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 N...

Author: Nowotny J.

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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


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-.

:

U 0

SPACE CHARGE

I = ! !

Z

l..ul

_3 4: l--l.ul _3 Lul

§

|

|

|

|

§

§ §

~

,,

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


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

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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.

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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

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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

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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

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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 _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

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T, ~ Figure 6. Non-ideal behaviour of Pt electrodes in potentiometric sensors at low temperatures. Air versus 5.01~ in N 2. After the heat treatment of electrodes at (O) 600~ (zx) 750~ and ()900~ in air. 15igure courtesy of F.T. Ciacchi. ity gases and oxygen, the non-ideal behaviour is considerably reduced [30]. Thus even for clean electrode/electrolyte interfaces, electrode morphology and interference from external factors can substantially modify the electrode performance.

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')

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

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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

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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

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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

99

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wt. % A1203

Figure 24. The effect of alumina addition on the resistivity of 5.9 mol% 5 C 2 0 3 - ZrO 2 ceramics. The data were recorded at 450~ [30].

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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Figure 27. Impedance spectra (recorded at 450~ in air) for 3.4 wt% MgO - ZrO 2 (MgOPSZ) before ((3) and after (o) annealing at the modest ageing temperature of 1000~ for l l0h. A substantial increase in the low frequency grain boundary arc appears to have occurred.

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', ~, 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.

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107

10 11 12 13

14 15 16 17

18 19 20 21 22 23 24 25

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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 and 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 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 azimuth [6].

closely with ICISS. Fig. 4.4 shows the Mg peak for the azimuth as a function of the angle of incidence. Due to the larger distance between the neighbouring atoms in the azimuth as compared with that in the 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 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) = 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 _

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o / /

-

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-

i

/

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/

Ir

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9

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=,' ~

t

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li~x").~

~ ~ "

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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

.':

...

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..

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:::

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-,b ~

-"I', I

1

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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].

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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+ ,

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 [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

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 (_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(.

I0 z

o=

Z (~ uu >. k1600~ 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 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)


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
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. 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

"/

>

(D

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03 (D c" LU 0 (D .~ x,,

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1610

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,

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1600

1800

2000

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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.



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)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 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 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 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_

D

~

16h

. . . . . . .

9

6

U

o

I

O.Ol Pore d i a m e t e r

I o.I (Fro)

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 ,
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


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 ( 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 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.


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 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, 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.

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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

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I

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,

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

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' 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 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 ! : ~

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\%,\\%\% "

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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

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

.....@ ~

i

i

9

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i

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.....

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

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

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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--

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

~-15

e

.

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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.


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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9

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500

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'

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'

'

'

'

i

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1000 1500 2000 V - IR (mV)

'

'

'

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'

2500

'

'

"

3000

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

0.0 540~ -0.1 0

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 . . . . .

300'

~

~ ( 3 )

9

L)

m

)

(2)

200' 9

1t

I

",1:1

]

100'

,

0

" ' ,

0

540~ " ' "

5

|

. . . .

,

. . . .

,

. . . .

,

. . . .

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~

OOoo w o

0.8

0

0.6

0.4 0.2 0.0

9

0

l

10

9

9

I

"

~

I

~

9

ii

"

9

I

"

9

li

20 30 40 50 60 Drawing Current (uA)

"

9

!

70

"

"

80

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 ~'

o=

0.80

ra~

12.6% Methane in Air 0.60

'1

100

"

"

9

"

J

i"

300 500 Temperature (~

'

'

'

'

7(

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

~o

....

60

A

40 -~

Effective Range

20 9

9

.

9

I

.!

200

|

.

9

I

.

.

400 Temperature (~

.

.

I

600

.

9

.

.

800

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.

1.00 >

0.75

9

0.50 "

= o

0.25 0.00 101

.

"

9

9

9 ,|,ll

9

9

9 | |i,,I

J

9 i

9 ,i i,I

|

10 2 10 3 10 4 Methane in Air (ppm)

9

9 J ii JJl

10 5

,

,,,

(470oc)

(610~

|i i,

10 6

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.

~

....

~"

680~

o

S

;~ 0.0 ~ -1.0 ~I::I 0

-2.0

o

.~.A~v

........

~

o

10 -2

'

........

10 -i

'

........

'

10 0 Oxygen (%)

........

101

10 2

Figure 9. The plots of open-cell potential versus the oxygen concentration for a sensor immersed in a O2-N2 gas mixture.

70

o

~. >.

60

9 40% N20,13% 0 2

~

50

=

40

0 ~,

30 20

O0

0

0

0 10

00 .

0 500

.

.

.

1

.

600

.

.

.

9 I

m

m

00 ,

.

9 I

700 800 Temperature (~

9

.

9 -

.

9 I

900

9 9

9

9

1000

100% N20

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

680oc '

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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