-......

o

z

.,t-.

t'v" ~

O v A

O O

O

r

.c2

I-.-.

2O

40

60 80 20/deg.

100

120

Figure 3. X-ray diffraction patterns of TiN films before and after the oxidation test. (a): Before oxidation, (b): After oxidation at 973 K for 15 ks.

633

Figure 4 shows the SEM micrograph of the cross sections of samples that oxidized at 1023 K for 18 ks. We can see the clear interface between formed product (TiO2) and the non reacted TiN layer.

Figure 4. SEM micrographs of the fracture cross sections of TiN films before and after the oxidation test. (a): Before oxidation, (b): After oxidation at 1023 K for 18 ks. The chemical reaction between titanium nitride and oxygen gives the following equation. TiN + 02 = TiO2 + 1/2 N2

(1)

We can calculate the reaction rate ( ot ) from the mass gain using the equation (1). In the case of specimens in Fig.4 ( b ), the ot was calculated as 45 % that reflecting the thickness ratio on oxide layer to total film thickness. This oxidation behavior was confirmed by EPMA compositional depth profiles of the oxidized samples with a thickness of 20/~m that oxidized at 1123 K for 4 ks. Comparing depth profiles before (Fig.5a) and after (Fig.5b) oxidation, the outer layer of the oxidized sample contained much oxygen corresponding to the TiO2 formation. The reaction from TiN to TiOz occurs with large volume expansion. This generates cracks at the oxide-nitride interface and consequently some spalling occurs. In the present experiments the oxide layer detached from TiN layer when the oxidation exceeded 50 % of the thickness. Figure 6 presents the parabolic plots of the mass gain as a function of reaction time at each oxidation temperature. The slopes of these lines give the parabolic rate constant, kp. From the results of Figs 4, 5, and 6, it suggests that the oxidation of TiN films is controlled by a diffusion process through the formed rutile layer. The diffusion of oxygen from the atmosphere to the reaction interface or the desorption of nitrogen from the reaction interface to the atmosphere is proposed as the primary diffusion mechanism controlling the

634

oxidation.

rll f

4.--

(19) Ti N Oxidized /+5%

TiN

ta)

z

~b-

o

20

3'0.

o

DISPLACEMENT, pm

I'o 20 30 DISPLACEMENT,pm

Figure 5. EPMA compositional depth profiles of the cross section of TiN films before and after the oxidation test. (a): Before oxidation, (b): After oxidation at 1123 K for 4.8 ks. The activation energy derived from the temperature dependence of the parabolic rate constants in Fig.6 was calculated as 136 kJ/mol. Tampieriet et al. [6] reported that the activation energy for the oxidation of hot-pressed titanium nitride ceramics in the oxidation temperatures ranging from 973 to 1373 K is 190 kJ/mol. 700

600 c~E500 '0 400

t'

1023 K 973 K

923 K

""300 ~l:::: 200 100

00

I

10

I

20

.....

I

30

t/ks

I

40

I

50

l

60

70

Figure 6. Parabolic plots of mass gain as a function of time for the oxidation of TiN films.

635

According to the study of the oxidation reaction on CVD TiN films the activation energy was found as 195 kJ/mol [5]. Many authors have studied on the high temperature oxidation of Ti metal. Kofstad [12] summarized that the oxidation kinetics of Ti metal in the temperature range from 873 to 1173 K is parabolic, and the activation energies of oxidation are 210-230 kJ/mol. Kofstad has concluded that the oxygen diffusion predominates the reaction below about 1173 K, while titanium ion diffusion is also important at a higher temperature. The TiN films used in our oxidation tests have a columnar microstructure with fine grains compared with CVD or hot-pressed TiN. In addition to microstructure, the deposition temperature was lower than oxidation temperature, so that the grain growth of TiN and formed oxide occurs during the oxidation test. The obtained lower activation energy in present study compared with that of monolithic TiN ceramics and CVD TiN films suggests that the oxidation of ion-plated TiN films is controlled by the oxygen or nitrogen diffusion that is enhanced by the grain boundaries along the formed oxide layer. In order to elucidate the grain boundary effect on oxidation, we carried out the similar oxidation test using TiN film with higher columnar structure (Fig.7) preparing it at zero substrate bias voltage. The parabolic rate constant of oxidation for this film was two or three times larger than that of standard types of TiN films in Figs.4 and 6. The activation energy (122 kJ/mol) for the oxidation of this film was slightly lower than that of the standard one. Chaze and Coddet [13] have studied on the role of nitrogen in the oxidation in air for Ti-A1 alloys, and proposed that the diffusion coefficient of nitrogen in rutile is higher than that of oxygen. From the present study and the previous reports, we conclude that the oxidation of ion plated TiN films are controlled by oxygen ion diffusion, which is enhanced by the grain boundaries of the formed rutile layer.

Figure 7. SEM micrograph of the fracture cross section of TiN film having a higher columnar structure. (a) Before oxidation, (b) After oxidation at 923 K for 15 ks.

636

3.2 High temperature oxidation of CrN Kawana and Ichimura [7, 8] have investigated the high temperature oxidation behavior of CrN films. They showed that high temperature oxidation of CrN films was also controlled by diffusion process as shown in Fig.8. The activation energy derived from the temperature dependence of the parabolic constants in Fig.8 was calculated as 225 kJ/mol. XRD analysis indicated that oxidized CrN films changed to Cr203 as shown in Fig.9. 200 ~ .,.. 150

I

1123K

I

/

~100

/

...................................................................................... ~..................

...........................................................................................................................

~'o

k

~'<: 5o

/ [ /

1048 K

/

./.,

HE

00

A

10

i

20

30 t/ks

I

40

5'0

6'0

70

Figure 8. Parabolic plots of mass gain as a function of time for the oxidation of CrN films.

,.-.. ,,o

(b)

Z' L. 1.3

~

i i

~

!

"~ f'4

_

20

40

....60

80

2O/deg.

~

"

1

~

=

100

120

20

40

60 80 2e/deg.

Figure 9. X-ray diffraction patterns of CrN films before and after the oxidation test. (a): Before oxidation, (b)" After oxidation at 1123 K for 10 ks.

100

120

637

The chemical reaction between chromium nitride and oxygen give the following equation 2 CrN + 3/2 0 2 - CmO3 + N2

-I-. t--

4-C

o ...i.

(a)

d~

t=.

tO>_"

CP

F-I==..l tz)

Z Ill F-Z

,.I::3 CL rO

CF

(2)

(c).

>--

Fe

F-i-i-i t/) Z Ill

F-Z

I==iI

0

, i

20 4O DISPLACEMENT, l]m

,4--. i.--,

0

2O 40 DISPLACEMENT,IJm

(b)

Er

I..

>:

l--

Z LI..I l--Z

I---I

_

0 i

20 40 DISPLACEMENT,l]m

i,i

Figure 10. SAM depth profiles of CrN films before and after oxidation test. (a) before, (b) 1123 K for 0.6 ks ( c~ =3%), (c) 1123 K for 480 ks (a: =47%). Oxygen, nitrogen, and chromium ion distributions in oxidized specimens were measured by Scanning Auger Microprobe (SAM). There is no oxygen ion in before oxidation sample as shown in Fig.10 (a). Figurel0 (b) shows that

638 the compositional depth profiles of the specimen that oxidized at 1123 K for 0.6 ks. It indicated that nitrogen concentration is almost same as that of before oxidation one. Although the oxidation rate ( ot ) of this sample was calculated as about 3 % from the equation (2), the oxygen penetrated to the interface between CrN film and the substrate. Figurel0 (c) presents the depth profiles of Cr, N and O in the specimen oxidized at 1123 K for 480 ks. The ot of this sample was calculated as 47 %. Nitrogen decreased drastically compared with that of specimens in Figs 10 (a) and 10 (b). These results suggest that the oxidation of CrN proceeds by the rapid oxygen diffusion along the grain boundaries. Kofstad and Lillerude summarized on the high temperature oxidation of Cr metal [14]. They concluded that the oxidation kinetics of Cr metal above 973 K are generally interpreted as parabolic and the reaction is controlled by the Cr ion diffusion in formed Cr203. The activation energies of self diffusion coefficient of chromium and oxygen ions in Cr203 were reported as 255 kJ/mol [14] and 424 kJ/mol[15], respectively. An activation energy obtained in this study is slightly lower than the value of Cr diffusion in Cr203, so that the oxidation of CrN in hot air is controlled by Cr ion diffusion through the Cr203 of reaction products. From the activation energy and the SAM depth profiles, it is concluded that the oxidation of CrN is controlled by the outward diffusion of Cr ion through Cr203 layer formed along the grain boundaries, The temperature dependence of the parabolic rate constants of oxidation for TiN and CrN films derived from the mass gain as a function of time is illustrated in Fig.11.

I

Temperature ( C ) 1000 900 800 700

10 -4

,r,~ 10-s

i '

..................

.....@ ~

i

i

9

~9m e t a l

'

600

i

i

.....................................................

. ..... . ..'.....-...~ .~ ................................................................. ......_ . __

"r'E 0_6 ....................

........

...~10 -7

.....

10 8

~ - . _ ...................

I

8

I

I

9

I,

...................

I

10

COK

I

I

11

Figure 11. Parabolic rate constants of oxidation reaction of ion plated TiN and CrN (this study), CVD TiN [5], Ti metal [12], and Cr metal [14].

639

The oxidation of CrN was two orders of magnitude smaller than that of TiN. Figure 11 also presents the rate constants of oxidation for CVD TiN [5], Ti metal [12], and Cr metal from the literatures [5]. The rate constants of TiN and CrN were almost same to that of Ti and Cr metal, respectively.

3.3 Aqueous corrosion behavior The corrosion resistance of coatings depends on not only chemical stability of deposited films but also the morphology of films, adhesion, and chemical stability of substrate materials. Particularly the pin-hole density of deposited films is very important for aqueous corrosion. In generally, the pin-hole densities of coatings decrease with increasing film thickness. Figure 12 presents the pin-hole densities as a function of film thickness for TiN and CrN coated SUS 304 steel specimens. The thinner TiN films have much pin-holes, so that the TiN films without pin-hole need to produce thicker coatings more than 15 Fm. This result was almost same with that of plasma CVD TiN reported by Kado et al. [9]. They also reported that anodic polarization current density of TiN/SUS304 at 0.33 V in 1 kmol/m 3 HCI solution decreases logarithmically with film thickness. The pin-hole densities of CrN coatings were remarkably less than that of TiN coatings as shown in Fig.12.

E

20

/ ~ ~

Ti N, CrN/SUS 304 Ferroxyl.test 9 oTiN

>.

"~ 10 o , - - .

n

5

10 15 Thickness , p rn

20

Figure 12. Pin-hole densities of TiN and CrN coatings, which measured by Ferroxyl test, as a function of film thickness. Figure 13 presents the weight change as a function of immerses time in 10 % HCI solution for TiN and CrN coatings having 5 Fm thickness that deposited onto SUS 304 stainless steels. The weight loss of specimens increased with

640

logarithmical time dependence after the incubation periods that depend on the coated films. Non coated SUS 304 steel specimen corroded severely compared with coated samples. In these corrosion experiments, coated films didn't almost dissolved into HCI solution. Therefore, observed weight loss is mainly due to the dissolution of substrate materials. The results of Figs.12 and 13 suggest that the dissolution rate of coatings in aqueous solution is determined by the pin-hole density.

03

E E

o:

CrN/ SUS304

(l) 03-5 c" t~ cO -10 (.-03 -15 -20 10

TiN/SUS304 I

I'

1

I

20

30

50

100

1

I

200 300

I

500

1,000

Immertsion time, t/hr Figure 13. Weight change as a function of immerses time in 10 % HCI solution for TiN and CrN coatings having 5 Hm thickness. In order to improve the aqueous corrosion resistance of TiN coating, Takizawa et al. [17] prepared the multi layer coated SUS 304 by ion plating, and evaluated the corrosion behavior in 5 % H 2 5 0 4 . They found that multi layer coating consisted of TiN/TiCN/Ti/SUS has higher corrosion resistance due to low pin-hole density. Meltecs et al. [18] investigated the corrosion wear properties of ion plated TiN films on M50 beating steel in 1N-H2SO4 at 298 K using potentio-dynamic polarization techniques. They reported that TiN coatings show essentially no corrosion action under rubbing loads and the current densities during testing are 3 to 4 orders of magnitude less than M50. Ichimura et al.[19] reported the high corrosion resistance of CrN in high temperature and high pressure water 9They compared the corrosion resistance of TiN and CrN coatings by using autoclave at 563 K and 80 kgf/cm 2 in water. After 1000 hours testing, the weight of TiN coating increased by 0.12 mg/cm 2 and very thin oxide film was detected by XRD. However it was no change for CrN coatings. Taguchi and Takahashi [20] investigated the corrosion behavior of chromium nitride films,

641

which were deposited by if-ion plating, in 1 kmol/m 3 H 2 5 0 4 solution at 373 K. They found that the Cr2N in the mixture phase consisted of CrzN and CrN was selectively corroded and the amount of dissolved chromium was about two orders of magnitude larger than that for the single phase of CrN.

ED

E oJ E

I00

121-5 c-" t13 cO -10 c'O3 (D -15

50 V "

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

300 -20

10

20

30

50

100

200 300

Immertsion time, t/hr

500

1,000

Figure 14. Effect of substrate bias voltage during deposition on corrosion resistance to 10 % HCI solution for TiN. The corrosion resistance also depends on microstructure of deposited films. It is well known that the morphology of ion-plated films is varied with substrate bias voltages during deposition [4, 21]. At low bias voltage the columnar structure is predominant, and it becomes more dense structure with increasing bias voltage. In order to study an effect of film morphology on corrosion resistance, we evaluated the corrosion resistance of TiN coatings that deposited at various bias voltages. As shown in Fig.14, the corrosion resistance of TiN coatings to HC1 solution increased with increasing bias voltage. Weight loss for low bias coating below 100 V was observed without incubation time. Corrosion resistance of coatings having much pin-holes such as thinner TiN coatings depends on the chemical stability of substrate material. Figure 15 shows the weight change of TiN coatings with 5 Fm thickness that deposited onto different two substrate materials; SUS 304 (Cr: 18, Ni:8) and more stable 310 (Cr:24, Ni:19) show in Fig.15. Figure 15 indicates that the corrosion resistance of TIN/310 is higher than that of TIN/304. Corrosion mechanism of coatings in aqueous solution is summarized as illustrated in Fig.16. The corrosive liquid such as HCI permeate through pin-holes or grain boundaries of deposited films, and attain to the interface

642

between the film and substrate ( I ). When corrosion at the interface proceeds ( II ), the substrate material is corroded and finally, coated films detach from the substrate ( III ).

~

.

E o

.

.

.

!

........

1 ~ - 5

...............................................

tt~ r--

0-10 _r 133

,. . . .

...................................................................................................................... .

~-15

e

.

.............................................................................................................

-20J

t-

10

,

,, ,

~

20

30

50

i

,

100

,__

,TiN/SUS3_, 10 I

200 3 0 0

500

1,000

Immertsion time, t/hr

Figure 15. Effect of substrate material of TiN coating on corrosion resistance to 10 % HC1 solution.

/

Corrosive

.!iquid

cotumner strctucture

\

~ ~ ~ ~ ~ "~ - ~ - -~ i __-=e;~

pin-hole

(1)

dissolved substrctte

-

z;

,I

grain boundaries detached film

-

(II)

(li~)

Figure 16. Corrosion mechanism of coatings in aqueous solution.

643

4. Conclusion

The high temperature oxidation and aqueous corrosion of TiN and CrN coatings that were prepared by cathodic arc-ion plating lead to the following conclusions: (1) The oxidation rates at temperatures ranging from 923 to 1173 K in air are found to fit well to a parabolic time dependence. (2) The rate determining step of the oxidation of TiN is the grain boundary enhanced diffusion of oxygen in the formed rutile layer. (3) The oxidation of CrN films is controlled by the outward diffusion of Cr ion through Cr203 layer formed along the grain boundaries. (4) The high temperature oxidation rates of TiN and CrN were same orders of magnitude to that of Ti and Cr metal, respectively. (5) The pin-hole densities of CrN films, which measured by Ferroxyl Test, were remarkably less than that of TiN films. (6) Aqueous corrosion resistance of coating having pin-holes not only depend on film properties such as pin-hole density, chemical stability and microstructure but also corrosion properties of substrate materials. (7) CrN coatings showed the high corrosion resistance in HCI solution. 5. References

1. 2. 3. 4. 5. 6. 7.

P.J.Burnett and D.S.Reckerby, Thin Solid Films 157, (1988) 233. E.Moll and E.Bergmann, Surface Coating Technol. 37, (1989) 483. S.Yamamoto and H.Ichimura, J.Mater.Res. 7, (1992) 2240. Y.Chiba, T.Omura and H.Ichimura, J.Mater.Res. in print. A.M~inster and G.Schlamp, Z.Physik.Chem. (Frankfurt) 13, (1957) 76. A.Tampieri, E.Landi and A.Bellosi, Br.Ceram.Trans.J. 90, (1992) 171. A.Kawana and H.Ichimura, High temperature corrosion of advanced materials and protective coatings, edited by Y.Saito, B.Onay and T.Maruyama, p267 (1992), Elsevier Science Publishers. 8. A.Kawana and H.Ichimura, Shigen to Sozai (in Japanese), 108, (1992) 868. 9. T.Kado, R.Makabe, S.Mochizuki, S.Nakazima and M.Araki, Bousyoku Gizyutsu (in Japanese), 36, (1987) 551. 10. H.Ichimura and A.Kawana, Submitted for publised in J.Mater.Res. 11. Phase diagram for Cemmists 1975 Supplement, E.M.Levin and H.F.McMurdiesds, The American Ceramic Society, 1975, p 135. 12. P.Kofstad, High-Temperature Oxidation of Metals, John Wiley & Sons Inc. New York, 1966, p 169. 13. A.M.Chaze and C.Coddet, Less Common Met. 124, (1986) 73. 14. P.Kofstad and K.Lillerude, J.Electro.Chem.Soc., 127, (1980) 2379. 15. W.Hagel and A.U.Seybolt, J.Electrochem.Soc., 108, (1961) 1146. 16. R.Lindner and A.Akeratrom, Z.Physk.Chem.(Frankfurt) N.F., 6, (1956) 162.

644

17. K.Takizawa, M.Fukushima and H.Imai, Hyoumen Gizyutsu (in Japanese), 42 (1991) 1255. 18. E.I.Meltics, A.Erdemir and R.F.Hochman, Ion plating and implantation, edited by R.F.Hochman, American Soc. for Metals, p173 (1986). 19. H.Ichimura, A.Kawana and Y.Chiba, Surface Modification of Large-size Parts and Members by Dry Process, edited by Japan Reseach and Develop Center for Metal, p75 ( 1990 ). 20. M.Taguchi and H.Takahashi, Nippon Kinzoku Gakkai-shi (in Japaneese), 56, (1992) 1221. 21. D.Wang and T.Oki, Thin Solid Films, 185, (1990) 219.

Science of Ceramic Interfaces II J. Nowotny (Editor) 9 1994 Elsevier Science B.V. All rights reserved.

645

DIRECT ENERGY CONVERSION IN SINGLE GAS MIXTURE BY AN OXIDE SOLID ELECI~OLYTE ELECTROCHEMICAL CELL AND ITS APPLICATIONS IN MULTI-GAS SENSING Da Yu Wang GTE Labs. Inc., Waltham, Massachusetts 02254 (*Now at GM, NAO R&D, Warren, M148090-9055) ABSTRACT The principle of direct energy conversion was evaluated and the conditions to create spontaneous emf on electrodes which were exposed to the same gas mixture were discussed. The demonstrations were carried out in CH4-O2-N2 gas mixture and N20-O2-N2 gas mixture. The oxide electrochemical sensor was made of yttria stabilized zirconia and had a thick-filmmulti-layer structure. Each sensor had two platinum electrodes which were unsymmetrical; one electrode was encompassed by the electrolyte membrane which had a small aperture open to the ambient atmosphere and the other electrode was directly exposed to the ambient atmosphere. Two ceramic heaters was used in each sensor to maintain the operation temperature. Immersed in CH4-O2-N2 and N20-O2-N2 gas mixtures, the electrochemical cell would produce spontaneous electromotive forces which are believed to be due to the gas oxidation reaction, CH4 + 202 = CO2 + 2H20, and the gas decomposition reaction, N20 = N2 + 1/2 02. The amplitude of the spontaneous emf of the nitrous oxide decomposition reaction was found to be proportional to the concentrations of N20 if the temperature of the sensor and the oxygen concentration of the gas mixture were kept constant. INTRODUCTION Utilizing the principle of direct energy conversion, solid state oxide electrolytes, such as fully and partially stabilized zirconia, have developed many interesting applications which include high-temperature fuel cells and various gas sensors.I, 2 In these applications, the cathode and anode are exposed to separate gases, and the difference of the free energy between the two gases is directly converted into electromotive force which is measured at the two electrodes. In electrochemistry, what happens at the electrodes is that the gaseous reactants chemisorb on the surface of the electrodes, and electronic charges and oxide ions are exchanged at the electrode-electrolyte interface. At high temperature, the electrode and the oxide-electrolyte are easily in equilibrium with the gases and separate gases are needed at the anode and cathode in order to produce the spontaneous emf effect. 1-6 At low temperatures, the catalytic effects of the metal electrode and the oxide-electrolyte are weak and it is not easy to form equilibrium between the gas, the electrode, and the oxide electrolyte.7-11 In such occasion, it is possible to have a direct energy conversion with a pair of unsymmetric electrodes exposed to the same gas-mixture without the involvement of a reference gas. 12 The electrode-gas reactions at the electrodes will be self modulated by the electrical fields of the spontaneous energy conversion and maintain a continuous spontaneous emf. 12 There is no restriction on the types of gas mixtures involved, as long as the gas mixture provides a negative free energy change and unsymmetrical electrodes are used for the anode and cathode.

646 In this paper, we will review this unique electrochemical phenomenon by using two different types of gas mixtures: CH4-O2-N2 and N20-O2-N2. THEORY In electrochemistry, an oxidation gas reaction such as methane-oxygen reaction1 CH4 + 202 = CO2 + 2H20,

(1)

would give a reversible potential Eth defined as Eth = -AG/8e,

(2)

where AG is the Gibb's free energy for the reaction (1). The corresponding half electrochemical reaction at cathode would be CH4 + 4 0 - -> CO2 + 2H20 + 8e,

(3)

and at anode, the other half reaction would be 202 + 8e-> 40-.

(4)

Since AG is negative for reaction (1), the direct conversion of the chemical energy to the electromotive force, Eth, is spontaneous. The Eq.(2) indicates that the two electrodes involved have to be unsymmetric in polarity. Traditionally, this is satisfied by supplying oxygen to the anode and methane to the cathode. If one has the two electrodes expose to the same CH4-O2-N2 gas, will a spontaneous emf to be observed? Since both electrodes are symmetric to each other here, one might think there should be no spontaneous emf. However, considering from thermodynamics point of view, one will recognizes that the driving force of Eq. (2) on reaction (1) is still there. If locally there is no electrode or electrolyte catalytic effects to compete with, the direct energy conversion for reaction (1) should prevail. However, the requirement of two unsymmetric electrodes is there. So, we have to make the two electrodes electrochemical-unsymmetric. With both electrodes exposing to the same ambient, this can be achieved by using different shapes or different materials for the two electrodes. 12-25 Once we have the spontaneous electrochemical reaction (1) on going, the measured open-cell potential can be less than what is given in Eq. [2]. This is because some of the available energy will be consumed by the electrode polarization and the IR drop in the electrolyte: emf+ 11 + IR =-AG/8e,

where R is the electrolyte resistance, 11 is the electrode overpotentials, and I is the current determined by the load resistance. Usually, the electrode overpotentials r I follow the Butler-Volmer equation:

(5)

647 I = Io [exp (etXarl/kT) -exp (-eOtcrl/kT)],

(6)

in which Io, t~a, Otc, e, and k are the exchange charge current, the anodic and cathodic chargetransfer coefficients, the electron charge unit, and the Boltzmann constant, respectively. 3,5-7 The values of t~a and O~cdepend on the type of reactions involved. Their values usually are 1.0 for the anodic and cathodic reaction of (4) if no slow surface-diffusion step is involved, and between 0.5 and 1 for the cathode and between 1.0 and 1.5 for the anode when the lateral -surface-diffusion of the chemisorbed oxygen becomes the rate control step. 3 Besides the oxidation type of gas reactions as shown in reaction (1), there is second type of gas reactions which involve oxygen reduction. A typical example is the reduction reaction of nitrous oxide, N20 = N2 + 0.5 02,

(7)

which would give an electrochemical potential, Eth = -AG/2e,

(8)

in which AG is the Gibb's free energy for the reaction (7). 1 The half electrochemical reaction at cathode would be O- -> 0.5 02 + 2e,

(9)

and the other half reaction at anode would be N20 + 2e -> N2 + O-.

(10)

The AG of reaction (7) is negative and the electrochemical reaction should be spontaneous. Following the same argument as in previous section, there should be a natural tendency for the nitrous oxide to directly convert its energy to emf even both electrodes are exposed to the same oxygen-nitrous-oxide gas mixture, as long as the electrodes have certain electrochemical unsymmetry built in. However, the situation for N20 reduction reaction is more complex. This is because the reduced oxygen can form oxygen molecular and released back to the ambient gas easily (it is difficult to maintain a high oxygen activity on metal surface). Under this situation, Eq. (8) is not suitable. One of the simplest modifications of Eq. (8) is obtained by assuming the difference of the two electrode electrochemical activity is controlled by N20 and 02 gas concentrations; therefore, the spontaneous emf E is defined as, E = E0 [CN20 ] (CN20 + AoCo2)],

(11)

in which, E0 is a constant, CN20 and CO2 are the gas concentrations of N20 and 02, A0 is an unsymmetric factor of the two electrodes and is a function of the differences of the gas species, electrode materials, and the electrode layout between the two electrodes (the degree of unsymmetry of the electrodes). For a pair of symmetric electrodes, one expects to have A0 ~>0 and obtains a zero spontaneous emf. Rearranging Eq. (11), one would have, (E0-E)/E = (AoCo2/CN20) = (A/CN20).

(12)

648

This equation suggests that if CO2 is kept constant, the plotting of Log[(E0-E)/E] against Log(CN20) would give a slope o f - 1 . 0 and the constant A0 could be calculated from the intercept. Because of the relatively small value of emf involved in Eq. (12), the effect of thermoelectric power (Seebeck effect) on the spontaneous emf measurement may become important. If we assume the electrodes are exposed to the same gas (Po2) and there is a temperature difference of AT between the two electrodes, the thermal electric power should be, emf = (-k/4e) [(Ln PO2)(AT).

(13)

in which, k and e are the Boltzmann constant and the electron unit.13,14

EXPERIMENTAL The electrochemical cells studied were made of yttria-stabilized zirconia. Solvent and binder were added to YSZ powder and the slurry was made into thick films. Platinum electrodes (Engelhard non-fritted platinum ink) were screen-printed on some of the films. Several films with and without Pt electrodes were thermally laminated together under pressure and made into the electrochemical cells.7,12 There were three types of electrochemical cells studied here. The trtrst type had symmetric electrode structure. The two Pt electrodes were encompassed separately by two electrolyte membranes which had individual aperture open to the ambient gas. The second type had a simple structure. It contained two electrodes with one having a gas chamber formed next to it (called the inner electrode) and had a small aperture open to the ambient gas. The other electrode was exposed to the ambient gas directly (called outer electrode). The third type of sensor was composed of three parallel gas chambers and three pairs of Pt electrodes. The two outer chambers had apertures open to the ambient gas and each chamber had one Pt electrode sit inside the chamber and another Pt electrode sit on the outside wall of the chamber. These two pairs of Pt electrodes were used for oxygen pumping, to control oxygen concentration inside the chambers. The inner chamber had two apertures open to the two outer chambers and had a pair of Pt electrodes, one was on the inside wall of the inner chamber and the other was on one side of the outside walls of the inner chamber (therefore this electrode was included in one of the outer chambers). The laminated samples were sintered at 1500~ for one hour. The sample density was close to 99% of the theoretical value. The electrode area was between 0.14 cm 2 and 0.16 cm 2. The thickness of the electrolyte between the electrodes was about 600 ktm. The aperture had a diameter of about 50 ktm and a length of 0.25 cm. The housing of the sample was made of stainless steel. To heat up the solid electrolyte cell, a pair of ceramic heaters were attached to the two sides of the sensor. A LFE microprocessor controller together with a Pt/Pt+10% Rh thermocouple were used to control the sensor temperature. The gas-control system was composed of several MKS gas flow meters, valves, and tanks of CH4, air, N20, N2, 02. The blended gas was introduced to a stainless steel testing stand which had the electrochemical cell mounted at the top and the testing stand which had a gas volume of 40 cm 3. The gas flow rate was controlled between 2,000 sccm to 10,000 sccm. The gas composition was determined from the readings of the gas flow meters.

649 Keithley 197 multimeters were used to monitor the voltage and the current of the electrochemical cells. For oxygen pumps, a HP dc power supply was used. An illustration of the experimental setup could be seen in Fig. 1 of reference 12. Detailed discussion of the electrochemical cell and its basic IV performance in nitrogen-air mixture could be found also in references 7 and 12. R E S U L T S AND DISCUSSIONS

Oxidation Types of Gas Dopants-To begin with, we used the first type of electrochemical cells mentioned in the experimental section to examine the criteria of strong and weak gas-metal and gas-electrolyte interactions for the oxide electrolyte system (which included Pt electrodes) studied here. These electrochemical cells had symmetric Pt electrodes encompassed by separate electrolyte membrane with separate aperture open to the ambient gas mixture. First we applied dc voltages to the electrochemical cell and studied the current as the function of the applied voltage, temperature, and gas mixtures. Since different fuel-air gas mixtures produced similar results, only results of the methane-air gas mixture are presented here. For gas mixtures with a methane concentration lower than the stoichiometric number (lean methane-air gas mixture), at low-applied voltages, a limiting current developed at the cathode of the cell, which was confirmed by destroying the cathode aperture and eliminating the limiting-current phenomenon. This low-voltage limiting current was found to be linearly proportional to the residual oxygen concentration in the gas mixture based on reaction (1). When we increased the applied voltages further, the cell would increase its current again until a second limiting current was developed. This high-voltage limiting current was found to be linearly proportional to the oxygen concentration of the gas mixture without considering the methane oxidation effect of Eq. (1).7,12 Typical examples are shown in Fig. 1, in which the dark circles are data for air electrodes and the open circles were data for air mixed with 6.1% methane. Both experiments were conducted at 610~ As shown in Fig. 1, the air electrodes gave one limiting current only. This is because the gas contained no methane and the current was limited by the cathode aperture at low and high voltages. At high enough applied voltage, the dark-circle data show the current would rise again; this is due to the electrolysis effect of the electrolyte which gave off an electronic current to the existing ionic limiting current (see the dark circles in Fig. 1).15 The open-circle data in Fig. 1 show two limiting currents, as we discussed. At low voltage, the methane and oxygen reacted at the cathode (see Eq. (1)) and the limiting current was lower than that of the air-electrodes (compare the open and the dark circles in Fig. 1). At high applied voltage, the high voltage stopped the reaction (1) (with V > Eth in Eq. (2)) and a second limiting current plateau was observed with the level slightly lower than that of pure air data because of the physical-dilution effect of the methane on oxygen in the gas mixture. No spontaneous emf was observed on this electrochemical cell at 610~ as the temperature was deemed too high for this to happen. When we lowered the cell operation temperature, similar phenomenon was observed, until the temperature was lowered below 450~ At this temperature, we started to observe the modulation effect of gas adsorption by the I-V experiments. A typical example is given in Fig. 2, in which the electrochemical cell was operated at 450~ and air and 6.1% methane doped air were used as the ambient gases. As clearly shown in Fig. 2, the I-V data of air-electrodes (dark circles) still behaved similarly as the corresponding data shown in Fig. 1, but the I-V data of 6.1% methane (open circles) show different results from that in Fig. 1.

650

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610~ o 6.1% Methane in Air 10-2 . . . . i . . . . i . . . . i . . . . ! . . . . , . . . . 0 500 1000 1500 2000 2500 3000 V-IR (mV)

Figure 1. The plots of I-V data of an electrochemical cell which was exposed to air (dark circles) and methane doped air (open circles). The IR contribution (electrolyte polarization effect) was determined by the ac complex impedance method.

10 ~ A A A ~

o.

9 9 9 9 9 9Oo..,,~s,q,q,-~ 0

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1000 1500 2000 V - IR (mV)

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Figure 2. The plots of I-V data of an electrochemical cell operated in two different gases: air (dark circles) and 6.1% methane doped air (open circles and open triangles). By comparing the open-circle data between Fig. 1 and Fig. 2, it becomes clear that the opencircle data in Fig. 2 had an initial I-V performance follow that of air-electrodes. Only after the expose of high enough voltage, the current would drop to the "normal" level and show limiting current behavior ("normal" in the sense that the reaction (1) became effective now). Afterwards, the I-V data showed "normal" behavior with two limiting currents; one corresponded to the effect of the reaction (1) and the other to the physical-dilution effect of the methane on oxygen in air. This results indicates that before the applied of the high dc voltages, the electrodes did not respond to methane in the gas mixture and oxygen was the main chemisorbed species at the surface of the electrode. After the application of dc voltages, the

651 electrode chemisorption behavior could be changed. This demonstrates for the first time the possibility of co-existence of oxygen-chemisorbed electrode and methane-chemisorbted electrode for the same methane-air gas atmosphere. The solid lines in Fig. 2 were the fitted results of Eq. (6) for the initial part of data before reaching the limiting current range. The parameters used in the fitting were Io = 3.0 I.tA, CXa= 3/2, and oct = 1/2. The value of Io was found not sensitive to the gas compositions. As we increased the methane concentrations for the I-V experiments beyond the stoichiometric point, we encountered the first evidence of direct energy conversion having the electrodes exposed to the same ambient gas mixture. The results are presented in Fig. 3, in which the methane concentrations were 8.3% (closed circles) and 9.8% (open circles) and the operation temperature was 540~

~ I

.o..-~s.s"88

0.4 03 ~9

0.2

000 000

0000 i

0.1

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500

1000 1500 2000 V- IR (mV)

2500

3000

Figure 3. The plots of I-V data obtained with two different gas mixtures" 8.3% methane doped air (closed circles) and 9.8% methane doped air (open circles). The two experiments were done at 540~ As shown in Fig. 3, we still had two limiting currents observed for each gas mixture; one for the reaction (1) and the other for the physical dilution effect of the methane on the oxygen. But the low-voltage limiting current occurred at the anode (not at the cathode as shown in Figs. 1 and 2) and was linearly proportional to the residual methane concentration based on the reaction (1). This could be doubly checked by the elimination of the limiting current after destroying the anodic aperture. What happened at the anode is that when the methane concentration of the methane-air gas mixture was higher than the stoichiometric number the reaction (1) switched from cathode to anode with the oxygen pumped from cathode side to the anode side to maintain a stoichiometric reaction of Eq. (1) and the current was controlled by the limited supply (diffusion) of methane through the anodic aperture. As the methane concentration was increased further, more oxygen would be needed at the anode and the supply of oxygen from the cathode could be limited by the cathode aperture. Then, the low-voltage limiting current became linearly proportional to the physically-methane-diluted oxygen concentration of the gas mixture. To illustrate this phenomenon, we plot the low-voltage limiting current data against the methane concentration in Fig. 4. In Fig. 4, the first group of data represents the low-voltage limiting currents having the methane concentration lower than the stoichiometric number and the limiting current was controlled by the cathode aperture and was linearly proportional to the residual oxygen concentration based on the reaction (1). The second group of data represents the low-voltage

652 limiting current with the methane concentration beyond the stoichiometric number. The limiting current was limited by the anode aperture and was linearly proportional to the residual methane concentration based on reaction (1). The third group of the data represents the low-voltage limiting current which was controlled by the cathode aperture with the limiting current linearly proportional to the physical-dilute oxygen. The high-voltage limiting currents shown in Fig. 3 had the same current-limiting origin but were generated in different experimental condition. Their plotting in Fig. 4 (see the open and close circles Fig. 4) demonstrates they and the third group of limiting-current data shared the same physical origin of the current limiting behavior (see the dotted line in Fig. 4). The first group of data (originated from the cathode aperture) and 400 !l-- ......... 9 "0 . . . . .

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. . . .

,

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10 15 20 25 Methane Concentration (%)

30

Figure 4. The plots of low-voltage limiting currents against the methane concentration in air. The open and enclosed circles represent the high-voltage limiting current data obtained from Fig. 3.

1.2 1.0

8.3% Methane Doped Air 540~

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Figure 5. The plots of overpotentials versus the currents drawn from an electrochemical device which had the two electrodes exposed to a gas mixture of 8.3% methane in air.

653 the second group of data (originated from the anodic aperture) shown in Fig. 4 had almost the same slopes but with opposite signs; this indicates the degree of symmetry of this two electrodes. Back to the I-V data shown in Fig. 3, what differs Fig. 3 form Fig. 2 and Fig. 1 is that when the applied voltages were decreased again from high values, the current did not follow the original low-voltage curve but became negative values (see the arrows in Fig. 3, which indicate the sequences of the data taking). The solid lines in Fig. 3 are the fitted results of Eq. (6) for the I-V data preceding the low-voltage limiting currents and the high-voltage limiting currents. The fitted parameters were Io = 0.02 mA, Ota = 3/2, and OCc--- 1/2. The data preceding the highvoltage limiting currents was fitted by the same parameters, with the base shifted by 940 mV. Negative current shown in Fig. 3means the electrochemical cell had developed a spontaneous electromotive force which had a value higher than the applied dc voltage, while the two electrodes were exposed to the same gas mixtures. This was confirmed by taking away the applied voltage and we observed an open-cell potential of 980 mV for the electrochemical cell we studied. This emf would last as long as the gas mixture was supplied. By coupling a resistor in series, electric currents could be drawn from the electrochemical cell. A typical example is shown in Fig. 5, in which the electrode potentials are plotted against the drawn currents. The electrochemical cell was operated at the temperature of 540~ and the methane concentration was 8.3%. It is worthwhile to point out that the observation of limiting currents in a rich methane-air gas mixture as shown in Fig. 3 indicated that the electrolyte was not in equilibrium with the electrode and gas; otherwise, the electrolyte should have been reduced by the methane-rich gas and no ionic current should have been observed here. When the methane-air gas mixture was kept constant, there was a temperature range within which the spontaneous emf would not disappear. At too low a temperature, the electrode and electrolyte polarization would be too big for the electrochemical energy to overcome (see Eq. (5)); at too high a temperature, the local equilibrium between the gas-electrode-electrolyte would reduce the solid electrolyte and the cell would be self shorted. A typical example is shown in Fig. 6, in which the methane concentration was kept at 12.6% and the active

1.20

1.00 & 9 ~'

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0.80

ra~

12.6% Methane in Air 0.60

'1

100

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300 500 Temperature (~

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Figure 6. The plots of non-zero open-cell potentials versus the operation temperatures. The solid line represents Eth based on Eq. (2).

654 temperature range was determined to be between 240~ and 540~ The top solid line in Fig. 6 represents the theoretical Eth values based on Eq. (2) for reaction (1). The deviation between the theoretical values and the actual data is explained by the Eq. (5) as the effects of electrolyte and electrode polarization. In Fig. 7, we present the effective ranges of methane concentration and temperatures within which non-zero spontaneous overpotential could be measured (see open circles). As seen in Fig. 7, the higher the methane concentration, the lower the effective-high-temperature point is; a reflection of the methane-reduction effect on the solid-electrolyte. Outside the boundary, no direct energy conversion was possible for the electrochemical cell studied here.

100 80

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Effective Range

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Figure 7. The plots of temperature and methane concentration ranges within which the nonzero open-ceU emf could be measured. The open circles represent data of the same cell which generated the data shown in Figs. 1-6. The dark circles represent data from an electrochemical cell which had unsymmetric electrodes built in. The dc voltages needed to start the emf effect (see Fig. 3) were temperature dependent; the lower the temperature, the higher the voltage required. At 335~ the dc voltage needed could be as high as 7000 mV. But once the spontaneous emf was on, the temperature and the methane concentration could be changed without destroying the emf, as long as they stayed within the effective ranges shown in Fig. 7. We believe the requirement of a dc voltage to start the spontaneous emf effects which are shown in Fig. 3-7 are due to the fact that the two electrodes of the electrochemical cell were quite symmetric in shape and the dc voltage was needed to help to establish the polarity of the spontaneous emf at the beginning. Once the polarity was built, the emf can maintain constantly by itself. When we used the second type of electrochemical cells (see Experimental section) which had two electrodes quite unsymmetric, we found the emf effect was indeed spontaneous and there was no need to use dc voltage to initiate the emf effect for the electrodes. The effective ranges of temperatures and methane concentrations were plotted in Fig. 7 (see dark circles). The interesting point is that as shown in Fig. 7, the dark circles present a larger effective ranges, which suggests that the effective range would be a function of the unsymmetry of the electrodes. The electrochemical cell which provided the "true" spontaneous emf results in Fig. 7 (the dark circles) always had the cathode for the "inner" electrode which had encompassed by an

655 electrolyte membrane with an aperture open to the ambient gas. This result may relate to the fact chemisorption of oxygen is easier than that of methane on Pt electrode (see Fig. (2)) and an open-structure electrode was better suited for the anodic reaction of Eq. (4). Thus far, in order to observe the direct energy conversion in the same ambient gas mixture, the methane concentration had to be larger than the stoichiometric value. However, this limitation could be overcome by adding oxygen pumps to the electrochemical cell. With the pumps pumping the oxygen out of the electrochemical cell, the spontaneous energy conversion could proceed within the electrochemical cell with the ambient methane concentration much less than the stoichiometric value (using the third type of electrochemical cells discussed in Experimental section). Fig. 8 represents a result of such study, the applied voltages on the oxygen pumps were 1800 mV and the open-cell overpotentials of the central pair of electrodes are plotted against the ambient methane concentration in air without counting the oxygen pumping effect. Two temperatures were experimented: 470~ (open circles) and 610~ (closed circles). The data shown in Fig. 8 are slightly deviated from what we expect; the edges of the data-curves are not straight. Another interesting point shown in Fig. 8 is that the operation temperature 610~ is beyond the high-temperature limit shown in Fig. 7 for the electrochemical cell without the oxygen pumps. The fact that at higher temperature, the cell detected a lower methane-concentration (100 ppm in air at 610~ versus 1000 ppm at 470~ can be explained by the effectiveness of oxygen pumping at higher temperature.

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Figure 8. The open-cell overpotential data plotted against the methane concentration in air with two oxygen pumps pumped at 1800 mV within an electrochemical cell studied here. Other than methane, we had experimented with ethane, propane, hydrogen and carbon monoxide as the dopants in air. Similar results as those shown in Figs. 1 to 8 were observed except the spontaneous emf level, the effective temperature ranges and the effective doping levels were slightly different from that of methane-air gas mixtures.

Reduction Type of Gas Dopants-We first tested the thermoelectric power effect on the electrochemical cell having a pair of unsymmetrical electrodes. This electrochemical cell was similar to the electrochemical cell which generated the dark circles data shown in Fig. 7 (the second type introduced in the Experimental section). It had one electrode directly exposed to the ambient gas and the other electrode encompassed within an electrolyte membrane which had an aperture open to the ambient gas.

656 The open-ceU potential of an electrochemical cell were measured in a series of O2-N2 gas mixtures. Since the gas mixture was nitrogen and oxygen only, we attribute any observation of emf to the Seebeck effect (see Eq. (13). The results are presented in Fig. 9, in which the spontaneous open-cell potentials are plotted against the logarithm of the 02 concentration. The electrochemical cell was operated at 680~ As shown in Fig. 9, a linear slope was obtained (see the solid line in Fig. 9); for each decade change of oxygen concentration, the open-cell potential changed by 1 mV less. Using Eq. (13), the temperature difference between the two Pt electrodes was estimated to be 17~ with the inner electrode had a temperature lower than that of the outer electrode. Based on this temperature difference, we could estimate the thermal conductivity of yttria stabilized zirconia of this electrochemical cell to be around 0.022 W/~ cm, which is comparable to the published result. 15 Based on the result of Fig. 9, the effect of thermoelectric power on the emf data was deemed too small for further consideration.

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657 Figure 10. The plots of open-cell-potential data versus temperature of an electrochemical cell operated in two different gases: 100% N20 (open circles) and 13% 02-40% N20-47% N2 gas mixture (closed circles). When the same cell was immersed in 100% N20 gas, spontaneous open-cell potentials with the values larger than 10 mV were observed. The spontaneous emf was found to be a function of the operation temperature also. Typical data are presented in Fig. 10 with the open-cell potentials plotted against the temperatures (see the open circles in Fig. 10). As shown in Fig. 10, there is deviation between the theoretical values and the actual values. Again the deviation could be explained as the effects of electrode and electrolyte polarization (see Eq. (5)). Compare Fig. 10 with Fig. 6, the open-cell potentials generated by nitrous oxide have lower emf values and a wider temperature ranges than that of methane-air gas mixtures. As shown in Fig. 10, similar results were obtained with the electrochemical cell immersed in 13% 02-40% N20-47% N2 gas mixture, except the emf levels were smaller (see the dark circles in Fig. 10). When the N20 concentrations in the gas mixture were changed, the values of the open-cell emf were found to be proportional to the N20 gas concentration but not in a linear way. Typical results are shown in Fig. 11, in which the open-cell emf data are plotted against the logarithm of CN20. In Fig. 11, the oxygen concentrations of the gas mixtures, CO2, were controlled at four different levels (see Fig. 11, the closed squares for 0.021% 02, the closed circles for 0.21% 02, the open circles for 1.0% 02, and the open triangles for 16% 02). We assume, the open-cell potential of E0 at 680~ is around 45.4 mV. Using this number and the data shown in Fig. 11 substitute into Eq. (12), we plot Log [(E0 - E)/E] against Log (CN20) in Fig. 12. As shown in Fig. 12, the plots give a slope of-1.0 for most of the data (see the solid lines in Fig. 12) but the data for CO2 = 16% (see the open triangles in Fig. 12). The values of the intercepts obtained from Fig. 12 (which we call A and should correspond to AoCo2 based on Eq. (12)) are listed in Table 1 and are plotted against CO2 in Fig. 13.

50

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Oxv~en Level 0.021% 9 0.21 % 9 1.0 % o 16 % A -

40 30

~ 2o 0 ~

10

~

0

-10 ....... ' 10 -3 1 0

2 .................

0 ....... ' . . . . . . . . '2"

10 1 10 101 Nitrous Oxide (%)

10

v

"'"" 10 3

Figure 11. The plots of open-cell potential versus N20 concentration for four different gas mixtures with four different 02 levels.

658 10 3

\

~m_ 10 2

680~

X ~

~" 10

o

10 -1 ........ , ........ , . . . . . . . . . . . . . . . . . . . . . . 10 -3 10 -2 10 -1 10 0 101 10 2

10 3

Nitrous Oxide (%) Figure 12. The plots of Log [(E0 -E)/E] versus Log (CN20). The symbols correspond to those shown in Fig. 11 Table 1 Constant A obtained from Fil~. 12 C02 (%)

2.1xlO -2

2.1x10 -1

1.0xl00

1.6xlO 1

Constant A (%)

3.2x10-1

2.0x100

1.3x101

3.4x102

10 3

680~

10 2 10 1 10 o 10 -1 10 -2

10 -1

10 0 Oxygen (%)

101

10 2

Figure 13. The plots of A listed in Table 1 against CO2. The solid line in this figure would give A0 = 15. As shown in Fig. 13, the constant A was linear proportional to CO2 and could be represented as A = A0 CO2 with A0 = 15. The solid curves shown in Fig. 11 are the fitting results of Eq. (11) with A0 = 15.

659 The result of A0 = 15 means the electrodes were 15 times more effective in interacting with the plain oxygen than with the nitrous oxide. We anticipate that the value of A0 could be altered by the use of different electrode materials as well as the use of different electrode layout. A good device should have A0 as small as possible. The electrochemical cells which provided the data shown in Figs. (10)-(14) had always the "inner" electrode more positive than the "outer" electrode. Again, this may reflect the fact that an open-structure electrode was less effective in reducing N20. In Fig. 14, the spontaneous emf data are plotted against the time while the concentrations of N20 were changed between 100 ppm and 900 ppm with an increment of 100 ppm. The electrochemical cell was operated at 680~ and the oxygen level was kept at 0.021%. Based on Fig. 14, the cell had a response time less than 20 sec.

10.0

~ 900 ppm / O p~m800 PPm _600 pp~ 500 ppm /,Z 400 ppm / . ~ 300 ppm

8.0 '-' 6.0 0

Background Oxygen is 0.021%

680oC

4.0

I'4

0 = o ra~

[,,Z 200 ppm ,a 100 ppm

2.0

'

0.0

'

'

I

I

500

0

,

9

~

m ....

,

,

I

I

1000 Time (sec)

I

I

'

I

,

1500

,

9

,

2000

Figure 14. The plots of spontaneous emf versus time of a typical electrochemical cell with the N20 concentration changed in an increment of 100 ppm. Built-in oxygen pumps were used to increase the sensitivity of the electrochemical cell to the small amount of N20 in high background concentrations of oxygen. Typical results are given in Table 2. The device was operated at 810~ and the gas mixture had 3.5% of oxygen in it with the N20 gas changed between 0.012% and 0.12%, and between 0.12% and 1.2 %. As shown in Table 2, the oxygen pumps improved the sensitivity of the device. Also we plot the typical spontaneous emf data versus time in Fig. 15. As shown in Fig. 15, the response time of the sensor was less than 20 seconds. Table 2 N20 gas sensing results of a device with the two built-in oxygen pump ; operated at 810~ Pumping Voltage Pumping Current Emf Differe ncc-(1) Emf Difference(2) (mV) (mA) (mA) (mV) 0.16 200 2.20 0.86 0.27 400 2.31 1.82 0.60 600 2.34 3.48 0.64 800 2.37 3.81" (1) the change of N20 was between 0.12% and 1.2% (2) the change of N20 was between 0.012% and 0.12% (*) its time response result is presented in Fig. 15 ,

,,

660

6.0[. ~'lnot-, '

4.0

9

[

lu1.1. K,/

"

"

'

"

"

'

~1.,,

(Background Oxygen is 3.5%)

"

"

Nitrous Oxide level changes between 0.12% and 1.2%

2.0

r~

0.0 0

30

60 Time (sec)

90

120

Figure 15. The time response curve of an electrochemical cell which had two built-in oxygen pumps pumped at 800 mV. The slow response time observed in Figs. 14 and 15 should not be taken as the final values for the electrochemical cells studied here. Because the principle of the spontaneous emf observed here is identical to that of traditional oxygen sensors, there should be enough room for further improvement.16 Finally, we like to point out that it is no strange to see the spontaneous emf effect on the electrochemical cells which have the electrodes exposed to the same gas mixture as we have seen in Figs. (3)-(8), and (10)-(15). It has long been known that the half reactions listed in Eqs. (3), (4), (8), and (9) could be made to happen by applying the electric powder from outside of the system as we had demonstrated in Figs. (1)-(3). 12, 17-19 What we had done differently here was to use the electrochemical energy provided by the system itself to force the reactions to happen and to operate the electrochemical cell as low temperatures so that there is no local electrode catalytic effect to compete with. Also the requirement of using unsymmetric electrodes can be extended to symmetric-shape of electrodes but built of different materials which have quite different catalytic effects. 21-25

SUMMARY We have demonstrated that in a single gas mixture, a direct conversion of energy is possible. The conditions to generate this phenomenon are (1) the electrochemical reaction has to have a favorable energy change (AG < o), (2) local electrode, and electrolyte catalytic effects are not stronger than the direct energy conversion effect, and (3) the electrodes have to be electrochemical-unsymmetric. This electrochemical unsymmetry can be achieved by using different materials or the shapes for the electrodes. ACKNOWLEDGMENTS The author thanks B. Ditchek, J. Gufstson, J. Proud, and J. Cote for their supports of this work.

661 REFERENCES Physics of Electrolytes, edited by J. Hladik, Academic Press, New York, 1972. Werner Weppner, "Surface Modification of Solid Electrolytes for Gas Sensors", Solid State Ionics 40/41,369-374, 1990. Da Yu Wang, "Electrode Reactions at the Surface of Oxide Ionic Conductors," Solid State Ionics 40/41,849-56, 1990. B.C.H. Steele, J.A. Kilner, and A.E. McHale, "Oxidation of Methane in Solid State Electrochemical Reactions," Solid State Ionics 18 & 19, 1038 (1986). T.H. Etsell and S.N. Flengas, "Overpotential Behavior of Stabilized Zirconia Solid Electrolyte Fuel Cell," J. Electrochem. Soc. 118, 1890-1900, 1971. M.V. Perfilyev, "Electrode Reactions in Solid Electrolytes," Solid State Ionics 9 & 10, 765-76, 1983. D. Y. Wang, "Low-Temperature Diffusion-Controlled Polarization of Pt Electrodes with Yttria-Stabilized Zirconia Electrolyte," J. Am. Electrochem. Soc., Vol. 137(11), 3660-6, 1990. D.S. Tannhauser, J.A. Kilner and B.C.H. Steele, "The Determination of the Oxygen Self-Diffusion and Gas-Solid Exchange Coefficients for Stabilized Zirconia by SIMS," Nuclear Instru. and Methods in Phys. Res. 218, 504-8, 1983 B.C.H. Steele, J.A. Kilner, P.F. Dennis and A.E. McHale, "Oxygen Surface Exchange and Diffusion in Fast Ionic Conductors," Solid State Ionics 18 & 19, 1038-44, 1986. 10

C.G. Vayenas, S. Bebelis, and S. Ladas, "Dependence of Catalytic Rates on Catalyst Work Function," Nature, Vol. 343,625-7, 1990.

11

S. Pancharatnam, R.A. Huggins, and D.M. Mason, "Catalytic Decomposition of Nitric Oxide on Zirconia by Electrolytic Removal of Oxygen," J. Electrochem. Soc. 122, 86975, 1975.

12

Da Yu Wang, "Direct Energy Conversion in Single Gas Mixture by an Oxide Solid Electrolyte Electrochemical Cell," Ceramic Transactions, Vol. 24, 387-96, 1991.

13

N.M. Tullan and I. Bransky, "Observation of Mixed Thermoelectric Power in ThO2," J. Electrochem. Soc., 118, 2, 345-49, 1970.

14

R.J. Ruka, J.E. Bauerle, and L. Dykstra, "Seebeck Coefficient of a (ZrO2)0.85(CaO)0.15 Electrolyte Thermocell," J. Electrochem. Soc., 115, 5, 497-501, 1968.

15

D. Y. Wang and A.S. Nowick, "Polarization Phenomena Associated with Reduction of a Doped Ceria Electrolyte," J. Electrochem. Soc. 127, 113-22, 1980.

16

W.D. Kingery, J. Francl, R.L. Coble, and T. Vasilos, "Thermal Conductivity of Zirconia, Zircon, and Spinel," J. Am. Ceram. Soc. 37, 109 1954.

662 17

C.T Young, "Experimental Analysis of ZrO2 Oxygen Sensor Transient Switching Behavior," SAE Paper 810380, Sensors SP.486, SAE, 29-41, 1981.

18

C. G. Vayenas, S. Bebelis, and S. Ladas, "Dependence of Catalytic Rates on Catalyst Work Function," Nature, Vol. 343,625-7, 1990.

19

S. Pancharatnam, R.A. Huggins, and D.M. Mason, "Catalytic Decomposition of Nitric Oxide on Zirconia by Electrolytic Removal of Oxygen," J. Electrochem. Soc. 122, 869-75, 1975.

20

C.G. Vayenas, "Catalytic and Electrocatalytic Reactions in Solid Oxide Fuel Cells," Solid State Ionics, 28-30, 1521-39, 1988.

21

N. Li, T. C. Tan, and H. C. Zeng, "High-Temperature CO Potentiometric Sensor," J. Electrochem. Soc., Vol. 140, No. 4, 1068-1072, 1993

22

N. Rao, C.M. van den Bleek and J. Schoonman, "Potentiometric NOx (x = 1, 2) Sensors with Ag+-b ''- Alumina as Solid Electrolyte and Ag Metal as Solid Reference," Solid State Ionics 52, 339-346, 1992.

23

G. Baler, V. Schule, and A. Vogel, "Non-Nemstian Zirconia Sensor for Combustion Control," Appl. Phys., A 57, 51-56, 1993.

24

E. Hafele, K. Kaltenmaier, and U. Schonauer, "Measurement of Ammonia with the Solidox-NH3 System," Sensors and Actuators B, 4, 529-531, 1991.

25

E. Hafele, K. Kaltenmaier, and U. Schonauer, "Application of the ZrO2 Sensor in Determination of Pollutant Gases," Sensors and Actuators B, 4, 525-527, 1991.

Science of Ceramic Interfaces 1I J. Nowotny (Editor) 9 1994 Elsevier Science B.V. All rights reserved.

663

TRENDS OF RESEARCH ON CERAMIC INTERFACES J. Nowotny Australian Nuclear Science and Technology Organisation, Program of Advanced Materials, Lucas Heights Research Laboratories, Menai 2234, Australia ABSTRACT The science of ceramic interfaces overlaps several existing disciplines such as solid state chemistry and physics, solid state electrochemistry, surface science, metallurgy and catalysis. Papers representing these disciplines constitute the present volume. This paper considers current research trends in the field of the science of ceramic interfaces and related applied aspects in the development of advanced ceramic materials. I. INTRODUCTION There is a growing awareness that many functional characteristics of ceramic materials are determined by interfacial layers, such as surfaces and grain boundaries, rather than by the bulk phase. Consequently, it is becoming increasingly important to utilise interfacial modifications in the preparation of advanced ceramic materials. In order to meet the requirements of modern day technology greater attention must be paid to these effects in addition to bulk properties alone. Consequently, better understanding of interface properties is a basic condition of further progress in materials science. The increasing awareness of importance of ceramic interfaces has resulted in accumulation a large amount of experimental and theoretical information concerning interface properties.

664

The specific experimental and theoretical approach to the study of interfaces of compounds has resulted in the formation of a new discipline: science of ceramic interfaces. This discipline deals with local properties of interfaces and their impact on materials properties. Progress in this discipline is important for the preparation of materials with enhanced properties. In contrast to the crystalline bulk, which is characterized by crystalline periodicity and homogeneous properties, the interface region is a very complex object involving strong chemical potential gradients and corresponding gradients in the electric potential within a region very limited in dimensions. These gradients have a substantial effect on the local transport kinetics [i, 2] and thus on reactivity of the material. It has also been observed that the segregation-induced enrichment of interfaces results in their structural deformations thus resulting in the formation of bidimensional structures which have outstanding properties not displayed by the bulk phase [i]. All these compositional and structural complications take place within a layer limited to 1-2 atomic layers in the case of metals and alloys [3, 4] and about 2-30 layers (or more) in the case of ionic solids of nonstoichiometric compounds [i, 5, 6]. This layer presents problems for qualitative evaluation of several properties which, within this layer, are functions of position. The science of ceramic interfaces combines several disciplines such as surface science, metallurgy, high temperature chemistry and electrochemistry. Surface science offers specialized techniques in studies of chemical composition and structure of the interface layer. Basic concepts of thermodynamics of solutes, thermodynamics of segregation and defect chemistry as well as grain boundary processes such as grain boundary diffusion have been developed mainly within metallurgy and high temperature chemistry. Finally, basic electrochemical concepts can be applied to consider electrochemical phenomena in the interface region, such as transport of charged defects in strong electric fields generated along segregation-induced compositional gradients. Accordingly, progress in the science of ceramic interfaces requires a collective effort of scientists representing all the above disciplines. It has been realized that interfaces of compounds, such as external surfaces and grain boundaries, are not limited to a bidimensional region demarcating individual phases or grains, as is the case for metals, but should be considered

665

rather as a tri-dimensional interface region [i, 5, 6]. Applied aspects of ceramic interfaces give the basis of another discipline: interface engineering. This discipline concerns modification of properties of ceramic materials through processing of interfaces and aims at the preparation of advanced materials with enhanced properties to meet the demands of rapidly developing technologies. The aim of the present volume was to address several important issues of the science of ceramic interfaces and related applied aspects. Scientists representing a multidisciplinary spectrum including experts from industry and academic organisations are contributing to the present volume. In considering the current trends in research and development on the science of ceramic interfaces and related applied aspects the following questions should be addressed: i. Which directions of the research on ceramic interfaces are the most important from the viewpoint of applications? Which areas are the most critical? 2. Which are the most important questions which require urgent attention? 3. Which are the most efficient ways in the modification of ceramic interfaces and how can the desired properties be achieved by engineering the interfaces? 4. Which topics, in the area of ceramic interfaces, should be taken as primary and secondary target of investigations? Basic conclusions arising from discussions during the present Workshop will be summarized below. 2. SYSTEMS Interfaces of ceramics and related applied aspects have been considered involving solid/solid, gas/solid and solid/liquid interfaces. Solid/solid interfaces involve mainly interfaces such as metal/ceramic, ceramic/ceramic, ceramic/polymer and ceramic/polymer/metal interfaces. Understanding of local interactions at these interfaces is important for the preparation of electronic components based on functional ceramic materials as well as for some structural ceramics. An extremely important area, from the viewpoint of electronic devices, concerns the metal/ceramic interface and related applied aspects in relation to the preparation of interconnections.

666

Gas/solid interfaces and related heterogeneous processes are important from the viewpoint of correct understanding of the effect of gas phase composition on processing of ceramics and their final properties. It has been realized that by changes of partial pressure of gas phase components during processing or subsequent standardization one may adjust the chemical activity of these components in a controlled way resulting in a desired reaction mechanism and, consequently, in desired material properties. Studies of the gas/solid system are of importance for the preparation of catalysts, chemical gas sensors and other functional materials such as dielectrics, semiconductors and superconductors. An awareness is growing that the properties of the surface active centres of catalysts, formed during the catalytic process, are not determined only by chemical composition and structure of the solid phase but also the gas phase composition. The liquid/solid system becomes increasingly important from the viewpoint of liquid corrosion of ceramics, such as erosion of ceramic materials used for immobilisation of nuclear wastes. Basic concepts and related applied problems concerning these materials have been extensively described in the review papers of Vance et al. [7], Turner et al. [8] and Blesa [9]. 3. MATERIALS 3.1. Oxide Materials Most of the literature data on properties have been reported for binary metal oxides and their solid solutions [i, i0, ii] which are the basic constituents of ceramic materials. Many of the material data have been reported for materials not well characterised for their impurity level. An awareness is growing that impurities may have a substantial effect on properties even if present at the level of several ppm. Therefore, the available material data require verification from the viewpoint of the effect of impurities. An important family of materials are perovskites which have been widely applied as basic components of electronic materials, semiconductors, superconductors and sensor-type materials. Most of the literature data has been accumulated for BaTiO 3 and SrTiO 3 which serve as model ternary oxides of perovskite structure. There has been observed a growing interest in the effect of segregation on interface composition in these compounds [12, 13]. Studies of the effect of

667

segregation on electrical properties of BaTiO 3 and SrTiO 3 are also important for correct understanding of the properties of dielectrics. YBa2Cu3Oz.x is a quaternary compound of perovskite-type structure and is the most widely studied oxide superconductor. It has been shown that segregation may be responsible for changes of its local composition resulting in the local critical temperature (Tc) and the critical current (Jc) to be different to those in the bulk phase [14]. One may expect that appropriate grain boundary engineering may result in fabrication of material with both T c and Jc above the level currently reported for the bulk phase. One may expect that further elevation of the critical parameters in the bulk phase will be very limited while appropriate preparation of interface structures may result in T c of 250 K and possibly well above this value [14]. 3.2. Effect of Impurities on Properties One of the most important feature of any material, having a direct impact on properties, is its chemical composition. Properties of compounds may be substantially changed by doping with aliovalent ions forming donors and/or acceptors [i0, ii]. Awareness is growing that in considering chemical composition one should also take into account the presence of impurities (unintentional dopants) involving both cations and anions. The effect of impurities on properties, especially of aliovalent ions, may be very strong. It has been shown that even if an impurity is present in the bulk only at a level of 'ppb' it may achieve substantial concentration at the surface, if the driving force for segregation is high [i, 16, 17]. Quantitative analysis of impurities is both difficult and expensive. Therefore, most of the materials described in the literature are not characterized for their impurity spectrum. One should, however, realize that publication of materials properties, without detailed impurity analysis, may result in apparent conflicts in the reported material data. For correct understanding of these apparently conflicting experimental data it is imperative that materials characterization involves detailed analysis of the impurities present at the ppm level or even below this level. The effect of impurities is especially important for crystals of relatively small stoichiometry such as AI203 and MgO [i0]. Then even a small amount of cation impurities may change the conductivity type from n- to p-type or turn an

668

insulator into a semiconductor. There is an urgent need to work with well defined materials which either are well characterized for their impurity level or are very pure. Only experimental data which were determined for pure specimens may be considered as reference material data. It has been shown that even if the impurities have negligible effect on bulk properties their impact on the interface layer may be significant and cannot be ignored [i, 16, 17]. The most widely studied and described effect of segregation on processing concerns the beneficial role of a small addition of MgO on sintering of alumina, resulting in the preparation of ceramic material of almost 100% density [1821]. Despite the impressive amount of work devoted to this material little is known of the role of segregated defects and surface states on densification, grain growth and sintering kinetics of powders. Important types of impurities are those coming from the gas phase. An important gas phase component, which usually exhibits high reactivity with materials, is oxygen. Its content in the bulk phase of oxide compounds is determined by oxygen activity of the gas phase during high temperature processing. During cooling the penetration of oxygen into the oxide lattice is limited, due to kinetic reasons, resulting in the formation of concentration gradients along interfaces. In the case of several electronic materials the formation of this gradient is intentional. Oxygen is also the lattice component in a wide range of materials: oxide materials. The relationship between oxygen partial pressure during high temperature treatment and the lattice oxygen content has been described for most binary metal oxides [i0]. Other important impurities, coming mainly from the gas phase, involve hydrogen and carbon. The effect of hydrogen on defect structure and related properties of nonstoichiometric oxides has been considered by Norby and Kofstad [22] and Waser [23]. 3.3. Low Dimensional Systems It has been observed that low dimensional systems exhibit outstanding properties which are not displayed by the bulk phase [i]. Therefore, investigations of new materials with enhanced properties, especially functional materials, should be focused on the following systems: i. Thick and thin films of both organic and inorganic materials deposited by both CVD and PVD. Understanding

669

2. 3.

4.

5.

of interactions between the film and the substrate is required for engineering of film properties. Multilayer systems ("nanoceramics') involving layers of different composition, structure and semiconducting properties, Heterogeneous system involving a dispersed minor phase within a basic phase. The procedure for the formation these heterogeneous systems, developed by Wagner et al [24], is termed 'heterogeneous doping'. Coated systems involving deposition of a surface layer of desired composition over the bulk phase. The procedure of coating becomes increasingly important in covering bulk metal materials with ceramic layers of superior properties. This procedure has also been recently applied in the preparation of gas sensors, to minimize their response time and to increase both sensitivity and selectivity. Homogeneous coating, which is generated e.g. by segregation-induced enrichment of interfaces, may be considered as a type of coating with species coming from the bulk phase.

3.4. Surface Active Centres Gas/solid heterogeneous processes play an important role in processing of ceramics. Also many function-related properties of ceramics, such as ionic conductivity, may be determined by the gas/solid heterogeneous kinetics. It has been a general assumption that the gas/solid heterogeneous kinetics is determined by bulk lattice transport while surface processes are very fast. It has been shown that surface reactions, such as adsorption, dissociation and incorporation into the surface layer [25], may play an important role in these processes and may be rate controlling for the overall gas/solid kinetics even at high temperatures. In this case the gas/solid kinetics may be enhanced by surface active centres which act as catalytically active centres for a given surface reaction. Better understanding of the catalytical properties of ceramics materials is required for the preparation of not only better catalysts but also electrochemical devices such as electrochemical gas sensors and fuel cells.

670 4. M E T H O D S

Most of the experimental methods for the determination of chemical composition and structure operate at room temperature. In addition most surface sensitive methods such as XPS, SIMS, AES, ISS and LEED operate under high vacuum. Then the analyses can be performed in conditions remote from the conditions of materials processing involving elevated temperatures and controlled gas phase composition. Accordingly, there is an urgent demand for the development of composition and structure sensitive experimental methods which could operate 'in situ' during the high temperature processes. It has been shown by Brongersma [26] that depth profiles may have complex structures involving layers of entirely different composition as a function of the distance from the interface. A similar sandwich-type model was proposed for the surface layer of yttria-doped zirconia [16]. Accordingly, correct determination of depth profiles, involving quantitative analysis of chemical composition of the outermost and also of deeper layers, requires parallel application of different techniques of various depth resolutions. It has been shown that bidimensional structures may be formed at the outermost interface layers, if segregationinduced enrichment exceeds a certain critical concentration value. Identification of these local low dimensional structures is important for reproducible preparation of these structures. It is important that properties of these structures are determined 'in situ' in conditions of their formation. Their properties may substantially change during cooling and quenching. Thus development of sophisticated surface techniques, which are adequate for 'in situ' studies of the bidimensional structures, is required. Recent efforts tend to perform surface studies of materials at elevated temperatures by using electron microscopy. Surface dynamics observed in this way allow the determination of the local transport kinetics data [27]. In many cases segregation may also lead to the formation of grain boundary nanophases such as those in zirconia [28, 29]. These phases have a strong impact on both electrical and mechanical properties. It has been shown that phase relations for the low dimensional systems are different to those for the bulk phase and must be considered as functions of position [30]. Therefore, engineering of phase relations at interfaces requires the determination of these local phase diagrams.

671

It is important to study interfaces both on an atomic scale as well as to determine their collective properties such as work function [31]. Most surface techniques have been applied to metallic surfaces. Studies of compound surfaces are subject to certain difficulties and require a certain methodical approach. This volume as well as previous volumes [i, 31] address this problem. Finally, more sophisticated theoretical methods are required for quantitative analysis of defect-related properties in crystals of nonstoichiometric compounds. 5. EXPERIMENTAL MATERIAL An accumulation of empirical data concerning interfaces is urgently required for constructing more general models and theories for description of fundamental interface phenomena such as interfacial segregation and grain boundary diffusion. Based on these models we may be able to construct more sophisticated models describing processes in ceramics such as sintering. Unfortunately, in contrast to the vast experimental data base on segregation and grain boundary diffusion which has been accumulated for metals little is known for compounds. The determination of thermodynamic data of segregation would allow one to predict the effect of segregation on the chemical composition of the interface region and thus, by adjusting the experimental conditions, to control this composition and, consequently, to engineer the related properties. In considering interface properties the basic principles of defect chemistry may be developed into the interface region involving quantities such as local defect formation energies and their mobility terms. Also kinetic data concerning both segregation and evaporation (sublimation) are required for correct analysis of the effect of segregation on interface chemistry. Data on grain boundary migration and the kinetics of transport of defects along and across grain boundaries is important to understand the processing of ceramics. A considerable experimental and theoretical data base in this matter has been accumulated for metals and alloys. Studies of grain boundary diffusion for compounds are, so far, limited to NiO [32], AI203 [32] and ZrO 2 [27]. Diffusion data for these compounds are based on the theoretical models

672

developed for metals [32, 33]. However, because of specific properties of the interface region of nonstoichiometric compounds, which are so different from those of metals, there is an urgent need to develop particular solutions of the diffusion equation for the ionic solids based on nonstoichiometric compounds. 6. GENERAL COMMENTS It has been realized that apparently homogeneous and uniform ceramic materials exhibit strong compositional gradients within the interface region. These local gradients have a substantial impact on the properties of the materials. Therefore, there is an urgent need for better understanding of the relationship between the local properties of interfaces and properties of ceramics and, consequently, to engineer this interface region in order to achieve the desired properties. Because of the obvious complexity of studies on interfaces involving different disciplines, progress in the science of ceramic interfaces may only be possible within cooperative research efforts between specialized research centers focusing their expertise on interface properties and interface phenomena. Cooperative research programs are also the only way to decrease the extensive costs of investigations resulting from the sophisticated and expensive techniques required in studies of interface properties. The costs involve the equipment itself, operating expenses and skilled technical services. Difficulties in studies of interfaces result mainly from very limited dimensions of the interface region which exhibit gradients of properties. Therefore, an objective picture of the interface region can only be obtained by using simultaneously several methods to examine the local chemical composition, structure and related properties. Again, this is possible through joint endeavors between laboratories with specialised expertise and related equipment. The distance between basic research on interfaces and related applied aspects is substantial. Bridging this distance is expensive and takes a long time. Therefore, it is important that studies on ceramic interfaces are carried out in close collaboration between academic scientists and industrial researchers.

673

There is a long list of industrial aspects of interfaces. The most important include resistors, dielectric materials, superconductors, ceramic gas sensors and fuel cells 9 Resistors have been manufactured, so far, in an empirical manner. Development of grain boundary engineering will enable one to gain a better understanding of the properties of the interface region and will result in the preparation of resistors with enhanced properties 9 Taking into account that a year production amounts to billions of resistors one may realize the extent of the economic impact of any progress in interface engineering which results in enhanced properties of these materials 9 ACKNOWLEDGEMENTS Support of the Commonwealth of Australia through the Department of Industry, Science and Technology (Grant # C91/02826) is gratefully acknowledged. In the preparation of this paper I have drawn on comments of several Workshop participants, mainly: S.P.S. Badwal, H.H. Brongersma, S.B. Desu, A. Glaeser, R. Hannink, A.E. Hughes, H. Ichimura, P. M~ller, K. Niwa, R.St.C. Smart, J.B. Wagner, Jr. and Da Yu Wang. REFERENCES 1 2 9 3. 4. 5. 6. 7. 8. 9.

J Nowotny, in: 'Science of Ceramic Interfaces' Elsevier, Amsterdam, 1991, p. 79 I. Kaur and W. Gust, "Fundamentals of Grain and Interface Boundary Diffusion", Ziegler Press, Stuttgart, 1988 J. Cabane and F. Cabane, in ref. [i], p. 1 McLean, "Grain Boundaries in Metals", Oxford University Press, London, 1957 W. Hirschwald, in: 'Surface and Near-Surface Chemistry of Oxide Materials', J. Nowotny and L.C. Dufour, Eds., Elsevier, Amsterdam, 1988, p. 61 P. Wynblatt and R.C. McCune, in ref. [i], p. 247 E.R. Vance et al., this volume, p. 431 Turner et al., ref. [i], p. 663 M.A. Blesa, A.E. Regazzoni and A.J.G. Maroto, in: ,External and Internal Surfaces of Metal Oxides' L C Dufour and J. Nowotny, Eds., Trans Tech Publications, Zurich, 1988, p. 31

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i0. ii. 12. 13. 14. 15. 16. 17 18. 19. 20. 21. 22. 23. 24. 25. 26. 27. 28. 29. 30. 31 32. 33.

P. Kofstad, "Electrical Conductivity, N o n s t o i c h i o m e t r y and Diffusion in Binary Metal Oxides", Wiley & Sons, New York, 1972 F. Kroger, 'Chemistry of Imperfect Crystals' , NorthHolland Publ. Co., Amsterdam, 1974 Y.M. Chiang and T. Touichi, J.Am. Ceram. Soc., 73 (1990) 3278, 3286 S.B. Desu and D.A. Payne, J.Am. Ceram. Soc., 73 (1990) 3391, 3398, 3407, 3416 J. Nowotny, M. Rekas, D.D. Sarma and W. Weppner, in ref. [i], p. 669 S.X. Dou and H.K. Liu, this volume, p. 239 J. Nowotny, M. Sloma and W. Weppner, Solid State Ionics 28/30 (1988) 1445 J Nowotny, in: "Grain Boundaries of Ceramic Materials" N. Kosuge, Ed., Publ. by Amer. Ceram. Soc., in print R.L. Coble, J.Appl. Phys., 32 (1961) 487 P.J. Jorgensen and J.H. Westbrook, J.Am. Ceram. Soc., 47 (1964) 332 S. Baik, J.Am.Ceram. Soc., 69 (1986) Cl01 S.M. Mukhopahay, A.P. Jardine, J.M. Blakely and S. Baik, J.Am. Ceram. Soc., 71 (1988) 358 T. Norby and P. Kofstad, J.Am. Ceram. Soc., 67 (1984) 786 R. Waser, J.Am. Ceram. Soc., 71 (1988) 58 J.B. Wagner, Jr., Mat.Res. Bull., 15 (1980) 1691 Z. A d a m c z y k and J. Nowotny, J.Phys. Chem. Solids 47 (1986) ii H.H. Brongersma, this volume, p. 113 M. Kusunoki, in ref. [17] S.P.S. Badwal, J. Dreannan and A.E. Hughes, in ref. [i], p. 227 A.E. Hughes, this volume, p. 183 J. Nowotny, this volume, p. 1 'Surface and N e a r - S u r f a c e Chemistry of Metal Oxides' J. Nowotny and L.C. Dufour, Eds., Elsevier, Amsterdam, 1988 E.G. Moya, this volume, p. 227 I. Kaur and W. Gust, 'Fundamentals of Grain and Interphase Boundary Diffusion', Ziegler Press, Stuttgart, 1988

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SUBJECT

INDEX

Aliovalent ions 9 A1203 295 Alumina 33, 46, 295 dense 46 surface transport 48 AIN Applied aspects 23 Bidimensional structures 15 Capacitors 613, 622 Charge neutrality 5 surface 9 Ceramic processing 399 Composites 593 CoO 295 Copper films 473 on metal oxide 473 Corrosion 629, 639 aqueous 639 Crack healing 50, 52, 56 CrN 629 oxidation 636 Cr203 grain boundary 295 diffusion 295 Cu

on MgO on CaO on TiO 2 on AI203 on ~-Fe203 on SrTiO 3 on ZnO Defect structure Densification Dielectric materials properties

478 480 487 499 507 509 512 4 413 615 618

Diffusion driven interface grain boundary induced phenomena Embrittlement Gas sensing Electrical conductivity 615, properties 71, Electrochemical cell systems Electrodes Electrode morphology reactions Energy conversion Etching Flux pinning Grain boundary amorphization diffusion 277, AI203 CoO Cr203 NiO film migration 46, roughening structure thin film wetting Grain alignment Grain boundaries superconductors weak links

356 277 353 593 645 617 622 645 71 613 75 73 645 40 256 388 353 295 295 295 295 393 395 392 253 371 372 254 246 241

676

High Tr s u p e r c o n d u c t o r s 25, 239 w e a k links 241 Hot p r e s s i n g 40 High R e s o l u t i o n T r a n s m i s s i o n Electron M i c r o s c o p y (HRTEM) 441 results 444 Implantation 555 Impurities 667 Interface engineering 26 phenomena 399 Interfaces electrolyte 88 e l e c t r o d e grains 88 ceramic substrates 341 grains i00 i n t e r m e d i a t e phases i00 microdesign 33 nuclear technology 267 precipitation 88 semiconducting 21 properties 21 synroc 431 zirconia 71, 183, 357 Interphase 79 Liquid immersion 399 Low energy ion s c a t t e r i n g (LEIS) 113 experimental .121 oxide surfaces 113 quantification 134 surface s t r u c t u r e 142 Low d i m e n s i o n a l structures 15 systems 668 Materials 666 Mechanical testing 594 Microdomain 97 Mobility e l e c t r o n i c carriers 7

Mullite 571 basic p r o p e r t i e s 574 microstructure 577 oxygen conductor 584 t h e r m a l shock 581 Multilayer capacitors 622 Nanodomain 441 Nickel films 473 on m e t a l o x i d e 473 Ni on MgO 480 on NiO 482 on TiO 2 498 on ~-AI203 504 on ZnO 519 NiO 295 Nonstoichiometry 4 effect of p(02) 5 Nuclear technology 267 waste ceramic 431 R a d i a t i o n effects 433 Sapphire 52, 56 Scandium tantalate 441 Segregation e l e c t r i c fields 16 enrichment 9 extrinsic 281 intrinsic 281 zirconia 79 Sensors 23 SnO 2 527 conductivity 527 electronic structure 537 doping 527, 547 implantation 555 Sb-doped 550 s u r f a c e states 543 surface s t r u c t u r e 529 thin film 547 Structures low d i m e n s i o n a l 668 Superconductors 239

677

Surface active centres conductivity equilibration defects layer segregation 9, 16, structure 15, 142, Surface vibrational spectroscopy Systems Thin Films 21, 371, Tin (IV) oxide TiN oxidation Transport effect of interfaces Trends of research Weak links Wetting 164, Zirconia 71, based ceramics composites grain boundaries interfaces 71, ZrO 2 71,

669 563 19 145 311 154 148 559 665 473 527 629 632 16 16 663 241 371 183 357 183 198 185 183

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