Functionally Graded Materials 1996

FUNCTIONALLY GRADED MATERIALS 1996 Organized and sponsored by THE FCM FORUM THE SOCIETY OF NON-TRADITIONAL TECHNOLOGY...

Author: I. Shiota | Y. Miyamoto

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FUNCTIONALLY GRADED MATERIALS 1996

Organized and sponsored by THE FCM FORUM THE SOCIETY OF NON-TRADITIONAL TECHNOLOGY Supported by SCIENCE AND TECHNOLOGY AGENCY, JAPAN

FUNCTIONALLY GRADED MATERIALS 1996 Proceedings of the 4th international symposium on Functionally Graded Materials, AIST Tsukuba Research center, Tsukuba, Japan, October 21-24,1996

Edited by

ICHIRO SHIOTA Department of Environmental Chemical Engineering, Kogakuin university, 2665-1, Nakanocho, Hachioji, Tokyo 192, Japan

YOSHINARI MIYAMOTO The Research center for Cyclic Loop Systems for Processing and iviaintenance. Joining and welding Research institute, Osaka university, ibaraki, Osaka 567, Japan

ELSEVIER Amsterdam - Lausanne • New York - Oxford - Shannon - Singapore - Tokyo 1997

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

ISBN 0 444 82548 7 © 1997, 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 u.S.A.-This publication has been registered with the copyright Clearance Center inc. (CCC), 222 Rosewood Drive, Danvers, MA 01923. information can be obtained from the CCC about conditions under which photocopies of parts of this publication may be made in the U.S.A. All other copyright questions, including photocopying outside of the U.S.A., 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

PREFACE

A formulated concept of functionally graded materials(FGMs) was proposed in 1984 by material scientists in Sendai area, Japan, as a means of preparing thermal barrier materials, and a coordinated research was developed in that country since 1986. The idea, that continuously changes in the composition, microstructure, porosity, etc., of these materials resulting in gradients in such properties as mechanical strength and thermal conductivity, has spreaded world-wide during the past ten years through the research. Aiming at opening channels among researchers working in state-of-the-art FGM topics and at discussing further developments in the FGM field, the first FGM symposium was held in Sendai, the birthplace of FGMs, in 1990 followed by the 2nd in San Francisco, 1992, and then the 3rd in Lausanne, 1994. Through these activities, the idea of graded structures and functions has attracted the attention of many scientists and researchers for its boundless scope in materials science and engineering. In this symposium, nearly three hundreds participants joined in order to exchange information which covers all aspects of functionally graded materials including their design, process and evaluation of structure, function, and integration, as well as applications. In particular, it should be noted that fifty five of these participants were from many countries in the worid. As a chairman of this symposium, I expect that this proceedings will be useful for FGM scientists and engineers, and will contribute to promote the research and development of FGMs in future.

V*^' Mitsue Koizumi Chairman

This Page Intentionally Left Blank

LIST OF ORGANIZING COMMITTEE

Chairman: Vice-Chairman:

Prof. M. Koizumi (President of the FGM Forum, Ryukoku University) Prof T. Hirai (Tohoku University)

Advisory Committee Dr. L. I. Anatychuk Prof F. Erdogan Dr. R. Ford Prof A. Glaeser Mr. T. Hirano Prof N. Ichinose Prof B. Ilschner Mr. R. Imoto Dr. M. Kamimoto Prof W. Kaysser Mr. T. Kurino Dr. A. Mazurenko Prof A. Merzhanov Prof Z. Munir Dr. Y. Nikolaev Dr. B. Rabin Prof S. Suresh Prof R. Watanabe Prof R. Yuan

Institute of Thermoelectricity, Ukraine Lehigh University, Pennsylvania, USA Materials Technology, New York, USA University of California, Berkeley, USA Daikin Industries, Ltd., Kusatsu, Japan Waseda University, Tokyo, Japan Swiss Federal Institute of Technology, Lausanne, Switzerland Science and Technology Agency, Tokyo, Japan Electrochemical Laboratory, Tsukuba, Japan German Aerospace Research Establishment, Koln, Germany The Society of Non-Traditional Technology, Tokyo, Japan Institute of General & Inorganic Chemistry, Kiev, Ukraine Institute of Structural Macrokinetics, Chemogolovka, Russia University of California, Davis, USA Research Institute of SIA LUCH, Moscow, Russia Idaho National Engineering Laboratory, Idaho Falls, USA Massachusetts Institute of Technology, Cambridge, USA Tohoku University, Sendai, Japan Wuhan University, Wuhan, China

Executive Committee Prof Y. Tada (Chairman) Nihon University, Chiba Dr. K. Eguchi National Aerospace Laboratory, Tokyo Dr. R. Fukuda Electrotechnical Laboratory, Tsukuba Prof Y. Miyamoto Osaka University, Osaka Dr. M. Niino National Aerospace Laboratory, Kakuda Dr. I. A. Nishida National Research Institute for Metals, Tsukuba Prof I. Shiota Kogakuin University, Tokyo Ms. S. Tsuda The Society of Non-Traditional Technology, Tokyo Dr. S. Yatsuyanagi National Aerospace Laboratory, Kakuda

Program Committee Prof. I. Shiota (Chairman) Dr. N.Cherradi Prof. M. Gasik Prof T. Kawasaki Dr. A. Kumakawa Prof Y. Miyamoto Prof J. Yoshino

Kogakuin University, Tokyo, Japan Swiss Federal Institute of Technology, Lausanne, Switzerland Helsinki University of Technology, Helsinki, Finland Tohoku University, Sendai, Japan National Aerospace Laboratory, Kakuda, Japan Osaka University, Osaka, Japan Tokyo Institute of Technology, Tokyo, Japan

Tsukuba Regional Committee Dr.R.Fukuda (Chairman) Dr.T.Fujii Mr.Y.Imai Mr.Y.Kasuga Dr.K.Kato Dr.I.Kojima Ms.M.Maeda Mr.A.Negishi Dr.I.A.Nishida Mr.Y.Nishio Mr.T.Ohta Dr.Y.Shinohara Mr.J.Teraki Mr.A.Yamamoto

ElectrotechnicalLaboratory,Tsukuba National Research Institute for Metals,Tsukuba National Research Institute for Metals,Tsukuba ElectrotechnicalLaboratory,Tsukuba National Institute of Materials&Chemical Research,Tsukuba National Institute of Materials&Chemical Research,Tsukuba The Society of Non-Traditional Technology,Tokyo Electrotechnical Laboratory,Tsukuba National Research Institute for Metals,Tsukuba Daikin Industries, Ltd.,Tsukuba Electrotechnical Laboratory,Tsukuba National Research Institute for Metals,Tsukuba Daikin Industries, Ltd.,Tsukuba Electrotechnical Laboratory,Tsukuba

CONTENTS

GENERAL TOPICS

FGM research programs in Japan -from structural to functional uses Mipmoto, Njino and Koizumi ............................................................................

1

Research program on gradient materials in Germany R&je/ and A. Neubrand.. ..............................................................................................

9

Lessons learnt in 7 years of FGM research at Lausanne

6.//schner .............................................................................................................................. PART I

15

STRUCTURAL MATERIALS

Design and Modeling

Local fields in functionally graded materials Y.D. Bi/otsky and M.M. Gasik ............................................................................................

21

Computer-aided process design for forming of pore-gradient membranes

c.w.

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

29

Mathematical model for axial-symmetrical FGM and C.C. ..................................................................................... X.D.

35

Stress analysis in a two materials joint with a functionally graded material y- Yang and D.Munz .........................................................................................................

41

Optimum design and fabrication of TiC/Ni,Al-Ni functionally graded materials Q- Shen, X.F. Tang, R. T", L.M. Zhang and R.Z. Yuan ..................................................

47

A mathematical model for particle distribution in functionally graded material produced by centrifugal cast B. Zhang, J. Zhu, Zhang, Z. Ying, H. Cheng and G.An ............................................

53

Modeling and measurement of stress evolution in FGMcoatings during fabrication by thermal spray S.Kuroda, y. Tashiro and Fukushima .........................................................................

59

Artificial neural network used for TiB,-Cu FGMdesign cao and C.C. .................................................................... Z C. Mu,Z.X.

65

Deformation analysis of graded powder compacts during sintering Shinagawa .......................................................................................................................

69

Simulation of the elasto-plastic deformations in compositionaily graded metalceramic structures: Mean-field and unit cell approaches H.E. Pettermann, E. Weissenbek and S. Suresh

75

Large deflections of heated functionally graded clamped rectangular plates with varying rigidity in thickness direction F. Mizuguchi and H. Ohnabe

81

Model investigation of ceramic-metal FGMs under dynamic thermal loading: Residual stress effect, thermal-mechanical coupling effect and materials hardening model effect O.J. Zhang, P.C. Zhai and R.Z. Yuan

87

Fractal geometry and it's implications to surface technology D.P. Bhatt, O.P. Bahl, R. Schumacher and H. Meyer

93

Database system for project of the functionally graded materials K. Kisara, A. Moro, Y.S. Kang and M. Niino 1-2

99

Fracture Analysis

Fracture mechanics of graded materials F. Erdogan Microstructural effects in functionally graded thermal barrier coatings M.J. Pindera, J. Aboudi and S.M. Arnold

105 113

Micromechanical failure criterion for FGM architecture studied via disk-bend testing of Zr02/Ni composites T. Ishizuka, Y. Ohta and K. Wakashima Thermomechanical response characteristics of ZrOg/Ni functionally graded materials: An experimental study to check model predictions T. Ishizuka, C.S. Kang and K. Wakashima Micromechanical approach to the thermomechanical analysis of FGMs S. Nomura and D.M. Sheahen Effect of gradient microstructure on thermal shock crack extension in metal/ceramic functionally graded materials A. Kawasaki and R Watanabe Thermal fracture mechanisms in functionally graded coatings K. Koklnl, Y.R. TakeuchI and B.D. Choules 1-3

123

131 137

143 149

Powder Metallurgical Process

Fabrication of AIN/W functionally graded materials K. Sogabe, M. Tanaka, T. Mlura and M. Tobloka Graded casting for producing smoothly varying gradients in materials B.R. Marple and S. Tuffe

155 159

XI

Gradient components with a high melting point difference M. Joensson, U. Birth and B. Kieback

167

Fabrication of pore-gradient membranes via centrifugal casting aw. Hong, F.Muller and P. Greil

173

Mechanical properties and microstructure of insituTiCp reinforced aluminum base FGM by centrifugal cast a Zhang, J. Zhu, Y. Zhang, Z Ying, H. Cheng and G. An

179

Dispersion and fabrication of ZrOj/SUSSIG functionally graded material by tape casting process J.G. Yeo, Y.G. Jung and S.C. Choi

185

Fabrication of ZrOj/Ni and ZrOj/AljOg functionally graded materials by explosive powder consolidation technique A Chiba, M. Nishida, K. Imamura, H. Oguraand Y. Morizono

191

Development of metal/intermetallic compound functionally graded material produced by eutectic bonding method S. Kihhara, T. Tsujimoto and Y. Tomota

197

Mechanical performance of Zr02-Ni functionally graded material by powder metallurgy J.C. Zhu, S.Y. Lee, ZD. Yin andZH. Lai

203

Fabrication of PSZ-SUS 304 functionally graded materials H. Kobayashi

209

Preliminary characterization of interlayer for Be/Cu functionally gradient materials - thermophysical properties of Be/Gu sintered compacts S. Saito, N. Sakamoto, K. Nishida and H. Kawamura 1-4

215

Deposition and Spray Process

Electrophoretic forming of functionally-graded barium/strontium titanate ceramics P. Sarker, S. Sakaguchi, E. Yonehara, J. Hamagami, K. Yamashita and T. Umegaki Processing and properties of electrodeposited functionally graded composite coatings of Ni-AI-AljOg K. Barmak, S.W. Banovic, H. M.Chan, L.E. Friedersdorf, M.P. Harmer, A.R. Marder, CM. Petronis, D.G. Puerta and D.F. Susan Functionally graded materials by electrochemical modification of porous preforms A Neubrand, R. Jedamzik and J. Rode! Thermal management of carbon-carbon composites by functionally graded fiber arrangement technique Y. Kude and Y. Sohda

221

227

233

239

Xll

Formation and properties of TiC/Mo FGM coatings T. Fukushima, S. Kuroda, S. Kitahara, K. Ishida and M. Sano Formation of a Ti-AlgOa functionally graded surface layer on a Ti substrate with the use of ultraflne particles A Otsuka, H. Tanizaki, M. Niiyama and K. Iwasaki

245

251

Oxidation-resistant SiC coating system of C/C composites N. Sato, I. Shiota, H. Hatta, T. Aoki and H. Fukuda

257

AljOg-ZrOj graded thermal barrier coatings by EB-PVD-concept, microstructure and phase stability U. Leushake, U. Schuiz, T. Krell, M. Peters and WA. Kaysser

263

Microstructure characteristic of plasma sprayed ZrOj/NiCoCrAlY graded coating Z Yin, X. Xiang, J. Zhu and Z Lai 1-5

269

Reaction Forming Process

Formation of functionally-graded materials through centrifugally-assisted combustion synthesis W. Lai, ZA. Munir, BJ. McCoy and S.H. Risbud

275

SHS - a new technological approach for creation of novel multilayered diamondcontaining materials with graded structure E.A. Levashov, LP. Borovinskaya, A.V. Yatsenko, M. Ohyanagi, S. Hosomi and M. Koizumi 283 Graded dispersion of diamond in TiB2-based cermet by SHS/dynamic pseudo isostatic compaction(DPIC) M Ohyanagi, T. Tsujikami, M. Koizumi, S. Hosomi, E.A. Levashov and l.P. Borovinskaya

289

Annealing of ceramic/metal graded materials fabricated by SHS/QP method A.N. Pityulin, Z.Y. Fu, M.J. Jin, R.Z. Yuan and AG. Merzhanov

295

Thermodynamic calculation and processing of TiBg-Cu FGM C.C. Ge, Z.X. Wang and W.B. Cao

301

Fabrication of Al-Cu system with functionally graded density profiles R. Tu, O. Shen, J.S. Hua, L.M. Zhang and R.Z. Yuan

307

AlgOgto Ni-superalloy diffusion bonded FG-joints for high temperature applications L IHeikinheimo, M. Siren and M.M. Gasik 1-6

313

Novel Process

Advances in the fabrication of functionally graded materials using extrusion freeform fabrication G.E. Hilmas, J.L. Lombardi and R.A. IHoffman

319

Novel routes to functionally graded ceramics via atmosphere-induced dopant valence gradients M Kitayama, J.D. Powers and AM Glaeser

325

The growth of functionally graded crystals by verneuil's technique M Ueltzen, J.F. Fournie, C. Seega and H. Altenburg

331

Excimer laser processing of functionally graded materials Y. Uchida, J. Yamada, Y.P. Kathuria, N. Hayashi, S. Watanabe, S. Higa, H. Furuhashi and Y. Uchida

337

Development of stainless steel/PSZ functionally graded materials by means of an expression operation K. Taka, Y. Murakami, T. Ishikura, N. Hayashi, S. Watanabe, Y. Uchida, S.Higa, T.lmura and D. Dykes 343 Microwave sintering of metal-ceramic FGM M A Willert-Porada and R. Borchert

349

Residual stress control of functionally graded materials via pulse-electric discharge consolidation with temperature gradient control H. Kimura and T. Satoh

355

Study on the composition graded cemented carbide/steel by spark plasma sintering A Ikegaya, K. Uchino, T. Miyagawa and H. Kaneta

361

Phase composition profile character of a functionally-graded AljTiOg/ZrOj-AljOa composite S. Pratapa, B.H. O'Connor and IM Low

367

The use of a functionally graded material in the manufacture of a graded permittivity element S. Watanabe, T. Ishikura, A. Tokumura, Y.Kim, N. Hayashi, Y. Uchida, S. Higa, D. Dykes and G. Touchard 373 1-7

Material Evaluation

Evaluation and modelling of the residual stresses generated on functionally graded materials - Two examples N. Cherradi, D. Delfosse and P. l\/loeckli Residual strains and stresses in an AljOg-Ni joint bonded with a composite interlayer: FEM predictions and experimental measurements B.H. Rabin, R L Williamson, H.A. Bruck, X.L. Wang, T.R. Watkins and D.R.Clarke Residual thermal stresses in functionally graded Ti-TiCx materials N. Frage, M.P. Dariel, U. Admon and A. Raveh

379

387 397

The effect of constituent and microstructure of composites on the residual thermal stress in TiC-NigAI FGMs J.H. Wang and LM Zhang 403 New application of FGMto identification of unknown multicomponent precipitates /. Itoh, H. Yamada, Y. Kojima, Y. Otoguro, H. Nakata and M. Matubara Evaluation of graded thermal barrier coating for gas turbine engine M Kawamura,Y. Matsuzaki, H. Hino and S. Okazaki Mechanical and electrical properties of multilayer composites of silicon carbide J. Hojo, F.Hongo, K. Kishi and S. Umebayashi

409 413 419

The effect of thermal shock on the thermal conductivity of a functionally graded material A J. Slifka, A Kumakawa, B.J. Filla, J.M. Phelps and N. Shimoda

425

Non-destructive evaluation of carbon fibre-reinforced structures using high frequency eddy current methods G. Mook, O. Koserand R. Lange

433

Thermal diffusivity measurement forSiC/C compositionally graded graphite materials J. Nakano, K. Fuji! and R. Yamada High-temperature ductility of TiC as evaluated by small punch testing and the effect of CrgCg additive L M Zhang, J.F. Li, R. Watanabe and T. Hirai Mechanical and thermal properties of PSZ/Ni-base superalloy composite S. Akama Processing-working stress unified analysis model and optimum design of ceramic-metal functionally graded materials P.C. Zhai, Q.J. Zhang and R.Z. Yuan

439

445 451

457

Evaluation test of C/C composites coated with SiC/C FGM, under simulated condition for aerospace application Y. Wakamatsu, T. Saito, F Ono, K. Ishida, T. Matsuzaki, O. Hamamura, Y.Sohda andYKude 463 Durability and high altitude performance tests of regeneratively cooled thrust engine made of ZrOg/Ni functionally graded materials Y. Kuroda, M. Tadano, A Moro, Y. Kawamata and N. Shimoda

469

XV

PART II 11-1

ENERGY CONVERSION, MATERIALS

ELECTRONIC AND ORGANIC

Thermoelectric Materials

Research on enhancement of thermoelectric figure of merit through functionally graded material processing technology in Japan T, Kajikawa

475

A design procedure of functionally graded thermoelectric materials J. Teraki and T. Hirano

483

Transport properties in multi-barrier systems Y. Nishio and T. Hirano

489

Theoretical estimation of thermoelectric figure of merit in sintered materials and proposal of grain-size-graded structures J. Yoshino

495

Computer design of thermoelectric functionally graded materials LI. Anatychuk and LA/. Vikhor

501

Anisotropic carrier scattering in n-type BijTejgsSeo 15 single crystal doped with HgBr^ I.J. Ohsugi, T. Kojima, H.T. Kaibe, M. Sakata and LA. Nishida

509

Percolation design of graded composite of powder metallurgically prepared SiGe and PbTe R. Watanabe, M. Miyajima, A. Kawasaki and H. Okamura

515

Design of multi-functionally graded structure of cylindrical Rl heat source for thermoelectric conversion system S. Amada, J. Terauchi and T. Senda

521

Fabrication of N-type polycrystalline Bi-Sb and their thermoelectric properties M Miyajima, G.G. Lee, A. Kawasaki and R. Watanabe

527

Development of functionally graded thermoelectric materials by PIES method A Yamamoto and T. Ohta

533

MIcrostructure and thermoelectric properties of p-type BiogSbigTea fabricated by hot pressing D.M. Lee, J.H. Seo, K. Park, I. Shiota and C.H. Lee

539

Microstructural and thermoelectric properties of hot-extruded p-type BiosSbigTeg J.H. Seo, D.M. Lee, K Park, J.H. Kim, I.A. Nishida and C.H. Lee •'• • Effect of dopants on thermoelectric properties and anisotropies for unidirectionally solidified n-BigTeg N. Abe, H. Kohri, I. Shiota and LA. Nishida Thermoelectric properties of arc-melted silicon borides L.D. Chen, T. Goto and T. Hirai

545

551 557

XVI

Graded thermoelectric materials by plasma spray forming J. Schilz, £ Muller, W.A. Kaysser, G. Langer, E. Lugscheider, G. Schiller and R.Henne • 563 Preparation of PbTe-FGM by joining melt-grown materials M Orihashi, Y, Noda, LD. Chen, Y.S. Kang, A. Moro and T. Hirai

569

Improvement and thermal stability of thermoelectric properties for n-type segmented PbTe S. Yoneda, H.T. Kaibe, T. Okumura, Y. Shinohara, Y Imai, LA. Nishida, T.Mochimaru, K. Takahashi, T. Noguchi and I. Shiota

575

Preparation and thermoelectric properties of IrSbg M Koshigoe, I. Shiota, Y. Shinohara, Y. Imai and LA. Nishida

581

p-n joining of melt-grown and sintered PbTe by plasma activated sintering Y.S. Kang, Y. Noda, LD. Chen, K. Kisara and M. Niino

587

Trial manufacture of functionally graded Si-Ge thermoelectric material T. Noguchi, K. Takahashi and T. Masuda

593

Microstructure and property of (Si-MoSigVSiGe thermoelectric converter unit J.S. Lin, K. Tanihata, Y. Miyamoto and H. Kido

599

Temperature dependence of the porosity controlled SiG/B4G+PSS thermoelectric properties K. Kato, A. Aruga, Y. Okamoto, J. Morimoto and T. Miyakawa Preparation of B4G-B system composites adding PSS and their thermoelectric properties A Aruga, K. Tsuneyoshi, Y. Okamoto and J. Morimoto

605

611

Joint of n-type PbTe with different carrier concentration and its thermoelectric properties Y. Imai, Y. Shinohara, LA. Nishida, M. Okamoto, Y. Isoda, T. Ohkoshi, T. Fujii, L Shiota and H.T. Kaibe 617 Effects of plasma treatment on thermoelectric properties of SigoGejo sintered alloys K. Kishimoto, Y. Nagamoto, T. Koyanagi and K. Matsubara Gontrol of temperature dependence of thermoelectric properties of manganese silicide by FGM approach T. Kajikawa, S. Suzuki, K. Shida and S. Sugihara Heat sensing device with thermoelectric film laid on insulated metal sheet T. Amano, N. Kamiya and S. Tokita 11-2

623

627 633

Thermionic IVIaterials

Recent developments in oxygenated thermionic converters J. L. Desplat

639

Development of refractory metal oxide collector materials and their thermionic converter performance R Fukuda, Y. Kasuga and K. Katoh

647

Thermionic properties and thermal stability of emitter with a (0001) oriented rhenium layer and graded structure M Katoh, R. Fukuda and T. Igarashi

655

Development of efficient thermionic energy converter T. Kato, K. Morimoto, K. Isogai, M. Kato, T. Fukushima and R. Fukuda

661

Radiation dose reduction by graded structures in the heat source of a ®°Sr radioisotope battery A Ohashi, K. Ueki and T. Senda

667

Output increase of thermionic energy converter due to the illumination of xenon short arc lamp Y. Shibahara and M. Kando

673

Hybrid mode concept of a thermionic converter with a FGM structured collector M Iwase and Y. Hirai

681

11-3

Electronic Materials

Thermoelectrically modulated/nanoscale multilayered gradient materials for application in the electromagnetic gun systems M A Otooni, J.F. Atkinson and LG. Brown Synthesis of In-Sb alloys by directional solidification in microgravity and normal gravity condition H. Minagawa, Y. Suzuki, K. Shimokawa, Y. Ueda, J. Nagao and J. Kawabata Full-colored zinc gallate phosphor with graded composition T. Endo, K. Uheda and H. Takizawa Synthesis and characterization of a model CuO/SnOg oxygen sensor P.J. Mailer, Z.S. Li, Q.L Guo

687

695 701 707

Fabrication of magnetic functionally graded material by martensitic transformation technique Y. Watanabe, Y. Nakamura and Y. Fukui

713

Characterization of single-crystalline Cu/Nb multilayer films by ion beam analysis S. Yamamoto, H. Naramoto, B. Tsuchiya and Y. Aoki

719

Enrichment of ^^Si by infrared laser irradiation T. Tanaka, I. Shiota, H. Suzuki and T. Noda

725

11-4

Natural, Organic and Intelligent Materials

Adaptive and functionally graded structure of bamboo S. Amada and N. Shimizu

731

XVlll

Learning about design of FGMsfrom intelligent modeling system in natural composites F.Nogata Development of the fire door with functionally graded wood H. Getto and S. Ishihara Elemental mapping of functionally graded dental implant in biocompatibility test F. Watari, A Yokoyama, F. Saso, M Uo, S. Ohkawa and T. Kawasaki Characteristics of epoxy-modified zirconium phosphate materials produced by an infiltration process AM Low, S. Yamaguchi, A. Nakahira and K. Niihara Preparation and properties of PVC/polymethacrylate graded blends by a dissolution - Diffusion method Y. Agari, M. Shimada, A. Ueda, T. Anan, R Nomura and Y. Kawasaki Preparation and properties of polyimide/Cu functionally graded material M Omori, A. Okubo, G.H. Kang and T. Hirai Smart functionally graded material without bending deformation J. Qiu, J. Tani and T. Soga

737 743 749

755

761 767 773

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

FGM research programs in Japan —from structural to functional uses Y.Miyamoto^ M.Niino^ and M.Koizumi'' ^ The Institute of Scientific and Industrial Research, Osaka University, Japan ^ National Aerospace Laboratories, Kakuda, Japan "^ Faculty of Science and Technology, Ryukoku University, Japan The FGM concept can be applied to various material fields for structural and functional uses. In Japan, several five-year programs have been conducted over the past ten years in order to develop the architecture of FGMs, and also to develop these materials for high temperature applications (e.g., components for the hypersonic space plane) and for functional applications (e.g., thermoelectric and thermionic converters). These programs are discussed with respect to the construction of FGM architecture and the future of FGMs. 1. APPLICATION OF THE FGM CONCEPT A functionally graded material (FGM) is a material in which the composition and structure gradually change resulting in a corresponding change in the properties of the material. This FGM concept can be applied to various materials for structural and functional uses. In order to create FGMs, the architecture of design, processing, and evaluation needs to be developed because no comprehensive study of such nonuniform materials has been carried out previously. The concept of integrating incompatible functions such as the refractoriness of ceramics and the toughness of metals with the relaxation of thermal stress, lead to a research project for the development of FGM architecture in 1987 [1]. In fact, it is possible to integrate a variety of dissimilar materials and properties if the thermal expansion mismatch or lattice mismatch can be relaxed and chemical compatibility can be maintained. Many applications exist that require high temperature resistance or thermal shock resistance, where the FGM concept can be applied. 2. THE DEVELOPMENT OF FGM ARCHITECTURE 2.1. For structural uses~the integration of refractoriness and toughness A five-year research program entitled "Fundamental Study on the Relaxation of Thermal Stress for High Temperature Materials by the Tailoring of Graded Structures" was established in 1987 with a total budget of 1,215 million yen under the auspice of the Science and Technology Agency. The goal was to develop the architecture of FGMs for structural uses and for high temperature components for the future hypersonic space plane. About 30 research organizations from national institutes, universities, and companies participated in the

Major Results of The FGM Program for 1987 -1991 CAD System: Inverse design model Selection of composition & microstructure Optimization of gradation Fuzzy function ,

Micromechanical Modeling: Correlation of graded microstructures & properties

Small Punch Test for Fracture energy

(3 Evaluation of Heat Shock: Xenon lamp irradiation. Burner Heating

immmm

r Stress Analysis by FEM J

Fractal & Percolation Theories: Quantitative analysis of gradation

Process Developments: CVD, PVD, PM, Plasma Spray, SHS, GalvanoForming, CVD/CVI, PM/CVD, SHS/HIP PS/GF

FGM Samples: Disk: SiC/C, PSZ/SUS, PSZ/Ni, AlN/SiC,TiC/Ni, Cr3C2/Ni, TiB2/Ni Nose cone: SiC/CC Rod: PSZ/Ni

Figure 1. Major results of the 1987-1991 FGM research program on the "Fundamental Study on the Relaxation of Thermal Stress for High Temperature Materials by the Tailoring of Graded Structures."

program as a member of one of three major groups: design, processing, or evaluation. Each investigation was coordinated for the purpose of developing the fundamental architecture of FGMs and their applications. Figure 1 illustrates the major results of the research program [2, 3]. For example, with respect to design and modeling, a CAD system using an inverse design model was developed that can produce an overall design architecture including selecting compositions and microstructures and optimizing the graded arrangement. Thermophysical parameters measured or calculated to minimize thermal stress both under process and service conditions were used for this optimization. A fuzzy function was used to combine different microstructures and properties smoothly, and a micromechanical approach to correlate graded microstructures and properties was established. Fractal and percolation theories were introduced for the quantitative analysis of the spatial change in graded microstructures, and FEM was used to model the distribution of internal stress. A number of processes were developed that use CVD, PVD, plasma spray, powder metallurgy, SHS, and galvanoforming. Several combined processes were also developed including CVD/CVI, PM/CVD, SHS/HIP, and plasma spray/galvanoforming. Various FGM samples were fabricated such as disks of SiC/C, AlN/SiC, PSZ/stainless steel, PSZ/Ni, TiC/Ni, Cr3C2/Ni, TiB2/Cu; nose cones of SiC/CC; and rods of PSZ/Ni. A small punch test was devised to evaluate the fracture energy of a thin FGM disk. Two methods were developed for the evaluation of thermal shock resistance up to 2000 K: irradiation by a strong xenon lamp and heating using an oxygen/hydrogen mixed-gas flame burner. Small combustion chambers for rocket engines made of SiC/CC by CVD/CVI and of Zr02/Ni by plasma spray/galvanoforming are undergoing combustion tests at the National Aerospace Laboratory. Although this program did not extend beyond fundamental research, it established the future direction for continuing FGM research worldwide. The FGM concept has been applied by several industries to a variety of products. To date, high performance cutting tools of TiCN/WC/Co, Ni FGM [4] and shaving blades of Al-Fe intermetallics/stainless steel FGM [5] have been commercialized. However, other commercial applications are still limited. 2.2 For functional uses ~ the direct conversion of thermal energy to electric energy Because the FGM concept was expected to be applicable to materials for functional uses as well as for structural applications, a new five-year project was initiated in 1993 with the aim of applying the FGM concept to the development of highly efficient thermionic and thermoelectric energy conversion materials. Both a themionic converter (TIC) and a thermoelectric converter (TEC) can produce electric power directly from thermal energy by the electron flow generated in space or in a solid under a high temperature differential. Figure 2 illustrates this ongoing program. In this Hybrid Direct Energy Conversion System, a TIC and a TEC are combined, and solar energy is used as the heat source to create a large temperature differential from ~2000K to ~300K. The design and optimization of the graded fields with respect both to the electronic and the elastic potential should lead to higher conversion efficiency with the relaxation of thermal stress. Thus the development of FGM architecture that would combine structural and functional properties is another goal of this program.

• C/C heat reservoir Mo radiation shield cyhnder TIC emitter Re graded coating TIC collector

SiGePbTeBi2Te3-

2.2.1. The design and processing of graded components for TICs and TECs In order to develop efficient and long lasting TICs and TECs, or combinations of these devices, an optimized system with lower heat loss and less degradation must be assembled using high performance TIC and TEC materials and devices. This will require solving various interface problems with respect to heat and carrier transportation, materials joining, thermal stress, electric contact, and insulation under extreme thermal conditions.

1) Graded C/C heat reservoir In order to achieve efficient heat accumulation and transfer from solar rays, Figure 2. A schematic illustrating the Hybrid Direct a composite FGM consisting of a 3-D Energy Conversion System. graded alignment of carbon fibers and pitch infiltration has been developed at Nippon Oil Company Ltd. [6]. Carbon fibers are highly anisotropic with respect to thermal conductivity along and across their length. Therefore, the graded alignment of fibers is designed to have a higher fiber density along the heat flux at the inner layer. A woven carbon fiber cup with a graded texture was infiltrated with pitch and hot isostatically pressed (HIP) to graphitize the pitch and densify the structure. Figure 3(a) shows the graded alignment of carbon fibers, and Figure 3(b) is a photo of the dense, graded C/C heat reservoir after HIPing. Solar rays are concentrated in this graded C/C heat reservoir by a large parabolic mirror, and the bottom and lateral sides are uniformly heated to 1680°C and 1380°C, respectively. The heat reservoir is covered with a radiat ion shield made of a highly polished Solar Rays

1380 x:

o ^^\\\\\\\

J680
(a) (b) Figure 3. A schematic of the graded ahgnment of carbon fibers in a heat reservoir (a), and a graded C/C heat reservoir (b).

cylinder of single crystal Mo (see Figure 1). A high heat flux of 1450°C can be transported from the bottom of the reservoir to the back surface of the TIC emitter electrode by heat radiation. 2) The graded TiC/MoAV/Re TIC emitter Titanium carbide is a promising material for the heat receiver of the TIC emitter because of its high melting point (3000''C) and its high emissivity (-0.9), which allows the efficient absorption of heat from a wide range of the solar spectrum at a high temperature. A graded coating of TiC/Mo with low thermal stress has been developed by using a double-gun plasma spray technique developed at the National Research Institute for Metals [7]. No cracks were observed in this graded coating after heating to 1800°C. The coating was formed on the back of the MoAV/Re emitter electrode [8]. W and Re were deposited on a Mo substrate by CVD. The compositionally graded layer formed through inter-diffusion of these elements at 2300°C by heat treatment. The W can act as a diffusion barrier for Re thereby creating a stable compositional gradation up to 2000°C. Figure 4 shows the graded cross section of this FGM emitter developed by Tokyo Tungsten Co. Ltd. The linear change of the thermal expansion coefficient for Re-W and the small change for W-Mo effectively relaxes the thermal stress.

Figure 4. Microstructural changes in the graded TiC/MoAV/Re emitter. The collector electrode was made at the Electrotechnical Laboratory by sputtering niobium oxide with a low work function of 1.38 eV on a Mo electrode [9]. A TIC device with a maximum output power of 8 W/cm^ obtained at an emitter-collector temperature differential of 1600°C-760°C and a Cs reservoir temperature of 330°C was assembled at Mitui Engineering & Ship Building Co.Ltd. [10]. 3) Graded PbTe thermoelectric material For TEC devices such as Bi3Te2 [11], PbTe [12], and SiGe [13], the theoretical calculation suggests that it is possible to improve the conversion efficiency by several percent by the gradation of dopants or compositions. Figure 5 compares the maximum output power of a graded PbTe having three different carrier concentrations (a-layer 3.51xlO^Vm^', b-layer 2.6xlO^Vm^ c-layer 2.26xlO^Vm^), fabricated at the National Research Institute of Metals, with that of homogeneous compositions of each layer. The PbTe with a Pbl2 dopant gradation was prepared by laminating powders with three different compositions of the dopant, followed by hot pressing. The output power was measured by holding the low temperature electrode at room temperature. It was shown that the FGM sample has a maximum power of 253 W/m for AT=486 K, which is 11% higher than the highest power of the non-graded sample with a layered composition.

0

100

200 300 400 500 Temperature difference A T/K

Figure 5. Effective maximum power as a function of the temperature difference for a three-layered FGM.

Figure 6. Photo of the graded symmetric MoSi2/ Al203/Ni/Al203/MoSi2 electrodes fabricated at Osaka University for a SiGe device.

4) A symmetrically graded MoSi2/Al203/Ni/Al203/MoSi2 electrode for a SiGe TEC In the case of the TEC device, MoSi2 is promising as a high temperature electrode for thermoelectric materials such as SiGe because it has a high melting point (2030°C), suitable resistivity (-LTxlO""^ Q-cm at 1000 °C), and excellent oxidation resistance. However, its mechanical toughness (3-4 MPa m^^^) is low. To make toughened electrodes, a symmetrically graded electrode of MoSi2/Al203/Ni/Al203/MoSi2 (shown in Figure 6) was fabricated at Osaka University by a combination of self propagating high temperature synthesis and HIP (SHS/HIP). This symmetrically graded structure produces a compressive residual stress of 100 MPa at the outer MoSi2 layer due to the thermal expansion mismatch between the outer and inner layers, resulting in reinforcement of the MoSi2 up to 5.8 MPa m^^^. The AI2O3 acts as a diffusion barrier against Ni at around 800 °C. A survey of experiments with Ni diffusion in an FGM structure at high temperatures suggests that at 800°C there will be little formation of a reaction layer for at least 10 years. This symmetric FGM electrode was joined to SiGe at Mitsubishi Heavy Industries Ltd. using interlayers of Ge powder and W and Zr -Ni foils at 1000°C [14]. 5) Graded AINAV radiative material For use of the TIC and TEC or their combined device in space, cooling by radiation is required. At Sumitomo Electric Industries, Ltd., a unique radiative material made of an AIN ceramic matrix composite containing a graded dispersion of fine W particles was fabricated by sintering at 1800°C. Dense AIN has a high thermal conductivity (-200 W/mK) and is transparent to wavelengths from the visible to the infrared. On the other hand, W has high emissivity (0.9) for the infrared region. Part of the heat flow conducted from the TEC to the AIN can be radiated by the W particles, as illustrated in Figure 7. A high emissivity of 0.97 was calculated and measured when a high concentration of W particles were dispersed near the interface with TEC, and the concentration was gradually decreased toward the opposite end.

As shown in Figure 8, the graded structure of AIN/W, which can satisfy both high thermal conductivity and high emissivity, has different characteristics compared with conventional materials [15]. In this program, further investigation will be carried out mainly on the following topics. 1) Modeling and evaluation of graded thermoelectric materials. 2) Synthesis of graded p and n type Bi3Te2, PbTe, and SiGe with higher conversion efficiency. 3) Fabrication of TEC segments and cascades with tough and low energy loss interfaces. 4) Fabrication of TICs with graded electrodes that have a conversion efficiency above 15%. 5) Stability of graded structures and properties at high temperatures.

Functionally Graded —T- Materials w

1.0 Dispersed particles: W Matrix: AIN Heat flow

-i

L_

^--4::::::::-::::::i»::::::::

0.8

t

0.6

a

Zr02 0 Y203 0

•

/

AI203 O ZrCO

00

Heat radiation

Si02 O

0.4

pq Al«

0.2

-Conduction-

MgOO Cu"

0.0

1 10 100 1000 Thermal conductivity , W • m"^ • K"^ Figure 7. Schematic of a highly radiative material Figure 8. The relation of thermal conductivity composed of an AIN matrix containing a graded and emissivity. dispersion of fine W particles. 3. THE FUTURE OF FGMs Although systematic research on the design, processing, and evaluation of FGMs has been carried out in Japan for the past ten years, it is only a first step for constructing the architecture for FGMs. Natural materials found in living organisms are composed of graded or nonuniform structures and textures. The organic combination of graded structures and properties can produce higher order functions than uniform structures. FGM architecture is expected to be a useful tool not only for integrating dissimilar materials and functions but also for creating synergistic effects in graded structures and properties like biomimetic and intelligent materials. Presently, FGM research is being actively conducted worldwide. New research programs were initiated in 1995 in Germany and in 1996 in Japan. In the German program 41 topics are being investigated in universities and national institutes with respect to processing and

modeling for developing FGM applications for the 21st century [16]. A new Japanese program is being promoted among universities on 67 topics relating to the chemistry and physics of FGMs. However, continuity and communication of individual and group research is of importance for developing FGM architecture. A new science and technology of non-homogeneous graded materials could result in the development of practical applications during the 21st century. To establish FGM architecture, the following research is needed: 1) Establishment of mathematical definitions and theories for graded structures and properties. 2) Development of a computer aided design and modelling system. 3) Development of cost effective processes. 4) Evaluation of graded properties. 5) Establishment of a data base for FGM architecture.

REFERENCES [I] T.Hirano, T.Yamada, J.Teraki, M.Niino and A.Kumakawa, Proc. 16th Int. Symp. on Space Tech. and Science, (1988) 375. [2] T.Hirai, Chapter 20,"Functional Gradient Materials", Mater. Sci. and Tech. vol.l7B, R.W.Chan, P.Hassen and E.J.Kramer(eds.), VCH, Weinheim, Germany( 1996)293. [3] Report on "Fundamental Study on Relaxation of Thermal Stress for High Temperature Materials by Tailoring the Graded Structure (Phase II, 1990-1991)", published by R&D Department of Science and Technology Agency (1992). [4] Pat. Japan, Application number: Hei7-179978 (1995). [5] NIKKEI Mechanical, No.469, December 11, (1995) 94. [6] Y. Kude and Y. Sohda, This symposium proceedings. [7] T. Fukushima, S. Kuroda and S. Kitahara, This symposium proceedings. [8] M. Katoh, R. Fukuda and T. Igarashi, This symposium proceedings. [9] R. Fukuda, Y. Kasuga and K. Katoh, This symposium proceedings. [10] T. Kato, K. Morimoto, K. Isogai, M. Kato, T. Fukushima and R. Fukuda, This symposium proceedings. [II] N. Abe, H. Kohri, I. Shiota and I. A. Nishida, This symposium proceedings. [12] Y.lmai, Y. Shinohara, I. A. Nishida, M. Okamoto, Y. Isoda, T. Ohkoshi and T. Fuji, This symposium proceedings. [13] T. Noguchi, K. Takahashi and T. Masuda, This symposium proceedings. [14] Y. Miyamoto, J. Lin and K. Tanihata, Y. Kang and M. Niino, S. Hirai and T. Umakoshi, Proc. Int. Symp. on FGMs, Ceramic Transactions, to be printed. [15] K. Sogabe, M. Tanaka, T. Miura and M. Tobioka, This symposium proceedings. [16] T. Rodel, This symposium proceedings.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

Research program on gradient materials in Germany J. Rodel and A. Neubrand Fachbereich Materialwissenschafl, Technische Hochschule Darmstadt, PetersenstraBe 23, 64287 Darmstadt, Germany The concept of the program on gradient materials, which was started in the fall of 1995, will be outlined. The program is structured in four subgroups, co-ordinating work in powder metallurgy, melt processing, coatings and functional materials. Since it is scheduled to run for 6 years, the early emphasis will be on processing, with some support in modelling and characterisation. There is no constraint on any group or combination of materials, but metal/ceramic combinations feature most prominent in the current arrangement. The prime constraint on accepting research into the core program results from the requirement that the gradient needs to determine the relevant properties of the gradient material and needs to extend over most of the active part of the component.

1. MOTIVATION Whereas in Japan, the research on functionally graded materials had gained a remarkable level already in the early 90's, activities in Germany were very restricted. German companies were holding a limited number of patents on manufacturing processes for FGM's and a few researchers at the German Aerospace Laboratory were developing FGM thermal barrier coatings. Although these efforts have to be acknowledged, their impact was limited and the knowledge and acceptance of gradient materials amongst German researchers and engineers was small. At the same time, a report of the Fraunhofer Society entitled „Technology at the beginning of the 21'* century" and the Delphi-Report of the German ministry for education and research both stated the potential of gradient materials in energy and transport technology, development of resources, medicine, food science and environmental protection. On the other hand, the reports asserted that the human capital and industry involvement in FGM technology required improvement in Germany. In this situation, a group of German scientists recognised the need for a national program as the complexity of a gradient material necessitates scientific collaboration of specialists from different fields such as metals, ceramics, characterisation and modelling. The broad horizontal integration across materials classes and disciplines in such a research program would also make it ideal for the education of young scientists. A program committee was formed, and a proposal was submitted to the German research society (DFG). Finally, in 1994 the senate of DFG decided to establish a research program on gradient materials to enhance and bundle research activities on FGM's in Germany.

10 2. STRUCTURE OF THE PROGRAM The German FGM program has a duration of 6 years. It intends to gather materials scientists, mechanical engineers, physicists and chemists from universities and research institutes. Its open structure allows to incorporate new research projects every year - which seems appropriate in a field where new ideas and potential applications are still emerging at a high rate. The program is not restricted to a certain class of materials or applications. Its intention is not to improve already existing technologies where gradients play a role such as case-hardened steels. Instead, the program supports development of new materials, in which the property gradient is essential for the function of the material. The gradient may serve to join two materials that would normally fail during production or service because of large differences in coefficients of thermal expansion and/or elastic modulus and/or poor adhesion. Alternatively, it may introduce a well-defined profile of a functional property that allows an optimum response of the component to an external field. The aim of the program is to improve the basic knowledge on the processing, properties and theory of gradient materials. Processing techniques allowing to produce continuous and multidimensional gradients for a variety of material combinations have preference within the program. Apart from already established production methods for FGM's relying on the addition of a gradation step to a conventional process (e.g. plasma-spraying), the development of special methods for the production of FGM's (e.g. using potential fields) is encouraged. The determination of the property profiles of gradient materials requires methods that have a sufficiently high spatial resolution - new methods offering such possibilities should be included in the program just as well as already existing techniques in order to be able to characterise the manufactured materials and to provide theoreticians with the information required for the prediction of FGM behaviour. Theoretical groups in the program have two tasks : First, they support the processing groups in improving their processing steps, e.g. drying or sintering. Second, they predict the FGM behaviour under load and calculate optimised property profiles for the FGM's produced by the program participants. Collaboration in the program is planned between participants with related topics and along the lines processing-properties-modelling. 3. STATE OF THE PROGRAM The German program on gradient materials started in November 1995. Already in the first year of the program there were 59 applications out of which 32 with a total funding of 4.1 Mio. DM were accepted. In the second year the program will have 41 participants and a budget of 4.5 Mio. DM. 26 projects cover the processing of gradient materials, out of which nine are powder techniques and seven coating techniques. 8 projects deal with theoretical aspects of gradient materials and seven projects are devoted to the characterisation and properties of gradient materials. In order to integrate such a large number of projects into the program and create links between the different participants, four discussion groups have been formed, which meet at regular intervals. The main topics of the groups are melt processing, sintering technology, coatings and thermomechanical aspects, and functional applications of gradient materials (Table 1). The melt processing group applies conventional casting techniques to the production of FGM's - the gradient is introduced by partial mixing or directed solidification of melts, sedimentation in melts or by the Vemeuil process. Graded fibre composites are produced by melt infiltration. The sintering technology group is dealing

11 with the problems of grading and densifying bodies using powder metallurgy. This includes the production of FGM's with various technologies including slip casting, centrifuging or hot extrusion. Sedimentation and sintering processes are treated theoretically and experimentally. The interest in the third group is focused on graded thermal barrier coatings and wear resistant coatings. Processing routes include PVD and spraying techniques. Experimental investigations of the behaviour of FGM's under static and cyclic thermal and mechanical loading is of major importance in this group. It is expected that modelling of the fracture mechanics, residual and applied stresses will help to improve the performance of the FGM's produced within the group. The last group gathers a number of participants which produce FGM's for specific functional applications. Emphasis hes on the production of materials for energy conversion and medicine. In contrast to Group I and II the goal of these projects is not the development of production technologies applicable to a large number of materials, but to achieve an optimum material for a specific application by improving production technology. Table 1 The German FGM program Group I Melt Processing

Group n Sintering Technology

Graded Single Crystals Directional Solidification Casting Technology Gas Pressure Infiltration

Modelling of Sintering Graded Porosity Slip Casting Sedimentation Powder Metallurgy

Group n i Coatings and Thermomechanical Modelling

Group rv Functional Materials

EB-PVD Thermal Barrier Coatings Chemical Vapour Condensation Method Wear Resistant Coatings Measurement of Residual Stresses Laser Thermal Shock Test Cyclic Loading of Gradient Materials Fracture Mechanical Modelling

Thermoelectric materials Solid Oxygen Fuel Cells Graded BaTiOs Biocompatible Composites Medical Implants Acoustic Characterisation

12

Figure 1

Important topics and links in the gradient materials program

13 At present 60 links between participants exist - a significant number of which has been established during the group meetings. Additional stimulation of collaboration between members of different groups is expected from yearly meetings of the participants, where the progress of the projects is presented. It is already obvious in the early state of the program that a number of projects in the characterisation, theoretical and testing field has gained an important support character for all participants of the program. Of particular importance in this context are the theoretical and experimental determination of stresses in FGM's, and the measurement of elastic and thermal property profiles. Any effort to describe the interaction between the participants must remain rudimentary - nevertheless Fig.l shows an attempt to sketch the most important links within the program.

4. OUTLOOK At the end of the first year of the program, the task of structuring the program is almost accomplished. Whereasfijndinga wide scope of methods and appUcations in the early phase of the program should awake academic interest in gradient materials, the second phase will bring a reduction of the number of projects. Nevertheless, a number of projects of high scientific quality and relevance as well as a few highly innovative works will be funded beyond the year 2000. However, the expected reduction in public funding should not lead to a general reduction in FGM research - it will only be transferred partly to industrial laboratories (Fig.2). A central task for the co-ordinators and participants of the program is therefore to stimulate industrial interest in FGM's. At present, industrial interest is only moderate as commercial applications of FGM's are still rare - the obstacles towards introduction of FGM's into components being manifold : Most production processes for FGM's are batch processes and difficult to scale up for mass production at a reasonable cost. With many processes, the production of components with complicated geometry and/or multidimensional gradation is impossible. The consolidation process of FGM's is frequently very sophisticated and considerable efforts are necessary to adapt it to new material combinations or shapes. Further, FGM's are presently not widely accepted by engineers because the effect of a gradation on the performance of a component is not known a priori. The above statements lead to a number of challenges for the participants in the German FGM program : 1. 2. 3. 4. 5. 6.

Elaborate processes suitable for mass production Ensure that the cost of the FGM does not interfere with the application in mind Prefer methods capable of near-net shaping Develop methods applicable to different material combinations without major changes Promote simple rules where a gradient will always be advantageous Determine a number of favourable gradation profiles for some important applications

Non-scientific efforts to awaken industrial interest are also made - a brochure in German language has been composed and distributed to companies manufacturing products where incorporation of FGM's may be of interest. Industrial participation in the yearly meetings is encouraged and a number of company representatives took part already in the first meeting. If a number of FGM's with application potential evolves from the program, it is expected that

14 the industrial interest will further increase within the next years opening possibilities for collaborations between participants and companies.

Industrial Interest Funding level by DFG awakening academic interest

increasing Industrial promise

1995

1998

2001

Figure 2 Expected development of the gradient materials program 5 CONCLUSION The research program on gradient materials in Germany had a good start. At present a broad spectrum of processes and applications is funded, and many groups report first interesting results a part of which is presented in these proceedings. The future will bring a reinforcement of industrial participation and a selection of processing methods and applications of technical relevance. It is believed that the basic work on gradient materials carried out in the framework of the program will result in an enhanced interest for FGM in research divisions of companies thus promoting the application of FGM's. If the program is successful, it is expected that German companies will produce components containing FGM's in the first decade of the next century.

ACKNOWLEDGEMENT The authors would like to thank the Deutsche Forschungsgemeinschaft (DFG) for establishing the program on gradient materials and for personalfinancialsupport.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 1997 Elsevier Science B.V.

15

Lessons learnt in 7 years of FGM research at Lausanne B. Ilschner Dept. of Materials, Swiss Fed. Inst, of Technology Lausanne

General experience gained during the research period 1988 - 1955 underlay the difficulties to make FGM economically visible ("marketplace"). The conclusion is the suggestion to focus future research on graded materials on 1) fundamental aspects ("Functions of Gradients in Materials") and 2) an effort to develop fast, low-cost, highly reliable fabrication routes for technically relevant components. Introduction In this paper, a number of reflections and conclusions are formulated which result from research work on functionally graded materials as performed by a small group at the Swiss Federal Institute of Technology at Lausanne, mainly in the period 1988 to 1995. It is hoped that these observations, after critical review by the "FGM community", may contribute to the planning and accomplishment of future research activities within this field. In order to explain these ideas, it is proposed to start with a retrospect. "Thermally graded materials" The origin of the author's preoccupation with graded materials was simple scientific curiosity, without an immediate application in mind: after 15 years of work in the field of high-temperature creep, performed with a group of young collaborators at the university of Eriangen-Nurnberg (Germany) between 1965 and 1980, discussion of the behavior of gas-cooled turbine blades led to the idea of investigating into the creep of axially symmetric, tubular model specimens [1]. The outer part of this type of specimen was externally heated to high temperature (up to 0.7 T J , while the central part was cooled by pressurized air, generating a radially symmetric temperature field T(r). These specimens were then subjected to a constant axial tensile stress, a, and the elongation e between markers on the specimen surface was measured optically. The corresponding constitutive law for steady state contains elements named after Arrhenius (T-dependence) and Norton (a-dependence), both of them being strong functions of their parameters: de/dt = C a" exp(-Qc/RT). Even a small temperature difference would therefore produce an appreciable difference in the creep rate for a given applied stress, if the specimen was to be considered as a set of independent coaxial shells. This is equivalent to the

16 Statement that the inner shells are "harder" than the outer ones, since TQ > Tj. As a matter of fact, the setup of the experiment imposes a constraint inasmuch as all coaxial shells have to have identical elongation at any moment, independent of their respective temperature T(r). One part of this elongation is due to the thermal expansion, the other part to the creep strain. We consider the material to be in a visco-elastic state. A transient stress distribution will therefore occur after each change of the applied stress and/or temperature profile. Only very small local deformations and, thus, short times are necessary to adjust local stresses to the general continuity condition. After the transition, the whole specimen will creep in tension under the action of a radial distribution of axial stresses a(r) which assures, respecting the creep rate equation, an equal creep rate for the whole specimen. From the viewpoint of continuum mechanics, a chemically homogeneous specimen with a radial temperature gradient is indeed a "graded material" inasmuch as each coaxial shell offers a different resistance to the applied stress and has a different time constant for relaxation. We may speak of a "thermally graded material". This work led to the idea of adapting the alloy composition to the temperature profile; this must be an iterative process, because the radial composition profile implies an inhomogeneity of thermal conductivity and is thus influencing the primary temperature profile. As a next step, it was felt that the mechanical behavior of a compositionally graded material at homogeneous temperature should be studied. Consequently, after taking up a new academic position in Lausanne, the author filed a research proposal with the Swiss National Funds [2]. This project was, incidentally, not approved in the first time, but later granted thanks to an initiative taken by Prof. V. Franzen, who at that time was the director of the National Research Program on "Materials for the Demands of Tomorrow". This was the first project in Lausanne on what later became known as FGM. The funding was modest and permitted to hire just one PhD. student, who was to be D. Delfosse, joining the lab on July 1,1985. Insights: The Japanese Model The fact that the above-mentioned Ph.D. project was part of a program on materials "for the demands of tomorrow" obliged us to go beyond the original guideline of obtaining valid scientific results and to seriously consider the potential for technical applications. Our primary concern was, however, to develop an experimental method for making fully dense continuously graded specimens in a sufficiently large number to allow for subsequent study of composition profiles, microstructures, and mechanical behavior. Although a number of alternatives were briefly studied, the process of "centrifugal powder metallurgy" as originally proposed in [2] emerged as the most efficient and reliable one. It was first published in [3,4]. At this stage, Prof. W. Bunk, at that time director of the materials division of the German Aeronautical Research establishment near Cologne (now DLR) established close contracts with the pioneering FGM research activities in Japan. These were centered at the National Aerospace Research Center (with Dr. Niino) and patronized by the Japanese Society for Non-Traditional

17 Technology. In the framework of these early contacts, a group of scientists from Japan under the direction of Prof. Koizumi visited Germany in 1988 and came subsequently also to the author's laboratory at EPFL Lausanne. This encounter proved to be of great importance for the future work of this group, The first insight gained from this visit was that the Japanese colleagues had a real national program, had a goal (at that time: thermal barrier coatings for future commercial space shuttles); they had generous funding, industrial resonance, an efficient organization - and all of this was practically nonexistent in Europe (and the rest of the world). One might illustrate this situation by depicting the Japanese FGM-program as a strong tree, and the other research activities in graded and layered structures as individual flowers distributed all over the world. By and large, this picture is still valid today. One exception - also going back to an initiative of W. Bunk - may be seen in the priority program on graded materials which has been launched in Germany in 1995 [5]. The Impact due to the discussion with the Japanese colleagues during their visit to Lausanne was reinforced by the insights which the present author gained during his participation in the first international conference on FGM in Sendai, 1990. It became clear that in order to achieve a visible result it was necessary to go beyond the limits of the traditional European university research style ("one topic - one thesis - one student"), which had, by the way, already been abandoned In other fields like nuclear physics or semiconductor research. Thus, in spite of serious difficulties to obtain funding, a small group of highly motivated young researchers [6] could be formed, which was since 1991 coordinated by N. Cherradi [7,8], who also made a very important contribution to the visibility of the FGM concept in his capacity as secretary general of the 3rd International. Symposium at Lausanne, 1994 [9]. On the other hand, the attempt to bring together a FGM Working Group on a national scale in Switzerland was not met by success. In retrospective, it appears that 2 necessary conditions for such a plan could not be established: a) to find a key person ("locomotive") of high reputation and sufficient political influence in the national community, being able to devote a major part of his energy and time to overcome the "activation barriers" during the incubation phase, and b) a fair amount of "seed money" which would enable the pioneer research group to produce preliminary results on which to base further funding applications. Again, both of these conditions were fulfilled in Japan! Functionally Graded Materials on the Marketplace During the 7-year period of FGM research in Lausanne, several other lessons had to be learnt. Among these is one which could already be sensed while working on the "thermally graded materials" project: The necessity of linking the experimental and microstructural aspects of graded (or layered) materials to the methods and results of continuum mechanics. Clearly, each inhomogeneous material is subjected to complex multi-axial stresses originating from the local differences in basic properties, in particular Young's modulus, the limit of elasticity, and the coefficient of thermal expansion. In this

18 field, it is therefore particularly important to coordinate expertise from both communities, such as demonstrated in a workshop in Davos organized in 1995 [10]. Probably the most important lesson which had to be learnt came in recent years only, and it came from outside of our community. As researchers and academic teachers in the field of materials we have to face an embarrassing fact: Materials and manufacturing actually are not in a dynamic growth phase comparable to the period 1960 to 1980. The so-called New Materials, in spite of the fascination originating in their often astonishing mechanical or electromagnetic behavior, do not find It easy to present themselves as an attractive value on the "marketplace". Too many of this kind, with sometimes exotic structures and compositions, have been announced with much ado, without finally living up to (exaggerated) expectations. The actual situation on the marketplace of our industrial society appears to be such that improved property values are no more considered to be a natural justification for Increased prices; this holds except for the field of telecommunication and for some "niche" applications which are not characteristic for the general status of materials science and engineering. Quite generally. Industrial leaders adhere more and more to a philosophy which states that the shareholder value - being their primary responsibility - is not essentially increased by research into new materials and technologies, but rather by skillful use of existing technology in combination with computerized design strategies and "lean assembly". The decisive battles, they say, are won or lost by commercial or financial moves. In parallel, governments (or parliaments ) are less and less Inclined to spend public money in scientific research as a "culture". This tendency is not likely to change in the next 10 years. As all sectors of our society, materials science and engineering must continuously discuss, redefine and justify its aims and ways. This general analysis may lead us to the following insight: If functionally graded materials, or FGM, are advertised as a wonderful new class of materials for the 21st century, they risk to be marginalized (as others before), except if they yield really spectacular results within the next few years. This is not impossible, but it Is not very likely either that this will happen. It has to be admitted that at the present stage of knowledge, graded materials with designed functions are too difficult and too expensive to make so that they cannot be produced in large quantities for industrial use; moreover, the means to ascertain their quality, reproducibility, reliability and lifetime by accepted standards with corresponding testing procedures are insufficient. There are no accepted design rules, and the community of design engineers (conservative as it is obliged to be) has little or no knowledge at all of the advantages, problems and limits of gradients in solid materials. The conclusion which we derived from this lesson at Lausanne is that the aforementioned limiting conditions reflect the reality and that a twofold strategy appears to be appropriate: I) To encourage basic studies of the physical, chemical and mechanical behavior of graded materials in general, of their cross

19 links with microstructure, and of the mechanisms likely to control possible fabrication routes. II) To consider the above-mentioned shortcomings as a challenge to resolve the associated technical problems by intelligent engineering, taking advantage of all the impressive knowledge on FGM which has been accumulated worldwide. In the following two sections, these two parallel strategies will be more closely described and discussed. Beyond FGM: The Function of Gradients in l\^ateriais There is no doubt that it can be of great advantage to design technical components using different materials for load-bearing, surface-protecting, electromagnetic and decorative functions. Fiber reinforcement and the whole surface treatment as well as joining technologies are excellent examples. Likewise, many scientific and manufacturing arguments can be brought fon/vard in favor of graded transitions inside such multi-material components, instead of abrupt property changes. The FGM concept has provided a new quality of understanding of these phenomena. In particular, it has demonstrated that compositional or microstructural gradients can not only serve to avoid undesirable effects (such as tensile stress concentrations) but can also serve to generate unique positive functions: focusing light in fibers, channeling heat in computer chips, implant-tissue transitions in biomedical engineering. Many others can be conceived. In a general way, the FGM concept has taught how to optimize concentration profiles. "Opening the lens" towards a broad view on "The Function of Gradients in Materials" may thus lead beyond the current scope of FGM as a special class of materials. The topic of gradients (and also multilayers) in solid materials presents itself as a promising perspective, very timely after a century of practical exercises and after 10 years of intensive work on FGM. The whole field Is urgently needing a coherent scientific infrastructure In all its chemical, electronic, mechanical aspects. Obviously, such a methodical approach envisages and enables applications in many traditional and non-traditional fields. Where a scientific background is in demand, there is no reason to shy away from topics such as welding and brazing, segregations after casting of liquid alloys, diffusion controlled hardening of steel surfaces and glasses; the analysis needs extension into polymer systems and natural materials such as wood, bone, teeth and shells. The Engineering Chaiienge: Fast, Cheap, Reiiabie - Conclusion The present shortcomings of graded materials from an engineering point of view have been listed above. Real progress can be achieved only if the following tasks are being fulfilled: Define applications which appear adapted for gradient solutions and establish a complete list of properties needed for their satisfactory function.

20 Specify candidate systems of graded materials/structures to comply with as many as possible (If not all) of the properties in demand. Evaluate feasible fabrication routes which yield either a semifinished product (strip, wire, etc.) or even near-net shaped components. Assess the materials systems and fabrication routes envisaged with respect to quality, reproducibility, available equipment and cost per piece. Define quantitative criteria for reliability (or admissible scatter) and durability (in terms of total service time or cycle number to failure). Find or design testing methods (preferably non-destructive, and as close as possible to currently standardized methods) which appear suitable to control the validity of the above-mentioned criteria. Evaluate possible environmental hazards as well as problems related to recycling. This package of tasks contains a considerable number of problems which belong to the modern field of manufacturing science. On the other hand, manufacturing science has not yet dealt at all with graded materials. So there is an "open sky" before us, and a rich and realistic source of motivation for young scientists and engineers to join the field which has been opened by the introduction of the FGM concept. Acknowledgment The author gratefully acknowledges financial support from the Swiss National Funds and the Priority Program on Materials Research (financed by BSFIT). Moreover, he thanks his collaborators for their successful work and many colleagues In the international FGM community for enlightening discussions. References [1] U.Engel, Z. Werkstofftech. 10 (1979) 243-248, see also U.Engel, B. Ilschner, Z. Werkstofftech. 17 (1986) 299-307 [2] B. Ilschner, Projet No. 4.834-(1985) du Fonds National Suisse [3] B. Ilschner, in: M. Yamanouchi, M. Koizumi, T. Hirai, I. Shiota (Eds.), Proc. 1st Internatl. Symposium on FGM, Tokyo 1990, pp. 101-106 [4] B. Ilschner, D. Delfosse, H.U. Kuenzi, Acta metall.mater. 40 (1992) 2219-2224 [5] J. Roedel, D. Neubrandt, Proc. 4th Internatl. Symposium on FGM, Tsukuba 1996 [6] The group consisted (in alphabetical order) of K. Barthel, M. Blumm, D. Delfosse, N. Desmonts, P. Li, X. Ding, K. Dollmeier, M. Probst-Hein, W. Thiele; important advice in theoretical and experimental questions is due to M. Cans, H.U. Kuenzi, and N. Merk. Most of their work hasa been published in the Proc. 3rd Internatl. Symposium FGM, Lausanne 1994. [7] N. Cherradi , A.Kawasaki, M.Gasik, Composite Engineering. Current Trends in Composites Research, Vol.4 (1994) 883-894 [8] N. Cherradi, D.Delfosse, B.Ilschner, A. Kawasaki, Rev. de MetallurgieCIT (1996) 185-196 [9] S. Suresh, F. Needleman (Eds.): Mechanics and Physics of Layered and Graded Materials (Proc. Of an Engineering Foundation Conference Davos 1995, Special Issue of J. of the Mechanics and Physics of Solids

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

21

Local fields in functionally graded materials Yevgen D. Bilotsky^ and Michael M. Gasik^^ ^Institute of Physics, 252650 Kyiv, Ukraine ^Helsinki University of Technology, 02150 Espoo, Finland

The structure and basic properties of the FGMs of different systems have been studied theoretically. There is a need for theoretical basis, which would allow to describe the structure and properties of materials from first principles, while remaining simple and easy to use. In this work, such theoretical principles are suggested and their applications to heat flow and respective stress fields in FGMs are considered as example. These principles can be spread out for other materials.

1. INTRODUCTION The functionally graded materials (FGMs) are characterised by a non-linear 3D-distribution of phases and corresponding properties [1,2]. They are distinguished from isotropic materials by gradients of composition, phase distribution, porosity, texture, and related properties (hardness, density, resistance, thermal conductivity. Young's modulus, etc.) [3-5]. The FGM is characterised not only by the presence and appearance of compositional or other gradients but also by the sophisticated behaviour of FGM component in comparison with conventional (macroscopically uniform) materials. In the simplest case, the structure of a material is represented or replaced by the model-like system of a matrix with embedded particles or grains. For such composites the microstructural fields are assumed to be homogeneous, whereas for FGMs they are heterogeneous. Due to the gradients in FGMs, the "normal" approximations and models, used for traditional composites, are not directly applicable to FGM. The situation becomes even more complicated, when an FGM has gradients on several levels, i.e. macro-, micro- and nano-scale, where defects such as vacancies and dislocations start to play an important role in the transfer processes and mechanical behaviour of the specimen. The main method actually used is based on the finite element approach (FEM) and its variations. Most numerical schemes are bases on discrete distribution models, not all of them taking into account possible non-isotropic phase distribution. When this distribution is not isotropic but complies with a certain law, the same observation of measurement process yields This work was partially supported by Technology Development Centre of Finland (TEKES) and the Commission of European Communities (COST-503 project)

22

different results depending on the way the sample is placed. It is obvious that the resources necessary to define, conduct, and interpret such an analysis, are prohibitive for complex structures [4]. Another opportunity would be in a modelling of the FGM structure and in a deducing of "structure-property" relationships, e.g. as micromechanical model. This, however, is limited to simplified structures with quite a lot of assumptions. The explicit description of any material (not only FGM) from the first principles can be obtained from its consideration over an arbitrary domain, where every inclusion, defect, anisotropy, etc., is precisely taken into account [6,7,15]. This task could be solved theoretically, but its application would be certainly useless in practice, since it will involve a vast number of calculations and measurements. In this connection, there is a need for theoretical basis, which would allow to describe the structure and properties of materials from first principles, while remaining simple and easy to use. In this work, such theoretical principles are suggested and their applications to FGM heat flow and respective mechanical stresses are considered as example. These principles, however, can be spread out for other materials.

2. THEORETICAL BASIS FOR FGMS A comparative analysis of such structures, involving the evaluation of their "effective" properties, e.g. with a "pointwise" homogenisation, with some micromechanical models [4], as well as numerical methods (combined models, FEM calculations, etc.) was reported elsewhere [4,6,7]. In this respect, the following issues might be formulated for an "ideal" theory that is able to describe and to solve these problems: (i) it should be a first-principle theory, involving as less as possible or none fitting assumptions or parameters, e.g. do not require an artificial representative volume element or grid introduction; (ii) this theory should be easily applicable for any material with the structure of any complexity; (iii) the solution procedure should be rather fast and reliable, and be asymptotically free (i.e. all asymptotic cases have to have finite solutions without singularities); the calculations themselves should be "error-resistant" (e.g., a small error should be compensated on the next iteration) [6,12]. This kind of a theoretical approach could be based on the local field analysis [6,7,10,15]. The external and local fields in materials can alter in a significant way such processes as the dislocations motion, solid-state reactions kinetics, sublimation, oxidation, etc. In many of the reported findings, the attempts were made to provide basic explanations for the experimental observations, but in general they were failed to give a consistent picture of the role of the fields in the various processes [9]. For instance, one of the main problem for polycrystalline specimens is in high non-linearity and inhomogeneity of the fields between the grains and near the defects. Resulting singularities usually do not allow the differential equations to be solved numerically. 2.1. Local stress and strain fields in FGMs Let's consider first a defect-free grain of one phase in a two-phase composite. The state of this crystal can be described by the equation of motion of elastic media [11]:

23

where vector f describes the density of the volume forces, applied to crystal, p is the crystal density, ui are components of the distortion (shift) vector

= l(^i^k-^^k^i) \^Xk

(2)

^^ij

and tensor Oj^i is bound with strains e^-^ by Hooke's lawCJ-j^ — A-^^^ £^^ . Let's consider now a crystal with a dislocation. In this crystal a single-valued vector of elastic shift u can be always introduced, where function u(r) will have a leap b on the surface SQ, laying on the dislocation loop or interface D:

5u = u^ - u ~ = b

(3)

where superscripts "+" and "-" refer to values of u(r) on upper and lower side of Sj) respectively. It is important that the same form of this equation is valid for a leap of u(r) on the grain boundaries and well as interfaces in the solid. In the latter case (3) does not specify a general appearance of the leap over S, whereas strain 8^-^ keeps its continuity and remains differentiable. Thus (3) transforms to

<5u = u^-u" = b + [Qr],

(4)

where b - translation vector, Q - rotation vector, and r - radius-vector of a point on surface S. When Q = 0, translation vector coincides with the Burgers vector. In a scalar elastic field, the distortion tensor can be substituted by vector h and stress tensor can be replaced by vector a : \i-grad

u ; diva

= -/

; a

= Gh .

(5)

These equations along with the Hooke's law constitute full system of equilibrium relationships for an elastic media with dislocations and interfaces. Let's suppose now that the piece of a two-phase material (grains of A within matrix of B) is assembled from these isotropic domains, where domain boundaries are some arbitrary distributed interfaces. For domains outside the A-phase grain and with conditions (3-5), it follows div u = 0. Thus, substituting distortion vector by u = UQ + grad cp , the respective Laplace equation for the distortion potential cp will be A(p(r) = 0. Consider now a jrain of phase A, which has several border interfaces with other A grains as well as with B-matrix; and the internal volume of this grain will be a positive direction (Fig.l). The picture on Fig.l is taken from the real structure of W-Cu infiltrated composite with a graded distribution of diamonds [4]. The equilibrium requirements stipulate that distortion potential on positive (inside) and negative (outside) sides of the interfaces must be coupled as follows, so far the material is not a subject of an application of external forces:

24

(p,(r)l = (p_ir)l /':s.^\ 3(p

l^nJ-

=

(6)

E(^ y^^JA

(7)

for every A grain in the composite, where n is outside normal vector to surface S, and ^ is a couphng coefficient.

Figure 1. Model of A-B FGM structure (see definitions in the text). Points schematically represent potential for positive and negative sides of interface SD. Here ^ is assumed to be the same for each interface of A grain, although its two principal values (for A-A and A-B interfaces) could be taken into consideration already in this simple approach. Solution of the Laplace equation A(p(r) = 0 in this case will have form:

^(r)dS, cpfrj = -u.r4- X 11 , I .

(8)

„„•' r-r -a

where analogue to charge surface density x(r) should satisfy the following integral equation:

1-^ f;?/^M V

^frj + j ; ^ J T f r ' j X

cos((r-r'-Si)'n,)dS, 27lj r - r - a

_

1 1-^

2711+^ u n.

(9)

Here n^ is outside normal vector in the point r, and the graded structure inhomogeneity is introduced by vector a. In general, a represents a distance between the centres of two A grains (Fig.l), where sum is taken over all existing grains (a(r) G [-©O; CO]). Thus, FGM structure parameters as well as grains topology and possible gradient of particle size distribution, are

25

taking into account. For homogeneous (non-FGM) materials the value of a could be simply taken equal to mean distance between A grains. The most essential feature of this approach is in exclusion of the singularities of the derivatives from consideration, since there are no singularities in the local domain by definition, and the integration removes singularities between the domains. These equations, however, cannot be apply for materials with very small grain sizes (20-50 nm and less), because these elasticity relationships (3-5) are not valid on atomic distances. 2.2. Dynamics of local Helds in FGMs Now let there be within a material a scalar field T(XQ), where XQ are space coordinates (|i =1..3). It could be temperature, electrostatic potential, concentration, etc. [6,8,11]. This field leads to the some form of transfer (heat, mass, charge, etc.) depending of its nature. In its turn, this transfer (flow) is one of the forms of motion, which must satisfy the principle of "least action" (61= 0): I = ^LdQ.

; dQ. = dh

(10)

where L(T,VjnT) is the Lagrangian density, and integration is taken over the body domain Q. The Lagrangian density L and the functional / must be invariant under any symmetry transformation (translational invariance in time and in space, and rotational invariance) in order to preserve the conservation laws (Noether theorem): T-> T + 5 T

,

8 T = Qe ,

(11)

where 8 is an infinitesimal parameter, and Q is a constant. The Lagrangian density itself should be also invariant (8L = 0):

^ST^4^ST^0.

(12)

After substitution of equations, the following form of (12) can be obtained [13]: dL dL d dT dx^ v^^.y

^^6r

dT^

0.

(13)

This equation is a form of the equation of motion and it virtually describes any variations and interactions. The Lagrangian L, which is generally expressed as difference between potential and kinetic energy [8,11], should be substituted by a respective expression, depending on the kind of problem considered (mechanical deformation, heat flow, diffiision, electrical conduction, etc.) These equations are valid over any domain of material, which does not contain inhomogeneties (inclusions, defects, dislocations, grain boundaries, etc.). Let's suppose now that a piece of material is assembled from isotropic domains, where domain boundaries are

26 arbitrarily oriented, but gradually distributed. The equations above are still valid in any domain, but they cannot be directly applied to the material due to singularities of derivatives on interfaces. In this case, the infinitesimal parameters e will no longer be constants but functions e(x), and the transformation (11) will be T ^ T+8T

,

8T=Q8fx^j

,

(14)

and the in variance of L will not hold anymore:

For retaining of the invariance of (12), a new field A(x) has to be introduced to cancel the right-hand side of (13). New Lagrangian L'(T, V„T, A) and the new considered transformation 5T(x) =E(x)Q,

5A = UAe(x)-tC^^

(16)

where U is a constant, will ensure that the functional will be invariant under transformation (8). Here A^ = C~^ A and the elements of the matrix C and its inverse C"^ here, and in the transformation (8), are [14]: C'C;' = 8 , C;' C" = 5; . In the case of a non-homogeneous, non-isotropic FGM, the gauges A(x) are considered to be uniquely determined by the structure of the material. Proceeding from the single domain to the bulk component, the non-steady field T could be evaluated from the invariance of L'(T, V QT, A) over the entire component by integration of the particular local field equations, taking into account coupling equations deduced above (6-7). This transformation gives new system of equations for any field distribution in a FGM of any complex structure from the single phenomenological point of view [6,15], where fields in domains and between them are considered within the same equations. Any external field, applied to the material with a certain temperature distribution, will not change the form of equations, since they will be t^en into account within the Lagrangian.

3. APPLICATION OF LOCAL FIELDS THEORY TO FGM With the theoretical principles above, it is possible to formulate their applications to the problems solution in FGMs. The sequence of the constructing of the systems of equations could be presented as a combination of the following stages: (i) obtaining structural information and construction of the "internal" fields Q, (ii) introduction of gauge fields A and description of the Lagrangian(s) L', (iii) introduction of "external" field(s) and analysis of the gauge and L', (iv) integration of the equadons and solution for unknown variables or fields. It is theoretically possible to determine the local stresses on every grain of one phase, depending on its distribution, structure, interfaces, etc. The exact description of the solution method is possible to formulate for particular cases, depending on the conditions applied

27 [6,15]. The following practical applications could be solved with this approach: (i) determination of the mechanical fields in FGMs and components under specific conditions (temperature, force and deformation); (ii) solution of the heat transfer problem in FGMs with an arbitrary defects distribution; (iii) prediction of the properties of materials and their behaviour in different fields (gravitation, electrostatic, electromagnetic, etc.). Although the equations shown do not often contain temperature explicitly, it may be introduced with the equation of motion, based on that thermal loading of a FGM component will be accompanied by propagation of elastic compression waves. This shows that interaction of mechanical and thermal factors in FGMs is one of the key issues of designing of reliable FGM components [6,12,15].

4. CONCLUSIONS The results of this work reveal that the general problems, such as "structure-properties" relationships, materials behaviour in different fields, etc. in FGMs, could be in principle solved through the construction and proper solution of integral equations in a scalar form for particular cases. Local fields of stress and strain, temperature, etc. as well as non-steady problems were considered. The approach suggested could be extended on fields of any nature, such as mechanical, concentrational, etc., whereas the form of equations and the method of their solution remain invariant to the kind of problem.

REFERENCES 1. M.Koizumi, Ceram. Trans.: Functionally Gradient Materials, Ed. B.Holt e.a., ACerS., Ohio, 34(1993)3-10. 2. M.Gasik and K.Lilius, Comp. Mater. Sci., 3 (1994) 41-49. 3. N.Cherradi, A.Kawasaki, and M.Gasik, Compos. Eng. 4 (1994) No. 5, 883-894. 4. M.Gasik, Acta Polytech. Scand., Ch 226 (1995) 73 p. 5. M.Sasaki and T.Hirai, J. Ceram. Soc. Jap. 99 (1991) 970-980. 6. Y.Bilotsky and M.Gasik, Compos. Eng., in special issue "Use of Composites in MultiPhased and Functionally Graded Materials", (1996). 7. Y.Bilotsky, B.Lev and P.Tomchuk, Proc. 2nd Int. Works. Non-linear Turbul. Proces. Phys., 1983, Kiev, Ukraine, Ed. R.Z.Sagdeev, Gordon & Berch Publ. Co., NY, 7 (1984) 701-704. 8. L.D.Landau and L.M.Lifschitz, Theory of elasticity. Compr. Cours. Phys., 7 (1986) 187 p. 9. Z.Munir and H.Schmalzried, J. Mater. Synth. Process. 1 (1993) 1,3-16. 10. E.Bilotsky and P.Tomchuk, Proc. 3rd Int. Works. EPMS-91, Berlin, Germany (1991) 410. 11. A.M.Kosevich, Physical mechanics of real crystals. Naukova Dumka, Kyiv (1981) 328 p. 12. M.Gasik, FGM News. J. FGM Forum Jap. 31 (1996), 6-9. 13. R.Utiyama, Phys. Rev. 101 (1956) 5, 1597-1607. 14. C.N.Yang and R.L.Mills, Phys. Rev. 96 (1954) 1, 191-195. 15. M.Gasik, Y.Bilotsky, N.Cherradi and A.Kawasaki, Proc. Europ. Powd. Metal. Conf EPMA'96, Stockholm (1996) 7.

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I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V, All rights reserved.

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Computer-Aided Process Design for Forming of Pore-Gradient Membranes C.-W. Hong Universitat Erlangen-Niirnberg, Institut fiir Werkstoffwissenschaften (3), Martensstrasse 5, D-91058 Erlangen, Germany

Abstract The DEM simulation method has been used to study the mechanics for the pore-gradient formation via centrifugal casting. A model system based on the dynamic similarity has been applied for the simulation. The simulation results show that pore-gradient membranes can be formed by using stabilized AI2O3 suspensions with a suspension height of 70 mm and a density of < 10% under a centrifugal acceleration of greater than 40 g. These design parameters have been applied and verified experimentally. 1. I N T R O D U C T I O N The Discrete Element Method (DEM) has been used to model the particle packing dynamics during forming processes of colloidal ceramic powder systems [1-4]. Hereby the response of each particle to mechanical and electrostatic impact from its surroundings has been described by the summation of all interactions, e.g. particle-particle, particle-medium, particle-boundary, and particle-external force field interaction. Long-range potential interactions, short-range repulsive hydration, steric stabilization through surface-adsorbed polymer as well as adhesion during solid state contact were taken into account. Using this DEM simulation method, the time-dependent particle packing processes and the resulting pore structure in the greenbody compact can be analysed. In addition to high temperature stability, ceramic membranes exhibit backflush capability, they are chemically and microbiologically resistant and have an excellent mechanical stability. Most of ceramic membranes are used in micro- and ultrafiltration applications (0.1 to 1.0 fiin and 1 to lOOnm., respectively) [5,6]. The extremely ultrafine pore size (< 0.1 iim) of the filtration (or inner) layer is produced by using the sol-gel method [5-9]. Typically, this filtration layer is formed on a support (or outer) layer which has larger pore size (1 to 6 f-irn) and is manufactured by slip casting [6,8]. Between the support layer and filtration layer, a discontinuous transition of material properties exists. Which, however, can be avoided by a gradient structure. It is well-known that large particles form the large pores, while small particles form the small pores. Multimodal suspensions will form greenbodies with pore gradient if the size segregation can be ensured during the forming processes. The developed DEM simulation program has been used to study the mechanics for the pore-gradient formation via centrifugal casting under diff"erent colloidal and process conditions. The necessary process parameters were designed and optimized by this DEM simulation tool. Hereby the fabrication of AI2O3 flat membranes was considered as a case study and will be shown in this paper.

30

increasing pore-channel diameter (a) asymmetric membrane

(b) pore-gradient membrane

Figure 1: Analogy between asymmetric and pore-gradient membrane. 2 PARTICLE PACKING A N D PORE-GRADIENT FORMATION Ceramic powders have usually a continuous size distribution. For well-stabilized and lessconcentrated suspensions, a size separation can be achieved through the optimized centrifugal casting processes. Once the size separation can be ensured, a continuous pore-gradient structure can be fabricated. The resulted pore structures will look like the random packing structures of nearly monosized spheres. Hereby the large particles form the large pores on the bottom of the sediment, while small particles form the small pores on the top layer of the sediment. This pore-gradient structure is analogous to an asymmetric membrane (see Fig. 1) which is prefered in the filtration applications. The resistance of an asymmetric membrane to mass transfer is determined largely or completely by the thin top layer. There are several factors which may influence the size separation and packing structure during the centrifugal casting. These factors include Zeta potential, ionic and solid concentration, centrifugal acceleration, viscosity of medium, size distribution and suspension height, etc.. The particle packing dynamics under simultaneous influences of these factors will be studied in the following sections by using the DEM simulation method. 3 D E M A N D D Y N A M I C SIMILARITY FOR PROCESS SIMULATION 3.1 D i s c r e t e element m e t h o d ( D E M ) The discrete element method regards each particle as an individual element, and where the motion of each particle is calculated according to the Newton's second law. At time t, the total response forces EF,- and moments EM,- of each element to its surroundings produce linear (a) and angular {to) accelerations: EMi = JCJ .

EF,

For the numerical computations, if the time step At is chosen small enough, then the accelerations (a.ij) could be assumed to be constant. At constant linear and angular accelerations, the linear v and angular uj velocities can be obtained by t-\-At

- /

idt ^ a At

and

UJ

I

t-{-At

Lodt ^ LO At

Using the trapezoidal rule, the Hnear and angular displacements have been calculated. At time t-|-At, these displacements are used to compute the new displacements at the end of the current time increment. This process is repeated until the response forces and moments become zero or when the process is stopped.

31 3.2 D y n a m i c similarity in centrifugal casting The DEM method can't simulate the system in full-scale because of computing limits, e.g. number of particles, computing time, etc.. Like many problems in the fluid dynamic, a smaller model system has to be developed and tested for the design purpose. Test results obtained from experiments with the smaller model system can then be transfered to the full-scale system. In centrifugal casting process, the centrifugal force (or accelerated gravity) is the only driving force for the size separation and mass transport, and only inertia and centrifugal forces play an important role. For such flow phenomena, the Froude number Fr = v/ s/gl should be the criterion of similarity [10]. The characteristic length / may be the suspension height, if the centrifugal acceleration g and the velocity of particle v are involved. For dynamic similarity between the model and full-scale system, the relationship can be expressed as (Fr)„, = (Fr)j

or

—^^^

= —L=

W hn 9'fn

.

(1)

\lfQj

If the velocities of particles at the corresponding points of both systems are assumed to be identical, i.e. ^;„^ = v/, then the conditions for the dynamic similarity reduce to

^m gm = yjhgs

or Im gm = hgf

•

(2)

For the case of free sedimentation, a sphere with diameter dp and density differing from that of the medium by A p moves at the Stokes settling velocity, Vo = ~ j ^ ^ g 18?/

,

(3)

where g is the centrifugal acceleration (or gravity) and rj the medium viscosity. Introducing the process times tm — lm/^\n and tj — Ij/vj and the model ratio l^ = Im/h ^^^ using the relationships in Eq. 2 and 3, the process time ratio can be expressed as

^ = k . i l ^ / , . . i ^ = /;2 . if

If

v,^-^

(4)

g.fy^

4 PROCESS SIMULATION A N D DESIGN 4.1 M o d e l s y s t e m For a full-scale system w^ith a suspension height of If — lOrmii^ a model suspension height with Im — 70 pm has been introduced. This corresponds to a model ratio of Ir = 10^. In the case of an applied centrifugal acceleration of gf — iO g (g ^ lOm/5^), the centrifugal acceleration for the model system gm. must be calculated according to Eq. 2, so that the dynamic similarity can be fulfilled. This results to g^ = 40 000 g. For the simulation, multimodal AI2O3 powders (p = 3.98 (//c??2"^) with the following discrete size distribution, 5 % with c/i = Ipni, 20% with do = O.Spm, 40% with c/3 = 0.6 pm, 2 5 % with (^4 = OApm and 10%) wdth c/5 = 0.2pm, will be used. The DLVO (Deryaguin-Landau-Verwey-Overbeek) potential interactions and medium influences described in [4] have been taken into account. A Zeta potential of ( == 40/7/1'', the Hamaker constant of AH = 4.76 x 10~^^J between AI2O3 particles in aqueous solution and a Debye-Hiickel parameter of AC = 10~^m~^ for the suspensions were assumed. In the following presentations, the process times for the full-scale system according to Eq. 4 have been used.

32 S 0.8|

(b)

pA = 5 % , t = 42.32mm

S

\(b) (a) initial state 0.61 \ (b) stabilized (C=40mV)| \ (c)flocculated(C=OmV)| 0.4

o

0.2

a ^

(c)-\

Bo PA = 5 % , t = 42.32 min

, 2 3 ^^^

(a)

^ 0.01 0-2 0.4 0.6 0.8 1.0 Normalized distance from left end •#•40 g

Figure 3: Particle packing dynamics and structures of a stabilized and a flocculated AI2O3 suspension and the corresponding size distributions of smallest particles. The initial state (a) is identical with the initial state in Fig. 2(a), and is therefore omitted here. of a stabilized and a flocculated suspension with PA — 5 %. In Fig. 3, the initial state is identical with the initial state in Fig. 2. The flocculated suspension shows obviously a worse size separation as the stabilized one. The agglomerates A (two particles with d^ = 0.2/utm and one particle with c/4 = OA/jm.) and B (three particles with d^ = 0.2 fim) in Fig. 3(c) will settle much faster than a single stabilized particle with d^ = 0.2 jim. This prevents the formation of a good pore-gradient. 4.4 Influences of solid concentration Using different concentrated suspensions, different thicknesses of sediments can be obtained. It is therefore desirable to know that can a higher concentrated suspension ensure also a good size separation for a good pore-gradient formation? In order to answer this question, suspensions with area density of pA = 7.5% and 10% have been simulated. There is no difference in the size separation effect to be observed in the simulation results in Fig. 4. The only difference is the thicknesses of each sediments. 5 CONCLUSIONS The DEM simulation results show that the pore-gradient membranes can be formed by using the stabilized AI2O3 suspensions with a suspension height of 70 mm and a density of < 10 % under a centrifugal acceleration of AOg. A process time of t > 6 hours is enough for a complete sedimentation of the finest particles with d^ = 0.2 pm. Using flocculated suspensions, the quality of pore-gradient will be negatively influenced. But using a higher centrifugal acceleration, a shorter process time, a better size separation and better pore-gradient can be expected. These simulation results have been used as process design parameters and verified experimentally [11]. AI2O3 flat membranes via centrifugal casting with a graded pore-size distribution from 40 to 250 nm across a membrane thickness of about 5 mm have been shown in [11]. A c k n o w l e d g m e n t : The author would Hke to thank the German Science Foundation (Deutsche Forschungsgemeinschaft, DFG) for the financial support.

33

(a) initial state (b) t = 25.27 min (c) t = 48.64 min

•P2J»? •».

(a) l.Oj

PA =

5% C=40mV

0.8

Kb);

0.6

:ic)

0.4

•N^

J^L-

0.2

(bj-—,P25

1(c)'

0.2

0.4

0.6

0.8

1.0

Normalized distance from left end

^40 g Figure 2: Pore-gradient forming dynamics of a stabilized AI2O3 suspension during centrifugal casting with an acceleration of 40 g and the corresponding size distribution of smallest particles. 4.2 Pore-gradient forming dynamics of stabilized suspensions It is believed that excellent pore-gradient structures can be obtained from well-stabilized and less-concentrated suspensions. Thus, a stabilized suspension with 250 particles and the above-mentioned discrete size distribution have been generated in a square box of 70jum X 20.53 ^ m without contacts for the simulations (see Fig. 2(a)). This corresponds to an area density of pA — 5%. Hereby periodic boundaries in the horizontal direction and walls in the centrifuging direction have been assumed. Fig. 2 shows the time-dependent pore-gradient forming dynamics during casting and the corresponding size distribution ratios of smallest particles with c^5 = 0.2 fim. Hereby the distribution ratio was calculated in the way that the counted smallest particles in every distance partition, in which 1/5 of all particles (in this case: 50) exist, was divided by the total number of smallest particles (in this case: 25). This size distribution ratio of smallest particles is taken as reference quantity for the size distribution and for the gradient formation. The initial state shows a relatively homogeneous size distribution. The size separation increases with the processing times. After t = 48.64 min, more than 80% of the smallest particles remain in the suspension. The smallest particle P25 needs at least 7 times of the past processing time, i.e. t — 1 x 48.64 min ^ 6 hours, to reach the sediment. Therefore, it can be expected that after 6 hours a thin top layer made of smallest particles can be formed and a pore-gradient can be obtained. There are only less than 20 % of the smallest particles are packed in interior of the sediment, because their free sedimentation ways are too short for a good segregation. 4.3 Influences of Zeta potential For the flocculated suspensions with C = OinV, the size segregation effect will be disturbed by the dominating attractive interaction between particles. This disturbance can be observed even in the less-concentrated suspensions. Fig. 3 shows the packing structures after a process time of t = 42.32 mi??- and their corresponding distributions of smallest particles

34

(b)

^

0.8

^

0.6

0.4

c= 40mV,t = 25.27 min (a)/>A== (b)pA ==

5% 7.5 % (c) PA=:10%

\^ vs. "^^

(b)

0.2

•^^^N

(c)

(c) C = 40 mV,t = 25.27 min

0.0 0.2

0.4

(a) ^ ^" - '^^' 0.6

0.8

1.0

Normalized distance from left end

^40 g

Figure 4: Particle packing dynamics and structures of stabilized AI2O3 suspensions with different solid fractions and the corresponding size distributions of smallest particles. The initial state (a) is identical with the initial state in Fig. 2(a), and is therefore omitted here. REFERENCES [I] C.-W. Hong and P. Greil, Discrete Element Modelling of Colloidal Powder Processing, Ceramic Transactions, Vol. 54: Science, Technology, and Applications of Colloidal Suspension, American Ceramic Society, (1995) pp. 235-249. [2] C.-W. Hong and P. Greil, DEM Simulation of Particle Packing Dynamics during Colloidal Forming Processes, Ceramic Transactions, Vol. 51: Ceramic Processing and Technology, American Ceramic Society, (1995) pp. 349-353. [3] C.-W. Hong, Computer-Aided Process Modelling of Colloidal Powder Forming{m German), Fortschrittsbericht der DKG 10 (1995) [4], pp. 225-241. [4] C.-W. Hong, Discrete Element Modelling of Colloidal Packing Dynamics During Centrifugal Casting, J. of the Ceramic Society of Japan, 104[9] 793-795 (1996). [5] A. Larbot, J.-P. Fab re, C. Guizard, L. Cot, New Inorganic Ultrafiltration Membranes: Titania and Zirconia Membranes, J. Am. Ceram. Soc, 72[2] 257-61 (1989). [6] K. K. Chan and A. M. Brownstein, Ceramic Membranes - Growth Prospects and Opportunities, Am. Ceram. Soc. Bull., 70[4] 703-7 (1991). [7] C. Guizard, A. Julbe, A. Larbot, L. Cot, Nanostructures in Sol-Gel Derived Materials. Application to the Elaboration of Nano filtration Membranes, Key Engineering Materials Vols. 61 & 62 (1991) pp. 47-56, Trans Tech PubHcations, Switzerland. [8] R. Kohl, G. Tomandl, A. Larbot, L. Cot and J. Gillot, Herstellung und Charakterisierung von keramischen Membranen aus Titanoxid zur Cross-Flow-Ultrafiltration nach dem Sol-Gel-Verfahren, Kurzreferatepp. 64-66, DKG-Jahrestagung, Bayreuth, 4.-7. Oktober 1992. [9] Q. Xu and M. A. Anderson, Sol-Gel Route to Synthesis of Microporous Ceramic Membranes: Thermal Stability of Ti02-Zr02 Mixed Oxides, J. Am. Ceram. S o c , 76[8], 209397 (1993). [10] D. N. Roy, Applied Fluid Mechanics, Ellis Horwood Ltd., Chichester, 1988. [II] C.-W. Hong, F. Miiller and P. Greil, Fabrication of Pore-Gradient Membranes via Centrifugal Casting, paper presented at the 4th International Symposium on Functionally Graded Materials (FGM), Tsukuba/Japan, October 20 - 24, 1996.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

35

Mathematical Model for Axial-Symmetrical FGM X.D.Zhang^ T.Q.Lm% C.C.Ge^ "Mathematics-Mechanics Faculty, ^Laboratory of Special Ceramics & P / M University of Science and Technology Beijing

Abstract Analytical procedure for empty cyUnder of Functionally Graded Material(FGM) with axicJ symmetry has been investigated. Based on thermcil elasticity theory and computational mathematical method, temperature distribution has been obtained and thermal stress of FGM in ring section has been calcidated.

1

INTRODUCTION

Ceramic-metal FGM is a new kind of heterogeneous composite material that consists of a gradient compositional variation from ceramic to metal from one surface to the other. These continuous changes result in property gradients. Because ceramic has good heat-resistance and metal has high strength. Ceramic-metal FGM can work at super high-temperature or under high-temperature difference field. Because of its good property , ceramic-metal FGM is expected to be used in hghtweight structure such as 2iircraft and spacecraft structures. It is the key problem to reduce thermal stress made from different effective materia properties for intermediate compositions of the FGM and prevent destruction by thermcil stress. This paper discusses temperature distribution 2ind thermal stress distribution in empty cylinder. This topic wiU be of great interest for fabrication of ceramic-metal FGM.

2

ANALYTICAL MODEL

Considering an empty cylinder with axial symmetry whose iimer radius is a, outer radius is b. The cylinder in ring section is made of n layers through the radius, which could represent a gradient of material distribution. CyUndrical coordinates are used and the z-axis is taken as symmetrical axis as shown in Fig.l. Such an ideaHzation requires uniform mixture of ceramic and met Ell phase and isotropic in each layer. It is assumed that the FGM is placed in a steady state temperature field. The temperature of inner and outer surface of FGM are Ti and T2 respectively, and two ends of model are insulation.

36

••

^

Figure 1. Analytical model. Figure 2. Mechanical model

3

DISTRIBUTION FUNCTIONS OF FGM CYLINDER

Assuming that the constituents of FGM comprise phase A(metal), phase B(ceramic). The volumes for contituents are express as VAandVB respectively. Let J^

-

VA+VB

' J^

-

VA-\-VB

here fAifs axe the volume functions for the metal and ceramic of FGM respectively, which satisfy the following equation /4 + / B = 1 Supposing r-coordinate of each layer is expressed as a = ro < 7*1 < • • • < Tn = 6 f\ denotes volume function of metcJ phase of FGM in i-th layer. According to hypothesis of analytical model, we have f\ = Ci{cost), Ti-x
fAir) = J2f/'l\{r),a
^^ ^

\ 0,

(1) * )^) is base function, the definition of which is as follows

< r < ri ri
(2)

37

l'A{r)=<

i^^^r, I 0,

1 0,

ri
ro
r„_i n-l

(3)

(4)

' U J^ ' ^ '

p is a form factor.

4

CALCULATION OF EFFECTIVE PROPERTIES

In Older to calculate the temperature distribution and the thermal stress distribution in the FGM cylinder, effective material properties such as heat expansion coefficient a(r), heat conduct on coefficient k(r), the volume modidus K{r)^ the shear elasticity G(r) and Young's modi V. E(r) for intermediate composition of the FGM are required. This paper assumes that ' ckch layer of FGM is simple macroscopically isotropic two-phase system with spherical particles.We utilize formula proposed by K.Wakashimat^Ho calculate effective properties in each layer.

5

TEMPERATURE DISTRIBUTION OF FGM

Because the empty cylinder of FGM is placed in steady state temperature field, two ends are heat insulation, the temperature distribution T is no relation with z, 9 and time t. One dimensional equation of heat conduction in r direction is given as: A_LZ^ =0

(5)

It satisfies boundary conditions T(a) = Ti,Tib) = T2 It is very easy to solve this equation by virtue of the assumption of FGM being homogeneous in each layer. Let us denote ki as the heat conduction coefficient at i-th layer. The temperature distribution can be expressed as follows by using local radius variables Vi : T(r) - Ti + C{T2 - Ti)( E MiiAiz^ + C =^l/{f2 H^»A'-0)

^-i^irJiA) (6)

Ts
6

THERMAL STRESS CALCULATION

According to hypothesis of analysis model, we can know variations of temperature AT(r) is symm-etrical with the z axis, and have nothing to do with the z coordinate. And thus displacement W along the z axis can be zero. Suppose taking the i-layer of the empty cylinder as an obj'ct of study as shown in Fig.2 , there are only three stresses (Tr{r), <7"^(r), ^^^(r). Owing to

38 symmetry of the center axis and equalization along the axis, three shearing stresses and shearing strains can be zero. crr(r) and (TQ(r) fit balance equation:

^

+ ^^ir)^

=0

(a)

Because of W = 0,6^ = 0 . So crz(r) = McTrir) + M^)) -

aiEiAT{r) (b)

Stress-strain relation is £, - (1 + Ui)aiAT{r)

== i ^

[
j^^Mr)]

£ « - ( ! + ./i)aiAT(r)

= i ^ K ( r ) - j^
Using u as redial displacment 2ind geometrical equation is

^r = t ; ee = ^

(d)

Prom (c) (Trj(T0 can be solved, and then return to balance equation (a) with (d), u as a residt can be solved.

u =i ± ^ ^

r AT{r)rdr + AiVi + —

\ — Ui Vi J a

(7)

Tj

SO

'^^^'^^ - 1 - ^i r? y„ ^^^'^^'^^ + (1 + .0(1 - 2^0 ^ (1 + "O'-l

(1 - "0

(9)

from (b) ,we have

'^^('•) - i r : 7 ; o ~ ^ (1+"0(1-2.0

^'°^

here Ai^Bi are certain integral constants, which can be defined by the boundary corditions and continuity condition. Two equations can be obtained by the boundary condition: (Tr{a) = 0,(7^(6) = 0.

39

On the basis of the continuity condition for the boundztries in every layer, u or a{r) along r direction is continuous, so

{(Tr)i = {(Tr)i-\.l t =

l,2,..-,(n-l)

According to hypothesis of analysis model, the empty cylinder has n layers . So there are 2(n-l) continuity conditions. Getting them together with two boundary condition, we have 2n equations. Thus 2n integral constants Ai^Bi can be solved by 2n equations. These equations are expressed as follows:

iP^ai

- l±i^ai+i)i r

I -fi

1 - I'i+l

aiEi at+iEi+il r^lEL_?^i±l^i±l)L

AT{r)r
ri Ja

p r

(13)

Ti

.^, . , ATMrdr

, +

, f

EjAj ^^

r ? l ( l + «/,)

Et+iAi+i ^ +i)(l-2i/i+i)^ ( l + «/Hi)^

^ '

For simphcity, let

Ci = -{-

ai -

Dj = —^(Bi

-

«i+i)-2 /

AT(r)r(fr

Bi^i)

i=l,2,...,(n-l) The equation (13) can be simpHfied as: Ai - Ai+i = Ci + Di

(15)

Similarly, let Fi = -{-

aiEi

ai^iEi^i

) /

r*

AT{r)rdr

Hi = z—-—^Gi l-2i/i ' l + i/i z = l,2,-..,(n-l) The equation (14) can be simpHfied as: GiBi - Gi^iBi^i

= Fi^ r^{GiHiAi

- Gi^iHi^iAi^i)

(16)

Thus Ai,Bi{i = 1,2,- -,71) can be solved by equation (11) (12) (15) (16). Finally, getting Ai^Bi{i = 1,2,- -,71) to (8) (9) (10), the expression of thermal stress in every layer can be obtained.

40

7

CONCLUDING REMARKS

ON the basis of the hypothesis analysis model, this paper gives temperature distribution and thermal stress distribution of FGM. By using above formulas to calcidate thermal stress in every layer, the optimum combination and distribution profile will be expected to obtain.

8

ACKNOWLEDGMENT

this project is supported by China National Natural Science Foundation, Doctorial Programe Foundation of State Education Commission and China National Committee of High Technology New Matericils.

References [1] K.Wakashima & H.Tsukamoto,Proc First Int . Symposium, FGM , Sendai JAPAN:1990 [2] Tohru H,Junichi T,Tomohiko Y,Proc First Int Symposium, FGM Sendai JAPAN:1990.

I. Shiota and M.Y, Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

41

Stress analysis in a two materials joint with a functionally graded material Y.Y.Yang and D.Munz Institute for Reliability and Failure Analysis, Karlsruhe University P.O.Box 3640,D-76021 Karlsruhe, Germany

Abstract In this paper the stresses in a joint with a functionally graded material (FGM) are analyzed. In the middle of a joint with FGM the stresses have been calculated analytically by using the plate theory. The effect of the thickness of the FGM layer and the effect of the transition function form on the stresses in the joint is discussed. Near the free edges of the interface in a joint with FGM, the stresses are described analytically by using the Mellin transform method. Some examples are presented to show the good agreement of the stresses calculated from FEM and with the analytical description in a joint with graded material.

1

Introduction

After a homogeneous change of temperature very high stresses exist near the free edges of the interface of a two dissimilar homogeneous materials joint due to the difference in elastic constants and in thermal expansion coefficients. For most material combinations with elastic material behaviour stress singularity occurs, i.e. the stresses are infinite at the intersection of the free edges and the interface. By introducing a FGM interlayer these stresses can be reduced significantly and the stress singularity at the free edge of the interface vanishes [1]. The functionally graded material is often used for reducing the thermal stress or residual stress in a dissimilar materials joint, also for another aims. The material properties in FGM depend on the coordinates. Therefore, it is very difficult to calculate the stresses analytically in such a joint, even for very simple geometry and simple loading. The Finite Element Method (FEM) is generally used to calculate the stress distribution in a joint with FGM. However, in the middle of a thin joint the stresses can be calculated analytically by using the plate theory. In the present paper the Mellin transform method is used to describe analytically the stresses in the edge range of a joint with graded material.

42

i

y A

y 1

,

Material 2

i

1

'

Material 2

I

FGM

h. ' f i

hi

1

yf

FGM

— • Xj

Material 1

hi

X

'

_.,.,

^x

(b)

(a) Fig.l The geometry of a joint with FGM.

2

Stresses in t h e middle of a joint with F G M

Using the plate theory, the stresses in the middle of the joints in Fig.l can be calculated from (1)

"•«' = T ^ ( ' ' - S - ^ ^ " ' « >

where the quantities e^ and R can be determined from the equilibrium conditions of the force and moment, i=l,2,...,N is the number of layers in the joint, E is Young's modulus, u is Poisson's ratio and a is thermal expansion coefficient. Fkt. 3

Fkt. 2

-^ ^^

Fkt. 4

4—e-^^ ^r;/y ^ = 0) qi^y = 0)

0 OQ 0.2

0.4

0.6

0.8

I 0.2

-O a^y = h,) 1 1.0

I 0.4

i 0.6

0.8

1.0

f

f

Fig.2 (left) Stresses versus f for hi/H=O.Ol {Ei = 200 GPa, u^ =: u^ = 0.3, c^i = 5 * 10-yK, E2 = 100 GPa, 0^2 = 10 * IQ-yK), Fig.3 (right) Stresses versus f for hi/H=0.05. In Fig.la, we assume that in FGM the transition function of E and a is the same, i.e., E{y) = g{EuE2, y),

a{y) = g{au 0^2, y)

(2)

where Ei, ai is the value of the material at y=0 and ui = iy2- We define a quantity

f = ^j\ifuf2,y)dy

(3)

43 with / i = 1 and /2 = 0. The quantity f is the average value of material 1 in the FGM. f = fi = 1 means that FGM is the homogeneous material 1, f = /2 = 0 means that FGM is the homogeneous material 2 and a linear transition function corresponds to f=0.5. The different value of f means that in FGM the transition function form is different. For hi/H =0.01 (H=/ii + /12), the stresses at different positions are plotted versus f in Fig.2 and for hi/H =0.05 in Fig.3. It can be seen that for a very thin FGM layer (thickness of FGM = 1% H) the transition function form in FGM has a negligible effect on the stresses, whereas for a thickness of FGM = 5%H the transition function form in FGM has a significant effect on the stresses. Generally, it is impossible to reduce the stresses at any position of the joint for thermal loading. Therefore, the optimization of the stress in a joint with FGM is not unique. There is no general rule for the optimization of the stress in a joint with FGM. The optimization of the stress in a joint with FGM may be: (a) reduce the maximum stress in the joint; (b) move the location of the maximum stress from the interface; (c) reduce the stress at a special location, e.g., at the upper surface, at the lower surface, at the interfaces, etc. Some studies on (a) and (b) have been given in [2]. As an example for the case (c), we use the transition function (coordinates see Fig.lb) E{y) = g{Ei, ^ 2 , y) = E2 - (^2 - ^1)

hi +

h2-y

(4)

where n can be varied to reach a optimization of the stress. Our aim regarding the optimization of the stress is that, for example, at the lower surface (y=0) or at the interface (y=^i) the stress in the middle of the joint equals zero. We can vary n and the geometry ^ 1 , ^2, hs. For a given geometry hi, /i2, /13, the corresponding values of n have been found and they are plotted in Fig.4 for cTxiv = 0) = 0 and in Fig.5 for cr^iy = hi) = 0, where U—hi + /12 + ^3- From these figures we can see that for some geometries it is impossible to reach the optimization, if this transition function form is used.

n 0.125

0.25

0.375

0.5 h,/H

0.625

1

r

0.75

0.875

1.0

0.375

Fig.4 (left) The possible n for a^{y = 0) = 0 {h2 = H - hi - /13). Fig.5 (right) The possible n for a^iy — hi) — 0.

0.5 h,/H

0.625

0.75

0.875

1.0

44

3

Stresses near t h e free edges of a joint with F G M

For thermal loading, the stress function $ in a homogeneous material must satisfy V^^ + V2(gT) = 0

(5)

aE for plane stress aE ^ ^I 7(1-^) ^ ^ for plane strain ' where q is independent of the coordinates and T is the temperature change. Now, in a graded material (i.e. q is not a constant, but q=q(x,y) =q(f, ^)), we can imagine that qT is the effective temperature change and the material is homogeneous. Under this assumption, we obtain the same equation as Eq,(5) for the stress function in a graded material. To get the solution for $, the Mellin transform method is used. The Mellin transform of a function (f, 6) is defined as _{

—\

roo

^s,e)=

/

^{r,e)f^'-^^df

(6)

Jo

where s is the Mellin transform parameter and f = r/L, L is a characteristic length of the joint. The basic equation Eq.(5) in the Mellin domain is [«' + ^ ] [ ( ^ + 2)' + ^]Ms,

with

0) + [is + 2f + ^]f{s

f{s + 2,6>) = / q(f, e)Tr^'-^^^df = To J

+ 2,9) =0

(7)

q{f, e)f^'^^Uf,

(8)

0

where the temperature change is T=To for f < RQ and T=0 for f > RQ. Its solution is

^,{s,e) = ¥,{s,o) + ^'(s,e) with

4^(5,9) = Aks''^ + AkC-''^ + Bke'^'"^^^^ + 5^6-^(^+2)^

4^^(s, 0) = — Lin{se) jfk{s

(9) ^^^O)

+ 2, e)cos{sO)de - cos{se) ffk{s -h 2,0)sin{se)de\ (11)

where the coefficients Ak and Bk can be determined from the boundary conditions (k=l, 2 for material 1 and 2). For the case of a polymonial as the transition function, i.e. E{f, e) = A-\- Bfsm{0) + Cr^ sin2(6>) + Df^ sm\0) + Er^ sin^(l9) + ... the stresses in the Mellin domain are obtained [3] as

(12)

45 where Sn is the solution of A*A2 = 0 and dijk{sn, 0) / 0, The stresses in a polar, coordinate system can be calculated from

Sn<7

-

V

^-

lim '^

Us - s ) ' ^

'^'^'^^'^^

r-'^+^H

(14)

For each s^ there is ^zjn(r, 0) = 7 ^ - [ ^ n ( ^ , i^n) + 9^Jn{0)ln{f)

+ /l.,n(^)],

(15)

where Un = —{sn + 2), the functions fijn{^^Kn),gijn{^)^ hijn{0) can be calculated analytically, but the factors K^ should be determined from one numerical method, e.g., FEM. Finally, the stresses in the edge range of a joint can be calculated by

^iiir,0) = J2 {{r/Lr-[fij„i0,K„) For the geometry 6i

+ gij„{e)lnir/L) + V W l } -

(16)

-0^ = 90° there is A2 = —sin{{s -h l)7r).

(17)

Its solutions are Sn = 0, ± 1 , ± 2 , ± 3 , ..,, ±n,,.. s = - l corresponds to the rigid body displacement and for 5 > —1, the displacements at r = 0 are infinite. Therefore, Sn — —2, —3, —4,... , -n, ... are the interested poles of dij{s,6). This means that there is uj^ — 0 , 1 , 2 , 3 , . . . . For A* we have A*-/(^,i/,50.

(18)

The solution of A* = 0 depends on the values of B and v. If ^ = 0, the solutions of A* = 0 and A2 = 0 are the same. If J? 7^ 0, s^^ can be arbitrary value, but 5^ < —2. EXAMPLES

(coordinates see the following Fig.)

The materials data for example 1 are Ex = 100 GPa, z^i = 0.25 = 1^2, o^i = 2.5 * l O ' V K , £;2 = 100 + 50f25m2(6>) ,GPa 0^2 = [2.5 -h 5f'^si'n?{e)] * 1 0 - 7 ^ ^ . The materials data for example 2 are El = 100 GPa, 1^1 = 1 = 1^2,^1 = 2.5 * I O ' V K ,

E2 = 100 + 50f5m(i9),GPa 0.2 = [2.5 + bf^sin^iO)] * 1 0 - V K .

For example 1 we have ^ = 0, therefore, the poles of aij{s, 6) are independent of u. They are s = -2, -3, -4, -5, -6, ..,, where s=-2, - 3 , -5 are the first-order poles of Gij{s,9) and

46 s=:-4, -6 are the second-order poles of dij{s,6). The stresses calculated from Eq.(16) and by FEM are compared and given in Fig,6. For example 2 we have B ^ 0, therefore, the poles of dij{s,6) depend on the values of v and B, They are s== -2, -3, -4, -4.1529, -5, ..., where s=-2, -4.1529 and -5 are the first-order poles, and s=-3 and -4 are the second-order poles of cFij{s, 9). The stresses calculated from Eq.(16) and by FEM are compared and given in Fig.7. From Figs.6 and 7 we can see that in the range of r/L < 0.1 Eq.(16) can describe the stresses very well. 0.030-

O.OlbO-i

0.020-

0.0100-

—™««niaBBB»»n»™«»»»R'^

0.00500.010H \

0.0000-

CO

o.oooH -0.0050-0.010-

...,.^V \

-0.0100o—

-0.020- TTT 1 lE-04

1 I II 1 lE-03

1 I I I 1 1 I Ir 1E-02 1E-01 1E+00 r / L along the line 8 = 0

p -0.01501E-04

"' "^'^1

1—TTT

1E-03

BBiuiiuuumMiii!™::::;^^

^

r-

m — 1 — 1 -I 11—^^t—1 1 11 1E-02 1E-01 1E+00 r / L along the line 8 = - 4 5

Fig.6 Stress distribution along ^ = 0 and 9 =-45 for example 1. U.UJUUt WHiHHHtHtfWHHBIfflgmffl^iqw^ge

0.0200-|

0.0250-

0.0150-

^

^

\

™ ^ 6

/

0.0100-

0.0200-

0.00500.01500.0000-

\

0.0100-

5^ a

-0.00500.0050-

-0.0100o

-0.0150- J

0.0000-^ -0.0050- W T

. 1E-03 2

t-r-n 1 r—r-n \ 1^• i - n 4 lE-02 2 4 1E-01 2 4 1E+00 r / L olong the line 8 = 0

C «•>,

lE--03 2

•W 1\ ^ 4

1E-02 2 4 lE-01 2 4 1E+00 2 r / L along the line 8 = - 4 5

Fig.7 Stress distribution along ^ = 0 and 9 = - 4 5 for example 2. Acknowledgement:The financial support of the Deutsche Forschungsgemeinschaft is gratefully acknowledged. The authors would like thank Mr. Schaller for some calculations.

References 1. D.Munz, Y.Y.Yang, Proc. of 3rd Int. Symposium on Structural and Functional Gradient Materials, 1994, pp.465-471. 2. Y.Y.Yang, D.Munz, Fracture Mechanics: Vol.26, ASTM STP 1256 (1995) pp.572-586. 3. Y.Y.Yang, Stress analysis in a two materials joint with a functionally graded material under thermal loading by using the Mellin transform, submitted to J. Solids & Structures.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

47

Optimum Design and Fabrication of TiC/NijAl-Ni Functionally Graded Materials Q. Shen , X.-F. Tang , R. Tu , L.-M. Zhang and R.-Z. Yuan State Key Lab. of Materials Synthesis and Processing, Wuhan University of Technology, Hubei, 430070 P.R. China

TiC/Ni3Al-Ni system was chosen for the potential use as the first wall materials of thermonuclear experimental reactors. The residual thermal stresses during fabrication were calculated by finite element inethod. Based on the consideration of minimum stress, minimum ratio of stress to fracture strength and proper distribution of thermal stress in pure TiC side, the optimum design with the distribution exponent P=1.6 was reached. According to the design result, TiC/Ni3Al-Ni FGM samples were then successfully fabricated.

L IISTRODUCTION In recent years functionally graded materials(FGMs) have received considerable attention. Ceramic/metal type FGMs are believed to be hold promise in applications for advanced technology, especially in aerospace and nuclear engineering, where the materials encounter high heat loads[l-2]. On the basis of the traditional cermet TiC-Ni, we chose TiC/NiaAl-Ni FGM system, which was composed of Ni as the metal phase, TiC as the ceramic phase and TiC-Ni3 Al composites with various Ni3Al ( has a lower thermal expansion coefficient and a better wettability with TiC[3] ) volume fraction as the graded interlayers.

2. EXPERIMENTS AND CALCULATION Pure metal Ni and TiC-Ni3Al composites with Ni3Al volume fraction of 0, 20%, 40%, 60%, 80%, 100% were sintered by HP method for 2h at ISOO^C and 30MPa under Ar gas protection( the same as the conditions for sintering TiC/'Ni3Al-Ni FGMs). The samples were then cut and ground into 36x4x3mm strips, the mechanical properties including Young's modulus, Poisson ratio and the fracture strength, were measured by a four-point bending test method. The thermal expansion coefficient was tested by a non-loaded thermal dilatometer. The tested v ^ e s are indicated in Table 1.

48 Table 1 Properties of Ni and TiC-Ni3Al composites with various Ni3Al volume fraction Ni3Al volume fraction (%)

0

Young's modulus (GPa) Poisson ratio Fracture strength (MPa) Thermal expansion coefficient (xl0"%-l)

20

40

60

320

318

340

267

0.195

0.195

317

587

0.225 1351 9.15

0.253 1261 9.46

7.40

7.55

80

100

100(Ni)

266

199

206

0.270 0.295 0.30 1468 1346 1322 11.46 11.90 13.30

The residual thermal stresses were calculated by a finite element computer programme Super-Sap. The model is 10mm in diameter, 6mm in thickness and has 11 graded layers. Due to the axisymmetric problem, an axisymmetric finite element modle is used and 1/4 part of the material is considered. The FE model includes 1200 elements and 1275 nodal points. The thermal load is raised from the sintering temperature 1300 ^C to room temperature. Moreover, an insulated heat condition is considered at the flank boundary. The compositional distribution of the metal and ceramic in the graded layers was assumed to take the form C=(x/d)P[4], where C is the volume fraction of Ni3Al, d is the distance to the graded layers, x is the layer location coordinate, and P is the distribution exponent. In the calculation, the material properties at the graded layers were obtained by using the tested values.

3. THERMAL STRESS ANALYSIS AND OPTIMIZATION 3.1. Thermal stress analysis of TIC-Ni two layers material It can be seen from Table 2 that the stresses at the metal-ceramic interface are extremely high. For the radial stress GJ^ and the circumferential stress 099, the metal side is in compressive state and the ceramic side in tensile state, which due to the difference thermal expansion coefficient. The stress state of the axial stress a^z in Z-axis direction is quite the contrary. Such large stresses at the interface can easily lead to rupture failure of the two layers material. Also, the experimental results prove to be the same. Table 2 Calculated results of the thermal stress for TiC-Ni t(vo layers material an- (MPa) Metal side Ceramic side

-1930 1550

099 (MPa) -1930 1550

GZ^

(MPa)

1870 -837

49 3.2. Thermal stress analysis and optimization of TiC/NijAl-Ni FGMs The thermal stresses of TiC/Ni3 Ai-Ni FGMs were calculated for different exponent values ranging from 0.6 to 2.2. Fig.l presents the tensile stress a^r and Cjri with respect to the distribution exponent P. From Fig. 1, it is seen that within the exponent values tfie thermal stress in any of the FGMs is relaxed compared with that in the two layers material. In particular, when P=1.8, the stress a^- is a minimum (280MPa), and the stress relaxation is a maximimi, up to 75%, whereas the stress Cjz reaches the minimum value (290MPa) at P=1.6, the stress 1000

1.2 1.6 diBtributtonttxponontP

2

24

Fig. 1 Relationship between tensile stress and the distribution exponent P

1.2

1.6

2.4

oiscnxjDon oixporioni r

Fig. 2 Relationship between the ratio of stress to fracture strength and the distribution exponent P

50 relaxation in FGM is nearly 80%. According to the minimum stress rule, P=1.6~1.8 can be the optimum points for the FGM compositional design. It k noticed that the FGM interlayers have differentfracturestrength. If the thermal stress occurs at any interlayer was not optimized, the FGM could be damaged. By examining the ratio of the stress to the corresponding layer fracture strength, an optimum P can be obtained. From Fig. 2, the ratio for c^ is found to reach its minimum value (0.27) at P=1.8, but for Oj^ it is at P=1.4. Therefore, the optimum result for the FGM compositional design is P=1.4~1.8. To verify the reasonableness of the FGM compositional optimum design, one must also check the stress state at the pure ceramic side since it is usually damaged first. Fig. 3 provides the relationship between the stress at the TiC side and the exponent P. From Fig. 3, it is observed that the residual stress at the pure TiC side decreases with increasing of P, and it experiences a transition at P=1.6 from tensile to compressive, where the stress is zero. This is favourable for FGM fabrication. Therefore, P=1.6 is indeed an optimum design for TiC/Ni3Al-Ni FGMs. 800

•soo

dtatributlonflxponontP

Fig. 3 Relationship beween the stress in pure TiC side and the distribution exponent P 4 FABRICATION Table 3 Constituent (vol%) and density (g/cm^) of FGM layers layer Ni3Al TiC Density

1 100(Ni) 0 8.90

2

3

97.5 92.4 2.5 7.6 7.47 7.23

4

5

6

7

8

9

10

11

85,4 76.9 67.0 55.8 43.5 30,0 15.5 14.6 23.1 33.0 44.2 56.5 70.0 84.5 7.13 6.94 6.67 6.52 6.21 5.85 5.47

0 100 5.03

51 According to the optimum design result P=1.6, the constituent and density of each layer is determined, as shown in Table 3. The mixtures were prepared by mixing a 5pm average sized Tie powder, 5S3puxk sized Ni3Al and Ni powder, fhen sintered under the same conditicms stated previously.

(a) Layers near the metal side (b) Layers near the ceramic side (c) the 4th and 5th layers

Fig. 4 SEM micrographs of the FGM Fig.4 gives the SEM micrographs of the FGM specimen. From Fig.4a,4b, the layered projSle is clear. The SEM micrograph of the FGM interface between the 4th and the 5th layer is shown in Fig. 4c, the transition crossing the layer interface is smooth, no cracks and pores are observed. It is indicated that the expected design has been achieved during the fabrication process.

5. CONCLUSIONS (1) Within the exponent values ranging from 0.6 to 2.2, the theimal stress in any of the FGMs is relaxed compared with that in TiC-Ni two layers material.

52 (2) By analysing the thennal stress and its distribution, an optimum design with the exponent P=1.6 was reached. (3) TiC/Ni3Al-Ni FGM samples were successfully fabricated via a HP method.

REFERENCES [1] M.Niino, T.Hirai, and R. Watanabe, J. Jpn. Soc. Compos. Mater., 13(1987),257 [2] Edited by B.Dschner and N.Cherradi, Proc. 3rd Int. Sym. on SFGM, Presses Pofytechniques et Universitaires Romandes,1995. [3] L.M.Zhang, R.Tu and R.Z.Yuan, Acta Materiae Compositae Sinica, 12(1995),22 [4] A.Kawasaki and R.Watanabe, J. Jpn. Soc. Powder Metall., 37(1990),253

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

53

A Mathematical Model for Particle Distribution in Functionally Graded Material Produced by Centrifugal Cast Zhang Baosheng Zhu Jingchuan Zhang Yongjun Ying Zhongda Cheng Hongsheng An Geying School of Material Science and Engineering, Harbin Institute of Technology, Compus Box 434, Harbin 150001,P.R.China 1. Introduction Functionally Graded Materials (FGMs), with gradiently distributed reinforcement particles, can be produced conveniently by centrifugal cast due to the segregation phenomena which is caused by the different specific gravities of reinforcement particles and alloy liquids J ^'^'^^ To produce FGMs with accurately distributed particles, the sedimentation procedure of reinforcement particles in centrifugal fields should be theoretically grasped. For this purpose, the paper applies the continuous theory and builds up the mathematical and physical model for reinforcement particle distribution by considering the influence of alloy viscosity and solidification on particle sedimentation as well as the interrelation of particles.^"^^ Then the author solves this model numerically and simulates the sedimentation procedure with aids of numerical analysis and computer graphics. At last, a SiCp/A356 FGM is prepared to rectify the results of theoretical analysis. 2. Mathematical Model Figure 1 shows the microelements filled with melt of reinforcement particles and alloy liquid in the cenrifugal force field. The initial concentration of this microelement is C and the densities of particles and alloy are Pp and pi, respectively. The gravity is omitted because it is * far smaller than the centrifugal force. Thus, the ^^^' ^ J^^^^ eiawfit reinforcement particles can only move in radius " ^^^ ^^ "*® ^ direction. If the particle concentration distribution is axisymmetrical, it can be considered as the function of radius r and time t. At any time t, in a microelement Ar, the volume concentration of particles is C, and the sedimentary velocity is Uc. The unit volume flux J of reinforcement particles in unit time is

J=CUc

(1 )

According to continuous theory, in time At, the change of concentration in the microelement is

[{J + dl/a-) - j]dt = -{dCla)drdt. Adjusting Equation (2), one can write * School of Material Science and Technology, Harbin Institute of Technology, Harbin , China, 150001

(2 )

54

[{J + dJ/a-) - J]dt = -{dCla)drdt.

(3 )

From Robison^^', the sedimentary velocity of particles with concentration C in centrifugal force field is

Uc=U^x{\-CY

(4)

with Kd^iPp - Or)

^

M where k is the shape coefficient of particles, d is the average diameter of particles, ^ is the viscosity, ps and PL are the specific gravities of particles and liquid, respectively, co is the rotation angular velocity, and r is the centrifugal radius. The change of melt viscosity caused by the increment in volume portion of reinforcement particles (Vp) is /^^P=A(I + | ^ ' . + 7 . 6 F / ] where |ii is the viscosity of alloy liquid. ^^^ Using equation (4) and (5), we may write equation (1) in the form aC o[Ar(l-CyC] —- + ——^^ —^ = 0

ar with

(6)

(7)

aA=

^^-—^-^co\

which is the mathematical model for particle sedimentation. Equation (7) is a nonlinear partial differential equation. According to conditions of convergence and stability, applying against wind differential scheme, it may be discretized as Ari C."-^' ^ ^ « | [ ^ r ( l - C ) ] . -\Ar{\-Cyc\ \ (0
( 10 )

55 which can be discretized as

C;''=C"o +

Ar{l-C,yC,\.

( 11 )

As to the outer element AtN, after the same derivation, we may get ( 12 ) 3.4 Mathematical Model of Solidification Let us still consider the microelement in fig.l, which filled with melt consisting of reinforcement particles and alloy liquid.

The temperature of melt and mold are To and T^, respectively. In order to

produce FGM with particles gradiently distributed in the direction of centrifugal force, the transfer of heat must be in the same direction. Hence, the heat conduction equation can be simplified as

^ ^T

d

( 13 )

dKK' ^%)^^'' dK) • "

dT

where p is the density of alloy, \ is the conductivity coefficient, L is the specific heat of solidification, and gs is the fraction of solid phase. A temperature picking up method is utilized to deal with the solidified specific heat

.

The

second class and third class boundary conditions are prescribed to describe the heat transfer in the inner surface and the outer surface, respectively.

The inner surface is perfectly insulted and the heat transfer of

outer surface can be described as Newton's law of heat transfer. is no interface heat resistance.

Between the casting and the mold, there

To solve equation (13), we can apply direct difference method.

4. Numerical Calculation of Mathematical Model The difference schemes for particles sedimentation in centrifugal force field are shown in equation (8), (11), (12), and the direct difference

format of equation (13). At a certain interval of time, the

particlesconcentration and temperature and the solid fraction of every element are calculated, and then the condition under which the particles stop sedimenting is judged. The sedimentation of particles terminates when (a) the fraction of solid reach to 100% due to the solidification procedure, and (b) the volume fraction of reinforcement particles reach beyond a critical value 0.74 or 74%. 5. Result of Numerical Simulation Fig.2 shows the changes of particle concentration distribution with time. normalized.

All the results are

The reinforcement particles is SiC with the density of S.2g/cm\ and the base alloy is Al alloy.

The viscosity of liquid is 1.19 X 10'^g/(mm.sec), and the centrifugal accelerate is 500g.

The average

concentration Co of reinforcement particles is 20%, and the initial concentration is uniform. Other parameters are selected from reference^^l

( a)

t=30 s

( b ) t=60 s

56

( c ) t=120 s

( d ) t=240 s

Fig. 2 Distribution of reinforcement particle concentration with centrifugal time

From Fig.2., in the case of uniform initial concentration, with the increment of centrifugal time, the concentration distribution of particles is obviously divided into four zones: particles enriched outer zone, uniformly distributed middle zone, gradiently distributed inner transition zone and the non-particle base zone. During process of centrifugal cast, the outer zone expands and its concentration increases, while the middle zone becoming thiimer gradually. The iimer zone moves outside and the base zone expands gradually. At last, after the meh has solidified, three zones - the particles enriched outer zone, the unformly distributed middle zone and the non-particle base zone ~ are formed, which ar^ shown as Fig.2(d). Generally, in the whole process of centrifugal cast, the distribution of reinforcement particle changes distinctively. The concentration curve varies from notching curve to straight line, and fmalUy becomes a convex curve, which is in accordance with and exponential function.

( a ) t=30 s

( b ) t=60 s

( c ) t=120 s

( d ) t=180 s

Fig. 3 Distribution of reinforcement particle concentration with centrifugal time

57 In fig.3, the initial distribution is not uniform while the result is the same as fig. 2. What different are that (a) the uniformly distributed middle zone is not even, and that (b) the changing process speeds up because of the uneven initial concentration. At last, it also develops to three zones~the particles enriched outer zone, the gradiently distributed middle zone and the non-particle base zone. 5.2 Influence of process parameter on distribution Fig.4 illustrates the influences of initial concentration and centrifugal accelerate on particle distribution. The centrifugal time is 180 seconds. When the initial concentration increases, the outer zone become thicker and the concentration distribution become of easier curve. When increasing the centrifugal accelerate G, the speed of sedimention increases and the time to stable state is shortened. The outer zone becomes thinner while its concentration increases.

( a ) 10 % SiC

C a ) G = 300

( b ) 20% SiC

( b ) G = 500

( c ) 30 % SiC

( c ) G^

Fig. 4 Influence of centrifugal process Parameter on particle graidet distribution

6. Fabrication of SiCp/A356 FGM 6.1 Experimental Method To verify the results of numerical simulation, the author designed a set of experiments to research the SiC particle distribution with different centrifugal time. A horizontally rotating centrifuger is applied, shown as fig. 5. The SiC particles, with concentration of 15%, were added through vacuum stirring method. The base alloy was A356. The temperature was 750 °C, and the centrifugal accelerate was 500g. The mold was iron crucible with insulted coating. After the experiment, the sample was cut through radically and axially. The volume fraction and distribution of SiCp particle distribution were determined by optical quantative analyzing technique.

58 6.2 Result of Experiments

±

O«=0(s) ^ « = 120(S) - t = 180(s)

B J^ 20 15(B—o 10 •

(b—<3)—^yg>—^-e—^—e—y

'

NormalizecJ distance

( a ) SiC/A356FGMmacrophotographst=240s

( b ) curve of SiC particle

Fig. 6 changes of distribution SiC particle

distribution

with different time

Fig.6 (a) shows the macrostructures of FGMs produced with different centrifugal time, and fig.6(b) shows the experimental concentration distribution of SiCp particles Which is in accordance with the simulated result perfectly. 7. Conclusion 1. In centrifugal force field, the mathematical model of reinforcement particle distribution is dC_

dr^ J-j,

^Ar{\-CfC\

a= ^/a

J-J^=^/a.

r^Kdjp-p).

)

( l
(i=i) (i=N)

2. The sedimentation process of particles in centrifugal field is dynamic. The static distribution of particles is in accordance with an exponential funtion. Therefore, the particle distribution can be controlled by selecting different process parameters. 3. Under the synphosized boundary conditions, the mathematical model for particle distribution is decretized and solved numerically to obtain the information of influences of different process parameter on concentration distribution. Thus, it is possible to control the particle distribution quantitatively and quaUtatively.

References 1. YJukui eto.Transactions ofthe Japan Society ofMechanical Engineas (C) 56 (1990) 67 2. Fukui, Y, JSME Int, J, Ser. 34-1 ( 1991 ) 144 3. Y.Fukui,Y.Watanabe,Transactions of the Japan Society of Mechanical Engineers ( A ) 58 ( 1992 ) 2472 4. Zhang Baosheng, Foundry vol 218. 3(1995) 6 (in Chinese) 5. T.M. Coulson, Chemical Engineering Volume Two Unit Operation Pergramon Press, 1978 6. Schowalter, W.R, Mechanics of Non-New-tonian Fluids p299 pergamon press 1978 7. Ohnaka , Computer Aided Solidification and Heat Transfer Analysis, 1988.3 8. Qian Miaogen, Metallogy, Shanghai Sceice and Technology Press, 1980 9. Mamoru Mizuno, Toru Matsuoka and Tatsuo Inoue, J.Soc. Mat. Sci, Japan, Vol. 42, No.480, pp. 1046-1051, Sep. 1993

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

59

Modeling and measurement of stress evolution in FGM coatings during fabrication by thermal spray S.Kuroda^> Y.Tashiro^ and T.Fukushima^ ^National Research Institute for Metals, Advanced Materials Processing Division, 1-2-1, Sengen, Tsukuba, Ibaraki 305, Japan ^Department of Materials Science and Technology, The Science University of Tokyo, 2641, Yamazaki, Noda, Chiba 278, Japan 1. INTRODUCTION In thermal spray, molten powder particles are sprayed onto a substrate to form relatively thick coatings of a wide range of materials[l] and the technique has been used to form functionally gradient coatings (FGM) for various applications[2-4]. Residual stresses in thus formed FGM is an important factor which can affect the durability of the FGM under service conditions. We have developed an instrument to continuously measure the curvature of a thin substrate in-situ during spraying to understand the mechanism of stress generation in sprayed coatings. With the understanding acquired through the experiments, a niunerical model has been developed to predict the residual stress distributions in thermal sprayed FGM. 2. EXPERIMENTS AND NUMERICAL MODELING 2.1 Spraying method and conditions NiCrAlY and 8%Y203-Zr02 (YPSZ) powders were plasma sprayed in air under the conditions listed in Table 1. As shown in Fig.l, two torches were operated simultaneously and the powder feed rates to each of the torches were varied as functions of time in order to form FGM coatings, whose typical cross section is shown in Fig. 2. 2.2 Curvature measurement Fig.l also shows the schematic diagram of the instrument to measure the curvature of a substrate during thermal spray[5]. A strip-shaped substrate is fixed onto a pair of knife edges by a pair of springs and the displacement at the center of the rear surface is measured by a displacement sensor. A thermocouple is spot welded to record the thermal history.

60 Table 1 List of spray conditions. Powder material 1 Spray parameters Arc current 1 Plasma gas 1 Flow rate 1 Voltage 1 Powder feed rate x 1 Feed gas flow rate 1 Substrate 1 Spray distance 1 Substrate temp.

NiCrAlY 10 - 44 tim 800A

YPSZ 10 - 44 ^lm 1200A Ar 451/min 30V 0 < X < 10 g/min | 0 < x < 12.5 g/min 2.5 1/min mild steel, 2 x 15 x 100 mm 100mm 470-670K spring

metal

tCU

\

3

B

substrate knife edge

ceramic'

1/R

displ. sensor!

twin plasma torch

Figure 1. Twin plasma torches to spray FGM coatings and the instrument to measure the curvature of a substrate during deposition. 2.3 Modeling The ntmierical model we used was originally developed by Gill and Clyne [6] and has been modified to handle multi-layered deposits. It is a 1-dimensional model and consists of two parts: thermal profile calculation and stress calculation. By regarding the torch motion as a fluctuation in the heat and mass flux onto a reference point on the substrate and assuming biaxial stress state, the program calculates both the through-thickness thermal profile and stress distribution during thermal spray as fimctions of time.

61

Table 2 Thermal and mechanical properties used for modeling. E: elastic modulus, kithermal conductivity, a:thermal expansivity, Gqi quenching stress. NiCrAlY

YPSZ (bulk)

YPSZ (deposit)

190 18

192 4 10 20

19 0.8 8.3 20

E(GPa) k(W/m/K) a(10-6/K) CTq(MPa)

12.2

250

Figure 2. Cross section of FGM coating. *—• 10|im ^

preheat

cool down

spray

^ o.3r — ^ -0.2

;=o.i 2

3 Timelmin]

400r _300 •^200 2

3 Time[min]

Figure 3. Measured curvature and temperature of a substrate during spra3dng of a NiCrAlY coating.

62 3. RESULTS 3.1 Experimental results Fig.3 shows the curvature and temperature records when NiCrAlY was sprayed onto a mild steel substrate. The high-frequency oscillations in the curvature and temperature traces are due to the heating and cooling by the plasma and cooling air jets. From the slope of the gradual curvature change with respect to time, i.e. the coating thickness, quenching stress in sprayed deposits can be evaluated[7]. Quenching stress arises because the thermal contraction of individual splats is constrained by the underljdng solid and is therefore always tensile. The value of quenching stress is independent of substrate material but depends on the substrate temperature. In the case of a substrate temperature around 600K, aq=250MPa for NiCrAlY and 20MPa for YPSZ were obtained. The values of aq as well as the elastic modulus of the deposits with several mixing ratios were also measured too. For the modulus determination, 3-point bending test was performed for deposits detached from substrates. Fig.4 shows the curvature and temperature data when a graded coating was sprayed onto a mild steel substrate by gradually changing the volume ratio of YPSZ from 0 to 1. It is evident that the slope, i.e. the quenching stress, decreases as more YPSZ is mixed. 3.2 Thermal and mechanical properties of sprayed deposits Table 2 lists the thermal and mechanical property data used for the model calculation. The thermal conductivity and elastic modulus of sprayed YPSZ are often an order of magnitude smaller than the bulk values. This is due to extensive microcracks and unbonded gaps between lamellae. 3.3 Model calculation Fig.5 shows the calculated curvature and temperature evolution for an FGM deposit with thickness of about 180|jm, which is consistent with the experimental results shown in Fig.4 except for the transient oscillations. Fig.6 (a) shows the calculated stress distributions in 2-layer and FGM deposits. The gradual stress variation in the FGM can be observed. In Fig.6 (b) effects of model parameters such as the substrate temperature and elastic modulus of YPSZ on the stress distribution in 2-layer deposits are demonstrated. As the substrate temperature is raised from 600 to 825K, the tensile stress in the NiCrAlY layer is significantly reduced. If a value of elastic modulus of 190GPa of a dense bulk material was used, the compressive residual stress in the YPSZ is excessively overestimated. This example clearly demonstrates the importance of using realistic values for modeling thermal and mechanical behavior of sprayed deposits. 4. CONCLUSION To predict the residual stress distributions in sprayed FGM, in-situ curvature measurement is necessary because the values of quenching stress are difficult to

63 preheat

YPSZ vol ratio 0 p400

cool doown

spray

0.1

0.2

0.3

0.4

0.5

0.6

0.7

6

7

8

9

0.8

0.9

1

|ii|iiMM^^^^

3*200'

0 1

2

3

4

5

10

11

12

Time[nnin]

Figure 4. Measured curvature and temperature of a substrate during spraying of a NiCrAlY/YPSZ FGM coating.

Time (min) ^uu

:

—

I

-

—

\

—

'

—

i

^

—

•

—

*

—

1

—

•

—

-

'

'^~

T

•

^

'

_ _^^^^^^^^ ^

._ _

300

'

1 1 t

U

j j 1

•

200

V

100 n

j

*-. ^

*

•

•

•

*»

1 -

1

1

1

1

1 — 1 — 1 .

8 Time (min)

•

J

1

•

n

•

•

1

12

Figure 5. Predicted curvature and temperature of a substrate during spraying of a NiCrAlY/YPSZ FGM coating.

64

200

l<-

Substrate Deposit

S

2-layer deposit FGM (91ayers)

100

li

CT)

2.0

1.0 Position (mm)

(a) 200

T

1

•

I

-r-

I

I

I

I

I

I

I

I

•

Ts=600K,Ed= 19GPa Ts=825K, Ed= 19GPa — ^ Ts=825K, Ed=192GPa

100

i

>-^

-^""""""lA^^

^ -100 -200 0.0

0.5

1.0

1.5

2.0

Position (mm) (b) Figure 6. Calculated residual stress distributions in deposits, (a) 2-layer deposit and 9-layer FGM, (b) effects of the substrate temperature and elastic modulus of the YSPZ deposit.

obtain theoretically. Proper evaluation of thermal and mechanical properties of sprayed deposits is also important for modeling because these can be remarkably differed from the values of bulk materials. Using numerical models, the influence of the property data on the stress distributions can be studied easily. REFERENCES 1. L.Pawlowski, The Science and Engineering of Thermal Spray Coatings, Wiley, Chichester, 1995. 2. T.Fukushima, S.Kuroda and S.Kitahara, Proc. 1st Int. Symp. FGM, Sendai, Oct., 1990, p. 145. 3. N.Shimodaet al., i6irf, p.l51. 4. H.StefFens, M.Dvorak, and M.Wewel, ibid, p. 139. 5. S.Kuroda, T.Fukushima and S.Kitahara, Thin SoHd Films, 164(1988)157. 6. S.C.Gill and T.W.Clyne, Thin Solid Films, 250(1994)172. 7. S.Kuroda and T.W.Clyne, Thin Sohd Films, 200(1991)49-66.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

65

Artificial neural network used for TiB2-Cu FGM design Z.C.Mu', Z.XWang', W.B.Cao', C.C.Ge' ^School of Automation and Information Engineering, University of Science and Technology Beijing, Beijing, 100083, P.R.China ^Laboratory of Special Ceramics and Power Metallurgy, University of Science and Technology Beijing, Beijing, 100083, P.R.China As FGM is a kind of complex heterogeneous composites, the underlying relationship among various factors (composition, constitution, processing parameters etc.) affecting its properties such as density, hardness, strength, fracture toughness, hot shock resistance, elastic modulus and grain size etc. of the bulk or certain layer of FGM are subtle, complicated and very difficult to describe as formulas or rules. Artificial neural network (ANN) does not require any explicit knowledge rules to construct, and can learn by itself to form "mapping" from inputs to outputs. In this work, ANN is first applied for density estimation of the FGM. 1. INTRODUCTION Knowing the effective material properties is essential for the FGM design especially in its earlier stages. Traditionally these properties are obtained mainly by using some classical rules of mixture. More recently some papers have presented the methods that use Expert System and/or Fuzzy Logic as alternatives for the property estimation. However, the FGM is a kind of complex heterogeneous composite, its properties are affected by various factors such as composition, constitution, and processing parameters etc. The underlying relationship among these factors are subtle, complicated, and very difficult to find and describe as formulas or rules. To solve these problems which lack existing solutions, we have applied the Artificial Neural Network (ANN) technology into property estimation of the FGM design. The result from this trial looks quite promising. 2. BACKGROUND OF THE ANN The ANN is mathematical model that simulates many characteristics of actual neurons in the brain. Generally, an ANN is a structurally multi-layered network which links a large number of nodes (the neuron-like computational elements) and operates dynamically. Although mathematical neurons were conceived as early as 1943, only recently have largescale real-world applications become practical. Unlike rule-based Expert System and Fuzzy Logic, the ANN can find relationships among the inputs and outputs of the network without the need for training by an expert. It does not require any explicit knowledge rules to construct. It applies a weight-adaption algorithm to a representative sample of training data to correlate inputs with desired outputs, i.e. it can learn by itself to form "mapping" from inputs to outputs. Another advantage of the ANN over Expert System is that it has the ability to generalize — to produce a reasonable response to data for which the system has not been explicitly trained. In a rule-based Expert System, if a situation occurs for which there is no applicable knowledge rule, it may be unable to give any response.

66 In short, the motive that we have applied the ANN into the FGM design was inspired by the ANN'S distinct characteristics-its noniinearity, ability of self-learning, and ability of generalization. 3. APPLICATION TO THE FGM DESIGN According to our application purpose, the development of the ANN can be described as follows: If defme x^eK" (i=l,2,...,n) as the network inputs which represent the factors that have significant influence on the properties of the FGM and yjGR"" (j=l,2,...,m) as the network outputs which represent the estimated properties, the goals of the work are to choose the network architecture, train the network with data from real experimental process, and make the network act mathematically as an underlying function Y=F(X) which maps network inputs to network outputs. density

hardness

gram size output layer hidden layer

Ti

input layer pressure

B particle size Fig. 1 The Architecture of ANN Network

Development of the ANN for the FGM design consists of two phases. In the first phase, or the training session, an ANN network is trained for the property estimation after its network architecture is determined. The architecture of the network is shown in Fig. 1 . It is a layered, feed-forward network. A supervised learning algorithm — Back-Propagation Algorithm (BPA) is used for the training. The algorithm compares the calculated outputs of the network to the expected outputs, and readjust the weights in the network so that the next time that same input is presented to the network, the network's output will be closer to the expected output value. The training error is defined as: 1

(1)

^^ k=\ f=i

where m is the number of output nodes; n is the number of training data sets; T^^, is the expected output value of the network which is measured from experimental process; and Oi,, is the calculated output value of the network. The network is trained step by step according to the equation : AW(t) = -y\VE{t) + a A ^ ( / -1)

(2)

where the arguments t and t-1 are used to indicate the current and the most recent training step

67 respectively, W represents the weights which interconnect the network nodes, r| is the learning constant, and a is an user-selected positive momentum. During the training session, the equation (2) is applied iteratively until convergence of the calculated and expected outputs. In general, the specific architecture of the network and the optimum value of y\ and a depend on the problem being solved, and there is no single solution suitable for different applications. Therefore, the number of hidden layers, the number of hidden nodes, learning constant r\ and momentum a must be chosen experimentally for the application during the training session. Experimental results have shown that the effectiveness and convergence of the BPA depend significantly on these factors. In the second phase of the development, the ability of property estimation of the network is examined with the data samples which are measured from the experimental process but never used in the training session. Based on the experimental data which were available for the trial, we trained and tested the ANN network to predict densities of FGM. Table 1 shows the comparison of measured and the network estimated densities of as-synthesized TiB2-Cu FGM using the Self-propagating High-temperature Synthesis (SHS) technique. The result indicates that the ANN network can give quite good estimations to the FGM properties. Apart from density, other properties of FGM such as hardness, strength, fracture toughness, hot shock resistance, elastic modulus and grain size etc. of the bulk or certain layer of FGM can be predicted and designed.

Table 1 The Comparison of Measured and BPA Estimated Densities estimated values Samples measured values 1 79.00 80.20 70.83 2 71.33 63.95 3 64.40 80.93 4 79.10 63.74 5 63.70 In order to further improve the ability of generalization of the network, a more complex ANN network was developed. This network combined two companion networks using the Double Back-Propagation (DBP) algorithm which improves the network performance by forcing the output to be insensitive to incremental changes in the input. The improvement is quite significant. The mean estimation error of the network decreases by 30% compared with the network using the BPA[3]. 4. CONCLUSIONS The ANN is an artificial intelligent technique that has several distinct advantages over rulebased Expert System and Fuzzy Logic. The technique has been shown to be a feasible technique that estimates the material properties for the FGM design. The estimation accuracy is satisfactory.

68 5. ACKNOWLEDGMENT This work is supported by China National Natural Science Foundation, The Doctorial Program Fundation of State Education Commission, and National Committee of High Technology New Materials. REFERENCES 1. T.Hirano et al. Proceedings, The First International Symposium, FGM, Sendai, (1990) 5. 2. T.Hirano et al. International Workshop on Artificial Intelligence for Industrial Application, (1988)245. 3. Z.C.Mu et al. Pattern Recognition and Artificial Intelligence, Vol. 8, No. 1 (1995) 51. 4. H.Drucker and Y.L.Cun, IEEE Trans. Neural Networks, Vol. 3, No. 6 (1992) 991.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

69

Deformation analysis of graded powder compacts during sintering K. Shinagawa Department of Mechanical Engineering, Anan College of Technology Minobayashi, Anan, Tokushima 774, Japan

A method for simulating the deformation behavior of graded powder compacts during sintering is proposed. A constitutive model for sintering of metal and ceramic powder mixtures is presented. The constitutive equation is applied to the viscoplastic finite element method. The shape change in graded powder plates during sintering is predicted. The effects of sintering properties of the powders and their combinations on warpage in the sintered plates are examined.

1. INTRODUCTION Powder metallurgy is one of the methods to produce functionally graded materials (FGM) with metals and ceramics. In general, shrinkage behavior of ceramics during sintering is different from that of metals. Difference in shrinkage rate often induces distortion or cracks in sintered bodies. Sintering process should be controlled to avoid such defects. Powders may be coordinated by selecting particle size or adding elements to get almost homogeneous shrinkage rate all over the layers in FGM powder compacts [1,2]. In another way, graded heating may be imposed to the powder compacts to adjust the sintering balance of the layers [3]. Conditions of the process, however, seem to be determined by trial and error. Thermal stress in FGM has been studied because it is important to avoid the fracture due to thermal shock. However, the defects in the powder compacts during sintering is caused by sintering stress, which arises from surface tension. Therefore, the structure of FGM made by sintering should be designed with both thermal stress and sintering stress taken into consideration. Analysis of the sintering process of FGM powder compacts has not been conducted. The author has proposed a constitutive equation for sintering of a single powder and applied it to the prediction of deformation behavior of powder compacts during sintering using the viscoplastic finite element method [4-6]. In this paper, a constitutive model for sintering of metal/ceramic mixed powder compacts is developed. The shrinkage behavior of metal and ceramic powder mixtures is expressed on the basis of the sintering behavior of each component powder and the mixing ratio. The constitutive equation is applied to the finite element method and the deformation behavior of cylindrical graded powder plates during sintering is simulated. The effects of the balance of sintering properties of two component powders and the size of powder compacts on the warpage in the graded powder plates are examined.

70 2. CONSTITUTIVE MODEL FOR GRADED POWDER COMPACTS 2.1. Constitutive equation for sintering Sintering can be described as the deformation process of viscous porous bodies under the action of the sintering stress, which is an apparent hydrostatic stress [4-6], Based this theory, the constitutive equation for sintering is given by = _1_ en = 2rj (1) hj=x,y,z 9/^ «=2.5, f=l/2.5fi7p, where rj is the viscosity, p the relative density, ^ , rr' are the strain rate and deviatoric stress, respectively, 6 is the Kronecker symbol, Om is the hydrostatic stress, as is the sintering stress. The sintering stress is evaluated from surface tension and pore geometry as follows;

^irl^'^^^^V^'^'^^H

as=p^-y, 1P(1-PO)I

(2)

'

N=6, A:=0.5, po=0.52,

where y is the surface tension,} is the effective pore radius, R is the radius of a powder particle, po is the initial relative density. If the mechanism of material transport in sintering is grain boundary diffusion, the viscosity can be expressed as a function of the temperature and the grain size. From a model for boundary diffusion controlled creep M2 1 proposed by Coble [7] and the temperature 70 C2 dependence of diffusion coefficients, we obtain 4U}

rj=ciT'exp(f\d\

where T is the temperature, d is the grain size and ci, C2 are constants. The grain growth during sintering is assumed to be given by d^=do^^C3t'exp(^],

I

/Ml 10 /

C/3

5

(4)

do=2R, where do is the initial grain size (= the diameter of powder particles ), t is the time (sec) and cs, CA are constants.

^Cl

15

(3)

n

j

/

j

-

C/ 1 i 600 800 1000 1200 1400 0.5 1.0 Temperature /°C Holding time /h Figure 1. Shrinkage curves.

2.2. Sintering behavior of metals and ceramics powder compacts Figure 1 shows shrinkage curves of metals Ml(Mo), M2(Mo with 0.1mass%Ni) and ceramics Cl(Zr02), C2(PSZ) calculated by Eqs. (1)'^(4), which are employed as components of FGM. The rate of heating up is 100°C/min. After the temperature reaches 1400 °C, the sintering is carried on at the same temperature for one hour. The material parameters used in Eqs. (1) ^^(4) are shown in Table 1. These component materials were referred to the research on molyb-

71 Table 1 Material constants. Metals

Ml

?;[Pa-s]

1.5628 x l O '

16415 d[m]

1.2794x10"'"

Ceramics

M2

7.3919x10'' 15426

CI

C2

5.4445x10'' 1.4454x10'' 47921

40043

1.2569x10"" 4.6148x10"' 2.5493x10"'

R[m]

-34804 1.0x10"^

-29041 1.0x10"^

-72213 0.025x10"^

-65313 0.025x10"^

)4mN/m]

1000

1000

1000

1000

denum / zirconia system by Watanabe et al. [1,2] and the parameters ci'-C4 were determined from their experimental data by the method of least squares. 2.3. Constitutive model for metal and ceramic powder mixtures Let us express the shrinkage behavior of powder mixtures MCI composed of Ml and CI, and MC2 composed of M2 and C2 by the shrinkage curve of each component shown in Figure 1. It is assumed that there is no reaction between the metals and the ceramics. The powder mixtures are described as two porous bodies mechanically engaging with each other. Using this assumption, the sintering stress of the mixtures may be given by Osmc=VmOsm-^VcOsc^

(5)

where Vmy Vc are the volume fractions of metal and ceramic, respectively (Km +Fc =1), Osmy Osc are the sintering stresses of metal and ceramic, respectively. When two kinds of particles with different diameters fill up a container, the bulk density may change with the ratio of mixture. Variation of the initial relative density with the volume fraction of ceramic is set as shown in Figure 2. The relative densities of metallic and ceramic parts of powder compacts are assumed to be given by p,=((1-a)Fc+a>-p , a=0.8, (6) where a is a parameter to express the reduction in relative density due to mixture. The viscosity of the mixtures is assumed to be evaluated by m=^mrjm-^Vcric, m

rjm rjc'

20 40 60 80 100 Volume fraction of ceramic /% Figure 2. Variation of initial relative density with ratio of mixture of metal and ceramic.

72

25 20 ^15

1400°C^

u

x^oo°c 1

I 10 c/5

5

1200°C

k

r O ^ — /\• _ 1 0

1

1

iioo°c[ &—-

0 -^ lOOO^C 1

1

1

20 40 60 80 100 20 40 60 80 100 Volume fraction of ceramic /% Volume fraction of ceramic /% (b) MC2 (a) MCI Figure 3. Variation of sintering behavior with ratio of mixture of metal and ceramic.

MCI: ^=1, MC2: y5=.0.0625+0.863 V^ The actual shrinkage behavior of the mixtures was also referred to the research of Watanabe et al. [1,2] and the initial relative densities in Figure 2 and the parameters a, fi were determined from their experimental data. Figure 3 shows shrinkage behavior of powder mixtures calculated by Eqs. (1)'^(7) with the heating rate of 100°C/min. The sintering balance of MC2 is better than that of MCI. The shrinkage does not follow the law of mixture, that is, the minimum shrinkage is obtained at the low volume fraction of ceramic [1,2]. This is considered to be because of the increase in the average radius of pores due to mixture. This phenomenon was expressed as the reduction in relative density in Eq. (6) in the present work. 3. FINITE ELEMENT ANALYSIS OF SINTERING PROCESS 3.1. Theory The finite element method with the sintering stress taken into account has been formulated [4,6]. In this case, the equilibrium equation is given by dXi

dXi

'

hj=x,y,z

(8)

where Sij is the sintering stress. From the principle of virtual work, the nodal forces for each element are given by

{P)=\[BY{o)dV^\

[B7{S)dV

(9)

73

Figure 4. Model of graded cylindrical powder compacts.

{a} = {ajc Oy Oz x^y Tyz T^Y where [B] is the matrix correlating the strain-rate components with the nodal velocity components. The stresses are given by (10) {a) = [D]{e) = [D][B]{u) s] = { ex

£y £z Yxy Yyz Yzx]

[D] = 3r/mao2''-l 0 0 0 0

0 0 0 c 0 0

0 0 0 0 0 0 0 0 c 0 0c

2c = i 9' 3 where {u} is the nodal velocity. Substituting Eq. (10) into Eq. (9) gives

a=f^^,b=f-

•.j[BnD][B]dV[u]^j[ {P]=\[BnD][B]dV{u]^\[BY{S)dV

(11)

Adding together the nodal forces of the surrounding elements at each nodal point gives simultaneous linear equations. The nodal velocities are obtained as solutions of these equations. 3.2. Sintering simulation Deformation of two types of graded cylindrical powder compacts having different diameters as shown in Figure 4 were analyzed in the same sintering conditions as Figure 1 (section 2.2). Figure 5 shows calculated distorted grid patterns of the sintered cylinders with the small diameter. The shape of MCI is not strait because the shrinkage of metal Ml is much smaller than that of ceramic CI. The middle of cylinder MC2 with the small diameter is slightly swelled because the initial density of this part is high and the shrinkage is small. However, a sound shape may be obtained in MC2 due to the good sintering balance of the component layers.

74 Figure 6 shows calculated shapes of the sintered cylinders with the large diameter. MCI, having an inferior sintering balance, is considerably warped because a difference in shrinkage exerts more significant influence on a large size. Even MC2 suffers from some warpage. During the process of MCI heating up, the cylindrical plate MC2 is warped once and deformed again in the opposite direction. This may produce a defect within the sintered body. These results suggest that the MC2 size effect should be taken into consideration for suitable design of FGM. 900X

1200X

ISOO'C

1400*C

4. CONCLUSIONS A constitutive equation to simulate the deformation of the graded powder compacts during sintering was proposed. The sintering stress of each layer in the powder compacts was evaluated from powder properties. The constitutive equation with the sintering stress was applied to the prediction of the deformation behavior of graded powder plates. Warpage in the plate during sintering was predicted by the viscoplastic finite element method.

UOOV (Ih)

Figure 5. Distorted grid pattern in cross section of cylindrical FGM with small diameter.

izoo'C 1300X:

MCI

MC2

Figure 6. Warpage in cylindrical FGM with large diameter.

REFERENCES 1. R. Watanabe, A. Kawasaki and N. Murahashi, J. Asso. Mater. Eng. Resour., (in Japanese), 1 (1988) 36. 2. R. Watanabe, Micromeritics, (in Japanese), 33 (1989) 76. 3. M. Yuki, M. Toshikazu, I. Toshio, A. Kawasaki, R. Wtanabe, J. Jpn. Soc. Powder Metall., (in Japanese), 37-7 (1990) 929. 4. K. Shinagawa, Trans. Jpn. Soc. Mech. Eng., A, (in Japanese), 62-539 (1996) 240, 246. 5. K. Shinagawa, JSME Int. J., A, 39-4 (1996) 565. 6. K. Shinagawa, Proc. 3rd Asia-Pacific Symp., AEPA'96, (1996) 439. 7. R. L. Coble, J. Appl. Phys., 34-6 (1963) 1679.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

75

Simulation of the Elasto-Plastic Deformations in Compositionally Graded Metal-Ceramic Structures: Mean-Field and Unit Cell Approaches* H.E. Pettermann^, E. Weissenbek"^ and S. Suresh^ ^Institute of Light Weight Structures and Aerospace Engineering, Vienna University of Technology, Austria, Gusshausstr. 27-29/E317, A-1040 Vienna, Austria ^Department of Materials Science and Engineering, MIT, 77 Massachusetts Avenue, Cambridge, MA USA 02139-4307 1. I N T R O D U C T I O N The elasto-plastic behavior of a compositionally graded metal-ceramic structure is investigated. The deformation under uniaxial loading is predicted using both an incremental Mori-Tanaka method and periodic as well as random microstructure extended unit cell approaches. The latter are able to give an accurate description of the local microfields within the phases. Furthermore, the random microstructure unit cell model can represent the interwoven structure at volume fractions close to 50%. Due to the high computational costs, such unit cell analyses are restricted to two-dimensional considerations. Computationally less demanding mean-field methods provide a tool to account for the out-of-plane constraints, but have the disadvantage of using phase averaged stress and stain fields. In the present work, an incremental Mori-Tanaka approach is employed, which is implemented as a constitutive material law in a finite element code. Both twodimensional and three-dimensional investigations are performed and the results are compared to the predictions of the extended unit cell approaches. 2. M O D E L I N G OF T H E F G M - S T R U C T U R E The investigated structure (cf. fig. 1) consists of a pure nickel top layer (1.2mm thick), an FGM zone with a linear variation of the volume fraction (2.2mm thick), and a pure alumina bottom layer (0.45mm thick). Material data of the constituents are given in table 1, where Eh denotes the strain hardening modulus. The boundary conditions force all cross sections perpendicular to the monolithic/FGM interfaces to remain plane, i.e. the lefthand and the righthand side (as well as the viewing plane for three-dimensional considerations). The uniaxial load (which causes extension and bending for the present *This work was supported by the Grant DE-FG02-93ER45506to MIT from the US Department of Energy. The post-doctoral study of EW at MIT was supported by an Erwin Schrodinger Fellowship from the AUSTRIAN NATIONAL SCIENCE FOUNDATION. HP'S visit to MIT was supported by a scholarship for

Overseas Scientific Study from the AUSTRIAN FEDERAL MINISTRY OF SCIENCE, TRANSPORT AND ART. ^present address: BMW Entwicklungszentrum Steyr, Hinterbergerstr. 2, A-4400 Steyr, Austria

76 Table 1 Material data for alumina '* and nickel ^ [GPa] 380 0.25

[GPa] [MPa] 214 0.31 148

[MPa] 668

case) is applied centered at and perpendicular to the left- and righthand side, respectively. All calculations are performed with the finite element code ABAQUS [1]. 2.1. Extended Unit Cell Models Periodic Packing In the field of metal matrix composites a number of models are used to characterize both, macroscopic behavior as well as local microfields. The present section deals with plane models, the three packing cases shown in [2] are used as the basis to extend the unit cell method for graded structures with hexagonal packing, square edge packing^ and square diagonal packing [3]. Nine subcells with different volume fractions (corresponding to linear grading) are combined to represent the graded part (fig. 1, left and middle). Due to the fact that these models can only represent a pure matrix/inclusion structure, a switch from nickel-matrix to alumina-matrix is necessary. To investigate the particular influence of the eflfective matrix phase for the center sublayer two diflFerent cases are considered, viz. the 50% layer shows a nickel-matrix (marked as Ni) or an alumina-matrix (marked as A1203). Random Micro structures Usually the microstructure of an FGM does not show a matrix/inclusion type topology throughout the thickness. Typically, in the range between 30 and 70 % volume fraction an interwoven type microstructure exists. To account for this, a further two-dimensional plane stress approach is used where the grading within the FGM layer is divided into 11 sublayers each having a fixed volume fraction. In each row the proper number of hexagonal "grains" are assigned to the nickel and alumina phases, according to the volume fraction. The locations with the pertinent row are chosen randomly (fig. 1, right). The bottom and top layers, consisting of pure nickel and alumina, respectively, show the same finite element mesh topology, viz. each grain comprises six triangular elements. Eight diflFerent randomly generated microstructures are studied. 2.2. Incremental Mori-Tanaka Method (IMT) As a complementary approach a mean-field method is used in combination with the finite element method to investigate the FGM. To compare the predictions with the periodic unit cell simulations, the FGM part is divided into nine sublayers. Each of them consists of two bi-quadratic 8-node plane elements over the thickness, and only one element is used in the horizontal direction. The parts of pure alumina and pure nickel are modeled by three and 12 elements, respectively. In each sublayer the volume fractions of the phases are constant. In addition, the center sublayer can be split into a metalmatrix and a ceramic-matrix half sublayer. The properties of the particular material on the meso-structural level within the finite element calculation are described via a con-

77

Figure 1. Extended unit cell finite element models; hexagonal arrangement and detail of center subcell (left); square edge arrangement and detail of center subcell (middle); random arrangement (right)

stitutive material law. This is an incremental Mori-Tanaka approach [4] (representing an matrix/inclusion topology) which is based on Eshelby's equivalent inclusion method. Herein, the Eshelby tensor is a function of the inclusion aspect ratio and the (instantaneous) material properties of the matrix phase. For a general behavior of the matrix material (i.e. elasto-plasticity or anisotropic elasticity) the Eshelby tensor is calculated numerically, based on an implementation developed by Gavazzi/Lagoudas [5]. The IMT method is implemented as a "user supplied material routine" [4] for the finite element package ABAQUS [1]. The mesoscopic, instantaneous material data of the microscopically heterogeneous materials are calculated individually for each integration point at each increment within the finite element analysis. In contrast to the extended unit cell models, the IMT-method does not provide detailed information about the micro-fields. The computational requirements, however, are much lower. Furthermore, this method can be employed to predict the complete deformation of a three-dimensional graded layer and to perform structural analyses for components made from composite materials. Within the present method a pure matrix/inclusion type micro-topology is assumed (no interwoven structure can be represented). Hence, a sudden change of the properties occurs, when the matrix material changes from alumina to nickel. The IMT-method as implemented performs a fully three-dimensional analysis on the micro level independently of the global finite element analysis option.

78

IQ

0)

U

CQ

•H

H

IMT reduced constr. (Ni) IMT reduced constr. (A1203) Square-diagonal (Ni) Square-diagonal (A1203) Square-edge (Ni) Square-edge (A1203) Hexagonal (Ni) Hexagonal (A1203) 2.0

Axial Strain (%)

Figure 2. Predicted axial strain response for the extended unit cell models and the corresponding incremental Mori-Tanaka approach (IMT); the matrix of the center sublayer being nickel (Ni) or alumina (A1203)

3. RESULTS The deformation is given in terms of axial strain and bending strain which are the mean value of the strains in loading direction and in proportion to the curvature of the layered structure, respectively. The global behavior of the periodic packing models with respect to the axial strain is displayed in fig 2, considering both an alumina and a nickel center sublayer. (The responses with respect to the bending strains are equivalent.) The square diagonal packing and the hexagonal packing show similar response, but the square edge packing model is stiffer. The results coincide with the findings in [2,6,7] for equivalent periodic packing arrangements. The overall stress vs. axial and bending strain prediction of the random microstructure models are displayed in fig 3. A slight variation for different randomly generated models can be observed; the global behavior, however, does not deviate markedly. Compared to the periodic unit cell approach the random microstructure models show a more compliant behavior. Mainly responsible for the overall stiffness of the FGM is the portion of interconnected alumina matrix. The periodic unit cell results are directly comparable to the IMT predictions, because both approaches represent the same matrix/inclusion type microstructure. However, such comparisons have to be done carefully since some assumptions regarding the finite element calculations are not equivalent for the extended unit cell approaches and the present mean-field method. The plane stress analysis of the unit cell models does not take into account the constraints in the out-of-plane direction. In contrast, within the present IMT formulation the inclusions are enclosed three-dimensionally by the matrix material. In contrast to the plane stress unit cell models, the constraint in the out-of-plane direction is accounted for. Accordingly, these predictions are denoted as full internal constraint. To overcome this internal constraint in order to "simulate" the plane stress model assump-

79 J

1

1

1

1

_i

1

1

L

(0

0) 0)

4.)

a^

"^St^^

•H

X

(d •d 0)

•H

n

0.0

^

1

4.0

,

1

8.0

^

Arrangement Arrangement Arrangement Arrangement Arrangement Arrangement Arrangement Arrangement 1 ' 12.0

8 7 6 5 4 3 2 1 r

16.0

Axial/Bending strain (%)

Figure 3. Predicted axial (stiffer) and bending (more compliant) strain response for eight different random microstructure models tions, the inclusions are introduced as infinite fibers in out-of-plane direction and the elastic material parameters of the ceramic phase in out-of-plane direction are reduced by several orders of magnitude. Such models are denoted as internally reduced constraint. A center layer with alumina matrix and nickel matrix, respectively, is modeled. It is well known that incremental mean-field methods show a tendency to over-predict the hardening behavior, but the general features are captured in a qualitative as well as quantitative way (cf. fig. 2). Additional calculations are performed where the center sublayer is subdivided. These results and the stress-strain behavior of a global plane stress and full internal constraint analysis (considering spherical inclusions and unmodified material data throughout) are displayed in fig. 4. A marked stiffening as the result of the full constraining on the micro level can be observed. Furthermore, the predictions of a generalized plane strain analysis are shown. This way, mesoscopic interaction between the sublayers is accounted for and a second set of strains can be calculated, describing the axial strain in the out-of-plane direction and the out-of-plane rotation (see fig. 4). These results show that a plate under the assumed loading is deforming to a saddle shape. 4. CONCLUSIONS The deformation behavior of a compositionally graded metal-ceramic structure has been investigated by numerical and (semi)analytical simulations. Random microstructure models are able to predict the response of an FGM-structure in a more accurate way than the other approaches. The interwoven structure in the "middle" of the FGM can be accounted for using this modeling strategy. For the extended periodic unit cell models the predicted stress strain response depends strongly on the micro-arrangement of the inclusions. Detailed information on the microfields of the stresses and strains can only be obtained by the extended unit cell models. The incremental Mori-Tanaka method

80 1

,

I

i

.

I

.

I

0)

u

GPE GPE GPE GPE PST PST PST PST

'

Id 0) •H H

& 1

'

1 -0.5

1 0.5

b e n d i n g / o u t of p l a n e a x i a l / o u t of p l a n e bending/in plane axial/in plane full c o n s t r . red.constr. r e d . c o n s t r . (Ni) r e d . c o n s t r . (A1203)

|

'

1

1 1.5

•

Axial/Bending strain (%)

Figure 4. Predicted axial strain response for plane stress (PST) calculations with internally full and reduced constraint and predicted axial and bending strain response for interally fully constraint generalized plane strain (GPE) calculations by the incremental Mori-Tanaka approach

shows a tendency to over-predict the hardening behavior, but, it is numerically less intensive. In addition, the three-dimensional interaction between matrix and inclusions can be considered, and it can be employed to predict the complete deformation of an FGM-structure. Acknowledgments The authors thank Drs. H. J. Bohm, M. Finot, A.E. Giannakopoulos, A. Needleman and F. G. Rammerstorfer for many helpful discussions during the course of this work. REFERENCES ABAQUS User's Manual; HKS Inc., Pawtucket, RI (1994). T. Nakamura, S. Suresh.: Acta metall.mater. 41, 1665 (1993). E. Weissenbek, H.E. Pettermann, S. Suresh: submitted to Acta.mater. H.E. Pettermann, H.J. Bohm, F.G. Rammerstorfer: In "Proceedings of the General Workshop in Materials Science and Engineering", 1996, Ed: M. Rappaz, European Commission, in press. A.C. Gavazzi, D.C. Lagoudas: Comput.Mech. 7, 13 (1990). E. Weissenbek., H.J. Bohm, F.G. Rammerstorfer: Comp.Mat.Sci. 3, 263 (1994). E. Weissenbek: VDI Fortschrittsberichte (Reihe 18, Nr. 164), VDI Verlag, Diisseldorf, Germany (1994).

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

81

Large Deflections of Heated Functionally Graded Clamped Rectangular Plates with Varying Rigidity in Thickness Direction F. Mizuguchi^^ and H. Ohnabe**^ ^Dept. of Mechanical Engineering, Japan Maritime Safety Academy, 5-1, Wakaba-cho, Kureshi, Hiroshima-ken,737 Japan. ^CompositeNfeterials Center, AeiD-Engine & SpaceOperations, Jshikawajima-Haiirna Heavy Inciustries Co., lid, 3-5-1, Mukodai-cho, Tanashi-shi, Tokyo, 188 Japan. The governing equations for the large deflection of heated functionally graded elastic, rectangular^plates with varying Young's modulus in thickness direction are derived by Berger and K i m ^ approach. We assume that the functionally graded Young's modulus varies symmetrically with respect to the middle plane according to the nth power of the nondimensional thickness. For a clamped rectangular plate affected by an arbitrary symmetrical temperature and load distribution about midpoint of the plate and prevented from inplane motions on the boundary, the governing equations are solved by means of the Galerkin method. As a numerical example, the cases as seen in aerodynamic heating are analyzed. The influence of functionally graded Young's modulus and temperature change on the large deflection are shown in the graphs. 1. INTRODUCTION Recently, the functionally graded materials (FGM) of the thermal relaxation type, adaptable to a super-high-temperature environment like a super and hyper sonic transportation and a space plane, have received considerable attention. An approximate method for investigating the large deflections of isotropic plates has been proposed by Berger [1] in 1955. Since then much research has been conducted on large deflections by Nowinski [2] in 1962 and on thermal buckling by Kamiya [3] inl976. Nowinski and Ohnabe [4] in 1972 pointed out that the relative exactness of the Berger method is closely associated with the restriction imposed on the in-plane displacement of the boundary conditions. The mathematical justification for this statement has been given by Schmidt [5] in 1974 using a proper perturbation method. Recently Horibe [6] in 1990 proposed a new iterative solution of the Berger equation. The solution is derived by utilizing both the idea of the Kantrovich method and the boundary integral equation method. Following Berger, Ohnabe and Mizuguchi [7,8] derived the governing equations for the large deflection and the non-linear vibration of heated non-homogeneous circular plates with radially varying rigidity. They were solved for the clamped boundary condition by using the Galerkin method. The governing equations for large deflections of heated functionally graded elastic, rectangular clamped plates with varying Young's modulus in thickness direction were also derived by Berger and Karman approaches. Assuming the functionally graded Young's modulus in thickness direction symmetrical with respect to the middle plane, for a rectangular simply supported plate due to a temperature distribution as seen in aerodynamic heating, they are solved by employing the Poincarc method [9]. In the present study, appling the same procedure, for a case of a rectangular clamped plates, they are solved by employing the Galerkin method. The influence of functionally graded Young's modulus and temperature change on the large deflection are shown in the graphs. The solutions for Berger and Karman approaches are compared.

82

2. GOVERNING EQUATIONS We consider a functionally graded, elastic, heated rectangular plate, with a lateral load q. Let the deformation be axisymmetric with respect to the center .We assume that the functionally graded Young' s modulus varies symmetrically with respect to the middle plane according to the nth power of and is given by According to Berger, eliminating the second strain invariant and applying the Euler-Lagrange variational principle for minimum potential energy, the Euler-Lagrange differential equations become the following governing equations for large deflection of heated functionally graded elastic, rectangular plates with varying Young's modulus in thickness direction [9] B V^ (V^w) + a( 1 + V ) i | (V^nu, + pV^m ') - K^v\ h^

- -5L ^0

(2) (3)

where fh/2 thil B = l + - ^ ^T-, , n u ,- - |I Tzdz,PT,= Tdz nu, Tzdz , PT = I| n.-3 ^ |_^2 T j_^^

1+-^, n+1 h/2 tV2

rh/2 tY\l2

1

h/2 equations with large ^^W.h/2 The von Karman-type deflection can also be12(1derived from the equations of equilibrium and compatibility referring to[9]

^ V ^ ^ ^ r ^ (^^"h-+P^^"^T')=^^'p) ^ ^ V"^ + aEQ ( v \ + PV^P^') = - 1 AEQh^w, w)

(4)

(5) ^^^

where the operator L applied to the functions w, F is 3x2 3y2 "^ ay2 9x2 ' ax9y dxdy

^^^

3. EXAMPLES OF SOLUTION OF GOVERNING EQUATIONS The approximate expression for w is selected directly from the linear theory of plates with small deflections, and is given by w(x, y) - f sin^ axsin^ Py /gx where f is the maximum deflection and ^

a,

*^ b

Boundary conditions for a plate clamped on all edges are, at X - 0, a

at y - 0, b

(9)

83

w

3w

^ 0

w

3w

^ 0

ax ay (10) The temperature distribution over the plate is assumed to be symmetrical with respect to the center of the plate and is given as oo oo

I

Px = h i: E Tij COS iax cos jpy ijieven Pj = h S E Tij cos iax cos jPy ijreven and the load is also given as

oo oo

I

m j - h^ | S Z Tpq sin pax sin qPy + Tg (p,q:odd ) m j ' = h^ ! S S Tpq sin pax sin qPy .+.T^ ;! \p,q:odd (11)

q = X S qij sin iax sin jPy ijiodd Substituting Eqs. (8) and (11) into Eq. (3) and integrating with respect to x and y, K^ is obtained _

(12)

Obtaining a general and a particular solution of Eq. (2) and using the boundary conditions of Eq. (10), we then determine the integral constants. Wefinallyobtain the following expression for maximum deflection

_2_ 128

T^C ^?K^!^I*^(^* J ^ t ^ C *^)Kf^lT(».*Poo)|;p,q:odd

p2 + >.V .= 64 £ % X2pq(p2 - 4)(q2 - 4) ^^-^j^J''^ ' "^^^ ' "^^ (14) ^=^

where

h, (15) We consider the case where the temperature distribution on the xy-plane of the plate takes a parabolic form and the temperature distribution through the thickness is linear and the temperature on the lower face is one-half on the upper face , that is, the temperature distribution is expressed as follows: T(x,y,z|)= T „ + T , (0 1L

•(vfi'-i^*" ^ a ' J L V b / JiM'^si ' 3h'

(16)

Then if a lateral load is uniform, q(x,y) - QQ , the above equation (14) becomes

128l-v2l »15

j,2| '^ j i2(i.v2)(^4\2'^'| 4 ( l - v ) P 5 ^ 2 H T , ^ 9 l U " ' • | ' '

-E_ Hfn2.4¥n2-4rh' '"'^'"^.;;idd^'(p^)^(p^-4X'^^-4)

_SQ_ - l (y)^i2-4)(j2-4) 7l4 . l l.i ij:odd

(17)

84 where Qo* - qob''/(Eoh'') stands for non-dimensional loading. In case of Karman, Eq. (14) becomes the following

^dJ 9 i + 2vx.^ + x'^ I 1 / n i + x"^, 4 , 1 I 1 (Ir^i ^128(l-v2) ^4 3 2 | 8 ^4 (^^^2)2 (^^^^^2)2 (4,;,2)2|J^

• { 4(T02 + PT'02) - (TO4 + pTo4)| + HTIO + PT'20) -(T40 + pT'4o)ll + ^ ^ / ^ ^0?^ " J 2(1+X2)

IX'

i |(T24+PT'24) ^ (T42 + PT'42)|

^-(ftlh)"^ 6

(Tpq-PTpq)

p,q:odd = 64 Ji2 U

i,j:odd

p2 + x V 5^ pq(p2-4Kq2-4)

'lij ij(i2-4)(j2-4)

(18) In case of /?= 0, this equation becomes the same equation as Eq. 26 of [10] and for aerodynamic heating with a uniform load, q(x,y) = qo, it becomes as follows 128(l-v2)

,4

12(l-v2)l;^4

^2

& JtMl+x2 = 21^ Z ^^I'l

i,j:odd

32|8

,4

(^^^2)2

(^^^^2)2

(^^^2)2|

U s 2 (l+J^)(Io^.l| + ^ | l + X U [4{l-v)(

5^2)17,

(t)25-TAr2ll^ L_._i 1 + 4)L2 4 + X^,

9l 4;t2| ^2J Z

P2.M2

(h)2^T,

* (iJ)^i2-4)(j2-4)

where qo* = qobVCEgh^) stands for non-dimensional loading.

(19)

4. NUMERICAL CALCULATIONS AND DISCUSSION OF RESULTS First we present the comparison with the solutions of the Berger and Karman equations for large deflections due to the temperature distribution through the thickness by means of the same Galerkin method in Fig. 1. The results by the Berger equation are thick lines and the solutions by the Karman equation are shown as thin lines. The solid lines are for the homogeneous elastic plate (/9 - 0) and the dashed lines are for the functionally graded plate (/9 - 1, n-2). It shows the relation between the maximum deflection(f ) and the temperature difference ((b/hfaT^) with TQ/TJ as the parameter and without a lateral load for a rectangular

85 clamped plate. The deflection due to heating according to the Berger method is lower than the Karman method. Nevertheless the two methods do show the good agreement over a wide range of temperature difference. When a temperature distribution as seen in aerodynamic heating which causes the deformation , f =0.7, at a plate with uniform elastic modulus (/? -0), is applied to the graded plate ( ^ = 1 , n=2), the amount of deformation can be reduced by 26% at To/Ti=0, 36% at 0.5, 40% at 1.0 and 49% at 2.0 in case of Berger approach. Figure 2 represents the maximum deformation due to uniform external loading QQ* using Berger and Karman. It may be seen that the deformation of the plate with the graded ekstic modulus in thickness direction is smaller by 20 % at (b/h)'aT^ = 0 and 15 % at (b/h)^aTi = 2.0 for QQ* =200 than the ones with a constant elastic modulus in case of Berger. It is observed from these figures that the graded elastic modulus can decrease the deformation of the plate.

5. CONCLUSION Assuming the Young's modulus and the temperature distribution through the thickness, the governing equations by Berger and Karman with the boundary condition of the edges clamped were solved by employing the Galerkin method. ^ A comparison with ^ e numerical results by the Karman -type non-linear governing equations showed sufficient agreement to verify the much simpler Berger approach. The graded elastic modulus can control the deformation of the plate. ACKNOWLEDGMENT The authors are deeply indebted to Dr. J. L. Nowinski, H. Fletcher Brown Distinguished Professor Emeritus, Department of Mechanical Engineering, University of Delaware for his helpful suggestions.

REFERENCES 1. H. M. Berger,. "A new approach to the analysis of large deflections of plates," J. Appl. Mech. 22 (1965) 465. 2. J. L. Nowinski, "Note on an Analysis of Large deflections of Rectagular Plates," Appl. Sci. Res., Sec. A, Vol. 11(1962)85. 3. N. Kamiya, "Large Thermal Bending and Thermal Buckling," ACTA TECHNICA CSAV, 1(1982)33. 4. J. L. Nowinski, and H. Ohnabe, "On certain inconsistencies in Berger equations for large deflections of elastic plates," Int. J. Mech. Sci. 14 (1972) 165-170. 5. Schmidt, R. 1974. "On Berger's method in the nonlinear theory of plates," J. Appl. Mech. 14:521. 6. T. Horibe, "Boundary Strip Method for Large Deflection Analysis of Elastic Rectangular Plates," Transaction of the Japan Society of Mechanical Engineers, 56(532) (1990) 140 (Japanese). 7. H. Ohnabe and F. Mizuguchi, "Large deflections of heated non- homogeneous circular plates with radicaUy varying rigidity," Int. J. Non-Linear Mechanics, 28, (4) (1993) 365. 8. H. Ohnabe and F. Mizuguchi, "Non-Linear Vibrations of Heated Non-Homogeneous Elastic Circular Plates with Radially Varying Rigidity," Proc. American Soc. Composites, 9th Technical Conference, (1994) 1147.

86 9. F. Mizuguchi and H. Ohnabe, "Large Deflections of Heated Functionally Graded Simply Supported Rectagular Plates with Varying Rigidity in Thickness Direction/* Proc. American Soc. Composites, 11th Technical Conference (1996) 957. 10. M. Sunakawa, ^Thermal Deformation of Clamped Rectangular Plate Subjected to Aerodynamic Heating," Transaaion of the Japan Society of Mechanical Engineers, Vol, 9, No. 85, (1961) 37 (Japanese).

A = 0, »/»0.3.qQ*-0 I

.

I

I

.

.

L

1 Fig. 1 Relation between maximum deflection( f ) and temperature difference ((b/h)^aTj) with TQ/TI ' '

\j,

Berger ^«0 - —-^/?-l,n-2 Karman /?.l.n-2

A=1,

y=0.3,

T/T=1 0

i—•—I—I—I—I—I

100

200

I

1

L

300

400

500

Fig.2 Maximum deformation due to uniform external loading QQ*

600

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

87

Model investigation of ceramic-metal FGMs under dynamic thermal loading: Residual stress effect, thermal-mechanical coupling effect and materials hardening model effect Qing-Jie Zhang, Peng-Cheng Zhai and Run-Zhang Yuan State Key Laboratory of Materials Synthesis and Processing Wuhan University of Technology, Wuhan 430070, P.R.China

ABSTRACT The analysis model for the response of ceramic-metal FGMs under dynamic thermal loading is investigated. Emphasis is put on the effects of the residual stress, thermalmechanical coupling and hardening model for the materials. It is shown that the three effects are significant when the materials' response is inelastic and should carefully be considered in constructing the analysis model.

l.THE STATEMENT OF THE PROBLEM The key issue in developing ceramic/metal FGMs is to identify an optimum compositional gradation or micro structure according to the response to service loads or environments. Since different responses may lead to different conceptions of the optimum design, it is important to establish a correct model for the response analysis. There have been some studies on the service stress and optimum design of FGMs under some service environments, but a detailed investigation of the analysis model is lacking. The present paper focuses on the model investigation of the service stress of the materials under dynamic thermal loading. The loading history considered is indicated in figure 1 which simulates a wide range of service environments. In figure 1, the thermal loading history consists of two phases: phase I corresponds to the cooling phase after sintering and this phase is steady one; phase II corresponds to the service phase and this phase is dynamic one, either thermal shock process or thermal fatigue process. Phase II is illustrated in more detail in figure 2. The model investigation involves three important effects: • the residual stress effect. The residual stress is produced in phase I and may exert a significant effect on the service stress in phase II. This should be considered in a correct way. • the thermal-mechanical coupling effect. The thermal-mechanical coupling results from phase II and is an important phenomenon for the dynamic thermal loading process. This effect depends on the thermal loading rate and the plastic deformation of the material.

88 • the material hardening model effect. For ceramic-metal FGM, there are two hardening models which can be used to describe the plastic deformation: one is the kinematic hardening model and the other is the isotropic hardening model. Under dynamic loading circumstances, the effect of the two hardening models becomes important since they may predict very different responses of the materials. 2.MODEL DESCRIPTION FOR THE THREE EFFECTS 2.1.The residual stress effect There are two methods that can be used to treat the residual stress effect: one is so-called separate analysis model and the other is an unified analysis model. The two methods can be explained in a simple way as follows: separate analysis model: the resultant stress = the residual stress in phase I only + the service stress in phase II only; unified analysis model: the resultant stress = the stress obtained through an unified analysis for phase I and phase II (namely the two phases are treated as one unified loading process). 2.2. The thermal-mechanical coupling effect The thermal-mechanical coupling model for suddenly-heated ceramic-metal FGMs has been developed by the present authors in Ref [1]. To consider the plastic deformation effect on the heat conduction in the materials, the coupled heated conduction equation in Ref [1] is now modified as: ^

^

JT

JT

d

- (A(z) —) - Q (z) - - -f 3/:, (z)a(z)r[3a(z) — + -

G (TTTTT)]

d

" ^:^

G

G

(TT

- -)

The signals in the above equation have been defined in Ref [1]. If the coupling terms(the second and third terms on the right hand side) in the above equation are abandoned, the coupled heat conduction equation is reduced to the common uncoupled heat conduction equation. The motion equation of the model remain the same as those in Ref [1]. 2.3.The materials hardening model effect The stress-strain relationship for FGMs is assumed to be a bilinear form and can be described by the isotropic hardening model and kinematic hardening model as: \G\- G + Hs for the isotropic hardening model J{G- HS^ ) - cr^, = 0 for the kinematic hardening model where H = EE^/(E - E^)\ errand s^ are the yield stress and effective plastic strain of the materials, respectively; E and E^ are the elastic modulus and the elastic-plastic modulus.

89 3.RESULTS AND DISCUSSION A TiC/Ni FGM is used as an example. The graded layers in the FGM are treated as homogeneous with effective properties. The properties, geometrical sizes and the compositional distribution function in the graded layers were given in Ref.[l]. The dynamic thermal load in the service phase is taken as q = 4 MW/m^ and t^ = Is for the residual stress effect analysis and the hardening model effect analysis, and as q = 6 MWlm^ and t^ = 0.5s for the thermal-mechanical coupling effect analysis, where q is the magnitude of the heat flux and IQ is the duration of one thermal cycle. The sintering temperature in phase I is taken as 1300A^. The compositional distribution exponent is taken SLS P = 1.6. The residual stress effect is considered by the separate analysis model and unified analysis model, and the results are shown in figures 3(a) and (b). Figure 3(a) corresponds to the elastic analysis and it is seen that the separate analysis model and unified analysis model give the same resultant stress; Figure 3(b) corresponds to the elastic-plastic analysis and it is seen that the resultant stresses found from the two models are different. This is because in the elastic analysis the two responses in phase I and phase II are both linear and the resultant response can be obtained from a direct superposition of the two separate ones. In the elastic-plastic analysis, however, the two separate responses are both nonlinear and the separate analysis model based on the direct superposition of the two separate responses is no longer suitable. The thermal-mechanical coupling effect is shown in figure 4. The difference between the coupling model and uncoupling model for the presented thermal loading is about 7%. Different thermal loadings have also been examined and the results indicate that the difference between the two models increases when the loading rate and plastic deformation increase. The material hardening model effect is demonstrated in figure 5. From figure 5, the kinematic hardening model and isotropic hardening model produce the same material response in first thermal cycle and different ones in the subsequent cycles. The first difference between the two responses after the first cycle is the compressive history at the ceramic surface: the compressive history predicted by the isotropic hardening model is longer than that predicted by the kinematic one. The second difference between the two responses is the maximum tensile stress at the ceramic surface: it is nearly the same in every thermal cycles for the isotropic hardening model, and it changes and increases in the subsequent cycles for the kinematic hardening model. ACKNOWLEDGEMENTS This work was supported by the National Science Foundation of China REFERENCES 1. Q.J.Zhang, G.H.Zhang and R.Z.Zhang, Proc. 3rd Int. Symp. on FGMs, 1994, 235-240

90 T(K)

time phase I

phase n

Figure 1: Thermal loading history of ceramic-metal FGMs phase I: cooling phase after sintering phase II: service phase

heat flux ,

1

X

ceramic graded layers

1

metal liquid cooling

tmie

Figure 2: Dynamic thermal loading simulating the service phase in figure 1

91 o

time (s)

^

o o (N

O

S o

-K3- unified analysis model —i^ separate analysis model

o o

a: elastic case o o

o o cd

T3

o a o (N B

I

o B o

• unified analysis model • separate analysis model

o o

b: elastic-plastic case Figure 3: processing-service stress obtained from the unified analysis model and separate analysis model

92

_ 1

a

uncoupled theory coupled theory

O rsi in

PLH

s rn C/1 d)

o 0^

^ ^

h

0

0.5

1

1.5

time t(s) Figure 4 : Effect of the thermal-mechanical coupling model

o o o

- isotropic hardening model • kinematic hardening model

o o o

Figure 5 : Effect of the materials hardening model

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

93

Fractal geometry and it's implications to surface technology D.P.Bhatf, O.P. Bahl^ R. Schumacher^ and H. Meyer^ ^Carbon Technology Unit, National Physical Laboratory, New Delhi 110 012, India ^Atotech Deutschland GmbH, Erasmusstrasse 20-24, 10553 Berlin, Germany The basic know-how is presented to simulate impedance diagrams of complex equivalent circuits by viewing the electrode surface through fractal patterns. The implications of this model for electrochemical surface technology are also reported in this paper. 1. INTRODUCTION In the recent years fractal geometry has made considerable advances in the surface problems of many scientific disciplines^'^. Fractals (a word coined by Mandelbrot in the seventies) which mean basically for either irregularity orfragmentation,are more than the topographical dimensions, or these are the geometrical objects of non-integer dimensions which always exhibit the rule of self-appearance characteristics in contrary to the conventional elements of Euclidean-geometry. One of the first steps of the development of systematic fractal geometry, including its computer graphic aspects have blossomed tremendously on the basis of a simple system of Z = x^+C in order to subsequently demonstrate how one could calculate Mandelbrot set^'"^ in a particular plane of Cp,q. Assemblance of ZiZ2,Z3 Zn, under varying conditions of the square of the hypotenuse of the p,q plane while put together in the plane by applying the simulation principles, could generate two categories of points viz. White and black, leading to the formation of the strange figure, nick-named as apple tree owing to its appearance (Fig.l). Any fiirther modelling of Cp,q points in different shorter ranges from the Fig. 1 could give rise to ensemble offinerdiagrams^

A

B

n^T;r-A>Q.

. ^ ^

3^

^-^

Brtjcr^tZt^

(a) Figure!.

(b) "^'^Ure

2.

94 Traditional electrochemical concepts of the interface usually proceed by assuming an ideal smooth electrode surface ( e.g. Mercury). In other words, the ideal homogeneous electrode surface plane is often the worst assumption on the part of the electrochemist to prepare their interfacial experiments. Decisive factors associated with irregular heterogeneous surfaces are governed in various processes for e.g. the growth of crystalline and amorphous materials, heterogeneous catalysis, galvanic processes, corrosion, general decomposition processes, surface rich adsorbance, surface enhanced Raman spectroscopy, etc. Distinct features of the fractal surfaces in comparision to those of the non-fractal ones have been exemplified in Fig. 2A-E^:SituationA: Here is a case where atoms or molecules (o-o) go alongwith the surface only. In case (a) having the smooth surface, the distance between the molecules or species is shorter within the same length of the electrode and hence could be considered as the better situation as compared to the situation with fractal surface (b) Situation B: In this situation atoms or molecules go to the surfaces by diffusion controlled mechanism. Fractal electrode (b) is better than the smooth electrode (a) because the ionic species have got relatively more paths in the former situation and could thereby possess more ionic mobility to get themselves transported to the electrode surface. Situation C: Edges, cracks, etc. (b) could place at disposal energy richer reaction centres than the surfaces with the normal number of free valencies (a). Therefore, the active centres can be distributed more in case of fractal surface (b) than the non-fractal surface (a). This is a prominent case applicable to electrocatalysed systems likefiielcells. Situation D: When the surface has many neighbouring atoms or species, one may then expect two neighbouring atoms or species to interact. In (b) situation, one has choice for the atoms to make more neighbours and hence the system possesses more neighbouring atoms. Because of the multiple choice, the interaction could,therefore, result more in the fractal electrode than in the case of smooth surface (a). Situation E: Both (a) and (b) planes here belong to the real surfaces (fractal). However, the dimension of the molecular species would be the detrimental parameter of the reactions in the real surfaces, i.e. having the same surface, either more or less number of molecules can be absorbed. The size of the molecules is apparantly smaller in the former one and hence more number of molecules can be accommodated in the structure (a) than (b). Fractal dimension, D is considered as an effective number that characterises the irregular electrode surface. The term has been related to physical quantities such as mass distribution, density of vibrational stages, conductivity and elasticity. If we consider a 2-D fractal picture in its self-similar multi-steps, one can draw various spheres of known radii at various points of its structure and may thus count the number of particles, N inside the sphere by microscope, following relation will then hold good : N(r) - r^ (1) where D = Fractal dimension and N = Number of lattice points inside a sphere of radius, r The plot of r vs. N could give rise to a straight line, the slope of which equals the value of fractal dimension. Equation (1) is, however, not valid for the 3-D geometry of the structure. A general simulation programme for the calculation of Nyquist plots with

95 respect to the 3-dimensional fractal electrode-electrolyte interface is included in the next section. ^4

/

0 STAGE 6=30.13 D = 2.50 1st STAGE — I \—I

^

>—^ ^^—11nd STAGE

Sc

.^JLAJ Figure 3.

^ A j X A - m r d STAGE Figure 4.

2. STRUCTURAL MODELLING OF THE NYQUIST CURVES FOR FRACTAL ELECTRODE/ELECTROLYTE INTERFACE The simplest model to describe the impedance behaviour of the metal/solution interface can be the Koch curve. Fig.3 illustrates the view of the irregular surface in terms of fractal geometry. Each interface consists of long parallel V-shaped grooves filled with electrolyte (black shaded region). Working electrode is depicted by the white portion. The metal/electrolyte interfacial boundary forms a generalised Koch-curve, the fractal-dimension of which can be measured by various methods described elsewhere^'^. Roughness of the surface is the measure of the fiinction of 0 (angle of the grooves). In this example, we consider a four-stage Koch-curve wherein 4 peaks of different sizes are present (Fig.4). Dilatational symmetry is one of the feature of the Koch-curve. Each peak could thus possess different impedance characteristic, which are named as Z4, Z3, Z2 and Zi according to the falUng sizes. Besides these peaks, there are 16 flat regions which show pure capacitive behaviour (Zo). As the impedances are all connected in parallel, the whole impedance, Zt is composed of the reciprocal values of the individual impedances. If the different regions of these peaks are taken into account, Zt may be written as : 1/Zt =16/Zo+8/Zi +4/Z2 + 2/Z3 +I/Z4 (2) 2.1. Calculation of the impedance of a flat region (Zo) If we permit charge penetration to the interface, the simplest equivalent circuit diagram is then a parallel circuit composed of R^ and Cai in series with Rbuik- Its impedance is calculated as : 1/Z = (jcoCdi + 1/Rct) (3) or Z = Rct/GcoCdiRct+1) = RcAo'Cdi 'Rct'+ 1) - j [(o)CdiRct')/(«'Cdi'Rct'+ 1) (4) Taking into consideration Rbuik gives :Z = Ret/(co'Cdi 'Rct'+ 1) - j [(cDCdiRct')/(«'Cdi'Rct'+ 1) + Rbuik

(5)

96 2.2.

Calculation of the impedance of smallest peak (Zi) Conical shape as divided into K number of slices is the drawing of the smallest peak [Fig. 5(a)]. Equivalent circuit of these different slices is an ensemble of branched resistance/capacitance ladders [Fig.5(b)]. The parameters like ro,ri,r2, r/; ro',ri',r2', n' and Co,Ci,C2, Ci are the K different values of bulk resistances, charge-transfer resistances and the capacitances, respectively. Wang^ did not consider these additional charge-transfer resistances. Considering the circled last parallel connection (Fig. 6), the impedance of the last but one slice could be calculated. This rule can be further followed to other equivalent circuit diagrams in a similar manner till one gets Zo which would be the total impedance of the peak. Following the analogy of the apple tree, in this situation too, one could do with a recursive calculation in order to formulate the following derivation to calculate the impedance of all the sHces : 1/Zi.i = l/{(ai +ri) + bj)} + l/r/.i + JCOCM (6) where i is the index for K and varies from K to 1. Substitution of uj = ai + n in equation (6) followed with suitable mathematical changes gives :1/Zi.i = Ui/(ui' + bi') - j [(bi/(ui' + bi')] + l/r'i.]jcoCi.i (7) Separating into real and imaginary parts gives :1/Zi.i = Ui/(ui' + bi') + l/r-.i + j[coCi.i - {bi/(ui' + bi')}] Substitution of Ci = Ui/(ui ^+ bi^) + l/r'i-i as real part and di = coCi. imaginary part gives :l/Zi.i = Ci+diJ or Zi.i = l/(ci + dij) = Ci/(Ci' + di') - j[di/(Ci' + di')]

Thus, we obtain > New real part: New imaginary part

ai-i = Ci/(ci^ + diO bi.i-di/(Ci' + di')

bi/(Ui' + bi')

(8) as

(9) (10)

(11) (12)

4=h^I3TCI]TCIh (b)

-^^Y^'fg"f(]

^^-o-

Figure 5. Figure 6. Now ri, r'i and Ci are to be calculated. The length and the depth of the electrode has been originally assumed as 1 cm. Capacitance of the interface may be calculated by multiplying the standard capacitance and the length of the Koch-curve. Both sides of the groove structure are considered here in constrast to the flat smooth surface. Using the geometrical principles, further steps with respect to the calculations of length of the fractal electrode, number of peaks, bulk resistance, etc. have been followed in the numerical

97 model. With this knowledge, it could be then possible to calculate the impedance of the cone (Zi) in any multiple number of the structure. 2.3.

Modelling of a rough electrode with equivalent electrical circuits Z2 is the impedance of the second smallest network as shown in Fig. 7(a). Parallel placements of three branches each having the Zi impedance could give rise to the resultant impedance of Zi/3 at the intersection. The latter, in turn, becomes the terminating impedance at the end of the network for the calculation of Z2 [Fig. 7(b)]. The branched resistance/capacitance network can be solved by using the same equation as employed previously for the calculation of Zi. The impedances of the next higher grooves such as Z3, Z4, Z5,.... Zn, can also be calculated in similar manner depending upon the branching configuration. In this work, a general programme has been thus developed using Turbopascal 5.0 language in order to simulate the impedance of the many peak irregular surface with n number of stages. The details of this software are to be published.

(a)

(b) 0

^00 200 REAL PART(ohm)

Figure 7. 3. PRESENTATION OF THE MODEL CALCULATIONS

Figure 8.

The assessment of the reported model for experimental passivation systems of say magnesium/magnesium perchlorate^^'^^ and titanium/or titanium dioxide/sulphuric acid has been made from the view-point of fractality of the electrode surface and a good agreement between the model calculated complex plane impedance (Fig. 8 being a typical simulated plot) and the measured Nyquist-plots is obtained. The charge transfer resistance (Ret) values in respect of magnesium AZ31 alloy in 2 M Mg(C104)2 have been measured as 4154 , 338, 238 and 150 ohm for the applied potentials of 0,10,30 and 50 mV, respectively^^. Increase of the passivation potential from 0 to 50 mV has brought about the gradual decrease of the radii of the semi-circle - an anomalous observation because in any normal circumstance, the Ret of the electrode surface should have increased with the rise in the passivation (a case of oxide film formation). In order to rationalise this unusual polarisation behaviour, one could imagine the interpretation based on the concept of anion adsorption in the surface oxide layer wherein afieldassisted transport of ions through the passive layer is introduced to account for the ion current density increase upon raising the oxidation potential. These interpretations are, however, not valid for systems containing no chloride in the base electrolyte. High corrosion rate accounts the reason of the non-employment of MgCb

98 electrolyte in the cells developed by us in India ^'*•^^ Different parameters such as the groove angle, number of slices, number of different peaks, capacitance, series/and parallel resistors and the thickness of charge transfer resistor have been varied during the simulation by the programme so as to obtain the series of semicircled impedance diagrams. The important findings are such that decrease of the groove angle brings about the decrease in the Ret which has been found valid for 10 < 9 < 75 and 5 < N (number of different peaks) < 20. It is thus envisaged qualitatively that during passivation, morphological changes in the electrode take place and this change becomes effective for the charge flux to diffuse which thereby explain the cause of anomalous behaviour. This model, however, requires much more refinement as the small lateral branches remain still unconsidered and the number of slices (K) must be selected so large that the impedance is constant i.e. Z ^ f (K). Also the simulation of the diffusion controlled interfacial processes has not been possible through the present model. These are some of the open areas wherein modelling research could be further pursued . Acknowledgements. Correspondence and discussions with M. Wicker is highly acknowledged. DPB is obliged to DST , Govt, of India and Klaus Hagen, Atotech, Tokyo 141 for providing partial financial aid in attending FGM96 symposium. Grateful thanks are accorded to the Director, NPL, New Delhi for giving permission to publish this paper. REFERENCES 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15.

L. Pietronero and E. Tosatti (eds.). Fractals in Physics, Elsevier Science Publishers B.V.,Netherland, 1986. M. Gardner, Scientific American, 238 (1978) 16. B.B. Mandelbrot, The fractal geometry of nature. Freeman W.H. and Co., New York, 1982. H.-O.Peitgen and P.H. Richter (eds.). The beauty of fractals - Images of complex dynamical systems. Springer Verlag, New York Inc., 1986. H.-O. Peitgen and D. Saupe (eds.). The Science of fractal images. Springer Verlag, New York Inc., 1988. P. Pfeifer, Chimia, 39 (1985) 120. D. Avnir and P. Pfeifer, J. Chem. Phys., 79 (1983) 3566. R.F. Voss, Physica Scripta, T13 (1986) 27. J.C. Wang, Electrochim. Acta, 33 (1988) 707. R. Udhayan and D.P. Bhatt, In International Conf on Magnesium alloys and their applications, Germisch-Partenkirchen, Germany (1992) 59. R. Udhayan, Ph.D. thesis, M.K. University, Madurai, India, 1991. M. Wicker, Ph.D. Dissertation, University of Kiel, Kiel, Germany, 1991. D.M. Drazic, S.K. Zecevic, R.T. Atanasoki and A.F. Despic, Electrochim. Acta, 28(1983)751. D.P. Bhatt and R. Udhayan, Indian Pat. 749/DEL/l 991. R. Udhayan and DP. Bhatt, J. Power Sources, 39 (1992) 323 ; J. Appl. Electrochem., 23 (1993) 393.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

99

Database System for Project of the Functionally Graded Materials K. Kisara, A. Moro, Y. S. Kang, M. Niino National Aerospace Laboratory, Kakuda Research Center, Japan. ABSTRACT This report will introduce the concept and status of a database system for a national project entitled: "A Study in the Development of Energy Conversion Materials through the Formation of Gradient Structures" , which project is being promoted by some thirtyodd industrial, academic and governmental organizations. While the database system will be, organizationally speaking, somewhat loose-knit, it is ejected that the project's initial goal will be realized within a fixed period of time. In order to provide a convenient means of comunication for researchers working in various organizations to exchange information, complete their work assignment in a short space of time and maintain a common perspective, the establishment of some method of information-sharing is essential. At the same time the report will describe measures of the need for maintenance of the database and a network to make use of it in real time so as to promote support for the project and, additionally, show that the use of the Internet is a viable method of operating the system at this present time.

1. INTRODUCTION The Fimctionally Graded Materials Database System was established to manage the entire body of data arising out of the Science and Technology Agency funded project: "A Study in the Development of Energy Conversion Materials throu^ the Formation of Gradient Structures" (Phase I: 1993-1995; Phase II: 1996-1997, hereafter referred to collectively as the FGM Part 2 Project), and to maintain such so as to facilitate effective access to said data. The project aims at dividing materials into thermal areas and developing the most efficient energy conversion elements for each individual area, all the while keeping to the fore the concept of a compound system that would convert thermal into electrical energy at a high level of efficiency. Some thirty-odd industrial, academic and governmental orgEinizations are participating in the project, the activities of which are carried out by subcommittees. As the fimction of these subcommittees is to make use of e)q)erimental data found in the published reports and minutes of seminars held in various fields of study, an information centre to collect

100 and maintain this data assumes an important role in enabling the project's participant organizations to share and effectively access such information resources. Additionally, the centre play s a vital role in preventing the loss of project assets. Also,fromthe very early stages of the project, maintaining and actually operating such a centre as a base for accessing information is vital to the project's receiving adequate support.

2. FGM PROJECT AND DATABASE The following is an outline of the role of the database within the project as a whole. The research being conducted under the project has as its principle aim the development of fundamental technological skills essential to the creation of new materials, these skills being developed by sharing results of research among the various specialist subcommittees dealing with planning, synthesis (creation of new materials) and evaluation into which the thirty-odd participating organizations are divided. By means of this co-operation, participating bodies become part of a loose-knit organization working together to achieve common goals and thus it can be seen that the adoption into the system of a groupware concept whereby these bodies engaged in research can mutually access each other's published data would be effective. The relative position of the database as seen against the background information outlined here is illustrated in Figure 1. As described above, within theframeworkof a looseknit organization where participant bodies work together to achieve common goals, the database needs to provide an operating system for the groupware, desiga the system platform that serves as its base and, at the same time, fulfill the primary function of the database, namely to prevent the loss of the fruits of project research so that individual researchers may mutually access this information. The different duties involved in this information distribution are classified variously as: data gathering and maintenance, standardization, access environment maintenance and educational.

^ ,

plaiming(designing)

.•^

estimate of characteristic estimate of performam design of FGM ^ 3 ^ 3 ^ ^ ^ Ithennoelectru material

Thermal andl Mechanical evalution total evaluation laracteristic of FGM iterial

S Ithermionic database material |(FGM emittei a te collector] radiator o insulator electrode heat collecter

, ^

synthesisofinaterials

'^.

Fig.l The relative position of the database

In order to accomplish all this, the FGM Database manages the entire body of project data and maintains a network for effectively accessing this data. The relationship between the database and the project is shown in Figure 2. Basically speaking, the database is made up of three groups of data shown in Figure 2.: data from written sources, that from electronic conferences and measurement datafromexperiments.

101 characteristic of FGM material Icharacter of thermal and | mechanical evalution FGM database (three groups) f Database server ^ database *«ct data management system store data

electronic filing report data (Optical memory disc) logging data of BBS (first class) communication

JI

provide data multi platform -interface—! subcommittees mtemet planning or modem [Network systeiji WWW e-mail ftp —news—

thermoelectric material thermionic generator synthesis of materials evaluation

-support system for planning-

knowledge base of FGM (example based)

theory, experence simulation technique of use of database] technique of designing

expert system

Fig.2 The relationship between the database and the project

3. OPERATION AND THE NETWORK 3.1. Design of the System Platform In the early stages of the project, a database system centering on personal computer communication was constructed to enable infonnation to be shared and accessed. At the present moment a different approach is possible with group ware making use of the Internet, and machine interface is possible by means of hypertext browsers such as the World Wide Web. As the database is constructed to gather and maintain information, questions such as the following are pertinent: What information is gathered in the database? How is it compiled? Who uses it? How do they use it? What is needed is a database that will make the optimum use of limited resources, in other words, one that will increase the efficiency of the brainwork of its users.

102 In the area of what information is gathered, the following data has been collated: information and actual experimental data on the physical properties, methods of making and measuring the performance of new materials being developed; graphs, tables and photographs published in seminar journals and subcommittee reports together with data in written form from monographs and the published minutes of subcommittee conferences. As to how it is compiled, keywords are assigned to assembled data from e^qjeriment measurement values, images, written works such as monographs and this data is maintained in electronic files as a means of managing project information. Also, the results of information exchanges using electronic bulletin boards are stored and accumulated as further resources of information. As to the envisaged users of the system, for the duration of the project, the idea is to provide timely information to researchers belonging to organizations involved in the project and, upon the project's completion, make this information available to public users in an open database. On the matter of method of access, direct access via the Internet has been chosen. At the inception of the project, a network was deployed using modems connected to the public telephone lines but, in response to changes in the network environment cause by the growing availability of the Internet, a switch was made to direct access via the net. In summary, the database was assigned the following three information management roles: To serve as a database in the narrow sense of the term - simply gathering, organizing and recording data. To provide a system for dissemination timely information - enabling participants to share groupware data. To maintain an environment for access and a users' service (support for researches on how to make use of the system). These three roles are dealt with morefrillybelow. 3.2. Maintaining an Environment for Information Access In order to keep the project moving forward it is vital that participant organizations not only share the same goal but also share information and maintain a common perspective. The project maintains a closed networksystem, excluding access to all but participant org3nizations, as a base for information sharing whilst protecting portions thereof that may need to be kept secret. Operation of the original service begun at the Japan National Aerospace Laboratory's Quaked Research Centre with the construction of a personal computer communication host computer (BBB) using four public telephone lines. Subsequently, with the availability of the Internet, a BBB using Window and Macintosh to support TCP/IP protocol was put into operation. At the present moment, a switchover is taking place in the mode of operating the database access environment from the old C/S (client server) system to a method employing the concept of a distributed database such as the World Wide Web using the Internet. A distributed database is being constructed hat uses data and monographs in the

103 possessionof the various individual participating organizations to create home pages, and then serves as a database centre, consolidating these pages into a single home page for convenient access. 3.3. The Functionally Graded Material Database in the Narrow Sense of the Term Gathering, organizing and recording data contributes to the accumulation of the project's immaterial assets. Data is principally assembled from the published monogr^hs of researchers and from subcommittee reports. These are recorded on electronic files. Data is organized at the time of recording according to information or keywords necessary to reference it. Additionally, in order to manage this information, a search application is provided to facilitate researchers' access to the data from their own personal computers. Also, as measurement data from e^eriments is used in designing new materials or inferring their physical properties, it is vital that such measurement data be kept up to date in a systematic fashion so that the database may fulfill its role in the narrow or restricted meaning of the term "database*' and the maintenance of such a body of experimental data increases the value of the database's support operations. 4JTJTURE PROSPECTS In these days when the dual concepts of the Intemet and group ware are being brou^t more and more to the fore, a base whereby project researchers can make optimum use of needed information is being put in place. Formerly, a WWW browser would read and display a specified file but, from the spring of 1995, the Hot JAVA browser developed by the Sun Microsystem Company has deployed a form of programming called "Applet" . By means of CGI (Common Gateway Interface) procedures org3nically linking h i ^ processing level database servers such as SQL together with the WWW server, the construction of a distributed database able to conduct sophisticated dialogue has become possible. Always adjusting to progress in the network environment, the Functionally Graded Materials Database will continue to gather data and support the project. It will also continue to frmction as an attractive centre for information pertaining to frmctionally graded materials by translating the tables and the abstracts of principal monographs into En^sh, thereby fiilfilling its role as an internationally valuable information source and contributing to the standardization of research data and, additionally, it is intended that it will serve as a news centre educating the public regarding the concept of fimctionally graded materials.

104

NAL KRC LAN Macintosh TCP/IP H

Macintosh

GatorBox Filing system

UNIX

Windows

LocalTalk

llU EtherNET

AppleTalk TCP/IP DataBase BBS;First Class

^ ^ ^ '

IBM/AT DOS/V Modem (4 ports)

Macintosh

TEL NTT

IBM/AT Windows

PC98/DOS

WWW databese homepage http://fgin.kakuda-splab.go.jp/

Fig.3 Network System for FGM Database

References 1) R. Watanabe ,Fimctionally gradient material(Japanese),The society of non-traditional technology

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

105

Fracture Mechanics of Graded Materials F. Erdogan Department of Mechanical Engineering and Mechanics Lehigh University, Bethlehem PA, 18015, USA In this article after a brief review of elementary principles of fracture mechanics, certain issues concerning the applications to graded materials are identified and some examples are given. 1. INTRODUCTION In recent past there has been a great deal of interest in the concept of material property grading as a tool for new material design. This is usually accomplished by suitably varying composition and/or microstructure of the medium. Thus far most of the work in the field has been on metal/ceramic composites, various intermetallics and electronic materials with current and potential applications as interfacial zones and coatings. From a mechanics view point the main advantages of material property grading appear to be improved bonding strength, toughness and wear and corrosion resistance, and reduced residual and thermal stresses. Some typical applications include thermal barrier coatings of high temperature components in gas turbines, surface hardening for tribological protection and graded interlayers used in multilayered microelectronic and optoelectronic components [1-3]. An important aspect that needs to be addressed in various engineering applications of FGMs is the question of reliability and durability in general and fracture related failures in particular. This article is concerned primarily with fracture mechanics as applied to structures involving FGMs. After a brief description of some basic notions of fracture mechanics, certain critical issues relating to FGMs are discussed and some examples are given. 2. FRACTURE MECHANICS The quantitative theories of fracture which are currently in use are based on a fundamental principle of continuum thermodynamics, namely the first law or the energy balance which states that dU_ _ dV^ dJ^ dD_ dt dt dt dt where t is the time U the external work, V the recoverable (elastic) energy, T the kinetic energy and D the sum of all irreversible energies associated with the creation of new fracture surfaces such as surface tension, plastic work and viscous dissipation. If the solid contains a

106 dominant flaw which may be represented by a planar crack having a surface area A(t) and if the fracture process is taking place in a quasi-static manner and A(t) can be characterized by a single length parameter a{t), then dT/dt = 0 and defining dD/da = Qc, (1) may be expressed as ±(U-V)=

Gc,

(2)

where the left hand side is the rate of energy available for fracture (also known as Q, the crack driving force or the strain energy release rate) and Qc represents the energy required for unit crack extension. If the fracture can be characterized as a low energy or brittle phenomenon, then it can be assumed that the size of the inelastic region around the crack tip (also known as the fracture process zone) where all the dissipative processes take place is small compared to the crack size a, Qc is independent of a and the energy flowing into the crack tip region comes from the elastic bulk of the medium and is insensitive to the details of the stress and deformation states in the fracture process zone. The significance of this last observation lies in the fact that a purely elastic solution may be used to calculate the crack driving force. For example, by observing that under normal opening or mode I loading the asymptotic values of the cleavage stress and crack opening displacement at the crack tip x = a, y = 0 are given by [4] ^^^^^'^^ ~ ~J2(i^'

^y ~ ^y = ^^y(^'^) = " ^

hy/2{a-x),

(3)

by using the concept of crack closure energy, the energy available for fracture can be expressed as 1 pa+da

d{U-V)

= -

J

ayy{x,0)I^Uy{x - da,0)dx,

Z Ja

ki = ^mJ2{x

—{U -V) aa

- a)ayy{x,0),

1 4- /c

= Qi = - ^ n k l Qfi

(4)

where ki is known as the mode I stress intensity factor, /x is the shear modulus and « = 3 - 4z/ for plane strain and « = (3 - i/)/(l +1/) for plane stress, i/ being the Poisson's ratio. Similarly under mode II, the in-plane shear and mode III, the anti-plane shear loading conditions, the stress intensity factors, the corresponding energy release rates and the total strain energy release rate for co-planar crack growth are given by /C2 = MiJ2{x

-a)(7a;y(x,0),

^// = ^ 7 r A : | ,

Giii = ^7rkl

ks = ]im^y2{x - a) ayz{x,0), Q = Qi+Qii + Qiii.

(5)

Referring to the general expression (2), in brittle fracture the critical value of Qc corresponding to the co-planar crack growth is known as Qic, the fracture toughness, with G = Qic being the fracture criterion. In practice, very often Kj = ^J^k\, Kjj = ^Jl^hi, Km = v^/cs and Kj = Kic are used as the stress intensity factors and the fracture criterion. In addition to their successfiil applications to fracture stability problems, the stress intensity factors have been widely used as correlation parameters in analyzing the subcritical crack growth rates da/dn (in fatigue) and da/dt (in corrosion), n and t referring to the number of

107 load cycles and time, respectively. In the presence of large scale inelastic deformations the general energy balance criterion described by (2) is still valid. However, in this case since Qc is no longer independent of the crack size, the fracture propagation can not be described by a single parameter criterion. In high energy fracture, the fracture process zone, and consequently Gc usually grows with the growing crack size. Hence, for fracture stability problems it becomes necessary to use a criterion based on variable Qc such as a crack extension resistance curve or a R-curve approach. Unlike the linear elastic fracture mechanics dealing with brittle fracture and subcritical crack growth, the tools of the so-called elastic-plastic or nonlinear fracture mechanics dealing with high energy or ductile fracture are not well-developed and universally accepted. In some cases the J-integral is used with some success to compute the crack driving force, Q = d{U — V)lda. In all cases the application of the criterion G = Gc requires extensive numerical and experimental work. 3. MAJOR ISSUES IN FGMS The principles of fracture mechanics described in the previous section are applicable to inhomogeneous as well as homogeneous materials. In FGMs the difficulties arise in the solution of elastic or elastic-plastic crack problems to evaluate G or ki, ki, k^ and in characterizing the material to determine Kjc, Gic or Gc where the fracture toughness Gic is no longer a material constant [5]. The definitions of stress intensity factors and expressions of the strain energy release rates given by (4) and (5) are still valid provided the elastic parameters /i and K are evaluated at the crack tip. Following are some of the major issues concerning the fracture mechanics of FGMs. (a) Elastic singularities. As long as the elastic parameters ^ and K are continuous ftinctions of the space variables with piecewise continuous derivatives, the stress state around the crack tips has the standard square-root singularity. For example, in plane isotropic elasticity problems for r -• 0 the leading terms of the stresses are given by [6-8] -

-'^^'"'^^[fci/iyW + fc2/2i,W], {ij) = ir,e)

(6)

where (r, 9) are the polar coordinates at the crack tip, ki and k2 are the modes I and II stress intensity factors, (f>{r,6) is a smooth function with
k^f

^-A^

aeeir, 0) c^ - ^ ( 1 -h ^^,, ^\.^ /3r\nr] + c^/^, y^V 87r(l-i.)

(7)

where c is a constant, /i(x) = fxocxip{/3x) for x < 0, /i = fiQ for x > 0 and the crack is located along 0 < x < a, y = 0. (b) Analytical methods/benchmark solutions. Even though there are no known closed form solutions for crack problems in FGMs, for simple property variations the formulations leading

108 to singular integral equations are straightforward and accurate solutions can be obtained [7]. (c) Computational methods. Finite element method is the major computational tool. However, to improve efficiency and accuracy the development of enriched crack tip and transition elements and ordinary inhomogeneous elements will be needed [10]. (d) Material orthotropy. In many cases the material orthotropy seems to be the consequence of processing technique. For example, FGMs processed by using a plasma spray technique tend to have a lamellar structure. Flattened splats and relatively weak splat boundaries result in an oriented material with higher stiffness and weaker cleavage planes parallel to the boundary [11]. On the other hand graded materials processed by an electron beam physical vapor deposition technique have usually a columnar structure giving higher stiffness and weak fracture planes in thickness direction [12]. These oriented materials can generally be approximated by an inhomogeneous orthotropic medium. (e) Inelastic behavior. Because of the length scales involved in FGM coatings and interfaces, in addition to conventional plasticity, one may have to use a microplasticity approach which accounts for the effect of strain gradients on strain hardening coefficients. The resulting nonlinear elastic-plastic crack problems require a numerical approach with special inhomogeneous elements. (f) Rheological effects. Invariably FGMs are used in high temperature environments. As a result the time-temperature effects may not be negligible and the material may have to be modeled as an inhomogeneous viscoelastic or viscoplastic medium. (g) Dynamic effects. Generally, high velocities in propagating cracks and high rates of loading (e.g., impact) in stationary cracks would necessitate the consideration of inertia effects in solving the fracture problem. However, even for the uncracked linear elastic inhomogeneous bounded medium, the stress wave phenomenon is not fully understood. The existing solutions are restricted mostly to one dimensional problems in materials with certain simple property gradings. (h) Material characterization. This is still the most important issue in studying the fracture mechanics FGMs. The knowledge of thermomechanical and fracture mechanics parameters of the material is essential for any realistic predictive reliability study of FGM components. 4. EXAMPLES In this section we will briefly discuss three groups of examples. The first two are concerned with FGM coatings on homogeneous substrates in which for simplicity it is assumed that the bond coat has the same thermomechanical properties as the substrate. The third group deals with the effect of material orthotropy on the stress intensity factors. 4.1. Surface Cracking In FGM as well as homogenous (ceramic) coatings the fracture related failures may take place in various ways. One way would be under cyclic mechanical and/or thermal loading the initiation of a fatigue crack at a surface defect, the subcritical growth of the crack in thickness direction, fracture of bond coat and opening an oxygen path to the substrate. This may happen if there are no weaker fracture planes in the coating and the coating/bond coat interface is

109 sufficiently strong. Such crack initiation and growth in thickness direction have been observed in FGM coatings by several investigators (e.g., [13,14]). A variation of this mode of failure would be multiple (or eventually, periodic) surface cracking. Multiple cracking is clearly in evidence in the work reported in [14]. In practice, because of the long hold times under high temperature, the crack growth process would be heavily enhanced by the environmental effects. Even in the simplest case of low temperature and relatively high cycle fatigue for which a simple two-parameter crack propagation model such as ^=C{AK)\ C = C{a), b = b{a) (8) an may be applicable, in FGMs the parameters C and b would be dependent on the material composition and the microstructure. This would mean that in surface crack problems C and b would be functions of the crack length a. For modeling and any quantitative analysis, these functions must be determined from the fatigue data on homogenous cupons with various composition. The surface crack problem is one of mode I and the determination of the stress intensity factor ki is sufficient for fracture stability and fatigue analysis. For a FGM coating on a homogenous substrate some sample results are given in Figures 1-3 [15] where hi and /12 are the thicknesses of coating and the substrate, respectively, c is the crack length, K is constant and the shear modulus of FGM is given by /i(x) = /ioexp(/3a:), ^0 being the modulus of the substrate. Figure 1 shows the effect of material inhomogeneity on /ci in a medium loaded by fixed grips or constant strain EQ. The normalizing stress is given by CTQ = 8/ii£o/(l + /^), /ii := /ioexp(-/5/ii). Figures 2 and 3 show ki for PSZ/Rene 41 FGM coating (/5/ii = 0.375) on Rene 41 substrate loaded by constant strain £0 or constant temperature change AT, respectively. Figure 2 shows the influence of the thickness ratio h2/hi on ki. The effect of the uniform temperature change AT is shown in Figure 3 where To corresponds to the stress-free state and T] and To are the surface temperatures. This is a special case of a general problem in which Ti 7^ To and the medium is under steady-state heat conduction.

. 7.0 6.0 ,

kl

i ; i !

y

h2/h,=0.5 h/h,=1.0 h/h,=2.0 h^h,=10.0

/-

5.0

y/

ic

\

^Oy 2.0 1.0

U—\

,^^^^^^:

-

-'-"'IT-

c/Zii

Figure 1. Mode I stress intensity factor for a surface crack in FGM coating, h\ —h^.

Figure 2. Mode I stress intensity factor for a surface crack in FGM coating, phi = 0.375.

4.2 Spallation Another mode of failure would be the transformation of the surface crack to a T-shaped crack at a relatively weak fracture plane parallel to the surface. This may be microcracks forming along the oxidized splat boundaries or the interface between the thermally grown

no oxide and the coating. There seems to be some evidence of such branching in the results given in [14]. For the T-shaped crack the stress state at the crack tip is one of mixed-mode. Therefore, in the fatigue model, for example, one would have to use A^ rather than Ai^ as the correlation parameter.

Tj=T,=5To Tj=T,=10To T,=T.=20T„

~».^

y

"^^

•r.^

^pL

N

0.025 o 0.020

TT

O^ 0.015

\

arxA"^ "*"-«.»

\

0.010

t\-

0.005

-^^ ^^ *^

1

0.003

^ ^

0.002

1

r_ f

'

0.4

0.6

0.8

1.0

Figure 4 Strain energy release rate for a T-shaped crack under uniform temperature change.

1

\CR2

f

1

LN

r

1

MR1

/^ '

1

0.2

CR1

f

0.001

•

\\.

0.0

- p=:CO

b/l

c/hi Figure 3. Normalized stress intensity factor under uniform temperature rise, ar = 8/ioQ:oTb/(l + K,)

0.004

0.000

p=0.2 p=0.5 p=1 p=3 p=5 p=8

MR2 .

1

1

.HM 1

1

1

"^

Figure 5 Normalized strain energy release rate due to uniform temperature change, ijhi = 5, ^0 = (1 - i^Dids^TfEs-Khi.

Figure 6 Mode I stress intensity factor in an orthotropic FGM subjected to uniform crack surface pressure o-ii(0, X2) = -po, z^ = 0.3

In [14] it was reported that spallation cracks develop in the graded region due to the change in residual stresses caused by the oxidation of the metallic lamellae. Similar observations were made in [16] where a candidate design for an abradable seal was tested 242 hours at 1000°C The seal consisted of a 0.13mm NiCoCrAlY layer, 2.54mm NiCoCrAlY/YSZ FGM region and a 1.27mm low density YSZ. The substrate was a MM247 superaUoy. The spallation occurred in FGM 0.5mm from the initial substrate surface. It was observed that most of the metallic phase in the spalled region was oxidized whereas the part of the seal remaining on the substrate was not. This appears to be due to connectivity of the oxidized region. The continuous oxide layer seems to create a weak fracture plane as well as preventing further oxygen diffusion. Figures 4 and 5 show the strain energy release rate for a T-shaped crack and symmetric edge cracks, respectively. In these examples the substrate is Rene 41 (E5 = 219.7 Gpa,

Ill 1/3 = 0.3, as = 1.67 IQ-^K ) the coating is FGM ( Rene 41 / YSZ, ^Jc = 151 Gpa, Uc = 0.3, Qc = 10"^/°K), h2/hi = 0.16, £/hi = 5,the loading is uniform temperature change AT and the normalizing strain energy release rate is QQ = {1 - u^)(asAT)^Es7vhi. For the FGM coating the modulus variation is given by E{y)

-{

Es, E, -h {Es - Ec)il + ihi/h2) - (y//l2))^

0
/12.

(9)

In Figure 5 a is the length of the edge crack and the results given correspond to homogeneous ceramic coating (IL, p = 00), ceramic-rich (CRl, p = 8 and CR2, p = 2.5), linear (LN, p = 1) and metal-rich (MRl, p = 0.5 and MR2, p = 0.2) FGM coatings, and homogenous metal (HM, p = 0). Note that in both examples Q increases with increasing ceramic content in the coating.

2.0 Figure 7 Variation of the normalized strain energy release rate with aa and 6 in a graded orthotropic medium under uniform tension /to = 1, i^ = 0.3, a22(a;i, ±oo) = po,Qo = np^a/Eo.

Figure 8 Variation of the normalized strain energy release rate with «o and ota in a graded orthotropic medium under uniform tension 0-22(3:1, ± 00) = po, 6 = 1,1/ = 0.3, Qo = Trpga/Eo.

4.3 Effect of Material Orthotropy As mentioned in section 3, because of the processing techniques used, the graded materials are seldom isotropic. Thus, in studying the mechanics of many of the FGMs the medium may be approximated by an inhomogeneous orthotropic continuum. In these materials there are basically two groups of crack problems. In the first the weak firacture plane, and consequently the plane of the crack is parallel to the direction of the material property variation and the related crack problem is one of mode I. The second is a mixed-mode problem in which the crack is perpendicular to the direction of property variation. The plane elasticity problem is solved by introducing the following four parameters for the plane stress problem to replace the engineering elastic constants En, E22, G12 and 1/12 : E = yjEnE22.

P = y/^^ni^x.

6^=En/E22,

KQ = {E/2Gi2) -1^-

(10)

Similar parameters may be defined for plane strain problems. The previous results have shown that the solution of crack problems in FGMs is not very sensitive to the Poisson's ratio. Thus, 1/ is assumed to be constant throughout the medium. In this study it is further assumed that in the graded materials the variations in Eu, E22 and G12 are proportional. Referring to

112 (10), these assumptions imply that «o, v and d are independent of the space variables xi, X2 and the inhomogeneity of the medium may be represented by the function ^(0:1,0:2). The results given in Figures 6-8 are obtained by assuming that xi and X2 are the principal axes of orthotropy and E = Eoexp{ax2). In the mode I problem a crack of length 2a is located along xi =0, -a < X2 < a, and figure 6 shows the influence of the shear parameter KQ and the inhomogeneity parameter a on the stress intensity factor fci. In this case ki turns out to be independent of 6 or the stiffness ratio. More generally, from the formulation of the problem it can be shown that (711(0,0:2) is invariant with respect to a 90-degree material rotation. In the figure «o = 1 corresponds to the isotropic material [17]. Figures 7 and 8 respectively show the effect of a and 6 and a and KQ on the normalized strain energy release rate in an orthotropic FGM with a crack of length 2a located along the xi axis. The problem is one of mixed-mode, the external load is a uniform tension po perpendicular to the crack plane and Go is the strain energy release rate for the corresponding homogeneous isotropic medium for which a = 0, KQ = 1, 6 = 1, ^' = 0.3[18].

REFERENCES 1. M. Yamanouchi, M. Koizumi, T. Hirai and I. Shiota (eds.), 1990, FGM-90, Proc. FGM90, FGM Forum, Tokyo, Japan (1990). 2. J.B.Holt, M. Koizumi, T. Hirai, and Z.A. Munir (eds.), Proc. FGM'92, Ceramic Transactions, Vol. 34, American Ceramic Society (1992). 3. B. Ilschner, and N. Cherradi (eds.), Proc. FGM'94, Presses Polytechniques et Universitaires Romands, Lausanne, Switzerland (1995). 4. H. Liebowitz (ed.). Fracture, Vol. 2, Academic Press (1968). 5. M. Saito and H. Takahashi, Proc. FGM'90 (1990), p. 297. 6. F. Delale and F. Erdogan, Int. J. Engng, ScL, 26 (1988) 559. 7. F. Erdogan, Tr. J. Engineering and Environmental Sci., 18 (1994) 185 8. F. Erdogan, A.C. Kaya and P.F.Joseph, ASME J. Appl. Mech., 58 (1991) 410. 9. P.A. Martin, J. Engineering Mathematics, 26 (1992) 467. 10. Y.D. Lee and F. Erdogan, Int. J. of Fracture, 69 (1995) 145. U . S . Sampath, H. Herman, N. Shimoda and T. Saito, M.R.S. Bulletin, 20 (1995) 27. 12. W.A. Kaysser and B. Ilschner, M.R.S. Bulletin, 20 (1995) 22. 13. M. Finot, S. Suresh, C. Bull, A.E. Giannakopoulos, M. Olsson and S. Sampath, Proc. FGM'94 (1994) 229. 14. M. Alaya, G. Grathwohl and J. Musil, Proc. FGM'94 (1994) 405. 15. M. Kasmalkar, Ph.D. Dissertation, Lehigh University (1996). 16. W.Y. Lee, Y.W. Bae, C.C. Bemdt, F. Erdogan, Y.D. Lee, and Z. Muntasim, J. Am. Ceram. Soc. (submitted for publication) (1996). 17. M. Ozturk and F. Erdogan, Int. J. Engng. Sci. (submitted for publication) (1996). 18. M. Ozturk and F. Erdogan, ASME J.AppLMech. (submitted for publication) (1996).

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 1997 Elsevier Science B.V.

113

Microstructural effects in functionally graded thermal barrier coatings M.-J. Pindera^ J. Aboudi^ and S. M. Arnold^ ^ Applied Mechanics, University of Virginia, Charlottesville, VA 22903, USA ^ Faculty of Engineering, Tel-Aviv University, Ramat-Aviv 69978, Israel "^ Structural Fatigue Branch, NASA-Lewis Research Center, Cleveland, OH 44135, USA Abstract The recently developed two-dimensional version of the higher-order theory for functionally graded materials (H0TFGM-2D) is employed to investigate the effects of microstructural architectures in graded thermal barrier coatings (TBCs) on the stress distributions in the presence of a through-thickness temperature gradient. In particular, the response of TBCs with different levels of functionally graded microstructural refinement and different arrangements is investigated, and the results for the through-thickness stress distributions are compared with those based on the standard micromechanical homogenization scheme. The examples presented illustrate the shortcomings of the standard micromechanics-based approach applied to the analysis of functionally graded TBCs, particularly if the presence of creep effects is included in the analysis. 1. INTRODUCTION Functionally graded materials (FGMs) are ideal candidates for applications involving severe thermal gradients, ranging from thermal structures in advanced aircraft and aerospace engines to computer circuit boards. In fact, this was the original motivation for developing this class of materials. In such applications, a ceramic-rich region of a functionally graded material is exposed to hot temperature while a metallic-rich region is exposed to cold temperature, with a gradual microstructural transition in the direction of the temperature gradient. By adjusting the microstructural transition appropriately, optimum temperature, deformation and stress distributions can be realized. This concept has been successfully employed by many researchers to enhance the thermal fatigue resistance and life of ceramic thermal barrier coatings [1-3]. Due to the many variables that control the design of functionally graded microstructures, full exploitation of the FGMs' potential requires the development of appropriate computational strategies for their response to combined thermomechanical loads. Presently, most computational strategies for the response of FGMs do not explicitly couple the material's heterogeneous microstructure with the global analysis. Rather, local effective or macroscopic properties are first obtained through homogenization, and subsequently used in a global thermomechanical analysis. As discussed previously, this often leads to erroneous results when the temperature gradient is large with respect to the dimension of the inclusion phase, the characteristic dimension of the inclusion phase is large relative to the global dimensions of

114 the composite, and the number of uniformly or nonuniformly distributed inclusions is relatively small [4]. As a result of the limitations of the uncoupled approach, a new higher-order micromechanical theory for FGMs (HOTFGM), that explicitly couples the local and global effects, has been developed [4-7]. The development of the theory has been justified by comparison with the results obtained using the standard micromechanics approach which neglects the micro-macrostructural coupling effects [8]. Herein, the two-dimensional version of the higher-order theory [6-7], briefly outlined below, is employed to investigate the effect of different levels of microstructural refinement on the internal stress field in a functionally graded TBC subjected to cyclic thermal loading in the presence of creep of the ceramic phase. 2. A BMEF OUTLINE OF HOTFGM-2D H0TFGM-2D is based on the geometric model of a heterogeneous composite occupying the region I jci I < «>, o < JC2 ^ H, 0 < ^3 < L (see Fig. 1). The loading applied to the composite in the X2 -x^ plane may involve an arbitrary temperature distribution and mechanical effects represented by a combination of surface displacements and/or tractions. The composite is reinforced in the JC2-JC3 plane by an arbitrary distribution of infinitely long fibers oriented along the xi axis, or finite-length inclusions that are arranged in a periodic manner in direction of the xi axis. The microstructure of the heterogeneous composite is discretized into Ng and Nr cells in the intervals 0<jc2 < ^ and 0<JC3 ^3 »

where the unknown coefficients T[f^] (/, m, « = 0, 1, or 2 with / + m + n < 2) are determined by: satisfying the heat conduction equation, as well as the first and second moment of this equation in each subcell in a volumetric sense; and imposing the continuity of heat flux and temperature in an average sense at the interfaces separating adjacent subcells, as well as neighboring cells, together with the thermal boundary conditions. This provides the necessary 4SNgNr equations for the ASN^Nr unknown coefficients of the form: KT = t

(2)

115

Figure 1. H0TFGM-2D model geometry.

where the structural thermal conductivity matrix K contains information on the geometry and thermal conductivities of the individual subcells (apy) in the N^Nr cells spanning the X2 and JC3 functionally graded directions, the thermal coefficient vector T contains the unknown coefficients that describe the thermal field in each subcell, i.e., r = [ T[\^^\ ...., T^^^^^ ] where jia^y) _ [ Y^^^^ Y^^^^^^ j^^^^^ Y^^^^^ Y^^^^^ Y^^^ j^apy)^ and the thermal force vector t information on the thermal boundary conditions. Once the temperature field is known, the resulting displacement and stress fields are determined by approximating the displacement field in each subcell of a generic cell by a quadratic expansion in the local coordinates jci" , J2 » ^^^ ^3 ^ follows:

«i«w)=HiA+xrMa)

(3)

(4)

4«w=w^«^ »

+xf M+^f

1 .,-(a)2

m\) + y(3-«

1 ,2

j4)m» +

i(3ir - |/.r^)w^A+|(3^f - ji^m^

(5)

where the unknown coefficients WJ(j5m) (i = 1, 2, 3) must be determined from conditions similar to those employed in the thermal problem. In this case, there are 104 unknown

116 quantities in a generic cell (q,r). The determination of these quantities parallels that of the thermal problem. Here, the heat conduction equation is replaced by the three equilibrium equations, and the continuity of tractions and displacements at the various interfaces replaces the continuity of heat fluxes and temperature. Finally, the boundary conditions involve the appropriate mechanical quantities. Application of the above conditions in a volumetric and average sense produces a system of l04NgNr algebraic equations in the field variables within the cells of the functionally graded composite of the form: KU=f+g

(6)

where the structural stiffness matrix K contains information on the geometry and thermomechanical properties of the individual subcells (aPy) within the cells comprising the functionally graded composite, the displacement coefficient vector U contains the unknown coefficients that describe the displacement field in each subcell, i.e., U = [ U[Y^\ .... , f/Jy?^^ ] where U^^^'^^ = [ W^QQ), . . . , 1^3(002) ]^r^^\ and the mechanical force vector / contains information on the boundary conditions and the thermal loading effects generated by the applied temperature or heat flux. In addition, the inelastic force vector g appearing on the right hand side of eqn (6) contains inelastic effects given in terms of the integrals of the inelastic strain distributions Efj^'^^'^\x^'^\x^^\x^'^^) that are represented by the coefficients R\f§[Jl,n), 1 1 1

n

Rm.n) = H ( „ p , ) ^ ^ i S i ^ ) ( i ^ I I I ef/«Wp,(C^"^)/'.(C^>)Pn(;^^)^Cf rf^f rf^?' (7) -1-1

where P/(-), Pm(')^ and P„() are Legendre polynomials of orders /, m and n. These integrals depend implicitly on the elements of the displacement coefficient vector U, requiring an incremental solution of eqn (6) at each point along the loading path. 3. RESULTS The configuration of the investigated TBC is shown in Fig. 2. It consists of a pure zirconia layer at the upper surface that is exposed to hot temperature and a pure CoCrAlY layer at the lower surface that is bonded to a steel substrate. The volume content of the CoCrAlY inclusions in the zirconia matrix gradually increases from the hot to the cold region of the TBC such that halfway through the TBC's thickness the roles of the phases reverse, with the CoCrAlY phase becoming the matrix phase below this point. The thickness H of the TBC is 1 mm, which is the same as the thickness of the steel substrate, while the length and the depth are sufficiently large to be considered infinite. The effect of several different microstructures of the TBC on the internal stress field was investigated under cyclic thermal loading, with the coarsest and the finest microstructures shown in Fig. 3. The thermal loading to which the TBC's top surface was subjected, while the bottom surface was held at 25°C, consisted of a ramp-up to 1200°C in 10 seconds, a hold period of 300 seconds, and a ramp-down to 25°C in 10 seconds. The inplane constraints imposed on the deformation of the entire configuration (TBC + the steel substrate) simulated generalized plane strain in the ^2 -^3 plane (i.e., the average stresses O22 and 033 were required to vanish). The material properties of the individual constituents are given in Table 1. Due to the low conductivity of zirconia, and the significant temperature drop in the pure zirconia layer,

117 C-C1

n

1 Figure 2. Functionally graded zirconia/CoCrAlY TBC bonded to a steel substrate.

Region 1

Region 1

Region 2

Region 3

Region 4

Region 5

Region 7

Region 6

Region 8

Coarse TBC Microstructure

FineTCB Microstructure

Figure 3. Examples of TBC architectures investigated. the response of CoCrAlY and the steel substrate was taken to be elastic with temperature independent elastic parameters. The inelastic response of the zirconia phase was modeled using the power-creep model generalized to multi-axial loading in the following manner: .vp

3F(a,) ,

(8)

118 Table 1. Material properties of the TBC constituents [2]. Material

E (GPa)

v

ajlQ-^/K)

K(W/m-K)

Zirconia top coat CoCrAlY bond coat Steel substrate

36 197 207

0.20 0.25 0.33

8.0 11.0 15.0

0.50 2.42 60.5

where a-^ are the components of the stress deviator, a^ = V3/2a-ya-y, and F{Oe)=Aaee~^^^, with A = 1.89x10"^, n = 1.59, AH = 277 kJ/mole, and R = 8.317 J/(mole K). Herein, we focus on the technologically important normal stress 033 which is responsible for microcrack initiation in the TBC's ceramic-rich portion during the cool down stage of a thermal cycle. This stress component becomes tensile upon cool down due to stress relaxation of the ceramic phase at the elevated temperature. Figure 4 presents the normal stress 033 distributions through the TBC's thickness in two representative cross-sections (see Fig. 2) of the coarsest TBC microstructure at t = 10, 310 and 320 seconds. Also included in the figure is the normal through-thickness stress distribution in a configuration with equivalent homogenized layers (see the regions in Fig. 3) whose properties were generated using the generalized method of cells micromechanics model. Figure 5 presents the corresponding normal stress distributions for the finest TBC microstructure. As observed in all the configurations, significant stress relaxation in the TBC's ceramic-rich region during the hold period when the top surface temperature is maintained at 1200°C leads to a significant tensile stress in this region upon cool down, while elsewhere in the TBC the residual stresses are very small. In the absence of such stress relaxation, the stresses would return to zero everywhere. Comparing the normal stress distributions generated with the higher-order and homogenized analyses we observe that in those regions where stress relaxation is not significant (i.e., in the interior) the homogenized results tend to provide average distributions about which the higher-order results oscillate. The exception is the transition region where the roles of the matrix and inclusion phases are not well defined. In this region, the homogenized analysis consistently underestimates the actual stress distributions. However, the most significant difference between the predictions of the two approaches from the point of view of this study occurs in the pure ceramic layer. In the absence of stress relaxation effects, the same stress distributions in this region are obtained using the two approaches. Although not shown, this result can be deduced by comparing the stress distributions at t = 10 seconds where little relaxation is observed. When stress relaxation occurs, however, substantially different stress redistributions in the pure ceramic region are predicted by the two approaches that may have significant implications with regard to microcrack initiation and propagation. In particular, at the end of the cool down cycle (t = 320 seconds) the homogenized analysis predicts slightly higher normal stress directly at the top surface which, however, decays much faster than its counterpart obtained from the higher-order analysis. Comparing the stress distributions in the coarse and the fine TBC microstructures generated with the higher-order analysis, we observe virtually no difference in the pure ceramic region. In the functionally graded regions, differences are expected due to differences in the microstructural scales. However, one common feature exhibited by the stress distributions in the two microstructures is the common envelope within which the stress oscillations in the

119

CO Q-

-750 1000

1250

1500

1750

2000

1750

2000

1750

2000

X2 (^m)

500

-I10 seconds 310 seconds 320 seconds

Q-

^ -750 1000

1250

1500 X2 (^m)

500 10 seconds 310 seconds 320 seconds

CO

a.

-750 1000

1250

1500 X2 (\im)

Figure 4. Through-thickness normal stress distributions in the functionally graded TBC with the coarsest microstructure: (top) cross section c-c 1 (H0TFGM-2D analysis); (middle) cross section c-c 2 (H0TFGM-2D analysis); (bottom) homogenized analysis.

120

CO

I -750 1000

1250

1500

1750

2000

1750

2000

1750

2000

X2 (^m)

500

CO QL

I -750 1000

1250

1500 X2 (fim)

500

-410 seconds 310 seconds 320 seconds

CO Q.

I -750

1000

1250

1500 X2 (^m)

Figure 5. Through-thickness normal stress distributions in the functionally graded TBC with the finest microstructure: (top) cross section c-c 1 (H0TFGM-2D analysis); (middle) cross section c-c 2 (H0TFGM-2D analysis); (bottom) homogenized analysis.

121 coarse and the fine microstructures occur. This common stress envelope could be useful in the design of functionally graded TBCs with different microstructural scales. 4. CONCLUSIONS Comparison of the coupled microstructural and homogenization-based analyses of functionally graded TBCs points to the significant impact of micro-macrostructural coupling on the stress field in the presence of stress relaxation in the TBCs ceramic-rich region. In contrast with the elastic analysis, where the coupled and homogenized-based approaches produced differences only in the TBCs functionally graded region, substantial differences in the pure ceramic region are now predicted by the two approaches when the ceramic phase is allowed to relax. This has significant implications on accurate modeling of crack initiation and growth due to stress reversal in the TBCs ceramic region during thermal cool down.

5. REFERENCES 1.) Jian, C. Y., Hashida, T., Takahashi, H., and Saito, M., "Thermal Shock and Fatigue Resistance Evaluation of Functionally Graded Coatings for Gas Turbine Blades by Laser Heating Method," Composites Engineering, Vol. 5 (7), pp. 879-889, 1995. 2.) Kokini, K., Takeuchi, Y. R., and Choules, B. B., "Surface Thermal Cracking of Thermal Barrier Coating Owing to Stress Relaxation: Zirconia vs. Mullite," Surface and Coatings Technology, Vol. 82, pp. 77-82, 1996. 3.) Kawasaki, A. and Watanabe, R., "Evaluation of Thermomechanical Performance for Thermal Barrier Type of Sintered Functionally Graded Materials," Composites: Part B (Engineering) (in press). 4.) Aboudi, J., Pindera, M-J., and Arnold, S. M., "Elastic Response of Metal Matrix Composites with Tailored Microstructures to Thermal Gradients," Int. J. Solids and Structures, Vol. 31 (10), pp. 1393-1428,1994. 5.) Aboudi, J., Pindera, M-J., and Arnold, S. M., "Thermo-Inelastic Response of Functionally Graded Composites," Int. J. Solids and Structures, Vol. 32 (12), pp. 1675-1710, 1995. 6.) Aboudi, J., Pindera, M-J., and Arnold, S. M., "Thermoelastic Theory for the Response of Materials Functionally Graded in Two Directions," Int. J. Solids and Structures, Vol. 33 (7), pp. 931-966, 1996. 7.) Aboudi, J., Pindera, M-J., and Arnold, S. M., "Thermoplasticity Theory for Bidirectionally Functionally Graded Materials," Journal of Thermal Stresses (in press). 8.) Pindera, M-J, Aboudi, J., and Arnold, S. M., "Limitations of the Uncoupled, RVE-Based Micromechanical Approach in the Analysis of Functionally Graded Composites," Mechanics of Materials, Vol. 20 (1), pp. 77-94, 1995. Acknowledgements - The authors gratefully acknowledge the support provided by the NASA-Lewis Research Center through the grant NASA NAG 3-1377.

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I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

123

Micromechanical failure criterion for FGM architecture studied via disk-bend testing of Zr02/Ni composites T. Ishizuka*, Y. Ohta and K. Wakashimat Precision and Intelligence Laboratory, Tokyo Institute of Technology 4259 Nagatsuta, Midori-ku, Yokohama 226, Japan

An investigation has been made to find a reasonable fracture criterion needed for a micromechanics-based approach to the architecture of ceramic/metal functionally graded materials (FGMs). Emphasis is placed on an equibiaxial plane-stress state of loading because it is essential in an anticipated service condition of FGMs for heat-shielding structural application. Disk-bending tests are carried out on a series of Zr02/Ni composites with different compositions. Results of the tests are analyzed by a micromecanical approach which takes into account the effect of plastic deformation of the ductile phase. A conceptually simple fracture criterion is derived in terms of the stress in the brittle phase.

1. INTRODUCTION Over the past decade since the concept of functionally graded materials (FGMs) emerged in Japan, many investigators have found their research interest in this novel type of composite materials [1-4], Among various important problems on FGMs from a solid mechanics point of view, the FGM architecture, i.e. a design methodology to find the so-called "optimal gradation" in composition and microstructure for FGMs of heat-shielding structural use, is an issue of prime importance, in which we have been involved [5-10]. The work reported here is also concerned with the FGM architecture, wherein we address an open question as to the means for identification of the location where failure is most likely to occur in a ceramic/metal FGM. Our previous approach to this identification was quite crude indeed; for example, refer to [10] for an outline. In what follows, we first describe an experiment done by "disk-bend" testing. This method of testing is advantageous to the matter concerned because it allows us to make an equibiaxial plane-stress state of loading to be met in individual "layers" constituting an FGM. However, its application to materials exhibiting inelastic deformation has not been established yet. Under the circumstances, we devise a new method for the analysis of test data.

* Present address: Technical Research & Development Institute, Japan Defence Agency, 2-2-1 Nakameguro, Meguro-ku, Tokyo 153, Japan. t Author to whom correspondence should be addressed.

124 2. DISK-BEND TESTING 2.1. Specimens and test method The ceramic/metal system consisting of partially stabilized zirconia (PSZ) and nickel (Ni) was chosen in this work. Using commercially available powders (ZiO^-l mol% Y2O3 Tosoh TZ-2Y and Mond carbonyl nickel Type-123), a series of non-graded specimens with different compositions were fabricated by hot die-compaction of powder mixtures at 1683 K for 1.8 ks under a compacting load of 35 MPa. The nominal compositions of the specimens were 0, 20, 40, 60, 80 and 100 vol% PSZ (to be referred to as Ni, 0.2PSZ, 0.4PSZ, 0.6PSZ, 0.8PSZ and PSZ, respectively). Typical microstructures of the composite specimens are shown in Fig. 1; it can be seen that the PSZ phase exhibits tablet-shaped appearances and is almost uniformly distributed in a matrix of the Ni phase. As will be described later, these PSZ particles are treated as oblate spheroidal ones with a fixed aspect (thickness-to-diameter) ratio of 0.4 in our micromechanical analysis of stress states. Apparent densities measured by the Archimedean method were slightly lower than the ideal ones, the resulting fractional void contents (f^) being 0.07, 0.02, 0.03, 0.04 for 0.2PSZ, 0.4PSZ, 0.6PSZ and 0.8PSZ, respectively. This is also taken into account in the analysis by assuming the void shape to be spherical.

Fig. 1 Microstructure of PSZ/Ni composites in sections parallel (left) and perpendicular (right) to the pressing surface. From the six different specimens above, test pieces for disk-bend testing were prepared; their dimensions were 20 mm in diameter and 0.5-1.5 mm in thickness. A specially designed loading fixture, as schematically depicted in Fig. 2, was employed for the testing on an Instron-type loading machine, whereby the test pieces were radially four-point bent at room temperature. The inner and outer spans were fixed: 2r~9 mm and 2r^=18 mm. Deflections w on loading were measured with an LVDT via a quartz rod-in-tube displacement transmitter, cf. Fig. 2. By assuming a spherical bending configuration in the center portion r < ^, the

125 equibiaxial strain e on the tension face was determined from a geometrical relationship: e = twlr^, where t is the thickness of test piece. A maximum bending stress o corresponding to an applied load P was calculated using a theoretical formula available in the literature [11]:

\{l-v)[rl-r^)/^] = |3P/(4jrr')|^

[+2(l+v)ln(r,/rJJ where 2R is the test-piece diameter and v the Poisson ratio. Note however that this stress o is an apparent one in that the formula above has been derived by assuming linear elastic deformation.

(lx)ad celf)

Fig. 2 Disk-bend testing device 2.2. Test results Generally, results of the disk-bend testing showed composition-dependent behavior. A typical set of test data is displayed in the o vs € diagram, Fig. 3. It can be seen that all but one specimen (monolithic PSZ) exhibit nonlinear stress-strain curves owing to the presence of ductile constituent. Highly ductile behavior (without fracture up to a strain of 4 %) is observed for monolithic Ni and 0.2PSZ composite specimens, Fig. 3(b); this is due to nonuniform plastic deformation like "deep drawing". Test-piece appearances after the tests are shown in Fig. 4. It is evident that the specimens displayed in Fig. 3(a) fracture in an essentially radial cracking mode. Observations by scanning electron microscopy revealed that this type of fracture occurs predominantly by cracking of the PSZ phase—not by interfacial debonding between the PSZ and Ni phases. In Fig. 5, the fracture strain e^ is plotted against the PSZ volume fraction f^^\ this result will be used in the following analysis.

Fig. 3 Observed in-plane apparent stress vs in-plane macrostrain in the tension face for specimens (a) fractured in equibiaxial tension and (b) extensively deformed.

126

(a) PSZ

^ ^ 9

(d) 0.4PSZ

(e) 0.2PSZ

(0 Ni

Fig. 4 Test-piece appearances.

'0.2

1.0

0.4 0.6 0.8 Volume fraction of PSZ,/psz

Fig. 5 Tension-face fracture strains for samples fractured in equibiaxial tension.

3. ANALYSIS OF TEST DATA 3.1. Method of the analysis As shown above, the ZrO/Ni composites examined by disk-bend testing are found to deform in a nonlinear manner, so that composition-dependent fracture strengths cannot be obtained directly from the stress-strain diagram in Fig. 3. Under the circumstances, we now make a micromechanical analysis to estimate actual stresses to be developed by plastic deformation of the ductile constituent on the basis of an established "mean-field" model [12]. In the following, the macrostress (a) is related to the microstresses {p)^^ and (0)^4 such that (a) = 7J>sz(<^)psz + /Ni(^)Ni' where ( ) indicates the volumetric average over an appropriate domain. Specifically, we are interested in a simple equibiaxial plane-stress state of loading: In view of the symmetry in microstructures, (a) = [ajp a22, 0, 0, 0, O] with o^^=o^^so. the macrostrain (e) to be linked with (a) will be such that (e) = [eii, ^22* ^33» ^» ^' ^] with ^11 = ^22 = ^' ^^t ^he out-of-plane component ^33 is immaterial for the present analysis. The in-plane macrostrain e is decomposed into the elastic and inelastic (plastic) parts, so that we write the increment de 2iS de ^ de"" + de^ ^ [S^ + S'^)da. Here, S^ is the plane-stress elastic compliance, given by 5^ = [C~^ )u\ '^ (^^)ii22 ^^ ^^^^^ ^^ ^^^ components of effective elastic compliance tensor C ^ Likewise, S^ is defined for the plastic part. Below, this plastic compliance under the specified state of loading is calculated by taking into account the effect of thermally induced residual stresses. For the present three-phase composites with aligned (spheroidal) PSZ particles and (spherical) voids in the Ni matrix, the mean-field model provides the following set of relations: (a)^i

= LXO) + M,(E^Z - ^Ni)»

(E)==C-\a) + N , E ; + N 2 8 ;

HPSZ = ^2(0)

+ M2(8*ps2 -

z^,)

ai.2) (2)

127 Here, z^- and Ep^^ are homogeneous inelastic strains defined in the respective phase domains; L^ and L^ are dimensionless constants satisfying f^.L^ + /pszLj == I (I • the identity tensor of rank 4); Mj and M2 are constants having the same dimension as the elastic modulus and satisfying /^jM, +/pszM2 = 0 ; N^ and N2 are also dimensionless constants relating to h^ such that I - N i = N2 = fpsz^l (^2' the transpose of L2). All the constants, as well as the effective elastic compliance C~\ are calculated once the shapes of the PSZ and void phases are specified together with the elastic constants and volume fractions of Ni and PSZ phases. Now, let us put BI =[a^,, a^,, a^„ 0, 0, o f ( r ~ r J + [£P., < , -2^;^,, 0, 0, o f and 0, 0, 0] ( r - 7 ; ) , where a denotes the coefficient of thermal expansion, T -T^ is the temperature change from a stress-free reference state, and e^. is a quantity measuring the magnitude of plastic strain in the Ni phase. And further, we suppose that the Ni phase is a von Mises' solid whose flow curve is expressed by Swift's equation: <^i^ = P\^+f^^2)

with constants /?, q and n. Here, a ^ and de^^ are the equivalent stress

and the equivalent plastic strain increment, now given as a ^ ^ K^ii)Ni ~(*-^33)Ni| ^^^ ^ C =2|d£P.|. From Eq. (1,), we have {on)^,={a,X

=«cr + M«Psz -aJ{T-T,)

+ ce^,

^Ni with constants a, b^ c, etc.; the other components of {o)^. all vanish.

Therefore, de^. =-k{a-a')da

+ k{b- b')[a^^ -

a^.)dr,

where \/k ^2s,s,[c-c^)p(q^2^\del^^~' with s,^{{aX-{a,X)f h X -{^A\ ^^^ ^2 = '^^Ni / k^Nil- Note that this de^. is directly linked with de^ using Eq. (2). Consequently, the plastic compliance 5*^ is determined and hence, our method of analysis has been formulated. In what follows, we are particularly interested in the microstress in the PSZ phase; it is calculated from Eq. (I2) without difficulty. 3.2. Computed stress-strain curves The in-plane macrostress a vs inplane macrostrain e curves computed with appropriate values of material constants are shown in Fig. 6. The corresponding in-plane and out-ofplane microstress components in the PSZ phase, (a,,)p3^ and {a^^) PSZ' ^ ^ ^

displayed as a function of the in-plane macrostrain e in Figs. 7(a) and 7(b), respectively. In these diagrams, solid and broken lines indicate results with and without consideration of thermal residual stresses to be induced in the cooling stage of hot consolidation; for full details of the computation, refer to [13].

Strain, e /10"^

Fig. 6 In-plane macrostress vs in-plane macrostrain curves in equibiaxial tension, computed with (broken lines) and without (solid lines) taking account of residual stresses.

128 0.8PSZ

0.6PSZ

0 -200i -400 -600 (b) Out-of-plane microstress in PSZ particle

-80ol (a) In-plane microstress in PSZ particle -400

0

1 Strain, e /10"^

1 Strain, e / lO^^

Fig. 7 In-plane and out-of-plane microstresses in PSZ particles vs in-plane macrostrain curves, (a) and (b), computed with (broken lines) and without (solid lines) taking account of residual stresses. 3.3. Evaluation of fracture Strengths Since the e values at fracture are determined experimentally (Fig. 5), we can read the fracture strength levels from Figs. 6 and 7. This is immediately done for the macrostress a. The result is shown in Fig. 8 (open and solid triangles), which provides a linear relation: Of /MPa = 200 + 400/ps2;. While this relation can be used as a fracture criterion, a better criterion can be deduced from Fig. 7, as described below.

800 -

600

«> 400 200

0.4 0.6 0.8 Volume fraction of PSZ,j^s2

Fig. 8 Fracture strengths, in terms of in-plane macrostress (triangles) and maximum "equivalent normal stress" in PSZ phase (circles), plotted against PSZ volume fraction. Solid and open markings refer to the cases with and without taking account of residual stresses.

In view of the particle cracking observed in most composite specimens, we focus attention on the microstress {ojp^^, ^^ ^^^ ^^^ phase, and suppose that fracture occurs when a latent microcrack starts to grow. Since {a)psz is triaxial ((c^ii)psz "(^22)PSZ ^^' (^33)PSZ ^^^^ we define the stress triaxiality /3 as j3 = (^33)^^^/ (<^ii)psz- ^ ^ further assume that latent microcracks are penny-shaped and randomly oriented. Under the circumstances, we adopt a mixed-mode crack extension criterion of the form: K^ + KI + Kf^^ /(l - v) 2: KI^ and define an equivalent normal stress a^ such that

a:=^al+4{r/i2-v)Y

(3)

129 where a^ and T are the normal and shear stresses acting on the crack surface (other symbols have the usual meanings). Then, the criterion is rewritten as

> K^c y where a

is the crack radius. Accordingly, our task now is to examine the maximum value of a^ given

a^ /((^ii)psz = Jsin^O + p^cos^6 + ih + 2{( 1 - p)/(2 - v)}'Isin'O cos'0

(4)

where 6 is the angle between the crack surface normal and the .Xg-axis. In Fig. 9, o^ /(^ii)psz is plotted against 6 for several different /3 values (in the range -1^/3 :s 1) with V = 0.25. From the diagram, we see the following: if -0.53 ^ j8 < 1, then (a^) = {o^^)^^ and a latent crack oriented at 0 = Jt/2 will extend; otherwise (-1 ^ /3 < -0.53), a mixed-mode crack extension results {ji/4<6 <jr/2). From Figs. 7(a) and 7(b), we also see that j8 is within the former range in most cases (-0.53 ^ p
Fig. 9 Diagram showing the variation of "equivalent normal stress" in PSZ phase with the orientation angle of penny-shaped crack {&) and the stress triaxiality (/S).

Crack orientation angle, 6 / rad

4. SUMMARY AND CONCLUSION Through disk-bend testing on a series of Zr02/Ni composite specimens fabricated by powder processing, we have examined the fracture behavior of ceramic/metal composites under an equibiaxial plane-stress loading, and derived, by making a micromechanical analysis of elastoplastic stress states, a brittle phase-controlled fracture criterion of the form, (^n^)max ^ ^^^^^'^ in terms of the equivalent normal stress a^. This criterion is conceptually simple and quite useful particularly for our micromechanics-based approach to the FGM architecture.

130 REFERENCES 1. M. Yamanouchi, M. Koizumi, T. Hirai and I. Shiota (eds.), Proc. of FGM'90, 1990. 2. J. B. Holt, M. Koizumi, T. Hirai and Z. A. Munir (eds.), Ceramic Trans., Vol. 34 (Proc. ofFGM'92), 1992. 3. B. Ilschner and N. Cherradi (eds.), Proc. of FGM'94, 1994. 4. M. J. Pindera, S. M. Arnold, J. Aboudi and D. Hui, Composites Eng., 4 (1994) 19. 5. K. Wakashima, T. Hirano and M. Niino, ESA SP-303, p. 97,1990. 6. K. Wakashima and H. Tsukamoto, Ref. 1, p. 19. 7. J. Teraki, T. Hirano and K. Wakashima, Ref. 2, p. 67. 8. T. Ishizuka and K. Wakashima, Ref. 3, p.279. 9. K. Wakashima, T. Ishizuka and H. Tsukamoto, Proc. of COMMP'93, p. 392, 1993. 10. T. Hirano and K. Wakashima, MRS Bull., 20 (1995) 40. 11. M. N. Giovan and G. Sines, J. Am. Ceram. Soc, 62 (1979) 510. 12. K. Wakashima and H. Tsukamoto, Mater. Sci. Eng., A146 (1991) 291. 13. T. Ishizuka, Doctoral dissertation, Tokyo Institute of Technology, 1995.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

131

Thermomechanical response characteristics of Zr02/Ni functionally graded materials: An experimental study to check model predictions T. Ishizuka*, C. S. Kang and K. Wakashimat Precision and Intelligence Laboratory, Tokyo Institute of Technology 4259 Nagatsuta, Midori-ku, Yokohama 226, Japan

Thermal loading tests have been carried out on Zr02/Ni FGM specimens fabricated by hot die-compaction of composition-graded powder layups. Specimens with four different gradations designed so as to possess the same heat-shielding capability are examined. During repeated heating over the entire surface of ZrOj skin layer, an indication of cracking is monitored by means of acoustic emission (AE). Results of the tests are analyzed in terms of the transient through-the-thickness distribution profiles of thermal stresses particularly in the Zr02 phase, which are computed by a micromechanics-based approach reported previously. Observed differences in damage tolelance of the specimens are reasonably understood.

1. INTRODUCTION In previous work [1-4] on ceramic/metal functionally graded materials (FGMs) for application to high-temperature heat-shielding structural components, we have developed a micromechanics-based computational approach to the FGM architecture, i.e. a design methodology to find the so-called "optimal gradation" in composition and microstructure. This paper reports a piece of our experimental work that has been done toward substantiation of the computational approach.

2. THERMAL LOADING TESTS 2.1. Specimens The ceramic/metal FGM system consisting of partially stabilized zirconia (PSZ) and nickel (Ni) was chosen in this work. Using commercially available powders (Tosoh TZ-2Y and TZ-3Y20A; Mond carbonyl nickel Type-123), disk-shaped specimens for thermal loading tests were fabricated by hot die-compaction of composition-graded powder layups at 1683 K * Present address: Technical Research & Development Institute, Japan Defence Agency, 2-2-1 Nakameguro, Meguro-ku, Tokyo 153, Japan. t Author to whom correspondence should be addressed.

132 for 1.8 ks under a compacting pressure of 35 MPa. Four different types of compositional gradation were dealt with; see Fig. 1. According to the supplier, higher-strength products can be obtained with TZ-3Y20A (3 mol% Y2O3-PSZ containing 20 mass% AI2O3) than with TZ-2Y; thus TZ-3Y20A was used only for the PSZ skin layer. These specimens were designed so as to possess almost the same overall heat transfer coefficient k = (^^, /AJ «2.6kWm~^K~\ where A. and t^ are the thermal conductivity and thickness of the ith layer. Thus, while their overall thicknesses are different, the heat-shielding capability is expected to be the same. For microstructural details of the specimens, refer to [5]. Briefly, the PSZ phase is present in the form of nonequiaxed (tablet-shaped) particles with a typical thickness-to-diameter ratio of 0.4 and the Ni phase is the matrix; this microstructural feature is even true for the PSZ-rich layers. 2.2. Testing method Thermal loading tests were carried out using an apparatus illustrated in Fig. 2. A spheroidalreflector infrared lamp heater with a quartz-glass radiation guide rod allowed nearly uniform heating over the surface of the PSZ skin layer of a specimen (dia. 20 mm), which was mounted with silicon grease on a stainless steel stage attached to a water-cooled copper heat sink. A thermal video system (thermography) was employed to measure a radial temperature distribution over the top surface, whereas the temperature at the bottom was estimated from the measurements with three equally-spaced thermocouples inserted into the specimen stage. A temperature drop within the intervening grease layer was then taken into account. Repeated thermal loading was applied to each specimen by controlling the power supplied to the lamp heater at four different preset levels of 0.6, 0.7, 0.8 and 0.9 kW. The power-on duration of each heating step was 60 s, and in between one heating and next a time over 600 s was consumed so that the specimen was cooled enough. An indication of cracking during the tests was monitored by means of acoustic emission (AE). PSZ 0.8PSZ 0.6FSZ 0.4PSZ 0.2PSZ Ni (4) Sp^ixmn l>

1 mm Fig. 1 (left) Four different gradations and (right) an example of graded structure (specimen B).

Fig. 2 Experimental setup for thermal loading tests.

133 3. RESULTS OF THE TESTS 3.1. Uniformity of surface heating In the present investigation we are particularly interested in the thermomechanical response characteristics of the specimens under conditions of uniform through-the-thickness heat flow. Thus, the uniformity of surface heating is quite important. Figure 3 shows an example of the observed radial temperature distributions over the upper specimen surface upon heating. While the temperature gradually decreases from the center to the periphery of the specimen, the maximum temperature differential AT^ in the radial direction is small in comparison with the temperature Tjj at the center. For instance, AT^/Tjj^O.ll and 0.03 after heating for 5 s and 60 s, respectively. Thus, we thought it reasonable to neglect all radial effects and assume the ideal condition of unidirectional heat flow. 3.2. Response to repeated thermal loading Changes in the upper and lower surface temperatures, Tjj and 7^, during repeated thermal loading are shown together with the count rates of AE event in Fig. 4. It can be seen that^u rises very sharply in a few seconds and then continues to increase gradually. In contrast, TJ^ initially remains almost unchanged but afterwards it rises gradually. This was because the specimen was cooled with a limited amount of water without circulation to minimize the noise in AE measurement. While both Tjj and 7^ increase at finite rates, a quasi-steady state will be attained when the difference Tjj - TL becomes almost unchanged; such a state is probably met when the surface heating is done at low-power levels, P = 0.6 and 0. 7 kW. (b)

0,9kW j

^^^^

oL

\

0.8kW 1

1 e

O.TkW 1

0.6kW 1

- . - U p p e r s n i t u * tt Ldwersof&KKte

-10

-5 0 5 Radial position, r /mm

300

r"'

K1

h---—T~^-~'--]

Fig. 3 Uniformity of surface heating. lOOOi

1

Power level of IR heater, P /kW

Fig. 5 Observed temperature differential. Fig. 4 Temperature and AE recordings for specimens (a) A, (b) B, (c) C and (d) D.

134

In Fig. 5, the observed temperature differential Tjj -TL at / =60S is plotted against the power level of heating. Note that the results are much the same for all the specimens, particularly for the data at P = 0.6 and 0.7 kW. This is what we have expected indeed. As already mentioned, all the specimens are designed to possess almost the same thermal response characteristics, i.e., they have nearly equal values of the overall heat-transfer coefficient ( k « 2.6kWm"^K"*). Nevertheless, the mechanical response characteristics of the specimens differ significantly. Note that AE signals are detected only in specimens A and D. In fact, no evidence of cracking was observed in specimens B and C after the testing.

4. ANALYSIS OF THE TEST DATA AND DISCUSSION To analyze the above test data in terms of thermally induced stresses, we have made a set of computations in the following manner. First, using a finite difference method based on the Crank-Nicolson time integration scheme, the temperature distributions within the specimens were estimated from the observed surface temperature profiles shown in Fig. 4. Next, the corresponding distributions of thermal stresses were computed with a micromechanics-based approach reported previously [3,4]. In this approach, we take into account plastic deformation of the ductile metal phase following an established "mean-field" model [6]. Thus, not only macro- but microstresses will be examined. The macrostress (a) is simply related to the volume-averaged microstresses (a)ps2 and {o)^. in the ceramic and metal phases such that (^) = /psz(^)psz + /Ni(^)Ni' where ( ) indicates the volume average over an appropriate domain. We are interested in a simple equibiaxial plane-stress state of thermomechanical loading: (a)=[aij, a22, 0, 0, 0, O] with o^^--o^^^o. In Fig. 6, transient through-the-thickness distribution profiles of the in-plane macrostress o upon the final heating at 0.9 kW, together with those on the subsequent cooling, are displayed for each specimen. The top and bottom edges of each diagram, i.e., z/h^X and - 1 , correspond to the upper and lower specimen surfaces, and such numerals as 0.1, 1, 60 indicate the elapsed times in seconds after the onset of heating. For full details of the computation, refer to [7]. As is evident from the figure, significantly different stress distribution profiles are simulated as a consequence of different compositional gradations. Note, in particular, that for specimens A and D the PSZ skin layer suffers large (equibiaxial) tensile stresses in its lower portion, whereas for specimen B the stress in that layer is always compressive everywhere. 0.9kW

(b) o"^6o

0 ~ ^ Macrostress, oMPa

1

Macrostress, o/MPa

(d)

Macrostress, o/MPa

L

Macrostress, o^MPa

Fig. 6 Computed distribution profiles of in-plane macrostress.

135

d

On healing

7\

On cooling

0 200 0 0 0 Out-ofplane microstress in PSZ jAase, o^/MPa

On heating

On cooling

0 200 0 0 0 Out-crf-plane microstress in PSZ pAiase, o^/MPa After testing J atO.SkW i

' - "^^

-L After testing J at 0 7kW /

" L. After testing! atO.ekw/

On heating

On cooling

-}..

/

£z

On heating

/

0.8kW

120 61] 601

'^•A

On cooling

im1 lA " ^

On heating

(

:

K360,1

'""•'^

0.1..1 60

I_

0.9kW

120

0.1 1 60

/

^•"^^

rr?'

0.1..,i ,60

^

Z_ U

0.7kW

12061

On cdoling

0.6kW

1

120^^60 1

On heating On cooling After fabricE tiai 0 200 0 0 0 In-plane microstress in PSZ phase, c^gj/MPa

(b)

J CL]

n heating ^

n

On heating

On cooling

0 200 0 Out-crf-plane microstress in PSZ piiase, a j:^/MPa

O.r 1 60

12q6L6n

Ar /J/ ..-On fieating

On coolin*

0.8kW i\60

/,!. On heating

On cooling

1

^%

fi

o.i'.oeo .On heating

O.U.I J

"^ ^

After fabrication On healing On cooling 0 200 0 0 0 In-plane microstress in PSZ 0iase, af^ /MPa

(c)

After fabrication 0 200 In-plane

On heating

)60.1

/X 4 120/ On cooling

IP ^On cooling

1 PSZ phase, a j's^/MPa

(d)

Fig. 7 Computed distribution profiles of in-plane and out-of-plane microstresses in the PSZ phase.

136

Fig. 8 Fracture strength data for PSZ. Temperature, r/K

In a separate investigation [5], we have examined the fracture behavior pertinent to the ceramic/metal FGM system concerned, and proposed a criterion for the initiation of cracking in the brittle ceramic phase. Basically, we consider a stress triaxiality p s a^^ /<^rez i^ ^^^ PSZ phase, which is generally of the form (a)ps2 = [<7Sz> ^sz» ^rez' 0, 0, O] , and presume that cracking occurs in the PSZ phase when one of the penny-shaped cracks latent in 3-D random orientation starts to grow.

The criterion is given as (a^)

= const, with an

"equivalent normal stress" a^ acting on a latent crack. For the above stress state, a^ is given as a^ ^a^Jsme+ p^os'S + 2[i5 + 2{(l - p)/[2 -v)}']sin'0cos'0, where 6 is the angle between the crack surface normal and the z-axis (taken in the thickness direction of the specimen). In Fig. 7, the computed distribution profiles of in-plane and out-of-plane microstresses in the PSZ phase, to be compared with those in Fig. 6, are shown (shaded diagrams); here, results on the previous thermal loadings at 0.6, 0.7 and 0.8 kW are also shown only for the in-plane component. As has been discussed in [5], if -0.53 ^ jS < 1, then a latent crack oriented vertically (0 = JT /2) will extend. Since this condition for /3 was found to be met in all the cases shown in Fig. 7, (p^) is now equivalent to the maximum tensile level of apsz- The constant in the above criterion is simply the fracture strength of the monolithic PSZ under in-plane equibiaxial loading, which can be determined by disc-bend testing; results of the testing are shown in Fig. 8. From this analysis, it was found that the integrity of the specimens under the repeated thermal loading can, in general, be reasonably understood. Specifically, we emphasize that large in-plane tensile stresses generated in the lower portion of PSZ skin layer are very critical for damage initiation; compare results for specimens A and D with those for specimens B and C. However, the observed AE events are not altogether ascribable to this skin layer cracking. In other words, the crack resistance due to the presence of the metal phase should be considered. This is definitely an inportant issue for further work. REFERENCES 1. 2. 3. 4. 5. 6. 7.

K. Wakashima, T. Hirano and M. Niino, ESA SP-303, p.97, 1990. K. Wakashima and H. Tsukamoto, Proc. of FGM'90, p. 19,1990. K. Wakashima, T. Ishizuka and H. Tsukamoto, Proc. of COMMP'93, p.392,1993. T. Ishizuka and K. Wakashima, Proc. of FGM'94, p.279,1994. T. Ishizuka, Y. Ohota and K. Wakashima, to appear in Proc. of FGM'96, 1996. K. Wakashima and H. Tsukamoto, Mater. Sci. Eng. A146 (1991) 291. T. Ishizuka, Doctoral dissertation, Tokyo Institute of Technology, 1995.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

137

Micromechanical approach to the thermomechanical analysis of FGMs Seiichi Nomura and Donna M. Sheahen

Department of Mechanical and Aerospace Engineering, University of Texas at Arlington, Arlington, TX 76019-0023, USA. ABSTRACT The thermoelastic fields in a functionally graded material (FGM) are derived semi-analytically by the use of the eigenfunction expansion method. The eigenfunctions are approxinaated by a linear combination of admissible functions each of which satisfies the homogeneous boundary condition. The eigenfunctions can also be used to construct the Green's function for the FGM which enables the handling of various boundary conditions and source terms. A numerical example is presented to show thermal stress distributions in a 2-D rectangular FGM under steady-state heat transfer subject to the first kind of displacement boundary condition.

1

INTRODUCTION

Functionally graded materials (FGMs) distribute the material functions throughout the material body to achieve the maximum heat resistance and mechanical properties ideal for spacecraft where one side may be exposed to extremely high temperatures and the other side may be exposed to extremely low temperatures. Because FGMs characteristically have continuously varying material properties, many analytical methods developed for conventional composites with distinct phases may not be directly applicable to FGMs. In this paper, a micromechanical approach is used to semi-analytically obtain the thermoelastic fields in an FGM by the eigenfunction expansion method. The physical fields in an FGM can be expanded by eigenfunctions defined for the material properties, and the geometrical boundary and the coefficients of each eigenfunction can be determined by the Galerkin method [1]. The eigenfunctions can also be used to construct the Green's function for the FGM in series form that can handle the change in the source term and the boundary conditions [2]. The eigenfunction and the accompanying eigenvalues represent the intrinsic nature of an FGM and can be used to characterize the FGM. The

138 eigenfunctions are approximated by a linear combination of permissible functions which satisfy the homogeneous boundary condition for the given geometry chosen from a pool of polynomials and transcendental functions. The implementation of the proposed method for multi-dimensional and heterogeneous bodies is novel. As an example of this method, the thermal stress distribution in a 2-D rectangular elastic body under steady-state heat conduction subject to the first kind of displacement boundary condition is shown.

2

ANALYSIS

The equilibrium equation of inhomogeneous linear elastic materials with temperature effects is expressed as (C7i,i/(r)uM(r))j - V • ( ( 7 ( r ) a ( r ) A r ( r ) ) - 0

(1)

with the prescribed boundary condition where C(r), w(r), a ( r ) and A T ( r ) are the elastic modulus, displacement, thermal expansion coefficient and temperature rise, respectively. In equation (1), the repeated index denotes summation and the symbol j denotes partial derivative with respect to Xj. As mentioned in the introduction, an FGM is characterized by continuously varying C(r) and a ( r ) . The displacement field, Uk{r), can be expressed by the eigenfunctions, t/jjf (r), as

«,(r) = ^ c ^ V . »

(2)

A

where ^j^(r) is the A-th. eigenfunction defined for the following eigenvalue problem: iC,jki{r)rl^tMh

+ A^V-.^Cr) = 0

(3)

where no summation is taken over A, X"^ is the A-th eigenvalue, and the eigenfunction, V'^(r), must satisfy the homogeneous boundary condition. It is noted that the eigenfunctions are mutually orthonormal as

i4>f,^f) = SAB

(4)

where SAB is the Kronecker's delta and (/, g) denotes the inner product defined as

{f,g) = J^fir)gir)dr The symbol 0 in equation (5) denotes the entire material points.

(5)

139 By substituting equation (2) into equation (1) and using equations (3) and (4), one obtains (V • (C(r)a(r)Ar(r)),

c' =

'

' '' '[r

V>^(r))

'

(6)

The A;-th component of the A-th eigenfunction, V'^(r), can be approximated by using the Galerkin method in which the eigenfunction, V'^(r), is expanded by a series of admissible functions, / ^ ( r ) , each of which satisfies the accomipanying homogeneous boundary condition. This can be shown as

(r) = E'^f/'W

(7)

where c?^ s are unknown coefficients to be determined and / ^ ( r ) ' s are admissible functions chosen from a complete set of polynomial or transcendental functions that satisfy the homogeneous boundary condition. By substituting equation (7) into equation (3), multiplying by P{r) on both sides and integrating over the entire material domain, the eigenvalue problem of equation (3) in partial differential equation format is converted into an algebraic eigenvalue problem as Ad + XBd = 0

(8)

where 4 ; = /ja,«(r)/,?(r))_J-(r)rfr

(9)

and B0^ = / f{r)r{r)dr.

(10)

For 2-D, equation (8) can be written as

where Aij and B in equation (11) represent matrices themselves and df is a vector. The components of equations (9)-(10) can be evaluated using a computer algebra system. It is noted that the Green's function for an FGM can be constructed from the eigenfunctions as [3] 5fcm(r,r) = 2 ^

~

(12)

140 where gkm{i^,T^') is defined as (13) with the homogeneous boundary condition and ^(r — r') is the Dirac delta function. Using the Green's function, the displacement field in an FGM subject to the Dirichlet type boundary condition {ui prescribed) can be expressed as «m(r) = /

ft,-,,(r,r')CiMr')utir')ds'

- f (C.-,«(r')a.-,(r')Ar(r')).,ff™(r,r')dr'

(14)

where 9 0 denotes the material boundary points. The steady-state heat conduction in an FGM can be handled almost in parallel to the procedure described above. The heat conduction equation for an FGM is expressed as

iKi,{v)Tir),),i+g{r)

=0

(15)

where Kij{r) is the thermal conductivity, T(r) is the steady-state temperature distribution and g{r) is the heat generation term. The temperature field, T ( r ) , can be expressed by the corresponding eigenfunctions following the steps above. It is noted that the use of conaputer algebra system is essential in evaluating A^jJ and B^^ [4].

3

EXAMPLE AND CONCLUSIONS

In the absence of available experimental data to be compared with the presented procedure, only a numerical example for a virtual FGM is presented.

Figure 1: von Mises' stress distribution in FGM

141

Figures 1 is an example for (non-dimensionalized) thermal stresses (von Mises' stress) under steady-state heat conduction in a rectangular body defined over {(a:,y),0 < x < 1,0 < 1/ < 1} subject to the Dirichlet type homogeneous boundary condition [T = 0 and Ui — 0) on the boundary with a uniform internal heat generation [5]. The thermal conductivity, elastic modulus and thermal expansion coefficient are all assumed to vary in the form oi ko -{- kix -{• k2y where ki^s are constants. For the purpose of illustrations, the values of all the material properties are taken to be unity. The use of computer algebra system is essential to manipulate the permissible functions. The proposed method will be used as a design parameter optimization for optimum material distribution in FGMs. The result will be reported subsequently.

ACKNOWLEDGEMENT Financial support by the Texas Higher Education Coordinate Board is greatly appreciated.

REFERENCES 1. S. Nomura and A. Haji-Sheikh, J. Eng. Materials and Technology, No.UO (1988) 110. 2. S. Nomura and D.K. Choi in Integral methods in Science and Engineering^ ed. by C. Constanda, Longman Scientific & Technical, Essex, UK, (1994) 45. 3. T. Mura, Micromechanics of defects in solids, M. Nijhoff, The Hague, 1982. 4. D.K. Choi and S. Nomura, Computers k Structures, No.43 (1992) 645. 5. S. Nomura and D. M. Sheahen, The Third International Conference on Composites Engineering, New Orleans (1996) 635.

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I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

143

Effect of Gradient Microstructure on Thermal Shock Crack Extension in Metal/Ceramic Functionally Graded Materials A.Kawasaki and R.Watanabe Department of Materials Processing, Faculty of Engineering, Tohoku University Sendai 980-77, Japan The thermal shock fracture mechanism of metal/ceramic functionally graded materials was studied by burner heating test. Two series of FGMs, having different microstructure of PSZ/stainless steel and PSZ/superalloy systems, were fabricated by powder metallurgical process. The fracture toughness increased with increase in the metal phase content, owing to toughening mechanisms of thermal-strainmisfit and crack deflection. The FGMs were joined on cooling substrates and applied for burner heating test. The crack formation was always observed on the ceramic surface during cooling due to large residual tensile stresses. Dependence of thermal shock crack propagation on gradient microstructure was discussed on the basis of fracture mechanics with reference to the determination of fracture toughness and the estimation of stress intensity factor in FGMs. By comparison between the fracture toughness and stress intensity factors, the depth of the parallel cracks beneath the surface may correspond well to a location of mode II stress intensity being equal to zero. The vertical-crack's arrest is considered to be due to decrease in stress intensity factor below the fracture toughness. 1. INTRODUCTION Functionally Graded Materials (FGM)[1], have the advantage of effectively reducing the thermal stresses to be generated at the interface between the ceramic coating[2] and the metal substrate and the authors have been demonstrated the thermal stress relief function along with the increase of bonding strength between the coating and the substrate[3]. The graded coating must be designed and processed to prevent the coating from causing delamination and spalling during thermal loading in service conditions[4]. However, the design criteria has not been well established at present, because of insufficient understanding of the fracture mechanisms of the FGM coating in actual thermal environments. For appropriate design of functionally graded TBCs to meet the above requirements, therefore, the failure mechanisms of the coating must be clarified. Reviewing previous studies on FGMs for high temperature use, it has hitherto been revealed that inelastic compressive strain is generated during heating in the top surface of a mechanically constrained FGM, when the top surface of the FGM is heated above the brittle to ductile transition temperature of the coating material[4-6]. This non-linear deformation causes a large tensile stress upon subsequent cooling, resulting in the formation of surface cracks in direction perpendicular to the surface, if the stress exceeds the fracture strength of the ceramic coating. This paper describes the crack extension behavior in the functionally graded thermal barrier coating, which has been studied by burner heating test. Two series of FGM coatings, having different microstructures, were fabricated by powder metallurgical processes. The effect of microstructure on crack extension behavior will be discussed on the basis of fracture mechanics. 2. EXPERIMENTAL PROCEDURE Disk-shaped FGMs of material combination of PSZ and SS( AISI type 304 stainless steel) were fabricated through the route of pressureless sintering and hot pressing. The detailed processing is found in elsewhere[7]. The specimens have 30mm in diameter with 4mm in thickness. The composition profile and the thickness of the graded layer were varied to see thermal response. PSZ/superalloy functionally graded coatings were fabricated through the route of slurry dipping and

144 HIP sintering processes[8]. Disk-shaped green compacts of the superalloy powders, having 14mm in diameter with 5mm in thickness, were prepared by die compaction and CIP as substrates for slurry dipping. The metal and/or ceramic powders were suspended in ethanol and milled with a tumbler ball mill to get a slurry having an appropriate viscosity for dipping. The substrate was dipped in the slurry, then withdrawn and dried. After drying, the coated substrate was CIPed again to settle the intended green layer. This process was repeated with slurries of different compositions to get a graded layer. The formed compacts were HIP-sintered for densification. Fracture toughness values at different compositions were determined by a conventional vickers indentation method on uniform non-FGM specimens which were prepared under essentially the same fabricating condition as in the case of FGMs. The specimens were polished to a mirror finish on one of the faces. The test was conducted at an indent load of 196N according to the standard indentation fracture procedure, JIS R1607. Crack lengths were measured by optical microscopy to evaluate fracture toughness values. Fig.l shows a schematic illustration of a burner heating test system[6]. An FGM specimen is brazed on a copper holder to assure it in a fully constrained mechanical boundary condition for the burner heating test. The surface of the specimen was heated by combustion flame of hydrogen/oxygen gas mixture, and the bottom side of the holder was cooled by flowing water. A shutter which cuts off the flame made the rapid heating and cooling possible. The surface temperature was monitored by an emission thermometer. The absorption band of infrared rays from the burning of the mixed gas was then optically filtered. An emissivity of 0.8 was used[5]. Three thermocouples spaced 3mm apart behind the specimen permit to determine the heat flux and to estimate the bottom surface temperature. Temperature difference is defined as the difference between the top and bottom surface temperatures of the specimen. An AE sensor is mounted on the body of a cooling chamber to detect the onset of cracking by using PAC's acoustic emission analyzer. The sequence of the bumer heating test includes heating up to a desired temperature, holding for two minutes and then cooling down. After testing, specimens were cut perpendicular to the surface and polished to observe the crack morphology in the cross section of every specimen by means of SEM and optical microscopy.

2000

O PSZ/IN100FGM OPSZ/lnco718FGM D (present study)

1800 ^

1600

g

1400

(0

^

vertical crack

-4^t^

1200

g. 1000 h12 +

800

crack free

600 Cu holder

Fig. 1 Test system and sample setting configuration for burner heating test. (1) test sample,(2) torch burner,(3)cooling chamber, (4) shutter,(5) protect plate,(6) AE sensor,(7) emission pyrometer, (8) thermocouple, (9) AE apparatus, (10) monitor, (11) regulate valve, (12) coohng water.

• °

400 200 200

400

600

PSZ/SSFGM PSZ spray coating 800 1000 1200

T bottom , (K) Fig.2 Fracture mode map

145 3. RESULTS AND DISCUSSIONS 3.1. Thermal shock cracking and crack morphology Fig.2 shows a typical fracture mode map[9] obtained previously, in which the present data are also plotted. Here, a vertical crack formation region and a crack free region are conveniently separated. For specimens in the crack formation region, a number of AEs were detected after the onset of cooling, whereas there was no signal during heating. The result shows that cracks were formed during cooling. The temperature of the first crack formation was defined as the critical surface temperature. The critical temperatures of the both specimens A and B were found to be around 1300K as shown in Fig.2 and to correspond well to the brittle-ductile transition temperature of PSZ at which the deformation behavior of PSZ changes drastically [11]. The damage formation was found to be very similar to the case of PSZ/ stainless steel FGMs[6]. Fig.3 and Fig.4 show typical damages having observed in the FGMs after burner heating test. The optical observation is carried out on the top surfaces as well as in the cross-sections of them. Fig.4 reveals that the cracks in specimen A are generated vertically in the top surface layer of PSZ. They deflect and propagate in the boundary between the PSZ and PSZ-25vol%IN100 layer in the direction parallel to the graded plane. The deflection of crack propagation causes a surface segment of PSZ to be spalled out by link-up of the cracks together. In contrast, in specimen B, the cracks generated vertically propagated through the whole graded layer without deflection, reaching the boundary between the graded coating layer and the substrate. In this case, no spalling will occur. 3.2. Fracture toughness of FGMs The results of indentation fracture test for PSZ/stainless steel FGMs is found in reference [10]. As shown in Table 1, the fracture toughness of PSZ/25vol%IN100 is about 5.5MPam^^ which is slightly larger than that of monolithic PSZ; however, PSZ/50vol%Inco718 shows a lower fracture toughness value of 2.7MPam^^^. In PSZ/INIOO system, the fracture toughness increases with increase in the metal phase content, while in PSZ/Inco718 system, there is no improvement in fracture toughness; its values are considerably lower than those of PSZ/INIOO system. From detailed observation of crack path in Fig.4, it can be seen that the crack propagates in the PSZ matrix or along interfaces between the PSZ matrix and the metal particle. This indicates that the crack is deflected and/or arrested by the metal particles. For this reason, in specimen A having finely mixed

100>i m

Fig.3 Cross section of PSZ/stainless steel FGM after burner heating test. FGM compositional profile variation: (a) 2mm, concave profile (n=3), (b) 2mm, linear profile (n=l), (c) 2mm, convex profile (n=l/3)

146

Specimen A: PSZ/IN100-FGM

Specimen B: PSZ/lnco718-FGM

25%IN100 50%IN10C

Fig.4 Top views and cross sections of PSZ/superalloy FGMs after burner heating test. cvj

1.0

o -

'••

specimen A

io.8 Table 1 Fracture toughness determined by IF method Materials PSZ Inconel718 PSZ/25vo/%IN100 PSZ/50vo/%IN100 PSZy50vo/%Inconel718

MPaml/2 5. 77. 5.4 11.4 2.7

vertical crack

(d/L)c/

K„=0/

/ \

'(/) SO.6 II c
2 0.4 V) -o

1 (PSZ/IN100FGM)

\

^y

1\

^\^

specimen B (PSZ/INC0718FGM)

(D

:y 0.2 (0 E o

Z 0.0

r 1

0

^ metal substrate

FGM layer .

1

.

1 — !

J

.

I

.

I

2 4 6 8 10 Relative crack depth, d/L Fig.5 Schematic diagram of expected R-curves with comparison of trends in nondimensional Kj for vertical and parallel crack.

147 microstructure, the fracture toughness increases with increase in metal phase content because of the high possibility of crack arrest. On the contrary, in specimen B having rather coarse microstructure, a crack propagates in the PSZ matrix with comparative ease and connectivity of metal particle is lower, resulting in the less possibility of arresting cracks. Thus, in this case, less improvement in fracture toughness results, in spite that the metal phase content increases. 3.3. Crack extension behavior in FGM coatings The formation and extension of surface cracks are closely associated with the thermal stress field in functionally graded coatings during thermal loading and cooling. It has been shown that stresses in the center part of the top surface during heating are bi-axially compressive and their values decrease in inverse proportion to the radial distance towards the periphery of the specimen. It has also been shown that the compressive stress, taking maximum at the top surface, decreases towards the bottom side along the center axis[5]. Thus, the axial stress causing delamination is rather small[5]. Typical distributions of transient thermal stress in the radial direction at the top surface of a PSZ/SS FGM during heating and cooling[6] shows that a large compressive stress is generated at 0.1 sec, while after 300sec the stress is relaxed in the center part. This is because the surface temperature reaches the brittle-to-ductile transition temperature and non-linear deformation of PSZ takes place. During cooling, the resulting inelastic strain causes the stress to turn into a tensile one drastically. It has been found that the large residual tensile stress after cooling is limited in a small depth beneath the surface. The fact that the temperature decreases sharply towards inside is associated with the non-linear deformation which is limited only in the shallow layer. The mechanism of the vertical crack formation has been elucidated as the following sequence[6]. During heating, the top surface of an FGM is in a large bi-axial compressive stress state. The stress causes non-linear deformation when the top surface is heated above the brittle-to-ductile fracture transition temperature of the ceramics. During cooling, the stress state converts into a tensile one, which is large enough to exceed the fracture strength level of zirconia, and hence causes the vertical crack. The surface region of the FGM specimen after cracking will be analogous to a plate with a crack in which residual tensile stress is acting in the surface layer. Recently, it has been reported that cracks tend to branch and deflect into a trajectory with mode II stress intensity , K,,, being equal to zero[12]. This behavior is understood to be noting that the shear associated K,, component of the crack tip field continually deflects cracks until the K, ,=0 condition is achieved[12][13]. The vertical crack propagation behavior will be discussed on the basis of the K, ,=0 criterion associated with the estimation of the stress intensity factor for the cracks perpendicular and parallel to the specimen surface[14][15]. Fig.5 shows the stress intensity factor for the vertical crack as a function of crack depth, where the fracture toughness values for specimen A and B are also plotted. The stress intensity factors are normalized by OQV^ , where GQ is the mean stress in the surface layer whose depth is L. In PSZ/Inco718 system, the fracture toughness does not increase with increase in metal phase content, and it is always lower than the stress intensity throughout the extension of the crack in the FGM layer. Thus, in this case, the crack will not be arrested until it reaches at the interface between the graded layer and the substrate which has sufficient fracture toughness to arrest the crack. This feature corresponds well to the actual crack propagating path as shown in Fig.4. In PSZ/INIOO system, on the other hand, the fracture toughness increases with increase in metal phase content. As a result, the vertical crack appears to be arrested at the depth where the fracture toughness value becomes larger than the stress intensity of the crack. The actual crack shown in Fig.4 is prevented from propagating into the inner side and is deflected toward the direction parallel to the specimen surface. As mentioned above, a vertical crack will deflect toward parallel to the surface at the depth where mode II stress intensity is equal to zero, provided the stress intensity does not exceed the fracture toughness level up to the depth of K, ,=0. Similar results have been found in the case of PSZ/stainless steel system[10]. If cracks are deflected toward parallel direction to the surface, the cracks coalesce and results in spallation and thus in the degradation of the heat insulating property of a thermal barrier graded coating.

148 Hence, the formation of parallel cracks must be suppressed. Vertical cracks will be beneficial in reducing thermal stress and they do not change the heat resistance of the graded coating. They are rather effective for the TBC performance, if they are stable. Therefore, the application of FGM structure for TBC offers the following three functions. The first is the thermal stress relief function. The second is to give the gradual increase of fracture toughness towards inside of the structure and a sufficient bonding strength between metal and ceramics. The vertical crack tends to be arrested as it extends towards inside owing to the increase of fracture toughness together with the decrease of stress intensity. The third is to give the capability of controlling vertical crack propagation, where it is emphasized that there is a possibility of stabilizing the vertical crack by optimizing the K, ,=0 location together with the improvement of fracture toughness which will be attained with controlling graded microstructure of the FGM. The vertical cracks are rather effective to reduce the thermal stress without changing the heat resistance of the FGM. 4. CONCLUSIONS The crack initiation and the crack extension behavior in functionally graded thermal barrier coatings was studied by burner heating test for two series of FGMs, having different microstructures. The effect of microstructure on crack propagating trajectory has been discussed on the basis of fracture mechanics. The fracture toughness is found to be dependent strongly on the microstructure following from the mixing conditions and the particle size of the raw material powders. In PSZ/INIOO FGMs, the fracture toughness increased with increase in the metal phase content, while in PSZ/Inco718 FGMs the fracture toughness was fairly lower than that of PSZ/INIOO FGMs, due to roughly dispersed metal phase in the PSZ matrix. By comparison between the fracture toughness and mode I stress intensity factor, initiated vertical cracks in PSZ/Inco718 FGMs are considered to extend into the interface of FGM/substrate without deflection. Although vertical cracks in PSZ/INIOO FGMs tend to be arrested in the FGM coating with the extension of the cracks into the graded layer, they deflect toward the direction parallel to the surface. The depth of the parallel cracks beneath the surface may correspond to a location of mode II stress intensity being equal to zero. Similar results have been found in the case of PSZ/stainless steel system.

REFERENCES 1. A. Kawasaki and R. Watanabe: Ceramics Intemational, 23(1997), 73. 2. N.Cherradi, A.Kawasaki and M.Gasik: Composites Engineering, 4(1994), 883. 3. A.Kawasaki and R.Watanabe: J. Japan Inst. Metals, 51(1987), 525. 4. K. Kokini, Y.R. Takeuchi and B.D.Choules: Surface and Coatings Technology, 82(1996), 77. 5. A.Kawasaki, A.Hibino and R.Watanabe: J. Japan Inst. Metals, 56(1992), 1450. 6. A.Kawasaki and R.Watanabe: Proc. 3rd Int. Symp. on Structural and Functional Gradient Mat., Ed. by B.Ilschner and N.Cherradi, (1994), 398. 7. A.Kawasaki and R.Watanabe: Ceramic Transaction, [Functionally Gradient Materials], J.B.Holt, M.Koizumi, T.Hirai and Z.A.Munir, Eds., Amer.Ceram.Soc, Westville, Ohio, 34(1993), 157-164. 8. H.Yamaoka, M.Yuki, K.Tahara, T.Irisawa, A.Kawasaki and R.Watanabe: Ceramic Transactions[Functionally Gradient Materials], J.B.Holt et al., Eds, Amer. Ceram. Soc, 34(1993), 165. 9. A. Kawasaki et al.: Materials Transactions, JIM, 37(1996), 788. 10. A. Kawasaki, C.H. Yeh and R. Watanabe: J. Japan Soc. Powder and Powder Metal., 43(1996), 295. 11. W.Jiang, J-F. Li, A.Kawasaki and R.Watanabe: J. Japan Inst. Metals, 59(1995), 1055. 12. M.D.Drory, M.D.Thouless and A.G.Evans: Acta Metall., 36(1988), 2019-2028. 13. M.D.Thouless, A.G.Evans, M.F.Ashby and J.W.Hutchinson: Acta Metall., 35(1987), 1333-1341. 14. A.Kawasaki and R.Watanabe: Proc. Intern. Workshop on Thermal Shock and Thermal Fatigue Behavior of Advanced Ceramic, (1992), Schloss Ringberg, G.Petzow & G.A.Schneider, Eds., Kulver Academic Publishers, Netherlands, (1993), 509-520. 15. A.G.Evans and E.A.Charles: J.Am. Ceram. Soc, 60(1977), 22.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

149

Thermal Fracture Mechanisms in Functionally Graded Coatings K. Kokini, Y.R. Takeuchi and B.D. Choules School of Mechanical Engineering, Purdue University, West Lafayette, IN 47907-1288, U.S.A. The mechanism of multiple crack formation at the surface of functionally graded thermal barrier coatings subjected to transient heating and cooling loads is studied. It is shown that the spacing of cracks is significantly affected by the temperature gradients in the coating and the temperature distribution on the surface. 1. INTRODUCTION The thermal fracture process of a multi-layered material subjected to a transient and nonuniform thermal loads is of significant importance to the development of material systems such as graded ceramic thermal barrier coatings whose purpose is to provide thermal protection for a metal substrate. These are critical to the next generation of high temperature machines such as diesel engines, gas turbines and aircraft engines [1-2]. However, the material discontinuities and applied thermal loads cause these coatings to fail after a finite amount of time. An understanding of the thermal fracture process is necessary in order to develop a criterion which can be used to design more durable coatings. Failure of FGM coatings has traditionally been defined as the spalling of the coating. However, the debonding process encompasses two phases; the initiation of a crack and the continued propagation until spalling results. The total life of the coating can be extended by increasing the time it takes to initiate a crack, or by reducing the rate of crack propagation. The objective of the previous work conducted by the present investigators was to extend the life of zirconia coatings by delaying crack initiation. Time dependent effects in the zirconia, at high temperatures, were shown to cause surface stresses which become tensile upon cooling[3-4]. Thus, it was shown that a low creep rate and a low thermal coefficient of expansion for the coating would prevent the stresses from becoming tensile. A material such as mullite, which satisfied these properties showed a significantly reduced crack initiation behavior[5]. Similarly, mullite coated pistons were also shown to result in a ten fold increase in the useful life of the coating in diesel engine tests [6]. However, under the proper thermal boundary conditions the mullite coatings do eventually crack along the surface. In general, coatings may fracture in different ways depending on the material behavior and different boundary conditions. Figure 1 is a schematic diagram of a surface crack which extends from the coating surface. As described above, one possible mechanism of surface crack initiation in zirconia coatings was determined to be the tensile stresses during cooling which resulted from stress relaxation during heating. Using the same criterion, it was shown that a thick FGM system could be designed which is as resistant to surface fracture as a thin single zirconia layer[7-8]. Other studies considered the continued surface crack propagation and showed that the crack growth could be retarded by the design of a functionally graded

150 coating[9-10]. The crack may continue to propagate after the tip touches the interface (Figure 1), where crack kinking may result. Another mode of fracture involves edge failures (Figure 1). These initiate due to stress singularities which are located at the free edges. Thus, under uniform loading conditions, interfacial fracture would most likely initiate at the free edges first and propagate toward the center [11]. By contrast, a different boundary condition, such as a concentrated heat flux may result in an interfacial crack away from the specimen edge (Figure 1). The propagation of an interface crack can cause the coating layer to eventually buckle under the compressive loading due to heating (Figure 1). Thus, the fracture behavior of a coating depends on the boundary conditions and the failure mechanisms. While these different modes of cracking are being investigated by the present authors, this paper presents the results of a study which was conducted to investigate the multiple surface cracking due to the application of transient thermal loads. 2. EXPERIMENTS The experimental thermal load consisted of a line heat flux focused along the center of the specimen. The specimen geometry was a multi-layered beam consisting of a mullite coating and intermediate layers which were composed of mixtures of mullite and CoCrAlY (the bond coat material). The resulting fracture mode was multiple surface cracking formed in the region near the thermal load. The coatings were plasma sprayed^ onto a 6 in. x 6 in. x 0.5 in. (15.24 cm x 15.24 cm x 1.27 cm) thick steel substrate. The manufacturing process was simulated to take place at a stress free temperature of approximately 417K, resulting in a residual stress in the specimen upon cooling. The coating consisted of a ceramic layer, two intermediate layers, and a bond coat. The intermediate material layers were composed of a mixture of bond coat and ceramic. The layers are indicated by the symbol m/n, where m is the percent weight of the ceramic and n represents the percent weight of the bond coat material (CoCrAlY). For the case of specimen used used in the present experiments, two layers with respective compositions of 80/20 and 40/60 were used. The configuration of one half of the specimen used is shown in Figure 2. The specimen dimensions were 1/4 in (6.35 mm) width and 1.5 in (38.1 mm) length with rounded comers which were chosen to prevent any comer edge failures. One advantage of the beam shape was that the depth of the surface cracks and the presence of the interface cracks could be visually observed along the sides. To increase the absorption rate of radiation during heating, the ceramic surface was coated with a thin layer of black paint. Each specimen underwent inspection before the experiment using a Bausch and Laumb microscope under 70x magnification. The thermal loads were applied to the coated specimen by first heating the surface for a certain duration followed by cooling. During the heating phase, a two dimensional 1. The coatings were plasma sprayed by Engineered Coatings, Inc., in Connecticut and provided to Purdue University by Cummins Engine Company, Columbus, IN, USA.

151 concentrated heat flux was applied to the coating surface, while the substrate was cooled on the bottom. During the cooling phase, the heat flux was turned off and an airjet was activated to cool the specimen back down to room temperature. A schematic of the experimental set up is shown in Figure 3. The concentrated heat flux was generated by two high powered infrared lamps with polished half-elliptical tube reflectors. This produced a maximum heat flux of 3.35 x 10^ W/m^ along a line at the center of the specimen. The lamps were controlled by a 5 volt control signal which was provided by a computer. The amount of heat flux generated by the lamps varied linearly with the magnitude of the control signal. The amount of heating and the duration of the thermal load were also controlled by the computer. The substrate was cooled by either contacting a water cooled copper plate (1300 W/(m^°C)) or by allowing still air was to be the cooling agent (9 W/(m^°C)). At the end of the heating cycle, the heat flux was turned off and the entire specimen was cooled back to room temperature via an air jet, where presurized air was forced downward on the specimen surface. The significantly larger heat transfer coefficient, in this case, resulted in a shorter time span for the specimen to cool back down to room temperature and a higher transient surface stress of the coating. A heat transfer calculation gave the magnitude of the convection coefficient to be 1300 W/m°C across the top surface of the coating. The sides of the specimen were insulated in order to prevent radiation heating of the sides of the specimen. This provided a temperature distribution close to the two dimensional model used in the analysis. Temperature measurements were made using a type K thermocouple which were attached along the ceramic surface and on the substrate. The temperatures measured over the duration of the experiment were used to calculate the thermal boundary conditions and the resulting stresses. 3. RESULTS A typical temperature distribution at the center of the specimen during heating and cooling is shown in Figure 4. It can be noted that the surface reaches a maximum temperature of 780°C while the bottom of the substrate remains at 120^C. It can also be noted that the cooling due to the airjet is relatively rapid, especially on the surface. The transient thermal stress without any crack present, at the center of the surface of the specimen corresponding to the transient temperature distribution is shown in Figure 5. Initially the surface is in residual compression. As the heating initiates, the compression increases, peaks and decreases to come to some steady-state value. The subsequent transient cooling causes a transient tensile stress on the surface. The experiments show that this loading causes first one crack to initiate at or near the center of the top layer. The location of this first crack somewhat depends on the location of the thermocouple which was placed at the surface to measure the temperature. Subsequently, the cyclic application of the thermal load results in the formation of a second crack followed by a third. This is schematically shown in Figure 6. The locations of these cracks is dependent on the stress distribution at the surface of the specimens. The distribution of the in-plane stresses on the surface during cooling is shown in Figure 7. In this figure, a typical location of the first crack is shown with the corresponding surface stress distribution. The location of the maximum stress indicates where the second crack is

152 expected. The stress distribution corresponding to the case when the second crack has formed is also shown. Again, the location of the maximum surface stress is where the third crack is expected to occur. Another interesting and important observation is that the magnitude of the maximum stress decreases as an additional crack is formed on the surface. This is to be expected since the presence of the cracks decreases the overall stiffness of the top coating layer. A summary of the experimental locations X2 of the second crack as a function of the first crack xj, for all the specimens tested is shown in Figure 8. Here, X2 is defined as X2/a(Tsurface~To)t where X2 is the actual location of the second crack, a is the thermal expansion coefficient of mullite, Tsurface is the temperature at the center of the surface of the coating immediately before cooling, TQ is the stress free temperature and t is the thickness of the substrate. It can be noted that the prediction of the location of the second crack agrees reasonably well with the experimental results. Another important result is that the location of the second crack varies with the temperature gradient across the entire specimen. 4. EFFECT OF SURFACE TEMPERATURE DISTRIBUTION It is important to note that these results were obtained with the experimental thermal boundary conditions. However, depending on the application, the temperature distribution at the surface of the coating is expected to be different. In order to study the effect of such differences on the formation of multiple cracks, the same coating configuration was subjected to the temperature distributions shown in Figure 9, before being subjected to the same convective cooling. The case of n = 15 results in a relatively flat temperature distribution, while at the other extreme, n = -15 causes a narrow peak near the center of the specimen. The corresponding non-dimensional stress distributions during cooling at the surface are shown in Figure 10. It can be noted that both the magnitude and the location of the maximum stress vary with the temperature distribution. The flattest temperature (n = 15) causes the largest stress with the location of the peak far from the center. The magnitude of the stress decreases as the temperature becomes more concentrated (n = 15). The locations of the maximum stresses and thus of the expected second cracks, for the different cases, are shown in Figure 11. It is clear that the smallest value of X2 is for the case of a linear temperature distribution (n = 0). 5. DISCUSSION AND CONCLUSIONS The experimental results and the analysis provided in this paper provide a framework for predicting the surface cracking behavior of multilayered coatings. It is shown that the successive formation of surface cracks reduces the surface stresses, thus making the initiation of the next crack more difficult. It is also shown that the temperature distribution along the surface and the temperature gradient within the coating system play a significant role in determining the spacing of multiple cracks. In addition, the experiments showed that none of the surface cracks propagated beyond the first layer. Therefore, the layering of the system serves the purpose of arresting the surface cracks. Consequently, it is clear that the grading and architectural design of the coating plays a significant role in the surface cracking process.

153 Acknowledgements: The support of the National Science Foundation and Cummins Engine Company are gratefully acknowledged. REFERENCES 1. W. Brindley (ed.), Workshop on Thermal Barrier Coatings, NASA Conference Publication 3312 (1995). 2. K. Kokini (ed.), Ceramic Coatings, ASME MD-Vol.44, ASME, N.Y. (1993). 3. Y.R. Takeuchi and K. Kokini, ASME Transactions, Journal of Engineering for Gas Turbines and Power, 116 (1994) 266. 4. K. Kokini and Y.R. Takeuchi, Journal of Materials Science and Engineering A, 188 (1994)317. 5. K. Kokini, B.D. Choules and Y.R. Takeuchi, Journal of Surface and Coating Technology, 82 (1996) 77. 6. Y.R. Takeuchi, K. Kokini and T.M. Yonushonis, Proceedings of the Workshop on Coatings for Advanced Engines, U.S. Department of Energy, (1992) 31. 7. K. Kokini and B.D. Choules, Composites Engineering, 5 (1995) 865. 8. B.D. Choules and K. Kokini, ASME Transactions, Journal of Engineering Materials and Technology, 118 (1996) in press. 9. G. Bao and L. Wang, International Journal of Solids and Structures, 32 (1995) 2853. IJ. F. Erdogan and M. Ozturk, AFOSR Report, 1994. 11. M. Case and K. Kokini, Ceramic Coatings, ASME MD-Vol.44 (1993) 149. 0.020-

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I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

155

Fabrication of AINAV Functionally Graded Materials K.Sogabe , M. Tanaka , T. Miura, M. Tobioka Itami Research Laboratories , Sumitomo Electric Industries Ltd. 1-1-1 Koya-kita, Itami, Hyogo, 664 Japan ABSTRUCT A highly radiative material has been developed comprising sintered aluminum nitride (AIN) in which tungsten (W) particles are dispersed in graded manner. Functionally graded materials (FGM) of tungsten and aluminum nitride was fabricated by pressureless sintering. The green compact comprised of different composition

layers (W:0.0-30.0wt%)

was

prepared and sintered in flowing nitrogen. Some of the physical properties of the AIN W FGM were discussed. 1. INTRODUCTION We have succeeded in developing Structurally graded AINAV (give an outline of AIN/W FGM). AlN/W FGM has a composition that changes in stages from the surface of the material to opposite side. To form graded Structure, AIN/W based materials have a high emissivity and a high thermal conductivity as opposed to conventional materials with uniform composition or microstructure. This concept of AIN/W FGM has the potential for application in a wide range of fields including optics, electronics and energy conversion. [1] 2. CONCEPT As shown Fig. 1, high-emissivity aluminum nitride materials use aluminum nitride as the matrix for transmittance and metallic tungsten particles as emissivity particles and have a structure in which the dispersion of the tungsten particles is changed in stages. When the graded structure shown in Fig. 1 is formed, emissivity is improved by the dispersion of particles in the matrix material. By employing a graded structure in which the dispersion of metallic tungsten particles is greatest near the heat source and is gradually

156 decreased, we can expect an increase in the amount of heat energy emitted over that in composite materials in which the particles are evenly dispersed throughout the material. 3. EXPERIMENT PROCEDURES Fig. 2 shows the experimental procedure to prepare AIN/W FGM. Aluminum nitride powder, yttrium oxide powder as a sintering aide and tungsten powder were used as the starting materials. Several mixtures of raw materials were prepared. Then, Green sheet , whose compositions were gradually changed, were formed by dip coating method. Green sheet were fired at 1800 C in nitrogen gas atmosphere (1 atm). In order to evaluate some properties , specimens were prepared in these processes. 4. RESULTS AND DISCUSSIONS 4.1. Evaluation of AINAV FGM Based on sturdy composite materials comprising metallic tungsten and aluminum nitride, we fabricated a AlN/W FGM approximately 0.5 mm thick in which the amount of metallic tungsten added was changed in stages by a normal sintering method in a nitrogen gas atmosphere ( at 1 atm ). We evaluated the properties of AIN/W FGM by perfonning the following evaluations : (1) (2) (3) (4) (5)

Identification of the composition by X-ray diffraction Analysis of the microstructure by transmission electron microscope Observation of microstructure by scanning electron microscope Analysis of the W distribution and dispersion condition by image analysis method Evaluation of radiant energy by FT-IR Spectrometer

4.2. Analysis of Composition by XRD We cut an AIN/W FGM sample (W content: from 0 wt% to 30wt%) and performed X-ray diffraction on the cut surface using the Ka waves of Cu at 1 mm intervals. The results are shown in Fig. 3. As a result of identifying the compound phases formed in the AIN/W FGM sintered body from the positions of the peaks shown in the X-ray diffraction patterns, we confirmed the existence of aluminum nitride and metallic tungsten. Fig. 3 shows the results of X-ray diffraction at all measurement points. The measurement points indicate the presence of

157 metallic tungsten and aluminum nitride. It was also shown that metallic tungsten increased in stages in one material. From the results of X-ray diffraction, we confirmed that the tungsten component in the AIN/W FGM existed in a metallic tungsten state.

4.3. Analysis of Microstructure by Transmission Electron Microscope From the results of observations by transmission electron microscope , we confirmed that the AlN/W FGM has a microstructure in which grains less than Ijiim in size are dispersed inside the matrix grains and grain boundaries. Further, from the resuhs of elemental analysis by EDX, tungsten was confirmed from the dispersed particles, aluminum was confirmed from the matrix grains and yttrium, which is a sintering assist, was confirmed from the triple point of the grain boundar\'. Evaluation of the microstructure of AIN/W FGM by transmission electron microscope showed that it was a nano-composite material in which metallic tungsten particles were dispersed inside the aluminum nitride crystal grains. 4.4. Observation of Microstructure by Scanning Electron Microscope After cutting a specimen of AlN/W FGM in the direction of the gradient, we polished the surface to a mirror finish and observed the microstructure by Scanning Electron Microscope. Photo 1 shows the results of backscattered electron image of the microstructure . The black distributed areas are metallic tungsten particles. The dispersed particles are also distributed inside the aluminum nitride crystal grain and grain boundary. The results of the image analysis of backscattered electron images were shown in Fig. 4 and Fig. 5. The number of dispersed particles was plotted against W content in Fig. 4. The number of dispersed particles increased with an increase in W content up to I0wt%. above which it was kept constant. From Fig. 5, The mean value of dispersed particles diameter increased abruptly up to about 10wt%, above which it increased gradually. There results suggests that agglomeration mechanism were operating in W content ranging from 10wt% to 30wt%. 4.5. Evaluation of radiant energy by FT-IR Spectrometer The radiant energy of several specimens were evaluated by the method shown in Fig. 6. Measured specimens were summarized in table 1. The specimen was attached the holder of heater and kept at 573K. An emission spectrum was measured with a FT-IR spectrometer

158 from 2fam to SOjiim. The spectral radiant flux of several samples were obtained from each emission spectrums by comparing with emission spectrum of graphite as a gray body which was a nonselective radiator( that is to say gray body). The results of AINAV FGM and graphite were shown in Fig. 7. By comparing with emission spectrum of graphite. The spectral radiant flux of AIN/W FGM was calculated. The results of spectral radiant flux were shown Fig. 8. The spectral radiant flux of AINAV FGM was higher than that of non-FGM ( AINAV composite ) over the whole wave length range measured. Radiant flux was obtained by integrating over the wave length of spectral radiant flux. The radiant flux of AlN/W FGM was approximately twice as that of AIN. 5. CONCLUSION By using a technique to form aluminum nitride and metallic tungsten into a graded structure, we were able to broaden the knowledge of high-radiant flux aluminum nitride materials. The development of new materials with a graded structure is gaining increasing importance as a means of achieving multiple functions. The high-radiant energy materials fabricated utilizing the technology of this research can also be applied to transmissive materials other than aluminum nitride. 6. ACKNOWLEDGMENT From the standpoint of resource conservation and reduced energy consumption, highemissivity material from various types of power sources and heat sources. This research was carried out as part of the research on technology of forming radiative structures by controlling the graded structure supported by the Special Coordination Funds of the Science and Technology Agency of the Japanese Government.

REFERENCES [1] The Society of Non-traditional Technology, Development of Energy Conversion Materials through Formation of Gradient Structures, (1994).

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 1997 Elsevier Science B.V.

159

Graded Casting for Producing Smoothly Varying Gradients in Materials Basil R. Marple and Stephane Tuffe National Research Council Canada, Industrial Materials Institute 75 de Mortagne Blvd., Boucherville, Quebec, Canada J4B 6Y4

ABSTRACT The recent interest in tailoring materials having various types of gradients has stimulated activity in developing processing routes for producing these structures. Graded casting offers the possibility of manufacturing relatively complex parts having continuous or stepwise structural or compositional gradients. The aim of this paper is to give a brief overview of the graded casting technique. Details are provided on the process and control strategy used to produce graded materials by this method. The importance of certain aspects relating to slip flow, mold design and mixing are discussed. Examples of some of the gradients that can be produced and possible approaches for producing them are presented.

1. INTRODUCTION The concept of functionally graded materials, first proposed several years ago, introduced a class of highly engineered structures tailored to enhance specific properties through controlled compositional changes within a body. Initially, a typical graded structure was one consisting of a metal and a ceramic in which the ratio between the two constituents varied in a gradual fashion through the thickness of a part. Graded materials now include those having a gradient in composition (chemical content, partitioning among various crystalline structures and/or amorphous phases) or a gradient in some aspect of the physical structure (grain size, grain orientation, porosity content). While the potential benefits of using materials engineered in such a fashion have been noted, models to predict their properties and behavior proposed, and possible applications identified, there remains a need to develop processing approaches for manufacturing components using this design strategy. As is the case with many processing technologies developed to produce a particular class of materials, the various production processes available for manufacturing graded structures tend to be complementary. Certain processes are suitable for small and/or simple shapes, while others can be used to produce large and/or complex shapes. It can be concluded, however, that many of the processing approaches previously proposed or in use are limited in the kinds of materials that can be processed, the complexity or size of the parts that can be fabricated, or the types of gradients that can be introduced. For example, with some processes only components having a stepwise change in composition through the part can be produced. While such a profile is sometimes preferred, there are apphcations for which

160 a continuous, smoothly varying transition in composition would result in enhanced performance. It was to address the need for a process to produce relatively complex-shaped components having a continuous gradient in composition that the graded casting process was conceived and developed [1].

2. GRADED CASTING PROCESS The graded casting process is based on the slip casting technique used for shaping ceramics. In the traditional approach to slip casting, a cavity in a plaster of paris mold is filled with a slurry containing particles of the material from which a part is to be produced. Through capillary force, the hquid is absorbed by the mold, leaving the ceramic particles, which are generally larger than the capillary channels, behind on the walls of the mold cavity. The casting kinetics normally follow a parabolic rate law, L = Kt^^'^, where L is the thickness of the cast layer, K is the casting parameter and / is time. Periodic additions of slip to replace the absorbed liquid may be required in order to ensure that the mold remains full. For hollow core casting, the excess slip is drained from the mold when the desired wall thickness has been reached. The green body is later removed from the mold, following a drying period. Layered structures, having a stepwise change in composition, have also been produced using the slip casting technique [2]. The graded casting approach introduced the concept of continuous flow of slip through the mold during the casting step. A schematic of the process is presented in Figure 1, and various other configurations have also been proposed [3]. The key elements of the apparatus are a porous mold having a cavity in the form of the part to be produced, reservoirs to contain the suspensions of the materials to be cast, flow pumps/meters to feed the slip through the mold, and a computer to control the flow from the various reservoirs. Mixing is enhanced through the use of in-line, stationary, spiral-shaped, plastic mixers to create turbulence during flow. The slip flowing from the mold can be collected in one of the reservoirs where it can be stored for later use or reintroduced into the slip stream. If recirculation is used, some means should be employed for stirring the slip in the collector reservoir.

••Cast layer H-Drain

Figure 1. Schematic of the process for producing graded materials by slip casting.

161 Using this apparatus, the composition of the slip in the mold can be changed in a preprogrammed fashion by changing the relative flow rates from the various reservoirs. The composition at any depth in the cast layer will be determined by the composition of the slip in the mold at the moment at which that layer is being deposited. Therefore, if the appropriate model and data exist for predicting the casting kinetics for the slips being employed, the process can be controlled to produce a selected profile.

3. PROCESS CONTROL A detailed description of the empirical model developed for producing materials having controlled gradients using the graded casting approach has been presented elsewhere [4]. The equations will not be reproduced in this paper; however, the process to produce parts having a given green thickness and a specific concentration profile can be summarized as follows. For any given system, the traditional constant composition slip casting technique is used to determine the relationship between the casting parameter, K, and slip composition, for the range of compositions to be produced in the part. This relationship, representing the change in K with composition, can be combined with the equation representing the composition profile through the part (change in composition as a function of wall thickness) to yield an expression showing how the casting parameter changes when casting the selected profile. Because the change in AT as a function of thickness is now known, the time to cast any element of thickness and the total time to cast the complete part can be calculated. The equation relating the thickness of the cast layer and time can be detennined for the graded material by curve fitting a graph of the thickness as a function of time. Combining this expression, relating thickness and time, with the equation for the composition profile, relating the thickness and composition, provides a relationship between the composition of the cast layer and time. Finally, this last equation can be used to determine the relative flow rates required from the reservoirs as a funcfion of time to produce the desired profile.

4. DESIGN CONSIDERATIONS There are several issues to be addressed when designing the apparatus and procedure for producing components by the graded casting approach. These play a key role in the successful application of the process for producing the desired part geometry having the required composition profile. 4.1 Recirculation The quantity of slip required to produce a given part will depend on both its size and the composition profile through the thickness. For small parts, the slip exiting the mold can be collected and stored for use in future production runs. For large parts, other possibilities can be explored. Figure 1 indicates that recirculation of the slip can be used to reduce the total quantity of slip required to produce a given geometry. It should be recognized that, in most cases, the slip entering the mold at any moment will be only slightly different in composition from that exiting the mold. Therefore, in theory, it would be possible to directly recirculate the slip while making minor adjustments in the slip composition by adding fresh slip from one of the reservoirs. However, because the flow rate of slip through the mold is nonnally faster

162 than the rate of absorption through the mold walls, excess slip will be produced and must be stored in the system. The design of the recirculation and storage system will depend to some degree on the composition profile being produced. For a linear composition profile, produced using two feedstock slips, the slip could be recirculated through the reservoir containing the slip composition to be cast first. For more complex composition profiles (e.g. parabolic or cychc profiles), for which the starting compositions in the various reservoirs will be required at various points during the casting step, a third collector reservoir may be used (as shown in Figure 1). Whichever approach is employed, equations can be developed to predict the composition of slip in the recirculating reservoir at any point in time. This can be accompUshed by knowing 1) the initial slip composition in the reservoirs and 2) how the various parameters (the flow ratefi-omthe reservoirs, the volume absorption rate of the mold and the total free volume of the mold) change as a function of time. These equations can be incorporated into the process control model outlined in the previous section. 4.2 Mold design Mold design is an important consideration when using graded casting. As is the case with traditional casting, one-piece molds are used when producing simple shapes such as tubes. This results in a cast piece having no marks caused by the mold join lines (defects that are nonnally evident on parts produced using multi-section molds). When producing large hollow parts, a central core or mandrel of non-porous material can be inserted into the mold to fill space otherwise occupied by slip. This will serve to reduce the total amount of slip required to produce the component by decreasing the volume of the intemal cavity of the mold. Therefore, less slip is needed to fill the mold initially and, during casting, a lower volume flow rate of slip can be used to achieve a given change in the composition of the slip in the mold at any point in time. For a given system, the volume flow rate of slip that must be fed through the mold will be determined by both the total intemal volume of the mold and the composition profile being produced in the part. For steeper composition profiles and larger mold volumes, higher flow rates are required to ensure the desired profile is produced. When designing a mold to produce a part by graded casting, consideration must be given to the path the slip will follow when travelling through the mold. It is important that the mold be properly vented and that the design not result in the creation of zones of stagnant slip. In some cases, it may be possible to redesign the part to help alleviate the problem. A design having multi-exit ports may also help to ensure continuous flushing of all regions of the mold cavity. When these approaches are not possible, other measures, such as those discussed in the following section on mixing, can be considered. Special consideration must be given when using graded casting to produce parts having a cross-sectional variation along the body. It has been shown that, for typical flow rates used in graded casting, the casting rate of a given slip remains unchanged from that measured with no flow [1]. However, when casting parts for which there are large differences in cross section along the component, the linear flow rate of slip at the slip/cast layer interface in regions of small cross section may become quite high. These high flow rates may interfere with the casting process and result in a non-uniform cast layer thickness along the part. 4.3 Mixing Mixing is an important element of the graded casting process. As is shown in Figure 1, mixing of the slips during flow can be aided by creating increased turbulence using in-line

163 "static mixers" to introduce a more tortuous flow path. When slip is recirculated, it is important that the apparatus include a means to mix the slip in the recirculating reservoir. This becomes particularly important if large volumes of slip are being used. Another aspect of mixing relates to slip introduction and mold design. For small, open-ended parts, the slip can be introduced at one end and exit at the other. For large pieces, the design shown in Figure 1, in which the slip is injected at various points along the length, can be employed. This may help to reduce gradients in the longitudinal direction. For complex shapes having zones where pockets of slip could remain stagnant, the injection of slip at various points in the mold cavity could also aid mixing. Continuous rotation of the inlet tube could also be used to further enhance the blending process. It may also be useful to cast with the mold oriented in the horizontal position while being slowly rotated. Such an approach could be particularly useful if casting was performed with the mold only partly filled with slip. This would improve mixing and help to ensure that the various areas of the mold were being drained and refilled constantly with new slip. To assist in preventing the introduction of defects in the cast body, defoaming agents could be added to the slip to aid in limiting bubble fonnation.

5. SYSTEMS AND APPLICATIONS The slip casting approach can be used with any material that can be placed in suspension in a suitable liquid. Typically, slips consist of a material, in powder form, dispersed in a carrier such as water or an organic liquid. Although it is important that the slip be stable (limited settling), the continuous slip movement involved with graded casting may aid in the casting of unstable suspensions. Traditionally, slip casting has been used to manufacture ceramic parts; however, further development has led to a modified slip casting-sedimentation process for producing graded metal-ceramic composites [5]. Gradients in materials may take various forms. These include gradients in chemical composition, in the density or porosity content and in the scale of the microstructure. The graded casting process has potential for producing materials having any of these various types of gradients. 5.1 Gradients in chemical composition Parts having a smoothly varying or stepwise change in composition can be fabricated by graded casting. The resulting material may be a particulate composite or, if desired, one of the phases can be in the form of whiskers or platelets. Multi-component systems could also be produced by using additional reservoirs and flow control pumps. Of course, normal principles of materials design and engineering must be observed when considering the production of parts having two or more components. Examples of linear and parabolic composition profiles produced in alumina-zirconia particulate composite tubes are shown in Figure 2. In these cases, the outer surface is rich in alumina, having the lower thennal expansion coefficient of the two components. This provides for the possibility of introducing a surface compressive stress when cooling the part after sintering. Using the same design strategy, the process could also be used to produce other ceramic-ceramic composites such as muUite-zirconia, silicon carbide-alumina and mullitealumina. In the case of muUite-alumina, the composite could be produced by co-casting muUite and alumina in a graded fashion and sintering, or by co-casting silica and alumina followed by reaction sintering [6].

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- 0 — Linear Parabolic

0.5

LLI

i_L

2

3

I I 1 I I I I I I I I 1 I I I I I I I I I I I I

0

1

Distance from outer surface (mm)

Figure 2. Composition profiles of alumina through two sintered alumina-zirconia tubes.

Figure 3. Cross-sectional view of two sintered alumina-zirconia wave guides.

Grading the composition using graded casting enables the tailoring of components for specific appHcations. Acoustic wave guides produced using this process are shown in Figure 3. These components consist of an outer AI2O3 layer, a central region of an Al203-Zr02 composite of a fixed composition, and an interface or transition zone in which the proportion of the two constituents changes linearly from that of the outer layer to that of the central region. 5.2 Gradients in porosity The graded casting approach could also be used to produce parts having a gradient in the level of porosity. For example, foaming agents could be added through one of the reservoirs into the slip stream. This would help in introducing gas bubbles into the cast layer. A second approach for increasing the level of porosity would be to add an organic phase in a controlled manner. For example, polymer spheres could be introduced and cast together with the ceramic. By controlling the size as well as the quantity of the spheres, the size distribution of the pores in the cast body, following polymer burnout and sintering, could be controlled. Another option for grading the porosity/density would be to control the presence of species that act to either promote or hinder the densification process. For example, in some cases, certain species will affect the relative rates of surface and bulk diffusion occurring during sintering. For a graded distribution, this could affect the resulting density of the various regions in the sintered part. In other systems, phase changes, decomposition or combination reactions may create porosity and affect densification. If the various species are present in a graded fashion in the green body, the sintered part may have a similar gradient in porosity. Yet another approach with potential for controlling the porosity within a part is by modifying the state of dispersion of a slip. By slowly adding suitable agglomerating agents or changing the pH during casting, the density of the cast layer could be varied. Of course, in all of these approaches, the integrity of the final product must be considered. If the process developed to introduce the graded porosity results in part deformation or cracking during heat treatment, it may be of little value.

165 Parts produced having a gradient in porosity could be used to fabricate composites with a gradient in composition. For graded ceramic-ceramic composites, the graded porosity in a partially sintered compact could be filled with a suitable hquid precursor to produce composites by the infiltration processing route [6, 7]. In the case of metal-ceramic composites, hquid metal infiltration of the sintered ceramic piece could be employed to introduce the metal phase [8]. 5.3 Gradients in microstructure Another aspect of a material that can be graded is the scale of the microstructure (grain size, whisker length or diameter). Graded casting enables the production of components having a continuous gradient in the microstructure. Various approaches exist for controlling grain growth in materials. For example, it is known that the addition of a second phase often limits grain growth during sintering. This effect has been shown for the addition of zirconia [9] or mulhte [6, 10] to alumina. An example of an alumina-zirconia particulate composite produced by graded casting and having a gradient in grain size is shown in Figure 4. More pronounced effects are possible; the magnitude of the change depends on the system under study and the materials being used. The above example involved a significant change in the chemical composition of the base material. This may not always be desirable and other methods of controlling the evolution of the microstructure will be required. In some systems, only very small amounts of a second material are required to inhibit or promote abnormal grain growth. For example, in alumina, MgO is effective at limiting grain growth [11], while small amounts of other species can cause massive grain growth [12]. The concentration of such additives could be varied in a controlled manner through the thickness of a cast body when fonning the part by graded casting. In this way, the amount of grain growth occurring during sintering would vary through the thickness. In cases where no additives can be tolerated, the size of the powder particles being cast can be varied during casting to yield a green body with a graded structure. This approach would pennit some limited control over the range of grain sizes present within the sintered body. Of course, with any of these techniques, the effect that the change in chemical content or particle size has on sintering and the quality of the final product will have to be evaluated.

Figure 4. Micrographs of different regions in a graded Al203-Zr02 composite. The distance, d, from the outer surface and the volume fraction of zirconia, Vf^, present is noted for each zone.

166 6. CONCLUSION The graded casting approach offers promise as a process for producing materials having various types of gradients. It introduces the concept of continuous slip flow to the widely used slip casting method for producing ceramics and uses an empirical model to predict the changes in slip composition required during casting to produce the desired gradient. The technique complements other processes for producing graded structures and it enables the fabrication of relatively complex shapes having smoothly varying gradients or stepwise changes through the thickness. However, as is the case with other processing routes, normal engineering and materials design principles must be observed and the physical and chemical effects of combining different materials and grading structures must be considered.

REFERENCES 1.

B. R. Marple and J. Boulanger, "Graded Casting of Materials with Continuous Gradients," J. Am. Ceram. Soc, 11 [10] 2747-50 (1994). 2. R. Moreno, A. J. Sanchez-Herencia, and J. S. Moya, "Functionally Gradient Materials by Sequential Slip Casting: Alumina-Yttria Tetragonal Zirconia"; pp. 149-56 in Ceramic Transactions, Vol. 34, Functionally Gradient Materials. Edited by J. B. Holt, M. Koizumi, T. Hirai, and Z. A. Munir. American Ceramic Society, Westerville, OH, 1993. 3. B. R. Marple and J. Boulanger, "Slip Casting Process and Apparatus for Producing Graded Materials," U.S. Patent No. 5,498,383 (1996). 4. S. Tuffe and B. R. Marple, "Graded Casting: Process Control for Producing Tailored Profiles," J. Am. Ceram. Soc., 78 [12] 3297-303 (1995). 5. J. Chu, H. Ishibashi, K. Hayashi, H. Takebe, and K. Morinaga, "Slip Casting of Continuous Functionally Gradient Material," /. Ceram. Soc. Jpn., 101 [7] 818-20 (1993). 6. B. R. Marple and D. J. Green, "Mullite/Alumina Particulate Composites by Infiltration Processing," J. Am. Ceram. Soc, 72 [11] 2043-48 (1989). 7. S. J. Glass and D. J. Green, "Fabrication of Multiphase Particulate Ceramics by Infiltration into Powder Compacts"; pp. 784-91 in Ceramic Transactions, Vol. IB, Ceramic Powder Science II. Edited by G. L. Messing, E. R. Fuller, Jr., and H. Hausner, American Ceramic Society, Westerville, OH, 1988. 8. C. Toy and W. D. Scott, "Ceramic-Metal Composite Produced by Melt Infiltration," J. Am. Ceram. Soc, 73 [1] 97-101 (1990). 9. F. F. Lange and M. M. Hirlinger, "Hindrance of Grain Growth in AI2O3 by Zr02 Inclusions," J. Am. Ceram. Soc, 67 [3] 164-68 (1984). 10. B. R. Marple and D. J. Green, "Graded Compositions and Microstructures by Infiltration Processing," J. Mater. ScL, 28, 4637-43 (1993). 11. R. L. Coble, "Sintering Crystalline Solids. II. Experimental Test of Diffusion Models in Powder Compacts," J. Appl. Phys., 32 [5] 793-99 (1961). 12. B. R. Marple and D. J. Green, "Mullite/Alumina Particulate Composites by Infiltration Processing: II, Infiltration and Characterization," J. Am. Ceram. Soc, 73 [12] 3611-16 (1990); correction, ibid., 74 [3] 641 (1991).

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

167

Gradient components with a high melting point difference M. Joensson, U. Birth, B. Kieback Institute of Material Science, University of Technology, Dresden 1 INTRODUCTION Gradient materials offer the chance to avoid stress singularities by bonding of different materials but open also a new field of material characteristics by combination of special application relevant properties. Applications which meet very different and extreme local demands in view of thermal, electrical, magnetical, corrosion and mechanical load require the development of graded bulk materials via the combination of materials with very different properties. Such materials often differ strong in there melting points. Because this difference requires variable processing parameters, the main problem in the development of gradients is the selection of suitable technologies.

[Material

tensile strains LOW MELTING IVIATERIALS thermai expansion]

HIGH MELTING MATERIALS hardeness

thermai conductivity

density

Young^s modulus

Figure 1 Comparison of properties of low and high mehing materials

W

" Cu"" 1

Tsrc

3410

Density, g/cm3

19,30

8,96

Hardness, HB

300-400

50,00

tensile strength, MPa Elastic Moduls, GPa Thermal Conductivity, w r c m Thermal Expansion, ppm/°C Poisson

1920,00 314,00

1083

411

145

174

403

4,50

17

0,28

0,34

Table 1 Properties of tungsten and copper

In the comparison of the properties of high and low mehing materials, especially metals, tungsten and copper become conspicuous because of their distinctive differences of properties. Therefore W-Cu composites are traditional of interest and many investigations are engaged in their processing and characterisation of properties /1;2;3;4/. Applications of such composites could be find in production engineering and electrical engineering. W-Cu is a promising material for thermal management applications such as microelectronics packaging /5/.Composite Cu-W sintered materials containing <15 Vol.-% W are used as spark erosion electrodes and as contact tips in gas metal arc welding guns /I/.

168 Gradients with very different melting points of material components could be used also in spaceplanes, gas turbines and fusion reactors. Especially W-Cu is suitable for beam targets 161. The combination of this two elements should provide a serviceable material and be suitable as a model system for the investigations. Their potential for high thermal and electrical conductivity and high hardness, high temperature stability and low thermal expansions coefficient can only be full exploit, if its possible to produce a gradient over the whole concentration profile. For materials with very different melting point powder metallurgy offers practicable ways for the composite processing. Traditional powder processing's can be adapted to the special composition of gradient materials and also new methods like sedimentation or centrifugal powder forming were added as steps for a continuous gradation before consolidation. The formation of powder mixing gradients as the base of the desired change in the material characteristics renders possible the combination of such unsolutable components like W-Cu. But due to the big difference in the melting point of such material combinations the characteristic sintering temperatures differ such extreme, that the sintering temperature of the higher melting component is much higher as the melting temperature of the lower melting material. In view of their efficiency in the differed concentration areas various powder technologies should be selected and combined for the development of a sintering gradient body with high density and a continuous change of concentration from 0-100%.

2. POWDER MIXING GRADIENTS 2.1. Layer pressing The well known and often used way to produce powder mixing gradients is the pressing of layers with stepped quantitative relationship of components. Only conventional technologies and equipment's can be used. So it seems to be a practicable and by optimising of the layer numbers also a simple way. But the different densities of W and Cu and the difference of their conventional be used particle size leads to difficulties in the mixing process. Van der Waals forces between the small tungsten particles often produce aggregation during classical dry mixing processes and by wet mixing the high specific gravity of W can result in segregation III. In our investigations we use tungsten powder with a relatively high particle size (mean size llyim) and different fractions of copper powder (<20nm, <32^m, <40|im and <75^m). The concentration changes from layer to layer in steps of 10 wt-% where the Cu-content cover an area from 10 wt-% to 90 wt-%. With the choise of nearby particle sizes of the components inhomogenities in the layers can not be avoid at all, because the fact of the density difference still stay. By accepting a decreasing of the sintering activity of the tungsten powder can only be realised a - for the solid phase sintering in the copper pure area important 111 - decreasing of the copper particle distances. For a low content of W (e.g. <15 Vol.-%) this problems can avoid with high energy milling, where the W powder is dispersed uniformly in the copper matrix III. This expenditure make this processing way more complicated as it seems of the first look.

2.2. Centrifugal Powder Forming (CPF) The CPF-technology 78,9/ offers the possibility to change continuously the mixing composition of up to four powder components. Defined amount of powders were supplied to a form, open on the top and rotating horizontally, fall on a distributor which spreads them uniformly to the wall. The powder package must be fixed with a binder because only the centrifugal forces kept them in position. In our equipment the forming process is combined with a liquid binder supply

169 via a controlled spraying process. Due to the vibration supported powder transport and the computer controlled microdosed powder supply, the combination of very different powders is possible. Differences in the powder characteristics, like density, flow ability or particle sizes were considered in the chose of the supply parameters and in the control program.

Figure 2 CPF-mixing gradient of W and Cu First investigations result in a W-Cu mixing gradient, where the change of the mixing is satisfied but the distribution of the Cu particles is not so good as expected. Therefore the optimization of the processing parameters is necessary. Especially the demands on the binder and its supply are very complex. The stability of the green body have to be guaranteed without an unwanted widespread of the particle distances by the binder or the agglomeration of the powder particles. The optimization of the binder content and its viscosity is the prerequisite of its good distribution which is necessary for the high quality of the part. On the other side there are the possibility for the transport of sintering additive via the liquid binder. Nevertheless as a result of this forming process the green density of CPF-parts is lower then of pressed parts (comparable with loose filled powder). Therefore after debinding a pressure supported consolidation process is inevitable, unless a porosity of the sintering part is wanted.

2.3. Consolidation The real problem in the development of a gradient of materials with very differed melting points is not so much the powder combination but the consolidation. Sintering investigations in hydrogen below (1050°C) the Cu melting point of pressed composite samples prove the influence of the Cu particle size but also the general problem of the sintering of a mixing W-Cu gradient. According to the green density the sinter density approaches the theoretical density only on the copper riche side. The micrograph of a layer pressed gradient (Cu: <30|im) support this fact. Especially the last W riche layer contrasts with high porosity.

170 sinter density [g/ccm] — <20 Mm —<30Mm —<75Mm — < 7 5 Mm activ. — < 7 5 Mm;1C360°C theor.

100 90

% theoretical density |-»-<20Mm •»-<30Mm -«"<40Mm •*-<75Mm |-iii-<75|jm activ.

80 70

0 5

20 40 60 80 100 content of Cu [wt.-%]

60

5 10 20 30 40 50 content of Cu [wt.-%]

60

Figure 3 Density of W-Cu composites in comparison with the theory h^n^. \ :^f *•'-'. ^''^ ^^^

\'

I

' >/

400pm

Figure 4 Layer pressing W-Cu gradient After Hot pressing at 1000°c with 55 kN in ArH6,5 the gradient areas with more than 20% Cu - in a W-Cu mixing layer combination with 5,10,20,30 and 40 wt-% Cu - have a clearly higher density as in the case of the pressureless sintering, but the densification below this content of copper is unsatisfied. It is known that in the case of a good wettability - this is relevant of the system W-Cu - a liquid phase sintering can lead to higher densities of composites with a low copper content. Our results correspond with other investigation of W-Cu composites with <25 wt.% copper / 2, 3, 4 /. The decreasing of the sintering temperature of W which is possible by doping with small amount (<0,5 wt.%) of additives - known as the Agthe-Vacek-Effect - can leads to better densification of W-Cu components. The activation of sintering is influenced by the content of copper and the sintering temperature. For W-Cu system are Co and Fe the most effective activators, in contrast of pure W where Ni the preferred additive is, because this elements in the liquid copper are not so solutable as Ni and can form a highdifiRisivity interboundary layer by segregation to the W grain boundaries. 73/ The use of this effect for a gradient material means a balanced doping only of the W riche side. An oversupply of additive elements in the direction of a higher Cu concentration can cause a serious worsening in the thermal conductivity.

171 3. INFILTRATION GRADIENTS 3.1. Porosity Gradient The key for an infiltration gradient of material combinations with very different melting points is the production of a skeleton of the higher melting component with a porosity gradient. Takahasi / 6 / uses the dependence of the sintering activity of the particle size to develop a porosity gradient and consolidate layers of different powder fractions. Based of the production of a porous preform with homogenous porosity a controlled electrochemical opening of the pores is under development / 10 /. By using the CPF different processes are possible. Rather investigations /11/ offer the opportunity to develop a porosity gradient by a mixing gradient of powder fractions different in particle size and shape. By combination of stainless steel powder fractions of <5nm up to <80^m a change of porosity between 10-50% can be realised if the fine powder has a spherical shape and the coarse powder is irregular shaped. Porosity [%] 63- At\ diffemce of porosity [%] 80

gradient dimention

irregular/spherical 35 45 55 65 difference of particle size Oim]

Figure 5 Change of Porosity by combination of powder fractions of stainless steel Because of the low density of CPF-green body's by using of a binder (filler) gradient for development of a porosity gradient more than 30% porosity difference can not be expected. As an other possibility we find out that the liquid phase sintering of a W-Cu gradient with 0-20 vol.-% copper and 0,25% Co as additive resuhs in an porosity gradient fi-om 30% to 75%. After infiltration it leaves to a W-Cu gradient with 30-80 vol.-% copper and a nearly linear gradient function over the whole dimension (3 mm) of the sample.

Figure 6 Porosity gradient

172 3.2. Infiltration Prerequisite for the infiltration is a open porosity. Concentration areas below 10 vol.-% of the lower mehing component needs in this case a below 10 % porosity an open pore structure. But this is not realizable because the content of closed pores into the direction of full density increase seriously from this value of porosity on. The area and the function of the concentration gradient are fixed by the processing of the porosity. On the other hand this process requires a good wettability of the porous material with the melt of the lower melting component. Therefor the infiltration of W with Cu is easy and is used conventional to produced W-Cu composites. For other combinations the wettability can improved in the same way as known from the composite development. 4. CONCLUSIONS The different discussed powder processing's offer satisfactory solutions only in partial concentration area of the gradient of materials with very different melting points. The problems are caused in the non possibility to consolidate a mixing gradient at once to a really dense fiill body. The infiltration of a skeleton with porosity gradient can avoid mixing and consolidation problems, but this is limited by the exist of only open porosity and on the other side by the form stability of the skeleton. The gradient area with more than 80% of the lower melting component can produced via solid phase sintering of a mixing gradient. The most problematical part of the gradient is there on the higher melting material riche side. Liquid phase sintering supported by additives can leave to gradients with high density in this area. A combination of various processing steps seems to be the only solution to produce a whole concentration profile included gradient of such material combination.

REFERENCES III Kaczmar,J. 111 German, R.M. 131 Johnson, J.L.; Gennan,R.M. /4/ Dm, T.-H.; Lee,S.-W.; Joo,S.-K. 151 German, R.M.; Hens, K.F.; Johnson,J.L. 161 Takahashi,M.; Itoh,Y.; Kashiwaya,H. 111 Lee, J.-S. /8/ Delfosse,D.; Ilschner,B. 191 Joensson,M.; Kieback,B. 710/ R6del,J.; Neubrand,A. /11/ Joensson,M.;Kieback,B.

Powder Metallurgy, Vol.32,1989, No.3, 171 Metallurgical Transactions A, Vol. 24A, Aug. 1993, 1745 Metallurgical Transactions A, Vol..24A, Nov. 1993, 2369 Powder Metallurgy, Vol.37, 1994, No.4, 283 ThelntemationalJoumal of Powder Metallurgy Vol. 30, No.2, 1994, 205 Proceedings of the First Intemational Symposiimi,FGM, Sendai,1990, 129 Dissertation, Institut fur Materialkunde der Universitat Stuttgart, 1983 Acta metall.mater. 40 (1992), 2219 Workshop „Gradientenwerkstoffe", Eds. Schilz,J.;Huelsmann, DLR,K61n-Porz, Institut fiir Werkstofforschung, 1993,17-1 Verbundwerkstoffe und Werkstofiverbunde, G. Ziegler (ed.),DGM Mormationsgesellschaft Verlag, Oberursel;1996,15 Proceedings of the second Intemational Conference on Composites Engineering (ICCE/2), New Orleans, 1995, ed. by D.Hui, Intemational Community for Composites Engineering and College of Engineering, University of New Orleans, 1995 ;3 65

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

173

Fabrication of Pore-Gradient Membranes via Centrifugal Casting C.-W. H o n g , F. Miiller and P. Greil Universitat Erlangen-Niirnberg, Institut fiir WerkstofFwissenschaften (3), Martensstrasse 5, D-91058 Erlangen, Germany

Abstract Using commercial AI2O3 powders, flat membranes with a graded pore-size distribution from 40 to 250 nm across a membrane thickness of about 5 mm have been fabricated by centrifugal casting. The scanning electron micrographs and the mercury-porosimetry data for such a pore-gradient membrane have been shown. Flocculated suspensions resulted to a narrower pore-size distribution without pores larger than 200 nm. Specimens prepared by stabilized suspensions with 10vol% solid content showed also a clear pore-gradient structure. The experimental results have been correlated with the results from the DEM simulation. 1 INTRODUCTION The use of ceramic membranes in the separation process instead of organic ones offers various advantages. In addition to high temperature stability, ceramic membranes exhibit backflush capability, they are chemically and microbiologically resistant, have an excellent mechanical stability and can be used in catalytic applications [1-3]. Membrane filtration processes are usually divided into filtration (pore diameters greater than lO'^nm), microfiltration (from 10^ to lO^nm,) and ultrafiltration (UF, from 1 to 100 nm) and reverse osmosis (up to I n m ) . Considering ultrafiltration and reverse-osmosis membranes, so far only solgel techniques can provide ceramic layers with ultrafine pores [2]. Typically, this filtration layer is formed on a support (or outer) layer which has larger pore size (1 to 6 fim,) and is manufactured by slip casting [3,4]. Between the support layer and filtration layer, a discontinuous transition of material properties exists. Which, however, can be avoided by a gradient structure. It is well-known that large particles form the large pores, while small particles form the small pores. Multimodal suspensions will form greenbodies with a pore gradient if the size segregation can be ensured during the forming process. As materials for the UF membranes, oxide ceramics, like Zr02 [2,5] and Ti02 [2,6,7] are primarily used. UF membranes made of 7-AI2O3 have lost their importance because of their low permeability and chemical and thermal stability [2,8]. Presently UF membranes made of a-Al203 could not be fabricated through the sol-gel techniques, because a calcination temperature of T > 1000° C will be needed for the transformation of 7- to a-Al203 [9]. This enhances the not-desirable grain and pore growth. In this work, the UF pore-gradient flat membranes have been attempted to fabricate by using the commercially available a-Al203 powders with a continuous Submicron size distribution via centrifugal casting.

174 —1

c

1

1

1

1

1

random spheres

CM

> 'M

1

ri/^ 0.0

0.2

0.4

0.6

insphere size [d/D] Figure 1: The insphere size distribution for randomly packed spheres. 2 PARTICLE PACKING STRUCTURES 2.1 Pore size distribution in m o n o m o d a l packing Besides the packing density, the pore structures belong to the important properties of particulate structures. They constitute an useful aspect of packed particles, since they control properties such as filtration, permeability, fluid trapping, etc.. Packing structures of powders with continuous size distribution are very complex. The pores in ordered packings of monosized spheres would seem to be the simplest case to study. The insphere sizes for four ordered packings are shown as follow [10]:

coordination number 12 relative insphere 0.1547 diameter

10 0.2647

8 0.4142 0.2247 0.1547

6 0.4142 0.3333

The multiple entries represent the variations in ways of calculating the insphere size depending on the pore placement. The range of normalized insphere diameters is from 0.1547 for a packing coordination of 12, up to 0.4142 for packing coordinations of 6 and 8. Alternatively, in a random packing the insphere size is distributed. Fig. 1 shows the frequency distribution of the insphere sizes from a random packing of monosized spheres [11]. There is a large peak at the close-packed insphere size of 0.1547; however, unlike the ordered packing, the random structure has a wide distribution with some relatively large openings. 2.2 Pore-gradient formation via centrifugal casting Hong [12] has shown in his DEM simulation results that pore-gradient membranes can be formed by using stabilized AI2O3 suspensions with a suspension height of 70 mm and a solid concentration of < 10% under a centrifugal acceleration of greater than 40 ^r. The powders have a discrete size distribution from 0.2 to 1.0 fim. His recent simulation results show that a higher centrifugal acceleration of 1600-g enhances the size separation effect and only a process time of about 10 min is needed for a complete sedimentation of the particles with d = 0.2/im. Once the size separation effect is increased, the quality of the pore-gradient structure can be improved. Ideally the pore-structure in each layer can be considered as the random packing structure of nearly monosized spheres. Hereby the large particles form the large pore on the bottom of the sediment, while small particle form the small pore on the top layer of the sediment. This pore-gradient structure is analogous to an asymmetric membrane (see [12]) which is prefered in the filtration applications. The resistance of an

175

Top surface of sintered specimen atl000^Cfor3h

Figure 2: Scanning electron micrograph of the top layer of a sintered body prepared by stabilized suspension with 5 vol% solid concentration. asymmetric membrane to mass transfer is determined largely or completely by the thin top layer. 3 EXPERIMENTAL PROCEDURES Commercial a-Al203 powders were used to prepare aqueous suspensions with 5 vol % solid concentration. These powders contain 70% Alcoa CT3000SG with d^o = 0.6 — 0.8//m and 30% Alcoa Premalox with 0^50 = 0.1 — 0.3 jwm. The powders were initially dispersed in distilled water at pH=2 and then adjusted to pH=4 by using HCl and were homogenized in a planet mill for 3 hours. The PVA binder (2 mass%) were used to prevent cracking during drying. Once prepared, the suspensions were filled in a plastic cylindrical container with a suspension height of about 70 m m and bottom diameter of about 32 mm. A centrifugal acceleration of about 1600 5^ were applied for casting for about 1 hour. After centrifuging, the supernatant was poured off and the consolidated greenbodies on the bottem plate were dried at room temperature for 72 hours. The dried greenbodies were then calcined in air at lOOO^C for 3 hours, where the heating and cooling rate of lO^C/min were applied. The sintered compacts with a thickness of about 5 mm were polished and thermally etched and examined by using a scaning electron microscope at five different positions across the thickness. Furthermore, the specimens were cut off from the bottom gradually. After each cut-off they were investigated for the pore-size distribution by using a mercury porosimeter. To calculate the pore-size distribution, a mercury surface tension of 0.48 N/m and a contact angle of 140" were assumed. For comparison, flocculated suspensions at pHiep and suspensions with 10vol% were prepared and investigated by the same procedures. 4 RESULTS A N D DISCUSSION Fig. 2 shows scanning electron micrograph (SEM) of the top layer of a sintered body prepared by stabilized suspension with 5 vol% solid content. Hereby the homogeneous distribution of pores with d <^ I fim can be observed. This range of pore size can be used for the UF filtration applications. The SEM pictures at five different positions across the thickness have been shown in Fig. 3. These pictures indicate a clear pore-gradient structure. In order to quantify this pore-gradient structure, the corresponding mercuryporosimetry data have been shown in Fig. 4(a-c). From (a) to (c) in Fig. 4, a pore-gradient can be observed. This pore-gradient shows the largest pores with d ^ 250 nm in the bottom layer and the smallest pores with d ^ 40 nm in the top layer. According to Kruyer [10] and Mason [11] (see 2.1), the particles which contribute to this pore size of d ^ 40 nm, have a

176

Sintered specimen atlOOO^^CforSh

Pore-gradient structure Figure 3: Scanning electron micrographs at five different positions across the thickness of a sintered body prepared by stabilized suspension with 5 vol% solid concentration. size distribution around 100 nm ( « 40/0.1547). This means that the finest top layer was formed dominately by the Premalox powders with dso = 0.1 — 0.3/im. Fig. 4(d) shows a much narrower size distribution and larger smallest pores than Fig. 4(c) because the flocculated suspension has a disturbed size separation during casting. Meanwhile there is no pores larger than 200 nm to be detected. This means that the large particles do agglomerate with another particles during casting and form the multimodal packing structure with smaller pores than the nearly monomodal packing structure will do (see [12]). The size distribution of specimen made by 10 vol% suspensions in Fig. 4(e) has a wide distribution like the specimen made by 5vol% suspensions in Fig. 4(c). These experimental results were also predicted by the DEM simulation results [12]. 5 CONCLUSIONS Using commercial AI2O3 powders, flat membranes with a graded pore-size distribution from 40 to 250 nm across a membrane thickness of about 5 mm have been fabricated by centrifugal casting. For a thicker pore-gradient membrane, a higher solid concentration in suspension can be used. Through the variation of grain size combination of different starting powders, pore gradients with different pore-size distributions can be controlled. Flocculated suspensions resulted to a much narrower pore-size distribution and larger smallest pores in the pore gradient. Therefore, it can be concluded that the not-stabilized suspensions are not suitable to be used as starting suspensions for fabrication of pore-gradient membranes with a wide pore-size distribution and stabilized suspensions have to be used if smaller pores on the top layer are wanted. Furthermore, the centrifugal casting can be considered as a promising forming technique for fabricating the flat pore-gradient membranes. The DEM simulation technique can be used for the further process optimization and for the understanding of the pore-gradient formation mechanics.

177

d in |im (log) (a) pH=4,5 vol%, 3 mm cut-off bJO S

•TTT

""^- "

'k a (D ° a ^ o

>

G S

o C o

1

1

•

I 1 1 III

L

_•

• ' • . • • •

/'

d in |Lim (log) (d) pHiep, 5 vol%, no cut-off L I

s

^

/l

f /

-

f/

1) J '

1/

o

P. \

>

1 /; /'

Mercury-porosimetry data for sintered specimens at lOOO^C for 3 h

d in jLim (log) (b) pH=4,5 vol%, 1 mm cut-off

-s?

^TL

1 1 ,1

s 6

1 ' 1' 1 11 1'

o -

a 3

11

o

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/( //

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cd

d in jLim (log) (c) pH=4, 5 vol%, no cut-off

OH

.

J. V

"

O

>

d in jim (log) (e) pH=4,10 vol%, no cut-off

Figure 4: Cumulative pore-size distribution of sintered specimens made by different suspensions with different degree of bottom cut-off.

178 A c k n o w l e d g m e n t : The authors would like to thank the German Science Foundation (Deutsche Forschungsgemeinschaft, DFG) for the financial support. REFERENCES [1] A. Grull, Prospects for the Inorganic Membrane Business^ Key Engineering Materials Vols. 61 & 62 (1991) pp. 279-288, Trans Tech Publications, Switzerland. [2] A. Larbot, J.-P. Fabre, C. Guizard, L. Cot and J. Gillot, New Inorganic Ultrafiltration Membranes: Titania and Zirconia Membranes, J. Am. Ceram. Soc, 72[2] 257-61 (1989). [3] K. K. Chan and A. M. Brownstein, Ceramic Membranes - Growth Prospects and Opportunities, Am. Ceram. Soc. Bull., 70[4] 703-7 (1991). [4] R. Kohl, G. Tomandl, A. Larbot, L. Cot, Herstellung und Charakterisierung von keramischen Membranen aus Titanoxid zur Cross-Flow-Ultrafiltration nach dem Sol-GelVerfahren, Kurzreferate pp. 64-66, DKG-Jahrestagung, Bayreuth, 4.-7. Oktober 1992. [5] C. Miiller, S. Gottschling, R. Kohl, G. Tomandl, Zr02 Ultrafiltration Membranes Processed by the Sol-Gel-Route, Ceramic Transactions, Vol.51: Ceramic Processing and Technology, American Ceramic Society, (1995) pp. 689-93. [6] R. Kohl, G. Tomandl, A. Larbot, L. Cot, Herstellung und Charakterisierung von keramischen Membranen aus Titanoxid zur Cross-Flow-Ultrafiltration nach dem Sol-GelVerfahren, Kurzreferate pp. 64-66, DKG-Jahrestagung, Bayreuth, 4.-7. Oktober 1992. [7] M. Zhou, G. Meng, D. Peng, G. Zhao, Studies on the Sol Gel Process for Ti02 Membrane Formation, Key Engineering Materials Vols. 61 & 62 (1991) pp. 387-390, Trans Tech Publications, Switzerland. [8] A. Larbot, J. A. Alary, C. Guizard, L. Cot and J. Gillot, New Inorganic Ultrafiltration Membranes: Preparation and Characterization, Int. J. High Technol. Ceram., 3, 143-51 (1987). [9] C. Guizard, A. Julbe, A. Larbot, L. Cot, Nanostructures in Sol-Gel Derived Materials. Application to the Elaboration of Nano filtration Membranes, Key Engineering Materials Vols. 61 & 62 (1991) pp. 47-56, Trans Tech Publications, Switzerland. [10] S. Kruyer, The Penetration of Mercury and Capillary Condensation in Packed Spheres, Trans, of Faraday Society, Vol.54, 1758-67 (1958). [11] G. Mason, A Model of the Pore Space in a Random Packing of Equal Spheres, J. of Colloid and Interface Science, Vol.35, pp.279-287 (1971). [12] C.-W. Hong, Computer-Aided Process Design for Forming of Pore-Gradient Membranes, paper presented at the 4th International Symposium on Functionally Graded Materials (FGM), Tsukuba/Japan, October 20 - 24, 1996.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

179

Mechanical Properties and Microstructure of in-situ TiCp Reinforced Aluminum Base FGM by Centrifugal Cast Zhang Baosheng, Zhu Jingchuan, Zhang Yongjun, Ying Zhongda, Cheng Hongsheng, An Geyin School of Materials Science and Engineering, Harbin Institute of Technology, Harbinl50001,P. R.China ABSTRACT The TiCp/2020Al FGM was fabricated by in-situ reaction in melted Al alloy and the following centrifugal cast, and was investigated by means of microscopy observation and mechanical property test. The hardness and bending strength of the FGM gradully varied corresponding to the composition gradient, and the wear resistance was remarkably improved due to the formation of TiCp-rich surface layer that is strongly bonding with 2024A1 substrate through the TiCp/2024Al graded interlayer. KEYWORDS functionally graded material, centrifugal cast, microstructure, mechanical property, TiC, Al alloy

1. INTRODUCTION Particle reinforced in-situ metal matrix composites (MMCs), with fme and clear in-situ reinforced ceramic particles generated by the exothermic reaction between elements and compounds, possess better mechanical properties than conventional composite materials^^^'. At present, carbides, borides, nitrides and oxides have already successfully been generated on bases of Al, Ti, Ni, Fe and other intermetallic compounds^^'^l By centrifugal cast, the reinforcement particles can be arranged gradiently along the direction of centrifugal force field due to their difference in specific gravity^^l In this paper, the TiCp/2024 FGM was fabricated by in-situ reaction in melted materials and centrifugal cast, and its microstructure and mechanical properties were investigated. 2. EXPERIMENTAL PROCEDURE Firstly, a TiCp/2024 compounded material was prepared by in-situ reaction in the melt. The raw materials are listed in Table 1. Then, the TiCp/2024 composite was put into a special iron crucible. After the composite melting and being refined, the crucible was taken out and

180 put into a centrifuger quickly. The centrifugal cast apparatus was shown in Figure 1, which maximum rotating speed was 5000 rpm and maximum centrifugal acceleration was 5180g. A layer of asbestos was placed in the crucible as insulating material to keep the melt from solidifying too fast for the TiC particles to move. Under the centrifugal force field, the crucible axis, which was initially vertical, became horizontal gradually. Due to the distinction of specific gravity between TiC particles and 2024 Al alloy, TiC particles can move in the direction of centrifugal force. The gradient distribution of TiCp can be controlled by adjusting such parameters as rotating speed, temperature, solidification condition and time of centrifugal cast.

Roter

Fig.l Schematic drawing of centrifugal cast

Table 1. The raw materials used in experiments material size (|im) Al Ti

C Al-Cu-Mg

29 100 45 147 0.5 45 Igot

Composition (wt%) Fe Cu Si H,0 O H <0.30 <0.17 <0.35 <0.15 <0.02

<0.06

0.15

0.15

Mg

Mn

Al

1.6

0.6

other

<0.04

The prepared TiCp/2024 FGM was cut apart radically and axially, and made into metallographic samples. The volume fractions of TiCp in various sections were measured by quantitative metallographic analysis method. The feature of fracture and its changes with microstructures were observed on a Hitachi S570 scan electron microscope (SEM). The distribution of hardness was measured on a HV120 Vickers Hardometer. The static Young's modulus, strength and extensibility of TiCp/2024 FGM were measured in the gradient direction, and as a comparison, the same properties of a uniform TiCp/2024 composite (non-FGM) were also measured on an electron mechanical universal material testing machine (Instron 1186). On a pin-plate wear testing machine, the wear experiments of uniform TiCp/2024 composite and FGM were compared under dry sliding friction condition. The material of pinplate was carbon steel with the hardness of HRC40~50. 3. RESULTS AND DISCUSSION 3.1. Microstructure of TiCp/2024 FGM Figure 2 shows the typical microstructure of TiCp/2024 FGM by centrifugal cast. Under

181 the centrifugal force field, the TiC reinforcement particles distribute in aluminum alloy dispersively, with volume concentration gradiently distributing in the direction of centrifugal force.

( a ) particle enriched zone

( b ) graded distribution zone

( c ) non-particle base zone

Fig.2 Microstructure of TiCp/2024 FGM by centrifugal cast 1

\

{^>^)

og=300L -^g-800 a

110 .^•^

^^iv^ j

/

4 0.2

0.4 0.6 Normal cSiatence

0.8

Fig.3 Change of hardness of TiCp/2024 FGM by centrifugal cast

' .-^^_

0.2

0.4 0.8 Normal disitance

0.8

Fig.4 Distribution of Young' s modulus in FGM by centrifugal cast

3.2. Distribution of Mechanical Properties The macro-hardness distribution in TiCp/2024 FGM is showen in Figure 3, which is similar to that of TiCp concentration. The maximum hardness of outer surface of the FGM reaches to 230Hv, 40 percent higher than that of TiCp/2024 non-FGM. Figure 3 also illustrates that, in spite of different centrifugal accelerations during FGM producing, the macrohardness increases with the rising of reinforcement particles' concentration. In the particle enriched zone, the maximum hardness exists and the curve is the steepest and vice versa. The macrohardness of the particle enriched zone in FGM is obviously higher than that of uniform composite, which makes TiCp/2024 FGM more suitable for working condition requiring higher wear resistance. Figure 4 presents the distribution of Young' s modulus of TiCp/2024 FGM, which is corresponding to the distribution of TiCp content as shown Figure 2.

182 The tension testing results of two typical FGMs in different parts along the direction of gradation are shown in Figures 5 and 6, respectively. Number 1, 2, 3, 4, 5 represent different parts of specimen used for test from inner to outer surface. In TiC particles reinforced FGM, with the increase of concentration, the strength and plasticity distribute gradiently. In the 1

2

3

Gradient direction \, '' • >. x^

!f

V 1 '°

1

. "K \^ \>

og=3Q0 -g=600 V

\

N-..L\ "'K\ 1 r^^

Normal (flatance

Normal dbttanca

Fig.5 bending strength change of TiCp/FGM

Fig.6 elongation change of TiCp/FGM

Load

Normalized distance (b)

( a)

700 600 MPa 500 400 300

III. I (c)

Fig.7 The change of bending strength in three type of FGMs (a)shematic diagram of the FGM; (b)three typical compositional distributions; (c) bending strength

183 particles enriched zone, the strength is extremely high, while the toughness of base is perfect. Comparing with the non-FGM, FGM exhibits better properties in strength and tougness. To study the effects of constitution distribution on the flexure strength, three typical FGMs (A, B, C) with different compositional distribution curves were selected. Their sizes, loading means and testing results are shown in Figure 7. The results reveal that different distributions have different influences on the strength. FGM A has the greatest strength, and distribution curve of A is beneficial for rising the strength of material. The fracture morphology of TiCp/2024 FGM bending specimens are shown in Figure 8. The fracture characteristics changes obviously from particle zone to the non-particle base zone, which corresponds to the distribution of mechanical properties as shown Figures 3-6. In the TiCp rich zone, the FGM displays macroscopically brittle but microscopically plastic fracture with small dimples due to the particle reinforcement. With the decreasing of particle content, the fracture mode gradually changes to the typically ductile fracture of metal materials.

( a ) particle enriched zone

( b ) non-particle base zone

Fig. 8 Fractographs of bending fracture of TiCp/2020Al FGM 3.3. Wear Resistance Figure 9 illustrates the weight loss of TiCp/2024 FGM with different centrifugal acceleration (G) as well as the relation uniform composite (G=lg) after dry sliding wear test under the load 84N. It indicates that the wear resistance of TiCp/2024 uniform material is very poor, and is remarkably improved by centrifugal cast. It mainly results from the gradient distribution of the TiCp, and may also be involved with the elimination of TiCp' s aggregation and 150 200 250 300 the enhancement of TiCp with the metal Sliding Distance (m) Fig.9 weight loss of TiC/2024 after dry sliding matrix under the centrifugal force. With the rise of centrifugal acceleration, the wear resistance further ameliorates due to the increase of TiCp content in the surface layer of the FGM.

184 Figure 10 shows the morphology of wearing surface of the different layers in TiCp/2024 FGM. It reveals the effect of TiCp content on the wear resistance more clearly. The surface layer is enriched with TiC particles and displays a smooth surface with very tiny grooves. The interlayer contains fewer TiC particles and exhibits rough surface with more deep groove. But the metal matrix without TiC particles displays very rough surface with unregularly large fillings. Therefore, the TiCp/2024 FGM has a good wear resistance and the high strength and toughness at the same time as was expected.

( a ) particle enriched zone

( b ) graded distribution zone

( c ) non-particle base zone

Fig. 10 Morphology of TiC/2024 FGM after sliding and wearing

4. CONCLUSIONS (1) TiCp/2024 FGM can be produced by in-situ reaction in melted materials and centrifugal cast, in which TiC particles disperse evenly in the base alloy under centrifugal force field and the aggregation phenomenon of TiCp in normal cast is avoided. (2) Physical and mechanical properties gradiently distributed in the TiCp/2024 FGM by centrifugal cast, which is in good accordance with concentration distribution. Thus, through controlling the distribution of reinforcement particles, FGMs with required physical and mechanical properties can be attained. (3) The wear resistance of TiCp/2024 composite can be remarkably improved by centrifugal cast. With the increasing of centrifugal acceleration, the TiCp content of the surface layer in the FGM rises correspondingly, which further ameliorates the wear resistance.

REFERENCES 1. 2. 3. 4. 5.

J. T. Moore, D. W. Readay, H. T. Feng, Kmonroe and B. Mishra, JOM, 1994, (9), 72. Z. A. Munir, Am. Ceram. Soc. Bull., 67 (1988) 342. H. J. Feng, J. J. Moore and G. With, Metall. Trans., 23A (1992) 2373. M. J. Koczar, K. S. Kumar, U.S. Patent No.480837, 1989. Mamoru Mizumo, Tom Matsuoka and Tatsuo Inoue, J. Soc. Mater. Sci. Japen, 42 (480) 1046.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

185

Dispersion and Fabrication of Zr02/SUS316 Functionally Graded Material by Tape Casting process Jeong-Gu Yeo, Yeon-Gil Jung and Sung-Churl Choi Dept of Inorg. Mat. Eng. Hanyang Univ., Seoul 133-791, KOREA

ABSTRACT Zirconia(Zr02) and stainless steel 316(SUS316) functionally graded material(FGM) was fabricated by tape casting method in an aqueous system. For the stable dispersion of ZrOi and SUS316, we observed zeta-potential of each phase with pH variation and investigated the effect of organic additives on dispersion. Tetragonal zirconia polycrystal(TZP) and monochnic zirconia polycrystal(MZP) could be dispersed with addition of polymethacryHc acid(PMAA) as dispersant. SUS316 could be dispersed with carboxymethylcellulose sodium salt(Na-CMC) as suspending agent. ZrOa and SUS316 were dispersed as the optimal condition estabhshed from the electrokinetic sonic ampHtude(ESA). Zr02/SUS316 FGM was fabricated by sintering at 1350°C in Ar/H2 atmosphere. As a results, the microstructure and the interface of the FGM showed the continuous compositional gradient and also, the adjustment of the particle size and the phase type of Zr02 made sintering defects reduced or eliminated.

1. INTRODUCTION When two materials are joined, it is likely to be a step-Hke structure at the heterogeneous interface. The intemal stress at interface and the difference in material properties are strongly related to thermal degradation of materials. It is necessary to minimize the sintering defects and the intemal stress caused by the thermal mismatch particularly in FGM. Therefore, we should control shrinkage and the sintering behavior over the entire compositional range, as follows; 1) The control of the particle size of starting materials, 2) The control of mixing condition for homogeneity, 3) The loading to specimens, 4) Inclusion addition, 5) Temperature gradient[l,2]. Among those methods, hot pressing has the advantage of minimizing the sintering defects which is caused by the different properties between ceramics and metal, but it has difficulties in controlUng the size and the shape of products and mass production. In this work, we try to minimize the sintering defects using MZP particle larger than TZP. So far, there are many ways to produce FGM. Tape casting method is favored owing to making large thin sheets of uniform thickness easily among them.

186 Table 1 Properties of Starting Materials Materials

Melting Point

. , , J

Thermal Expansion coefficient

(t:)

^«^

(io-*/°c)

Modulus of Elasticity (GPa) '--

^•^^

9.6(20~400r) 11.8(20~1000°C)

p .

™"°

Mean Particle size (jm)

Specific surface (m^/g)

i86r20r^ 186(200

0.31

0.15

13.46

TZP*

2719

MZP^'

2677

5.56

6.5(20°C)

210(20r)

0.3

0.367

4.90

SUS316"

1375 1400

8.06

15.9(0-100 °C) 16.2(0-315 °C) 17.5(0~538°C)

260(20 °C)

0,3

6.83

0.11

*3 mol% YiOs-doped Tetragonal Zirconia Polycrystal, ^Monoclinic Zirconia Polycrystal, ^Stainless Steel 316

Table 2 Properties of Organic Additives Materials

Company

Chemical Analysis (wt%)

Average Molecule Weight

Function

Na-CMC*

Yakuri Chem. Co.

Na-CMC 98.5% Water < 1.5%

—

Binder

PEG^'

Aldrich Chem. Co.

PEG 100%

400

Plasticizer

PVA*

Aldrich Chem. Co.

8 7 - 8 9 % hydrolyzed

13,000-23,000

Binder

PVA*

Yakuri Chem. Co.

99% hydrolyzed

66,000

Binder

Daxad-34

Hampshire Chem. Co.

PMAA 25% Water 75%

10,000

Dispersant

D-SK

Dongnam Synthesis Co.

—

—

Antifoaming agent

*Carboxymethylcellulose Sodium Salt, ^poly(ethylene glycol), *Poly(vinyl alcohol)

In general, this process has been performed in nonaqueous system[3,4]. In this work, we intend to manufacture FGM using the tape casting method in an aqueous system. The stable dispersion of slurry is required to casting process. Thus, we investigate the interaction between starting materials and other organics by measuring electrokinetic sonic amplitude(ESA, MBS-80(X), Matec AppUed Sc, U.S.A) and the rheology of slurry with viscometer(Brookfield DV-U). Therefore, the optimal condition for dispersion was determined. The compacted green sheets with various composition are sintered and then characterized.

2. EXPERIMENTAL PROCEDURE We used the starting materials such as 3 mol% Y203-doped tetragonal zirconia

187 polycrystal(TZP, 99.9%, Tosho co., Japan) and monoclinic zirconia polycrystal 1 (MZP, SC30, 99%, Imperial polychemicals 1 TZP 1 |SUS316 TMP 1 1 1 Dispersant CO., U.S.A.) as ceramics, and stainless 1 i 1 steel 316(SUS316, Anval co., Sweden) as Mixing metal. The properties of starting materials 1 are shown in table 1. TMP is arbitrarily Plasticizer, Mixing Binder denoted as powder mixture with various Antifoaming agent ratios of TZP and MZP. A variety of 1 Defoaming organics was added for dispersion and 1 casting and the properties of organic De-agglomeration additives are shown in table 2, 1 As shown in the experimental procedure Tape Casting of figure 1, in a vacuum chamber we 1 evacuated for 2 hours to remove foams in Drying Stacking | slurry. To prevent interparticle agglomeraPressing tion, 1 hour-sonifmg(Sonifier450, Branson ultrasonics, U.S.A.) and sieving were 1 Bum-out done. Dispersed slurry was casted on flat 1 mylar film and was slowly dried at humid Sintering room in order not to be induced crack resulting from the different drying rates. 1 Characterization Then the stacked sheets were pressed at 20 MPa and burned-out below 500 °C. Figure 1. Experimental procedure FGM was sintered at Ar/H2 atmosphere at 1350 °C for 2 hours. Microstructure of the fabricated FGM was observed by optical microscope(Reichert Metaplan 2, Leica, Austria) and SEM(JSM-5200, Jeol, Japan). Also, the continuity of interface was identified by SEM and wavelength dispersive spectrometer(WDS, JAX-8600, Jeol, Japan). M.ZIP]

Mixing

Various r iri<w, 12hrs-haU milling

TZP

Various ratios, Ihr-hall milling

:>24hrs-hall milling

in vacuum chamber

Ihr sonifing and sieving

Preparing FGM

2hrs at 1350 t in latm Ar-H, atmosphere

3. RESULTS AND DISCUSSION 3.1. Dispersion of ZrOi and SUS316 TZP and MZP powders are dispersed in deionized(DI)-water. Figure 2 shows zeta-potential of TZP, MZP and SUS316 with pH variation. The dispersant of TZP is PMAA(Polymethacryhc acid) which has been used frequently in dispersing ceramic powders[5,6]. Isoelectric point(IEP) of TZP and MZP shifts toward the acidity as the weight percent of the PMAA increases. Thus, we choose the optimal dispersion condition as TZP + 0.5 weight percent(wt%) PMAA and MZP + 0.3wt% PMAA at pH 8.0. In the case of SUS316, the dispersion effect of organic additives is rarely shown. We can know that the dispersion of SUS316 is, therefore, effective by means of steric hindrance rather than electrostatic repulsion. We can obtain the optunized dispersion condition as SUS316 + 0.7wt% carboxymethylcellulose sodium sak

188 \ \

40

^ ^ .

> ^30

•~...

"*"*--^.

£io (0

1"

a. -10

* A »v *

— TZP —X— 72P(50vol%)+MZP(50vol%) ----TZP+0.1PMAA •—-•—^ * TZP+0.5PMAA V , ^ - ^ M2P+0.1PMM V.^^- <^ MZP+0.5PMAA

«

*^.

*

•-

""'\ ^ ** v^ <^

"•'-,

\ "«,

^^ ^ ^

— • — SUS316

; " '

" •^•"."x'X

»

-

'""•'*>'v^

(0

«.

v__

•

- •-

SUS316-f0.1PMAA

*

SUS316-fO.SPMAA

- '-

SUS316-f0.1Na-CMC

-

SUS316-»0.5Na-CMC

«-

''^•"•~'•-.'••..-.-.^^

.••

•••...-••-.'• '^

'\v.

-30

"'~"-v-,,,^^^^-^*'' -40

*-o..^ 3

4

5

^0-°'*' 7

PH

(a)

8

9

PH

(b)

Figure 2. Zeta potential of (a) TZP and MZP, (b) SUS316 with pH variation respectively

(Na-CMC) + 0.015wt% PMAA at pH 8.0. When all the additives are added, the viscosity of slurry is about 4,000—5,000 cps. At that time, soHd loading of TZP, MZP and SUS316 are 25, 40 and 80 volume percent(vol%), respectively. Heating schedule is determined by TG-DTA data[7], from which most of the binder is bumed-out below 500 °C.

3.2. Control of the sintering defects & optimization of compositional gradient In the fabrication of FGM, it is essential to control the sintering defects. So, the adjustment of particle size and phase type of ZrOi, that is, TMP/SUS316 FGM(rather than TZP/SUS316) can prevent defects generated in FGM. This result is ascribed to two things, one is the larger particle size of MZP, which plays a role in releasing the difference of shrinkage rate between ceramics and metal. The other is 3—5 vol% expansion on cooling[8]. From the data about the influences of metal volume fraction and sintering temperature on shrinkage rate[4], we can estabhsh the suitable number of stacking layer and compositional gradient. The number of stacking layer is 16.

3.3. Microstructure and continuity of interface Some discontinuity created frequently in multilayer materials, are investigated at the interface of the fabricated FGM but the continuous compositional change throughout FGM is observed with optical microscope as shown in figure 3. As it were, the composition changes continuously from the Zr02-rich region to the SUS316-rich region. We observe the dispersive structure in the metal side which ceramic particles are dispersed in the metal matrix. Dispersive structure is shown in the ceramic side to the contrary, in other words, metal particles are dispersed in the ceramic matrix, as if dispersing particle and matrix are woven each other. For intermediate composition, network structure is observed, which two phases are Unked respectively as the volume

189 ZrOi

SUS316

Figure 3. Multilayer and interface structure

a) lOOjm thick, pressureless sintering b) 200/M thick, pressure sintering c) 400/m thick, pressure sintering Figure 4. Photograph of FGM(16 layers)

SUS316

ZrOz

'/«y#iiliftlfeM a)

b)

Figure 5. Distribution of a) Fe and b) Zr in ZrOz/ SUS316 FGM with WDS

fraction of each phase increases. These microstructure are consistent with the model designed theoretically and proposed by R. Watanabe et al[2]. Because of the different soUd loading of each phase in each layer, it shows the different shrinkage rates. Therefore, it is Hkely to cause the various flaws in the sintered sample such as warping, crack, delamination and so on. Besides, ZrOz particle is smaller than SUS316 particle, therefore, shrinkage of ZrOi is larger, so the warping to ceramics happens easily. As shown in figure 4, this is relieved appreciably by pressing and increasing the layer number. Pressure sintering just means that sintering is performed under some load to specimen. In pressureless sintering, warping appears toward the cerainics-rich region although we control the difference of shrinkage and the sintering behavior between ZrOi and SUSS 16. It is thought that it is effect of the difference of soHd loading. This is verified by the observation of optical microscope as well. Consequently the difference of shrinkage rate and sintering

190 behavior between each layer can be decreased by the adjustment of the particle size and phase type of ZrOz so that sintering defects are minimize. Figure 5 shows the distribution of Zr and Fe.

4. CONCLUSION Zr02/SUS316 FGM was fabricated by tape casting method in an aqueous system. The dispersion state of ZrOa depended on the electrostatic stabiUzation and that of SUS316 depends on the steric stabilization, respectively. We could control the sintering defects due to different shrinkage rates of starting materials by adjusting particle size and phase type of Zr02. It was a key factor to press specmiens and increase the layer number, specially in problem of warping. Furthermore, the residual stresses induced on FGM were relaxed as the thickness and the number of compositional gradient layer was increased. Also, the continuity of interface and the microstructure of Zr02/SUS316 FGM were observed.

REFERENCES 1. R. Watanabe, Powder processing of functionally gradient materials, MRS Bull., (1995) 3 2 - 3 4 2. R. Watanabe and A. Kawasaki, Development of FGMs via powder metallurgy. Powder and Powder Met., 39(4) (1992) 279-286 3. M. Takemura and M. Tamura et al., Mechanical and thermal properties of FGM fabricated by thin sheet lamination method. In Proc. 1st. Int. symp. FGM, edited by M. Yamanouchi, M. Koizumi (1990) 97-100 4. A. Kawasaki and R. Watanabe, Fabrication of sintered functionally gradient materials by powder spray forming process. In Proc. 1st. Int. symp. FGM, edited by M. Yamanouchi, M. Koizumi, (1990) 197-202 5. M. Itoh, The fluidity of zirconia slurries as a function of ammonium polyacrylate molecular weight. Ceramic powder science vol. 4, (1991) 251—256 6. M. Hashiba et al. Dispersion of ZrOi particles in aqueous suspensions by ammonium polyacrylate, J. Mat. Sci., 24 (1989) 873 — 876 7. P. Calvert and M. Cima, Theoretical models for binder burnout, J. Am. Ceram. Soc, 73[3] (1990) 575-579 8. A. H. Heuer, Transformation toughening in ZrOi-containing ceramics, J. Am. Ceram. Soc, 70[10] (1987) 689-698

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

191

Fabrication of Zr02/Ni and Zr02/Al203 Functionally Graded Materials by Explosive Powder Consolidation Technique A. Chiba, M.Nishida, K Imamiira, H.Ogura and Y. Morizono Faculty of Engineering, Kumamoto University, Kumamoto City 860 Japan Zr02/Ni and Zr02/Al203 functional gradient materials(FGMs) were fabricated by our newly developed underwater-shock consolidation technique. The experimental assembly consists of three parts, i.e. explosive container, water tank and powder container from top to bottom. The energy of shock wave increased by the convergence of shock wave due to the reduction of the cross section area of water tank. The increased shock pressure and long shock duration facilitated the consohdation of difficult-to-consohdate powders to fuU density. The pressure level obtained was estimated to be 7 GPa. The 5 to 10-layered FGMs for both systems were fabricated in an overall cross section without any discontinuity and defects. The obtained FGMs could withstand the thermal stress by cyclic differential thermal load testing. 1. INTRODUCTION Functional gradient materials characterized by the materials having continuously varying material property from one surface to the other, have been prepared via several methods such as CVD, PVD, ion plating, plasma spraying, sintering and self-propagatuig high temperatures synthesis[l]. Shock consolidation is known as a technique to produce the bulk materials from powders. However, it was hitherto difficult to obtain sound specimens without any cracks and/or central hole by using an axisymmetric explosive consolidation technique. The authors have so far developed the dynamic consolidation of difficult-toconsolidate powders such as Zr02 and Si3N4 ceramic powders, utilizing underwater-shock wave generated by the detonation of explosive as a pressure medium of water. And they have reported that sound specimens can be fabricated without using any sintering additives[2]. The advantages of the shock consolidation are that the compact with full density can be fabricated within nano second order and lots of defect induced and fresh surface exposed by shock pressure can reduce the post-sintering temperature in comparison with the conventional process. The purpose of the present study is to establish the technique for fabricating two kinds of FGMs, i.e. Zr02/Ni with 10 layers and Zr02/Al203 systems with 5 layers, by using the underwater-shock consohdation technique mentioned above and to investigate microstructures and thermo-mechanical properties.

192 2. EXPERIMENTAL PROCEDURE 2.1. Materials Partially stabilized zirconia powder (TZ-3Y)was supplied from Tosoh Co., Tokyo, Japan. The content of Y2O3 in the powder was 3.64 mass% and the average particle size was about 0.3 \im. Ni powder was supplied from Nilaco Co.,LTD. Tokyo J a p a n and the average size was about 5(jim. AI2O3 powder was supplied from Sumitomo Chemical Industry LTD. Tokyo, Japan and the average size was about O.S^im. 2.2. Underwater-shock consolidation assembly The assembly used in the present study consists of explosive container, water tank and powder container from top to bottom as illustrated in Figure 1. The mixed powders with prescribed ratio were tapped into the powder container as illustrated in the right hand side of Figure 1. Hard steel powder in the top side was used to prevent the scattering of FGM powders from the powder container and SUS304 powder in the bottom side was used for momentima trap.

water

hard steel powder (100 i/m)

SU$304 powder •AlaOipavfiti 30ZrG2-70AJ2d3, ;5fnm-50M2O3 TOZirdi-aOAliOJ

ZrOz powder SUS304 powder

no[Mixing ratio:vol% , Unil:mm]

Figure 1. Schematic illustration of underwater-shock consolidation assembly. All containers are made of mild steel. The magnitude of the shock wave associated with the explosive detonation increases by convergence of shock wave due to the reduction of cross section area and reflection on the conical wall of the water tank. The pressure level and shock duration can be easily controlled by adjusting the mass and detonation velocity of explosive and the conical angle of the water container. The explosive used in the present study was SEP (provided by Asahi Chemical Industry Co.LTD.) mainly consisting of nitric ester and the detonation velocity was 6900m[/s. The Zr02/Ni FGM was sintered at 1673K for 2h. The Zr02/Al203 FGM was firstly degassed at 1673K for 2h and then sintered at 1823K for 4h.

193 2.3 Estimation of shock pressure The shock velocity at the powder container was measured by the ion gap method, and the shock pressure was calculated using the equations by Cole [3] and Penny and Dasgupta [4] expressing the relationship between the shock pressure and velocity of water. The shock velocity was measured by arrival time difference of shock wave between pins with different height. The velocity of the shock wave was measured to be about 4600m/s, so t h a t the shock pressure obtained by this assembly was estimated to be 7 GPa in the powder container. 2.3. Microstructural investigation and thermomechanieal testing Microstructural charaterization of the obtained FGMs were performed by optical microscopy, scanning electron microscopy (SEM:JSM-6100) with EDX analysis (JED-2000). Thermomechanieal property of Zr02/Ni FGM was investigated by the cyclic differential thermal load testing. 3. EXPEraMENTAL RESULTS AND DISCUSSION 3.1. Microstructures of as-consolidated state Figure 2 shows the cross sections of 10-layered Zr02/Ni and 5-layered Zr02 /AI2O3 FGMs cut parallel to the shock wave direction.

EISBl

W'Ma Smm "^ *^

Figure 2. External views of FGMs. The thickness of the consohdation state reduced to about 11/19. in Zr02/Ni, and 11/16 in Zr02/Al203 system compared with the tapping state. From macroscopic observations, no cracks were found in the specimens and continuous compositional change were confirmed. Figure 3 shows the results of EDX analyses of the two systems. Both the back electron images in (a) and Ka-line images of constitutional elements clearly indicate that the continuous compositional changes were achieved in each system. It is concluded that the FGM produced by the underwater-shock consolidation has fiilly densified microstructure eveiywhere in an overall cross section without any discontinuity and defects. 3.2 Microstructures of annealing state Scanning electron micrographs of bonding interface of the Zr02/Al203 FGM are shown in Figure 4. All of the interface completely bonded. In the 70%Al2O3-30%ZrO2 layer, the both grains of AI2O3 and Zr02 are fine about 2 \IWL in diameter due to the inhibiting the grain growth each other by

194 'M

mZt~K a X<

(i:)Hi-lvw X-rcn imiun:

Figure 3. EDX analyses of Zr02/Ni and Zr02/Al203 FGMs.

10 ff m

Figure 4, SEM microgi'aphs of interface of Zr02/Al203 FGM, (a) Al2O3-30ZrO270Al2O3, (b) 30ZrO270Al2O3-50ZrO250Al2O3 (c) 50ZrO250Al2O -70ZrO230Al2O3, (d) 70ZrO230Al2O3-ZiO2

195 pinning the minor phase grains. The mechanical properties such as fracture toughness of the Zr02/Al203 FGM are now under study. 3.3. Thermomechanical properties of Zr02/Ni The cyclic differential thermal load testing was performed to the Zr02/Ni FGM by using an infrared ray furnace. The surface of Zr02 side was heated to 1273K for 2 min. On the other hand, the surface of Ni side was cooled with water flow and maintained to 573K, therefore the temperature difference was 700K After the 30 thermal cycles, there were no cracks and teas in specimens. A micro-crack was formed on the surface of Zr02 vertically after 50 cycles. However, this crack did not propagate to the ZrO2/90ZrO2-10Ni boundary as shown in Figure 5. This crack might be formed during cooling the specimen. This fact suggests that the residual thermal stress between monoHthic Zr02 and Ni was fairly relaxed by the presence of the composition gradient layers.

Figure 5. Optical micrograph of Zr02/Ni FGM after cyclic differential thermal loading test(50 cycles). In summary, obtained results in the present work can be stated as follows; 1. Zr02/Al203 and Zr02/Ni FGMs were fabricated successfully by the underwater-shock consolidation technique and they were densified fully without any cracks and warps. 2. By cyclic differential thermal load testing with the temperature difference about 700K between the top and bottom of the Zr02Wi FGMs, no cracks appeared after 30 thermal cycles. The Zr02/Ni FGMs showed good thermal durability. This work was partially supported by a Grant-in-Aid for Scientific Research(C) (1994) from the Ministry of Education, Science and Culture, Japan. We thank Mr. Y.Ishitani of Kumamoto University for assistance with the shock consolidation experiment. Steel materials for making the assembly were kindly provided by Godo Steel Ltd. REFERENCES 1. A.Chiba,M,Nishida,KImamura,T.Anraku, and C.Moon,Proc. of FGM'94,Lausanne,(1994),21. 2. A.Chiba, M. Fujita, M.Nishida, R.Tomoshige, Shock waveand high-strain rate phenomena in materials, Marcell Dekker, New York,(1990),415, 3. R.H.Cole, Underwater Explosion, Prenceton University Press, Princeton, New Jersey 1948. 4. W.G. Penny and H.K Dasgupta, Underwater Explosion Research vol.l,Office of Naval Research,(1950),35.

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I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

197

Development of metal/intermetallic compound functionally graded material produced by eutectic bonding method. S. Kirihara, T. Tsujimoto and Y. Tomota Faculty of Engineering, Ibaraki University, 4-12-1 Nakanarusawa, Hitachi 316, Japan

1. INTRODUCTION In the future, structural materials which are light weight and have a high resistance to heat, oxidation and shock will be needed. But it is difficult to produce homogeneous materials which have all these properties. Functionally graded material includes a prominent concept for realizing multi-functional material [1]. However, many problems exist before putting FGMs to practical use in the current production way. This cause is attributed to that the present manufacturing process are very complicated and costly, because the process is by material accumulation method in which the material of different composition are accumulated by turns. And in the current candidates of FGMs, the compositional gradient is difficult to be improved after production. Eutectic bonding method which is a new way of FGM production was invented to solve these problem inherent in the conventional FGMs and their production way. Purposes of this study are to produce FGM by the eutectic bonding method and to prove effectiveness of this method [2].

2. PRINCIPLE On production of FGM by the eutectic bonding method, a metal and an intermetallic compound are combined using a eutectic reaction which exists between both the substances. An important advantage in this combination of the substances is that relieving thermal stress is easy. Because the difference in thermal expansion coefficients between metals and intermetallic compounds is much smaller than that between metals and ceramics. A concrete production way is as the following : Two plates of the metal and the intermetallic compound are placed in contact and they are heated above the eutectic point. The eutectic reaction occurs at the contact plane and eutectic melt is formed. When the melt is solidified by cooling, the metal and the intermetallic compound are bonded and at the same time compositional graded part is formed in the eutectic bonding part, in which a ratio of the metal and the intermetallic compound changes almost continuously. While structure of the bonding part becomes confuse on the solidification by rapid cooling, good compositional gradient is obtained in the case of slow cooling. This means that the structure of FGM, i. e., the compositional gradient can be controlled by changing cooling rate.

198 In this study, using Ti as light weight and shock resistance material, TisSn and TisSia as light weight, heat and oxidation resistance material [3], Ti/TiaSn and Ti/TisSis FGM were produced by the eutectic bonding method. Ti-TisSn and Ti-TisSis binary phase diagram are shown in Figure 1 and Figure 2 [4].

3. EXPERIMENTAL PROCEDURE 3g button ingots of pure Ti, TisSn (Ti-25at%Sn) and TisSis (Ti-35.5at%Si) intermetallic compounds were produced by arc-melting in an Ar atmosphere. TiaSn and TisSis buttons were homogenized for 6 hours at 1400°C and for 12 hours at 1200°C in an Ar atmosphere, respectively. The bottoms of buttons were polished and put together so that their bottoms might be in total contact, and fixed by wrapping Ta foil. The pair was eutectic bonded by heating above the eutectic temperature and then cooling. Heat treatment patterns shown in Figure 3 and Figure 4 were used in order to control compositional gradient part. Microstructure of compositional gradient part in fabricated FGM was observed by SEM with BEI mode.

4. RESULTS AND DISCUSSION Microstructures of Ti/TiaSn and Ti/TisSis FGM are shown Figure 5 and Figure 6. In these structures, existence of the eutectic bonding part between the metal and the intermetallic compound, and change of structural gradient in the eutectic bonding part with changing cooling rates can be confirmed. In Ti/TisSn FGM, structure of the eutectic bonding part becomes a mixture of grates and small dendrites in the case of rapid cooling, continuous structural gradient consisting of three steps in formed in the case of slow cooling. The three steps region consists of TisSn primary region which has large TisSn dendrites in Ti matrix containing Sn, Ti/TiaSn eutectic region which has small TisSn dendrites dispersed in Ti matrix, and Ti primary region composing of small Ti grains. Formation mechanism of compositional gradient on the slow cooling is schematically described in Figure 7. As the eutectic melt which has compositional distribution between TisSn part and Ti part is cooled, on both the side of the melt TisSn primary crystal and Ti primary crystal are formed, respectively. The rest melt is eutectic solidified after cooling under eutectic point. In the eutectic bonding part of the Ti/TisSis, a mixture of net and drop shape TisSis dendrite is formed on rapid cooling, and homogeneous structure where only needle shape TisSis dendrites exist in the matrix is formed on slow cooling. Small TisSis precipitates in Ti part and Ti alloy precipitates along TisSis grain boundary are observed. As whole the material is constituted three layers structurer. The formation mechanism of structural gradient in this case is shown in Figure 8. When the eutectic melt that has compositional distribution is cooled, TisSis dendrite appears from Ti side of the melt and grows to TisSia side. After cooling under the eutectic point, diffusing Si and Ti atoms difftising into the button of solid state precipitate as TisSis phase in Ti part and as Ti alloy phase in TisSis part. Both Ti/TisSn and Ti/TisSis FGM have good thermal stress relief function, because those

199 has continuous structural gradient. In fact, thermal stress cracks did not appear on FGM fabrication and repetition test of heating and cooling the fabricated FGMs.

5. FURTHER STUDY The eutectic bonding method can be applied to a combination of shaped metal parts and intermetallic compound powder, and results in coating of the shaped metal part by intermetallic compound. We intend to call this technique "eutectic coating method".

6. CONCLUSIONS 1. Ti/TisSn and Ti/TisSis FGMs with controlled structure were produced by eutectic bonding method. 2. It was verified that thermal stress cracks did not exist in FGM which has good structural gradient. 3. From above resuks, it is concluded that the eutectic bonding method is effective for production of FGM.

REFERENCE 1. "Functionally Gradient Material", The FGM Forum, The Society of Non-Traditional Technology, Kogyo Chosakai Publishing Co., Tokyo, 1993. 2. Kirihara, M. Takeda and T. Tsujimoto, Scripta Met. Mater., 35(2), 1996, p. 157. 3. Murata, T. Higuti, Y. Morinaga and N. Yukawa, JIMS, 6, 1991, p.627. 4. J. L. Murray, Phase Diagrams of Binary Titanium Alloys, ASM, Ohio, 1987, p.289, p.294.

200 1800-

1670°C

L

2200

1670°C

I I I I I 'i I I I I I I I I I I I I I I I I I I

0 Ti

5

10 15 Sn at%

20

I I I I I I I I I I I I I I I I I r r

25 TisSn

Figure 1. Ti-TisSn binary phase diagram.

0 Ti

10

20 Si at%

30

40 TisSis

Figure 2. Ti-TisSis binary phase diagram.

1350jCahour

1640°Cahour /

/ _ _l6p5°C_ _ Eutectic temperature 1570°C

_L3_3Q°C_ . Eutectic temperature 1290°C

FGM-1

FGM-2 FGM-3 Ti/Ti3Sn FGM

Fig.3 Heat treatment patterns for fabricating Ti/TisSn FGMs by eutectic bonding method.

FGM-1 FGM-2 Ti/Ti5Si3 FGM Fig.4 Heat treatment patterns for fabricating Ti/TisSis FGMs by eutectic bonding method.

201

TisSn TisSn primary crystal Ti/TisSn eutectic crystal Ti primary crystal Ti Schematic explanation of the micrographs

Figure 5. Micrographs of eutectic bonding parts in Ti/TisSn FGMs.

Ti5Si3 Ti precipitate Ti/Ti5Si3 eutectic crystal TisSiB precipitate Ti Schematic explanation of the micrographs

Figure 6. Micrographs of eutectic bonding parts in Ti/TisSis FGMs.

202

Ti/TisSneutectic crystal

Figure 7. Schematic illustrations of the formation mechanism of Ti/TisSn FGM.

Ti precipitate

O

Ti/Ti5Si3 eutectic crystal TisSiB precipitate Figure 8. Schematic illustrations of the formation mechanism of Ti/TisSis FGM.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

203

Mechanical Performance of ZrOi-Ni Functionally Graded Material by Powder Metallurgy J. C. Zhu''^ S. Y. Lee^ Z. D. Yin' and Z. H. Laf 'School of Materials Science and Engineering, Harbin Institute of Technology, Harbinl50001,P. R.China ^Department of Materials Engineering, Korea Listitute of Machinery and Materials, 66 Sangnam-Dong, Changwon, Kyungnam 641-010, Korea

ABSTRACT The mechanical performance of a ZrOi-Ni functionally graded material (FGM) developed by powder metallurgy was investigated by means of three-point bending test. It was shown that the mechanical behavior of Zr02-Ni system strongly depends on constitutional variation, exhibiting typical behavior of elasto-plastic deformation similar to metallic material in the Ni-rich composition and typical behavior of elastic deformation and macroscopic brittle fracture in the Zr02-rich composition. The mechanical properties of Zr02Ni system also display various gradient distributions corresponding to constitutional change. With the rise of Ni content, Vickers hardness decreases, but fracture toughness increases remarkably. Young's modulus changes corresponding to the distribution of porosity. Bending strength is quite sensitive to microstructural factors, and its peak and valley point are relevant to the dispersion strengthening of Zr02 particles and the distributional transition of components, respectively. Keywords

functionally graded material, mechanical property, microstructure zirconia, nickel

1. INTRODUCTION Ceramic/metal functionally graded material (FGM) has been developed as a super heatresistant material used in ultra-high temperature and great temperature gradient environment since 1987 ^^' ^\ A FGM usually comprises of two or more different material components and has a constitutional gradation over macroscopic distances^^l The gradual change of constitution between the two sides of FGM can eliminate the traditional joint interface and relax the thermal stress induced by the temperature variation both in fabrication and in service.

204

Since material properties strongly depend on the composition and microstructure, the property distribution in a FGM should be properly characterized so as to optimize FGM design and improve the FGM's performance. We have developed ZrOi-Ni FGM by powder metallurgy^"^'^^, and the purpose of the present work is to investigate its mechanical performance and its microstructural control factors, especially the variation of basic mechanical properties with the compositional change and microstructural transition. 2. EXPERIMENTAL PROCEDURE Partially stabilized zirconia (PSZ, doped with 3 mol% Y2O3) powder (average diameter of 0.5 |Lim) and nickel powder (4 |im) were chosen as the raw materials. After being blended with different Zr02/Ni ratio, the powders were stacked sequentially with stepwise changes in mixing ratio and compacted in a steel die. Then the green compact was hot pressed at 1350 °C under a pressure of 25 MPa for 1-2 h in nitrogen atmosphere, and the Zr02-Ni FGM plate (60x60x5 mm^) was gained finally. The Zr02-Ni uniform composite (non-FGM) plates with various mixing ratios corresponding to each layer of the FGM were fabricated as well under the same conditions. Samples for microstructural inspection and mechanical tests were cut with a diamond saw, and their surfaces were ground and polished. The Vickers hardness of Zr02-Ni composite was determined by indenting with a load of 5 kgf The Young's modulus, bending strength and fracture toughness were measured by threepoint bending test with an electron mechanical universal material testing machine (Listron 1186). For elastic modulus and bending strength tests, the specimens with 3x4x36 mm^ in size and a jig with a span of 30 mm were used. Strain gauges were mounted on the tensile surface of bending specimens in order to determine the stress-strain curve and Young's modulus. Single edge notched beam (2x4x22 mm^) and 16 mm span were adopted for testing fracture toughness. 3. RESULTS AND DISCUSSION 3.1. Microstructure The microscopic observations^"^'^^ demonstrate that the microstructure in Zr02-Ni FGM distributes with a stepwise gradient from nickel side to zirconia side as shown in Figure 1, in which the grey phase is nickel and dark one is PSZ. Li the region rich in Ni, nickel serves as matrix phase and displays typical network structure with dispersed PSZ particles (Figure la,b). hi above 60 vol% PSZ region, the continuity transition of the FGM components can be found (Figure Ic, d). Nickel phase begins to change from connective distribution to dispersive one, but PSZ phase alters from dispersive to connective conjugately. Furthermore, the microstructure of each layer in sintered Zr02-Ni FGM body is quite fine and homogeneous. 3.2. Stress-strain behavior The stress-strain curves of Zr02-Ni system in three-point bending test are presented in Figure 2. It can be found that the Zr02-Ni system shows quite different deformation and fracture behavior with the change of constitution. The nickel-rich material displays typically elasto-plastic deformation behavior, which is controlled by the continuous matrix nickel phase and similar to metallic material. During bending test, the specimen can be bowed to "V" shape

205

Figure 1. Microstructure of interlayers in Zr02"Ni FGM with different compositions: (a) 20, (b) 40, (c) 60 and (d) 80 vol% PSZ

1 — 20 vol%PSZ 2 — 40 vol%PSZ 3 — 60 vol%PSZ 4 — 80 vol%PSZ

0.0

0.2

0.4

0.6

0.8 1.0 1.2 Strain, %

1.4

1.6

1.8

2.0

Figure 2. Stress-strain curves of ZrOi-Ni system with different compositions. with a bending deflection of 2.8 mm but free of rupture. On the other hand, when the content of PSZ reaches to above 40 vol%, the dispersed PSZ particles begin to contact one another

206 and form a network structure (refer to Figure 1), so that the plastic deformation of nickel phase is restricted within microscopic area. Therefore, materials containing PSZ from 40 to 80 vol% mainly exhibit typically linear elastic deformation behavior, having brittle fracture characteristic without macro plastic deformation and a little rupture strain less than 0.4%. It should be noted that the material with 60 vol% PSZ behaves as non-linear elastic deformation after the linear stage, which is related to the microstructure transition around this composition(Figure 1 c). The alternation of matrix phase results in loose connectivity of the two components, and leads to non-linear elastic deformation behavior. 3.3. Mechanical properties Figure 3 illustrates the relationship between Vickers hardness and composition in ZrOi-Ni system. The hardness decreases remarkably with the rise of Ni content and conforms with linear rule of mixture approximately. However, the slope of the hardness curve changes at 40 vol% Ni, which seems to reflect that the matrix phase alters from PSZ to Ni at about that composition. 250

20 40 60 go Volume FfaictimiL of Ni, %

im

Figure 3. Relationship between Vickers hardness and composition in Zr02-Ni system.

10

oEftested) - E(Reiissnile)

20

40

60

80

Volume Fraction of Ni, %

Figure 4. Variation of Young's modulus with composition in Zr02-Ni system.

Figure 4 shows the variation of Young's modulus with composition in ZrOi-Ni system. It reveals that the elastic modulus in Zr02-Ni system does not monotonically change with composition, and disagrees with the Reuss rule of mixture on the whole (refer to the dashed line in Figure 4). This may be attributed to the effect of porosity as presented in Figure 4. Clearly the varying trend of Young's modulus is just opposite from the distribution of porosity in ZrOi-Ni system, that is, the higher porosity corresponds to the lower elastic modulus and vice versa. It can be noted that at 60 vol% Ni content the ZrOi-Ni material is nearly frill densitified so that its elastic modulus conforms to calculated value through the Reuss rule. As illustrated in Figure 5, the distribution of bending strength with compositional change in ZrOi-Ni system is very complex. With Ni content increasing, the bending strength first

207 16

TraasitioBi t*- Brittle ^ 1 I— Ductile —I 20 40 6© VoloMC Fracticmi of ]

im

Figure 5. Change of bending strength with composition in ZrOi-Ni system.

20

40

t

6(D

80

ICO

VoloMe Fracticm of Ni, % Figure 6. Variation of fracture toughness with composition in Zr02-Ni system.

decreases Unearly, then ascends obviously in the range of 40-60 vol% Ni, and finally descends again with the further raise of Ni content. In a FGM, the material strength not only depends on its composition, but also is sensitive to its microstructure including porosity and the distribution of its components. On the one hand, the existence of pores will strikingly weaken the material strength, which descending tendency is corresponding to the rise of porosity (refer to Figures 4 and 5). On the other hand, microstructural transition is the more essential reason for the abnormal distribution of bending strength, hi PSZ-rich compositions (<40 vol% Ni), the PSZ component exists as continuous matrix phase with dispersed nickel component (Figure 1 d). Of course, the flexural strength decreases with the rise of Ni content due to the fact that Ni phase behaves here as pores and defects for its low strength, which is the so-called pore controlled strength characteristics'^^ At about 40 vol% Ni, the alternation of matrix component in ZrOi-Ni system occurs according to the fractal analysis'''^ and microstructural observations as above (Figure 1 c), which causes the rise of porosity due to the loose connectivity of the two components and leads to the minimum strength value. However, in the higher Ni content region, the nickel serves as matrix phase with dispersed PSZ particles (Figure 1 b) and the strength characteristic of dispersion strengthening is exhibited. As a result, the peak strength value appears at 60 vol% Ni. The further increase of Ni content will weaken the effect of dispersion strengthening and induce the decline in the bending strength again. Figure 6 indicates the variation of fracture toughness with composition in Zr02-Ni system. From Figures 1 and 6, it can be concluded that the ZrOi-Ni system manifests a transition of fracture behavior from brittle to ductile with the constitutional change. The PSZ-rich material behaves as typically macroscopic brittle fracture with the lower fi-acture toughness. With the increase of Ni content, the fracture toughness gradually ascends and the fracture behavior begins to transit to ductility. However, the Ni-rich material displays typically ductile fi-acture behavior and its fracture toughness grows swiftly.

208 It should be noted that the pores and microcracks, which remarkably weaken elastic modulus and bending strength, are not disadvantageous to the fracture toughness. Even for the 40 vol% Ni material with maximum porosity, the fracture toughness trends towards rising, which may reveal the existence of microcrack toughening mechanism^^l Moreover, the effect of phase transformation toughening^^^ cannot be ruled out in Zr02-Ni system. But for the experimental conditions in this paper, the fracture toughness may be mainly controlled by ductile phase toughening^ ^^l 4. CONCLUSIONS (1) The mechanical behavior of Zr02-Ni system strongly depends on constitutional variation. The Ni-rich materials exhibit typical behavior of elasto-plastic deformation and ductile fracture similar to metallic material. The materials containing PSZ from 40-80 vol% mainly presents typically linear elastic behavior and macroscopic brittle fracture. However, the material with 60 vol% PSZ behaves as non-linear elastic behavior after the linear stage. (2) The dependence of mechanical behavior on constitution in ZrOi-Ni system results from the variation of microstructure and its distribution, hi the regions rich in Ni or PSZ, the mechanical performance is controlled by continuous matrix component and displays elastoplastic or linear elastic characteristics, respectively. The non-linear elastic behavior at 60 vol% PSZ is related to the connectivity transition of matrix component. (3) The mechanical properties of ZrOi-Ni system also display various gradient distributions corresponding to constitutional change. Vickers hardness decreases with the rise of Ni content. Young's modulus changes corresponding to the distribution of porosity. Bending strength is quite sensitive to microstructural factors, and its peak and valley point are relevant to the dispersion strengthening of Zr02 particles and the distributional transition of components, respectively. Fracture toughness is controlled by ductile toughening of Ni-phase, and increases remarkably with the rise of Ni content.

REFERENCES 1. 2. 3. 4. 5.

M. Niino, T. Hirai and R. Watanabe, J. Japan Soc. Comp. Mater., 13 (1987) 257. M. Koizumi and K. Urabe, Iron and Steel, 75 (1989) 887. T. Hirai, Functionally Gradient Materials, Ceramic Transactions, 34 (1993) 11. J. C. Zhu, Z. D. Yin, Z. H. Lai, J. Mater. Sci. Technol., 10 (1994) 188 Z. D. Yin, J. C. Zhu, Z. H. Lai, X. D. Li and D. Z. Yang, Proc. of the Second Pacific Rim hiter. Conf on Adv. Mater, and Processing, ed. by K. S. Shin, J. K. Yoon and S. J. Kim, The Korean Listitute of Metals and Materials, 1995, 1731. 6. K. Hirani and T. Suzuki., Proc. of the First hiter. Symp. on FGM, ed. By M. Yamanouchi, M. Koizumi, T. Hirai and L Shiota, Sendai, Japan, 1990, 313. 7. K. Muramatsu, A. Kawasaki, M. Taya and R. Watanabe, Proc. of the First hiter. Symp. on FGM, ed. By M.Yamanouchi, M.Koizumi, T.Hirai and LShiota, Sendai, Japan, 1990, 53. 8. A. G. Evans, Advances in Ceramics, Vol.12, Science and Technology of Zirconia II, American Ceramics Society, Columbus, Ohil, 1984, 193. 9. F. F. Lange, J. Mater. Sci., 17 (1982) 225. 10. G. Evans, J. Am. Ceram. Soc, 73 (1990) 187.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

209

Fabrication of PSZ-SUS 304 Functionally Graded Materials H. kobayasi CoUeg of Industrial Technology 1-27-1, Nisikoya, Amagasaki 661, Japan The fabrication of P S Z - S U S 304 functionally graded materials ( F G M s ) h a s been investigated by powder stacking and pressureless sintering processes in a vacuum. PSZ,mixed PSZ and SUS 304,and SUS 304 powders with both a binder and a deflocculant were laminated into a mold,and then were compacted by the vibration pressing method under a pressure of 50 MPa. Green specimens were sintered pressurelessly for one hour in a vacuum at 1623 K. Many sintered FGM specimens with the graded distributions of 2.5,5,10,and 20wt.% showed crakes in the PSZ layer,or in the middle layers between PSZ and SUS304. These cracks are presumablly due to a tensile or compressive stress caused by the difference of the sintering shrinkage of each monolithic layer. However,some FGMs specimens with both the large graded distribution and a diflfemt thickness of each layer were free from cracks,although the specimens were slightly curved in thin ones. 1. Introduction The processing of functionally gradient materials has been proposed by several methods such as CVD,^^ ion beam mixing process,^^ plasma spraying,^^ self-propagating high temperature synthesis, ^^ and green forming and sintering process. ^^ The powder metallurgical processing with powder stacking and sintering is superior to the other processing for a large scale shape. In this ordinary processing, FGMs are fabricated by laminating and pressing many green compacts with a controlled microstructural gradient in a mold. The compacts are formed by a slip-casting ^^ or a powder pressing method.^^ However,FGMs can be fabricated by laminating and compacting in a mold at one process,if the powders are used as starting materials without green compacts. In this work, the vibration pressing method with wet powders was used for the fabrication of FGMs. This method can reduce to applied pressures for compacting in contrast to a conventional die pressing method with dry powders,because the binder content was much more sufficient to allow the freedom of the flow of powders in the state of vibration.®^ Therefore,the vibration pressing method is effective for the stress relaxation in compacts. The green compacts with three dimensional gradient distribution can be obtained. Thus,the fabrication of PSZSUS 304 FGMs has been investigated by this wet vibration pressing method.

210 2. Experimental 2.1. Starting materials The PSZ powder (TZ-3Y-E, Tosoh Corp.) used in this investigation was composed of partially stabilized zirconia containing 3 mol % of YgOg with a mean particle size of 2.90Mm. SUS 304 powder(SF-SUS 304L, Nippon Atomize Kakoh Co.,Ltd)was 6.29 A6m in diameter. These chemical compositions are shown in Tables 1 and 2. Wet powder materials containing PSZ, SUS 304,and mixtures of PSZ and SUS 304 were prepared by blending the powders and a binder in a mortar. The binder was an aqueous solution with a 12 wt.% emulsion tj^e binder (Seruna WE-518,Chukyo Yushi Co.,Ltd) and 8 wt.% dispersant. Table 1. Characterization of PSZ powder. Chemical composition (wt.%) Y2O3 L.O.I FeaOg SiOs NagO TiOg A l A 5.03 0.68 0.002 0.007 0.012 tr 0.244

Mean particle size (Atm) 2.90

Table 2. Characterization of SUS 304 powder.

Ni 10.71

Chemical composition (wt.%) Cr C O 18.23 0.032 0.670

Mean particle size (/im) 6.29

2.2. Experimental apparatus and procedure The vibration pressing method with wet powders was used for the forming of specimens in this investigation. This is a new plastic pressing method with vibration,and it could reduce to applied pressures for compacting in contrast to a conventional die pressing method with dry powders,because the binder content was sufficient to allow the freedom of the flow of powders. The vibration pressing machine for the forming of specimens is shown in Figure 1. This machine consists of a vibration top plate with an oil cylinder and a bottom plate,both of which are fixed with 4 poles, and 4 top and bottom springs supported with 2 plates. The mechanical vibrator with a vibrating force of 300 kg is attached to the bottom of a vibration plate,and it consists essentially of 2 contra-rotating shafts with an out-of-balance weight at each end. The die set in vibration plates possesses a cavity of 30 mm in diameter and 90 mm in depth. The vibration and pressure were applied simultaneously by the top and the bottom punches. PSZ,PSZ and SUS 304 mixture,and SUS 304 wet materials were laminated into a die,and then were compacted by a vibration pressing method. 120 sec is taken to reach a fixed pressure after the die was set in vibration plates. The forming vibration time after a fixed pressure was 60 sec,amplitude was 0.9 mm,

211

® ® (D ® (D ® ®

Spring Vibraition plate Vibrator Die Punch Oil cylinder Powder materiales

Fig.l. Schematic layout of vibration pressing machine. freequency of vibration was 60Hz,and applied pressure was 50 MPa. The obtained PSZ-SUS 304 gradient green specimens were measured as 30 mm in diameter and approximately 4 to 20 mm in height. The green specimens were dried at 373 K, and then sintered without pressures for 1 hour at 1623 K in a vacuum. The relative density and the linear shrinkage of monolithic sintered specimens were measured. The microstructure of monolithic and multilayer sintered specimens was microscopically examined. The gradient distribution was examined by an X-ray line analysis of Fe and Zr. 3. Results and discussion 3.1. Linear shrinkage of monolithic sintered composites The density and the apparent porossity of monolithic specimens after compacted and sintered without loads for 1 hour at 1523 K to 1623 K in a vacuum were investigated. The relative density increased with the sintering temperature. The apparent density of specimens sintered at 1623 K decreased in direct proportion to the content of PSZ. The sintering temperature was determined from these results and the melting point of SUS 304. Figure 2 shows linear shrinkage after sitering. This shrinkage possesed the minimun value for the composition of 30 wt.% PSZ-70 wt.% SUS 304,and an irregular curved line of a V character with some inflection points. It is thought that in case of directly bonded two different monolithic compositions within the ranges of 0 to 30 wt.% PSZ,a tensile stress occures at the side of enriched SUS 304 of the bonding boundaries while a compressive stress does at the enriched PSZ side. Within the ranges of 30 to 100 wt.% PSZ,the either side of the boundaries possesses a reverse stress. A tensile or compressive stress exceeding its own strength limit invites cracks or exfoliations in layers with weak strength or in boundaries. Therefore,the strenth of each layer should be investigated for the fabrication of FGMs.

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Fig.3. Bending strength of monohthic composites.

3.2. Bending strength of monolithic sintered composites The bending strength of monolithic PSZ and PSZ / SUS 304 composites pressurelessly sintered is shown in Figure 3. The bending strength decreases as the content of PSZ increases to 30 wt.%,and then slightly increases with its content. 3.3. Fabrication of PSZ-SUS 304 functionally graded materials Firstly,many multilayer PSZ-SUS 304 FGMs with the graded distribution of 5, 10 and 20 wt.% were fabricated. The typical illustrations in section of sintered spesimens are shown in Figure 4. These FGMs with the graded distribution of 5tol0wt.% possessed large cracks in boundaries and in weak layers,although the FGMs were expected to possess smaller cracks than FGMs with 20 wt.% distribution for the small difference of the linear shrinkage,thermal expansion, thermal conductivity of monolithic composites.^^ The FGMs with 2.5 wt.% gradient distribution over 90 wt.% PSZ also possessed cracks. Nextly,some FGMs with large graded distribution were studied excluding the monolithic composite of 30wt.% PZT-70wt.% SUS 304 possessed the minmum value of linear shrinkage. The comoposition with the minmum linear shrinkage was changed from 30 wt.% PZT-70wt.% SUS 304 to 40 wt.% PZT-60 wt.% SUS 304,and then the monolithic composites with large difference from its curved line of V character shrinkage and with weak bending strength were excluded. These composites possessed slight vertical cracks in a PSZ layer a l o n e . Therefore,some FGMs with large graded distribution and with a different thickness of each layer were studied.

213

Consequently,for example,the FGMs fabricated from multilayers with the compositions of PSZ,65 wt.% PSZ-35 wt.% SUS 304, 60 wt.%PSZ-40 wt.% SUS 304, 40 wt.% PSZ-60 wt.% SUS 304 and SUS 304,or with the near compositions of that were free from cracks, although their FGMs were slightly curved in thin ones. Figure 5 shows the microstructure of a sintered PZT-SUS 304 FGMs with five layers. The FGMs had the relatively smooth graded distribution of Fe and Zr in the middle layers from the X-ray line analysis profile obtained by EPMA on the cross-section of that, although had the substantially different distribution at the both sides of enriched PZT and SUS 304. The FGMs were free from cracks after re-firing at 1473 K with a speed of 30 K per min.

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Fig.5. Microstructure of PSZ-SUS 304 FGMs

214 4. Conclusions The fabrication of PSZ-SUS 304 functionally graded materials was investigated by powder stacking and pressureless sintering process. The laminated green compacts was formed by the vibration pressing method for the stress relaxation in compacts. Many FGMs with small graded distribution possessed cracks in boundaries and weak layers. Some FGMs with large graded distribution and with a different thickness of each layer were free from cracks. References 1. M.Sasaki, Y.Wang, T.Hirano and T.Hirai, J.Ceram.Soc.Jpn., 97, 539-43 (1989). 2. M.Sato, S.Tanaka, I.Hashimoto, E.Setoyama, T.Gejo and T.Sato, Hitachi Hyoron, Vol.68, 821(1986). 3. S.Kitaguchi, H.Hamatani, N.Shimoda, Y.lchiyama and T.Saito, Proceedings ofthe Fourth symposium on FGM, (1991) pp. 149-52. 4. A.Takahashi, K.Tanihara, Y.Miyamoto, M.Oyanagi, M.Koizumi and O.Yamada, Funtai oyobi Funmatsuyakin, 37, 263-66 (1990). 5. R.Watanabe, Proc. Shinzairyo Sosei Toronkai, No.6, Jpn. Inst. Metals (1988) pp. 9-13. 6. H.Takebe, T.Teshima, M.Nakashima and K.Morinaga, J.Ceram.Soc.Jpn., 100[4]387-91 (1992). 7. M.Sasaki and T.Hirai, J.Ceram.Soc.Jpn.,99[10]1002-13 (1991). 8. H.Kobayashi, Trans. Materials Research Soc. Jpn., Vol.5, 109-18 (1992). 9. A.Kawasaki and R.Watanabe, 3rd FGM Symp.(1989) pp. 35-48.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

215

Preliminary characterization of interlayer for Be/Cu functionally gradient materials - Thermophysical properties of Be/Cu sintered compacts Shigeru Saito", Naoki Sakamoto\ Kiyotoshi Nishida^ and Hiroshi Kawamura^ ' Japan Atomic Energy Research Institute, Oarai-machi, Higashi Ibaraki-gun, Ibaraki-ken 311-13, Japan ' NGK INSULATORS, LTD., 1 Maegata-cho, Handa-city, Aichi-ken, 475, Japan At present, beryllium is one of candidate materials for plasma facing materials of fusion experimental reactor. And copper alloy (dispersion strengthened (DS)-copper, Cu-Cr-Zr and so on) is proposed as heat sink material behind plasma facing material. In this application, it is necessary to develop reliable bonding techniques between beryllium and copper alloy because plasma facing components are exposed to high heat load and high neutron flux generated by plasma. Then, we started the bonding technology development of beryllium and copper alloy with functionally gradient materials (FGM) to reduce thermal stress due to the difference of thermal expansion between beryllium and copper alloy. As the interlayer for FGM, eleven kinds of sintered compacts in which the mixing ratio of beryllium powder and oxygen free powder is different, were manufactured by the hot press (HP) and hot isostatic press (HIP) method, and thermophysical properties of these compacts were measured in order to estimate thermal stress of bonding interface. In this paper, the results of measurements on thermophysical properties and of metallographical observation were reported. 1. INTRODUCTION At present, beryllium is regarded as one of candidate materials on plasma facing materials of fusion experimental reactor, because of many advantages such as low Z[l], relatively high thermal conductivity, low activation and so on. Among the different divertor design options, the duplex structure where beryllium armor is bonded with heat sink structural materials (DS-copper, Cu-Cr-Zr and so on) is under consideration[2]. Not only plasma facing components are exposed to high heat load[3] and sputtered by plasma particles, but also the components will be heavily irradiated by 14MeV neutron which is generated by I>T reaction. When plasma facing components are irradiated by 14MeV neutron, many unfavorable effects will be caused in the materials, for example, irradiation damage and He atom production by nuclear transmutation. Under these conditions, it is considered that bonding strength between beryllium and heat sink materials will be decreased. Therefore, it is necessary to develop a reUable bonding techniques between beryllium and heat sink materials in order to fabricate plasma facing components which can resist the many unfavorable effects. Tlien, we started the bonding technology development of beryllium and copper alloy with FGM mterlayer so as to reduce thermal stress of bonding

216 interface due to difference of thermal expansion between beryllium and copper alloy. The concept of divertor mock-up with FGM interlayer is shown in Fig. 1. As FGM interlayer, eleven kinds of sintered compacts in which the mixing ratio of beryllium powder and oxygen free copper powder was different, were manufactured by powder metallurgical method, that is, HP and HIP. In this study, to estimate thermal stress at joining interface, thermal conductivity and thermal expansion coefficient of Be/Cu sintered compacts were measured by laser flash method and laser interferometry method, respectively. The characterization on these compacts was performed by using SEM (Scanning Electron Microscope) to investigate distribution of intermetallic phases on these compacts. Mixture matrix (at.%)

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Fig.l Concept of divertor mock-up with FGM.

Fig.2 Manufacturing flow of Be/Cu sintered compacts.

2. EXPERIMENTS 2.1 Specimens As FGM interlayers, eleven kinds of sintered compacts in which the mixing ratio of beryllium powder and oxygen free copper powder was different, were manufactured by powder metallurgical method, that is, HP and HIP. Detailed manufacturing flow of Be/Cu sintered compacts is shown in Fig.2. At first, beryllium powder and oxygen free copper powder were mixed by mechanical mixing process. Those particle size was -200mesh. And mixture powder was sintered by HP and HIP methods. These processes were performed at 600°C in order to avoid the formation of Be-Cu( /:/ )phase (620°C). Tlien, sintered blocks were cut to disk shape for which is convenient for their characterization. Mixture matrix of those powder is shown in Fig.2. The dimension of each compacts was 8mm in diameter, 2mm in thickness and the sintering density of each compacts was almost 100%T.D.. 2.2 Procedure In this study, thermal diffusivity and specific heat of Be/Cu sintered compacts were measured by laser flash method. In those measurement, specimens were loaded with laser which had constant energy under vacuum. The degree of vacuum was less than 1 X lO^Pa in order to avoid the oxidation of specimens. Thermal diffusivity and specific heat of these compacts were

217 obtained by elevated temperature on the specimen using a thermocouple and an infrared ray sensor. Thermal conductivity was obtained from calculation of those measured values. Thermal expansion coefficient of these compacts were measured by laser interferometry method. The degree of vacuum was less than 1.3X10^Pa. Both of the measurement were performed at temperature from RT to 700 °C. SEM observations were also performed to investigate distribution of intermetallic phases on these compacts. Here, accelerating voltage was 15kV, and electron beam diameter was about (f> 10 fi m. 3. RESULTS AND DISCUSSION At first, the thermal diffusivity, a , and the specific heat, Cp, on sintered compacts of Be/Cu mixture are shown in Fig. 3 and Fig.4, respectively. From these results, it appeared that the thermal diffusivity on sintered compacts which contained less than 50at. %Cu was lower than that of lOOat. %Be. The thermal diffusivity increased with increase of Cu containing ratio and the specific heat gradually decreased with increase of Cu containing ratio.

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To

0 10 20 30 40 50 60 70 80 90 100 Composition of Be/Cu sintered compacts, at.%Cu

Fig. 3 Thermal diffusivity of Be/Cu sintered compacts. The thermal conductivity, K , of sintered compacts was obtained by calculation from the following equation:

0 10 20 30 40 50 60 70 80 90 100 Composition of Be/Cu sintered compacts, at.%Cu

Fig.4 Specific heat of Be/Cu sintered compacts. 300 r E

K = a X Cp X f)

The p is dencity. This is shown in Fig.5. The tendency of the thermal conductivity was similar to the thermal diffusivity. Particularly, the thermal conductivity of sintered compacts which contained more than 50at.%Cu was higher than that of lOOat %Be. Therefore, it is considered that these sintered compacts are advantageous to apply those to FGM

0 10 20 30 40 50 60 70 80 90 100 Composition of Be/Cu sintered compacts, at.%Cu Fig.5 Thermal conductivity of Be/Cu sintered compacts.

218 interlayer. On the other hand, thermal conductivity of sintered compacts which contained less than 50at%Cu was lower than that of 100at.%Be. It is considered that this phenomenon was due to decrease of thermal conduction with free electron by formation of Be-Cu intermetalHc phases or lattice strain. Thermal conductivity of lOOat.% Cu was about 270W/m/K. This value agreed with measured value in literature [4]. (0 E The thermal expansion coefficient of sintered compacts are shown in Fig. 6. From 0 10 20 30 40 50 60 70 80 90 100 these results, between 300°C and 400''C, it Composition of Be/Cu sintered compacts, at.%Cu appeared that thermal expansion coefficient of these compacts gradually increased with Fig.6 Thermal expansion coefficient increase of Cu containing ratio and the of Be/Cu sintered compacts. compacts which contained less than 60at.%Cu was lower than that of DS-copper and higher than that of HP-Be. The bonding interface will be used at this temperature range. Therefore, it is possible to reduce thermal stress of bonding interface by applying those sintered compacts to FGM interlayer.

Fig. 7 SEM photographs of as manufactured Be/Cu sintered compacts.

219

Fig. 8 SEM photographs of Be/Cu sintered compacts used for the measurement of thermal expansion coefficient. The measurement were performed at RT to 700 °C .(The arrows indicate cracks.)

From the SEM observations on sintered compacts of Be/Cu mixture, it became clear that the sintered compacts contained residual beryllium, copper and two kind of intermetallic compounds, namely, Be->Cu( o ) and BeCul/ ). SEM photographs of the as manufactured compacts , which composition are (a)90at.%Cu (b)50at.%Cu and (c)10at.%Cu, are shown in Fig. 7. These compacts were composed of those four phases and the ratio of these phases varied with varying the composition. Fig. 8 shows SEM photographs of the compacts used for the measurement of thermal expansion coefficient. The measurement were performed at RT to 700°C.From this observation, many cracks were observed at the interface between BeCu( / )phase and Cu phase in these compacts. It is considered that these cracks were generated by the transformation of BeCu(/j)phase into BeCu( / )phase and Cu phase with large shrinkage[5,6]. Therefore, the fomiation of Be-Cu(/j )phase, occurred at 620°C, should be avoided when beryllium and DScopper would be joined by diffusion bonding or HIP bonding method.

220

4. CONCLUSIONS In this study, thermal diffusivity and specific heat of Be/Cu mixture sintered compacts were measured by laser flash method, then thermal conductivity was obtained from calculation of those measured values. And thermal expansion coefficient was measured by laser interferometry method. These thermophysical properties were measured in order to characterized those compacts as interlayer between beryllium and copper alloy used in the plasma facing components. The obtained results are as follows. (1) From the results of thermophysical characterization, thermal conductivity of Be/Cu sintered compacts which contained more than 50at.%Cu were higher than that of 100at.%Be. And thermal expansion coefficient of compacts which contained less than 60at.%Cu were lower tlian that of DS-copper and higher than that of HP-Be. Therefore, it appeared that Be/Cu sintered compacts which contained 50-60at. %Cu were advantageous to apply those to FGM interlayer. (2) From the metallographical observations, it became clear that Be/Cu mixture sintered compacts contained residual beryllium, copper and two kind of intermetallic compounds, Be2Cu( d^) and BeCu( / ). And the ratio of these phases varied with varying the composition. (3) SEM photographs of the compacts annealed to 700°C show many cracks due to the transformation of Be-Cu( /^ )phase into BeCu( / )phase and Cu phase with large shrinkage. Therefore, joining with beryllium and DS-copper by diffusion bonding or HIP bonding method should be performed below 620 °C in order to avoid the formation of BeCu( i3 )phase. And the bonding interface also should be used below 620°C. From these results, thermal designs of beryllium and copper alloy bonded plasma facing components using Be/Cu sintered compacts as FGM interlayer became possible. However, it is necessary to give full consideration to effect of brittle intermetalhc compounds in FGM interlayer on bonding strength of beryiiium/copper alloy interface. Mechanical properties tests will be performed in order to estimate themial stress at bonding interface.

REFERENCES [1] D.W.Doll, etal., J.Nucl.Mater., 85-86(1979)191. [2] T.Kuroda and G.Vieider, ITER Plasma Facing Components, ITER Documentation Series N30 (IAEA, Vienna, 1991). [3] K.Nakamura, etal., J.Nucl.Mater., 212-215(1994)1201. [4] C. Carmichael (ed.), Kent's Engneere's Handbook, (John Wiley&Sons, Inc.) [5] N.Sakamoto, H.Kawamura and R.Solomon, to be published in Proc. I9th Symposium on Fusion Technology, Lisbon, 1996. [6] N.Sakamoto, S.Saito, M.Kato, R.Solomon and H.Kawamura, to be pubUshed in Proc. 5th international workshop on Ceramic Breeder Blanket Interaction, Rome, September 23-25, 1996.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 1997 Elsevier Science B.V.

Electrophoretic Forming of Functionally-Graded Barium/Strontium

221

Titanate Ceramics

Partho Sarkar Ceramic Research Group, Department of Materials Science and Engineering McMaster University, Hamilton, Ontario, Canada L8S4L7 Sumie Sakaguchi, Eiko Yonehara, Jun-ichi Hamagami, Kimihiro Yamashita and Takao Umegaki Department of Industrial Chemistry, Tokyo Metropolitan University Hachioji, Tokyo

Abstract Single phase BaTiOa (BT) and SrTiOs (ST) have sharp Curie Point at 393K and lOK respectively. ST forms solid solution with BT and shifts the Curie Point toward the lower temperature. In a multilayer sample varying the BT/ST ratio among the layers it would be possible to fabricate a fiinctionally-graded laminated dielectric composite which will exhibit a board transition temperature and as a result it will have a low temperature coefficient and high dielectric constant in a wide temperature region. Laminated BT/ST composites thick films and bulk samples were synthesized by electrophoretic deposition (EPD). Two types of solvents were tested for the EPD suspensions, alcohol/acetylacetone mixed solvent and ethanol. Microstructure of the sintered samples was characterized by optical and SEM. Dielectric properties as a fiinction of temperature were measured using an impedance analyzer. /.

Introduction EPD is an effective technique to synthesize monolithic as well laminates thick film [1-8] and bulk samples [9-13]. Sarkar and Nicholson [9,10] are thefirstto demonstrate the EPD can be successfiiUy use for fabrication of laminated and fiinctionally graded materials. There are few reports on fabrication of BaTiOj monolithic thick film by EPD [1,5,6]. Yamashita et al [7] is the only one used EPD to synthesize BaTiOj/SrTiOj laminated composite with varying BT/ST ratio and observed board Curie temperature response in the sample. EPD technique uses electrostatically stabilized suspension. It is a combination of two processes, electrophoresis and deposition. Electrophoresis is the motion of a charged particle in a suspension under the influence of an electric field. It was discovered around 250 years ago by the Indian Scientist, G.M. Bose. The Russian Scientist, Reuss, was the first to observe electric field-induced motion of clay particles in the water. The second process is deposition, i.e., the coagulation of particles into a dense mass. In the present study alcohol/acetylacetone mixed solvent is used for fabrication of thick films. It was found that mixed solvent is good for making thick film but the deposition voltage is high, the rate of deposition is low and also suspension stability is low. To overcome these problems a new suspension, i.e., titanates/ethanol system is being investigated.

222 //. Experimental Procedure Starting powders are BaTiOj (Kanto Chemical Co., Tokyo, Japan). SrTiOj powder used with mixed solvent is from High Purity Chemicals Laboratory. Since this SrTiOj has a large particle size (average size ~ 10//m), therefore for ethanol solvent system SrTiOa (average Size ~ l / / m ) form a different source (Transelco Division o f Ferro Corporation, ^ U S A ) is used. Suspensions are characterized 's by measuring their electrophoretic mobility "^ using Coulter DELSA 440SX and deposition ^ rate. Microstructure of the sintered samples g* is characterized by SEM and dielectric properties measured as a function of temperature using HP impedance meter. HI Results and Discussion 0 20 40 60 80 100 Figure 1 shows the deposited weight Acac PrOH Composition (v/o of Acac) o f BT, ST and 80w/o BT-ST mixed powder Fig1 Deposited weight of BT, ST and BT/ST mixture as in mixed solvent as a function of an a function of solvent composition of the suspension. alcohol/acetylacetone ratio. Depositions were conducted using constant voltage (400 V/cm) for 180 sec from a suspension o f 10 g/1. ST forms good deposit in pure alcohol but deposited weight decreases as the alcohol/acetylacetone ratio decreases in the solvent and no deposition occurs in pure acetylacetone. B T does not deposit either of the pure solvents but deposited weight increases departure from pure solvent and shows maximum deposition rate in 30-50 v/o acetylacetone. Mixed powder deposition behavior is similar to BT. Mixed powder deposition has some anomaly, it does not deposit in pure alcohol though ST has a maximum deposition rate. Pure alcohol ST particles positively charged. It is speculated here that in pure alcohol BT particles are slightly negatively charged and as a result of that the heterogeneous coagulation occurs between ST and BT. Since B T

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223 powder is finer (BT -l/zm, ST ~10//m), therefore BT form an envelope on ST particles BT/ST multilayered thick film. and as a result of this mixture powder behave like BT and form no deposit in pure alcohol. Figures 2a and 2b are the ^ ^ A • electrophoretic mobility of BT and ST powder in ethanol as a function of pH of the A ^ • • 1kHz ^ ^ A A suspensions. Both the materials below pH 9 *^ A A A 10 kHz show positive mobility ie. particles are • • • • • - • -100kHz positively charged. The BT and ST has a basic 100 150 surfaces. Since the absolute value of the Temperature ("C) mobility is higher at low pH (positive mobility) than at high pH (negative mobility), therefore it is decided that deposition will be conducted Fig. 3. Temperature dependance of capacitance of from a positively charged suspension. These Bi.ySJ FGM multilayered thick film. suspensions provide good deposition in the pH range between 5 and 4. The temperature dependence of the dielectric of the multilayer thick film deposited from titanate /mixed solvent (alcohol/acetylacetone) system is shown in Figure 3. This multilayer (5 layers) FGM sample was made using the suspension of composition 0, 30, 46, 63 and 100

Fig. 4. Cross-sectional SEM micrographs of multilayer BT/ST bulk sample: (a) low magnification is showing multiple BT(dense) and ST(porous) layers and (b) high magnification is showing the grains structure of dense BT layers.

v/ of BT. It is clearfromthisfigurethat there is no sharp Curie point and as a result of that the sample has a smaller temperature coeflScient than that of pure BT. Bulk multilayer BT/ST and as well as thick films were fabricated from tatinates/ethanol suspension (100 g/1). The deposition was conducted at pH'-4 to 5 at constant current. The sample was sintered at HOOT for 2 H in air with a 300 °C heating/ cooling rate. Figures 4 is a cross sectional micrograph of a poUshed and thermally etched (1300°C for IH) multilayer of 100% BT/100% ST sample. In Figure 4(a) porous layers are SrTiOj and dense layers are BaTiOj. The thickness of the ST

224

Fig. 5. Cross-sectional SEM micrographs of non-planner laminates. Sequentionally deposited from BT-ST/ethanol suspensions of composition 100, 75, 50 and 25 v/o of BT. (a) is a low magnification micrograph showing thefibre-electrodesposition and (b) is a high magnification micrograph showing the details of layers microstructure. layer is ~20//m and BT -SO/zm respectively. Figure 4(b) is the high magnification micrograph of the BT region showing high density of the region with average grain size - 18//m. Non planner laminates were also fabricated using Nicalon fibres as a depositing

Fig. 6. Cross-sectional SEM micrographs of non-planner laminates with 3-fibre electrodes. Each electrode has 37 layers, (a) is a low magnification micrograph showing the fibre-electrodes position and (b) is a high magnification micrograph showing the details of layers microstructure.

electrode. Four suspensions of composition 100 v/o BT, 75 v/o BT, 50 v/o BT and 25 v/o BT are used. There are deposited in the following sequences 100, 75, 50, 25, 50, 75, 100, 75.... etc. Figures 5a and 5b are the microstructure of a nonplanner laminates where fibre electrodes

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Fig. 7. Normalized capacitance as a function of c- © r^ • temperature of monolithic B„S.T samples ^'^ ^- J""" temperature of Ba,.Sr,TiO:'3 as a containing 85 (BT85), 95 (BT95) and 100 ^^'^^'^'^ of composition. (BTIOO) v/o of BT. are placed in a plane. In this sample each fibre has 31 layers. 25 v/o BT has the maximum porosity and lOOv/o BT has nearly no porosity. Between two porous layers (25 v/o) there are 5 layers (50, 75, 100 (thick), 75, 50 v/o BT). Figures 6a and 6b are the BT/ST FGM Thick Film on Pt 6000 microstructure where 3-fibres electrodes were used, two of them in a 5000 plane and one ofthem out of plane. In 'fe A IkHz 4000 this case each fibre has 37 layers. To determine the Curie 3000 + 10 kHz temperature, monolithic samples of the 42000 composition 75v/o, 80 v/o, 85 v/o, 90v/o, 95 v/o and 100 v/o BT were H 1000 made and their capacitances were . I • I • I • I • I • I • I • I • I • I • I • I • I • I • I 0 measured as a function of temperature. 100 120 140 160 180 60 80 20 40 F i g u r e 7 Temperature (C) is a plot of normalized capacitance of 85, 95 and 100 v/o BT. All these Fig. 9. Capacitance/Dielectric constant of a 160/^m multilayer thick fihn deposited from 100, 75, samples show sharp transition. Figure 50 and 25 v/o BT suspension. 8 is a plot of Curie temperature as a function of composition. 75 v/o sample shows a Curie point - 16°C. Figure 9 is a dielectric response of a multilayer thick fihn on a Pt substrate made from titanate/ethanol suspensions of composition 100, 75, 50 and 25 v/o BT. This sample shows broad transition temperature (80°120°C). Although pure BT has a transition temperature ~ 120T and next nearest transition temperature by 75 v/o BT(~ 16°C). This indicates inter-layer diffusion of the cation resuhed the broadening as well as shifting of the peak towards lower temperature. Dielectric constant at transition temperature is -5,000 in IkHz. This preliminary results indicate that by chosing appropiate suspension composition, individual layer thickness and sintering time and

I

226 temperature it would be possibe to control the position and broadening of the transition temperature. IV Conclusion Mukilayer thick films of BaTiOj, SrTiOj and their mixture are fabricated by EPD technique. It was demonstrated mixed solvent system is good for thick film deposition. Bulk and thick film multilayer samples were fabricated from titanates/ethanol system. Multilayer samples with varying composition shows board Curie temperature and as a result of that it has low temperature coefficient. References 1. V.A. Lamb and H.I. Salmon, "Electrophoretic Deposition of Barium Titanate" Am. Ceram. Soc. Bull. 41 (1962) 781-782. 2. P. Sarkar, S. Mathur, P.S. Nicholson and C.V. Stager, "Fabrication of Textured Bi-Sr-CaCu-0 Thick Film by Electrophoretic Deposition", J. Appl. Phys. 69 (1991) 1775-1777. 3. P. Sarkar and P.S. Nicholson, "Magnetically Enhanced Reaction Sintering of Textured Yba2Cu30x", Appl. Phys. Lett. 61 (1992) 492-494. 4. S. Sugiyama, A. Takagi and K. Tsuzuki, "(Pb, La)(ZrTi)02 Film by Multiple Electrophoretic Deposition/Sintering Processing", Jpn. J. Appl. Phys. 30 (1991) 21702173. 5. S. Okamura, T. Tsukamoto and N. Koura, "Fabrication of Ferroelectric BaTiOj Films by Electrophoretic Deposition", Jpn. J. Appl. Phys. 32 (1993) 4182-4185. 6. M. Nagai, K. Yamashita, T. Umegaki and Y. Takuma, "Electrophoretic Deposition of Ferroelectric Barium TiTanate Thick Films and Their Dielectric Properties" J. Am. Ceram. Soc. 76 (1993) 253-255. 7. K. Yamashita, E. Yonehara and T. Umegaki, "Dielectric Properties of Electrophoretically Layered Barium Strontium Titanate Films" K. Yamashita, E. Yonehara and T. Umegaki, to be published in the IEEE proc. 8. P. Sarkar, J. Hamagami, K. Sakaguchi, K. Yamashita and T. Umegaki,"Multilayers BaTiOj/SrTiOj Thick Films and Bulk Ceramics", proc.of the 16th Electronic Division Meeting by Electronic Division of Jpn. Ceram. Soc, pp.43-44 (1996). 9. P. Sarkar, X. Huang and P.S. Nicholson,"Structural Ceramic Microlaminates by Electrophoretic Deposition", J. Am. Ceram. Soc. 75 (1992) 2907-2909. 10. P. Sarkar, X. Huang and P.S. Nicholson,"Zirconia-Alumina Functionally-Gradiented Composites by Electrophoretic Deposition Techniques", J. Am. Ceram. Soc. 76 (1993) 1055-1056. 11. P.S. Nicholson, P.Sarkar and X. Huang,"Potentially Strong and Tough ZrOj-Based Ceramic Composites ^ 1300°C by Electrophoretic Deposition", Science and Technology of Zirconia V; Edited by S.P.S. Badwal, M.J. Bannistar and R.H.J. Hannik, (Technomic Publishing Company, Inc. (1993)) pp.503-516. 12. P. Sarkar, O. Prakash and P.S. Nicholson, "Micro-Laminate Ceramic/Ceramic Composites (YSZ/AI2O3)", Ceramic Engineering and Science Procedings 15 (1994) 1019-1027. 13. M. Whitehead, P. Sarkar and P.S. Nicholson,"Non-Planar AI2O3ATSZ Laminates by Electrophoretic Deposition using AI2O3 Fibre Electrodes", Ceramic Engineering and Science Procedings 15 (1994) 1110.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

227

Processing and Properties of Electrodeposited Functionally Graded Composite Coatings of Ni-Al-A^Oa K. Barmak^ S. W. Banovic^ H. M. Chan^ L. E. Friedersdorf, M. P. Ha^ner^ A. R. Marder\ C. M. Petronis^ D. G. Puerta^ and D. F. Susan^ ^ Materials Research Center, Lehigh University, 5. E. Packer Ave.,. Bethlehem, Pennsylvania 18015, USA ^ Energy Research Center, Lehigh University, 117 ATLSS Dr., Bethlehem, Pennsylvania 18015, USA

1. ABSTRACT Single- and dual-particle, uniform and graded composite coatings of Ni-Al-AUOs, with Ni as the matrix and primarily Al and AI2O3 as second phase particles, were fabricated via electrodeposition. During the electrodeposition process, Ni was plated from an electrolytic bath to which the particles had been added. For single particle baths, the codeposition of AI2O3 was more strongly affected by current density and bath particle content than was the codeposition of Al. In the mixed particle bath, codeposition of AI2O3 was suppressed at low current densities, whereas codeposition of Al was not affected at any of the current densities studied. When coatings containing Al were annealed, the reaction of the two elements resulted in the formation of either single phase y solid solution or two phase y-y', in agreement with the equilibrium phase diagram. In addition to current density and bath particle content, the shape and size of the particles was found to affect the electrodeposition process. Here, angular nickel aluminide particles resulted in porous coatings. The microhardness of the Ni-Al and NiAI2O3 coatings showed a complex behavior with volume fraction of particles as a consequence of the effect of these particles on the microstructure of the Ni matrix.

2. INTRODUCTION The introduction of gradients of chemical composition, phase distribution or microstructure represents a now fervently pursued concept in the design of engineering components for optimum performance [1]. The occurrence of the Fourth International Symposium on Functionally Graded Materials (FGMs) bears witness to this fact. FGMs can be made by a variety of techniques - see, for example, the proceedings of the

228

previous international symposium on these materials [1]. The fabrication of graded coatings by electrochemical methods has the advantages of versatility, ability to coat complex shapes, and low cost. In addition, particles that are added to the electrolytic bath codeposit with the electroplated metal, thus creating a metal matrix particulate composite. The volume fraction of codeposited particulates depends on many parameters. These include the nature of the electrolytic bath, the current density and the type (metaUic, ceramic, etc.), size, shape and amount of particles. Here we report the effect of a variety of these parameters on processing and the resultant microhardness of uniform and graded, single- and dual-particle composite coatings of Ni, with primarily Al and AI2O3 as the second-phase particles.

3, EXPERIMENTAL The details of our electrochemical deposition process are given elsewhere [2]. Here we give a brief summary. The substrate for all the coatings was pure Ni 200 plate. The electrodeposition bath was sulfamate type and contained 400 g/l of nickel sulfamate tetrahydrate, 30 g/l boric acid, 5 g/l nickel chloride hexahydrate, 0.5 g4 sodium lauryl sulfate and 0.1 g/l Coumarin. For most experiments a-alumina powder, with an average size of 0.60.8 )Lim, and Al powder, with majority of the particles in the 1-4 |im size range, were added to the electrolytic bath for codeposition with Ni. In some experiments large nickel aluminide particles (with average size 7.1 fim) were used to study the effect of particle shape and size on the codeposition process. The bath was sonicated prior to the deposition run and was mechanically agitated during the run. The starting pH of the electrolyte was 4 ± 0.2 and the bath temperature was maintained at 50 ± 2°C. The volume percent of the incorporated particles was measured through quantitative image analysis of the pohshed cross sections on a LECO 2001 image analyzer. The microhardness of the as-deposited and annealed coatings were measured using a LECO-M 400FT microhardness tester in accordance with ASTM E384 standard using loads of 25 g on polished cross sections.

4. RESULTS AND DISCUSSION Figure 1 presents scaiming electron micrographs of the second phase particles used for producing the coatings on the left and optical micrographs of the cross-section of the resultant coatings on the right. As can be seen in the figure, the smaller particles resulted in dense coatings, while the large, angular nickel aluminide particles gave a porous coating. Thus, at given deposition conditions and electrochemical cell set-up, the shape and size of the particles Fig.l - (Opposite page) Second phase particles used for producing the composite coatings and the resultant coatings. Scanning electron micrographs of (a) alumina, (c) aluminum and (e) nickel aluminide particles. Optical micrographs of the cross sections of the resultant coatings (b) Ni-alumina with 225 g/l alumina in bath, (d) Ni-Al with 225 g/l Al in bath and (f) Ni-nickel aluminide with 200 g/l aluminide in bath. All coatings were deposited at 5 A/dm^.

229

1^ ^t,»M

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230

can have a significant impact on the quaUty of the coating, with dense coatings being obtained for "rounder", smaller particles, particularly when the particles are metallic. Figure 2a is a plot of volume percent alumina in Ni-Al203 coatings as a function of volume percent in the bath for a range of current densities. It can be seen that increasing the amount of alumina in the bath resulted in a steep increase in vol.% alumina in the coating. A maximum of 39 vol.% alumina in Ni was achieved at 1 A/ dm^ for a bath loading of 5.3 vol.%. This value is almost twice the maximum reported by Ding et al. [3] for Ni electrodeposits that contained 2.7 jim a-alumina particles. The decrease in volume percent as a function of current density for our coatings followed a similar trend to that found by Ding et al. [3] for Ni-a alumina and Cua alumina and Celis et al. [4] for Cu-y alumina deposits. Compared with AI2O3, the codeposition of Al was less strongly affected by current density and bath particle content. The amount of Al in the coating ranged from 5 to 17.5% only, as can be seen in Fig. 2b. Comparison of the structure of the Ni matrix at high current densities (> 10 A/dm^) with and without the presence of Al showed that codeposition of Al resulted in refinement of the structure. Comparable studies in Ni-Al203 coatings were not possible, because the etchant preferentially attacked the interface [2]. However, from the variation of hardness of the coatings (see below) with particle vol.% we believe that at lower current densities the alumina particles result in a coarsening of the Ni grain structure. Note that grain structure of Ni for sulfamate bath becomes finer with decreasing current density. [2] In mixed particle baths, the codeposition of Al was not affected by the presence of alumina whereas the codeposition of alumina was suppressed at lower current densities so that the two lines shown in Fig. 2a for low and high current density regimes collapsed onto the latter. The reason for this behavior, we believe, is related to the distortion of the field lines around metallic (conducting) versus ceramic (insulating) particles during deposition. Graded coatings of Ni-Al-Al203 were produced by varying the bath particle content at a fixed current density of 5 A/dm^. A light optical micrograph of one such coating is shown in Fig. 3a and the annealed structure of the same coating in Fig. 3b. When coatings containing Al

2

4

6

8

Vol.% Alumina In Bath

10

2

4

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Vol.% Aluminum In Bath

Fig. 2 - Volume percent (a) AI2O3 (b) Al in coatings as a fiinction of particle volume percent in the bath for a series of current densities in the range 1 - 25 A/dm^.

231

Fig. 3 - Light optical micrographs of a graded Ni-Al-Al203 coating, (a) From bottom to top, the three layers in the coating contain approximately 12-0, 9-16, 9-24 Al vol.%-Al203 vol.%., respectively, (b) Same coating as in (a) annealed at 635° C for 1 hr.

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232 were annealed, the reaction of the two elements resuhed in the formation of either single phase Y sohd solution or two phase y-Y, in agreement with the equilibrium phase diagram. The sample in Fig. 3b, annealed for 1 hr at 635 "C, was a two phase mixture of y-Y, with the alumina particles residing in the y phase. Figures 4a and 4b present the hardness of single-particle, uniform coatings of respectively Ni-Al203 and Ni-Al as a function of current density. Two points are worth noting. One, at high current densities ( > 10 A/dm^) the "soft" metalhc Al particles resulted in greater hardening than the "hard" ceramic AI2O3 particles. Two, at lower current densities that the incorporation of alumina resulted in a smaller increase in hardness than at higher current densities, even though the volume percent of the incorporated particles was larger at the lower current densities. In other words, for both types of coatings the hardness did not follow a simple rule of mixtures. The reason for this, we believe, is the change in microstructure of the Ni matrix when the second phase particles are incorporated.

5. CONCLUSIONS Electrodeposition offers an attractive method for the fabrication of uniform and graded particulate composite coatings. Here we demonstrated the use of this method in the fabrication of single- and dual-particle Ni-Al-Al203 coatings. In doing so we delineated the effect of current density and bath particle content on the volume percent of particles in the resultant coating. The hardness of the single-particle coatings shows significant deviations from the rule of mixtures. This was believed to be a result of the effect of the particles on the microstructure of the matrix during deposition. This work was made possible by research subcontract DE-FC21-92MC29061 sponsored by the U.S. Department of Energy - Morgantown Energy Technology Center through a cooperative agreement with the South Carolina Energy Research and Development Center at Clemson University. The authors thank A.O. Benscoter for help with metallography and microscopy, and X. M. Ding and Profs. B. Ilschner and R. Chaim for helpful discussions.

REFERENCES 1. B. Ilschner and N. Cheradi FGM '94, Proc, of the 3rd international symposium on structural andfunctional gradient materials, (Ed. by B. Ilschner and N. Cheradi, Presses polytechnique et universitaires romandes) pp. V-IX (1995). 2. K. Barmak, S.W. Banovic, C. M. Petronis, D. F. Susan and A. R. Marder, J. Micros., (in press). 3. X. M. Ding, N. Merk and B. Ilschner, FGM *94, Proc, of the 3rd international symposium on structural and functional gradient materials, Ed. by B. Ilschner and N. Cheradi, Presses polytechnique et universitaires romandes, 365 (1995). 4. J. P. CeHs, J. R. Roos, and C. Buelens, C, J. Electrochem. Soc. 134, 1402 (1987).

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

233

Functionally graded materials by electrochemical modification of porous preforms A. Neubrand, R. Jedamzik, and J. Rodel Fachbereich Materialwissenschaft, Technische Hochschule Darmstadt, PetersenstraBe 23, 64287 Darmstadt, Germany A novel method to produce gradient materials based on the infiltration of refractory porous preforms with a molten metal or polymer has been developed. The porosity gradient in the preform is created by electrolysis. For this purpose, a gradient of the electrochemical potential is set up inside the porous preform, leading to a gradient in the rate of electrochemical dissolution or deposition of the preform material and thus to a graded porosity. A macroscopic electrokinetic model of the gradation process is developed, and the influence of experimental parameters like current density, electrode and electrolyte resistivity and geometrical factors on the gradation profiles are discussed and compared to experimental observations. W/Cu graded materials have been produced and some of their properties have been determined as a function of position.

1. EVTRODUCTION In the last decade, a variety of processing methods for functionally graded materials including powder processing, thermal spraying and deposition processes have been developed. Infiltration techniques are a very promising method for material combinations with very different melting points. A preform of the more refractory phase possessing a porosity gradient is produced and infiltrated with the melt of the lower melting component at elevated temperatures. The main advantages of the method are the low porosity levels that can be achieved without the need of developing sophisticated densification techniques like sparkplasma or microwave sintering, which have to be adapted to the material combination and gradient. The main challenge lies in the production of suitable preforms in a one-step process. An ideal preform should have a large porosity gradient with minimum closed porosity and sufficient mechanical stability. The lower porosity limit for a preform with equiaxed phase elements is about 8% (below this limit most porosity will be closed). The upper porosity limit is given by the need of mechanical stability of the preform - using fibres a porosity as high as 95% is achievable. Thus, infiltration processing can produce gradient materials with smooth gradients from 92% to 5% volume fraction of the more refractory phase. A variety of methods has been used to prepare graded preforms. For example sintering of powder compacts with a graded grain size leads to the formation of a porosity gradient. However, in tungsten preforms large deformations due to uneven sintering shrinkage have been observed unless different layers of the compact were pressed separately to different green densities [1]. Porosity gradients can also be introduced in the compact by admixing variable

234

amounts of non-equiaxed elements. As an example, Al2Ti05/Al203 preforms with graded porosity have been produced by sequential casting of slips with different content of short alumina fibres [2]. Plasma spraying has been employed for the production of graded AI2O3 preforms [3]. The method yields materials with some closed porosity and only moderate gradients in open porosity of 5-14%. Another method uses open-celled polymer foams as precursor materials. The polymer foam is coated with a ceramic slip while rotating in a centrifuge. Afler burning out the polymer a ceramic foam with a porosity gradient is retained [4]. In the present study an electrochemical process is used to introduce a porosity gradient into a homogeneous preform. The method is able to produce a large variety of gradient materials in a simple two-step process consisting of gradation and infiltration thus making it attractive for large-scale industrial processes.

2. THEORETICAL BACKGROUND Electrochemical gradation of porous preforms is based on an electrochemical reaction taking place in a porous electrode where the solid phase of the electrode takes part in the electrochemical reaction - it is either dissolved or deposited during the process. The apparatus for the gradation process is sketched in Fig. 1.

Galvanostat

Referenceelectrode

Electrolyte Porous preform

Insulation

Fig. 1 Experimental setup for the electrochemical gradation process

235

For tungsten as porous anode material the anodic half reaction is W + 80H-

-^

w o / " + 4H2O + 6e'

During the reaction current passes through the pores of the electrode. According to Ohm's law this leads to a potential gradient in the electrolyte. In simple cases, the current density j depends on the overpotential rj (the difference between actual potential and equilibrium potential) according to the Tafel relation j=j^J^

(1)

For the anodic dissolution of tungsten in alkaline solution Jo = 5.8*10"^ A/m^ and k= 38.4 V "^ [5]. This means that a change of 60 mV in overpotential causes the reaction rate to increase by a factor of 10. Such potential differences occur in the electrolyte within a few millimetres distance at current densities of a few mA/cm^, which are easy to realise for the electrode reaction mentioned above. Using the Tafel relation, the distribution of the current (and thus reaction rate) inside the anode at the beginning of the experiment can be calculated analytically [6]. In the present paper a refined numerical model was used for the calculation of the dissolution rate which can use experimentally determined current-potential-relations and can take into account the change of pore radius during the dissolution process. The model shows that the current distribution and thus the gradation profile is controlled by a number of experimental parameters like initial porosity, current density, conductivity of electrolyte and electrode material, kinetics of the electrode reaction and temperature. By changing these and the geometry of the experiment, a variety of different gradation profiles can be produced.

3. EXPERIMENTAL Tungsten preforms of 16 % and 42 % porosity containing about 1% nickel (Tridelta AG, Hermsdorf, Germany) were used for the gradation experiments. The area of the preforms was 30 X 24 mm^, the thickness was either 5.2 mm or 5.8 mm. The preforms were first weighed and then graded by using them as anode in an electrolytic cell. For this purpose the preform was infiltrated in vacuum with the electrolyte and contacted on the back. The platinum cathode was mounted parallel to the anode surface at a short distance. It had the same area as the anode giving rise to a nearly uniform electrolyte potential at the front surface of the anode. The cell current was kept constant during experiments by means of a galvanostat. The electrolyte was commercial grade NaOH. Under the experimental conditions tungsten was dissolved at the anode while hydrogen was generated at the cathode. After completion of the experiment the tungsten anodes were flushed several times with water in order to remove sodium hydroxide and reaction products, dried in a drying chamber and weighed. The resulting graded preforms were infiltrated in vacuum with either molten copper or the molten alloy CuNi2Si. An overview of the processing sequence is shown in Fig. 2. The volume content of the different phases was determined from optical micrographs using an image analyser. EDX analysis of the elemental composition was performed on some of the samples. Hardness tests were carried out using a Vickers indenter using a load of 49N.

236

anodic dissolution >

porous tungsten

infiltration with copper

tungsten with porosity gradient

tungsten/copper gradient material

Fig. 2 Overview of the processing sequence 4. RESULTS A porous tungsten sample (A) was electrochemically graded under experimental conditions given in Table 1. Its microstructure after copper infiltration is shown in Fig. 3. As a result of the electrochemical gradation process a one-dimensional gradient in tungsten content (dark areas) from the side away from the cathode (left) to the side close to the cathode (right) is observed. This is in accordance with theory which predicts that the electrochemical reaction will always be faster on the side close to the counter electrode if the electrode material has a higher conductivity than the electrolyte. The gradient is continuous and accompanied by an increase in the size of the copper ligaments (light areas). A slight step is only observed at the former outer surface of the tungsten preform on the right of the micrograph. No signs of porosity are observed even at higher magnifications.

Fig. 3 Microstructure of W/Cu FGM

237

Table 1 JExgerimental conditions for Sample Initial Porosity A 16% B 42%

electrochemical gradation Electrolyte Duration of Anodisation IMNaOH 23 5h 2MNaOH 23 5h

Current density 6.25mA/cm2 6.25mA/cm2

Current yield 103 % 103 %

The chemical composition of the tungsten-copper FGM as a function of position as determined by optical micrography is shown in Fig.4a. From weight loss data a current yield close to 100% was determined for the experiments - showing that there were very little side reactions during the electrochemical dissolution of tungsten. Modelling the gradient is therefore straightforward and a theoretical prediction of the composition profile is in good agreement with the observed one, if one assumes that the BET surface remains constant during the dissolution process. In reality one would expect that the BET surface first increases and then decreases with material removal. It appears that the assumption of no change in time average still yields good results. The gradient in composition is also reflected in a change of Vickers hardness. However, close examination shows that the tungsten content shows little variation on the side away from the cathode, whereas the hardness decreases significantly. The reason for this behaviour is yet unclear, but macrostresses at the tungsten-rich side of the graded interface may be responsible for the apparent hardness increase. In general, the observed hardness values compare favourably with ungraded W/Cu composites (possessing a Vickers hardness of 2.8 GPa at 85% tungsten content) indicating a low porosity level of the produced FGM. Theory predicts shallower composition gradients for electrolytes with higher conductivity. This has been demonstrated by grading a second sample (B) under different conditions - in this case the composition profile (determined by optical microscopy and EDX) extends over several millimetres (Fig.4b). Again the composition gradient is well predicted apart from a small increase in tungsten content close to the surface which existed already in the untreated preform. The gradient is again reflected by the Vickers hardness. As can be seen from Table 1 the only parameters changed were the initial porosity of the sample and the concentration of the electrolyte. The increased porosity and electrolyte conductivity leads to smaller potential a)

b) 60

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Cathode -

238 differences in the liquid phase when the cell is operated, which in turn leads to shallower gradients. 5. CONCLUSION Electrochemical gradation of porous preforms is a versatile and cost-effective method to produce gradient materials. It is suitable for the production of gradient materials from two components with very different melting points. The method yields graded tungsten-copper composites of high density in a simple two-step process. The electrochemical gradation process of tungsten is adequately described by a macroscopic model of the electrode kinetics.

ACKNOWLEDGEMENT This work was supported by the German Research Society (DFG) as part of the „Schwerpunktprogramm Gradientenwerkstoffe". We also thank Tridelta AG for kindly supplying the tungsten preforms.

REFERENCES [1] M. Takahashi, Y. Itoh, M. Miyazaki, H. Takano, and T. Okuhata, p. 17-28 in Proceedings of the 13 th International Plansee Seminar, H. Bildstein and R. Eck (eds.),Metallwerk Plansee, Reutte 1993 [2] W. Henning, C. Melzer and S. Mielke, Metall 46 (1992) 436-439 [3] W. Schultze, S. Schindler, F.-U. Deisenroth, German patent DD 300 725 A5 (1990) [4] Y. Miura, H. Yoshida, Y. Takeuchi, K. Ito, German patent DE 3527 872 Al (1986) [5] J.W. Johnson and C.L. Wu, J. Electrochem. Soc. 118 (1971) 1909-1912 [6] A. Neubrand, B. Kastening and J. Rodel, p.488-497 in Elektrochemie der Elektronenleiter, GDChMonographienBd.3, F. Beck(Hrsg), GDCh-Verlag, Frankfurt 1996

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

239

Thermal management of carbon-carbon composites by functionally graded fiber arrangement technique Y. Kude and Y. Sohda Central Technical Research Laboratory, Nippon Oil Company, Ltd., 8, Chidori-cho, Naka-ku, Yokohama, 231, Japan Next to mechanical properties, the most important characteristics of a carboncarbon composite(C/C) are thermal conduction and thermal expansion. In this paper, several investigations have been made into carbon fiber arrangement relationships for different carbon-carbon composite materials. Pitch-derived carbon matrix-carbon fiber composites have been used, processed by means of the hot isostatic pressing (HIP) technique for converting pitch into a dry carbon fiber preform. Repeated HIP cycles are required to build the composite matrix up to an acceptably high density/low porosity for deployment in severely ablative environments. The effects of heat treatment temperatures on thermal conductivity, thermal conductivity at high temperatures and thermal expansion behavior have been studied. At room temperature, the value of thermal conductivity for unidirectional (UD) carbon-carbon composites is 700 W/m-K. In the case of three-dimensional (3D) carbon-carbon composites, this value is determined by the volume of the fiber arrangements. On the other hand, the thermal expansion of carbon-carbon composites in the fiber axial direction is chiefly governed by the thermal expansion of the fiber. On the basis of this fundamental research, a functional graded fiber arrangement technique has been proposed which presents the opportunity to 'tailor' thermophysical properties into carbon materials. 1. INTRODUCTION A lot of investigations have been made in the field of high performance carbon/carbon composites for aerospace applications. Most of these composites w^ere associated with their superior mechanical properties at high temperatures. In addition, thermal conduction and thermal expansion of carbon-carbon composites are also important for some applications, e.g., re-entry vehicles or aero-engine components. On the other hand, it is expected that the positive thermal management technology of carbon/carbon composites will provide excellent heat receivers or heat radiators, especially in fields that require severe thermal shock resistance. In this paper, first the thermal properties of carbon fibers are described, and second the thermal properties of carbon-carbon composites are presented. Finally, the application to sunshine heat receivers by a functionally graded fiber arrangement technique is also described. One of the important aims of this research is to prove that carbon-carbon's

240 properties are capable of being tailored to specific applications by this functionally graded fiber arrangement technique. 2. THERMAL PROPERTIES OF PITCH-BASED CARBON FIBERS Graphite shows very anisotropic thermal properties as a result of its crystal structure. This is because graphite possesses two-dimensional hexagonal network structures and the layers are held together very loosely by weak forces. For example, chemical vapor deposited carbon, which is manufactured with heat treatment at 3000°C after the deposition and possesses almost ideal graphite structures, has a thermal conductivity of 2,000 W/m'K (at room temperature) parallel to the layers and 10 W/m-K in the perpendicular direction as shown in figure 1. This value in the parallel direction is approximately 4-5 times more than the value of silver or copper, which are typical high thermal conductivity metals. Carbon fibers, which utilize the preferred orientation of the graphene layers, show not only high modulus and high strength but also high thermal conduction and low thermal expansion along the fiber axis. On the basis of these properties, fiber reinforced materials present the opportunity to design the thermal properties into materials. Figure 2 shows the thermal conductivity of pitch-based carbon fibers for each modulus grade (value [XN-**] means its tensile modulus) in comparison with other high thermal conductivity materials. High modulus carbon fibers, that is, carbon fibers which have high degrees of preferred orientation of the graphene layers, show high thermal conductivity. Those values are more than the value of silver as a typical high thermal conductivity metal or silicon nitride as a typical high thermal conductivity ceramic. Incidentally, pitch-based carbon fibers show much higher thermal conductivity than PAN-based carbon fibers with the same modulus. Carbon fibers show negative thermal expansion behavior at temperatures between 20°C and around 500°C as shown in figure 4. This behavior depends on each fiber's grade. By utilization of this negative behavior, materials whose coefficient of thermal expansion is zero can be created when quasi-isotropic laminates are controlled with an optimum fiber volume fraction and are incorporated with matrices which have positive coefficient of thermal expansion. In practice, those materials can be used in satellites or space telescopes which demand severe thermal environment resistance. 3. FABRICATION METHOD OF CARBON/CARBON COMPOSITES In this work, all carbon-carbon composites were fabricated by using pitch-based carbon fibers and pitch-derived carbon matrices. Two types of preforms, were mainly used. One is unidirectional (UD) preforms, and the other is 3-dimensional (3D) fabrics. 3D fabrics normally contain 55 vol.% of fibers that are introduced in three directions. Each volume fraction of fibers is 40 vol.% in the x-direction, 10 vol.% in

241 the y-direction, and 5 vol.% in the z-direction. After matrix impregnation, the pitch was converted to carbon by a process of pyrolysis in high isostatic inert gas pressure. This densification process was repeated several times. After densification processing, the matrix was graphitized by controlled heating to temperatures above 1700°C. 4. THERMAL PROPERTIES OF CARBON/CARBON COMPOSITES 4.1. Thermal conductivity Thermal conductivity was evaluated by the product of the density, specific heat, and thermal diffusivity. Thermal diffusivity was measured by the Laser Flash method. Figure 5 shows the thermal conductivity of UD-C/C compared with that of 2D-C/C. At room temperature, the thermal conductivity value is 700 W/m-K parallel to the fibers in the case of UD-C/C. The values of conductivity rise significantly as the final heat treatment temperature increases. This is not only because the size of the graphite crystallites became bigger, but also because the orientation of the crystallites along the fibers increased. UD-C/C possesses roughly twice as great a value of thermal conductivity as 2D-C/C. This result also proves that the matrices in UD-C/C greatly contribute to the thermal conductivity because the orientation of the matrices' crystallites along fibers increased. Figure 6 shows the temperature dependence of the thermal conductivity for UDC/C in comparison with copper. As the temperature is increased, the thermal conductivity falls. This is because the amount of phonon scatter is due to the vibration of the crystal lattice. The thermal conductivity data of 3D-C/C are presented in figure 7 correlated with the fiber volume fraction for each direction. The direction with 40 vol% of the fiber has a thermal conductivity of 470 W/m-K. On the basis of the good relationships between the fiber volume and the thermal conductivity, it is considered that the design and the regulation of the heat flow in C/C would be possible. In this case, there are two noteworthy points for the design as can be seen from 3D-C/C type A and type B in Figure 7. First the density of C/C greatly contributes to the thermal conductivity. Second, the fiber volume valance influences the contribution of the matrices to the thermal conductivity more than the actual fiber volume dose. In other words, the more isotropic the fiber volume valance in C/C is, the more dependent on the fiber itself in C/C the thermal conductivity is. 4.2. Thermal expansion Thermal expansion was measured by the laser light interference method. Figure 8 shows thermal expansion behavior of UD-C/C with the heat treatment at 3000°C. There are large differences in the thermal expansion between the fiber axial direction and the fiber vertical direction. Figure 9 shows the relation between the coefficient of thermal expansion and the

242 measurement temperature in UD-C/C. The value of the coefficient in the fiber axial direction is almost zero. In contrast, the value of the coefficient in the fiber vertical direction is close to the value of the coefficient in copper. This fact means that the joining between copper and C/C is easy when limited to this direction. The thermal expansion in the fiber axial direction is chiefly governed by the thermal expansion of the fiber, and the thermal expansion in the fiber vertical direction is chiefly governed by the thermal expansion of the matrix. Figure 10 shows the coefficient of thermal expansion in 3D-C/C. There is a small difference for each direction, but all of the values are almost zero. Therefore, the joining between copper and C/C is expected to be quite difficult in the case of 3DC/C. 4.3. Application to sunshine heat receivers of functionally graded fiber arrangement technique As mentioned previously, the heat flow through C/C composites by conduction can be controlled by the design of the fiber architecture. The most suitable design method is by using the functionally graded fiber arrangement technique. Figure 11 shows a C/C composite cavity manufactured with the functionally graded fiber arrangement technique for a sunshine heat receiver. This C/C cavity is manufactured from 3D carbon fiber fabrics which are designed for getting the optimum heat flow. Solar rays are irradiated by a large parabolic mirror into the inner surface of this C/C cavity. Since the heat is utilized from the back bottom wall, efficient heat flow to the back bottom wall is required. At the side, the fiber volume of the inner side in the bottom direction is larger than the outer side. At the bottom, the fiber volume of the center to the back bottom direction is larger than the outer. The fiber volume of each intermediate part is graded. This cavity was tested by the solar ray concentration equipment at Tohoku University. The results are shown in figure 12. C/C composites made with the FGM method have higher performance than isotropic graphite. 5. SUMMARY Carbon-carbon composites made with the functionally graded fiber arrangement technique present the opportunity to 'tailor' thermo-physical properties into carbon materials. In this paper, the changing of the fiber architecture is the method for FGM. Fibers or matrices are other options for FGM. This functionally graded fiber arrangement technique can be applied to a wide range of materials processing. ACKNOWLEDGEMENT The authors would wish to thank Prof. Arashi and Dr. Naitou of Tohoku University for evaluating the C/C composite cavity.

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Figure 1 Difference in thermal conductivity of graphite between a-axis and c-axis.

Figure 2 Thermal conductivity of carbon fibers and various materials.

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Figure 5 Thermal conductivity of C/C composites at room temperature.

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Figure 6 Thermal conductivity of UDC/C composites as a function of temperature.

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1900 1800 1700 1600

Isotropic graphite

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1500 1400

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Figure 11 C/C composite cavity manufactured by ftinctionally graded fiber arrangement technique for sunshine heat receiver.

Figure 12 Temperatures of the back wall in a FGM C/C cavity heated by solar ray concentration equipment.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

245

Formation and Properties of TiC/Mo FGM Coatings T. Fukushima ' \ S . Kuroda ' \ S . Kitahaia '\K.

Ishida '' and M. Sano ''

1) National Research Institute for Metals, Tsukuba, Ibaraki, Japan 2) National Aerospace Laboratoiy, Chofii, Tokyo, Japan ABSTRACT Tie and Mo powders were sprayed by the plasma twin-torch process to form an FGM layer over a substrate to be used as an emitter in a thermoionic energy converter. The sprayed TiC coatings showed high absorptance for solar radiation. As compared to Tie monolayer coating and TiC/Mo two-layer coating formed on Mo substrate, FGM coatings performed better in a thermal cycling test. 1. INTRODUCTION Thermionic energy generator to utilize condensed solar radiation is now being developed through the FGM national project in Japan. One of the components needed for the generator is the emitter electrode with high solar absorptance. Since the emitter material is usually a refractory metal such as Mo, its surface needs to be coated with a material with high absorptance. We have selected TiC for its high absorptance and high melting point ^^ since the operating temperature of the emitter is expected to be 2000K. In order to reduce the thermal stress in the coating, possibillity to form graded TiC/Mo coatings by thermal spray has been investigated. TiC and Mo powder ( particle size : 10 to 45 (uim ) were sprayed onto Mo substrates to obtain coatings with high absorptance. Furthermore, the laminated structure, the thermal stability, the thermal resistance and the durability of the obtained coatings were investigated. 2. EXPERIMENTAL PROCEDURES TiC/Mo FGM coatings were sprayed onto Mo substrate by the plasma twin torch spraying system ^^ shown in Figure 1. In this system, each plasma torch to spray | Mo Ti C powder particles of Mo metal or TiC ceramics was operated at the same time under spraying conditions shown in Table 1. \ Torch 1 \ / Torch 2 / These conditio n were searched to \ ^"^'^^^ C^^"^'^ / obtain higher deposition efficiency of the two materials. When forming FGM L:Spraying disrance coatings, the powder feeding rate to the 8 :Spraylng angle two torches were controlled in a stepwise manner. t T-T r.-n. =. r-^r -^. ^ - = - = - - -r_ •=-.-:^T =. r:r-jvr_j oproy ie coating } } Substrc ite Thermal radiative properties of the obtained coatings in the visible wave Fig. 1 Scheme of tliemial spraying by twin torch

246 length (350 to 700 nm) were measured by spectrophotometer. The laminated structure of the coatings was observed by SEM, and their thermal stability and thermal resistance were examined by heat treatment and repeated thermal cycles, respectively.

Table 1 Spraying conditions

Operating current(A) Operating gas flow rate (l/min)(Ar) Spraying angle (° ) Spraying distance(inin) Substrate

Spraying materials Mo TiC (10 to 53;t^in) (10 to 44/zm) 800 1200

45 70 100

3. RESULTS AND DISCUSSIONS Mo (3 mmt) 3.1. Formation of TiC coatings Operating conditions to spray TiC ceramic with the high melting point ( 3373K ) were searched. Since TiC coatings were thermal sprayed in air, a slight amount of oxide such as TiO, Ti02 were produced in the coatings. The formation of oxides was suppressed by introduction of hydrogen gas into the argon plasma gas. However, the oxides in the TiC coatings formed under operating conditions shown in Table 1 had no influence on energy absorptance. Spectral absorptance of TiC coatings coatings was determined from the ratio of reflected radiation to the irradiatedradiation by the spectro-photometer.

O

Br O c/j

<

0

4

5 6 Wave length (xlO 'nm)

Fig. 2 Absorptance of TiC coatings Figure 2 showed the spectral absorptance of TiC coatings sprayed with two kinds of powder particles of different size (A powder : 10 to 44fxm, B powder : about 3fxm ). It was found that the spectral absorptance of the coatings with A powder indicated the value of 89 to 92 % and it was higher than that with B powder (80 to 88 %). From these results, it was suggested that deposition particle size affected surface roughness of the coatings. In measured results of surface roughness, mean roughness (Ra) was 41.4pLm in TiC coatings with A powder and was 19.4^Lm with B powder.

247 Furthermore, the temperature dependence of emissivities of TiC coatings and Mo substrate without coatings were measured in Ar atomsphere by using a sample-moving type emissiometer ^^ . Figure 3 shows measured emissivity in the temperature range from 1073 to 2073 K. In this figure the emissivity was expressed as the intensity ratio of emission at materials to emission at black body. The emissivity of TiC coatings decreased gradually with an increase in temperature, however, it maintained higher values (0.72 to 0.87) than that of Mo substrate (0.5) in the experimental temperature range. 1.

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20

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(xlO'K)

Fig. 3 Emissivity of TiC coating and Mo substrate 3.2 Formation of TiC/Mo FGM coatings Both TiC and Mo powder feeding rates were controlled for the formation of TiC/Mo composition graded coatings. The macrostructure of cross section of TiC/Mo FGM (coating thickness : 500 '^ 600^m) sprayed under those operating conditions is shown in Fig. 4. To compare with the characteristics of TiC/Mo FGM coatings, TiC coatings Fig. 4 Structre of TiC/Mo FGM coating (coating thickness : 320 ~ 330juLm ) and TiC-Mo two layer coatings ( coating thickness :320 ~ 330fxm) were formed onto Mo substrate. 3.2.1. Thermal stability and thermal resistance of coatings TiC, TiC-Mo two layer and TiC/Mo FGM coatings were heat-treated under conditions of 1473K, 16h and 72h, in a vacuum (2 X 10 "^ Pa) .

248 Firstly, figure 5 shows the typical cross sections of TiC and TiC-Mo two layer coatings observed by SEM. As shown in this figure, some cracks were found in the heat-treated TiC coatings and the TiC layer peeled from Mo substrate after the heat-treatment.. In the heat-treated two layer coatings, there were no cracks in both TiC and Mo layers and the interface between TiC and Mo layers was sound. These results were caused probably by thickness difference of TiC layers between TiC coatings and combination coatings as shown in Fig. 5.

Fig. 5 Structures of coatings by heat treatment Secondly, figure 6 shows the distributions of Ti, C, O and Mo near the interface of the TiC and Mo layers in the two layer coatings for the as sprayed and the heat-treated coatings. Elements Ti, C and Mo did not diffuse to the other layer. From the results and observations of micro structures, it was found that TiC and Mo showed no chemical reaction in the two layer and FGM coatings. Horiguchi et al. carried out diffusion bonding of TiC with Mo and reported a kinetic expression for the diffusion between the two materials ^^ . Extrapolating their results to the present conditions of heat treatment, it is expected that a diffusion layer of at least 5 jU m thickness is formed at the TiC/Mo interface possibly resulting in a reaction product such as Mo 2 C. Lack of diffusion or reaction layer between the TiC and Mo in the sprayed specimens may be because sprayed Mo particles were covered with thin oxide film . Further study is necessary to confirm this hypothesis, however. Finally, figure 7 shows the scheme of a solar heating test carried out by Prof Arashi's laboratory at Tohoku Univ. In this test, the specimen was vacuum-encapsulated in a quartz cell and the coating surface was heated by a condensed solar radiation. The surface of FGM coatings reached 1775K where as the rear surface reached 1725K during the test. No micro-crack and peeling between the coatings and Mo substrate were observed after the test.

249 3.2.2. Thermal cycle resistance of coatings Thermal cycle resistance of Tie, TiC-Mo and TiC/Mo FGM coatings were investigated by the thermal cycling test of temperature range between room temperature and 1223K in 5 times ( heating, cooling speeds were 100 K/s ) in air . These experimental results were summarized in Table 2. The cracks in TiC layer and peeling between TiC layer and Mo substrate were observed in TiC coatings and the peeling was observed at some part of the interface between Mo and TiC layers in the TiC - Mo coatings. In TiC/Mo FGM coatings, there were no recognizable defects. From these results, it was found that thermal stress , caused by different expansion coefficients between TiC and Mo ^^' ^^ , was significantly reduced in TiC/Mo FGM coatings. As mentioned above, TiC/Mo FGM coatings formed by thermal spraying in the air were effective for the emitter element of the thermionic generator.

Mo-TiC two layer coating | i I" • As sprayed

Heat-treated

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.--.

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Solar radiation Thermocouple Ho-TiC two-layer coating

100 jum

Pyrometer Fig. 7 Scheme of a solar heating test

TiC coating

Fig. 8 Surface and cross-sectional structures of coatings after thermal cycle test

250

Table 2 Results of thermal cycle test of coatings FGM coating TiC/tto coating TiC coating

Substrate/ coating Sound Sound Peeling

TiC/Mo coating boundary Sound Peeling

Coating surface Sound Sound Crack

Evaluation

0 A X

4. CONCLUSIONS Formation of graded coatings (FGM) was carried out with the plasma twin torch spraying method in air. TiC and Mo powder (particle size ilO to 44 fi m) were sprayed onto Mo substrate to obtain coatings with high thermal absorptance in order to improve the performance of the thermal energy conversion system. The results are summarized as follows. 1) Graded coatings are obtained by stepwise control of each powder feeding rate. 2) The compositional gradient in the coating is nearly smooth and linear. 3) After the heat treatment of the sprayed coatings in a vacuum ( at 1473K, 16h, 10 ~^ Pa ), there are no recognizable reaction products between Mo and TiC in the coatings. 4) There are no recognizable defects in the FGM coatings after the thermal cycle test (from R.T. to 1223K, 5 cycles) while there are defects in the TiC single coating onto the Mo substrate. REFERENCE 1) Journal of the Society of Materials Science : An Extreme Situation and Materials, Syokabo, 134, 1987 (in Japanese ) 2) T. Fukushima : Proceedings of the first Int. Sympo. on FGM, 145, 1990 3) Y. Watanabe : Journal of the Japan Society for Aeronautical and Space Sciences Vol. 42, No. 482, 141, 1994 ( in Japanese ) 4) A. Horiguti : The Materials, Vol. 35, No. 388, 35, 1986 the Japan Society for Aeronautical and Space Sciences, vol. 42, No. 482, 141, 1994 (in Japanese ) 5) Journal of the Society of Materials Science : An Extreme Environment and Materials, Syokabo, 123, 1987 (in Japanese ) 6) National Astronomical Observatory : Chronological Scientific Tables, Pyhs.49 (463 ), 1985 (in Japanese)

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

251

Fonnation of a Ti-Al203fimctionallygraded surface layer on a Ti substrate with tlie use of ultrafme particles A. Otsuka, H. Tanizaki, M. Niiyama and K. Iwasaki Steel & Technology Development Laboratories, Nisshin Steel 7-1 Koyashinmachi, Ichikawa, Chiba, 272 Japan A Ti-Al203 functionally graded surface layer of about 0.5 mm in thickness is formed by dryjet spraying of ultrafine particles produced by radio-frequency plasma onto a cylindrical Ti rod. The spraying is made by continuously changing the ratio r = Ti/(Ti+Al203) from 1 to 0 or from 0.5 to 0 in the outward radial direction. The obtained green composite is sintered in the temperature gradient condition, where the Ti-rich and the Al203-rich sides are sintered at about 1400 K and 1800 K, respectively. The ratio r in the sintered composite is found to change gradually from 1 to 0 independent of the starting r value. The r region omitted in the spraying is considered to be compensated by abnormally rapid diffusion of Ti. The size of voids found in the layer is much smaller when the starting r value is set at 0.5 than when it is set at 1. The number of cracks in the layer is reduced by adjusting the sintering conditions and by adding Zror Ti-hydride powders to the substrate. The layer has a fine metallographic structure and the adhesion strength between the substrate and the layer is measured to exceed 60 MPa. 1. INTRODUCTION Though the sintering of ultrafine particles (UFPs) are known to proceed at low temperature and to give rise to fine-grained materials, their industrial applications are still very few. By taking these advantages of UFPs into account a trial to fabricate a functionally graded material (FGM) with the use of them is undertaken in this work, where an artificial tooth root is chosen as an example. It consists of a cylindrical core rod of Ti and a Ti-Al203 surface-coated FGM layer. The composition ratio, r = Ti/(Ti+Al203), is set to decrease gradually in the outward radial direction from 1 to 0. The mechanical strength of the tooth root is sustained by the Ti rod, while the AI2O3 surface serves as a biochemically stable (bioinert) layer. If bioactivity is required, the AI2O3 surface is fijrther coated with hydroxyapatite (HAP) [1-3]. Since metals and ceramics exhibit quite different sintering behavior from each other, a special

252

care must be taken when a green compact of their mixture is sintered. Otherwise there ^sually appear many defects such as voids and cracks due to sintering unbalance [4]. One of the most popular methods to suppress the defect formation is to use processes operated at high pressure and temperature such as hot pressing (HP) and hot-isostatic pressing (HIP) [5]. The cost performance of these methods, however, is very low and the size and the shape of the composite are limited. If it comes to the sintering of an FGM where the sintering temperature should be varied corresponding to the composition, the difficulty in sintering increases greatly and the HP and the HIP methods are not any more appropriate. In order to cope with these problems temperature gradient sintering as will be described in Sec. 2 below is tried in this work. 2. FABRICATION PROCEDURES Figure 1 shows schematically a closed system used for the production and the deposition of UFPs. Raw powders of Ti (average diameter = 27 jam) and AI2O3 (10 fim) are fed into the plasma flame with two mutually independent dispersion feeders, where Ar gas is used as carrier. By adjusting the feeding rates of the two feeders the ratio r of the raw powders can be adjusted at any value between 0 and 1. The total feeding rate is kept at 2 g/min. As for the details of UFP production, see refs. [6-8]. The produced UFPs are homogeneously mixed in an aerosol form, cooled in the cooler and fed into the deposition chamber, where the UFPs are deposited onto the cylindrical surface of the substrate rod by dry-jet spraying [9]. The velocity of the spraying gas is about 200 m/s estimated at the exit of the nozzle. A newly developed L-shaped nozzle is used to increase the deposition efficiency [10]. The cylindrical substrate rod is made of Ti powder of 27 |im and Ti- and Zr-hydride powders of 4 fim in average diameter. The powders are compacted in a polyurethane die at 200 MPa by cold isostatic pressing (CIP) to obtain the substrate of about 3 mm in diameter and about 40 mm in length. It is located in front of the nozzle, where the distance between them is kept at 15 mm. In order to make the spraying homogeneous it is rotated at 30 rpm and simultaneously moved back and forth along the rod length direction with the amplitude of 30 mm and at the frequency of about 0.05 Hz. The total spraying time is about 30 min, during which the ratio r of the UFPs changes gradually from 1 to 0 or from 0.5 to 0. The total thickness of the deposited layer is about 1 mm, which shrinks to about 0.5 mm after sintering. The deposited specimen is sintered in a vacuum (10"^ Pa) furnace shown in Fig. 2. It consists of a W-mesh heater to heat the atmosphere, YAG laser to irradiate the specimen surface, and a thermotracer and a thermocouple to measure temperatures. The YAG laser is

253 first switched-on at the beginning of the sintering to irradiate the outermost AI2O3 surface of the specimen that needs higher temperature for sintering than the Ti core. Then the W-mesh heater is additionally switched-on to heat the whole part of the specimen. It should be emphasized here that the heating sequence is very important. If it is reversed, that is, if the W-mesh heater is switched-on first and then the YAG laser irradiation follows, the FGM layer often flakes off from the substrate. The outermost AI2O3 surface of the specimen is heated to about 1800 K and the central core is heated to 1400 K in this way. The temperature difference of about 400 K is attained between the outer and the inner part of the specimen. The specimen rod is rotated at about 10 rpm and the laser is scanned along the rod axis at 2000 mm/sec to make the irradiation homogeneous.

Carrier gas

Thermocouple

Plasma gas

W-mesh heater

Optically flat window 1.06//m band gap filter

YAG laser

Deposition chamber

Specimen holder

Plasma chamber

Fig. 1 Schematic view of the system for UFP synthesis and deposition.

Specimen

Fig. 2 Schematic view of the temperature gradient sintering furnace.

3. EXPERIMENTAL RESULTS AND DISCUSSION Figure 3 shows the difference in the size of voids (black area) between the two cases where the deposition starts from r = 1 (a) and from r = 0.5 (b). It is clearly seen that the void size is much smaller in the latter than in the former. Since the Ti-rich region with r ranging from 1 to 0.5 showed very large shrinkage during sintering in the previous preliminary experiments, the decrease in the void size is considered to be mainly due to the elimination of this region in the deposition process. The ratio r of the sintered specimen deposited from r = 0.5, however, changes from 1 to 0 as shown in the results of electron probe micro analyses (Fig. 4). The

254

region of r in the range between 1 and 0.5 is considered to be constructed by abnormally rapid diffusion of Ti from the substrate to the FGM layer. Diffusion of this kind is supposed to take place along the surface of the small voids.

FGM surface FGM surface

*'''^' Substrate-FGM boundary

Substrate-FGM boundary

) 00 fi m

100 ;em

(a) r changed from 1 to 0 (b) r changed from 0.5 to 0 Fig. 3 Optical photos of the cross section of FGM perpendicular to the rod axis.

FGM surface

Substrate-FGM, boundary

Signal intensity(Arbitrary unit) Fig. 4 Results of EPMA across the FGM layer. Relative size changes of the substrate and the AI2O3 UFP compact during the sintering are shown in Fig. 5(a). The AI2O3 compact is taken as a representative component of the FGM layer. Temperature profiles during the sintering are shown in Fig. 5(b). The substrate and the AI2O3 compact are sintered at 1400 K and 1800 K, respectively, in consideration of the temperature gradient sintering. The final value of the relative shrinkage of the Ti substrate is about 4 %, which is about half of that of the AI2O3 UFP compact (dotted line). If 5 % hydrides are added to the Ti substrate the amount of shrinkage increases and comes very near

255

to that of the AI2O3 UFP compact, where Zr-hydride seems to be more effective than Tihydride. It should be noted here, however, that the reproducibility of their concentration dependence is rather poor. Though the amount of the relative shrinkage has a tendency to increase with increasing hydride addition up to about 5 %, the relation between them becomes quite complicated above 5 %. This is partly due to the difficulty of uniform compaction because of the high hardness of the hydrides powders. Then it is difficult to determine the proper amount of hydride addition exactly. Since the appearance of cracks in the FGM layer perpendicular to it after sintering is considered to be mainly due to the difference in the relative values of shrinkage between the substrate and the FGM layer, their agreement by the addition of hydrides is expected to be very effective for the suppression of the number of cracks. This is actually seen in Fig. 6, where the decrease in the number of cracks is clearly observed by comparing (a) with (b) as expected.

AI2O3 UFP

Ti+5%Ti-hydride Ti+5%Zr-hyderide

100 Sintering time / min

100 Time/min

(a) Relative size change (b) Temperature profile Fig. 5 Shrinkage behavior during sintering.

(a) Without Zr-hydride (b) With Zr-hydride by 5% Fig. 6 SEM photos of the cross section of FGM perpendicular to the rod axis.

256 As shown above, a fine-grained FGM is obtained with the use of UFPs and by the temperature gradient sintering. Advantages of the present method lie in the high speed formation of an FGM layer and one-process sintering at low temperature and low pressure. The important factors are the distribution of r during the spraying, the sequence of heating with a W-mesh heater and YAG laser, and the adjustment of the difference in the relative shrinkage between the substrate and the FGM layer. The procedures developed here are expected to be extend to the production of various kinds of FGMs. ACKNOWLEDGMENT This work was conducted in the program "Advanced Chemical Processing Technology" consigned to Advanced Chemical Processing Technology Research Association (ACTA) from New Energy and Industrial Technology Development Organization (NEDO), which was carried out under the Industrial Science and Technology Frontier Program enforced by the Agency of Industrial Science and Technology. Authors would like to express their gratitude to these institutions for supporting this work. Dr. S. Koura, Messrs. K. Ohsaki and T. Tanaka and Ms. E. Kobayashi of the authors' laboratory are also acknowledged for their help in performing this work.

REFERENCES 1. K. Ohsaki, H. Tanizaki, K. Iwasaki, M. Ueda, T. Kameyama and K. Fukuda, Proc. 7* Symp. Plasma Sci. Mater., (1994) p.83. 2. T. Kameyama, M. Ueda, K. Onuma, K. Ohsaki, H. Tanizaki and K. Iwasaki, Proc. 14* Int. Thermal Plasama Conf ,(1995) p. 187. 3. H. Tanizaki, A. Otsuka, M. Niiyama, K. Iwasaki, M. Ueda and A. Motoe, Trans. Materi. Res. Soc. Jpn., (1996) in press. 4. R. Watanabe, Materi. Res.Soc. Bull. (Jan. 1995) 34. 5. R. Watanabe, J. Jpn. Industri. Materi.,38(Dec. 1990) 39. 6. S. Koura, H. Tanizaki, M. Niiyama and K. Iwasaki, Materi. Sci. Eng. A208(1996)69. 7. H. Tanizaki, A. Otsuka, M. Niiyama and K. Iwasaki, Materi. Sci. Eng. A, in press. 8. A. Otsuka, S. Koura, H. Tanizaki, M. Niiyama and K. Iwasaki, Nisshin Steel Tech. Rep. 72(1995)11. 9. S. Kasyu, E. Fuchita, T. Manabe and C. Hayashi, Jpn. J. Appl. Phys., 23(1984) 1900. 10. Jpn. Patent H6-336271, (Dec. 1994).

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

257

Oxidation-Resistant SiC Coating System of C/C Composites N.Sato'^, I.Shiota'\ H.Hatta'^ T.Aoki'\ H. Fukuda'^ 1) Kogakuin University,Tokyo,Japan 2)The Institute of Space and Astronautical Science,Kanagawa,Japan 3) Science University of Tokyo,Tokyo,Japan ABSTRACT C/C composites are easily oxidized to evaporate above 500°C. Anti-oxidizing SiC coating on a surface of C/C composites is known to be an effective measure against oxidation. However weakness in bonding strength often become a serious problem in actual applications. In this paper optimal fabrication technique of SiC coating on C/C composite was explored based on two kinds of criteria; bonding strength and oxidation resistance. Thus a parametric study was carried out experimentally in terms of thicknesses of the coating layer and the conversion FGM layer and CVD temperature. It was shown that the optimal condition of the coating is low CVD temperature, thick conversion FGM layer, and intermidiate coating thickness. 1. INTRODUCTION C/C composites are unique material which possess exceptional high heat resistance along with lightweight, high stiffness and high strength [1,2]. For this reason, C/C composites have been expected to be applied to high temperature structures, in such as aerospace and nuclear fusion industries. However C/C composites are known to possess a serious shortcoming; easily oxidized to evaporate above a temperature as low as about 500''C[3,4]. To overcome this defect, anti-oxidizing SiC coating on the surface of C/C composite was a quite effective measure in an oxidizing environment up to about 1700''C[5]. However weakness in bonding strength and coating cracks due to thermal mismatch between the substrate C/C composite and the SiC coating still often become a serious problem in actual applications. In this paper optimal fabrication technique of SiC coating on C/C composite using CVD (Chemical Vapor Deposition) process was explored based on two kinds of criteria; bonding strength and oxidation resistance. 2. EXPERIMENTAL 2.1 Material (1) C/C composite

258 The examined C/C composite was fabricated via. preformed yarn method[6]. The reinforcing fiber, fiber volume fraction, stacking sequence, and dimensions of specimens of it are Toray M40, 50%, 0790°, and 30mm x 30mm x 3mm, respectively. In the laminated C/C composite, periodical cracks along the fiber axis direction, transverse cracks (TCs), frequently appear as shown in Fig. 1. The surface layers of the TCs especially affect characteristics of the coating on C/C composites. (2) SiC coating The SiC coating was composed of two layers; a thin SiC conversion layer (several \xm) and a thick CVD coating layer. The conversion layer was formed by chemical reaction of gas phase Si with carbon in the substrate. The gaseous Si was provided by thermal decomposition of bubled SiCl^, in which carrier gas was H,. The aim of the conversion layer is to make bonding strength of the coating higher. The CVD layer was thermally deposited at three kinds of temperatures, 1800, 1600 and 1200°C, named hereafter CVD-18, -16 and -12, respectively. The SiC coating with three kinds of nominal thicknesses, 60, 100, 200!LAm were formed.

a^g^TO"

Fig. I Cross sectional view of CVD-SiC coated C/C composite with 0/90 stacking sequence

Fig.2 SEM micro-photograph of surface of SiC-coatin5

Typical cross sectional and top views of the SiC coating are shown in Fig. 1 and Fig. 2. It is obvious from these figures that a lot of cracks appeared in the coating. These cracks were formed in the cooling down process from CVD to room temperatures due to thermal expansion mismatch between the SiC coating and the C/C substrate. SEM and optical microscope observations revealed that the coating cracks appeared first just above the TCs to the direction parallel to the TCs and then perpendicular branch cracks developed from them. The apertures of the former cracks are normally larger, about lOjim, than the latter, several \xm. 2.2 Oxidation Test Oxidation tests have been carried out under xenon lamp heating and natural convection of air, where weight loss under constant temperature was monitored. The temperature was monitored by an infrared thermo-viewer and was calibrated with that of the tungsten-rhenium or platinum-rhodium thermo-couples. Temperature range examined was from 600 to 2300°C. Main features of this apparatus are that a large size specimen can be exposed and equilibrium constant temperature is rapidly attained (about 20s).

259 2.3 Bonding Strength of SiC Coating Evaluation method for bonding strength between a hard and thin cx)ating and the C/C composite of the substrate has not been established. In this study the method illustrated in Fig. 3 was tried to be applied, i.e., shear stress is directly applied to the interface between a coating and a substrate by the plunger. We hereafter call the method as "shear load method" and abbreviates as SLM. A typical load-displacement curve obtained by the SLM is shown in Fig. 4. It is obvious from this figure that shear fracture, at the maximum load point, can be easily attained even for thin brittle coating. However several problemes were anticipated in this method; 1) How to apply precisely compressive load along an interface 2) The sample must be processed into a quite thin plate but whether residual stress in the thin sample is same as the thicker original one. To confirm the formar problem, the interfacial strength was evaluated as a function of loading place, X in Fig. 5. The correct value is considered to be the value at the clearance. Therefore from this figure it is concluded that 0.03mm is sufficient clearance between the interface and the plunger edge. Thus in the present study it is tried to maintain the clearance <0.03mm.To reply the 2nd question, the interfacial strength was evaluated as a function of specimen thickness as shown in Fig. 6. If the stress state was not changed when specimen had been cut, the interfacial strength should be linear function of the specimen thickness and crossing origin and this is case for Fig. 6. Thus rather precise strength was considered to be obtainable by the SLM. In the present study

Plunger , Specimen

, Specimen thickness

Support stage

Clearance (X mm)

Fig.3 Schematic drawing of apparatus to measure bonding strength of the coating

Shear fracture point

Specimen thickness;2mm 0

0.02 0.04 0.06 0.08 0.1 Displacement (mm)

0.12

0.14

Fig.4 Typical load-displacement curve obtained in shear strength test

0

0.05

0.1

0.15

Clearance (IDDI)

Fig.5 Shearing strength of C/C composite as a function of clearance between the sample and plunger edge

260 2mm was chosen as the specimen thickness due to covenience of test. Top sarfacc

3. RESULTS AND DISCUSSION 3.1 Characterization of the SiC Coating The cracks in the coating are said to have pivotal influence on the oxidation rate and the interfacial strength of SiC-coated C/C composites. This is because the oxdation mainly tends to occur through the cracks and appearence of cracks lower the thermal stresses due to thermal expansion mismatch between the coating and the substrate. Actually, as clearly observed, severe oxidation is observed in the CIC substrate just beneath the coating crack. Thus characterization of cracks in the SiC coating becomes important even to qualitatively evaluate these phenomena. To characterize coating cracks, crack density, crack opening distribution and distances between cracks were evaluated. Typical surface figures of the SiC coated C/C composite are shown in Fig. 7(a), top view, and (b), side view, where surface of the SiC coating was polished to clarify distribution of coating crack. It is obvious from these figures that there are two kinds of directions. Thick and sparse cracks run paraDel to the surface fiber direction and thin and dense cracks parpendicular to it. The formar cracks formed above the transverse cracks in the surface layer of the C/C composite. The opening width of these cracks tend to reduce with temperature increase and close at CVD temperature, as shown in Fig. 8. This tendency manifests that the coating cracks appear due to the thermal mismatch strain. In Fig. 9, the sum of the mean

t '

^

Coating thickness; 60 // m CVD temperatureaSOOt:

L:^-

i

1_____ 1

i

1

i

1

1

Specimen thickness (min)

Fig.6 Bonding (shearing) strength of interface between SiC coating and C/C composite as a function of sample thickness

Fig.7 Optical micro-photograph for top and side surface of the SiC-coated C/C composite after SiC polishment

261 values of the crack opening width in a 1mm distance along the coating surface after the oxidation test are plotted as a function of oxidation temperature. It follows from this figure that crack gaps increase with the CVD temperature and are not affected by heating of the oxidation tests. In this figure, thermal mismatch displacement calculated from the difference of thermal expansion [7] of both materials is also shown by the dashed line. The observed total gap is lower than the calculated thermal mismatch displacement. It is due to incomplete relaxation of the thermal mismatch strain by the formation of the coating cracks. The residual, not relaxed, portion of the thermal stress in the cracked coating is represented by the length between the solid and dashed line. Hence it follows from this figure that residual thermal stress increases with temperature.

400

600 800 1000 1200 1400 1600 Temperature ("C)

Fig.8 Variation of crack gaps in the SiC coating on C/C composite as a function of temperature

10 H

• o

CVD-1811 1 CVD-121 CVDtB calculated v^ue:8.15/tm

8 i

p §

E

4

i

) <

C

3.2 Strength of the Interface The shear (bonding) strength of interface between the SiC coating and a substrate C/C composite as a function of CVD temperature and coating thickness is shown in Fig. 10 and 11. As shown in Fig. 10, apart from CVD-16, the bonding strength decreases with increase of the CVD temperature and the bonding strength for the top surface is lower than the side surface. These results clearly reflect the effect of thermal residual stress; the thermal residual stress in the room temperature increases with the CVD temperature and the thermal stress in the top surface should be lower than that for the side surface because thermal expansion mismatch between the C/C composite and SiC coating in the through the thickness direction is smaller than that in the inplane direction.

i.

^

J..J L..

CVD12 : 5.5^jum 4

E2

1 L^

m i

C

>

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-

T

2

i i _J500 1000 1500 Oxidation temperature (°C)

-L0

2000

Fig.9 Total crack width in 1mm length of SiC coating after oxidation test

7

60pim

coating

- « — CVD-18 (Top surface) - e — CVD-18 (Side surface) - • — CVD-16 (Cwiversion 5h) - a — CVD-16 (Conversion 3h) —A— CVD-12 (Top surface) -^^— CVD-12 (Side surface) —•—Bare C/C

a

I 3 60

^0

500 1000 Coating temperature

1500 ("C)

Fig. 10 Bonding strength of SiC coating on C/C composite treated at various CVD temperature

2000

262 The results for CVD-16 represent the effect of conversion layer. It is obvious from these results that within the present experimental condition longer conversion time is preferable and comparison with other data should be done on the base of the conversion time 3hrs. The bonding strength also depends upon coating thickness as shown in Fig. 11. This A Conversion time;5h | l e o o t : coating • 2.5 tendency can be explained qualitatively if we O A Conversion tiine;3h I consider the disturbance of the stress near the interface due to application of the coating; ^ 1.5 disturbed portion of stress increases with increase of the coating stiffness, which " 0.5 increases with the coating thickness. This tendency is believed to be predictable 0 50 100 150 200 2J Coating thickness (ju m) theoretically and now we are trying to clarify by Fig. 11 Bonding strength of SiC coating on use of a commercial FEM cord. C/C composites as a function of coating thickness

4. CONCLUDING REMARKS Thicker and low temperature CVD coating was shown to be preferable from the view point of oxidation protection performance[8]. On the other hand, a combination of thinner and low temperature CVD coating and thicker conversion FGM layer is advantageous in terms of keeping high bonding strength of interface between a coating and a C/C composite. Thus the compromise must be done with respect to the coating thickness. From the view point of oxidation protection, the coating thickness should be determined by the consideration of diffusion of oxigen through the SiC coating and coating cracks. The latter is ordinary resolved by use of vitreous crack sealing material. Thus the former determines the coating thickness and experimentally the 50-lOOpim thickness of the coating has been untilized [9]. REFERENCES [1] G. Savage, "Carbon-Carbon Composites", Chapmanhall, 1993. [2] C. R. Thomas, "Essentials of Carbon-Carbon Composites", Roy. Soc. Chem., 1993. [3] J. E. Sheehan," Carbon-Carbon Materials and Composites", NASA Ref. Pub., 1992. [4] J. R. Strife, J. E. Sheehan, Am. Ceram. Soc. Bui., 1988. [5] NASDA, Design Concept of HOPE, TK-S03017., 1992. [6] Okura, A.,"Studies on a New Manufacturing Process of C/C Composites.", 1991. [7] E. Yasuda, S. Kimura, Y. Sibuya, Trans. JSCM., 1980. [8] H. Hatta, Y. Kogo, T. Yarii, "Oxidation Behavior of SiC coated C/C Composites", Proc. 7th Japan-US Conf. on Cmpos. Mater., Composites '95 : Recent Advances in Japan and the United States., 1995. [9] Hoffman, R., Heitzer, J., Kromp, K., Compos. Sci. Tech., 1994.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

263

AI2O3 - Zr02 Graded Thermal Barrier Coatings by EB-PVD - Concept, Microstructure and Phase Stability U. Leushake, U. Schulz , T. Krell, M. Peters, W.A. Kaysser DLR - German Aerospace Research Establishment, Institute of Materials Research, 51140 Cologne, Germany Graded alumina/zirconia thermal barrier coatings (TBCs) offer a great potential to extend the application on aeroengine turbine blades integrating low oxygen diffusion layers of alumina. Since the thermal conductivity of alumina is higher than that of zirconia the design of the TBCs for the use in stationary temperature gradients has to be extended to same heatflux coating thickness. As a basis a processing route for alumina ingots was developed. Evaporation of alumina was carried out characterizing phase content and morphology. Co-evaporation of alumina/zirconia was used to fabricate lateral graded TBCs. 1 INTRODUCTION The aim to increase efficiency of aeroengines led to the demand for increasing combustion chamber temperatures, which can best be realized by protection of turbine blades with thermal barrier coatings (TBC). Nowadays application of TBCs in stationary gas turbines has also become of interest, although conditions concerning application temperature and lifetime are different from aeroengines. State of the art thermal barrier coatings on nickel-based superalloys consist of a bond coat (i.e. MCrAlY) and a yttria partially stabilized zirconia thermal barrier coating on top. TBCs can be produced in two different ways. Electron-Beam Physical Vapor Deposition (EB-PVD) processed coatings show superior lifetime compared to plasma sprayed coatings due to high thermal shock resistance of their columnar microstructure. The TBC system is thermodynamically unstable and during application at elevated temperatures several effects take place, limiting the coating system's lifetime. To increase lifetime and include TBCs in the turbine design several tasks had to be solved. Some solutions may be found by the application of the PGM concept [1,2].

2 ALUMINA/ZIRCONIA GRADED TBCs Critical for the coating system lifetime is the formation of a thermally grown oxide scale (TGO) at the interface BC/TBC during service, which mainly consists of alumina [3]. The possible oxygen diffusion in zirconia itself and the open columnar structure of the ceramic coating allows oxidation of bond coat aluminum. The scale's growth and the difference in

264 material properties causes failure of the thermal barrier coating system at the interface BC/TBC or close by in the zirconia layer. According to [4] the lifetime of a coating system can be estimated by the thickness of this TGO layer, correlating failure with a certain TGO thickness. The concept of a graded thermal barrier coating is to combine the thermal insulation potential of zirconia with the low oxygen diffusivity of alumina in order to reduce the TGO's growth rate. The concept also includes controlled formation of the TGO/TBC interface including higher adhesion and a reduction of interfacial stresses. From the large variety of the alumina polymorphs, a-A^Os is the most promising because of the lowest oxygen diffusivity. The advantage of a continuous grading compared to a single alumina interlayer is believed to be the avoidance of sharp interfaces and their stress concentrations. First steps in EB-PVD processing of graded TBCs have been taken, hoping to combine the conceptual advantages of graded TBCs with the processing advantages of EB-PVDtechnology. Unlike plasma spraying, where only suitable powders have to be available, a more complex processing of the vaporizing material is necessary. Semi- or industrial size coaters work with a continuous feeding system based on the use of ingots (cylindrical shaped ceramic), which will be melted on top in a bottomless crucible. Since only yttria partially stabilized zirconia ingots are commercially available, a new processing route for ingots had to be developed first [5] to achieve graded coatings. 3 EXTENDEND DESIGN CRITERIA FOR TBCS The design is one of the essential points of FGMs. In case of TBCs the design was mainly carried out for the fabrication. Stresses during application may be different and need to be integrated into the design process. This design process should include the calculation of thermal fields and their stresses and also residual stresses during fabrication. By variation of the grading function parameters an optimum state for given loads can be reached. But changing the gradient function also effects the FGMs overall thermal conductivity and therefore a characteristics of the thermal field. This means that in case of constant coating thickness the heat flux or free temperatures depend on the grading function parameters. This makes the comparison of calculated stresses more complicated and the design process has to be extended to take this into account. Since the thermal conductivity of alumina is higher than that of zirconia the thickness of the graded coating has to be modified. A way to solve this problem is to define the effective thermal resistivity R* (see Figure 1). In combination with different rules of mixtures for the thermal conductivity of composites it can be solved either analytically or numerically [8]. In Figure 1 the effect of profile parameter on the effective thermal resistivity is calculated to estimate coating thicknesses of same heat flux. Since ratio of substrate/coating thickness is effecting stresses, stresses in coatings with same thickness and different heat fluxes can not be compared easily. Some results of calculations [9] with modified thicknesses can be seen in Figure 2. Compared to a non-graded TBC of 200 |Lim ZrOi the heatflux of a linear graded coating with the same thickness is more than twice as high, causing also higher stresses. The reduction of the integral thermal conductivity of the coating can be overcome by increasing the overall thickness of the graded layer or by only increasing the zirconia layer thickness. Both ways reduce thermal stresses in the interface.

265

R*(Zr02)

1

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linear rule of mixture

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200 urn FGM. d|=4-2 MWm^

300

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.

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ZTO2/M2O3 - TBC

i\ i\ [\\ \'i

linear graded linear rule of mixture infinite plate bi-axial stress

I n . — Tsurface=1200°C / N Tbutton=900°C \ \

« e|=2MW/rrf

1 \ '\ ** complete graded '*'* \ \ei=2MWm^ 200 \m\ ZrOj. dj=2 MWm^ '. 200 Mm FGM+150 \m ZrOz

()

'X

0.5

1

1.5

2

2.5

Coating Thickness [mm]

Figure 1: Effec. thermal resistivity for differ- Figure 2: Effect of corrected coating thickent rules of mixtures and profile parameters ness on stress distribution 4 EXPERIMENTAL PROCEDURE 4.1 Ingot Processing and Evaluation The developed processing route for ingots is based on the use of bimodal powder mixtures and is adaptable to different powders or materials. In this context only the fabrication of alumina ingots with a diameter of 50 mm will be described. Two alumina powders of different particle size distributions were selected. The fine powder is characterized by an average particle size d5o= 0.6 jim, the coarse one by d5o= 22 |Xm. The coarse one was generated from a finer powder by presintering of a loose powder pile at 1450°C/lh. The powders were wet mixed (Isopropanol) in a high speed homogenizer, with the addition of 6 wt% solved binders, dried in vacuum destilation equipment, sieved and agglomerated in a tumble mill. Compaction was done by uniaxial pressing in a double action die of an inner diameter of 55 mm, vacuum capsulation in a rubber form and cold isostatic compression. If necessary, the green body was machined to the exact diameter of 50 mm using a turning lathe. Finally the binders were removed by separate heating. The evaluation of the vaporizing behavior was carried out in an 60 kW single source EB-PVD coater with video control of the melting pool. 4.2 Evaporation of Alumina Evaporation was carried out in a 150 kW single gun, two sources jumping electron beam coater with separate loading and preheating chambers. As substrates aluminium free chromium steel (X8Crl7, 1.4016) plates (25mm*25mm) were used. Before coating the specimen were preheated at different temperatures. Coating was carried out under different conditions (Table 1) with non-rotating specimens. Specimens were sequentially annealed after coating in a vacuum furnace (
266 1 Substrate Temperature 1 2 3 4 1 preheating [°C] 960 1000 1090-995 1010 coating [°C] 960-890 800 3*1060-800 950 1 annealing [°C] 995°C/5 min Table 1: EB-PVD coating conditions for AI2O3 and Al203/Zr02 layers

1

1

rial mixing ratio allows to achieve a laterally graded coating. Coating conditions are listed in Table 1 as Run 4. Chemical composition was analyzed by X-Ray fluorescence with standards. 5 RESULTS AND DISCUSSION 5.1 Ingot Fabrication and Evaluation The required thermal shock resistance is realized in case of commercially available ingots by a certain level of relative density and phase composition. The fabrication process includes several sintering processes and is therefore specific for the given powders and materials. A more flexible way can be realized by the fabrication in a powder metallurgical process without sintering. The level of 60% relative density can seldom be reached by compaction of single size ceramic powders. The use of bimodal powder mixtures provides a way to improve green density by filling the voids between the large particles with smaller ones. Green density is not the only criteria which has to be considered, green strength of the green body, sintering during evaporation and evaporation behavior has to be taken into account, too. The influence of processing parameters like content and type of binders, axial and isostatic pressure were investigated for the fabrication of zirconia ingots and TZM/PSZ FGMs and could be transferred to alumina [5,6]. Ingots were compacted by an axial pressure of 40 MPa and an isostatic pressure of 200 MPa. Taking into account the results from the vaporizing experiments, a ratio of 65% coarse powder and 35% fine powders was found to be suitable. 5.2 Evaporation of Alumina The phase formation and morphology of coatings depend on process parameters during evaporation, mainly influenced by the deposition temperature. In the common coating temperature range of 900-1100°C the formation of the V- zirconia phase is fixed by addition of stabilizing elements and only the morphology is influenced by the process itself. In case of alumina, phase transformations take place in the temperature range of deposition [7]. The substrate temperature is the equilibrium temperature during the evaporation, being heated by radiation from the melt, reflected electrons and vapor condensation. Since the melting point of alumina is remarkablly lower than that of zirconia, the substrate temperature during evaporation is lower than in the case of zirconia. Vaporization was carried out under three different conditions (Table 1). In Run 1 a high electron beam energy was used and it is characterized by spits and a high deposition rate. Run 2 and Run 3 were carried out at a lower electron beam energy level. In addition Run 3 was coated three times with an intermediate heating between the coating cycles to increase substrate temperature. The phase constitution in the as-coated condition and also after thermal annealing is summarized in Table 2. The identification of alumina phases is not simple, since there are some polymorphs. Strong textures, typical for EB-PVD processed coatings, make it even more complicated.

267 1 alumina as coated phase Run a 0 8

1 2

3

X

1000°C

X

1200°C

Ih 1 2 3

9h 1 2 3

Ih 1 2 3

9h 1 2 3

Ih 1 2 3

X

X

X

X

X

X X

1100°C

X

X X

X

X

X X

X

X

X X

X

X

X

X X

X

X

9h 1 1 2 3 X X X 1

X X

X

X X

~~l

Table 2: Phase content of EB-PVD AI2O3 coatings after annealing The observed peaks of Run 2 and 3 were not as sharp as those observed in monolithic ceramics and Run 1. It is assumed to be caused by remaining un- or not fully crystallized areas. Run 1 shows only a-AliOs and no change in phase formation after annealing was observed. Run 2 consists only of 6-AI2O3, while Run 3 had formed 6 and 0. Annealing at 1000°C has little effect, only the peaks are getting sharper. Formation of a-A^Os is taking place at temperatures above 1100°C, whereby diffraction pattern is different to in-situ deposited a-A^Os. Run 3 shows lower tendency to phase transformation, since the driving force is lower than in Run 2. After annealing at 1200°C the coating consists mainly of a-alumina, only small amounts of other phases are found. The microstructure of the different runs can be seen in Figure 3. In case of Run 1 a dense coating can be observed. Run 2 and 3 show a columnar microstructure. Comparing the microstructure and phase contents of Run 1 with the others, it seems that substrate temperature is not the only criteria for coating microstructure. 5.3 Co-evaporation of Zirconia / Alumina The local dependence of the chemical composition of the co-evaporation of zirconia and alumina is given in Fig.4. The diagram shows clearly a lateral gradient of coating compositions. While the weight content of AI2O3 decreases rapidly with decreasing distance to the Al203-ingot, the weight content of Zr02 rises more slowly. The reason for that is the difference of melting points and evaporation behaviour of Zr02 and AI2O3. As a consequence the density of the zirconia vapor cloud is much higher than the density of the alumina vapor cloud. Furthermore it can be assumed that the re-evaporation rate of alumina adatoms is much higher than the re-evaporation rate of zirconia because of its lower melting and boiling point. For both reasons alumina is the minor component in all coatings.

ZA'^f^W

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Run 1 Run 2 Run 3 Figure 3: Microstructure of EB-PVD Processed Alumina Coatings

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1 OW

75

Figure 4: Lateral gradient of chemical coating composition 6 SUMMERY AND CONCLUSIONS 1. It could be pointed out that the concept of graded alumina/zirconia TBCs is a promising way to extend life time of turbine blades. 2. The design of TBCs for the use in temperature gradients has to take into account the change of overall thermal conductivity by the grading function. The effective thermal resistivity offers possibilities to calculate coating thicknesses of same heatflux. 3. A powder metallurgical processing route for the fabrication of ingots was developed only based on the use of bimodal powder mixtures with satisfactory vaporizing behavior. 4. Evaporation of alumina and co-evaporation of alumina/zirconia was successfully done as a necessary processing step for graded TBCs. 5. Phase formation and morphology of EB-PVD alumina coatings were characterized. 6. Laterally graded TBCs were produced by co-evaporation of alumina and zirconia. ACKNOWLEDGMENTS The authors would like to thank C.J. Kroder, J. Brien, H. Schurmann and W.-D. Zimmermann for their help. The work was supported by the Deutsche Forschungsgemeinschaft (DFG) as part of the German FGM project. REFERENCES 1. U. Schulz, K. Fritscher, M. Peters, W.A. Kaysser, 3rd International Symposium on Structural and Functional Gradient Materials, FGM 94, Lausanne, Switzerland, 441. 2. K. Fritscher, W. Bunk, The First International Symposium, FGM, 1990, Sendai, 91. 3. K. Fritscher, M. Schmiicker, C. Ley ens, U. Schulz, Materials Science Forum, in press 4. S.M. Meier, D.M. Nissley, K.D. Scheffler, T.A. Cruise, ASME J. Eng. Gas Turbine Power, 114(1992), 258. 5. Application for a German Patent, Application number: 19623587.1 6. U. Leushake, Master thesis, RWTH Aachen, 1994 7. A.L. Dragoo, J.J. Diamond, J. of the American Ceramic Society, 11(50), 568. 8. U. Leushake, unpublished results 9. Multitherm VI.4, M. Finot, S. Suresh, MIT

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

269

Microstructure characteristic of plasma sprayed ZrOi/NiCoCrAlY graded coating Zhongda Yin, Xinghua Xiang, Jingchuan Zhu and Zhonghong Lai School of Materials Science and Engineering, Harbin Institute of Technology, Campus Box 433, Harbin 150001, P.R.China ABSTRACT The plasma sprayed ZrOi/NiCoCrAlY graded coating exhibits excellent gradually compositional distribution. The relative density of the graded coating is 90% or so, and contains a few pores inevitably. ZrOi component contains some c, t and t' phases, as well as a little m phase, and the NiCoCrAlY component is mainly consists of Y(Ni) and Y'(Ni3Al), as well as a little a-Cr phase and s-Co phase. During spraying process, NiCoCrAlY reacts with oxygen in air, and some NiO, CriOs, AI2O3 film formed on the particles surface. KEYWORD plasma spraying, Zr02/NiCoCrAlY graded coating, microstructure 1. INTRODUCTION Ceramics/metal graded coating is a new type of thermal barrier coatings(TBCs)^^\ which compositions varied gradually from metal bond layer to ceramics surface layer,and eliminate the macro-interface between ceramics and metal that lies in common duplex TBCs, so the thermal stress could be relaxed, and the thermal shock property could be largely improved. Due to Y2O3 partially stablized Zr02 has a relatively low thermal conductivity and large coefficient of thermal expansion compared with other structural ceramics^^^, and MCrAlY(M=Ni,Co,Fe etc.) alloy has excellent anti-oxidation property at elevated temperatures^^^ plasma sprayed Zr02-Y203/MCrAlY coatings has become the most important thermal barrier coatings(TBCs). Since plasma spraying process has the characteristics of rapid solidification and rapid cooling, the Zr02-Y203/MCrAlY graded coating has complicated microstructure, which need to be deeply investigated^"^^. But this work is unsurfficient, because people are usually more interesting in the usability of TBCs^^'^^. Li this paper, microstructure plasma sprayed Zr02/NiCoCrAlY graded coating was discussed preliminarily. 2. EXPERIMENTAL PROCEDURE The substrate material was TC4(Ti-6Al-4V), and the sprayed powders were 8wt%Y203 partially stabilized Zr02 and Ni-4.5Co-20Cr-4Al-lY alloy, the ZrOi/NiCoCrAlY graded coating was fabricated by air plasma spraying process^^^. Microstructure characteristic of plasma sprayed Zr02/NiCoCrAlY graded coating was investigated by means of Optical Microscopy(OPM), Scanning Electron Microscopy(SEM) and X-ray Diffraction analysis(XRD), Transmision Electron Microscopy(TEM).The density of graded coating was measured by Archimeds method.

270 3. RESULTS AND DISCUSSION 3.1. Distribution of microstructure Fig. 1 is the cross-sectional microstructure of Zr02/NiCoCrAlY graded coating, in which the left side is TC4 substrate and the other is the graded coating. Li the coating, the white phase is NiCoCrAlY alloy and the gray one is Zr02. From substrate to coating surface, the amount of ZrOi increase gradually, and that of NiCoCrAlY changes oppositely. The Zr02/NiCoCrAlY graded coating exhibits excellent gradually compositional distribution. The coating was formed by the melted particles heaped up layer by layer, and displays typical laminar structure.

Fig.l Optical microstructure of Zr02/NiCoCrAlY graded coating 3.2. Density of graded coating Fig.2a and Fig.2b is the surface morphology of Zr02 and NiCoCrAlY component respectively. It could be found that the sprayed powders deformed severely, hi spraying process, the melted powders impact on the surface of substrate or deposited layer at very high speed, and lead to the particles flow and deform along the surface. It could be also found that some pores and undeformed spherical particles with very small size remain at the interface between deformed particles.

Fig.2

SEM images of the coating natural surface (a) Zr02 layer and (b) NiCoCrAlY layer

271 Fig. 3 displays the density of the graded coating, which reveals that the density of graded coating decreases in linear way with the rise of Zr02 vol% due to the lower density of Zr02 in comparsion with the Ni-base alloy. The relative density of graded coating is 90% or so. The coating formed by deformed particles heaped up layer by layer, and the melted particles contract during cooling process. Meanwhile, some undeformed particles which include unmelted powders remain in the coating. So some pores lie in the coating inevitably. 100

0

20

40

60

80

100

Content of ZrO in graded coating(vol%)

Fig.3 Density and relative density of the graded coating

60 2e,deg

Fig.4 XRD pattern of NiCoCrAlY bond layer

3.3. Phase compositions of the graded coating 3.3.1. Phase compositions of NiCoCrAlY component Fig.4 is the XRD pattern of NiCoCrAlY bond layer. The result shows that NiCoCrAlY alloy includes alloy phases and some NiO, Cr203, AI2O3 etc, which indicates that during spraying process, NiCoCrAlY alloy particles react with oxygen in air, and form some oxide films on the particles surface.

Fig.5

SEM images of NiCoCrAlY alloy (a) y(Ni) phase, (b) y'CNisAl) phase

272 The microstructure of NiCoCrAlY component was investigated by SEM, which shown in Fig.5. The NiCoCrAlY alloy is mainly consists of y(Ni) phase(Fig.5a) and Y'(Ni3Al) phase(Fig.5b). y(Ni) phase exists with fine column grain structure, and displays the characteristic of rapid solidified microstructure. -/'(NisAl) phase displays typical cellular structure, and y phase distributes along the interface between y' phase grains, which indicates that the reaction of Ni and Al occurred during plasma spraying, and forms NisAl phase. TEM observations reveal the phase structure of NiCoCrAlY alloy more clearly. Fig.6a is the morphology of y(Ni) phase. It could be found that a lot of dislocations exist in y(Ni) phase, which twine each other and form net structure inside the grain, and squeeze to form dislocation walls at the boundary(shown by the arrow). A little a-Cr phase and pure Co phase were found in NiCoCrAlY component. Fig.6b is the morphology of a-Cr phase, it could be found that a-Cr phase exists with cellular structure, which forms during solidifying process. Fig.6c displays that^ome dislocations lie in pure Co phase, Fig.6d is the electron diffraction pattern from [01 TO] direction of pure Co phase, which indicate that some element Co transformed from fee to hep structure during rapid cooling(i.e. y-e), and formed s^Co phase.

Fig.6 TEM micrographs of plasma sprayed NiCoCrAlY alloy (a) y(Ni) phase, (b) a-Cr phase, (c)morphology and (d) EDP of s-Co phase The microstructure of plasma sprayed NiCoCrAlY alloy is very complicated, and is different from the common Ni-base superalloy. This results from the complex forming process

273

of the microstructure. On the one hand, because of the rapid solidifying and cooling process, the solubilizing process of alloy elements such as Co, Cr, Al and Y into Y(Ni) phase is unsufficient, so some pure element phase such as a-Cr etc. remained to room temperature. On the other hand, owing to the alloy particles went through rapid solidification and cooling, largely mechanical stress and thermal stress formed in the coating, high density dislocations form in NiCoCrAlY alloy, and y->s phase transformation occurred in Co phase which indues tha s-Co phase separated from Y(Ni)phase^[8] 3.3.2. Phase compositions of plasma sprayed ZrO2 component plasma sprayed ZrOi is consists of c, t, t' and a little m phases^^^(shown in Fig.7). Fig.Ta is the morphology of c phase which exhibits fine column grain structure and has the characteristic of rapid solidified microstructure. t'phase exhibits twins structure^^^^(Fig.7b), and t phase exists with fine grain structure (0.2|Lim)(Fig.7c). Fig.7d exhibits the morphology of m phase which has the characteristic of "N" type, and shows that the t->m transformation has the characteristic of rapid transformation^^ ^^.

Fig.7 TEM micrographs of plasma sprayed ZrOi component (a) c-Zr02 phase, (b) t-Zr02 phase, (c) t'-ZrOi phase, (d) m-Zr02 phase Owing to the different ZrOi particles or the different zone of ZrOi particles have different cooling rate, some c-Zr02 phase occurrs c ^ t diffusive transformation and some c-Zr02 phase occurred c-^t' nondiffusive transformation. Meanwhile, due to the Y2O3 content is

274 high(8wt%), increasing the stabilization of ZrOi, some c-ZrOi retained to room temperature. Moreover, only a little t-»m martensite transformation has taken place in some t phase grain with larger size^ . 4 CONCLUSIONS 1. The plasma sprayed Zr02/NiCoCrAlY graded coating exhibits excellent gradually compositional distribution, and displays typical laminar structure. 2. The relative density of the graded coating is 90% or so, and contains a few pores inevitably. 3. Owing to the rapid cooling process after spraying, the microstructure of the coating is very complex. Zr02 component contains some c, t and t' phases, as well as a little m phase, and the NiCoCrAlY component is mainly consists of y(Ni) and y'(Ni3Al), as well as a little a-Cr phase and s-Co phase. During spraying process, NiCoCrAlY reacts with oxygen in air, and some NiO, Cr203, AI2O3filmformed on the particles surface. REFERENCES 1 2 3 4

Musil J, Fiala J. Surface and Coatings Technology, 1992, 52: 211 Taylor R, Brandon J R, Morrell P. Surface and Coatings Technology, 1992; 50:141 Stecura S, NASA85" 15678 Gudmundsson B, Ejacobson B. Materials Science and Engineering, 1989, A108: 73-80, 181-187,105-115 5 Wu B C, Chang E, C.H. Chao C H. Journal of Matrials Science, 1992, 25: 1112 6 Steffens H D, Babiak Z, Fischer U. Surface Engineering, 1987, 6: 471 7 Xiang Xinghua, Zhu Jingchuan, Yin Zhongda and Lai Zhonghong. Surface and Coatings Technology, 1996, (to be pubUshed) 8 D A Porter, K E Easterling. Phase Transformation in Metals and Alloys. Sweden, Van Nostrand Reinhold Co. 1981, 410 9 Xiang Xinghua, Zhu Jingchuan, Yin Zhongda and Lai Zhonghong. Journal of the Chinese Ceramic Society, 1996, (to be published) 10 Zhou Y, Lei T C, Sakama T. J Am Ceram Soc, 1991;74(3):633 11 Dai Zhurong, Li Baocheng, Wu Houzhen. Acta Metallurgica Sinica, 1991 ;27(3):A246 12 Lange F F . Transformation toughening. J mater sci, 1982; 17:225

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

275

FORMATION OF FUNCTIONALLY-GRADED MATERIALS THROUGH CENTRIFUGALLY-ASSISTED COMBUSTION SYNTHESIS W. Lai, Z. A. Munir, B. J. McCoy, and S. H. Risbud Department of Chemical Engineering and Materials Science, University of California, Davis, CA 95616, USA. ABSTRACT The use of a centrifugal force during combustion synthesis to obtain functionally-graded materials was investigated. Composites formed by the reaction 2A1 + 3 CuO + x Cu = AI2O3 + (3 + x) Cu were synthesized in a centrifuge. The effects of diluent content, x, relative density of the reactants, and the particle size of CuO were investigated. Graded zones between the ceramic and metallic phases were obtained under a given set of these parameters. Phase separation times were calculated from sedimentation theory and discussed in light of experimental observations. 1. INTRODUCTION The use of self-propagation high-temperature synthesis (SHS) as a method to form functionally-graded materials (FGM) has been investigated previously [1-5]. Other means through which FGM structures can be formed have also been reported [6]. Of these, the use of a centrifugal force to coat the inside surfaces of steel pipes with a corrosion-resistant layer is of special interest to this work [7,8]. Although the objective in such studies is to form an adherent coating, the formation of a compositionally graded interface, even on a very limited scale is a desired outcome. The most widely investigated FGM systems are those in which the product is a graded ceramic/metal composite. If the combustion temperature is higher than the melting point of the metallic component, then, in principle, the application of centrifugal force during combustion can lead to a non-homogenous phase distribution. The effect of a centrifugal force on SHS reactions has been investigated with a primary focus on the influence of this force on the dynamics of the combustion process [9-11]. In this paper we present the results of an investigation on the formation of a functionally-graded ceramic/metal zone in a composite formed by combustion synthesis under the influence of a centrifugal force. 2. EXPERIMENTAL MATERIALS AND METHODS Several reactions systems, primarily thermite reactions, were investigated in this study [12]. However, the focus of this short paper is on the formation of FGM structures in the composite AI2O3 / Cu. The composite was synthesized via the thermite reaction, 2A1 + 3 CuO + xCu - • AI2O3 + (3 + x) Cu

(1)

In Eq(l), X represents the amount of copper diluent added to the thermite reaction, and had values of 4, 6, 7, and 8. The aluminum powders used were 99.5% pure and had a sieve designation of -325 mesh. Two powders of CuO were used. The first had a purity of 99.99% and a particle size range of 43-841 |Lim, while the second was 99+ % pure with a particle size of <5 |Lim. The copper (diluent) powders were 99% pure and had a sieve designation of -325 mesh. For the stoichiometric ratios indicated in Eq(l), powders were mixed and the samples were

276 prepared in two ways. In one method the powders were pressed to form cyHndrical pellets (1 cm in diameter and 1 cm high) with a relative density of 55%. The resulting pellets were placed in a graphite cylinder, which was subsequendy placed inside a furnace. The second method of sample preparation involved pouring loose powder mixtures directly inside the graphite cylinder. These powders were either shaken to give a relative density of 37%, or were left as poured, resulting in a relative density of about 20%. The graphite cylinder containing the sample in either form (pellet or loose powder), was placed inside a specially designed furnace. The sample assembly was covered with insulation and placed inside a large ceramic container equipped with an inlet and oudet for argon gas flow. The entire set-up was positioned inside the working arm of a centrifuge. The centrifuge used for this study has a one meter-long arm holding the "working bucket". A second arm provided the counter-balance. The maximum centrifuge acceleration, g, is 100 g^, where g^ is the gravitational acceleration. The entire sample enclosure was purged with argon gas before starting a centrifuge experiment. When the acceleration reaches the desired g level, the process of heating sample is initiated. The temperature was raised to the ignition point, which for Eq(l) is around 900°C. Argon gas flow is maintained until the sample temperature has decreased to near room temperature. After being removed from the centrifuge, samples were examined and analyzed by optical microscopy, scanning electron microscopy (SEM), x-ray mapping, computerized image analysis, and x-ray diffraction analysis (XRD). 3. RESULTS AND DISCUSSION When the reactants of Eq(l) were in the form of a compacted pellet (55% relative density), the product contained non-segregated AI2O3 and Cu phases dominated by an inter-connected ceramic structure. However, when the reactants were in the form of a loose powder (37% relative density), the results were significantly different. For x = 6 or 7, the product contained three regions. The bottom region was dense Cu, the top region was relatively, porous AI2O3, and in between these there was a transition region showing a graded composition. Figure 1 shows an SEM micrograph of the graded region between the ceramic and metallic phases for the reaction (Eq(l)) with X = 6. Nearly spherical AI2O3 particles are distributed according to size, with the largest near the boundary between the transition region and the ceramic region. The transition region and the metallic region (Cu) are nearly fully dense while the ceramic region contained porosity, as can be seen in Figure 1. A microprobe analysis of the compositional change between the upper and lower ends of the FGM zone (transition region) is shown in Figure 2. The atomic concentration of Cu increases from zero at the upper interface of the FGM zone to nearly 100% at the lower interface. The aluminum concentration has an opposite trend, decreasing from 100% at the upper interface to nearly zero at the lower interface. A graded distribution of the ceramic particles for the case of x = 7 was also obtained. However, a major difference between the two cases relates to the size of the alumina particles. The difference between the two systems can be better demonstrated by plots of AI2O3 particle size distributions for the two cases. These are seen in Figures 3(a) and (b) for x = 6 and 7, respectively. The plotted results were obtained by a computerized image analysis where the average particle size was determined in layers sectioned successively within the FGM zone. In both cases, the particle size decreased by a factor of two from the ceramic side to the metallic side in the FGM zone. However, the AI2O3 particles in the x = 6 system are significandy larger and

277 their distribution varies inversely with distance cubed, as contrasted to a nearly linear function for the case with x = 7. The difference in particle size between the systems ( x = 6 and x = 7) is attributed to two related factors: combustion temperature and alumina content. For x = 6, the adiabatic combustion temperature is 2694K and for x = 7, the corresponding value is 2511K. Obviously, the AI2O3 content is also lower for the case of x = 7. Both of these factors favor the growth of larger alumina particles for the case of lower dilution (i.e. x = 6).

Figure 1. SEM micrograph of graded AI2O3/CU region (Note: Dark particles are AI2O3)

uo - _ Q

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80-

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.

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1

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1

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1

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200 400 600 600 1000 1200 1400 1600 1600 2000

Distance ( ^ m )

Figure 2. Microprobe analysis of Al and Cu across the compositionally graded region.

278

< e I-

s es

o 1400

300

400

500

600

Distance (^m) Distance ( ^ m )

Figure 3(a) AI2O3 particle size distribution within the graded zone, x = 6.

Figure 3(b) AI2O3 particle size distribution within the graded zone, x = 7.

As was pointed out earlier, when compacted powders (55% relative density) were used, the product showed a network of inter-connected ceramic particles and no compositional gradient. With powders having a relative density of 37%, the product contained an FGM zone as described above. With the use of powders with a still higher porosity (relative density of about 20%), the product consisted of totally segregated ceramic and metallic phases. It is not clear whether these observations are related to the possible role of porosity in sweeping the ceramic phase to the top or they are the consequence of the larger ceramic interparticle distance in the product. At this point the role of porosity in the process of formation of FGM's through a centrifugally-assisted method is not clear. 4. ANALYTICAL EVALUATION OF FGM FORMATION IN A CENTRIFUGAL FIELD The analysis of the AI2O3 particle size distribution within a matrix of copper resulting from a centrifugally-assisted thermite reaction was made utilizing the population balance equation [13].

df ^ d{vf) ^ d(af)^^^ dt dz dr

(2)

where f is the particle size distribution function, v is the partcicle terminal velocity, Q. is the growth rate of the particle, A is the net generation rate of particles, t is time, r is the particle radius, and z is distance. Application of Eq(2) is based on the following simplifying assumption: (a) the separation process begins immediately following the initiation of the reaction, (b) smallsized nuclei form in the homogeneous low density pellet first, then the sample shrinks to the high density product. Thus the nucleated particles are assumed to experience no growth during the densificaiton stage because of its short duration. We consider the formation of larger particles to be the consequence of particle agglomeration. With these assumptions, a Laplace transformation

279 of the population-balance model gives the relationship of the initial particle size distribution function, f^^, to the final distribution function, as: f (r, z, t) = f,(r, z) [ l - U ( x ) ]

(3)

where x = t - z / v. The term U(x) is a step function with U(x) = 1 for x > 0 (t > z / v) and U (x) = 0 for X < 0 (t < z / v). In Eq(3), v is the terminal velocity of the particle, t is time, and z is distance. Assuming f^ to be in the form fo(r) = br^

(4)

the expressions for the volume fraction of particles in the sample, F^(z, t = o) and F(z, t) are: (5)

(6) Normalization of the volume fraction (at any t) relative to the initial value gives

where r,^ and rj^ are the maximum and minimum particle size, respectively, and r^ is defined by

where L is the length of the graded zone, t is time, and (j) is defined by Stokes' law as (9) The terminal particle velocity, v, is in turn defined by v-^(P.-P.)-r'

(10)

where p^, p^. are the densities of the solid (AI2O3) and liquid (Cu) phases, respectively, g (= agj is the centrifugal acceleration, |LI is the liquid phase viscosity, and r is the particle radius. The normalized volume fraction, E, can be calculated from experimental results. The calculation requires knowledge of F(z, t) and F^ (z, t = 0) at any given z value. The former can be obtained from image analyses of sections of samples (i.e. at various z values). However, F^ cannot be determined experimentally but an approximate value can be calculated from the initial

280 stoichiometry, Eq(l), assuming the product to be a fully dense mixture of AI2O3 and Cu. For x = 6 and 7, F^ is 28.74 and 26.63% by volume AI2O3. Thus assuming the 2-dimensional image analyses to represent volumetric distributions, E values are calculated as a function of z, as shown in Figures 4(a) and (b) for systems with x = 6 and 7, respectively.

Best-^hline •

-T—I 0.06

O.i

Distance (cm)

Figure 4(a) Normalized volume fraction of AI2O3 particles with the graded region, x = 6

0.03

1

1

1

0.04

Experiment

I

• •

0.05

0.06

0.07

Distance ( cm)

Figure 4(b) Normalized volume fraction of AI2O3 particles with the graded region, x = 7.

Through a least-squares fit of the E values, two experimental parameters of the separation process can be calculated from Eq(7). These are the particle size exponent, a, and the separation time, t. The last parameter is implicit in the definition of r^ in Eq(7). The calculated values for "a" and t for x = 6 are -1.8 and 0.61s, respectively. The corresponding values for x = 7 are 2.8 and 0.27s. The calculated times are the durations of the separation process for the two x values. The separation process, of course, takes place only when the copper is in the liquid phase, and thus the total time when the sample temperature is at or above 1083°C is important. Attempts to measure the temperature profile during the centrifuge experiments were not successful. Determinations of temperature profiles made at 1 g^ and in a non-flowing argon atmosphere showed that the duration when T > 1083°C is 12 and 7s for the systems with x = 6 and 7, respectively. These times are higher than the calculated separation times by factors of about 20 and 26, respectively. A complete phase separation would take place if the copper remained in the molten state for the times indicated by the 1 g^ temperature profiles. However, a simple heat transfer analysis [12] shows that in the presence of a flow of argon gas, convective heat loss could reduce the times by a factor of about 30 [14]. When taken into account, heat loss would reduce the length of the separation process to 0.4 and 0.21s for the cases of x = 6 and 7, respectively. These approximately calculated values are in general agreement with those obtained from Eq(7).

281 ACKNOWLEDGMENTS This work was supported by a grant from the National Science Foundation (Division of Materials Research).

REFERENCES 1. N. Sata, K. Nagata, N. Yanagisawa, O. Asano, and N. Sanada, Proceeding of the First USJapanese Workshop on Combustion Synthesis, Tokyo, Japan, Y,. Kaieda and J. B .Holt, (eds.), 1990, p. 139. 2. S. E. Niedzalek and G. C. Stangle, /. Mater. Res., 8 (1993) 2026. 3. Y. Miyamoto, H. Nakanishi, I. Tanaka, T,. Okamoto, and O. Yamada, Proceeding of the First US-Japanese Workshop on Combustion Synthesis, Tokyo, Japan, Y. Kaieda and J. B. Holt, (eds.), 1990, p. 173. 4. Y. Matsuzaki, H. Hino, J. Fujioka and N. Sata, Proceeding of the First US-Japanese Workshop on Combustion Synthesis, Tokyo, Japan, Y. Kaieda and J. B. Holt, (eds.), 1990, p.89. 5. Z. Fu, R. Yuan and Z. Yang, Proceedings of the First International Symposium, FGM, Sendai, Edited by M. Yamanouchi, M. Koizumi, T. Hirai and I. Shiota, 1990, p. 175. 6. J. B. Holt, M. Koizumi, T. Hirai, and Z. A. Munir, Editors, "Functionally Gradient Materials", Ceramic Transactions, vol. 34, American Ceramic Society, 1993. 7. O. Odawara, J. Amer. Ceram,. Soc, 73 (1990) 629. 8. O. Odawara, K. Nagata, K. S. Goto, Y. Ishii, H. Yamasaki, and M. Sato, J. Jpn. Inst. Met., 52(1988) 116. 9. A. G. Merzhanov and B. I. Yukhvid, Proceedings of the First US-Japanese Workshop on Combustion Synthesis, Tokyo, Japan, Y. Kaieda and J. B. Holt, (eds.), 1990, p i . 10. B. B Serkov, E. I. Masksimov, and A. G. Merzhanov, Combust. Explos. Shock Waves, 4 (1968)349. 11. S. A. Karataskov, V. I. Yukhvid, and A. G. Merzhanov, Fiz. Gor. Vzryra, 6 (1985) 41. 12. W. Lai, MS thesis. University of California, Davis, CA, 1996. 13. D. M. Himmelblau and K. B. Bischoff, "Process Analysis and Simulation-Deterministic System", Wiley, 1968, Chap. 4. 14. R. B. Bird, W. E. Stewart, and E. N. Lightfoot, Transport Phenomena, Wiley, N. Y., 1960.

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SHS - A NEW TECHNOLOGICAL APPROACH FOR CREATION OF NOVEL MULTILAYERED DIAMOND-CONTAINING MATERIALS WITH GRADED STRUCTURE E.A.Levashov^, I.P.Borovinskaya^, A.V.Yatsenko^, M.Ohyanagi^, S.Hosomi^, M.Koizumi*' ^SHS-Center of Moscow Steel and Alloys Inst, and Inst, of Structural Macrokinetics RAS, Leninskypr., 4, Moscow, Russia ^Russian Ministry of Science and Technical Policy, Tverskaya str., 11, 103905, Moscow, Russia ^'Ryukoku University, Yokotani 1-5, Seta, Ohtsu City 520-21, Japan ^Tomei Diamond Co. Ltd., Joto 4-5-1 Oyama 323, Japan

L INTRODUCTION Functionally gradient method is that same process allowing the solution of three problems simultaneously while manufacturing diamond tools and development of new materials [1-4]. Additional demands are made to the production of multi-layered diamondcontaining and functionally gradient materials with a gradually from layer-to-layer changing diamond concentration. Those demands include: the increase of a material impact resistance and strength, reduction of the expensive diamond powder input, the increase of the expensive diamond powder input, the increase of the diamond boundary concentration in the working layer. The SHS-method allowed to produce 6-layer composites with (Ti,Mo)C ceramic binder and from layer-to-layer changing diamond concentration from 0 to 12 % with a step of 3 % and from 0 to 25 % with a step of 5 %. Regardless the composition of the exothermal mixture there is a boundary concentration of diamond powder in the mixture above which the SHS-process in the layer can't proceed. This can be explained by the fact that being an inert diluent diamond possesses a relatively high coefficient of thermal conductivity (X) as compared to the charge X. The growth of its concentration results in the increase of heat losses from the diamond heating up and from the heat transfer through the diamond grains to the environment and finally the combustion process is interrupted. However it's possible to produce a thin (1-2 mm) diamond layer with the diamond concentration up to 90 % when diamond is introduced into the charge layer of the metal powder with a melting temperature much low than the combustion temperature of the mixture in the diamondless layer. After the metal melting in the diamond layer in course of the combustion process the melt saturates the porous skeleton of the capillary forcer. The diamond layer reduces and the diamond concentration sharply grows because the metal binder leaves the layer. The authors of the present work studied FGM in the system (Co+diamond)/(TiC+Co) obtained under various ratio of the layer masses. The behaviour of natural diamond in the combustion wave of various SHS-systems (Ni-Al, Ti-Mo-C, Ti-Al-C, Ti-B, Ti-B-Si) was studied also for the production of bi-layered compositions with a FGM structure.

284

2. EXPERIMENTAL PROCEDURE The following powders were used in the experimental: carbonyl cobalt of a dispresity less than 40 |j,m; titanium powder produced by NPO "Tulachermet"(Russia) with the size of the dominant fraction in the range 63-7-160 (xm; brown amorphous boron 98,5 %; molybdenum - less than 50 ^m; carbonyl nickel of a dispresity less than 50 |im; aluminium - less than 10 fxm; carbon (lamp soot) with the particles measured about 0,2 |im; synthetic diamond AC 20 brand 160/125 jim (Russia); synthetic diamond IRV brand 150/125 |im produced by Tomei Diamond Co., Ltd (Japan); natural diamond A5 brand 250/200 [im of a medium strength as much as 38 N (Russia). Initial powders were dried out in specific drying chambers under the temperature 85-95° C. The reactant powders were weighed out in the proper stoichiometric proportions keeping a constant equimolar ratio of Ni-Al; 80 % (Ti-C) + 20 % (mass) Co; 84 % TiB + 16 % Ti; (Ti - aC) + 30 % (mass) Mo. Reactive mixtures of various compositions were prepared in ball mills of a volume of 6 litres. The necessary amount of diamond (from 3 to 25 % mass.) was introduced into the charge and was mixed up with it without grinding balls. The ready mixture was pressed into bi-layered and multi-layered pellets of a diameter of 48 mm and of a height 10-20 mm with a relative density of 50-60 %. Nine compositions with the diamond concentration of 3, 5, 7, 9, 10, 12, 15, 20, 25 mass % were mixed with the charge Ti-C-Mo to produce multi-layered semi-products. The ready mixtures were placed layer-by-layer into a pressform in the following order: diamondless layer weighing 25,5 g; 3 mass % diamond layer, weighing 10 g; 5 % layer 10 g; 7 % layer - 10 g; 9 % layer - 9,9 g; 12 % layer - 9,9 g. After densification pellets were obtained 48 mm in diameter with the thickness of the layers 5.0, 2.0, 2.0, 2.0, 2.0, 2.0 mm correspondently. Multilayered pellets with the diamond concentration from layer to layer as much as 0, 5, 10, 15, 20, 25 % mass were prepared similarly. The final pellet was placed into a reactional mold. An SHS reaction was initiated from the lateral face of the cylindrical pellet by a tungsten spiral. After accomplishment of the combustion reaction and propagation of the combustion synthesis wave, the hot SHSproducts were compacted in a hydraulic press at P > 400 MPa for no more than 10 s. The time of exposure to pressure was chosen dependent on the combustion temperature and reology of the products, e.g., on their plasticity and the amount of the liquid phase formed. Usually this time is 0.5 -^ 2 sec. SHS-products were cooled at the room temperature. To produce FGM with cobalt varied concentration a mixture was prepared of the following composition: 64 % Ti + 16 % C + 20 % Co. The mixture weighing 56 g was placed into a mold. Then cobalt powder was added. Three pellets were obtained with the diameter of 48 mm with various mass ratio of the mixture and cobalt layers: 13/56 (0.23); 20/56 (0.36); 28/56 (0.5) correspondently. The relative density of the mixture layer was 0.58; of cobalt layer - 0.65. The SHS-densification was carried out in the reactional mold with the values of the delay time ti = 2 -r 5 sec; pressure Pk = 30 MPa and time of exposure t2= 5 -r 10 sec. Concentration profile of cobalt distribution throughout the sample thickness were constructed by means of micro-X-ray-spectral analysis (MXSA). The regime with the optimal correlation of the parameters: mco/m(Ti-c); Pk; ti; t2 was determined. A complex of parameters was considered as the optimal one when the cobalt layer was melted at the expense of the heat of chemical reaction Ti + C + Co -> TiC + Co and

285 all the melt penetrates through the synthesis products TiC + Co. The next series of experiments when diamond was introduced into the cobalt layer with the concentration equal to 10 and 20 % mass was carried out under the optimal regimes. The methodics used in the experiments with the natural diamond A5 was similar to that in the paper [3], The ratio of the masses mi/m2 varied (mi - the mass of the exotthermal mixture with diamond concentration equal to 25 % vol.; m2 - diamondless layer), Pk = 30 MPa, t2 = 1 sec, ti = 2 -=- 5 sec. The phase analysis was carried out on an X-ray diffractometer DRON-3M (CuKa and FcKa radiation) with the rate of X-ray photographing 2 °/min. The morphology of the products synthesised and diamond grains was studied on electron scanning microscopes JSM-35 (JEOL), and JSXA-733 (JEOL). The strength of layer-by-layer recuperated diamond grains was determined by the standard method by crushing. The wear resistance test was performed in comparison with the volume loss of the diamond grinder against the volume loss the samples, using a resin bond diamond grinder. The ratio in the grinding test, volume loss of the sample test piece, was measured as one of the index for the wear resistance degree. The ratio in samples against that is cemented carbide (WC-Co, K-10) was evaluated as the wear resistance index. The test was performed under wet grinding condition (wheel speed: 1500 m/min, table speed 5.0 m/min, down feed: 0.02 mm/pass).

3. RESULTS AND DISCUSSION 3.1. Multilayered version Figure l,a and b present the distribution profiles of diamond concentration and strength through the thickness of multilayered samples. The figure shows that the diamond strength grows from 3 % layer to 12 % layer (fig.la) and is the maximum one in the layer with the diamond concentration equal to 15-20 % (fig.lb). The analysis of the synthesis products' mixture as well as the conclusions of the proper [3] allowed to explain the scientific results. Low values of the diamond strength in the synthesis products when its' concentration in the layer is small can be explained by a powerful heat stroke on every grain from the diamondless layer. The heat transfer from layer doesn't produce any noticeable effect on the velocity of the combustion wave propagation in the diamond layer. The heat stress on every diamond grain decreases with the growth of the diamond concentration and the diamond strength reduces inconsiderably. There is one more important feature which is the diamond protection against oxidation by the heat-resistant synthesis products (Ti,Mo)Ca, that limit the access of oxygen to its' surface. However when diamond concentration in the layer exceeds 20 % mass then the convention rate in the combustion process reduces (initial reagents are present in the layer) because of the high heat losses from the diamond heating up and heat transfer through the diamond grains into environment. The porosity of the layer with the diamond concentration more than 20 % is raised. The diamond isn't protected from the oxidation and the grain strength decreases. Such an approach - the production of multi-layered composites - allows to determine the limit of combustion with the growth of diamond concentration in the mixture and to construct multi-layered diamond containing materials for tools of various destination. Besides, the multilayered materials are characterised by the raised impact viscosity and strength as comparing to the homogeneous and bi-layered diamond-containing composites. The mentioned materials find their applications in industry for the production of cutting and grinding tools.

286

3.2. FGM - a version in the system Co+diamond (TiC+Co) Figure 2 presents the concentration profiles of cobalt distribution through the thickness of FGM samples, produced by the SHS-densification technology with the 12

P, N

mass % dia 10-

6H

2H

0+ 8

9

10

' thickness, mm

a) mass % dia

thickness, mm b)

Fig. 1. Distributions of diamond grains AC20 (160/125 |im) concentration and diamond grains strength (P) in the SHS products with (Ti,Mo)Ca ceramic binder on the thickness of FGM-compositions with 12 %(a) and 25 % (b) of diamond.

287 various ratio of the masses of the charge layers of pure cobalt and of the exothermal mixture Ti+C+Co. When the mass ratio is equal to 0.5 (curve 1), cobalt is not melted completely, there is only its' partial fusion on the division boundary between two layers. The decrease of the cobalt charge layer mass (curve 2) results in its' complete melting at the expense of the heat of the chemical reaction. Cobalt penetrates into layer of the synthesis products by two ways: capillar impregnation and migration. The capillar impregnation of the porous hot combustion products by cobalt starts in the process of the combustion wave propagation before the application of densification pressure and proceeds during the stage of the products' densification (after the pressure application) while the existence of capillar-porous space. The consequent migration of cobalt possibly occurs by the migration mechanism. In the third case (curve 3) cobalt penetrates through the reagents and the layer of pure cobalt hot formed in the final products. The mass proportion equal to 0.23 was taken as the optimal one. In such a regime the diamond-containing FGM-samples were produced with diamond concentration in the cobalt layer equal to 90 % volume. No softening of diamond brands AC 20, IRV, IMS, A5 can be noticed. Such a FGM approach is recomended for manufacturing grinding tools. 100

I

I { I I I I I I I I I I I I I I I » r r I I

0,0 1.0 2.0 3.0 4.0 5.0 6.0 7.0 8.0 J^J^ Fig.2. Cobalt distribution on the thickness of diamond containing FGM with (TiC+Co) binder, produced at the different relationship of mass layers mco/mTi+c+coi 28/56(1); 20/56 (2); 13/56(3).

3.3. Application of natural diamond One of the peculiarities of the natural diamond is a considerably lower content of admixtures Ni, Mn, oth. as compared to the synthetic diamond. It is this feature that determines its' raised resistance to the action of high temperature in the combustion wave. Figure 3 shows the dependencies of the recuperated synthetic and natural diamond strength on the mass proportion of the charge layers mi/m2 with diamond concentration equal to 25 % vol. on the example of the bi-layered composite with the ceramic binder (Ti,Mo)Ca. The strength of diamond grains is also affected by mi/m2 is the composites with the binder of NiAl, TiB+Ti, TiC+TiAl, TiBz+Si.

288 The tests of the produced materials for the abrasive wear showed that the composition 20 % TiB2 + 25 % Si-natural Dia possesses the highest wear resistance. The wear resistance index is equal to 206.

50

P,N 40 30

initial n a t u r a l diamond

initial synthetic diamond

20: 10- J — I — 1 — 1 — I — I — I — 1 — I — I — I — I — I — r

0.0

0.8

0.4 ml/inZ

Fig 3

Dependencies of the diamond grain strength in the FGM SHS-composition on the mass ratio of the Ti+Mo+C-f25 vol % diamond (mj) and Ti+Mo+CC/Wj) layers

CONCLUSIONS The paper presents the new technological opportunities of SHS for the production of diamond containing ceramic multilayered and FGM composites.

ACKNOLEDGEMENT The present work was carried out within the frame of the joint investigation program between the Moscow Steel and Alloys Institute, Ryukoku University and "Tomei Diamond" Inc., and supported by the Russian Ministry of Science and Technical Policy.

REFERENCES 1. E.A.Levashov, I.P.Borovinskya, A.S.Rogachov, M.Koizumi, M.Ohyanagy and S.Hosomi. Intern. Journal of SHS, vol.2, 2 (1994) 189. 2. E.A.Levashov, A.V.Trotsuk, I.P.Borovinskaya, M.Ohyanagy, M.Koizumi and S.Hosomi. Abstr. of the 3rd Intern. Symp. on Structural and Functional Gradient Materials, Laus'atine, Switzerland, October 10-12 (1994) 57. 3. E.A.Levashov, B.V.Vijushkov, E.V.Shtanskaya, I.P.Borovinskaya, M.Ohyanagy, S.Hosomi and M.Koizumi. Intern. Journal of SHS, vol.3, 4 (1994), 287. 4. M.Ohyanagy, T.Yoshikawa, M.Koizumi, S.Hosomi, E.A.Levashov and I.P.Borovinskaya. Intern. Journal of SHS, vol. 4, 4 (1995) 387.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

289

Graded Dispersion of Diamond in TiB2-based Cermet by SHS/Dynamic Pseudo Isostatic Compaction (DPIC) M.Ohyanagi^), T.Tsujikami^), M.Koizumil), S.Hosomi^), E-A-Levashov"^) and I.P.Borovinskaya^) l)Dept. of Materials Chemistry, Ryukoku University, Japan. 2)Dept. of Mechanical and System Engineering, Ryukoku University, Japan. 3)Tomei Diamond Co., Ltd., Japan. 4)Center of SHS, Moscow Steel & Alloys Institute, Russia. 5)Russian Academy of Sciences, ISMAN, Russia

Graded dispersion of diamond in TiB2/Si cermet (70vol% diamond layer / 40vol% diamond layer / matrix) was carried out by dynamic pseudo isostatic compaction (DPIC) just after self- propagating high temperature synthesis (SHS). The DPIC was performed using commercial casting sand as the pressure transmitting medium for the densification of cermet. The process enabled to simultaneously synthesize and densify the cermet matrix within a few minutes. Diamond (an average particle size, approximately 30)j,m) mixed with the reactant was fixed in the matrix produced after the SHS. The maximum combustion temperatures were controlled to be approximately 2000 K to prevent the diamond to graphite transformation. X- ray diffraction patterns and Raman spectra indicated that the diamond was embedded in the matrix mostly with no damage. The diamond particles were strongly fixed in the matrix even after lapping with a diamond abrasive. The primitive calculation for residual stress based on graded structure of diamond in the matrix was also performed. 1. INTRODUCTION Diamond being optical, high thermal conductive, semi- conductive and very hard materials, itself is widely expected to become an industrial materials of twenty first century. Metal alloys and ceramic materials containing diamond with such a useful feature will be also developed as new materials. The bonding is one approach to fabricate diamond composite materials. Instantaneous bonding of diamond and metal has been studied by new cost- effective SHS method. Another approach is to disperse diamond particles into the material matrices. ' We, herein, perform the latter approach. In case of using SHS process, the reaction does not occur in diamond- highly concentrated reactant without external energy support because the diamond works as the reaction diluent. The diamond- locally dispersed matrix reactant containing the side layer without diamond as an energy supply source is required to synthesize the material. However, in two layer system consisting of diamond highly- dispersed and the matrix layers, there would be residual stress in the interface. So, graded dispersion of diamond m the matrix reactant is required to fabricate diamond-containing materials. On the other hand, diamond with a meta- stable structure usually transforms into graphite by exposing for long time above 1800K even under an inert gas atmosphere or in vacuum. Accordingly, the fabrication of diamond- dispersed ceramics with high melting temperature such as TiC, TiB2 is considered to be difficult because the sintering and the densification are often performed over 2000K, conventionally using hot press and hot isostatic press. However, the cost- effective short processing, one of the advantages in SHS, is very effective to prevent diamond to graphhe transformation in course of the SHS processing even if the the maximum combustion temperature raises up to over 2000K. The SHS products

290 in highly dense form can be also fabricated using a combination technique of this SHS and an external pressure such as hot pressing, hot isostatic pressing (HIP), pseudo- HIP, explosive consolidating, and high-velocity forging. ^^ Dynamic pseudo isostatic compaction (DPIC) was applied for the hot and partially molten samples after the SHS of the matrix materials. The DPIC technique using commercial casting sand as the pressure transmitting medium was developed for the densification of cermets by Russian scientist and was also applied for the fabrication of diamond- dispersed cermets by the DPIC apparatus newly developed. In the equipment, a slender sheet of carbon ribbon as the heat device only for ignition is embedded with a sample in commercial casting sand, which is contained in a pressure vessel. The compaction was performed by quickly pressing the sand containing the sample just after the SHS reaction. A pseudo- HIP using sand as the pressure medium is well- known to cause pseudo isostatic pressing. Similarly, herein, the pressing of the sand by high speed auto-pressing machine was applied to perform the DPIC. One of the objective of this research is to fabricate diamond- gradually dispersed cermets by the combination technique of SHS for short time processing and following dynamic compaction for densification. The other is to support for the fabrication by the calculation of residual stress based on graded structure of diamond in the matrix. 2. EXPERIMENT PROCEDUE 2.1 Evaluation of residual stresses using finite element method The finite element method (FEM) has been used to evaluate thermal residual stresses at interface of Diamond/TiB2/Si composites. Axisymetric cylindrical specimens were used, allowing two dimensional models to be employed. A model system composed three layers, Diamond, Diamond/(TiB2/Si) and TiB2/Si. The finite element analysis was performed using the original developed software SACOM for composite materials^^'-^^. In this simulation, thermal residual stresses, considering only elastic behavior were calculated, and the Diamond/riB2/Si composites was cooled from the assumed high temperature service (2000K) to room temperature (293K). Time and temperature dependent properties were neglected. Table 1 shows physical and mechanical properties of Diamond, TiB2 and Si relevant to the calculation. The specimen's dimensions were 5 nmi long and 16 mm in diameter. Constitutive properties for the composite material mterlayer were computed using a rule-of-mixtures.

Diamond TiB2 Si

Table 1 physical and mechanical properties of Diamond, TiB2 and Si Elastic modulus E Coef. of thermal Poisson s ratio v (GPa) expansion (K' ) 0.3 1050 1.00X10"^ 6.39X10"^ 0.2 365 4.2X10-^ 105 0.3

2.2 Materials and procedure The reagents used were elemental powders of Ti (an average particle size: approximately 22.5 |j,m, >99.5%, Osaka Titan Inc.), B (- 325 mesh, >99%, High Purity Chemicals Laboratory Inc), Si (-10 |am, >99.9%, High Purity Chemicals Laboratory Inc) and C (diamond: artificial, an average particle size: approximately 30.0 )j-m, >99.9%, Tomei Diamond Co. Ltd.). The reactant powders were weighed out m the proper stoichiometric proportions, keeping a constant equimolar ratio of Ti- 2B. The mixing ratio of Ti- 2B/Si was kept at 70/30 in vol%. Diamond powder was added so that it occupied 40 to 70 vol% of the reactants in the locally dispersed layer. The powder batches were mixed dry by auto agate mortar for half an hour. Cylindrical compacts (approxunately 16 mm in diameter and 25 mm long) were formed in a stainless steel die with double- acting rams so that diamond powder was dispersed in the 1/3, 1/5, 1/8 bottom layer of the compacts. In case of the graded dispersion, the composition of each layer was mixed individually, then layered in the steel die in the green pellet. The compacts were pressed uniaxially at the pressure of

291 approximately 5.0 MPa. DPIC was performed by a special SHS/DPIC equipment.^'^^ A stainless steel pressure resistant vessel of 30 mm inside diameter, 100 mm outside diameter and 60 mm deep was filled with commercial casting sand. An ignition heating device made of the slender carbon ribbon was placed on top surface of the sample. Each compact was ignited by a passage of current through the carbon heating ribbon under an atmosphere. In the delay time for 1 to 3 sec after the reaction, the sand containing the sample was pressed by a piston from top by using high speed auto- pressing machine (a piston moving velocity: 60 mm/s). The total applied pressure was approximately 255 MPa. The pressure was maintained for 10 sec after the pressing. A temperature profile was measured by a thermocouple (W- Re5%AV-Re 26%)- ^he vohage outputs from the thermocouple and the transducer indicator for pressure were monitored using a data acquisition recorder (OMUNIACE RT3200, highest sampling rate: 200kHz, NEC). This recorder made h possible to measure and store the data during the SHS/DPIC. The product surfaces lapped using a diamond abrasive were observed by SEM (JSM T- 330A, JEOL) and identified by X- ray diffraction equipment (RAD- C system, Rigaku Inc). 3. RESULTS AND DISCUSSION 3.1 Evaluation of residual stresses using finite element method The finite element method (here, triangle linear element method) has been used to evaluate thermal residual stresses at interface of Diamond/TiB2/Si composhes. Axisymmetric cylindrical specimens were used, allowing two dunensional models to be employed. A model system composed three layers. Diamond, Diamond/(TiB2/Si) and TiB2/Si. Figure 1 shows the scheme of the Diamond/riB2/Si composites. Figure 2 shows the axisymmetric mesh model for the Diamond/TiB2/Si composites. The finhe element mesh consisted of 12857 elements and 6610 nodes. At first, the effect of varying the diamond volume fraction of the top layer were studied. Different diamond volume fractions of top layer were changed from 50% to 100% for the evaluation. Constitutive properties for the composite material interlayer were computed using the rule- of- mixtures. The diamond volume fraction of middle layer is 40%, and the TiB2 volume fraction is 50% in TiB2/Si matrix in all layers. Figure 3 shows the behavior of maximum radial, axial and shear stresses of the top layer whh varying the diamond volume fraction. It is found that the stresses become larger with increasing diamond volume fraction. The contour plot in Figure 4 shows the distribution of the radial stress when the diamond volume fraction of top layer is 70%. Since the thermal expansion coefficient of diamond is smaller than that of Diamond/rriB2/Si, there are compressive in the Diamond and tensile in Diamond/riB2/Si. The contour plot in Figure 5 shows the distribution of the axial stress. The maximum stress concentration occurs at the free surface of the top layer. The contour plot m Figure 6 shows the distribution of the shear stress. The maximum stress concentration occurs at the interface between the top layer and the middle layer on the free surface side. The results suggest that the stresses become larger as a result of the property mismatch.

Figure 1. Scheme of Diamond/TiB 2/Si composite.

Figure 2. Finite element mesh for Diamond/ TiB2/Si composite.

292 3500 3000 ^2500 a.

D . • : Radial stress A , A : Axial stress : Shear stress

^2000 1^1500

11000 (D 500 •S

0

i-500 m E-IOOO

1-1500 ^2000 -2500 -3000 -3500

50 60 70 80 90 Diamond volume fraction of tfie top layer

Figure 3. The maximum axial, radial and shear stress of diamond layer with varying the diamond volume fraction.

Figure 4. The distribution of the radial stresses for the material when the diamond volumefractionof the top layer is 70%.

Figure 5. The distribution of the axial stresses for the material when the diamond volume fraction of the top layer is 70%.

Figure 6. The distribution of the shear stresses for the material when the diamond volume fraction of the layer is 70%.

Next, in two layers system consisting of diamond- dispersed layer and only matrix, effect of thickness of the diamond layer on the residual stresses was evaluated. In radial stress, the maximum compressive and tensile stresses in the diamond layer and in the interface of the layers became smaller in the thinner layer. In three layers system consisting of diamond- high, low dispersed layers (40 and 70 vol%) and matrix, influence of insertion of middle layer on the maximum stresses was compared with the two layers system. The residual stresses in each layer seems to become reduced by the insertion of middle layer. 3.2 Graded dispersion of diamond in TiB2/Si-diamond system by SHS/DPIC The SHS/DPIC of diamond- gradually dispersed TiB2/Si cermets was performed. In case of fixing diamond in cermet matrix, there are two methods. One is the physical fixing just like a diamond ring, and the other is the chemical fixing by covalent bond between diamond and metal in the cermet. In this work, we studied the latter system and considered that a metal carbide is very suitable as an interlayer to form the covalent bond.. There are many metals in periodic table to form the metal carbide. However, the

293 metals are desirable to have low melting temperature for suppression of the graphitization of diamond and acceleration of the densification in rather low temperature. Consequently Si were chosen as a metal portion for the cermets. The combustion maximum temperatures {Tmax) were controlled to be approximately 2000 K to prevent the diamond to graphite transformation. 255 MPa of the pressure was applied to the sand filled in the pressure vessel in the delay time for 1- 3 sec after the reactions. The pressure was maintained for 10 sec after the pressing. The sample could be taken out from the reaction vessel within a couple of minutes after ignition. The sand works as the pressure transmitting medium,

Figure 7. SEM photograph of lapped surface of diamond layer (TiB2/Si/diamond70vol%), BEL

(a) Cross section.

Figure 8. SEM photograph of lapped cross section in two layersTiB2/Si-Diamond(70vol%) and the matrix, (ratio of each thickness: 1 to 4)

(b) Interface between diamond layers, 70 and 40vol%.

(c) Interface between diamond layer, 40vol% and the matrix Figure 9. SEM photograph of lapped cross section in three layers of TiB2/Si- diamond (70vol%) and TiB2/Sidiamond (40vol%), and the matrix.

294 which suggests that Tmax to be measured is difficuh to reach the adiabatic temperature {Tad) of the reaction system. Figure 7 shows SEM photographs of the lapped surfaces of the specunens in TiB2/Si/diamond(70vol%), (a) SEI, (b) BEL The darkest part in the each SEM corresponds to diamond. The adhesion between diamond and the each matrix seems to be smooth and good. In EPMA analysis, the X- ray Kp spectrum of Si measured on the diamond surface in lapped cross section of TiB2/Si/diamond cermet agreed with the spectrum of SiC. The spectrum of Si in the cermet matrix also agreed with that in Si wafer. These suggest the formation of strong covalent bond layer, SiC between diamond and Si. The Raman spectrum based on diamond shows an intense and sharp peak in 1335 c m - 1 . The X- ray diffraction patterns also indicated in any cases that the diamond was embedded in the each matrix mostly with no damage. Figure 8 shows SEM photograph of lapped cross section in two layers composite of TiB2/Si- Diamond (70vol%) and the matrix (ratio of each thickness: 1 to 4). The lateral crack along the interface and almost perpendicular crack were observed in the specunen. Figure 9 shows also SEM photograph of lapped cross section in three layers of TiB2/Si- Diamond (70vol%), TiB2/Si- Diamond (40vol%), and the matrix (ratio of each thickness: 1 to 1 to 6). In each interface, the crack-based on residual stress was not observed. 4. CONCLUSION Graded materials of diamond- dispersed TiB2/Si composite were fabricated by SHS/DPIC method. Each diamond fraction in the graded composite was 0, 70vol% in two layers system, and 0, 40, 70vol% in three layers system. In the two layers composite, some cracks occur in the matrix and also in the interface of the layers. On the other hand, in the three layers, crack did not occur in the diamond layers and did not often occur in the interface between middle layer containing 40vol% of diamond and matrix. The tendency corresponds to the results of evaluation of residual stresses using finite element method. ACKNOWLEDGMENT This work was performed in High- Tech. Research Center (HRC) of Ryukoku University. One of the authors, M.Ohyanagi thanks for partial supports by Grand- in Aid for Scientific Research on Priority Area Physics and Chemistry of Functionally Graded Materials , The Ministry and Education, Science and Sports and Culture, and also by the Science Research Promotion Fund from Japan Private School Promotion Foundation.

REFERENCES 1. M. Ohyanagi, M. Koizumi et al., Am. Cer. Soc. Bull., 72, 86 (1993) 2. E.A. Levashov, LP. Borovinskaya, A.S. Rogachov, M. Koizumi, M. Ohyanagi, S. Hosomi, Intern. J. SHS, 2, 189 (1994) 3. E.A. Levashov, B.V. Vijushkov, E.V. Shtanskaya, LP. Borovinskaya, M. Ohyanagi, S. Hosomi, M. Koizumi, Intern. J. SNS, 3, 287 (1994) 4. M. Ohyanagi, M. Koizumi, S. Hosomi, E.A. Levashov, K.L. Padyukov, LP. Borovinskaya et al.. Trans. Mat. Res. Soc. Jpn., 14A, 685 (1994) 5. A G . Merzthanov, Ceram. International, 21, 371 (1995) 6. O.Yamada, Y.Miyamoto, and M.Koizumi, Am. Ceram. Soc. Bull, 64, 319 (1985) 7. Y. Miyamoto, Am. Ceram. Soc Bull, 69, 686 (1990) 8. P.H. Shingu, K.N. Ishihara, F. Ghonome, T. Hyakawa, M. Abe and K. Tagushi, Pro. of 1st US- JAPAN Work-shop on Combustion Synthesis (Tsukuba), 65 (1990) 9. L.J. Kecskes, T. Kottke, and P.H. Netherwood, J.Am. Chem.Soc ,73, 383 (1990) 10.J.C. LaSalvia, L.W. Meyer, and M.A Meyers, J.Am.Chem.Soc ,75, 592 (1992) 11. M. Zako, T. Ttujikami, Development of Personal Computer Program of Stress Analysis for Composite Materials, Journal of the society of material science, 38, No.438, (1990) 12. M. Zako, T. Ttujikami, M. Hibino, M. Ichikawa, M. Uemura, Development of Structure Design System for Composites, Reinforced plastics, 38, No.2, (1992) 13. M.Ohyanagi, M.Koizumi, S.Hosomi, E.ALevashov and I.P.Borovinskaya, Intern. J. SHS, 4, 387 (1995)

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

295

Annealing of cermic/metal graded materials fabricated by SHS/QP method A. N. Pityulin% Z. Y. Fu^ M. J. J i n \ R. Z. Yuan^ and A. G. Merzhanov' ^ Institute of Structural Macrokinetics, Russian Chernogolovka, 142432 Moscow Region, Russia

Academy

of

Sciences,

^ State Key Lab of Advanced Technology for Materials Synthesis and Processing, Wuhan University of Technology, Wuhan 430070, China

SHS/QP is an efficient method for preparing FGM. But in synthesis process, high temperature and high velocity may result in some results not expected by us, which as a consequence will affect the FGM's performance. Al so, repeatability of the concentration distribution is not good. In this work, FGMs were prepared by SHS/QP and heat treating, which had very thin gradient layer and wide range concentration distribution. The manufacturing was carried out as following: the pellet including mixture-metal layers was processed by SHS/QP, and then the product was treated longer than two hours at 700 "C -1200 °C . Quantitative and qualitative analyses were carried out to different points in the gradient layer. The result indicates that the thickness of the gradient layer is decided by the heat treatment to a great extent. The chemical and phase composition, and the connection between grains show no great change.

1. INTRODUCTION Modern industry needs hard alloys with both good wear-resistance and highstrength. Ceramic/metal layered materials, with the structures as shown in Figure 1, are very promising. However, because of the difference of the heat expansion coefficients between ceramic and metal, it is very difficult to sinter the two parts together by traditional methods. SHS compaction (hydraulic pressing) or SHS/QP method with its features shows potential in the fabrication of such layered materials [1,2]. In their experiments, Pityulin and his coworkers obtained two kinds of layered structure: symmetrical profiles (SYGMA-1) and nonsymmetrical profiles (SYGMA-2) [1]. SYGMA-2 materials, one layer is pure metal and the other side is ceramic (Figure 1), can have very good wear-resistance on the ceramic side and good overall ductility owing to the metal part. In this paper, SYGMA-2 with Ti as the metal side and TiB-45wt%Ti

296 as the ceramic side made by SHS/QP was heated at different temperatures. Variations of structures and properties of the sample with annealing temperatures were studied.

2. EXPERIMENTAL PROCEDURES Ti (grain size < 60|im, purity > 99%) and amorphous B (grain size < lOfam, purity > 94%) powders were used in the experiment. Thoroughly mixed powders with determined composition were pressed into plate (70mmx70mmxl0mm) with 50% relative density, which was then put into a special die as shown in Figure 2. Between the raw plates and die, there are Si02 powders with average particle size 0.5mm, which serve both to protect the die and to transform the mechanical force from the hydraulic press to the sample in a pseudo-isostatic way. The SHS reaction was ignited by a tungsten coil with a short electric pulse. Thermocouples were used to determine the reaction temperatures and propagating time. Immediately after reaction, the hot product was pressed by a lOOOkg/cm^ force. By this way dense layered sample with determined structure, one side Ti and the other side TiB-45wt%Ti, as shown in Figure 1 was produced, which was then cut and polished into small strips (5x5x30mm). The small strips were heated in a vacuum stove (vacuum degree 10'^ Torr). Maximum heating temperature and time were 1500 °C and 5 hours respectively. Hardness and ultimate bending strength were tested by standard procedures. Structure and element distribution were analyzed by EPMA and SEM.

3. RESULT AND DISCUSSION Ti and the hard alloy linked well with each other in the two-layer sample made by SHS/QP as shown in Figure 3. Porosity of Ti side is 0.5%.There is no pore in the intermediate layer. XRD proves that the two layers are composed of TiB-Ti and Ti respectively. Gradient intermediate layer between Ti and the hard alloy is lOOfam in thickness as shown in Figure 4, which is independent of thickness ratio of the two layers. The sample was etched with 36% HCl. Etched structures of the sample were shown in Figure 5. Structure of the hard alloy is quite uniform, in which TiB grains are in a long-flake shape with maximum length up to 30)Lim. Structure in the gradient layer is not uniform. Thickness of the gradient layer increases with annealing of the sample as shown in Figure 6. There are three regions in the chart. At low-temperature (< 900 °C ), thickness of the gradient layer (about lOOfim) does not change with annealing temperature. Thickness of the gradient layer will increase to 350~400|im, when the annealing temperature rised to 1200 "C , but the sample's shape does not change. When the annealing temperatures are higher than 1300 °C , the thickness of the gradient layer will increase obviously. At an annealing temperature 1 500 °C , Ti layer will

297 totally melted and migrated into the hard alloy, which forms a 4mm gradient layer and makes the sample shrink to a certaim extent. Distribution of element concentration of the sample treated at 1500 °C is shown in Figure 7. The gradinet distribution is similar to ordinary FGMs made by other methods [3,4]. Hardness variations across the sample are shown in Figure 8. The sample presents a sudden change in hardness distribution, when it is heat treated at 700 °C . Hardness changes gradiently along the sample, when the sample is heat treated at 1500 °C . Strength of the sample increases with treatment temperature. The ultimate bending strength can be rised by 25% as shown in Figure 9.

4. CONCLUSIONS Annealing of the SYGMA-2 type sample can change structure and thickness of the gradient intermediate layer. 1500 °C annealing results in the migration of melted Ti totally into the hard alloy, which forms a 4mm gradient layer. Sudden and gradient variations in hardness distributions across the sample were observed after 700 °C and 1500 °C teratment, respectively. Ultimate bending strength of the sample increases with annealing temperature.

REFERENCES 1. A. N. Pityulin, Y. V. Bogatov and A. S. Rogachev, Inter. J. SHS, 1(1992)111 2. Z. Y. Fu, W. M. Wang, H. Wang, R. Z. Yuan and Z. A. Munir, Inter. J. SHS, 2(1993)307 3. M. Koizumi, Ceramic Transactions, 34(1993)3 4. Z. Y. Fu, R. Z. Yuan and Z. L. Yang, Proceed. 1st Inter. Sym. on FGM, Sendai, (1990)175

298

Reactants

Gradient layer

'Mh Ceramic

Thermocouple

Metal SiOa powers

Figure 1. Schematic representation of ceramic/metal layered materials

Figure 2. Schematic representation of SHS/QP die

3 alloy

0

Figure 3. Structure of two-layered sample by SHS/QP

0. 2 0. 4 0. 6 0. 8 1. 0 X>mm

Figure 4. Element concentration profiles in two-layered sample

299

Figure 5. SEM micrographs of sample. a. Gradient region b. Region 1mm from the gradient layer c. Region 3mm from the gradient layer

lOOi

300 600 900 1200 1500 T/C Figure 6. Dependence of thickness of gradient layer on annealing temperature

Figure 7. Element concentration profiles in sample treated at 1500 °C

300 ^

10 o

»

1

y\

8

r-H

X CO

> X

• —700°C 4

• ^ ^ l

t

. 1

o—1500°c1

. 2 X)inm

3

4

Figure 8. Hardness variations across samples treated at different temperatures

0

300

600

900

1200 1500

Figure 9. Effect of treating temperature on ultimate bending strength

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

301

Thermodynamic calculation and processing of TiB2-Cu FGM C.C.Ge, Z. X. Wang and W. B. Cao Laboratory of Special Ceramics & Powder Metallurgy, University of Science and Technology Beijing, Beijing 100083, P. R. CHINA Thermodynamic calculation and SHS/HIP were successfully used for making TiB2-Cu FGM without macro-defects from element and diluent powders. In this case, SHS/HIP was used not only to create "a chemical oven" for densification of FGM, but also to combustion synthesize the foundmental constitute of FGM~TiB2. 1. INTRODUCTION Since 1984 Japanese scientists proposed the idea of "Functionally Graded Materials" (FGM), research on FGM is developing rapidly. For commercial applications of FGM, problems on design and processing have to be solved. In designing of FGM, not only the thermal stress in FGM under service conditions should be minimized, but also the thermal stress during processing should be reduced in order that macro-defects of FGM can be prevented. SHS/HIP(self-propagating high-temperature synthesis/hot isostatic processing) is an advanced technology proposed by Y. Miyamoto and M. Koizumi[l] for making FGM, which has following advantages: 1. high reaction rate and short duration at high temperatures of SHS/HIP process are very beneficial for keeping the designed constitute gradient in FGM; 2. Simultaneous synthesis and densification are realized; 3. Product with large dimensions or complicated shapes can be made; A series of FGM including TiB2-Ni, TiC-Ni, MoSi2-SiC/TiAl, Cr3C2-TiC etc. has been successfully made with SHS/HIP[2]. In most of these works ,the exothermic reaction of SHS was used to form a chemical oven for densification of FGM, while the fundamental constitutes (such as Cr3C2 in Cr3C2-Ni FGM, MoSi2, SiC in MoSi2-SiC/TiAl FGM etc.) in FGM were used as commercial raw materials which had been beforehand synthesized. The problems of forming and preventing macrodefects in FGM samples have not been reported. This work was undertaken for making TiB2-Cu FGM without macro-defects from element and diluent powders through thermodynamic calculation and SHS/HIP. SHS/HIP was used in this case not only to create "a chemical oven" for densification of FGM, but also to combustion-synthesize the fundamental constitute of FGM-TiB2. 2. EXPERIMENTAL PROCEDURES Processing of TiB2-Cu FGM was based on our work of non-graded homogeneous TiB2-Cu composites by SHS/HIP technology for investigating the effects of processing parameters on the microstructures, constitutes, properties and sinterability of as-synthesized products[3].

302

Commercial powders of Ti (~ 42 jn m), amorphous B( ~ 5 jn m) and Cu( ~ 7 ja m),and selfmade SHS TiBj powder as diluent are used as raw materials. As the first step, graded green compacts with a dimension of (^ 17mm x lOmm and a relative density of - 60% of the theoretical value were made by stacking of mixed powders in the die in proper order of Cu content. Green compacts were combustion-synthesized in Ar under 5MPa in a combustion chamber of a self-designed SHS/HIP assembly (SHA). Preliminary experiment in making 6-layered FGM in SHS assembly led to frustum samples, expanding in TiBj-rich side and shrinking in Cu-rich side, while in making 8- and 11-layered FGM serious cracks, warping and delamination occurred in samples. These defects of samples were attributed to high thermal stress due to great difference in combustion temperatures between layers. Thermodynamic calculation was carried out before the following experiments with the aim to adjust the combustion temperature of different layers for reducing the thermal stress during processing and preventing the macro-defects of products. On the basis of our previous work on analysis model and mathematical calculation for design of thermal-relaxed TiB2-Cu FGM[4], graded green compacts with dimensions of cj) 17 X 10mm and cj) 26 x 10mm and with - 60% of the theoretical density were made by stacking of mixed powders in the die in proper order of Cu content, according to the optimized parameters: Cu content changes from 0-100%, thickness of graded layers t = 8mm, thickness of surface layers = 1mm (TiB2 layer and Cu layer respectively), thickness of each graded layer was designed with the volume distribution function / = (

-y,

while

constitute distribution factor p=0.8. Both green compacts with number of layers n=ll and n=15 were pressed and encapsulated with glass and embedded into a Ti ignition agent in a graphite crucible under nitrogen pressure. In order to make the process less-expensive, the gas pressures of 5MPa and 1 IMPa were used for different samples. The synthesized TiB2-Cu FGM samples were longitudinally cut, polished and observed with SEM. The distributions of element Ti and Cu were line-scanned and area-scanned with EDX. The densities of layers were measured with image-analysis method. 3. RESULTS AND DISCUSSION 3.1. Thermodynamic calculation While the SHS process has high reaction rate and short duration, it can be regarded as an adiabatic process in our calculation:

40 60 80 Cu content, wt%

100

Fig. 1 Variation of adiabatic temperatures with Cu content for TiB2-Cu system

303

Fig. 1 is the calculation result for variation of adiabatic temperatures with Cu content in Ti2B-Cu system : Ti + 2B + bCu = TiB2 + bCu

(1)

where b is the mole content of Cu. During SHS, this exothermic reaction consists of following processes: (1) melting of Cu (endothermic) (2) melting of Ti(endothermic) (3) Phase transformation of Ti (endothermic): a -Ti(hcp) ^^^^\ p -Ti(bcc); (4) Formation of TiB2 (strong exothermic) Fig.2 is the variation of enthalpy with temperature for Ti-2B-Cu system from our calculation according to:

-A//;,,^^=£;c,(5yr+*Atf„,, (2)

where: -AH^ j ^ .-the formation enthalpy of TiB2 at 298K; C {s) -heat capacity under constant pressure for TiB2(solid state); A//^ ^„ -melting enthalpy of Cu.

Ti (P) + 2B( S)

r^-'Tif^

+ 2B (S)

o

e

Ti (a) + 2B (S)

>> u. (0

ri

TiB2+Cu T

\ Cu(S)

^^-^ Cu(l)

A Hf"

W

At,Ti

im.Cu

Tm,Ti

Temperature, K

Fig.2 Variation of enthalpy with temperature for Ti-2B-Cu system For adjusting the combustion temperature of different layers in FGM, we use the reaction product phase-TiB2 as diluent, then: Ti + 2B + aTiB2 + bCu=(l+a)TiB2 + bCu and

(3)

304

where a is mole content of TiB2. From equation (4) it is shown that through adjusting the metal and diluent content of the reaction mixture, the control of combustion temperature can be realized.

I Metal Cu Content , wt%

Fig.3. Composition design of TiB2-Cu FGM compositions at different Tc

H

20

40

Time, sec

Fig.4. Tc profiles of some kinds of compositions according to Tc=2000k Fig.3 is our calculation results for composition design of TiB2 -Cu FGM at different combustion temperatures. Fig.4 is the experimental results for combustion temperature profiles of some kinds of compositions according to Tc = 2000K and corresponding design of compositions of reaction mixtures. Fig. 5 is the comparison of measured temperatures(Tm) and designed temperature (Td). It is proved that through thermodynamic calculation and adjust of constitutes ratio of graded layers, control of combustion temperature and composition design can be achieved. 3.2. Processing of TiBj-Cu FGM in self-designed combustion chamber 11-layered and 15-layered TiBj-Cu FGM samples with dimensions (|) 17 x 10mm free of macro-defects were successfully processed with SHS/HIP technology under gas pressure of

305 5MPa in self-designed combustion chamber. It is seen that through reasonable designing the constitutes ratio of each layer, SHS of different layers could be controlled in the same combustion temperature region and similar shrinking behavior between layers , this led to SHS/HIP TiB2-Cu FGM which is free of macro-defects.

p

H ?3

1 Metal Content, wt%

Fig. 5. Comparison between measured values(TJ and designed values (Td) of combustion temperatures. Fig.6 is the area-scanning distribution of element Cu in the polished longitudinal crosssection of TiB2-Cu FGM samples processed with SHS/HIP. 3.3. Processing of TiB^-Cu FGM in HIP apparatus R120 Samples with dimensions cj) 26 x 10mm were processed with SHS/HIP under gas pressure of llMPa in HIP apparatus R120. The dimensions of the high-pressure chamber is cj) 120 x 350mm. Several large samples can be processed in each batch.

Fig.6 Area-scanning of distribution of element Cu. 3.4. Comparison of relative densities of TiBj-Cu FGM Image analysis for density measurement was made on the longitudinal cross-section of 11layered TiB2-Cu FGM processed under 5MPa (in self-designed combustion chamber) and processed under 1 IMPa in HIP apparatus R120. The results are shown in table 1. It is noticed that relative density of >94% is achieved at the rich-Cu end of samples processed under llMPa in HIP-apparatus.R120. The porosities of these samples are 50% lower than the porosities of rich-Cu end of samples processed under 5MPa. For the low-

306 Image analysis for density measurement was made on the longitudinal cross-section of 11layered TiB2-Cu FGM processed under 5MPa (in self-designed combustion chamber) and processed under 1 IMPa in HIP apparatus R120. The results are shown in table 1. It is noticed that relative density of >94% is achieved at the rich-Cu end of samples processed under llMPa in HIP-apparatus R120. The porosities of these samples are 50% lower than the porosities of rich-Cu end of samples processed under 5MPa. For the lowtemperature end of thermal-relaxed FGM, higher density favors heat conductivity and strength of the material. Though the relative density at the rich-TiB2 end of samples processed under 1 IMPa increased inconsiderably comparing with samples processed under 5MPa, yet for thermal-relaxed FGM some porosity is beneficial for heat insulation ability of the ceramic end. The processed TiB2-Cu FGM may fulfill the requirement of relaxation of thermal stress for certain applications. Table 1 Measured porosity is various regions along the longitudinal cross-section

Cu Content low

1

high

Porosities in Synthesized Porosities in Synthesized! Products Made in HIP(%) Products Made in Selfmade Apparatus(%) | 51.84 50.89 35.82 41.25 13.85 25.88 5.95 13 17

4.CONCLUSIONS Through thermodynamic calculation for adjusting the combustion temperatures of different layer of TiB2-Cu FGM, SHS/HIP was successfully used both for combustion-synthesis of the fundamental FGM constitute and for densification of FGM in one step, and TiB2-Cu FGM samples without macro-defects has been obtained. 5. ACKNOWLEDGMENT This work is supported by China National Natural Science Foundation, The Doctorial Program Fundation of State Education Commission, and National Committee of High Technology New Materials. REFERENCES 1. Y. Miyamoto and M. Koizumi, Proc. Int. Symp. on Sintering '87, Tokyo, Elsevier (1988), 511. 2. Y. Miyamoto, Int. J. SHS, V.l, No. 3, 1992 3. Z. X. Wang, Dissertation of Univ. Sci. Tech. Beijing, 1995. 4. Z. X. Wang, C. C. Ge, W.B. Cao and X. D. Zhang, Proc. 4th. Int. Symp. on FGM, Tsukuaba, Elsevier (to be published).

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

307

Fabrication of Al-Cu system with functionally graded densitj^ profiles* R. Tu^ Q. Shen^ J.-S. Hua^ L.-M. Zhang^ R.-Z. Yuan^ ^ State Key Lab. of Materials S>^thesis and Processing, Wuhan University of Technology, Wuhan, P. R. China, 430070 ^ Laboratory of Shock Wave and Detonation Physics Research, Southwest Institute of Fluid Physics, Chengdu, P.R.China, 610003

To lower the temperature-enhancement inside the target material and achieve higher pressure and velocity is important in dynamic high-pressure technology. It can be realized through quasi-isentropic loading by density functionally graded materials (DFGM). A kind of .^-Cu DFGM was hot-pressed by powder stacking with adjusting soaked temperature, load interval and pressure et al. The FGM has quasi-continuous density variation along the thickness direction.

1. INTRODUCTION Recently, researches in the world have extended progresses in thermal stress relaxation FGMs[l] and energy conversion materials [2]. But they all can not be put into practice at present time. In the same teims, a new kind of FGM with density gradient(DFGM), which can be used in dynamic high-pressure technology has been prepared in America [3]. The DFGM was applied to the shock-wave loading technology in order to carry out quasiisentropic loading on target materials [4], from which can offer extreme experimental conditions of pressure or velocity^ for thermodynamics and dynamic physics study. In this paper, the Al-Cu system which has large density difference was chosen to prepare DFGM on a trial basis.

2. EXPERIMENTS 2.1. Materials designation The volume fraction C of gradient layers of Al-Cu FGMs was defined as the form [5]: C=(x/d)P

(1)

Where, d is the total thickness of FGM, x the location coordinate of any gradient layer, and P the distribution exponent. In order to gain the linear density distribution, P is fixed at 1.0. This work is supported by National Natural Science Foundation of China.

308 According to the phase diagram of Al-Cu system, it is noticed that many intermetallics will be formed in the range of 15-50 weight^'b.Al. The high brittleness of the intermetallics [6] and their phase transformation wiU degrade the materials. Therefore, the layer in the above range, i. e. the layer of 50 vol% .\1 + 50 \T)l?b Cu was eliminated from the design for avoiding intermetallics as much as possible. In addition, the eutectic point of Al-Cu is much low (821K) and a great deal of liquid phase would occurr at the temperature, so the sintering temperature was fixed at 800K or so. The temperature 800K is too low^ for the densification of pure copper powder. In order to protect liquid from flowing out and avoid the difficulty to sinter pure Cu layer densely, a solid Cu disk and Al disk were employed to replace the pure copper and pure aluminium powder layer respectively. The designed compositional distribution of materials is given in Table 1. Table 1 Composition of the .Al-Cu system FGM layers 1 Al voI% 0 .\latom% 0

2 10 7.31

3 4 5 6 7 8 9 10 20 30 40 (50) 60 70 80 90 100 15.06 23.31 32.10 41.49 51.55 62.66 73.94 86.46 100

2.2. Materials preparation Commercialh^ available average particle size 75 j.im and high-purit>^ 99.5% aluminium and copper powders were used as raw materials. The ground .Al-Cu mixtures were loaded in the grapliite mould as designed and hot pressed. The thickness of each layer is 0.5mm except that the Cu and M layers are 1mm. Previous expeiiments showed that melted metal flowed out when hot-pressed at 900K for 1.5 hours directfy. UTiile at 800K, it can not be sintered densely. To avoid the above phenomena, the experimental procedures as Fig.l was proceeded. It can be seen that the specimen was first heated to 900K and soaked for hatf an

300

1.0

1.5

2.0

t(hours)

Figure 1 Diagram of experimental procedure. hour with no applied pressure. Then the temperature was decreased to 800K and soaked for one hour with the pressure ot 15 Nfi^a. The as-hot-pressed FGM was cut along the diameter and polished. A scanning electron microscopy (S£M) was used for microstructure of the materials, and a electron probe microanahsis (EPMA) for its elemental macro linear

309 distiibution. Under the same sintering conditions, each layer of the FGM was prepared. Their densities were measured by the water-immersion technique.

3. RESULTS AND DISCUSSION Fig. 2 is the cross-section of Al-Cu densit>^ functionally graded material. It shows that there were no micro cracks, and the transition between layers was in a good state. Fig. 3 gives the result of elemental macro linear analysis. From Fig. 3, it can be seen that the Cu element content increases along FGM's thickness, while that of Al decreases. The relative density of gradient layers are low but the solid aluminium layer and copper layer are high, which due to the oxidization of aluminium powder on the surface [7]. It can reject the reaction of aluminium and copper. The detailed anah'sis of oxide content in aluminium particles will be

Figure 2 Cross-section of .Al-Cu DFGM. 8000

0

1

2

3 A D(mm)

5

6

7

8

Figure 3 Compositional distribution of Al-Cu DFGM along its thickness. performed in future. For investigating the densit\' variation of FGM, each layer was hot pressed under the same conditions. Theii* densities were measured and was shown in Fig. 4. It can be found that the density^ of FGM along the thickness direction increases with copper content. But relative densitv^ of the layer containing 10 vol^^b Al decreases obviously. Except for the oxidation of aluminium powder, it could be thought that aluminium and copper did

310 not fbrm the eutectic compound in that composition range. At this time, the densification temperature between aluminium and copper is so large that the relative densit\' decreases. When content of aluminium increases to 60 vol% (51.55 atom%), the relative density increased. Especiall>^ when the content of aluminium is over 70 atom%, low-melting-point compound is formed between aluminium and copper, and the relative densities of the layers increase. The FGM has quasi-continuous density- variation.

^0 60 Cu Content (vol*/.)

Figure 4 Designed and practised density of graded layers of .\1-Cu DFGM. One of the future works is to solve the problem of aluminium oxidation so that the graded layers can be sintered denseK. Adding some active metal, e. g. Mg and Sn, to the Al-Cu mixture may be a feasible ways.

4. CONCLUSIONS (1) A kind of Al-Cu density FGM coinciding with designed composition is prepared. (2) The low relative density of graded layers is due to the oxidation of aluminium powder on the surface.

REFERENCES

3. 4.

L. M. Zhang, M. Oomori, R. Z. Yuan and T. Hirai, J. Mater. Sci. Lett., 14(1995) 1620 M. Niino, M. Koizumi, in proceedings of 3rd International Symposium on Structural and Functional Gradient Materials, Lausanne, Switzerland, 1994, 601 L. M. Barker and D. D. Scott, SAND 84-0432 L. C. Chhabildas, in BuUetin of the 1995 APS Topic Conference on "Shock Compress of Condensed Matter", Scatter, Washington, 1995

311 5. 6. 7.

A. Kawasaki and R. Watanabe, J. Jpn. Soc. Powder Metall., 37(1990) 253 L. F. Mondolfo, Aluminium Alloys: Structure and Properties, Butterworths, Boston, 1976 C. N. Cochran and W. C. Sleppy, J. Electrochem. Soc, 108(1961) 984

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I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

313

AI2O3 to Ni-superalloy diffusion bonded FG-joints for high temperature applications* Liisa Heikinheimo^ Mika Siren^, Michael M. Gasikt* ^Technical Research Centre of Finland - VTT Manufacturing Technology, FIN-02150 Espoo, Finland ^'Helsinki University of Technology, FIN-02150 Espoo, Finland

The aim of this study is to manufacture alumina-superalloy joints for high temperature applications using direct diffusion bonding process with metallic interlayers and with functionally graded plasma spray coatings. In the study for the first graded joints with interlayers the bending strength values were found to be three to seven times higher than for the direcdy bonded materials. The funcdonally graded joints (FGJs), the transient produced coating, are characterised by a non-linear 3D-distribution of phases and corresponding properties, the results show a great potential in respect with the high temperature properties.

1. INTRODUCTION Manufacturing of ceramic-metal joints (AI2O3 to Ni-superalloy) for high temperature applications in power generation can be performed using solid state bonding technologies, either high temperature brazing or direct diffusion bonding (DB) with interlayers. In this study DB-technique that is suitable for joining of highly mismatching materials is developed. First joints (with Ti and Cr active layers combined with Ta and Nb interlayers) were prepared at 1050-1150°C for 1 - 3 hours under axial pressure of 5 - 20 MPa. The graded layers in these preliminary experiments were formed during the bonding cycle. Second, specimens with graded layers were manufactured by plasma spray (PS) coating procedure, where alumina content was decreased in the coating layers and the metal phase content was simultaneously increased. It was shown that use of functionally graded materials (FGM), characterised by a non-linear 3D-distribution of phases and corresponding properties, allows to decrease the residual stresses and to improve the properties of the joint [1-3,5]. The joint integrity was examined using mechanical testing with four-point bending test. The joint microstructures were studied using LOM and SEM+EDS methods. The results of strength measurements show that specimens with graded layer exhibit higher fracture strength three to seven times more than those without the layers.

* This study has been supported by the European Commission through the project BE-7249 under the contract BRE2-CT94-0928.

314 2. ALUMINA-NICKEL ALLOY JOINTS FABRICATION The alumina-superalloy joints are intended to withstand high service temperatures (700 1000°C) and thermal cycles typical in power generation processes. The fabrication of a joint should provide a relatively good strength but more important that it will remain relatively stable in the service conditions. Therefore, experimental studies for the joint design optimisation should be carried out and the data for novel procedures and for modelling the FGstructure and FE-analysis should be created. 2.1. Fabrication of joints by diffusion bonding The DB-procedure was optimised in respect with the kinetic requirements and the hightemperature mechanical properties of the Ni-superalloy. From the kinetic point of view, the bonding temperature should be over 1000°C when alumina and transition metals are directly bonded [6]. The bonding procedure was always carried out in high vacuum, better than 2-10'^ mbar (0.2 mPa). The typical thermal and axial compression cycles are presented in Fig.la. It was experimentally found that the ambient bonding temperature is llOO^C or less due to the fast creep of the superalloy beyond this. The compression for the tests was selected as 10 MPa in ceramic-metal joints and 20 MPa in ceramic-ceramic joints [6]. The initial approach in this study was to demonstrate the use of metallic interlayers and/or coatings in the bonding procedure. For the first layer (intended for the metallurgical bonding of the ceramic), Ti and Cr were found to be the most promising metals. Here both coatings (by PVD, 1-10 |Lim thick) and layers of foils (25 |Lim thick) were used. The function of the second Ta layer of 25 |Lim thickness was to suspend the diffusion of Ti between the foils. In all of the experiments the third layer of niobium of 2 mm thickness was applied for thermal residual stress relief. The multilayered joint microstructure is shown in Fig. lb [7]. The bonding process and joint structure optimisation are resulting in a four-point bending strength (sample 12x12x60 mm^) of 7.5 MPa/4.3 MPa (tested at 25''C/400°C) for Al203/Nb/Al203-joints, 22.5 MPa/- for AlzOa/Cr/Ta/Nb/Ta/Cr/AlzOa-joints and 50 MPa/27.3 MPa for Al203/Ti/Ta/Nb/TayTi/Al203-joints. In AI2O3 to Ni-superalloy (PK33) -joints the strength values decrease about to one fourth of the above listed ones: Al203/Cr/Ta/Nb/PK33 5.6 MPa and Al203/Ti/Ta/Nb/PK33 12.6 MPa at 25 ""€. Thus the combination of Ti-, Ta- and Nb-layers is considered to be most promising, reported in details in [8]. However, the formation of the thermal residual stresses in ceramic-metal joints and the diffusion at the aimed service temperature leads to the deterioration of this potential joint configuration at the ceramic interface [9]. The use of functionally graded layers instead of the pure metallic ones seems to be the optimal solution. However, it includes two manufacturing processes, namely thermal spraying and diffusion bonding, which must be optimised for the materials and conditions considered. The first benefit claimed is in a gradient in mechanical properties providing the minimum of thermal stresses. The second one provides more stable composition at high temperatures, in comparison to the metal layers, such as Ti or Cr. 2.2. Fabrication and studying of FG-joints In this study several graded joints were fabricated. All of them were produced by low pressure plasma spray method at 800-900°C over alumina substrate (A-479, Kyocera Co.+ powder Metco 105SFP). In the experiments, variations of concentration of metal component, namely Ni-20%A1 (Metco 404NS) and NiCoCrAlY (Amdry 995) alloys were applied [6].The resulting profiles of graded coatings are shown on Fig.2. The microstructure of these coatings

315 was studied by optical microscopy and SEM. The majority of the specimens were of a good quality, although some surface cracks were observed in specimens in series 58, 59 and 63. The most possible origin of these cracks was due to higher residual thermal stresses after the spraying.

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\

-400

I

\ \

rw vA, •N

_

Time [15 min/div]

a)

b)

Figure 1. The typical thermal cycle for diffusion bonding of AI2O3 ceramic to Ni-superalloy (a) and the joint microstructure with the optimised joint configuration using interlayers (b).

100

'^n '=:o

oojOy

80 h

61,62 63 65

o

>

60 h

iS 40 20

V ^ y < ^ ^ ^ -

1

2

1

1

3

4

5

Relative coating thickness Figure 2. Profiles of metal-alumina FGM coatings produced for joining experiments. Numbers indicate different experimental series: Ni-20%A1 (58,61), and NiCoCrAlY (59, 62-65).

316 The crack-free coated alumina specimens were subjected to a diffusion bonding (DB) procedure, described above. The parameters of the process were same as for non-graded specimens. After the bonding to IN-738 superalloy the joints were examined by their appearance and microstructure. For some specimens it was found that good, crack-free microstructure does not guaranteed high mechanical properties, in particular at elevated temperatures under external mechanical load. A complete series of mechanical testing of these joints shall be made after the whole array of the data will be obtained. In order to disclose the general peculiarities of FG-joints, the calculations of their basic properties were made by a micromechanical model [2,4].

3. CALCULATION OF PROPERTIES OF FG-JOINTS In these calculations of properties of FG-joints, the micromechanical model was applied as developed by Gasik e.a. [2-4]. The calculated equations of temperature dependence of source materials data were integrated in "FGM for Windows"-program files to use them for analytical properties evaluation instead of expensive FEM methods. The following properties were calculated: elastic and shear moduli, CTE, thermal conductivity, specific heat, density, thermal diffusivity, etc., versus temperature at 20 - 820°C and volume fraction of alumina in the "metal-ceramic" graded composite. Taking into account the particular geometry of the joint, the following initial conditions were established (Fig.3): (i) FG-joint has only one-dimensional gradient (in X-direction), (ii) the composition in the joint will follow the rule Yue = x^, VAI2O3 = 1 - x^, where Vi - volume fraction of component (NiAl, NiCoCrAlY or alumina), x = X/L is relative coordinate, p - anisotropy coefficient (0.7...1.3), (iii) only thermoelastic behaviour is considered and temperature distribution in FG-joint is assumed to be uniform and steady, and the joint has no external forces applied.

Figure 3. The model for FG-joint calculations.

The results are partially shown here as contour plots for NiCoCrAlY-alumina joints (Fig.4). These values are the mean ones of the respective tensor in certain direction (X or Y/Z). As far as the specimen is assumed to be free, the graded composition in X direction will be sufficient

317 for the relaxation of the whole specimen in order to satisfy thermal and mechanical equilibrium criteria.

200

400

600

800

Temperature, °C

200

400

600

Temperature, °C

200

400

600

800

Temperature, °C b

800

200

400

-r 600

800

Temperature, °C Figure 4. Calculated properties vs. temperature and anisotropy coefficient (p): elastic modulus, MPa, along Y/Z-axes (a) and its relative difference, %, to X-axis (b); thermal conductivity, W/mK, along Y/Z-axes (c) and its relative difference, %, to X-axis (d). The main bottleneck in this case will be in relaxation on the Y-Z plane (perpendicular to gradient), since there will be the most property mismatch and restricted movement of the parts of the whole specimen. Interesting, that relative anisotropy in elastic module between X and Y/Z components is not large (2-3%), but depends on temperature and anisotropy coefficient in a complicated way (Fig.4a and b). On the other hand, differences in values of thermal conductivity between X and Y/Z are almost the same for different anisotropy, but change strongly with the temperature (Fig.4c and d). The results for Y/Z-plane could be summarised in Table 1.

318 Table 1. Properties of the FG-joint at the Y/Z plane. Increasing of... Elastic modulus, GPa Anisotropy p Increases Temperature T Decreases

Thermal conductivity, W/mK Decreases Decreases

CTE, x 10^ 1/K Decreases Increases

These calculations show sources and values of the possible properties mismatch in properties of the graded joints. For instance, such large differences and anisotropy in thermal conductivity confirm that heat flow in non-steady conditions would affect the temperature distribution in the joint quite significantly. In this case additional thermal stresses could be generated by internal gradients of temperature.

4. CONCLUSIONS Two types of approaches to produce FG-joints by the DB-method have been presented, the use of metallic interlay ers and the use of graded PS-coatings. The results show that quality joints can be obtained with the optimisation of interlayers and the bonding process. However, the joints with the graded coatings/layers are evident to meet the high service temperature requirements. The calculations and experiments also reveal the origin and the magnitude of possible properties mismatches in the graded joints.

ACKNOWLEDGEMENTS The authors would like to thank the BE-7249 project consortium and especially Dr. G. Kleer, Fraunhofer-Institut IWM Freiburg, Germany, and Dr. M.Tului, GSM Rome, Italy.

REFERENCES 1. M.Koizumi, Ceram. Trans.: Functionally Gradient Materials, Ed. B.Holt e.a., ACerS., Ohio, 34(1993)3-10. 2. M.Gasik and K.Lilius, Comp. Mater. Sci., 3 (1994) 41-49. 3. N.Cherradi, A.Kawasaki, and M.Gasik, Compos. Eng. 4 (1994) No. 5, 883-894. 4. M.Gasik, Acta Polytech. Scand., Ch 226 (1995) 73 p. 5. M.Gasik, FGM News, 31 (1996) 6-9. 6. L. Heikinheimo (Ed.). Report No.VALC154 for project BE-7249, VTT Manufacturing Technology, Espoo (1995), 60 p. 7. Report No.VALC253 for project BE-7249, VTT Manufacturing Technology, Espoo (1996), 42 p. 8. IWM-Report No. V54/96 for project BE-7249. IWM Freiburg (1996), 38 p. 9. L.Heikinheimo, Thesis Dr., VTT Publications 218, Espoo, Finland, (1995), 166 p.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

319

Advances in the Fabrication of Functionally Graded Materials using Extrusion Freeform Fabrication Greg E. Hilmas, John L. Lombardi, and Robert A. Hoffinan Advanced Ceramics Research, Inc. 851 E. 47th Street, Tucson, Arizona, USA ABSTRACT The Extrusion Freeform Fabrication technology (EFF), developed by Advanced Ceramics Research, Inc. (ACR) offers tremendous potential for net shape rapid prototyping of polymers, metals, and ceramics, as well as their hybrids such as functionally graded materials (FGMs). Two unique EFF systems capable of rapid prototyping monolithic polymer, metal, and ceramic parts have been developed and entail the sequential deposition of layers of self supporting viscous suspensions or highly loaded thermoplastics using a computer controlled extrusion head. In particular, the extrusion head builds up a 3D body by sweeping out a path based on a CAD virtual image. For the fabrication of FGMs, the EFF systems are modified to contain two extruders which dispense dissimilar materials into a small mixing head. The composition of the bi-component extrudate is controlled by proportioning the raw material feedrate from the two extruders. In this manner, the composition can be continuously graded to produce a FGM component. The ACR EFF technique offers the advantages of being able to form the body into almost any shape which can then be processed through traditional powder metallurgical or ceramic firing routes. This approach is inexpensive and potentially feasible for grading between any thermodynamically compatible ceramic-metal, ceramic-ceramic, or metal-metal material combination. 1. INTRODUCTION Functionally Graded Materials (FGMs) are currently receiving considerable attention from the materials science community, particularly in Japan, where the concept originated[l]. FGMs consist of a synergistic combination of different materials and are typically composed of ceramics graded to metals. Unlike conventional coated materials and composites, FGMs have a continuous grade in composition between their respective end members. FGM materials therefore take advantage of the properties of two different materials within the same body. The graded composition eliminates many of the problems associated with the presence of discrete interfaces in conventional composites such as poor mechanical integrity, transport losses due to low interfacial adhesion, and can also eliminate problems associated with thermal expansion mismatch which is a significant problem for many conventional high temperature composites. The use of highly loaded slurries or thermoplastic formulations combined with state-of-the-art freeform fabrication technologies also enables the rapid prototyping of FGM components.

320 A host of industrial and military applications can benefit from FGMs, consequently there is substantial interest in devising an inexpensive and versatile process for their fabrication. A unique process for FGM fabrication is under development at ACR and is an extension of Solid Freeform Fabrication (SFF) techniques which have been previously developed for the rapid prototyping of monolithic and composite ceramic components [2-5]. SFF is a rapidly developing technology destined to have a large commercial effect on the manufacturing industries. It is a computer controlled process where the desired part being prototyped is first reduced to geometric sections through the use of Computer Aided Design (CAD) software and then built up sequentially, layer by layer, out of its raw material(s). The method for transferring the CAD design to the fabrication of an actual component is quite complex and dependent on the particular SFF technology being utilized. The SFF field has rapidly progressed from producing simple models to producing complex functional prototypes. Prototypes which were once solid freeformed using waxes can now be made from high strength structural materials such as thermoplastics, thermosets, metals, ceramics and discontinuous fiber reinforced composites. ACR is actively involved in developing its own SFF technology known as Extrusion Freeform Fabrication (EFF) which has been shown to be a rapid and flexible prototyping and manufacturing process [6,7]. Two in-house systems have been developed which successfully freeform CAD designed complex parts using polymer and ceramic engineering materials including AI2O3, Zr02, Si3N4 and SiC, as well as filled and unfilled PEEK and polycarbonate thermoplastics. The next technological breakthrough lies in gaining the ability to rapidly EFF fabricate FGM prototypes for use in the design of potential mass produced FGM components. When a successful FGM composition is found, direct application of the technology can be utilized to prototype functional 3D parts. The goal of this study was to develop a rapid, flexible, and precise fabrication method for producing and evaluating potential FGM compositions. Nine different ceramic-to-metal graded compositions were successfully prepared during this study, resulting in a method which appears promising as a low-cost, high pay-off approach for fabricating and screening potential FGMs. 2. EXPERIMENTAL PROCEDURE 2.1 CAD/CAM Capabilities Similar to other rapid prototyping techniques, ACR's EFF process begins with a 3D drawing (AutoCAD, .DXF, etc..) of the component to be fabricated. The file is imported into ACR proprietary software and sliced into the individual layers and fill patterns which will be utilized to build the part. The fill pattern in each successive layer ultimately becomes an extrusion path for the EFF machine to follow while extruding the chosen material in the shape of the component being fabricated. The EFF machines utilize a 3-axis gantry with a piston extruder mounted on the z-axis. Extruder motions are driven by stepper motors and are indexed with a 4-axis motion control card which drives the x-axis, y-axis, z-axis, and one proportional axis. The 3D drawing of the desired component is converted to indexer code in an AutoCAD environment using proprietary ACR subroutines. During fabrication the extrusion piston is indexed at a rate proportional to movement on the x-y planes.

321 A second extrusion piston was required in order to fabricate FGMs, subsequently requiring precise control over a fifth axis or second proportional axis, one for each material required to form the gradient. This required multiple software and hardware modifications beyond just the creation of a dual extrusion head. The modifications included calculations associated with the ratio of material simultaneously extruded from each of the two heads. Code was written to calculate the proportions and write the values to each path definition along with code for x-y-z movement. Once the code was generated using these programs, it was downloaded to the controller card. In order to run two proportional axes on the AT6400 controller, a custom operating system was required which substituted a proportional axis for the z-axis. Movement on the z-axis was then controlled by a stand alone controller which raised the extrusion head (z-axis) a predefined amount when triggered by a statement in the indexer code. 2.2 Extrusion Freeform Fabrication (EFF) The processes for manufacturing FGMs at ACR are based on the deposition of self supporting viscous suspensions ('liquid feedstocks') or highly loaded thermoplastics ('solid feedstock') from a computer controlled moving head. The extrusion head sweeps out a path while depositing either viscous liquid slurries or a ceramic or metal loaded thermoplastic strand to fabricate the desired 3D body. In order to produce FGMs, the EFF machine is configured with dual extrusion cylinders which control the flow of two materials into a small mixing head containing an in-line static mixer connected to a deposition needle. In order to produce a 3D part, a virtual image of the desired final body is drawn in CAD. The image is then sectioned into layers and extrusion paths are generated to sweep each layer. Ultimately, the desired composition at each point can be controlled in the CAD package by proportioning the rate of feedstock flow from the two extruders utilizing the indexer code. The viscous suspensions used in the 'liquid feedstock' EFF process are thermally polymerizable acrylate gel casting formulations, very heavily loaded suspensions (>50 vol.%) of ceramic or metal particulate in polymerizable monomer solutions [8-10]. The suspensions are prepared at sufficient viscosity to maintain the shape of the body during the forming process while still able to be extruded at low pressures (50 to 150 psi). The gel casting suspensions are loaded to such an extent that very little shrinkage occurs during thermal curing thus, the shape of thefi-eeformedbody is maintained. The ceramic or metal particulate loaded thermoplastics used in the 'solid feedstock' EFF process are similar to formulations utilized in conventional powder injection molding processes and contain >50 vol.% solids. The solid feedstock approach drastically increases the variety of materials that can be used to fabricate FGMs using the EFF technique, however the extrusion process requires considerably higher pressures (500 to 1000 psi). A large number of particulate raw materials can be blended with a thermoplastic and extruded in controlled manner since the rheology of the mix can be precisely regulated by the temperature and pressure utilized during extrusion. For either EFF approach, the free formed bodies are processed through traditional powder metallurgical or ceramic firing routes. The main advantages of the ACR's EFF techniques over other FGM fabrication processes are that it has the ability to control the composition of the body in both the horizontal and vertical orientations, the ability to prepare complex shapes directly, and that

322

the process is amenable to a large variety of materials systems. Any material system that can be prepared as a gel casting formulation or blended with a thermoplastic can be used to fabricate FGMs using the ACR EFF process. 2.3 Fabrication of Functionally Graded Materials by EFF Ten different functionally graded material combinations have been currently fabricated using the ACR EFF process including the following: AI2O3 to NiAl, Zr02 to NiAl, AI2O3 to 304 S.S., Zr02 to 304 S.S., AI2O3 to Inconel 625, Zr02 to Inconel 625, WC to NiAl, TiB2 to NiAl, Tie to Inconel 625, and AI2O3 to tungsten. The majority of the compositions were fabricated in this study as flat billets to demonstrate the EFF FGM techniques and for preliminary mechanical property evaluations [10]. However, the AI2O3 to tungsten FGMs were being fabricated as W-AI2O3-W rings for potential insulating columns for heavy ion fusion accelerators containing in-situ electrodes [11]. 3. RESULTS The 'liquid feedstock' (low pressure) and 'solid feedstock' (high pressure) EFF machines utilized in ACR's FGM fabrication processes are shown in Figures 1 and 2, respectively. Figure 3 shows a scanning electron micrograph (SEM) of the cross-section of a typical FGM billet fabricated on the ACR EFF machines as square billets for mechanical testing purposes. The AI2O3-304 Stainless Steel billet shown was hot pressed for 1 hour at 1250°C and a 25 MPa load. It can be seen from the SEM micrograph that the EFF process is capable of producing a uniform transition between the ceramic and metal end members.

Figure 1. 'Liquid feedstock' low pressure EFF machine

Figure 2. 'Solid feedstock' high pressure EFF machine

323

^ 304 SS

A\p,

Figure 3. SEM micrograph of a cross-section of the AI2O3-304 Stainless Steel FGM billet. The majority of the FGM compositions were linearly elastic to failure when tested in four-point bending, however the 304 S.S. and Inconel 625 containing compositions exhibited high strengths and nonlinear fracture behavior. Figures 4 shows the load-deflection curves for the AI2O3-304 S.S. FGM system. The bars were tested separately having both the ceramic and metal side placed in tension in the four-point bend test fixture. With ceramic side in compression, the ceramic actually spalled off the compressive side of the bars prior to failure during many of the flexure tests. With the ceramic side in tension, the crack would pop in on the tensile side of the bar at a low load but would then be deflected several times by the ceramic-metal graded layers. The latter tests resulted in low strengths but extremely high work-of-fracture. In the end, these bars were visibly bent and cracked but remained intact (see inset of AI2O3-304 S.S. flexure bar).

o

2 3 Crosshead Displacement, mm

Figure 4. Load-deflection curves for AI2O3-304 S.S. four-point bend test bars. The bars were tested individually with the 304 S.S. side and AI2O3 side of each bar placed in tension.

324 4. CONCLUSIONS This study demonstrated that ACR's EFF technique is a versatile method for the rapid prototyping of functionally graded materials. The myriad of ceramic-metal FGMs produced shows that the technology is a viable method for both screening and producing potential FGM systems and components. In addition, preliminary mechanical property measurements on the FGM compositions demonstrated both high strength and high toughness with some unique failure characteristics. The FGM systems developed in this program and many other systems which have yet to be fabricated have a large number of potential commercial and government applications which may be realized through this technology. 5. ACKNOWLEDGMENTS The authors wish to graciously acknowledge the Ballistic Missile Defense Organization for their support of this research under grant #DAAH04-95-C-0049. REFERENCES 1. M. Koizumi, Ceram. Eng. Sci. Proc, 13 (1992) 333. 2. P. Jacobs, "Fundamentals of Stereolithography," Society of Manufacturing Engineers, Dearborn, MI. 1992. 3. P.M. Dickens (ed.), Proc. Third European Conf on Rapid Prototyping and Manufacturing, University of Nottingham, England, 1994. 4. H.L. Marcus, J.J. Beaman, J.W. Barlow, D.L. Bourell, and R.H. Crawford (eds.), SoHd Freeform Fabrication Symposium, University of Texas, Austin, TX, 1994. 5. H.L. Marcus, J.J. Beaman, J.W. Barlow, D.L. Bourell, and R.H. Crawford (eds.). Solid Freeform Fabrication Symposium, University of Texas, Austin, TX, 1993. 6. P. Calvert, R. Crockett, J. Lombardi, J. O'Kelley, and K. Stuffle, pp. 50-55, in Solid Freeform Fabrication Symposium, H.L. Marcus, J.J. Beaman, J.W. Barlow, D.L. Bourell ,and R.H. Crawford (eds.). University of Texas, Austin, TX, 1993. 7. K. Stuffle, A. Mulligan, P. Calvert, and J. Lombardi, pp. 60-63, in Solid Freeform Fabrication Symposium, H.L. Marcus, J.J. Beaman, J.W. Barlow, D.L. Bourell, and R.H. Crawford (eds.). University of Texas, Austin, TX, 1993. 8. M. A. Janney, Method for Molding Ceramic Powders, U.S. Patent No. 4 895 194 (1990). 9. A. C. Young, O. O. Omatete, M. A. Janney, and P. A. Menchofer, J. Am. Ceram. Soc, 74 [3] (1991) 612. 10. G. E. Hilmas, J. L. Lombardi, R. A. Hoffman, and K. L. Stuffle, pp 443-450, in Solid Freeform Fabrication Symposium, D.L. Bourell, J.J. Beaman, H.L. Marcus, R.H. Crawford, and J.W. Barlow (eds.). University of Texas, Austin, TX, 1996. 11. DOE Contract # DE-FG03-95ER82105.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

325

Novel Routes to Functionally Graded Ceramics via Atmosphere-Induced Dopant Valence Gradients M. Kitayama, J. D. Powers and A. M. Glaeser Department of Materials Science and Mineral Engineering, University of California, & Center for Advanced Materials, Lawrence Berkeley National Laboratory Berkeley, CA 94720-1760, USA

Alumina compacts, when doped with Ti, can be made to develop a dopant valence gradient during or after firing through control of the sintering/annealing atmosphere. The dopant valence, Ti^+ (vacuum) versus Ti"^"^ (air), has a pronounced effect on the resulting microstructure, and on the rate of grain boundary migration. It is possible to generate microstructures in which a transition from fine equiaxed grains to elongated facetted grains occurs. Grain boundary migration characteristics have been investigated by monitoring the growth of an oriented single crystal sapphire seed into Ti-doped and undoped AI2O3. An enhancement of the grain boundary mobility, relative to behavior in undoped AI2O3, is indicated for Ti^+-doped AI2O3. Opportunities for more widespread use of the furnace atmosphere as a means of producing microstructurally graded ceramics suggest themselves.

1.

INTRODUCTION

A broad range of materials can be described by the general term functionally graded material (FGM). Hirai [1] has recently provided a review summarizing the various types of FGMs. One form of FGM involves a continuous or nearly continuous variation in microstructure and properties that is achieved by continuously or nearly continuously grading the phase contents. An example of such a structure might be one in which a transition from a pure metal to a pure ceramic is achieved. Similarly, a gradient in the volume fraction of fiber reinforcement may be developed to optimize the performance (and reduce cost). When ceramic materials are used, such gradients are generally established in the green structure. There are other classes of material in which there are chemical discontinuities and discontinuities in selected physical properties, but continuous or nearly continuous variations in a specific property. An example of such a material would be a thermal barrier coating in which the thermal expansion coefficient is graded. Such gradients could be achieved by varying the coating composition during deposition, for example from the vapor phase. An additional category of FGM is one in which the chemical composition is essentially constant, as is the phase content, but a microstructural gradient develops that induces a property gradient. Processing in a temperature gradient is one route to achieving such microstructurally graded materials. Such materials are also referred to as fine composites [1]. The range of processing techniques that can be employed to produce FGMs is also broad [1]. Vapor-phase methods {e.g.y CVD, CVI, and PVD methods), liquid-phase methods {e.g.y electrodeposition, sol-gel, plasma spraying and molten metal infiltration methods), and a variety of solid-phase methods based on powder metallurgy are available. The solid-state methods include powder stacking techniques, powder infiltration techniques, slurry techniques {e.g.y sedimentation

326 and electrophoretic deposition methods). In contrast to the vapor-state and liquid-state methods, which yield a final product, the solid-state methods generally lead to green structures with builtin gradients that must be retained during subsequent firing and densification. One interesting variation of powder metallurgy methods is that reported by Rosier and Tonnes at FGM '94 [2]. In this work, a microstructural gradient was produced by introducing a spatial variation in the Cr content of a TiAl powder. Subsequent processing was isothermal. One can anticipate that similar variations could be produced in ceramics. Another approach, one that has not been explored extensively, is to vary the valence 0^2. multivalent impurity through control of the sintering atmosphere. The use of valence gradients is likely to provide substantial opportunities for microstructural design because one can anticipate that the local valence state will affect the local solubility, grain boundary diffusivity, surface diffusivity, lattice diffusivity, and grain boundary mobility, and thereby, the density, grain size, grain size distribution, and grain shape. Several prior studies of sintering and grain growth attracted us to exploration of Ti-doped AI2O3. Bagley et al. [3] showed that the addition of Ti"^"^ led to significant increases in the sintering rate; apparently the much more soluble isovalent Ti^+ form had no interesting effect on densification. Brook [4] proposed a defect model in which Ti4+ substitutes for AP"*", and introduces Al vacancies as a charge compensating defect; Ti^+ would not produce a similar defect. Horn and Messing [5] have studied grain growth in high-density aluminas containing between 0.15 and 0.4 wt % Ti02. Normal grain growth, anisotropic grain growth, and abnormal grain growth occurred within specific ranges of temperature, composition, and time, and models linking this to the anisotropy of the grain boundary energy were proposed [6]. Work by Glaeser and coworkers has shown that the morphological stability of surfaces in alumina [7], and the Wulff shape of alumina [8] are changed by Ti doping. Thus, Ti-doped alumina provides an interesting and challenging model system for study, with the potential for using the atmosphere to produce microstructural and property gradients. The present work was undertaken to isolate the effects of small amounts of Ti on sintering behavior and the anisotropy in grain boundary motion. Low levels of Ti dopant (<2000 ppm) were incorporated into extremely clean alumina powders (99.997%) to eliminate potential effects of liquid phases [9]. Sintering and grain growth behavior were observed as a function of dopant level, processing temperature, and atmosphere. Controlled studies of grain-boundary mobility were performed to quantify the effect of doping on the "average" grain boundary mobility, but also to ultimately gain insight on the anisotropy of the mobility, and its relationship to the development of elongated grain structures. Highlights of this work are presented in this paper.

2.

EXPERIMENTAL PROCEDURES

Samples were prepared using a high-purity alumina powder (Showa Denko UA-5101). This powder is 99.997% pure, with principal impurities: 1 ppm Mg, 2 ppm Ca, 6 ppm Si, and 15 ppm Na. The average particle size is -0.4 |im. The powder was dispersed in HPLC-grade (high-performance liquid chromatography grade) ethanol (Aldrich) with 4-aminobenzoic acid (PABA), and then processed to remove hard agglomerates [10]. A portion of this powder was doped with -500 ppm Ti by controlled hydrolysis of 99.999% pure titanium isopropoxide. Undoped and Ti-doped powders were cold isostatically pressed at 240 MPa to produce compacts. These compacts were packed in additional UA-5105 powder, placed in a 99.8% pure AI2O3 crucible (Coors), and then fired for 2 h at 600°C in air to burn out residual organics (and convert Ti to Ti^"*"). All subsequent firing steps were also performed with the compact packed in powder. Sintering studies were performed both to evaluate the sintering behavior of the aluminas, and also to identify conditions that produce fully dense compacts for subsequent study. Sintering was performed in both vacuum furnaces and in air furnaces for varying times at temperatures between

327 1350° and 1550°C. Densities were measured as a function of sintering time at temperature using the Archimedes method. After sintering, samples were polished to a 1 |im finish, thermally etched, and microstructures were subsequently examined in the SEM. Selected near-theoretical density samples were hot isostatically pressed to full density (1400°C, 180 MPa, 1 h). Compact purity at this stage of processing was assessed using wavelength-dispersive x-ray fluorescence; a detection limit of <1 ppm for the impurities of interest is claimed. No loss of purity relative to the starting powder was detected for the undoped material. Chemical analysis of the Ti-doped material indicated a Ti level of ~700 ppm, with a factor of two uncertainty. This suggests no substantial Ti loss occurs during firing. The fully dense samples were used for two experiments. In the first, the growth behavior of matrix grains was examined. This yields information on the evolution of the average grain size, and thereby provides a measure of the spatially averaged grain boundary mobility (Mb, the grain boundary velocity per unit driving force). The evolution of grain shape provides a qualitative indication of the degree of orientational anisotropy of the grain boundary mobility. In the second, dense polycrystals were diffusion bonded to oriented single-crystal sapphire seeds, and the growth rate of the seed into the polycrystalline material was investigated. This experiment yields more specific orientation-dependent information on the grain boundary mobility, and allows an independent assessment of the effect of Ti-doping on the grain boundary mobility.

3.

RESULTS AND DISCUSSION

Figure 1 compares the time dependence of the density in undoped and Ti-doped aluminas fired in air at 1400°C. Under these conditions, Ti should be in the 4+ state. Pairs of data points for both the undoped and 500 ppm Ti-doped material represent measurements taken on different samples, not measurement error. It is evident that both the undoped and doped alumina sinter to

Figure 1.

Undoped —••—- 500 ppm Ti-doped • 2000 ppm Ti-doped 652 3 Time (hours)

Plot of density versus sintering time for undoped, 500 ppm and 2000 ppm Ti-doped alumina. Firing was performed in air at 1400°C. The air atmosphere is expected to yield Ti"^"^. The enhancement in densification, relative to the behavior of undoped AI2O3 is quite modest at the 500 ppm doping level. At 2000 ppm Ti^+, a much more substantial effect is evident.

328 high density in relatively short periods of time, and that doping with 500 ppm Ti'^^ does not substantially enhance the rate of densification. The data for undoped and 500 ppm doped material overlap when the density is <90% of theoretical. Data for 2000 ppm doped material is included; here the effect of Ti-doping is quite pronounced. At 1400°C, the solubiUty limit of Ti4+ in AI2O3 is <2000 ppm [11], and thus, precipitate pinning may have an effect on sintering. The interesting microstructural differences arise when compacts, processed in air at 600°C to yield Ti4+, are then sintered at high temperature in vacuum where the Ti4+ is reduced to Ti3+. Samples were cut and polished after firing. Figure 2 compares microstructures that are produced after 2 h vacuum sintering of undoped AI2O3 and 500 ppm Ti-doped AI2O3. In the undoped material (fired at 1500°C), there is no difference in microstructure between (a) the edge and (b) the center of the sample. The grain sizes are similar, and the grains are equiaxed. In Ti-doped AI2O3, the grains at the sample edge were equiaxed. Fig. 2c; the edge of the sample had become pink, a characteristic feature of Ti3+ doped Ai203. In the center of the sample, the onset of grain elongation, and facetting is evident; here the material was "white". X-ray fluorescence measurements have confirmed the presence of Ti^+ on the outer rim of the material, and the absence of Ti^+ in the inner regions of the sample. It is clear that Ti^+ and Ti^^ have distinct effects on microstructural evolution. The observations show that microstructural gradients can be induced due to a dopant valence gradient.

Figure 2.

Comparison of microstructures developed in a, b) undoped and c, d) 500 ppm Ti-doped alumina. At the edges of both the a) undoped and c) Ti-doped sample, equiaxed microstructures develop. At the centers of the b) undoped and d) Tidoped alumina, the microstructures are different, with facetted and elongated grains developing in the Ti-doped material.

329 Many factors will influence the nature of the structure that is produced. The chemical form of the dopant will affect the initial valence state of the dopant, and the doping level will influence the potential for second phase formation, and the potential for full charge compensation of the aliovalent impurities in the "undoped" material. The dopant valence can be manipulated during a low temperature prefiring treatment, during which little densification occurs. The dopant valence gradient will be impacted by the temperature and atmosphere conditions/cycle used during sintering. A broad range of microstructural types may be accessible if such dopant gradients are combined with intentional codoping with impurities of fixed valence. One aspect of particular interest was the effect of the impurity addition on the grain boundary mobility, M^. The development of facetted and elongated grains suggests substantial anisotropy of the grain boundary mobility. To explore this issue, experiments mimicking those conducted earlier by Rodel and Glaeser on undoped and MgO-doped aluminas [12] were performed. In the present case, sapphire substrates with ^-plane and c-plane surfaces were bonded to dense polycrystals of undoped and Ti-doped AI2O3. As in the prior work, the displacement of the sapphire/polycrystal boundary was monitored as a function of time, temperature, and atmosphere. The most complete results to date are for /^-plane and c-plane sapphire migrating into undoped and 500 ppm Ti-doped AI2O3 in vacuum at 1600°C. Figure 3 shows results for growth of ^-plane sapphire seeds into two 500 ppm Ti-doped samples and the undoped material of this study. Results from prior work [12] utilizing a different starting powder are also shown for comparison. In general, as the purity of the system is increased, M^ is expected to increase. The slightly larger displacements observed in the undoped case when using a somewhat higher purity undoped powder than in previous work is thus consistent with expectations. This effect was seen for both <^-plane and c-plane seeds. In general, dopant additions decrease the grain boundary mobility. Horn and Messing reported small increases in Mb as a result of Ti^+ doping. The pronounced effect of adding Ti^+ observed in the present case is clearly unexpected, since it suggests that doping leads to an increase Figure 3.

1000

160(fC; vacuum

gSOO

§600

o400 u

^200

" • X • • Prior work (undoped) — ^ — Undoped Ti-doped (sample 1) Ti-doped (sample 2)

10 15 Time (hours)

20

25

Plot of grain boundary displacement versus anneal time for growth of aplane sapphire seeds into 500 ppm Tidoped and undoped aluminas at 1600°C. In comparison to results obtained previously with a lower purity powder, mobilities in undoped material are increased. However, the most striking feature is the increased migration rate for the Ti-doped material.

330 in the extent of grain boundary migration. When the data are converted to mobiUties, the difference persists. The values of the boundary mobiUties for ^-plane and c-plane seeds growing into the Ti-doped material are -5-1 Ox those of identical orientation seeds growing into the undoped polycrystal. It is possible that a thin layer of glassy phase has formed, but this appears inconsistent with the equiaxed grain morphology observed in the Ti^+ doped regions of our samples {see Figure 2c), and seems unlikely since the starting materials are of high purity, and the processing retains this high level of purity. An alternative explanation is that the Ti^+:Ti'^+ ratio in these materials is such that it charge compensates for aliovalent impurities such as Na, Mg, and Ca. Experiments in which the Ti level is being varied, and in which the atmosphere is being changed to air are being conducted to answer these questions. Although not directly related to processing of FGMs, the ability to induce high grain boundary mobilities may find application in areas where the growth of single crystals by solid-state processes is desired.

SUMMARY AND C O N C L U S I O N S

AJ2O3, when uniformly doped with Ti, can be made to develop a dopant gradient during or after firing through control of the ambient atmophere. The dopant valence gradient induced in this fashion has a pronounced effect on the resulting microstructure and rate of grain boundary migration. Elongated, facetted grains and equiaxed grains can be generated in the same sample, depending upon the local Ti valence state. An unanticipated increase in the grain boundary mobility is suggested for Ti^+-doped alumina; work is in progress to understand the underlying cause of this enhancement. More generally, new opportunities for using atmosphere-induced valence gradients to broaden the range of graded structures that are available suggest themselves.

ACKNOWLEDGEMENTS This work is supported by the National Science Foundation under Grant No. DMR-9222644. We also acknowledge an NSF Equipment Grant No. DMR-9119460 which allowed the acquisition of hot pressing equipment used in this work. REFERENCES

1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12.

T. Hirai, Chapter 20 in Volume 17B, Materials Science and Technology, A Comprehensive Treatment, R. J. Brook, Ed., VCH Verlagsgesellschaft mbH, Weinheim, Germany (1996). J. Rosier and C. Tonnes, pp. 41-46 in FGM *94 Proceedings, N. Cherradi and Ilschner, Eds., Presses polytechniques et universitaires romandes, Lausanne, Switzerland (1995). R. D. Bagley, I. B. Cutler, and D. L. Johnson, / Am. Ceram. Soc, 53 [3] 136-41 (1970). R. J. Brook, /. Am. Ceram. Soc, 55 [2] 114-5 (1972). D. S. Horn and G. L. Messing, Mat. Set. andEng. A, A195 169-78 (1995). W. Yang, L-Q. Chen and G. L. Messing, Mat. Sci. andEng. A, A195 179-87 (1995). J. D. Powers and A. M. Glaeser, / Am. Ceram. Soc., 7G [9] 2225-34 (1993). M. Kitayama and A. M. Glaeser, pp. 285-92 in SINTERING TECHNOLOGY, R. M. German, G. Messing and R. G. Cornwall, Eds., Chapman-Hall, London England (1996). P. E. D. Morgan and M. S. Koutsoutis, / Am. Ceram. Soc., 68 [6] C-156-C-158 (1985). M. Kitayama and J. A. Pask, /. Am. Ceram. Soc, 79 [8] 2003-11 (1996). R. W. Grimes, / Am. Ceram. Soc, 77 [2] 378-84 (1994). J. Rodel and A. M. Glaeser, /. Am. Ceram. Soc., 73 [11] 3293-301 (1990).

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

331

The Growth of Functionally Graded Crystals by Verneuil's Technique M. Ueltzen«, J.-F. Fournie^, Ch. Seega^, H. Altenburg^ « SIMa Steinfurt, FH Miinster, Stegerwaldstr. 39, D-48565 Steinfurt, Germany ^ FH Miinster, Dept. of Chemical Engineering, Germany

The potential of the Verneuil process in the fabrication of functionally graded materials is discussed in detail. The powder feeding system of a conventional Verneuil apparatus was changed in order to enable variations of stoichiometry in the growing crystal. The possibilities of this technique are shown with the system ruby-sapphire. Several graded crystals based on alumina, hke rubyleukosapphire and blue sapphire-leukosapphire-ruby, were grown. A model shows the way, in which the results can be transferred to other materials. An outlook on the possibihties of the VerneuU technique relating to other materials including non-oxide crystals is given.

1. INTRODUCTION

In the field of crystal growth, a lot of work was done in order to avoid gradients of stoichiometry. Such gradients in the composition are caused by segregation phenomena during the crystaUisation process. In the case of the Verneuil technique, it is possible to influence the composition of the hquid phase in situ. So, gradients in composition can be avoided or adjusted during the growth process. The main part of a conventional Verneuil apparatus is a vertical inverted oxyhydrogen burner. The burner consists of several concentric tubes [1, 2]. The inner tube suppUes oxygen and crystal powder to the flame, which burns into a ceramic muffle. A seed crystal is positioned in such a manner that only a thin film of hquid phase is on the top surface of the crystal. If the powder supply, i.e. the feeding of molten droplets into the molten layer, and the crystal growth rate are well balanced, the crystal will grow under constant conditions. The principal design and operation of a Verneuil flame fusion apparatus was described earher [3].

332 Because of the small volume of liquid phase, the Verneuil process is predestined to the growth of graded materials. The development of the fluxVerneuil technique [4] enables the variation of the material support to the growing crystal. In this way, a specific influence on the stoichiometry of the liquid phase is possible.

2. A P P A R A T U S

An Adamski type 5ring burner was used [2]. The powder supply system is shown in fig. 1. Each powder supply channel consists of a container with a rotating brush on a sieve. The brush is driven by a gearbox motor from the upper part of the construction. This system allows the in situ variations of three different types of powder. Because the inner tube of the burner suppHes gas as well as powder to the flame, no technical equipment against a flame kickback into the material support system is possible. It has to be ensured that the whole system has an overpressure of some 10 Pa (some mm H2O). Therefore, the seal of the dome is very important for a secure operation of the apparatus. A water ring seal (see fig. 1) was used. The cables for electric power supply and control of the motors as well as the controlled oxygen stream are passed through the water ring seal into the powder supply system.

Fig. 1: Powder supply system 1 - gearbox motor, 2- dome, 3 - rotating brush, 4 - funnel, 5 - water ring seal, 6 - oxygen and powder supply to the burner

333 3. CRYSTAL GROWTH EXPERIMENTS

Alumina is the standard material for the Verneuil process, produced in the range of 200 t per year [5]. Therefore, the possibilities of the process in the field of functionally graded materials are shown at the system ruby-sapphire. Alumina crystals with varying amounts of dopants are investigated. The gradients are limited by the thickness of the molten layer on top of the growing crystal. Free flowing powder was used. It was prepared by the alumn process at CKB Bitterfeld for the undoped material or at ESCETE Enschede for the chromium doped alumina powder, especially. Typical gas flows were 30 1/min hydrogen and 16.5 1/min oxygen. An alumina crystal of a) jCrKa(cts) 2-2 mm2 upper area was used as a seed. The growth velocity was around 1 cm/h. Tj^ical dimensions of the grown crystals are 10 mm thickness and several cm length. Crystals with sharp transi100 Mm tions of the dopant as well as b) crystals with extended transition regions in the cm range were grown. 4. CHARACTERIZATION

The crystals were embedded in epoxy, cut longitudinal and polished using diamond spray down to a grain size of 3 ^m. Fig. 2b shows a SEM image of a crystal with a gradient in the chromium doping. In this case a crystal with a sharp transition was investigated. In SEM micrographs as well as in line scans of the chromium content, the graded region was examined. Fig. 2a shows a line scan of the chromium content across the ruby-sapphire

Fig. 2: SEM image (b) and line scan (a) of chromium content at the rubysapphire boundary (Length: 175 |mi)

334 border in the central part of the crystal. The noise is caused by the low chromium percentage. The transition was sharp (8 p.m) in the central region of the crystal (fig. 2a) and extensive (80 ^m) in the outer region. Up to now, there are only experiments with small crystals of up to 10 mm diameter. It is well known, that the phase boundary is not flat in this case. So, these high differences between outer and inner zones of the crystal should be diminished by the growth of larger crystals.

5. MODELING

In order to ensure the transfer of knowledge from the alumina system to other materials, a mathematical model was developed. This model is based on the balance of the dopant, e.g. chromium, in the molten layer: An increase of chromium concentration in the suppUed powder leads to an enrichment of the melt with chromium and by this way to an increasing amount of chromium in the crystallizing material. This case is written in the following equation: dmc. = A'S'Pj

dCi -\- A-dz'p^ • c^

(1)

where dm is the amount of powder (suppUed during a short time dt), a is the initial concentration of the dopant, e.g. chromium, in the suppUed material, A is the base surface and S is the height of the molten layer, c and p are the concentration and the density of the soUd phase (s) or the liquid phase (0, especially, z is the coordinate from the phase boundary into the liquid phase (fixed to the laboratory system in this case of Verneuil geometry) while z*=z-^vt is the coordinate fixed to the growing crystal, v is the growth velocity. This becomes with dm = ps'A'dz\ introducing the segregation coefficient k = Cs/ciy which ratio is given by the phase diagram or even by the effective segregation coefficient for more complex transport phenomena, and at ideal and immediately mixing in the molten volume — -^^-^—Z'^sdz' Pi'O

=^

(2)

Pi'O

In the case of an abrupt change in the initial concentration in the suppUed powder a at the time to, this equation results in an exponential change of the concentration in the grown crystal:

Csi^') = C,

1 - exp - ^ • - • ( z ' - z ; )

(3)

335 Assuming ps - pi and k = 1, the exponential decreasing varies with the thickness S of the molten layer on top of the growing crystal. This is a reasonable result. It is discussed in detail because of the wide variation range of the thickness of the molten layer for different materials. In the case of alumina a thickness of ca. 20 ^m is discussed [6, 7, 8], while at strontium titanate the molten layer reaches a thickness in the mm range [9]. So, the maximum gradient of the concentration differs strongly depending on the material. This fact should be taken into account when developing new methods for new materials. In reality, the process may be much more difficult as described in the model given above. Further problems are evaporation losses of the dopant material at the way through the flame or even from the molten layer, hampered mixing and transport in the small molten volume, inhomogenieties of the liquid phase as well as a non-flat phase boundary. The results given in chapter 4 are difficult to discuss in the framework of this model. A thickness of the molten layer in the ^m range seems not reasonable. A large shrinking during the melting process (ps » pi) or a big distribution coefficient (k » 1) of chromium are not known in the case of alumina. The strong dependence of the melt thickness from the radius shows that there are other effects taken into account. It seems that at larger diameters of the crystal and at a more flat phase boundary as well as for substances with a thicker melting layer, the model fits better. This hopeful assumption should be confirmed by growing thicker crystals of alumina and by growing other substances like spinel or titanates.

6. OUTLOOK

With the development of the Low-temperature Flame Fusion technique (LTFF) [10, 11], the Verneuil process works in the temperature range between 800...2500 °C. A brief review of the possibiUties of the Verneuil process was given some years ago [3]. A lot of different oxide materials, e.g. metal oxides, rare earth oxides, and multi component oxides Uke aluminates, ferrites, silicates, and garnets were grown by this technique. Furthermore, the introduction of other heat sources, like rf plasma, radiation furnace or sun light, enables the growth of non-oxide materials by this technique. So, Si, Nb, W, Mo, ZnSe, TiB2 as well as carbides Uke TiC, HfC, UC were grown by modffied Verneuil techniques (review see [3]). A lot of different materials can be grown by this technique - even with gradients of dopants. Further developments in this special field of crystal growth are hampered by a strong prejudice against the Verneuil process concerning the quality of the crystals. In the just pubUshed Handbook of Crystal Growth [12], the Verneuil process is only mentioned marginaUy and with regard to the historical aspects only. There are two reasons for this situation. Firstly, other techniques of crystal growth consumed a high amount of man power and money in the era of the

336

development of the semiconductor industry. Secondly, the techniques of temperature control in electric furnaces developed strongly during this period. Today, the development of gas flow control techniques during the last years enables good conditions for the further automation of the Verneuil process as well. So, the way of further technical developments to grow high quality crystals by the Verneuil technique is clear. The new possibilities of functionally graded materials should promote these developments. ACKNOWLEDGEMENTS

This paper is based on a project sponsored by the Deutsche Forschungsgemeinschaft DFG. Prof. Lodding and Mr. Uphoff, FH Miinster, are acknowledged for making the SEM investigations, Mr. Droste from ESCETE Single Crystal Technology B.V., Enschede/NL for valuable discussions. The students Th. Briiggenkamp, W. Seifert, M. Franke and C. Straniero are thanked for their assistance in the experiments. REFERENCES 1. A.V. Verneuil, Annales de chemie et de physique Huitieme Serie 3 (1904) 20. 2. J.A. Adamski, J. Appl. Phys. 36 (1965) 1784. 3. M. Ueltzen, J. Crystal Growth 132 (1993) 315. 4. M. Ueltzen, VDI-Fortschrittberichte R5 (1995) Nr. 391 (in German). 5. H.J. Scheelin[12],Vol. la. 6. J.A. Adamski, Rev. Sci. Instrum. 40 (1969) 1634. 7. E.M. Akulenok, Kristallografiya 25 (1980) 653. 8. J.G. Grabmaier, J. Crystal Growth 5 (1969) 105. 9. J.G. Bednorz, H.J. Scheel, J. Crystal Growth 41 (1977) 5. 10. M. Ueltzen, T. Briiggenkamp, M. Franke, H. Altenburg, Rev. Sci. Instrum. 64 (1993) 1089. 11. M. Ueltzen, C. Grause, H. Altenburg, J. Lons, Cryst. Res. Technol. 28 (1993) K69. 12. D.T.J. Hurle (Ed.), Handbook of Crystal Growth, North-Holland, Amsterdam, 1993.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

337

Excimer Laser Processing of Functionally Graded Materials Yoshihisa Uchida^, Jun Yamada^, Y.P.Kathuria^, Niichi Hayashi^, Shigeo Watanabe^, Shuntaro Higa^, Hideo Furuhashi^ and Yoshiyuki Uchida^ ^Department of Electrical and Electronic Engineering, Aichi Institute of Technology, 1247 Yachigusa, Yakusa-cho, Toyota 470-03, Japan E-mail:[email protected] Fax:+81-565-48-0070 ^ Laser X CoLtd., Shinbayashi-cho, Chiryu, Aichi 472, Japan This paper reports on some characteristic properties of Functionally Graded Materials (FGMs) processed with pulsed KrF excimer and Nd-YAG lasers. In these experiments, a new production method — a progressive lamination with continuous gradation by filtration using a procedure resulting in a mechanical separation of sohds and liquids — is used for producing thick blocks of FGMs. Ceramic-metal FGMs were produced by mixing Korean kaolin (Si02-Al203) and ferric oxide (Fe203) powders at 900 °C of gas calcination temperature. Surface treating and cutting of some layers of FGMs were performed.

1. INTRODUCTION Functionally Graded Materials (FGMs) have become the object of pubHc attention for various apphcation fields [1-3]. FGMs have the properties of the two raw materials which are mixed together and the component distribution is graded continuously. For example, one of the FGMs produced using ceramic and metal has the property of metallic tenacity and yet it is heat proof and anti-corrosive Hke ceramic. It can also be used as a material to withstand thermal stress [1,4-5]. The authors have proposed a new production method using a method resulting in a mechanical separation of sohds and hquids and they have succeeded in producing thick blocks of FGMs using this method. Functional gradation of FGMs was accompHshed by progressive lamination of the boundary layers through continuous gradation using filtration. Until now, microprocessing for micromachine apphcations, such as cutting, driUing and surface treating of polymers, ceramics and metals has been accomphshed using

* This research was supported in part by a Grant-in-Aid for Scientific Research from the Ministry of Education, Science, Sport and Culture, Japan.

338 pulsed excimer [6] and Nd-YAG [7] lasers. This paper reports on some characteristic properties of the FGMs processed with KrF excimer and Nd-YAG lasers.

2. EXPERIMENTAL PROCEDURE 2.1. Functionally Graded Materials FGMs in this experiment were produced by a newly developed method as shown in Figure 1 by using mechanical filtration, consolidation, calcination and sintering fi*om mixing composite powders of uniform granular diameter in distilled water. Mixing composite powders | Lamination Mechanical filtration Consolidation Calcination

I

Sintering Figure 1. Production process of FGMs. Ceramic-metal FGMs were produced by mixing Korean kaolin (Si02-Al203) and ferric oxide (Fe203) powders at 900 °C of gas calcination temperature. The mixing powders were progressively laminated in the order of the heavy weight of ferric oxide and the light weight of Korean kaolin at a constant ratio.

Kaohn

Iron

Figure 2. Scanning Electron Microscope photograph of FGMs with a kaolin rich second layer.

339 Each layer, therefore, is the same weight. The FGMs thus produced have ten layers with a composite thickness of 11mm. Figure 2 shows an example of a scanning electron microscope (SEM) photograph of FGMs. From the SEM photograph, it was observed that grains of iron and kaolin exist together in each layer. However grains separated from each other occur locally. Also, some cavities caused by steam were observed. From the SEM photographs of the side wall of FGMs, functionally graded characteristics of the grains were observed. 2.2. Exdmer laser processing A Mitsubishi Excimer Laser Work System MEX-24 was used to study the characteristic properties of the FGMs. The experimental conditions used are as foUows: a laser medium of KrF, a wavelength of 248nm, multi reflection optics, a synchronous scanning system, a transcribing magnification of 1/4.5 and an energy density of 6J/cm^. Six laser beams in multiple scanning mode were applied on the functionally graded side wall of the FGMs.

1st layer

2nd layer

3rd layer

1st layer

2nd layer

Figure 3. SEM photographs of FGMs with kaolin rich layers.

340 2.3. Nd-YAG laser processing In another experiment, an Nd-YAG laser with a wavelength of 1.06 // m from a Lumonics model JK704LD/Laserdyne 780 system was used for cutting FGMs. The beam was focused by a lens with a focal distance of 200mm. 3. RESULTS AND DISCUSSION 3.1. Excimer laser processing Figure 3 and 4 show the SEM photographs of the FGMs processed by the excimer laser. From the SEM photograph of FGMs exposed by laser, different characteristic properties of each layer were observed. They are described as foUow: Figure 3 shows the SEM photograph of the surface of kaohn rich layers of FGMs; the frrst layer, second layer and third layer and Figure 3(b) shows an enlarged boundary of first and second layers. An irregularity of the surface of the kaohn rich layers was observed due to non-uniformity of powder particles and imperfections of mixture.

10th layer 9th layer

10th layer

9th layer

Figure 4. SEM photographs of FGMs with iron rich layers.

341 Figure 4 shows the SEM photograph of a surface of the iron rich layers of FGMs processed; the tenth layer, ninth layer and eighth layer. Figure 4(b) shows an enlarged boundary of tenth and ninth layers. Good quality of the surface of iron rich layers was observed due to the high thermal conductivity of molten iron relative to the laser conditions when processing. From the SEM photograph of the boundary of second and third iron rich layers, it was observed that grains of iron had grown covering the grains of kaoUn. 3.2. Nd-YAG laser processing Figure 5(a) shows a photograph of a section of the FGMs processed with a Nd-YAG laser. The kaolin side shows on the top and iron side shows on the bottom. The laser beam was radiated from the kaolin side of the FGMs. An irregularity of the kaolin layer was observed. Figure 5(b) shows the topographic view of a section of FGMs taken by an Opton Moire 3D Camera. The surface heights are shown with sixth ranks in each area of 200 // m x 200 fi m. Each rank has a pitch of 30 ji m. The red color shows the lower level and the green color is the higher level. An undulation of 180 ii m was found for the graded layers.

(b) Figure 5. Photograph (a) and topographic view (b) of the section of Nd-YAG laser processed FGMs.

342

It was considered that an oscillating mode change and non-uniformity of the laser beam was the cause of undulation. Irregularity of the laser beam direction was observed due to a different thermal conduction of each material. The heat conducted a period of laser pulse uniformly in the iron layer because the thermal conduction coefficient of iron is higher than that of kaoUn. So the plane surface of the iron side is smoother than that of the kaolin side. Because of the low thermal conduction coefficient of kaohn and the multiphotons, non-thermal photochemical process, we can consider that an irregularity of the surface of the kaolin side was due to the change in intensity of the laser beam. Due to the high thermal conduction coefficient of iron and the thermal process, the surface of the iron side had smoother character. The porosity, the mixture conditions and the microstructure of the iron and kaohn powders must also be considered when evaluating the flatness of the processed surface. 4. CONCLUSIONS It was experimentally demonstrated that some thick Korean kaohn-ferric oxide Functionally Graded Materials could be produced by a newly developed filtration method, and surface treating of the FGMs could be processed by a KrF excimer laser and cutting of the FGMs could be processed by a Nd-YAG laser. It was concluded that the characteristic properties of the FGMs surface when processed with the lasers depends directly on the functionally graded characteristics. REFERENCES 1. S.Watanabe, N.Hayashi, Y.Uchida et al., Proc. 4th International Conference on Properties and Applications of Dielectric Materials, 1(1994)185. 2. S.Watanabe, N.Hayashi, Y.Uchida et al., Proc. International Symposium on Electrical Insulating Materials, (1995)129. 3. S.Watanabe, N.Hayashi, Y.Uchida et al., Proc. Seventh International Conference on Dielectric Materials, Measurements and Apphcations, (1996) in press. 4. N.Hayashi, S.Watanabe, Y.Uchida et al., Proc. 2nd China-Japan Joint International Conference on Filtration and Separation, (1994)353. 5. N.Hayashi, S.Watanabe, Y.Uchida et al., Proc. Thailand-Jap an Joint Conference on SoUd-Fluid Separation Technology, (1995)10. 6. H.Furuhashi, N.Hayashi, S.Watanabe, Y.Uchida, Y.P.Kathuria et al.. International Ceramic Monographs, 1(1994)1234. 7. Y.P.Kathuria, Y.Uchida et al., Proc. Japan Industry Apphcations Society Conference, (1994)173.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

343

Development of stainless steel / PSZ functionally graded materials by means of an expression operation K.Taka^^ Y.Murakami ^^ TJshikura^^ N.Hayashi ^^^, S.Watanabe^^ Y.Uchida^^ S.Higa ^M.Imura ^^ D.Dykes ^^' (a)Aichi Institute of Technology, Yachigusa Yakusa Toyota 470-03 JAPAN (b)Yokkaichi University, Kayo Yokkaichi 512 JAPAN Recently, with advances in technology, materials for the construction of machinery are now being utilized in various kinds of severe environment. It is in order to cope with difficulties of this kind that recourse to "functionally graded materials" (FGMs) has been proposed in Japan. However, although various techniques are currently available for the manufacture of FGMs, it is very difficult to produce FGMs of the greater thickness widely required in manufacturing industry. This paper attempts to present a method of production for FGMs by means of an expression operation resulting in a mechanical separation of solids and liquids, using the two processes vacuum filtration and expression consolidation. It examines how this expression technology can be used to produce FGMs of greater thickness more simply and cheaply than has hitherto been possible. In this way, compressed FGM cakes materials could be manufactured using ZrO 2 as the ceramic material and stainless steel powder as the metallic material. FGM cakes produced in this way were air dried and then baked in an electrically powered reducing furnace. FGM cakes produced in the above described way observed the case that it use an addition agent and an effect of the graded layer steps and an effect by the thickness of it.

1. INTRODUCTION Recently, with advances in technology, research and development of new materials is being carried on continuously in many countries. In a sector such as aircraft engines, the demands made on construction materials are extremely exacting, so the durability of materials is becoming an increasingly important consideration. Two ways of dealing with this requirement are by smiace or coating treatments, and by the layered combination of materials possessing different properties. However, the usual methods of treating materials may prove inadequate in extreme cases. It is in order to overcome this difficulty that recourse to functionally graded materials (FGMs), has been advocated in Japan^^

344

Various techniques are currently available for the manufacture of FGMs, in particular the chemical and physical vapour deposition methods^^\ However, using these methods, it is very difficult to produce FGMs of the greater thickness widely required in manufecturing industry. This paper presents a method of production for FGMs by means of a mechanical separation of solid and liquid, using the two processes of vacuum filtration and mechanical expression^^^ Compressed FGM cakes have actually been produced in this way using electrically powered reducing furnaces.

2. EXPERIMENTAL PROCEDURES Figure 1 show the compression permeability cell used in the experiment. The vacuum filtration apparatus comprises a principal base, a cylinder of 60mm diameter and a perforated plate. Figure 2 shows the vacuum filtration equipment. To produce the filtered cakes of FGM, a mixed slurry consisting of two components was sucked continuously by a vacuum pump through a filter paper extended over the perforated plate. Before the start of the filtration, a number of mixed slurries were prepared in glass vessels, each containing the same two component materials progressively graded in ratio. The filtered cakes of graded material were made by means of a staged addition method, in which the slurries were introduced in order into the cylinder of the apparatus. r=^^

PERFORATED DISC FILTER PAPER

^

^

UPPER PART OF CYLINDER

BOTTOM PART OF CYLINDER

Figure 1. Schematic diagram of a vacuum filtration apparatus and a compression permeability cell In the experiment reported here, slurry was added in six stages. After a graded filter cake had been formed, a piston was inserted into the cylinder of the apparatus, producing the compression permeability cell shown . The cake was compressed by applying mechanical pressure to the piston. The compressed cakes were air dried and then baked in an electrically powered reducing furnace.

345 CYLINDER WATER TRAP /) VACUUM TANK

VACUUM PUMP

Figure 2. Systematic diagram of vacuum filtration For the experiment, materials were deliberately chosen that are common and readily obtainable. The ceramic material used was partially stabilized zirconia (PSZ)(50 IJL m)and the metallic material was stainless steel(2 ^ 10 /i m). PSZ has the chemical formula Zr02-Y203. An addition agent used for binding was Baindoseramu WA-320(Maruto Ltd) and the dispersing agent was Seruna D-305(Maruto Ltd).

3. EXPERIMENTAL RESULTS AND DISCUSSION 3.1. Production of the expressed cakes The most significant advantage of the proposed production method is that it should result in cakes which are homogeneously graded in every layer. To ensure that this is so, the effect of adding a dispersing and binding agent to twocomponent mixed slurry were measured. The presented results concern the cermet material SUS:PSZ produced from a 20%:80% mixed slurry. This is one of the most difficult graded materials to manufacture. 3.1.1 Production without the use of an addition agent Photograph 1 shows a cake of cermet material of the type SUS:PSZ=20%:80% that was manufactiured without the use of an addition agent. The photograph confirms that the SUS304 and PSZ were segmented at one layer in the material. It was not possible to produce a uniformly graded cake, as the settling rate of the SUS304 particles was greater than that of the PSZ ones. 3.1.2 Production using an added dispersing agent Photographs 2 and 3 show compression cakes to which a dispersing agent has been added in proportions of 0.5% or 1 % to the weight of the powder. No effects are evident in the photographs from the addition of the dispersing agent. The agent has a dispersing effect on the PSZ particles, but not on the SUS304 particles. This is no doubt due to the difference in specific gravity between the two.

346 3.1.3 Production using an added binding agent Photographs 4 and 5 show compression cakes to which a binding agent has been added in proportions of 0.5% and 1% to the weight of the powder. The photograph 4 reveals that the materials are not so sharply dissociated as in 3~ 1-1 and 3 - 1 - 2 above. However, a completely homogeneous cake has not been achieved. Photograph 5, on the other hand, shows that an obvious tendency towards homogeneity in the cermet material in the case where a binding agent has been added in the propotion of 1% weight. It is conceivable that the binding agent might cause aggregations of SUS304 and PSZ particles to form in the slurries, elimmating the difference in settling rates. 3.2 Production of FGMs Where different materials use gradedly laminated together, the coefficient of contraction. Therefore, the FGM may bend after baking. It is now examined how far this bending can be relaxed, by varying the layer steps and the thickness of the FGM. 3.2.1 The effect of varying the layer steps Tow types of FGM were examined, consisting of respectively 6 grade stages (20% rate of change in slurry concentration ), and 11 grade stages (10% rate of change). The behaviour of each type was observed. Photographs 6 and 7 show the FGM laminated in 6 and 11 grade stages respectively. It is clear that the 11-stage material displays a greater bend than the 6-stages material. This is because the thickness of each layer is smaller. 3.2.2 The effect of thickness In this test, the number of stages was fixed at 6 , while the thickness of the FGM was varied, and the effect on bending was observed. Photographs 6 and 8 show the after- baking bend in two FGMs, each produced in 6 grade stages and having a thickness of 6 mm and 12 mm respectively. It can be seen that bending is better controlled in the 12 mm material than in the 6mm one. This is no doubt because the difference in the coefficients of contraction is less influential when the layer thickness is small.

4. CONCLUSION The addition of a binding agent during the manufacture of FGMs is advantageous, as a more uniform mixture is slurry obtained. For an FGM of given thickness, a composition of 11 grade stages allows better control of bending deformation than a composition of 6 stages. With a layer thickness of 2mm, bending can be better restrained than with a layer thickness of 1mm.

347

Photograph.l PSZ:SUS304 80%:20% / Expression Operation lOMPa / Addition Agent Owt%

Photograph.2 PSZ:SUS304 80%:20% / Expression Operation lOMPa / Dispersing Agent 0.5wt%

Photograph.3 PSZ:SUS304 80%:20% / Expression Operation lOMPa / Dispersing Agent lwt%

Photograph.4 PSZ:SUS304 80%:20% / Expression Operation lOMPa / Binding Agent 0.5wt%

Photograph.5 PSZ:SUS304 80%:20% / Expression Operation lOMPa / Binding Agent lwt%

348

Photograph.6 Graded Layer Steps 6 Stages / Expression Operation lOMPa / Binding Agent lwt%

Photograph.7 Graded Layer Steps 11 Stages / Expression Operation lOMPa / Binding Agent lwt%

Photograph.8 Graded Layer Steps 6 Stages / Expression Operation lOMPa / Binding Agent lwt%

Literature cited (1) M, Niino. A, Kumagawa : Kikai-no-kenkyu. in Japanese Vol. 44 No. 4, PI, (1992) (2) Keisha-kinou-zairyou kenkyukai: The report of a functionally graded materials, Mitou-kagaku-gijutu-kyoukai. in Japanese, Pll, (1993) (3)N,Hayashi. Y,Murakami. M,Shirato. et al: Proceedings of Thailand-Japan Joint Conference Solid-Fluid Separation Technology, P13, (1995)

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

349

Microwave sintering of metal-ceramic FGJVJ M.A. Willert-Porada and R. Borchert Department of Chemical Engineering, Div. of Materials Science, University of Dortmund, 44221 Dortmund, Germany 1. INTRODUCTION In the fabrication process of metal-ceramic FGM powder metallurgy is an attractive method, particularly when densification can be achieved by pressureless sintering. In order to achieve homogeneous densification across the compositional gradient, an adjustment of the local sintering rate to the microstructure and composition of the gradient structure is required. In case of conventional sintering techniques, full densification by pressureless sintering Is only accomplished, when the sintering temperature of the metallic and the ceramic part of the FGM is adjusted, e.g. by using fine grained ceramic powders with a significantly increased sintering activity [1]. However, this method is limited by the availability of ceramic as well as metallic powders with a sufficiently high or low sintering activity. Furthermore, adjustment of sintering activity by the grain size is mainly used to lower the sintering temperature of the ceramic phase, which on the other hand also decreases the temperature range of application of such FGM. As an alternative method not based solely on powder properties, application of microwave heating (2.45 GHz frequency) offers a solution to this problem, due to the sensitivity of the microwave heating behaviour to microstructural and compositional differences within the gradient material. In previous studies on metal-ceramic or semiconductor-insulator-mixtures the effect of composition on dissipation of microwave radiation into heat turned out to be much more pronounced than the influence of grain size or agglomeration on the heating rate [2 ] as long as no nanoscaled powders are applied. However, only few metalceramic and ceramic-ceramic-systems were investigated, e.g., AI2O3-AI [3], AI2O3Steel [4 ] as well as WC-Co [3], A^Os-TIC [5] and SiC-Zr02 [6]. The present study is therefore devoted to three different metal-ceramic-systems of a particular practical importance, e.g. A^Os-steel, ZrO2-Ni80Cr20 and AI2O3-M0, with the aim on an experimental proof of the general applicability of microwave sintering for metalceramic -FGM. 2. EXPERIMENTAL Commercial powders and additives were employed for FGM-green processing. The following materials with a grain size dso as specified were used: AI2O3 a-Al203 99.7%; 0.3Mm (A16 SG, Alcoa); 8Y-Zr02; 0.3 pm (TOSOH); NiCr 80/20; < 45 |jm (H.C. Starck); Stainless Steel 1.4401, (FeCr18Ni10Mo3), < 45|jm (Goodfellow) and Mo, 99%, 3|jm (ChemPur).

350 Al-stearate (Fluka), PEG (Union Carbide) and EtOH were used for the adjustment of green density within the compositional gradient upon CIP. The best method to decrease the density of the metal-rich zones of an FGM in terms of easy removal is a liquid filler with low boiling point, like e.g. ethanol [7]. The compositional gradient materials were prepared by milling different mixtures of ceramic and metallic powders in EtOH, drying and grounding/sieving prior to form the gradient sample by sequential filling of different mixtures into a cylindrical silicone mould ((j) = 50 mm, H = 30 mm). The gradient consisted usually of 11 layers with a thickness according to (1) and a thickness after sintering between 8 - 1 6 mm. Cceramic(^) = Co —

H

(1)

c: concentration, d: thickness of the layer, H:FGM total height, X exponential factor; X=1 linear gradient The composition of the powder mixtures used was changed with 10Vol% steps. Based on exponents between 1 and 2.5 a large variety of metal-ceramic FGM is obtained using the same basic powder mixtures. The amount of a powder mixture necessary to achieve a certain compositional gradient was controlled be weighing the sample during gradient formation. The wt% were calculated from the required volume ratios and the density of the particular powder mixture. Besides an unsymmetric ceramic-metal FGM a symmetric ceramic-metal-ceramic FGM was prepared in case of M0-AI2O3. The gradient samples were CIPed at 150-300 MPa. The green density differences within a particular compositional gradient strongly depend upon the applied pressure. Therefore the CIP-conditions were adjusted to the material system. Milling of the metal-ceramic mixtures is not favourable for alumina-steel-mixtures, because of work hardening. The highest green strength is achieved in the Mo-alumina system, as shown in Figure 1. 20.0

Al203/Steel

ZrO2/NiCr8020

AI2O3/M0

Figure 1: Compressive strenght of 40wt% metal-ceramic dispersions after CIP Microwave sintering was performed in an "in house built" multi mode cavity operating with 2.45 GHz microwave radiation at a maximum power level of 2.5 kW. This cavity has movable walls, which enable the adjustment of the microwave field pattern to the load. Sintering is performed in Ar-5%H2 atmosphere, using a susceptor-free casketing [4]. In order to maximise the temperature gradient formation, the FGM green parts are introduced into a tight casket of alumina fibre board insulation. By this method, heat exchange due to convection and heat loss by radiation is significantly suppressed. Pyrometry is used for temperature measurements. Usually

351 temperature is detected at the "cold", metallic side of the gradient pointing outwards the casket to allow for radiant cooling. Comparable emissivity was established for different metal-ceramic FGM by covering the metal side of the sample with a < 1mm layer of alumina powder. Hardness was estimated by Vickers indentation on sintered gradient samples, thermal expansion coefficient measurements and elastic behaviour was tested on the different dispersions. Preliminary mechanical testing (single edge notched beam, 3PB-SENB) and oxidation tests were performed on sintered FGM's. 3. RESULTS AND DISCUSSION Among the metal-ceramic systems in this study, AbOs-steel is the one most difficult to sinter, as visible from the physical data in Figure 2. ^7~.

25.0

O

20.0 '

c o cCO

15.0

Q. X LU

10.0

"co E

5.0

^^Al203/1.4401 ^^Al203/M0 -°-Zr02 / NiCr802d

1—tt—^—^^

sz

\-

0.0

^___^___^^__ _ ^ 20

40

^ 60

80

10

Vol% Ceramic

Figure 2: Coefficient of linear thermal expansion (measured at 20 - 750°C) Besides the highest thermomechanical stress build up due to the different thermal expansion behaviour, this system requires the largest temperature gradient over the compositional gradient for pressureless densification. Therefore a systematic investigation of the microwave heating behaviour of different AbOs-steel mixtures was performed in order to calculate the thermal gradient achievable due to differences in microwave dissipation as a function of composition and density within the FGM [4]. The results obtained from microwave heating of different compositions at subsequent stages of the densification process show, that at 20°C only regions of the FGM with quite a high metallic content and a low density are effectively heated by microwave radiation with a frequency of 2.45 GHz. By heat conduction the ceramic part of the FGM is heated until it starts to dissipate microwaves into heat. In Figure 3 a typical sintering profile is shown for the microwave sintering of a linear A^Os-steel-FGM with 12-15 mm thickness. The high heating rate up to 700°C is attributed to the preferential heating of the metallic powder. With increasing densification and grain growth of the metal phase microwave dissipation in this part of the FGM decreases. At the same time the ceramic phase starts to dissipate microwaves effectively, with a major contribution to the heating process at T > 1000°C.

352 Based on these experimental data the thermal gradient developing during microwave heating of the FGM was calculated from a general heat equation, assuming relaxation of thermal gradients [7]. As shown In Figure 4, a steady state thermal gradient is developed during the microwave sintering process of an A^Os-steel FGM, with a temperature difference between the metallic and the ceramic phase of 150°200X. i1400

14UUdielectric loss of the ceramic

1200

1200

|iooo

1000

j 800 800 ohmic lo/6s of the/natal

600

Total Power

H

?00 0

u j 200 y

Absorbed MW-Power 20

40

Q.

j 400

400

(3

1 600

p

60

80

]

100

0

120

t[min]

Figure 3: Typical microwave sintering profile for a A^Oa-steel FGM sintered with microwaves of 2.45 GHz frequency. Temperature readings are taken at the metal side of the FGM.

Qi+2

Q+i

Metal

Alumina [%Vol]

Figure 4: Thermal gradient developed upon microwave sintering of an A^Os-steel FGM The thermal gradient is established in a similar way for the 8Y-ZrO2-NiCr8020 system. Because of the low heat conductivity of the Zr02-rich regions of the FGM, the compositional gradient needs to be adjusted in thickness to the thermal conductivity profile. When a variable exponent from 1.3 to 2.5 is used, no melting of the Ni80Cr20-powder occurs. The spatial distribution of Ni80Cr20 established within the green part is preserved upon microwave sintering with a thermal gradient, as indicated in Figure 5 from the NiKa-concentration profile of the sintered part.

353

i O (D

S 3

o

Figure 5: Compositional gradient in a microwave sintered 8Y-ZrO2-NiCr8020 sample as compared to the green part (top). Microwave sintering yields dense metal-ceramic FGM, as confirmed by mechanical testing (3PB-SENB), hardness measurements and from oxidation experiments. When a 8Y-ZrO2-NiCr8020 FGM is notched only in the ceramic region, the brittleness of the dense ceramic prevents stable crack growth in a controlled crack growth experiment. As shown in Figure 6, stable crack growth is achieved with specimen notched down to the metal-ceramic region of the FGM. Notch in the 50-60% metal zone: toughening effect

Notch in the 100-90% ceramic zone: brittle fracture

/I iiiiiilii

A

„400l ^300 ] § 200 _i

100

liiii i;i|i|i||||iii

Ill

iilii

iiiiil' lilK

III l l i i i i i l

IIIpiiii l i l l l I I llMpI mil iiiiil

100

iliiiii

200

Displacement [|jm]

111111111111111 iiliiii

80 1

BHI iliiiii 11

60 1iiiiiiipif liiiiiiil8!i& o 40 1 l l l l l l i l l ill

•D CD

90 -1llpllWBmSmmmmm

^yJ

li

iiH:0,2

1

i

0,4

m

H i l l ^'T™-^-^ 0,6

Displacement [mm]

Figure 6: 3PB-SENB-experiments on 8Y-ZrO2-NiCr8020 microwave sintered specimen An exponential gradient, which reduces the thermomechanical stress in the ceramic region of an FGM by increasing the thickness of this zone, could therefore be detrimental for the crack resistance behaviour of the FGM. Vickers-hardness measurements of 8Y-ZrO2-NiCr8020 and of AI2O3-M0 reveal as compared to results presented earlier for AbOa-steel a percolation threshold-like profile for 8Y-ZrO2-NiCr8020 whereas a smooth profile is obtained for the refractory metal-brittle ceramic system AI2O3-M0. Preliminary oxidation resistance testing on the 8Y-ZrO2-NiCr8020 reveals the necessity of additional components of an FGM for

354 thermal barrier coating, as shown in Figure 7. Because of the significantly increased contact area between the 8Y-Zr02 and the metal in a FGM, oxidation of the metal is enhanced as compared to a conventional thermal barrier coating, particularly because of the high ionic conductivity of 8Y-Zr02. Opposite to the successful use of metallic interlayers based on Ni-Cr-AI-Y as oxidation resistant barriers in conventional TBC's, the addition of a second metallic phase to a ZrO2-NiCr8020-FGM would not allow to protect the large ceramic-metal interface from oxygen migration. Therefore a ceramic material with a very low oxygen diffusivity and good adhesion to the metal particles was added to the metal powder used for the intermediate metal-ceramic zone of a ZrO2-NiCr8020-FGM. By this method direct contact between the metallic and the 8Y-Zr02 grains is avoided. As a result, a better oxidation resistance for the FGM is achieved as compared to the plain 8Y-ZrO2-NiCr8020 material.

1100X

with Silicate

delamination after 22 h T

20

i

30

'

40

50

Time [h]

Figure 7: Preliminary testing of the oxydation resistance of dense, microwave sintered TBC

Financial Support of DFG, SFB 316, TP A6 is gratefully acknowledged. 4. REFERENCES [1] R. Watanabe, „Powder Processing of Functionally Gradient Materials", MRS Bull. XX [1], 32-34(1995) [2] M. Willert-Porada, T.Gerdes, H. Kolaska, K. Rodiger; „Einsatz von Mikrowellen zum Sintern pulvermetallurglscher Produkte" in „Pulvermetallurgie in Wissenscliaft und Praxis Band 11", Hrsg. H. Kolaska, DGM-lnformationsgesellschaft mbH, S. 177-209 (1995) [3] T. Gerdes, M. Willert-Porada,, „Microwave Sintering of Metal-Ceramic and CeramicCeramic Composites", MRS Proc. Vol. 341, S. 531-537 (1994) [4] R. Borchert, M. Willert-Porada; „Eigenschaften pulvermetallurgisch hergestellter und drucklos mittels Mikrowellen gesinterter metallisch-keramischer Gradientenwerkstoffe, in „Verbundwerkstoffe und Werkstoffverbunde\ Hrsg. G.Ziegler, DGMlnformationsgesellschaft mbH, S. 73-77 (1996) [5] T. Gerdes, „Mikrowellensintern von Metall-Keramik-Dispersionswerkstoffen", Ph.D. Thesis University of Dortmund, 1995 [6] M. Willert-Porada, "Reaction Rate Controlled Microwave Processing of Ceramic Materials", Ceram. Trans., 36 , S. 277-286 (1993) [7] R. Borchert, unpubl. results (1996)

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

355

Residual Stress Control of Functionally Graded Materials via PulseElectric Discharge Consolidation with Temperature Gradient Control Hiroshi Kimura and Tatsuya Satoh Department of Mechanical Engineering, The National Defense Academy, 239 Yokosuka, Japan

This article describes how one can control a residual stress profile of a functionally graded material via pulse-electric discharge consolidation with temperature gradient control, using the graphite die with specially designed outer shape. For a five layered materials graded from TiAl intermetallic to partially stabilized zirconia(PSZ) with the compositional distribution of n=l, the temperature gradient necessary for a constant sintering rate is experimentally derived. Then, the stepped die, having a temperature gradient larger than the ideal one, can be used to show an increase in vickers hardness from the prediction of the mixture rule at the PSZ layer and the layers with high volume fractions of PSZ. This increase strongly indicates the occurrence of a residual compressive stress during consolidation up to full density.

1. INTRODUCTION Functionally graded materials(FGMs) are becoming one of promising advanced materials for the feasibility of next generation technologies such as space planes. In order to provide a route to the design and development of FGMs with required performance, the temperature gradient controlled electric discharge consolidation method using a die with a specially designed outer shape have been proposed by the author and his coworkers[l-5]. This paper describes a basic procedure of controlling a residual stress profile to develop a strengthened FGM.

2. EXPERIMENTAL PROCEDURE Figure 1 illustrates the instrumented method of pulse electric discharge resistance consolidation with temperature gradient control for the control of residual stress profiles of FGMs. This instrumented apparatus provides means of controlling temperature, applied pressure, electric field, and plunger displacement. A temperature gradient profile at the surface of the die is measured in heating by thermography with data acquisition system.

356

Oscillograph

Displacement

Temperature

Temperature gradient

System controller

Figure 1 The pulse-electric discharge consolidation method with temperature gradient control for the residual stress control of functionally graded materials. 3. PROCESS DESIGN FOR RESIDUAL STRESS CONTROL OF FGM Figure 2 shows the basic procedure of controlling a residual stress profile of a Compositional gradient , Height(hj>, FGM at will via temperature gradient Residual stress profile controlled electric discharge resistance ( 2 ) Die E)esign consolidation. The demand placed by Derivation of ideal temperature profile service requirements on the application TFOM VS h ,TFGM (h) vs T of FGMs is to realize the optimized compositional gradient, the specimen Determination of outer shape of die height(hf) and the stress field profile. diameter difference, slope The first step is the design of the suitable ( 3 ) Process Control outer shape of the die based on an ideal Consoldating process variables temperature profile necessary to obtain a e = f(a,T, o ,A) constant consolidation rate along a graded composition. The next step is to Real-time measurements of relative density and die temperature gradient set up the process condition using Ty, DvsT 1 consolidation variables of applied pressure(a), temperature(T) and current Calculation of temperature difference density(A), and then to measure the ATFGM V S T relative density(D) of a FGM compact ( 4 ) Evalua ion of Residual Stress and the temperature gradient during Measurement of vickers hardness | consolidation. Finally, the residual Hy vs h 1 stress profiles involved in the FGM are Figure 2 Process design for the control evaluated by the exact measurements of hardness across height and radius. residual stress field of FGMs. ( 1)

Material Design

Demands

placed by service requirements

357 4. RESULTS AND DISCUSSION 4.1. Die design Figure 3 shows the temperature gradients necessary to obtain the constant sintering rate along the graded composition consisting of pure layers of amorphous TisoAlsoand PSZ separated by layers mixed in the proportion 3:1,1:1,1:3 in the case of hf=6.504. This FGM is characterized by n=l for the compositional distribution function of f(h)=hn where f(h) is a volunne fraction of PSZ vs distance from the bottom of FGM and n is an exponent. The result is derived from the experimentally obtained temperature dependence of consolidation rate of each powder to full densification under 49 MPa[3]. Figure 4 is a schematic of the stepped die with external diameters of 55 and 40 m m along loading axis. Figure 5 shows a change of the temperature profile in heating as a function of distance from the bottom of the stepped graphite die in the case of one way loading. Figure 6 shows the difference(TY-TY25) of the surface temperature estimated from the temperature gradient across a radius of the die and the temperature slope(a) for the cylinder with the diameter of 55 m m as a function of surface temperature at Y=20 m m . The temperature profile along the graded composition in the case of the stepped die is depicted in Fig.3. It can be seen that the temperature shows an increase at the PSZ layer and a decrease at the TiAl layer. These deviations from the ideal temperature gradient necessary for constant consolidation can be used to control a sintering rate of each layer, which is conducive to a suitable residual stress profile of a five layered FGM of TiAl intermetallic and PSZ system.

r ><

^ .<-04(

055 p^20

>

mm

y

>k 1

Tds

0

2

4

6

8

10 12

Distance , h / mm Figure 3 Temperature gradients along a 5 layered TiAl/PSZ FGM with n=l, necessary for constant sintering rate and temperature profile of stepped die.

Td Ts

Tc

t

40 25

± 1 20

1 >

> X

Figure 4 The stepped graphite die, used in this study, for the control of residual stress profile of FGM via pulse- electric discharge consolidation.

358 800 - ^ 4 0 mmStepped die ^

(0 55 , 040 m m )

1100

TY-TIS y=Y-25

600 1000 400 y=5.7 mm

0

10

20

30

40

500

Distance from the die bottom , Y/ mm

Figure 5 Temperature profile at the surface of the stepped graphite die in heating, taken by thermography.

700

900

1100

1300

1500

Surface temperature at Y=20 mm , T.20/ K

Figure 6 Temperature differences and slope at the surface of consolidated bulk as a function of temperature(Y=20 mm).

4.2. Process control Figure 7 shows densification of five layered compacts graded from amorphous TiAl to PSZ via temperature gradient controlled electric discharge resistance consolidation using the stepped graphite die as shown in Fig.4. The consolidation of FGMs having hf = 7.40 and 6.52 under two different temperature gradient show

600 800 1000 1200 0 Surface temperature , Tsio/ K

100 200 Time , t / s

300

Figure 7 Densification of 5 layered TiAl/PSZ FGMs during heating and at the constant holding temperatures. This figure includes the case of amorphous TiAl.

359 1000

^ 8-

800 h

600 CD

S 400 h 05 QJ U U

200

1100 1300 Temperature at Y=20 mm , Tsz Figure 8 Differences of the surface and top temperature between the top and bottom during electric discharge resistance consolidation of 5 layered TiAl/PSZ FGM. definitely densification rate higher than that of mechanically alloyed amorphous TiAl powder[6] as depicted in Fig.7. Figure 8 shows the differences of the surface and center temperature between the top and bottom of five layered compacts as a function of temperature at Y=20 m m . The FGM with hf=7.4 m m undergoes a temperature difference of 900 K at the m a x i m u m under TTiAi=1376 K, while the FGM with hf=6.52 m m is close to the temperature difference necessary for constant rate consolidation[3-5]. Note that the difference of the center temperature of the TiAl/PSZ FGM is larger than that of surface temperature. 4.3. Evaluation of residual stress profile Figure 9 and 10 shows the vickers hardness as a function of height(h) from the bottom of TiAl intermetallic layer at the center of five layered FGM produced by a relatively large temperature difference and nearly equal to constant sintering rate, and hardness at the surface of FGM. The vickers hardness(Hv) for TiAl/PSZ FGM, produced by amorphous TiAl powder under the almost constant sintering rate, is fairly well expressed by the following equation[3]: Hv = HvTiAi Vf + HvPsz(l-Vf) (1) where Vf is a v o l u m e fraction of TiAl, HvXiAi and HVPSZ are the vickers hardness for TiAl intermetallic and PSZ respectively. W e see that the vickers hardness for FGM shows great increases from the variation of hardness according to a mixture rule at the center of the PSZ layer and layers with a high volume fraction of PSZ. These increases indicate the occurrence of the residual compressive stress field during electric discharge resistance consolidation with a relatively large temperature gradient as shown in Fig.8.

360 2500

2400 5 layered TiAl/PSZ FGM

:z; Q 2000

—

QJ 1600[-

mo • a AA TV

5 layered TiAl / PSZ FGM

TiAl 3:1 1:1 1:3

• O Psz

1376 K-i

e

^^^1 1200

k

Tr\

/

1

i

y _ v_.i j

1325 K

r

1"

=Hm,V,+Hr,z(l-V,)

Jn 800 h:--:^. Center

> 400

1

1

1

1 2

1 3

1 4

1 5

1

Height, h / mm

Figure 9 Vickers hardness for 5 layered TiAl/PSZ FGMs, prepared by different process conditions, versus height.

2

3

4

Height, h / mm

Figure 10 Vickers hardness at the surface and center of FGM, consolidated with a larger temperature gradient.

5. CONCLUSIONS This article describes how one can control an internal stress fields of FGM at will during pulse electric discharge resistance consolidation having temperature gradient control using the graphite die with specially designed outer shape. This methodology is applicable to the design and development of high strengthened functionally graded materials in widespread applications, when using graphite dies with optimizing outer shapes for various compostional gradients and heights[7-8].

REFERENCES 1. H. Kimura and K. Toda, Japanese Patent pending. No. 117716,287321(1993) 2. H. Kimura and S. Kobayashi, J. Jpn Soc. Powder Powder Metall., 39(1992) 287. 3. H. Kimura and S. Kobayashi, J. Jpn Inst. Met., 57(1993) 1346. 4. H. Kimura and K. Toda, Proc. Conf. Euro PM'95(Birmingham UK, 1995 Oct.), 1(1995) 647. 5. H.Kimura and K. Toda, Powder Metall. 39(1996) 59. 6. H.Kimura, Phil. Mag A, 73(1996) 723. 7. H.Kimura, J. Metals Technology(Kinzoku), 66, No.2(1997), in the press 8. H.Kimura and Y. Tsushima, to be published.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

361

Study on the Composition Graded Cemented Carbide/Steel by Spark Plasma Sintering Akihiko Ikegaya ^ , Katsuya Uchino ^ , Tsugio Miyagawa " and Hidenori Kaneta "" ^ Kami Research Laboratories, Sumitomo Electric Industries, Ltd. 1 - 1 - 1 Koya- Kita, Itami 664, Japan ^ Japan Research and Development Center for Metals, 1 - 26- 5 Toranomon Minato- ku, Tokyo 105, Japan "" Technology Research Center, Japan National Oil Corporation, 1 - 2 - 2 Hamada Mihama-ku, Chiba 261, Japan Abstract: Steel coated with high wear- resistant and high corrosion- resistant hard materials are widely used, but these composite materials do not always achieve the aimed properties, mainly because coating materials with high hardness could not be applied thickly and also could not be jointed firmly as desired. To overcome these problems, the new functionally graded materials were studied. To relieve the thermal stress, Co composition graded cemented carbide powders stacked in multiple layers on steel substrate were sintered and joined on the steel by the Spark Plasma Sintering. Because of low- temperature and short- time sintering, the graded composition of cemented carbide was maintained as desired and a crack- and peel- free coated product could be achieved by this process. From this study, we comfirmed that it is possible to realize a high wear- resistant cemented carbide/steel composite material for practical use. 1. INTRODUCTION Steel coated with high wear- resistant and high corrosion— resistant hard materials like stellite and cemented carbide are widely used. But these composite materials do not always obtain the properties which coating materials originally have. Thick and dense application, especially of high- hardness cemented carbide with low metal content, could not be done, as well as achieving desired firm jointing. This is the result of cracking of coated product and the peeling of coated layer caused by thermal stress owing to the mismatches of the thermal expansion co- efficient between the steel and cemented carbide. If the high- hardness cemented carbide could be densely and thickly coated, it is possible to realize the excellent properties which the hard material originally have. To overcome these problems, a method of sinter- bonding composition graded cemented

362 carbide powders on steel substrate was proposed. But it is difficult to maintain the graded composition, because the Co moves easily during liquid phase sintering. So it needs to sinter-bond two pre-sintered bodies with different Co compositions by inserting the same cemented carbide powder as one side body [1]. Lately, the spark plasma sintering has been used gradually in practical applications. The characteristics of this new sintering process are low- temperature and short- time sintering. The features of this process are expected to solve the above problems. Therefore, we tried to sinter- bond the composition graded cemented carbide powders on steel substrate by spark plasma sintering, and evaluated the structural and mechanical properties of this composite material. This paper reports the results of these evaluations. 2. EXPERIMENTAL PROCEDURE WC and Co powder were used as starting materials, and were combined to three mixtures of WC-10, 15 and 20%Co by weight, wet- mixed in a ball mill for 7 hours. Low carbon steel substrate of 0 30 x 5mm ^ was prepared. Figure 1 shows the e x perimental procedure. As shown, after the carbon die and steel substrate were set, WC - 20%Co powder was poured on the steel substrate and pre- pressed at 1.4MPa. The same operation was repeated using WC- 15%Co and WC- 10%Co powder in that order. The quantity of each powder was adjusted to 0.5mm thickness after sintering. Then the green compact of graded composition on steel was set in the spark plasma sintering equipment and was sinter- bonded with the steel substrate. The load of 50MPa was gived to the green compact during sintering. The temperature was measured by an optical pyrometer at the point of the carbon die surface located the interface between the cemented carbide and steel substrate. This temperature was regarded as the sintering temperature in this experiment and selected between 1243K to 1323K. In order to confirm the effect of stress relief, single layer of WC- 10%Co and WC- 15%Co sinterbonded on steel substrate were prepared, respectively. Moreover, the same graded green compact was sintered by the conventional sintering process at the temperature of 1673K under vacuum. To evaluate the structure of the sintered cemented carbide, a cross section of the sample was ground with a #200 GC grindstone, then mirror- polished using a #3000 diamond grid. In order to eliminate the compressive stress introduced by grinding, it Pre-pressing powder of c e m et^d n tcarl ^rbide^_^

Spark Plasma Sintering

steel load^4MPa , ^ ^i.^trode substrate ^^R^^~^20%Co^l5%Co••lo%Co r n r n j ^ U—I -L
363 was polished from the surface to the depth of 60 Id m. The sample was observed using optical microscopy and SEM. EPMA analysis was also performed to determine the Co distribution. The residual stress was measured in both the horizontal and the vertical directions to the bonded interface of the sample by the x - ray sin xjj ^ method. For the mechanical characteristics, the hardness distribution was measured from the surface to the inside of sample using a micro Vickers hardness gauge. 3. RESULTS AND DISCUSSION 3.1. Effect of sintering temperature on densification of cemented carbide The microstructure of samples sintered at various temperatures were observed by optical microscopy, and the micrographs are shown in Photo. 1. As shown, the number of pores tends to reduce as the sintering temperature increases. As to Co content at the same sintering temperature, the number of pores tends to reduce as the Co content increases. Relative density was estimated from micrographs using classification standards for apparent porosity of cemented carbides [2]. The results are shown in Fig. 2. Both WC- 10%Co and 15%Co are densified completely at more than 1283K. WC-20%Co is densified completely at more than 1273K. The liquid Co flows out of the carbon die at more than 1373K. Because lower sintering temperature is preferable to reduce thermal stress, the suitable sintering temperature seems to be 1283K. 3.2. Hardness and Co content distribution in the graded cemented carbide/steel The hardness distributions by spark plasma sintering at 1283K were measured, and the results are shown in Fig. 3. For comparison with the conventional process, the results sintered at 1673K are also shown in the same figure. As to the sintered sample by the conventional process, the hardness is nearly uniform although the green compact has graded composition. On the contrary, the hardness of the sintered sample by spark plasma sintering changes in 3 steps corresponding to graded green compact. The mi-

liquid Co flows out!

WC-10%Co 100/zm Photo. 1 Optical micrographs of the Co content graded cemented carbides by Spark Plasma Sintering as a function of sintering temperature

-ih1240

1280 1320 1360 1400 Sintering temperature(K) Fig. 2 Effects of sintering temperature on relative density of the Co content gradient cemented carbides 0

364 20

10%Co

Steel

CO

10%Co

o 5^ 15

iS7rCo

CO

c/:

10

conventional sintering ^'V/V-

207rCo 0 -500

207rCo'

''S

0 500 1000 1500 2000 D i s t a n c e from i n t e r f a c e (/z m)

Fig. 3 Profiles of microhardness on Co content graded cemented carbide-steel by Spark Plasma Sintering (1283K) compared with conventional sintering(1673K) steel

»^emented carbic]|g

(a)

(b)

lO/ini

Photo. 2 Microstructures of each section of 3 layers stacked cemented carbide: (a) Spark Plasma Sintering; (b) conventional sintering steel 10%Co steel 15%Co

steel

20%Co

^^)|flHH (b). ^^^^: (c) 'A.

100/im

Fig. 4 Variations of Co, W and Fe contents of 3 layers-stacked cemented carbide sintered by Spark Plasma Sintering at 1323K

Photo. 3 Optical micrographs of the composite materials: (a) WC-10%Co and steel; (b) WC-15%Co and steel; (c) 3 layers-stacked cemented carbide and steel

crographs of each part of stacked layers sintered by both sintering processes are shown in Photo. 2. As to the spark plasma sintering, the Co exists differently, corresponding to graded green compact. On the contrary, the Co exists uniformly in the cemented carbide by the conventional process as expected. Co, Fe and W content distributions in the graded cemented carbide/steel were measured by EPMA, and the results are shown in Fig. 4. The Co shows two sharp changes stepwise corresponding to graded composition of green compact. From these results, the hardness distributions are believed to be due to the graded composition being maintained after sintering.

365 3.3. Effect of graded structure on thermal stress To evaluate the effect of the graded structure on thermal stress, the cross section of the bonding interface of samples was examined. The micrographs by optical microscopy are shown in Photo. 3. Cracks were found near the interface in the cemented carbide of the sample sinter-bonded with single layer of WC- 10%Co. There are no cracks in the samples sinter-bonded with graded cemented carbide and single layer of WC- 15%Co, however. For the same above mentioned samples, the residual stress from the cemented carbide to the steel was measured in both the horizontal and the vertical directions to the bonded interface. The results are shown in Fig. 5. The residual stress of cemented carbide in the vertical direction is shown in the circle. Because the values are nearly zero, the compressive residual stress of cemented carbide layers observed in the horizonta direction is due to the difference of the thermal expansion c o - efficient between the steel and the cemented carbide. The residual stress of graded cemented carbide near the interface is small compared with the two kinds of single layer cemented carbide. In case of the single layer cemented carbide, the compressive residual stress of c e mented carbide is larger in the inside near the interface than in the surface. However, the compressive stress of graded cemented carbide is smaller in the inner layer than in the outer layer. This seems to be the result of the graded structure. Although the compressive residual stress of WC- 10%Co single layer is smaller than that of WC- 15%Co single layer, it is because the WC- 10%Co layer was relieved of the stress by the occurrence of cracks as shown as Photo. 3. From the above results, it is confirmed that the Co composition graded cemented carbide/steel composite sinter- bonded by spark plasma sintering is sintered densely and also relieved of thermal stress. Therefore this composite material is expected to display the excellent properties as desired.

CO OH

O

03

Steel

^ cemented cemented

-2000 -1000 0 1000 Distance from interface (/z m) Fig. 5 Profiles of residual stress on the composite materials: (a) horizontal direction ; (b) vertical direction

Photo. 4 SEM micrographs of WC-15%Co cemented carbides: (a) sintered by Spark Plasma Sintering (at 1283K); (b) sintered by conventional sintering (at 1673K).

366 3.4. Characteristics of microstructure sintered by spark plasma sintering SEM micrograph of WC- 15%Co cemented carbide sintered by spark plasma sintering at 1283K is shown in Photo. 4 compared with the same one sintered by the conventional sintering process. Co dispersion of cemented carbide sintered by spark plasma sintering is not as good as that sintered by the conventional method, and the Co exists in the shape of a pool. As to WC particles, WC particles sintered by spark plasma sintering is much smaller than those sintered by the conventional process, and the shape is round, while that sintered by the conventional process is square. We believe these phenomenon to be the suppression of Co dispersion and WC grain growth due to lowtemperature and short- time sintering of the spark plasma sintering. It is known that the densification of cemented carbide is possible by solid state sintering without a liquid phase, but it is thought to be difficult to obtain full density at 50 MPa press load and 6 - minute sintering time. This is supported by studies examining the effect of sintering temperature and pressure in the HIP process on densification of cemented carbide [3]. Combined with the results above mentioned, the real sintering temperature seems to be considerably higher than 1283K, and to be slightly above the temperature that liquid phase appears. 4. CONCLUSION In order to overcome the problem of cracks and peeling of hard coating layer, the composition graded cemented carbide sinter- bonded on steel using the spark plasma sintering was studied. The following results were obtained on this new material. : (1) By using spark plasma sintering, the Co composition graded green compact of c e mented carbide sinter-bonded on steel substrate could be sintered densely and the d e sired graded structure was maintained after sintering. (2) The thermal stress of graded cemented carbide near the interface is smaller than non- graded cemented carbide, and it was confirmed to achieve a crack- and peel- firee coated product by a graded structure. (3) The microstructure of cemented carbide sintered by spark plasma sintering is suppressed Co dispersion and growth of WC particles compared with the conventional process, and it is attributed to low temperature and short- time sintering of this process. From this study, we comfirmed that it is possible to realize a high wear- resistant cemented carbide/steel composite material for practical use. This paper was refered to another paper [4] written by the same co- authors.

REFERENCES 1. S. Kamota, A. Sakai, et al. , Japan Patent 7-3306. 2. ISO 4505 (Hardmetals Metallographic determination of porosity and uncombined carbon) 3. H.Suzuki, ed. , Cemented carbides and sintered hard materials, Maruzen Japan, 1986, p. 30 4. K. Utino, Journal of the Japan Society of Powder and Powder Metallurgy, 43(1996), 472

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

367

Phase Composition Profile Character of a Functionally-graded Al2Ti05/Zr02-Al203 Composite

S. Pratapa, B.H. O'Connor, and I.M.Low Materials Research Group, Department of Applied Physics, Curtin University of Technology, GPO Box U 1987, Perth, WA, Australia 6001 Abstract. The phase composition character of a functionally-graded Al2Ti05/Zr02-Al203 composite, produced by infiltration involving an Al203-Zr02 (90:10 by weight) green body and a solution containing TiCls, has been determined by x-ray diffraction. The Rietveld "whole pattern" refinement method was applied to diffraction patterns of the sample which were collected from the surface and at depths of 0.1, 0.3, 0.4, 0.8, 1.2, and 1.5 mm. The sub-surface measurements were made after polishing away the material to the designated depth. Absolute phase composition levels were estimated using diffraction data from an a-AliOs external standard. The results show that the weight fraction of Al2Ti05 is approximately 44.5% at the surface and reduces linearly to 9.5% at 0.3 mm, and then to 5.3% at 1.5 mm. The AI2O3 content increases with depth in a complementary manner. The results also indicate the presence of an amorphous phase of mean value 3 wt% in the specimen. The fiinctionally-graded profile was verified by qualitative energy-dispersive x-ray microanalysis. INTRODUCTION Microstructures of functionally-graded materials (FGMs) may be graded according to composition, fibre-orientation, crystal structure, porosity, particle size, etc. Most FGMs produced are of graded-composition type^^l Various methods have been applied to produce FGMs. For example, CVD, electro-deposition, sol-gel, and slip casting methods have been used to produce ceramic-ceramic FGMs. Liquid infiltration into porous ceramic green bodies has also been used to produce ceramic-ceramic FGMs^^'^l A fiinctionally-graded Al2Ti05/Zr02-Al203 synthesized by this technique is reported here. There are several ways to reveal the graded character of FGMs. For example, Marple and Green made use of quantitative electron-probe microanalysis to characterise the compositional change in their alumina/mullite FGM system^^l In this study, the graded profile of phase compositions in functionally-graded Al2Ti05/Zr02-Al203 was characterised using the Rietveld x-ray diffraction analysis, with mass attenuation corrections being performed using xray emission spectrometry. Energy-dispersive x-ray microanalysis was utilised to qualitatively verify the graded-composition character of the material. EXPERIMENT Zirconia-alumina powder was obtained by wet ball milling 90 wt%> a-Al203 (AlOOOSG grade Alcoa, USA, of median particle size 0.39 jim) and 10 wt% monoclinic zirconia (SF Ultra Z-Tech, Australia, of median particle size 0.40 |^m). The slurry was then dried and sieved until

368 free-flowing (45 [xm grid-size). The powder mixture was pressed uniaxially in a metal die to a pressure of 37 MPa to yield a bar sample with dimensions of 5 mm x 12 mm x 60 mm. Partial sintering at 1000° C was used to increase the strength and retain the porosity of the green body prior to infiltration. Infiltration of porous preforms was conducted at room temperature by completely immersing in a 30 wt% TiCb-contained solution (BDH Limited Poole, England) for 24 hours. The infiltrated preforms were then dried at room temperature for 24 hours. They were finally sintered in a Ceramic Engineering (Model HT 04/17) high-temperature fiimace at 1°C min"^ to 450° C for 30 minutes, followed by 5° C min"^ to 1550° C for 3 hours, and then the fiimace was naturally cooled. An alumina-zirconia (90:10 by weight) uninfiltrated control sample was prepared following the same sintering condition. The central part of the sample (10 mm x 10 mm in surface dimension) was obtained from the as-fired sample for composition graded characterisation. X-ray diffraction patterns and xray emission Compton scatter intensity of the specimen at depths of 0.0, 0.1, 0.3, 0.4, 0.8, 1.2, and 1.5 mm were measured using a Siemens D500 x-ray diffractometer (CuKa tube) and a Siemens SRS200 wavelength-dispersive x-ray emission spectrometer (MoKa tube), respectively. A high-purity a-alumina powder (Praxair Surface Tech. Inc.) was used as an external standard specimen in the x-ray diffraction measurement. The Compton scatter intensity was converted into the corresponding mass attenuation coefficient value using a calibration curve which was developed from a set of "standard" alumina-zirconia compact powders containing 0 wt% to 20 wt% zirconia. The sub-surface measurements were made after polishing the material to the designated depth. Further details of the measurements are described elsewhere^^l Rietveld "whole pattern" refinement was applied to each pattern to determine the scale factors of all detectable phases. The goodness-of-fit indices for the refinements, (R^p/Rcxp)^^, ranged from 0.02 - 0.03 wliich is adequate for phase composition determination. The weight fraction of phase / at each depth was determined by an external standard 'ZMV' approactf^'^^

where 5, and Ss denote the Rietveld scale factors of sample / and the external standard, respectively. Z, is the number of formula unit of phase / with mass M in the unit-cell volume Vj. The mass attenuation coefficient of the standard /^ was calculated whereas that of the specimen ju* was determined by Compton scattering measurement. Amorphous phase content was determined by subtracting the total concentration of crystalline phases from unity (see reference 8 for example), WG=l-SWi.

(2)

1=1

Energy-dispersive x-ray microanalysis was used to qualitatively verify the graded profile. A cross-sectioned sample, polished to a l|am finish, was prepared. The measurement was conducted from the near-surface region to the center of the sample with step size of 50 |im using a JEOL 35C scanning electron microscope. The x-ray emission intensities for TiKa, AUCa, and ZrLa were collected at each point.

369 RESULTS AND DISCUSSION

Ji. "\k4x\x 'u

AT

'i?

AT A

?

a

.

cj

Figure 1. X-ray diffraction patterns of the functionally-graded aluminium titanate/zirconia-alumina composite at various depths. Labels: AT=aluminium titanate, A=a-alumina, Z=zirconia (m: monoclinic, t: tetragonal). CuKa radiation was used. Figure 1 shows the x-ray diffraction patterns of the FGM measured at several depths. The aluminium titanate (P-Al2Ti05) in the sample is believed to form through the sintering reaction a-AlzOs + Ti02 (rutile) -> P-Al2Ti05

(3)

at approximately 1390° C^^l The figure shows that the intensity of aluminium titanate (AT) lines decreases gradually from the surface to the center of the sample. On the other hand, the intensities of a-alumina peaks increase with depth. The tetragonal zirconia peak intensity appears to increase slightly with depth. The figure shows no other titania-related peaks which indicates that Ti02 (rutile) had reacted completely with a-alumina to form AT. The quantitative phase analysis results are shown in Figure 2 and Table 1. As can be seen from the figure, the amount of AT is 44.5 wt% on the surface and reduces Unearly with depth to 9.5 wt% at 0.3 mm, and then to 5.3 wt% at 1.5 mm. By contrast, the a-alumina content increases linearly with depth from 44.4 wt% at the surface to 80.2 wt% at 0.3 mm, and then 85.7 wt% at 1.5 mm. This suggests that the kinetics of infiltration are time-dependent and thus the amount of infiltrant reduces with depth. Clearly, liquid infiltration is a useful method to produce FGMs, as also indicated by other researchers^^'^l

370 100

80

-

60

h

40

li

I

I

I

I I I

+ Alumina oAT

I

20

5 0.0

5

i

1

0.2

0.4

5

5 1 0.6

1 0.8

1 1.0

1 1.2

5 1

1.4

Sample Depth (mm)

Figure 2. Weight ifraction of aluminium titanate (AT) and a-alumina according to depth in the functionally-graded aluminium titanate/zirconia-alumina composite. Error bars indicate 2x estimated standard deviations.

Table 1 Weight Fraction of Minor Phases and Mass Attenuation Coefficient of Aluminium Titanate/Zirconia-Alumina Composites as a Function of Sample Depth. Depth (nun) Wt(t-Zr02) Wt(Amor.) MAC 1 Wt(m-ZK)2) % % % cmV' i 5.3(3) 1.0(2) 5(2) 0.0 56.2(12) 5.8(2) 0.7(1) 0.1 52.0(10) 3(2) 7.0(2) 0.3 2(2) 41.6(8) 1 1.2(1) 7.0(2) 0.4 41.2(9) 5(2) 1.0(1) 0.8 6.8(2) 4(2) 41.2(9) 1.1(1) 6.5(2) 2.4(1) 1.2 0(2) 40.2(5) 6.7(2) 2.1(1) 0(2) 39.4(5) 1.5 Wt : weight fraction. m-Zr02 : monoclinic zirconia. t-Zr02 : tetragonal zkconia. Amor. : amorphous phase. MAC : mass attenuation coefficient of specimen at CuKa wavelength. Parenthesised figures represent the estimated standard deviation in terms of the least-significant figure to the left.

371 There is some indication of a marginal increase in the weight fraction of tetragonal zirconia (Table 1) with depth. In the alumina-zirconia (90:10 by weight) control sample, the content of the t-phase is approximately 5 wt%. On the surface of the FGM sample, where AT is approximately 45 wt%, the weight fraction of t-phase is 1.0% and this value increases up to 2.4% at a depth of 1.2 mm. It is suggested that the presence of AT has induced tensile residual stresses which are responsible for enhancing the t-->m phase transformation. The graded character is also shown by the change with depth of the mass attenuation coefficient (MAC) values (Table 1). The theoretical MAC values of AT, alumina, and zirconia are 72.9, 30.4, and 101.5 cmV\ respectively^^^l This suggests that AT has a significant contribution to the final MAC value at each depth. Since the amount of AT decreases with depth, the MAC value should also decrease with depth. The Rietveld 'external standard' method has allowed the amount of amorphous material to be computed. The presence of this phase is probably due to the incomplete crystallisation of the infihrant precursor during calcination. Table 1 depicts the content of this phase at each depth with average value of 3 wt%. Calculation shows that the MAC value of this material at each depth is between 69.8 and 125.8 cmV^ Since the theoretical MAC values for Ti02 and AliTiOs are respectively 127.3 and 72.9 cmV^ ^^^\ the detected amorphous phase could be either amorphous Ti02 or Al2Ti05^^^l Energy-dispersive x-ray microanalysis was used to qualitatively verify the graded character. Figure 3 shows the plot of the x-ray emissions of TiKa, AlKa, and ZrLa versus sample depth. Titanium emissions gradually reduce with depth whereas those of aluminium and zirconium are fairly constant. The titanium emission reduction agrees with the composition measurement using x-ray diffraction. This suggests that the infiltration has led to the formation of an FGM. Similar graded profiles were obtained for muUite/alumina system by infiltration^^l Therefore, these resuhs complement the result of x-ray diffraction quantitative phase analysis.

15

120

Al - ^ 96

12

H72

rV ^^

48

124

1

1

1

0.2

0.4

1

1

1

1

1

0.6

0.8

1

1.2

1.4

1

Sample Depth (mm)

Figure 3. X-ray characteristic emissions of TiKa, AlKa, and ZrLa with depth of FGM sample measured using energy-dispersive microprobe analysis. Error bars indicate 3x estimated standard deviations.

372 SUMMARY It can be concluded from the study that: 1. A functionally-graded aluminium titanate/zirconia-alumina composite has been synthesised by a liquid infiltration. 2. X-ray diffraction Rietveld analysis, with attenuation corrections applied using x-ray emission Compton scatter measurements, provided a powerful means for revealing the graded composition character of FGMs. Its ability in analysing phases rather than elements makes this method superior over the electron-probe microanalysis technique, especially for determining tetragonal zirconia and the amorphous phase contents. 3. The presence of aluminium titanate appears to reduce the content of tetragonal zirconia possibly due to the formation of AT-induced residual stresses in the microstructure. ACKNOWLEDGMENTS One of us (S.P.) is very grateful to the Australian Agency for International Development (AusADD) for scholarship support. We thank our colleagues Prof Deyu Li and Arie van Riessen for useful discussion. REFERENCES L Hirai, T. (1996). "Functional Gradient Material." In: Processing of Ceramics (Part 2). Ed. Brook, R.J., VCH Verlagsgesellschafift mbH, Weinheim. 2. Marple, B.R. and Green, D.J. (1989). "Mullite/Alumina Particulate Composites by Infiltration Processing". J, Amer. Ceram. Soc. 72[11], 2043-2048. 3. Low, I.M., Skala, R, Richards, R. and Perera, D.S. (1993). "Synthesis and Properties of Novel Mullite-Zirconia-toughened Alumina Composites". J. Mater. Sci. Lett. 12, 19851987. 4. Low, I.M., Skala, R.D. and Zhou, D. (1995). "Synthesis of Functionally-gradient Aluminium Titanate/Alumina Composites". J. Mater. Sci. Lett. 15, 345-347. 5. Pratapa, S. and Low, I.M. (1996). "Synthesis and Properties of Functionally-gradient Aluminium Titanate-Mullite-ZTA Composites". J. Mater. Sci. Lett. 15, 800-802. 6. Pratapa, S., O'Connor, B.H. and Low, I.M. "Use of Compton Scattering for Attenuation Corrections in Rietveld Phase Analysis". In preparation. 7. Jordan, B., O'Connor, B.H. and Li, D. (1990). "Use of Rietveld Pattern Fitting to Determine the Weight Fraction of Crystalline Material in Low Quartz Specimens". Powder Diffraction 5(2), 64-69. 8. Latella, B.A., Burton, G.R. and O'Connor, B.H. (1995). "Use of Spodumene in the Processing of Alumina-matrix Ceramics - Influence on Microstructure and Mechanical Properties". /. Amer. Ceram. Soc. 78[7], 1895-1899. 9. Pratapa, S. and Low, I.M. (1996). "The Effects of Spodumene Addition on Properties of Functionally-graded Aluminium Titanate/Zirconia-toughened Alumina Composites". In: Proceedings of the 2nd International Meeting ofPacific Rim Ceramic Societies. 10. O'Connor, B.H. and Thomas, A.G. (without year). X-ray Analysis Toolkit. Version 3.0. 11. Feltz, A. and Schmidt, F. (1990). "Preparation Study of Amorphous Al2Ti05". J. Eur. Ceram. Soc. 6, 107-110

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

373

The Use of a Functionally Graded Material in the Manufacture of a Graded Permittivity Element S.WATANABE*. T. ISHIKURA*, A. TOKUMURA*. Y. KM**. N. HAYASHl*, Y. UCHIDA*, S. HIGA*, D. DYKES***. G. TOUCHARD**** * Aichi Institute of Technology, Yachikusa Yakusa-cho Toyota 470-03 Japan. ** Yongwol Institute of Technology, Yongwol Kangwondo. 230-800 Korea. *** Yokkaichi University, Kayo-cho, Yokkaichi, 512 Japan. ****Universit6 de Poitiers, 40 av. du Recteur Pineau, Poitiers 86022 France.

1. ABSTRACT The authors of this paper have previously developed a vacuum filtration technique for the manufacture of functionally graded materials (FGMs) by a progressive lamination method. For this, they have been granted a U. S. patent. Using this method, it is possible to manufacture FGMs with thicknesses ranging from several millimetres to several centimetres. The authors have already produced an iron ( m ) oxide-kaolin FGM. the graded condition of which they have verified by means of a scanning electron microscope. In addition, they have performed measurements to determine this material's electrical properties viz: conductivity, relative permittivity and magnetic permeability. The present paper reports an experiment to create a different type of FGM, characterised by graded permittivity. The constituent materials used are titanium oxide and kaolin. The graded condition of manufactured specimens was investigated by mean of scanning electron microscope photographs and measurements of relative permittivity. Relative permittivity was found to vary between 2 and 5, while the photographs confirmed that the specimens were smoothly graded. On the basis of these results, it seems probable that graded permittivity elements can be manufactured using the authors' method.

2. INTRODUCTION Technology has recently been developed in Japan to manufacture materials which combine two or more constituent substances in graded proportions, as a means to achieve thermal relaxation. Such materials are known as functionally graded materials (FGMs). Application fields for such materials are found in mechanical, chemical, biological and

374 electrical engineering. In electrical engineering, they have utilisation potential in feeler sensors, resisters, magnetic shields, lossless optical fibres and superconductors. Methods of manufacture vary, but include the chemical and physical vapour deposition methods, the electrolytic deposition method, the atomised metal spray method, and a method in which powdered material is first melted in a plasma jet and then deposited ^\ But none of these methods or other previously existing ones, allowed the manufacture of a comprehensive range of FGM thicknesses extending from less than 1 millimetre to several centimetres. The authors have developed a new method of FGM manufacture, for which a US patent has been granted ^^. Using this method, it has became possible to produce FGMs across the whole range of thickness from a few millimetres up to several tens of centimetres. The authors have manufactured functionally graded materials consisting of iron ( m ) oxide-kaolin and copper-kaolin, using a successive layer accumulation method. The present research aims to develop a type of element permitting electric field relaxation, consisting of constituent materials graded for electric permittivity. The constituents used are titamium oxide and kaolin. 15 combinations of titamium oxide and kaolin in differing relative proportions were produced by a vacuum filter prcess, and tests were conducted to measure the permittivity and conductivity of each. A simple field distribution calculation was then performed for the assumed case of an element composed of these 15 combinations in graded sequence.

3. METHODS OF MANUFACTURING FGMs In order to manufacture this kind of FGM, Korean kaolin of uniform granular diameter is mixed with titanium oxide and dissolved in distilled water, agitating well. The mixture is put into a cylinder and then vacuum filtered. The filtering (extraction) rate is 120 1/min. When the first cake has been formed, fresh materials are put into the cylinder and thus successive layers are added. After the final cake is formed, the whole FGM mass is subjected to applied pressure for 24 hour. For a 60 mm diameter cylinder, the pressure applied is 3. 6 kg/cm ^. The FGM is then dried naturally, and baked to firmness in a reducing furnace. The furnace temperature is regulated in accordance with JIS R8101 1959^^. The sintering temperature itself is determined by the use of a Seger cone and a test piece. The titanium oxide used for the present experiment was first grade experimental TiO 2. The Korean kaolin used was a clay primarily consisting of kaolin ore, having the chemical formula SiO 2 -Al 2 0 3.

4. THE EQUIPMENT USED FOR THE MANUFACTURE OF FGMs The apparatus was constructed in bronze and comprised four parts, an upper and a lower cylinder, a piston and two perforated plates. The dimensions of the upper cylinder were

375 oxide-kaolin FGM after sintering. The layered manufacture of the FGM is recognisable from the photograph. A scanning electron microscope photograph of the same material is shown as Photograph 2. The graded condition of the titanium oxide and kaolin particles in successive layers can be verified in this photograph.

Photograph 1 The titanium oxide-kaolin FGM

:«?*#'i-«; .,^-- .••J -S.'H

.i^

Photograph 2 Scanning electron microscope photograph of the titanium oxide-kaolin FGM

s J

Fig. 3 Electrodes for measurement of relative permittivity.

6. MEASURED PERMITTIVITY OF FGMs In view of the possible use of FGM elements for electric field relaxation, the permittivity

376 60 X 130 X 95 mm. The lower cylinder was 40 mm in length, with an outlet port for the extraction of filtered water. The piston was 60 mm in diameter with a port in its upper portion for the extraction of air. The two perforated plates were 5 mm in thickness. The cylinder plate had a diameter of 52 mm, the piston plate a diameter of 66 mm. A schematic diagram of this apparatus is seen in Fig. 1. Fig. 2 shows the pressing operation after the successive accretion of layers.

cuum ump

Fig. 1 Apparatus for the manufacture of FGMs

5. METHODS OF SINTERING THE FGMs For sintering, the FGM was heated in a furnace using butane fuel. Following this, a Seger cone was placed inside the furnace to determine the desired temperature, the gas pressure was set at 0. 05 kg/cm ^ and the furnace was ignited. The gas pressure was raised gradually in steps of 0. 05 kg/cm ^, and at 30 minute intervals measurements of the furnace temperature were taken. C o m p re s s ion 3. 6 K g f / c m ^ Layer 1

Natural

Si n t e r e d Temp.lSOOC

dry

f rWYYvWYVJ

UMMMM

m

mMMMM

--

mMm^

Fig. 2 Manufacturing operation Two sintering methods could then be used. In the reducing method, the flow of air into the furnace was cut off after the furnace temperature reached 900 °C. The sintering temperature was determined by observing the melting down of the Seger cone. After the sintering was completed, the furnace was left to return very gradually to normal temperature while a plentiftil air circulation was assured. Photograph 1 shows the titanium

377 e r of the materials was measured using a Q-meter. A schematic representation of the electrodes used is to be seen in Fig. 3. The measurement frequency adopted was 50 MHz, The FGM sample to be tested was placed between the electrodes, and measurements were taken of the distance between the electrodes (L) and the electrostatic capacity Cm. The sample was then removed and the electrostatic capacity Co was measured for the same value of L without the test piece. The relative permittivity was obtained from a comparison of Cm and Co. The results of this comparison are shown in Fig. 4 for the various types. It can be seen that the relative permittivity changes from about 2 on the kaolin side to about 5 on the titanium oxide side. O CO

6

>

o

i

4 0

\— (U

>

^ ^ ^

2

05

1

1

Cd

2

4 6 8 10 1 2 1 4 Samp 1 e No.

Fig. 4 Relation between relative permittivity and titanium oxide-kaolin ratio in samples

7. CALCULATION OF THE FIELD DISTRIBUTION For calculation purposes, the field distribution was taken for the case of a uniformly composed (non-FGM) insulating material. The model used for the calculation is shown in

Potential Fig. 5

OV

Simulation model for numerical analysis

378 Fig. 5. The equipotential line was found by means of a conformal mapping procedure, using the following formula. 4C(1-C) 3C-1-—^ ~ b _ , 2 C + C - i b ^ _^ C-C -] + —-ja Z=—[cos — —sm C+i 2 7C a C+i where C=a/(a ^ +2b ^ ) and a=2b. For numerical analysis, the Y axis was divided into 18 segments, and the X axis into 12. The results are shown in Fig. 6.

^

f9900 V f7500 V 5000 V

Potent iai Fig. 6 Results of electric field calculation The symbols O, @, O and • represent the equivalent potentials for 2500 V, 5000 V, 7500 V and 9900 V respectively. It is to be anticipated that with the use of an FGM composed of materials of differing relative permittivity, the same kind of difference between relaxation and concentration areas will be found within a single piece of material. If an FGM of especially high relative permittivity is used, a greater relaxation effect ought therefore to be attainable. In future experiments, the authors plan to manufacture FGMs of higher relative permittivity, and to perform detailed measurements in order to verify this hypothesis.

8. REFERENCES 1. The Functionally Graded Material Forum and The Society of Non-Traditional Technology : Functionally Graded Materials, p. 351, Kogyochosakai, Tokyo, 1993. 2. U.S. patent. No. 5167813(1992). 3. Japan Industrial Standard: JIS-R-8101(1959). 4. Prinz(Masuda, Kouno, trans.) : "The Calculation of Electric Fields" pp. 37-168, Asakurashoten, Tokyo, 1974.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

379

Evaluation and Modelling of the Residual Stresses Generated on Functionally Graded Materials -Two examplesN. Cherradil, D. Delfosse^ and P. Moecklil iSwiss Federal Institute of Technology, CH-1015 Lausanne, ^Swiss Federal Laboratories for Materials Testing and Research, EMPA, CH-3602 Thun, Switzeriand

1. ABSTRACT A parametric study was carried out to determine the influence of the compositional gradient on the residual stress distribution. The calculations were based on a cylindrical geometry for WCyCo samples and a rectangular geometry for CuNi samples with stepwise compositional variation at the interlayers. The effects of the gradation size and composition profile of the graded materials were investigated by a visco-elasto-plastic finite element analysis using the ABAQUS code. It was found that the degree of residual stress is mostiy determined by the compositional distribution and its thickness, but not by the thermal history. The calculated stress values were compared with those measured experimentally either by Xray diffraction on graded WC/Co specimens or by deflection measurements during electrochemical removal of subsequent layers for graded CuNi samples. The comparison with experimental methods showed good agreement, thus validating the results obtained by the parametric finite element study.

2. INTRODUCTION A functionally graded materials (FGM) is an engineered composite which is designed to optimize materials properties for use under complex loading conditions by local control of composition and microstructure. The intentionally introduced constitutional gradation can be tailored for specific requirements. However, the different material combinations will generally have dissimilar thermal expansion coefficients that can lead to the generation of significant residual stresses, whenever the part is exposed to a thermal cycle (e.g. during processing). Residual stresses are commonly considered a nuisance or even a potential danger to the integrity of the part. In certain cases, however, they may have a beneficial effect on the performance of a component. The FGM concept offers an altemative method to design a part with a well defined, built-in stress state. By judiciously tailoring the composition and the microstructure, thermal stresses can be either dispersed or minimized during both the processing cycle and the in-service use. The purpose of this research work is to develop a fundamental understanding of the effect of a graded structure on stress distribution and thus an efficient tool for optimized profile design.

380 The theoretical results were compared with those obtained from the experimental techniques using two different FGM systems, namely Cu-Ni and WC-Co, as examples. 3. STUDIED MATERIALS The materials used in this study were fabricated following two different techniques (figure 1). The WC/Co graded samples were fabricated by a stepwise compositional control It consists on layering the mixed powders with different composition ratios in a die. Then, they were

Stepwise compositional

Continuous

control

compositional control

i

I

Stepwise stacking

Dosage system

Pre-compaction

Preform

I

1

1

Pressureless sintering

1

1

CIP/Compaction

Hot pressing

compacted uniaxially and sintered. Figure 2 shows an optical micrograph of the polished cross section at the interface between two different layers. The samples under consideration are a bi-layer part 75WC/25Co 95WC/5CO and a tri-layer part 75WC/25Co 85WC/15CO - 95WC/5CO.

|

Hotlsostatic Pressing

The Cu/Ni graded samples were fabricated by a continuous compositional control based on the centrifugal method described elsewhere [1]. An optical micrograph of the microstructure shown in the figure 3 exhibits a smooth variation of the composition.

Figure 1. FGM production by P/M routes

Figure 2. Bi-layer sample 75WC/25Co-95WC/5Co

Figure 3. Cu/Ni graded sample

4. FINITE ELEMENT ANALYSIS For the finite element analysis, the following assumptions were made: - The geometrical details of the bodies: For the WC/Co it was a cylinder with 10 mm diameter and 6 mm height, and for the CuNi, it was a rectangular plate with 3 mm thickness, 14 mm

381 width and 32 mm length. The mesh configuration was modeled by a finite element strip of isoparametric eight-node elements with 4 integration points. We have considered a fixed geometry for each sample and changed either the composition of the gradation and its size. The model is divided into a number of geometric elements which are in contact with one another and considered as layers, each layer being assigned slightly different materials properties. For the 1-phase-system (Cu-Ni), the composition of the Cu-Ni-alloy changes by 10% from layer to layer, whereas for the graded 2-phase-systems (WC-Co), the step-wise nature of the gradient is reproduced in the analysis. - We used an elasto-plastic analysis for WC/Co and visco-elasto-plastic analysis for Cu/Ni samples. Then, the physical and mechanical properties as function of the temperature and composition were compiled from the literature or determined by mechanical tests. - The boundary conditions were imposed following the symmetry taken. - The temperature was considered to be uniform over the whole sample at each calculation step, and the sample was considered to be stress free at the starting temperature. It was set up at 800°C. In fact, above this temperature, Cu, Ni as well as Co are too soft to build up a load and residual stresses disappear by local creep/stress relaxation. - Numerical solutions are obtained using the ABAQUS code [2]. 4.1. WC/Co analysis - The axi-symmetric specimen geometry allowed two-dimensional models to be employed. In order to illustrate the stress generated in a WC/Co part,figure4 shows contour plots of the radial and shear stress distribution in a three layer sample based on 95, 85 and 75 wt.% of WC. The highest concentration of the radial stresses are between the 95 and 85% layers in comparison with the interface between the 75 and 85% layers. This is due to the plastic deformation which could occurs in the Corichregion, whereas in the 95% layer, the relaxation mechanism is effectively prohibited due to the high WC content. This layer is essentially elastic even at high temperatures. Also, since this composition exhibits high elastic moduli, a small displacement in these areas generates a large stress. Regarding shear stresses in the left figure, we can notice that the peaks values are mostly located at the free surface between the 95 and 85 layers.

Figure 4. Contour plots showing the radial and shear stress distribution

382 Figure 5 shows the peak values of the axial andradialresidual stresses for different types of profiles. As expected, the peak stresses depend strongly on the gradient profile, but they are always larger in the non-graded material as a result of the large property mismatch.

Profile A

Profiles

Thickness

^

Profile A

Profile B

Profile D

Thickness

Thickness

Profile C

Profile D

Figure 5. Predicted peak values of radial and axial residual stresses within different WC/Co FGM parts. Figure 6 summarizes the numerical results of the axial stresses for different types of profiles. As we can see, the evolution of the stresses are higher at the edge interface between 95 and the other compositions, and that is true for all the profiles. However, the evolution of the stress through a graded sample from 90_85_80_75 is lower. In fact, if the upper limit of the WC content is lower then 95% (e.g. 90%), then the residual stresses are reduced. This is due to the relaxation mechanism which is more efficient. 800t

^9f'5fi[

^75

i

rhrh ^ ^75

4.

u I I t I I I I I I I II

I I I I I I I I

2 3 4 Thickness (mm)

5

^75

D=^

Thickness

Thickness

^.9:

E^

•90!

i r h ^ ^751 I M l ^ Thickness

Figure 6. Influence of the profile shape on the axial stress distribution

Thickness

383 4.2. Cu/Ni analysis For this analysis, symmetrical parts have been considered. Figure 7 shows the contour plot of Oy stress. As expected the Ni-rich layers at the surfaces are under compression and the Copper -rich layers in the central part under tension. The peak values varies between 100 and -100 MPa. From the contour plot of the plastic deformation shown in figure 10, we can see that all of them occur in the pure metal layers in both Ni and Cu. This is explained by the low yield strength of pure Cu and Ni which are soft metals as compared to the solid solutions hardened CuNi alloys. 4 = -100MPa 10 = 0MPa 14 = 100 MPa

epl

2 = -0.07% 4 = -0.01 % 6 = +0.05%

3 = -0.04 % 5 = +0.02% 7 = +0.08%

Figure 7. 2D analysis for a graded Ni/Cu/Ni sample (lower right quarter of the sample is shown) Figure 8 shows the calculated stress distribution in NiCuNi part. The peculiar form of the curve is due to the solid solution strengthening effect that occurs while changing from a soft, pure metal to an alloy. In the pure metal layers, the thermal stresses are relaxed by visco-plastic deformation. In the adjacent layers, the yield stresses of the alloys (for example Cu-20Ni and Ni-20Cu) are higher and thus are the residual stresses.

-2001

0.00

0.50

1.00

1.50

2.00

2.50

3.00

Position through thickness (mm) Figure 8. Residual stress a n in graded Ni/Cu/Ni

384 5. RESIDUAL STRESS MEASUREMENT 5.1. WC/Co sample The chosen technique was X-ray diffraction that is widely used for non-destructive surface measurement of applied and residual stresses. Stress analysis relies on the determination of the lattice strain using the interplanar spacing as a gauge by measuring the peak shifts in a fixed O direction for different \j/-tilt of the sample [3]. Stresses are calculated from measured strains using diffraction elastic constants which were calculated theoretically. As the Co phase takes up a certain amount of W and transforms after cooling into a solid solution with a variable W content, the measurements were limited to the WC phase. The X-ray measurement is confined to the surface of the specimen. Therefore, in order to investigate the residual stresses trough the thickness, we applied a method that measures the radial stresses at the outer surface of a specimen and tried to correlate them to the intemal stress state [4]. It is known that for a materials with two or more phases the stress field is the superposition of stresses at two levels : Macroscopic stresses which exist between the different layers and resultfi*omthe intemal force balance through the whole material. Microscopic stresses which appear between grains or phases in the material. Thus, the micro residual stresses stemming from the two-phase system have to be added to the results from finite element analysis (where only macro residual stresses are determined) allowing direct comparison with the total stresses experimentally measured. Figure 9 shows the macro residual surface stresses from the numerical analysis for the two and three layer specimens. One can see that the resuhs from Xray measurements agree fairly well with the predicted values. 1000 1

Two-layer-spec imen

^500

K^ 1 1-500 -1000J 0

,

1 2

,

,.

\

1

3 4 5 Thickness (mm)

1

6

Macro stresses X-ray measurements

-1000

2 3 4 Thickness (mm)

"• Total stresses for WC phase V X O Samples

Figure 9. Comparison between the predicted and the measured total residual stress within the WC phase of a two and three layer WC/Co part. X-ray measurements of surface stresses offer a reasonable tool to check the validity of numerical predictions. The method is, however, not fine enough to pick up all the stress

385 maximas and minimas that occur over small distances. It is also not possible to extrapolate from measurements of residual surface stresses to the stress state in the interior of the part. For a non-destructive, more in-depth determination of the residual stress state, neutron diffraction method with its much higher penetration depth has to be employed [5]. 5.2. Cu/Ni sample Copper and nickel exhibit a simple one-phase diagram characterized by complete solid solubility. Thus, it was not possible to apply the X-ray diffraction method. The residual stresses were determined by continuously removing thin layers of materials by an electrochemical technique while monitoring the resulting deformation of the sample [6]. The experimental device consists mainly of a clamp in which a copper cathode, the specimen and a linear voltage displacement transducer (LVDT) are fixed. The amount of bending of the specimen is measured with the LVDT and therefrom, the original stress distribution calculated. Figure 10 shows the measured deflection of some investigated gradient samples. As we can see, the deflection sign changes following the material that is removed at the surface. IDU

NCN-part

£^100 ^

50

X/i

S 0 "I -50

—•—FE-Analysis Experim. 1 Experim. 2 Experim. 3

1-100 0.1

0.2

0.3

0.4

0.5

Removed thickness (thickness of original specimen = 1) Figure 10 Measured deflection of three different Cu/Ni parts as function of the thickness of the removed layer

-150 0.00

0.50

1.00

1.50

2.00

\L • \>s*^^ ^SST *^ 2.50

3.00

Position through thickness (mm) Figure 11. Comparison between the predicted and measured residual stress

In figure 11, the residual stress distribution in the "NCN" part obtained from the analysis is plotted against the experimental results from the electrochemical thinning method. The agreement between the results is highly satisfactory and validates the visco-elasto-plastic approach taken in the FE-analysis as well as the values of the input data. 6. CONCLUSION Residual stresses are often impossible to avoid as a result of manufacturing operations. The graded concept shows, however, that a desired stress state can be designed allowing the dispersion or even optimization of these stresses. Furthermore, the finite element analysis tums

386 out to be a useful tool for mapping residual stresses in the bulk and at the surface of the components non-destructively, providing information which can be used for manufacturing process optimization, analysis of structural integrity, improving mechanical behavior and for service life prediction. Residual micro stresses have to be taken into account for the comparison with experimental results, if two or more phases are present within the FGM. ACKNOWLEDGMENT I would like to express my sincere gratitude and thanks to Professor B.Ilschner who introduced the FGM concept and initiated the work in Switzerland. He gave me the opportunity to create my own group and work in such an exciting field. Also, my thanks go to all my colleagues and collaborators for their help and assistance, the Swiss National Fund and the Swiss Priority Program on Materials Research for their financial support.

REFERENCES [1] Delfosse D. and Ilschner B., "Pulvermetallurgische Herstellung von Gradientenwerkstoffen", Materialwissenschaft und Werkstofftechnik 23,1992,235-240. [2] Hibbitt, Karlsson and Sorensen, Computer Code ABAQUS, Inc., Providence, RI, 1994 [3] Noyan I.C. and Cohen J.B, "Residual Stress, Measurement by Diffraction and Interpretation", Material Research and Engineering Series, Springer-Verlag, New York, 1987 [4] Delfosse D., Cherradi N. and Ilschner B., "Numerical and Experimental Determination of Residual Stresses in Graded Materials", Comp. Eng., Special issue, 1997 [5] Williamson R.L., Rabin B.H. and Byerly G., "Residual Stresses in Joined CeramicMetal Structures: FEM Studies on Interlayer and Creep Effects", in 3rd International Symposium on Structural and Functional Gradient Materials, ed. Ilschner, B. and Cherradi, N., PPUR, Lausanne, (1994), 215-221. [6] Delfosse D., Kiinzi H.-U. and Ilschner B., "Experimental Determination of Residual Stresses in Materials with a One-Dimensional Gradient of Composition", Acta Metallurgica et Materialia 40, (1992), 2219-2224.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

387

Residual Strains and Stresses in an AliOa-Ni Joint Bonded with a Composite Interlayer: FEM Predictions and Experimental Measurements Barry H. Rabin, Richard L. Williamson and Hugh A. Bruck Idaho National Engineering Laboratory, Idaho Falls, ID Xun-Li Wang and Tom R. Watkins Oak Ridge National Laboratory, Oak Ridge, TN David R. Clarke University of California, Santa Barbara, CA Abstract A cylindrical Al203-Ni joint bonded with a 4.0 mm thick composite interlayer of 40 vol.% AI2O3-6O vol.% Ni was fabricated by powder processing and the residual stresses and strains in the specimen were studied experimentally using neutron diffraction, x-ray diffraction and optical fluorescence spectroscopy. Experimental measurements were compared with finite element method (FEM) modeling results obtained using a variety of different constitutive assumptions. The predicted residual strain distribution within the AI2O3 along the center of the specimen was in excellent agreement with neutron diffraction measurements. Alternatively, the predicted peak strains and stresses within the AI2O3 along the specimen surface were significantly higher than those measured, suggesting stress-relief occurred near the free-edge during cooling. The mechanisms of stress-relief are uncertain, however localized plasticity and damage within the composite interlayer are believed to play a role. Introduction Ceramic-metal joining has recieved considerable recent attention [1,2]. Numerous studies have focussed on residual stresses in ceramic-metal joints [3-5]. Residual stresses have also been measured experimentally in a several cases, however these studies involved joints fabricated using brazing techniques [6-11]. The goal of this work was to experimentally measure the residual strains and stresses in a ceramic-metal joint that is more characteristic of typical FGM structures, and to compare the results with modeUng predictions. In this study, a ceramic-metal joint bonded with a thick composite interlayer was produced by powder processing. Strains and stresses were measured using a variety of experimental techniques. Experimental results were compared with FEM predictions from a detailed elastic-plastic, temperature-dependent model. Al203-Ni was selected as a model system because the properties of the pure materials are well characterized, there is a large mismatch in properties, and Al203-Ni composites are amenable to diffraction methods of residual strain measurement. Rather than examining a complicated, multi-layered FGM structure, this model study was conducted using a cylindrical Al203-Ni joint containing a single, homogeneous composite interlayer approximately 4.0 mm thick. Experimental Specimen fabrication involved powder processing techniques described in detail elsewhere [12, 13]. The steps included powder selection, incorporation of sintering aids and binder, powder mixing, die compaction, sintering, and final consolidation by hot isostatic pressing (HIP). The cylindrical specimen examined in this study is shown in Figure 1, along with photomicrographs showing the microstructure of the composite in the regions near the ceramic-composite and metal-composite interfaces. The 40 vol.% AI2O3-6O vol.% Ni composite interlayer was approximately 4.0 mm thick.

388 The neutron diffraction measurements were conducted at the High Flux Isotope Reactor of Oak Ridge National Laboratory using a triple-axis spectrometer operated in the diffractometer mode. The details of the experimental method have been described elsewhere [10,14]. Slits of dimensions 0.8 X 4 mm^ and 0.8 x 30 mm^ were inserted before and after the specimen, which together defined a sampling volume of approximately 2,6 mm^. Radial and axial strain (Crr and e^z) distributions were investigated in this study. For each strain component, a series of measurements was taken along the specimen axis of symmetry through the AI2O3 and composite layers. For the purpose of illustration, the specimen orientation for the measurement of e^z is shown in Figure 2. To improve the ability of detecting the predicted steep strain gradient, an overlapping sampling volume was employed, with a step size of 0.25 mm between analysis points in the vicinity of the interfac. The AI2O3 (3 0 0) reflection (20-72.0°) was used for strain determination according to e=

Figure 1. The powder processed specimen examined in this study, consisting of pure AI2O3 joined to pure Ni using a 40 vol.% AI263-6O vol.% Ni composite interlayer approximately 4.0 mm in thickness. Micrographs of the composite and the interface regions are also shown.

d-do _ sin 60 ^0

sin 0

where d is the lattice spacing and 20 the diffraction angle of AI2O3 (3 0 0). do and 20o are the corresponding stress-free values. In general, strains measured with diffraction methods are a superposition of macro- and microstrains [15]. Since in this study we were solely concerned with the macrostrain distribution resulting from the thermal expansion mismatch between the bonded dissimilar materials, the data points farthest from the interface were assumed to correspond to zero macrostrain. In this way, effects of any microstrains due to the thermal expansion anisotropy within the single phase AI2O3 were removed. However, data in the composite layer still contain a microstrain contribution due to the thermal expansion mismatch between AI2O3 and Ni phases, and this contribution was not taken into account in the analysis. Details of the x-ray diffraction strain measurements can be found elsewhere [14]. Measurements were performed on a 4-axis goniometer* in the "Q-goniometer geometry" [15]. Cu Ka^ radiation (X = 1.54060 A) was used. A pinhole collimator with a 0.5 mm opening was used to define the sampling volume, and soller slits were and to minimize sample displacement errors. The (1 4 6) AI2O3 reflection at -136° 20 was step-scanned using three azimuthal angles ((|) = 0, 45, and 90°) and seven tilt angles, {\\f = 0, ±28.2, ±42, and ±55 °) corresponding to sin^Xj/ values of 0, 0.22, 0.45, and 0.67, respectively. In order to map the strain gradients along the length of the specimen, the sample was translated and the above measurements were repeated at various points along the cylindrical axis of the sample. The penetration depth of Cu Ka^ radiation in a-Al203 depends upon the tilt angle; a depth of * PTS goniometer, Scintag, Inc., Cupertino, CA.

(1)

389 about 35 |im was calculated and can be taken as an average value for the current experiments. The sample was rotated about its cylindrical axis to improve particle statistics for some of the measurements. Careful goniometer and sample alignment procedures were employed. The average strain free interplanar spacing, d^, was determined from annealed powder the same as that used to fabricate the specimen. Strains and stresses were calculated using the procedure of Winholtz and Cohen [16].

Figure 2. Schematic illustrating the specimen orientation for neutron diffraction axial strain measurements, and the overlap of sampling used to improve spatial resolution of the measurements.

Optical Fluorescence Spectroscopy Details of the optical fluorescence technique used to measure residual stresses in polycrystalline AI2O3 have been published previously [17, 18]. A portion of the sample was excited using an argon ion laser and the fluorescence peak of the Cr^^ impurity present in the AI2O3 is detected using a Raman spectrometert. The shift of the Rl and R2 fluorescence lines in Cr^^ doped AI2O3 can be related to the stress state in the excited volume using the relation

Av = (2n,-fn,)(^ii±^|i±^) where Av is the average frequency shift, and Hy are the piezospectroscopic coefficients relating frequency to stress. The numerical values of the piezospectroscopic coefficients Fla and lie have been determined previously by direct experimentation [17, 18], and are 2.7 and 2.15 cm"^ GPa~^, respectively for the R2 line, which was used in this study. For these experiments, a high spatial resolution laser probe was used, having a diameter of approximately 50 |im. The excited volume from which stress information was obtained is determined by the penetration depth of the argon ion laser in polycrystalline AI2O3, which has been measured experimentally to be on the order of 50 fim. FEM Modeling Strains and stresses were computed for the joined specimen cooled uniformly to room temperature from an assumed stress-free elevated temperature using numerical models described in detail previously [19, 20]. The coordinate system and an example of the finite element mesh utilized are shown in Figure 3. Elastic-plastic response was permitted in both the Ni and Al203-Ni composite materials; a von Mises yield condition and isotropic hardening were assumed.

t Model T64000, Instruments SA, Inc.

(6)

390 Calculations were performed for four different cases as described below. In all cases, pure AI2O3 was assumed to remain elastic with a temperature independent Young's modulus of 380 GPa and Poisson's ratio of 0.25. The thermal expansion coefficient of AI2O3 decreased linearly from 9.4 x 10"^ K'^ at 1100 K to 5.4 x 10'^ K"^ at room temperature. For pure Ni, all cases assumed the same temperature dependent Young's modulus, which decreased from 208 GPa at room temperature to 166 GPa at 900 K, and the same Poisson's ratio of 0.31. The thermal expansion coefficient of Ni decreased linearly from 17.8 x 10"^ K"^ at 1100 K to 13.4 x 10"^ K"^ at room temperature. Two different sets of Ni strength properties were examined, corresponding to properties taken from the literature for both fine grained and coarse grained microstructures [21, 22]. These simulations are referred to in Table I as Case 1 and Case 2, respectively. The temperature dependent yield strength and the flow strength at 2% strain are listed in Table I. Linear hardening behavior was assumed in both cases. For Case 1 and Case 2 the properties of the composite material interlayer were computed as follows. The thermal expansion coefficient of the composite was estimated using a simple volume fraction based mixture rule. The temperature dependent composite stress-strain curves were constructed using an modified rule-of mixtures approach first proposed by Tamura, et al. [23]. Also listed in Table I are the properties used in a calculation referred to as Case 3. This calculation was performed using a set of material properties intended to be representative of the actual material behavior. Most of these properties were determined through direct experimental measurements made on bulk Ni and composite specimens, although some of the temperature dependent values that were not measured directly were estimated by extrapolation based on trends observed in literature data, as well as unpublished research. The notes provided along with Table I describe in detail the origin Figure 3. Coordinate system and example of these data. of FEM mesh used to model the specimen. Calculations for Cases 1 through 3 were all performed assuming a stress-free temperature of 1100 K. Recent model calculations performed including creep deformation constitutive laws for an Al203-Ni bimaterial joint indicated that creep strains in the pure Ni exceeded the plastic strains for all temperatures above about 700 K [24]. Since creep deformation in the Ni at high temperatures effectively lowers the stress-free temperature, a final calculation was performed assuming a stress-free temperature of 700 K in an effort to account for this effect. This calculation, referred to as Case 4, was performed using the same material properties as in Case 3.

391 Table 1. Temperature Dependent Mechanical Properties of Ni and the 40% Al2O3-60% Ni Composite used in the FEM Calculations for Three Different Cases. 900K

300K Case

Material

E

E

V

(MPa) 1

2

3 a. b. c. d. e. f. g. h.

(MPa)

1

V

(MPa)

(GPa)

(MPa)

Ni

208.0'

0.31'

148.0'

161.0'

166.0'

0.31'

69.0'

72.0'

interlayer

254.0'

0.29'

148.6'

230.7'

214.8'

0.29'

69.4'

134.0'

Ni

208.0'

0.31'

25.0''

53.7'

166.0'

0.31'

11.9'

25.6'

interlayer

254.0'

0.29'

25.r

133.2'

214.8'

0.29'

12.0'

94.8'

Ni

208.0'

0.31'

18.0^

75.0^

166.0'

0.31'

15.3^

63.8^

interlayer

219.0^

0.29'

80.0^

192.0«

171.0^

0.29'

56.0^

134.4^^ 1

from [221 from [21] linear rule-of-mixtures modified rule-of-mixtures for fine grained Ni [22] and q = -4.5 GPa modified rule-of-mixtures for coarse grained Ni [21] and q = -4.5 GPa temperature dependence assumed to be same as that of fine grained Ni from direct experimental measurement estimate based on unpublished research

Experimental Results The results of neutron diffraction strain measurements made along the axis of symmetry and within the vicinity of the interface between the AI2O3 and the composite are shown in Figure 4 as a function of axial position within the sample. Both the radial strain (EJ and axial strain (EZZ) are shown, along with errors estimated from the standard deviations of the least-squares fitting of the recorded diffraction profiles. The errors in this experiment were dominated by the unfavorable scattering intensity due to the small sampling volume used. Within the experimental precision, the experimental data provide evidence of a steep strain gradient across the interface, extending to a distance of approximately 2 mm on either side of the interface. In general, the magnitudes of the measured strains are quite small, on the order of 10' . In the AI2O3 layer, E^ becomes increasingly compressive as the interface is approached. A maximum compressive strain of approximately 3 x 1 0 ' was measured within the AI2O3, approximately 1 mm from the interface. The axial strain (E^^) is compressive in the AI2O3 layer, becoming tensile only when the interface is approached. Measurements of £22 across the interface were not attempted because in this measurement geometry, an artificial peak shift was anticipated when the sampling volume was partially buried in the AI2O3 layer [25]. This artifact leads to an apparent strain and adds ambiguity to the determination of e^z. The results of x-ray diffraction strain measurements made along the specimen surface near the interface between the AI2O3 and the composite are shown in Figure 5 as a function of axial position along the sample. In this figure, the axial strain and hoop strain components are shown. The peak axial strain value was quite small, approximately 2 x lO"^ tensile within the AI2O3 and very little variation with distance from the interface was observed. The hoop strain values exhibited considerable scatter and larger errors, but the peak value appears to be compressive within the AI2O3 and occurs within 1 mm from the interface.

392 The results of stress measurements made along the specimen surface using optical fluorescence spectroscopy are shown in Figure 6 as a function of axial position along the sample surface. In this figure, the stress reported is the sum of the three principal stresses. As a result of the high spatial resolution (-50 ^im diameter spot size) and small errors (estimated to be ±20 MPa), the entire stress distribution within the AI2O3 as the interface is approached is well characterized. The stress increases smoothly and gradually as the interface is approached. The maximum stress measured was on the order of 100 MPa and occurred at a location approximately 1 mm from the interface. —

1

—

1

—

1

—

(

—

Modeling Results FEM modeling results showing the predicted strain and stress distributions within the pure AI2O3 are J 11 '1 li r ' also shown in Figures 4-6 for each of the different cases (i.e. different constitutive assumptions) examined. Significant differences in the predicted peak strain and [ 'NI^T stress values were observed depending upon the '\ J constitutive assumptions. In all cases, the magnitude of the predicted peak strain/stress values decreased in the order Case l>Case 3>Case 4>Case 2. Case 1 resulted in the highest predicted strain and stress values whereas Case 2 resulted in the lowest values. These results can be explained by recognizing Figure 4. Neutron diffraction results showing the distribution of that in these analyses the thermal strains and stresses in the AI2O3 were dictated by the plastic flow properties of the volume averaged (a) radial the metal and composite. Case 1 used properties taken strain and (b) axial strain as a from the literature for fine grained Ni (approximately 50 function of axial position along ^im grain size) resulting in a relatively high value for the the interior of the specimen. assumed yield strength, and since the composite properties were calculated using the modified rule-of-mixtures, the composite also exhibited a relatively high yield strength. In contrast, Case 2 used Ni properties characteristic of a coarse grained Ni microstructure (greater than approximately 1 mm grain size), resulting in very low estimates of the flow strength for both the pure metal and the composite. It is clear that the flow strength of the metal and composite play an important role in determining the magnitude of the residual stresses in the joint. The flow strength of Ni is determined primarily by grain size. In contrast, the flow properties of a composite material are generally influenced by additional microstructural aspects, such as amount and size of second phase, its spatial arrangement, interfacial characteristics, etc. These effects are considerably more difficult to assess, and this information is not readily included in simple, empirical models such as the rule-of-mixtures. portion of the specimen experienced significant grain growth during elevated temperature processing and had a final grain size of approximately 1 mm. Its properties were thus expected to be similar to those assumed in Case 2. In contrast, within the composite interlayer the presence of the AI2O3 particles limited the extent of grain growth and the final grain size was on the order of the interparticle spacing, approximately 35 |Lim. Therefore, when estimating the composite properties using ^ e rule-of-mixtures approach it would make sense to utilize the fine grained pure Ni properties assumed in Case 1. Although this calculation was not performed, it would be expected that the results would fall between those of Case 1 and Case 2. The point is that use of a single set of Ni flow properties for estimating the 1

1

1

Cm 1

•

1 . 1 1

—, / IH / 1

nnriron dlffraairon

flu

: 1'" AI2O3

1

1

^

60Ni-40Al203

D l i l o n e i from Intarfact ( m m )

393 behavior of both the pure metal and the composite would not be expected to result in accurate residual stress predictions. Note that near the r Microstmctural examination of the powder processed materials indicated the pure Ni adial freesurface (Figures 5 and 6) the differences in the peak stresses within the AI2O3 between the Case 1 and Case 2 predictions were significantly greater than within the specimen interior (Figure 4). This result is due to the fact that the plastic strains are predicted to be considerably larger at this location, and differences in the assumed flow characteristics of the composite therefore become more obvious. Case 3 was a calculation performed using properties actually measured for the bulk Ni and composite materials. As expected, the predicted strain and stress values fell between the Case 1 and Case 2 results. Case 4 used the same properties as in Case 3, except the assumed stress-free temperature was lowered from 1100 K to 700 K. Reducing the simulated temperature change during cooling resulted in a concomitant reduction in the peak strain and stress predictions. This result clearly indicates the importance of understanding and accounting for possible stress-relief mechanisms operating in any of the materials within the joint during cooling from elevated temperatures.

52 /«m from rodlol frw ijrfoc« 1 ^ ^ — Can 1 .

—

'1

—

•

f

Cait 2

y /

xray (Xffriiotisn

___-iiresf5i^'

. \\, . . . 1

/N\I

' %

^"1\

iK

-1

1 r

I

AI^Oj

i

1 Composite

Dlitonct from InUrfae* (mm)

52 /im trom radial frat surfaiiB

Di*tone« from Inltrioc* (mi

Figure 5. X-Ray diffraction results showing the distribution of (a) axial strain and (b) hoop strain as a function of axial position along the surface of the specimen, at a fixed depth of 52 fim.

Discussion Figure 4 compares the strains measured using neutron diffraction with the FEM predictions. Due to the large sampling volume used in the neutron diffraction experiments it was necessary to manipulate the FEM results in order to directly compare the measured strains with those predicted by the model. This was accomplished by calculating an average value of strain from the elements that would have contributed to the diffracted intensity. The predicted strains shown in Figure 5 are the volume averaged results. Note that, within the experimental errors, both the magnitudes of the experimentally measured maximum strains, as well as the shape of the strain distribution, are in reasonable agreement with the FEM predictions for all of the numerical cases examined. Overall, these results suggest that the strain distribution within the interior of the specimen is relatively insensitive to differences in material properties, and therefore reasonably accurate modeling predictions can be expected using relatively simple constitutive assumptions. This can be explained by the fact that limited plasticity occurs in this region of the specimen, therefore the strain distribution is governed primarily by the thermal expansion coefficients and the elastic modulii of the materials, which are relatively independent of microstructure (i.e. to a first approximation they depend only on volume fraction). Of course, this conclusion is only valid for this particular specimen geometry since the fonnation and spreading of plastic zones during cooling is highly geometry dependent.

394 Due to the larger plastic strains predicted locally close to the interface near the surface of the specimen the situation is significantly different. Consequently, less satisfactory agreement between the measured and predicted strains and stresses was observed. Figure 5 compares the strains measured in the AI2O3 along the surface of the specimen using x-ray diffraction with the FEM results. The FEM results shown correspond to the strains predicted within a row of elements at a fixed distance of 52 |im from the radial free surface. This depth Composite compares reasonably well with the calculated - 4 - 3 - 2 - 1 0 penetration depth of Cu Ka^ radiation in Oistanco from interfoce (mm) AI2O3, approximately 35 |im. The FEM predictions of the axial strain appear to overFigure 6. Optical fluorescence spectroscopy predict the strain values measured using x-ray results showing the sum of the principal diffraction for all numerical cases studied. stresses as a function of axial position along The stresses measured in the AI2O3 the surface of the specimen, at a fixed depth along the surface of the specimen by optical of 52 |im. fluorescence spectroscopy are shown in Figure 6, along with the FEM predictions. These FEM results are also presented for a depth of 52 fim from the radial free surface, comparable to the penetration depth of the laser in AI2O3, which has been measured to be --50 jim. Note that the shape of the stress distribution is well predicted by the FEM model, and the measured distribution is in reasonable agreement with the FEM predictions for Case 2 or Case 3. However, the peak value of stress is overestimated by the model for all cases, and the location of the maximum stress is predicted to be much closer to the interface between the AI2O3 and the composite. The calculations predict a significantly steeper strain gradient immediately adjacent to the interface than was observed experimentally. These results suggest that a localized stressrelief mechanism operates in this region of the specimen in response to the large concentration of strains and stresses (including significant plasticity) near the intersection of the interface with the free-edge. There are several possible strain/stress relief mechanisms that could be considered to explain these results. For example, microcracking in the AI2O3, partial debonding at the interface, or damage accumulation within the composite, each could result in lower peak strain and stress values in the AI2O3. No microstructural evidence for any of these mechanisms has yet been found within as-fabricated joints. However, mechanical property studies carried out on bulk 40 vol.% AI2O3-6O vol.% Ni composites has demonstrated that damage does accumulate during large-scale plastic deformation, and that models incorporating microstructural damage (in the form of fractured particles, separation of contiguous AI2O3 particles, or particle-matrix interfacial decohesion) can adequately explain the observed deformation and fracture behavior of this material [26]. It is difficult to quantitatively assess the effects of any potential stress-relief mechanisms using the continuum models presented here, since the edge-stress concentration and its effects occur on a size scale comparable to the microstructure. Additional experimental work is needed to characterize the evolution of the residual strain and stress state in this area of the specimen during cooling, and new modeling methodologies incorporating detailed microstructural information will be required to allow local material response to be predicted. 52 /im from rodiql free surface

395 Summary Residual strains and stresses were studied experimentally using a variety of techniques in a model ceramic-metal joint containing a thick, homogeneous composite interlayer. An elastic-plastic FEM model was used to investigate the residual strains and stresses in the joint, and to investigate the role of the assumed interlayer properties on the predicted strains and stresses. The elastic strains measured in the specimen interior using neutron diffraction were found to be in excellent agreement with model predictions made using a simple, modified rule-of-mixtures approach to estimate the joint interlayer properties. It is therefore demonstrated that, excluding edge-effects, FEM models can be used to reliably predict strain and stress distributions within composite and graded interlayer joints, provided reasonably accurate material property estimates are available. Using two independent measurement techniques (x-ray diffraction and optical fluorescence spectroscopy) it was shown that the strains and stresses in the Al2C)3 measured along the specimen surface are lower than what was predicted by the FEM model, suggesting that stress-relief occurs during cooling from the joint fabrication temperature. No experimental evidence for the occurrence of any stress-relief mechanisms has been found, however, it is believed that large, localized plastic strains may have induced damage within the composite interlayer. Acknowledgments Research sponsored in part by the U. S. Department of Energy, Office of Energy Research, Office of Basic Energy Sciences, under DOE Idaho Operations Contract DE-AC0794ID13223, and in part by the Assistant Secretary for Energy Efficiency and Renewable Energy, Office of Transportation Technologies under DOE Oak Ridge Operations Contract DE-AC05-96OR22464. References A. H. Carim, D. S. Schwartz and R. S. Silberglitt (eds.). Joining and Adhesion of Advanced Inorganic Materials, Mater. Res. Soc. Symp. Proc, Vol. 314. Materials Research Society, Pittsburgh, PA, 1993. A. J. Moorhead, R. E. Loehman and S. M. Johnson, Structural Ceramics Joining II, Ceramic Trans., Vol. 35. The Amercian Ceramic Society, Westerville, OH, 1993. T. Suga, K. Mizuno and K. Miyazawa, "Thermal Stresses in Ceramic-to-Metal Joints"; pp. 137-142 in Metal-Ceramic Joints, Proc. MRS International Meeting on Advanced Materials, Vol. 8. Edited by M. Doyama, S. Somiya and R. P. H. Chang. Materials Research Society, Pittsburgh, PA, 1989. D. Munz, M. A. Sckuhr and Y. Yang, "Thermal Stresses in Ceramic-Metal Joints with an Interlayer," /. Amer. Ceram. Soc., 78 [2] 285-290 (1995). H.-Y. Yu, S. C. Sanday and B. B. Rath, "Residual Stresses in Ceramic-Interlayer-Metal Joints," /. Amer. Ceram. Soc., 76 [7] 1661-1664 (1993). 1. M. Kurita, M. Sato, I. Ihara and A. Saito, "Residual Stress Distribution of CeramicMetal Joint"; pp. 353-362 in Advances in X-Ray Analysis, Vol. 33. Edited by C. S. B. e. al. Plenum Press, New York, 1990. S. Tanaka and Y. Takahashi, "Effects of X-ray Beam Collimation on the Measurement of Residual Stress Distribution in a Si3N4/Steel Joint," ISIJInternational, 30 [12] 10861091 (1990). O. T. lancu, D. Munz, B. Eignemann, B. Scholtes and E. Macherauch, "Residual Stress State of Brazed Ceramic/Metal Compounds, Determined by Analytical Methods and Xray Residual Stress Measurments," /. Amer. Ceram. Soc., 73 [5] 1144-1149 (1990). L. Pintschovius, N. Pyka, R. Kubmaul, D. Munz, B. Eigenmann and B. Scholtes, "Experimental and Theoretical Investigation of the Residual Stress Distribution in Brazed Ceramic-Steel Components," Mater. Sci. Eng., A177 55-61 (1994).

396 10. X.-L. Wang, C. R. Hubbard, S. Spooner, S. A. David, B. H. Rabin and R. L. Williamson, "Mapping of the Residual Stress Distribution in a Brazed Zirconia-Iron Joint," Mat ScL Eng., A211 45-53 (1996). 11. H. Li, L. Z. Sun, J. B. Li and Z. G. Wang, "X-ray Stress Measurement and FEM Analysis of Residual Stress Distribution Near Interface in Bonded Ceramic/Metal Compounds," Scripta Materialia, 34 [9] 1503-1508 (1996). 12. B. H. Rabin, R. L. Williamson, R. J. Heaps and A. W. Erickson, "Powder Processing of Nickel-Aluminum Oxide Gradient Materials"; pp. in press in PM' 92, Proceedings of the 1992 Powder Metallurgy World Congress. Edited by Metal Powder Industries Federation, Princeton, NJ, 1992. 13. B. H. Rabin and R. L. Williamson, "Design and Fabrication of Ceramic-Metal Gradient Materials"; pp. 145-154 in Processing and Fabrication of Advanced Materials III. Edited by V. A. Ravi, T. S. Srivatsan and J. J. Moore. The Minerals, Metals and Materials Society, Warrendale, PA, 1994. 14. B. H. Rabin, R. L. Williamson, H. A. Bruck,. X.-L. Wang, T. R. Watkins and D. R. Clarke, "Residual Strains and Stresses in an Al203-Ni Joint Bonded with a Composite Interlayer: FEM Predictions and Experimental Measurements," to be published in /. Amer. Ceram. Soc, 1996. 15.1. C. Noyan and J. B. Cohen, Residual Stress, Measurement by Diffraction and Interpretation, a) p. 101-2, b) p. 118, Springer-Verlag, New York, 1987. 16. R. A. Winholtz and J. B. Cohen, "Generalized Least-squares Determination of Triaxial Stress States by X-Rays Diffraction and the Associated Errors," Aw^t. /. Phys., 41 189-99 (1988). 17. Q. Ma and D. R. Clarke, "Stress Measurement in Single-Crystal and Polycrystalline Ceramics Using Their Optical Flourescence," / Amer. Ceram. Soc., 16 [6] 1433-1440 (1993). 18. Q. Ma and D. R. Clarke, "Piezospectroscopic Determination of Residual Stresses in Polycrystalline Alumina," /. Amer. Ceram. Soc., 77 [2] 298-302 (1994). 19. R. L. Williamson, B. H. Rabin and J. T. Drake, "Finite Element Analysis of Thermal Residual Stresses at Graded Ceramic-Metal Interfaces, Part I: Model Description and Geometrical Effects," / Appl Phys., 74 [2] 1310-1320 (1993). 20. J. T. Drake, R. L. Williamson and B. H. Rabin, "Finite Element Analysis of Thermal Residual Stresses at Graded Ceramic-Metal Interfaces, Part II: Microstructure Optimization for Residual Stress Reduction," /. Appl. Phys., 74 [2] 1321-1326 (1993). 21. W. Betteridge, Nickel and its Alloys. Ellis Harwood, Ltd., West Sussex, UK, 1984. 22. M. A. Meyers and K. K. Meyers, Mechanical Metallurgy: Principles and Applications, p. 345, Prentice Hall, Inc., Englewood, NJ, 1984. 23.1. Tamura, Y. Tomota and H. Ozawa, "Strength and Ductility of Fe-Ni-C Alloys Composed of Austenite and Martensite with Various Strength"; pp. 611-615 inProc. 3rd Int. Conf. Strength of Metals and Alloys. Edited by Institute of Metal and Iron, London, 1973. 24. R. L. Williamson, B. H. Rabin and G. E. Byerly, "FEM Study of the Effects of Interlayers and Creep in Reducing Residual Stresses and Strains in Ceramic-Metal Joints," Composites Eng., 5 [7] 851-863 (1995). 25. S. Spooner and X.-L. Wang, unpublished research, Oak Ridge National Laboratory, Oak Ridge, TN, 1995. 26. H. A. Bruck and B. H. Rabin, "Deformation and Fracture Modelling of Nickel-Alumina Composites for FGMs," Acta Metall. Mater., submitted for publication (1996).

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

397

Residual Thermal Stresses in Functionally Graded Ti-TiCx Materials. N. Frage^, M.P. Dariel^, U. Admon and A. Raveh Department of Materials Engineering, Ben-Gurion University, Beer-Sheva, Israel, Nuclear Research Center-Negev, P.O.Box 9001, Beer-Sheva, Israel. ABSTRACT Graded Ti-TiC nanoscale multilayers transform after an appropriate interdiffusion anneal into a functionally graded (FG) region between a TiC hard coating and a Ti substrate. Using the available database regarding the properties of the TiC^ solid solution, it was possible to develop analytical expressions for the residual thermal stresses that develop in the FG-material. It was shown that by appropriate design of the graded concentration in the initial multilayer, it is possible to generate compressive stresses in the vicinity of the ceramic-like edge of the FGtransition zone.

1. INTRODUCTION AND BACKGROUND A functionally graded material (FGM) is a composite, consisting of one or more phases, with its composition varying in some spatial direction. The design is intended to take advantage of certain desirable features of each of its constituents. We are focusing our attention on the titanium (metal) - titanium carbide (TiCx) combination. According to the Ti-C phase diagram, the carbide phase has a wide homogeneity range that extends from TiCg 43 to TiCg gg. The extended composition range of the TiCx phase is a key element in the approach we suggest for interface tailoring and which consists of diffusion anneal of a graded Ti-TiC multilayer. The multilayer is produced by sputter-coating on an appropriate Ti substrate of Ti-TiC bilayers with constant thickness. Whereas the thickness of each bilayer is kept constant, the thickness of the individual components in each bilayer, namely that of Ti and of TiC, varies according to the required design, as shown schematically in Fig. la. A hypothetical but plausible outcome of a short diffusion anneal will result in a different configuration of the two phases, as shown in Fig. lb. During the cooling stage after the diffusion anneal, thermal stresses, that may impair the performance of the system, are generated within the thin layer coating. The object of this study was to examine whether an appropriate design of the FGM region may alleviate the effect of the residual thermal stresses. Consider an ordinary substrate-ceramic coating combination. Such a composite system, prepared at elevated temperatures and subsequently cooled to room temperature, will be thermally stressed due to the usually large difference in thermal expansion and elastic moduli of the substrate and coating. These stresses often exceed the fracture strength of the ceramic component, particularly in regions close to free surface near the interface. This leads to either cracking of the ceramic part or to failure at the substrate-coating interface.

398 component, particularly in regions close to free surface near the interface. This leads either to cracking of the ceramic part or to failure at the substrate-coating interface. The residual thermal stresses in a substrate-coating combination are due to the interfacial forces arising from the thermal expansion coefficient mismatch between coating and substrate and from the presence of a lending moment [1]. The former are uniformly distributed over the film thickness while the laUe'" arises from the requirement to balance the external bending moment induced by the interfacial force in a coating-substrate combination and varies across the film thickness [2]. Functionally graded materials by virtue of the gradual change in the thermal expansion mismatch over the transition region offer a solution and can minimize the thermal stresses arising from cooling or heating (a) Multilayer Coating (n layers)

[Substrates, k y y y y y yl

KTiC fcoatingi

(b) Single phase Ti(C) solid solution -• •Two-phase region

Ti Substrate H

Single phase Tie region

[Tie

3

[Coating j

't A rt rt rt rt rt

Fig.l. (a) Schematic view of graded multilayer. The two basic components of;the multilayer are Ti and TiC layers. The basic unit of the multilayer is a juxtaposed double layer of Ti and Tie. This basic unit has a constant L width. The relative width of the two components of this basic unit varies, however, as one proceeds from one end of the multilayer to its other. In the vicinity of the Ti substrate, the Ti component makes up most of the basic unit; on side, this basic unit consists mostly of the TiC component, (b) Schematic drawing of the possible outcome of the diffusion anneal on the microstructure of the functionally gradient region. 2. THE MODEL We have used one-dimensional model [3] to calculate the residual stresses in the system involving a ceramic-like (TiC) coating connected to a metal-hke (Ti) substrate by an inter-mediate PGM region, consisting (ultimately after the interdiffusion anneal) of titanium carbide with a carbon concentration that varies with distance. The PGM plate has a thickness 2c, unit dimension in depth (z-direction) and is infinitely long in the x direction, as shown in Fig. 2. The composition in any xz plane is held constant. The carbon concentration, actually the titanium carbide composition, x in TiC^, varies within the transition region in the y direction according to a given functional form. We have used such a function after Wakashima et al. [4], Eq.(l), where y^ and y2 are border regions of pure phase 1, Ti, and phase 2, TiC, respectively. This function has the ability, depending on the single parameter, N, of being either "concave upward" and "concave downward".

399

c,max

/(y) =

[y2-yi J

(1)

FGM

Knowledge of / (y), and the composition-dependent microstructure, allows to determine the y-dependence of the effective values of the coefficient of Tie thermal expansion and Young's Figure 2. A schematic view of the modulus. These, in turn, can be used to Functionally Graded Material system calculate the stress distribution across the transition layer. In principle, this allows to establish a linkage between FGM design and performance, where/ (y) is related to design, while the calculation of the residual stresses is related to performance. We have adapted this function to a single-phase FGM based on nonstoichiometric titanium carbide TiCj^. It is possible to treat the nonstoichiometric carbide as a solution of titanium carbide of high carbon content (TiCo.98) and added titanium. Note that x^Mc/Mxi where M stands for the number of moles, and JA, the molecular weight is equal to 48, 12 and 59.76 g/mole for Ti, C and TiCo.98, respectively. The value of x can written as :

59.76 AS

xO.98

4704(1-m^.) 48 + 11.76m„.

(2)

59.76

where rrij.. is the titanium wt. fraction in Ti - TiCo.98 mixture (in the pre-diffusion anneal stage, as shown in Fig.la). When mj.= 0.456, x=0.48, and mj.=0, jc=0.98. Assuming that wt. fraction of titanium in mixture varies according to the function (3), we obtain for the dependence of x in the y direction (eq.4), within FGM region: S my.. = 0 4561

^

•^'"°

(3)

1-0.456 [ y-ynnn 1 ^ jc=47. 04|

b

-V-I

I 48 + 5.363 - ^ - ^ \

L -'max

-^i

I ^1

jJ

(4)

The dependence of m xi and x on the distance (>') for a transition region 10 /<m is shown in Fig.3, for several values of N (0.2, 1,5), values that determine both the curvature and the steepness of the dependence. The relevant thermophysical parameters (expansion coefficient and Young modulus) of the TiC phase vary with the carbon content and, therefore, along the yaxis in the FGM region. Using reported data values in the literature [5], one can plot the xdependence of these parameters that can be approximated by the following expressions: a=(9.61-2.09jc)xl0"^^"'

(5)

E=l94 + 293.6x,GPa

(6)

400 and using expression (4) for x, we can deduce the y-dependence of a and of E,

f

147.04

a{y) = 9.61-2.09^

u

iri

y-Vm

Li>«Hil2™2iJi_^ ^ 10^j^-' f y-y^n 1 48+5363 147.0411 - ; - ^ ^"^

£(j)= 194.4+293.61

\

[..^min

1)'^^

^m

- KGi'^ 48+5.363' ^ ^™ ' L-^min

0'>

(7)

(8)

^max J

-0—m(Ti), N=1 1 -x--x{y), N^Tl --B--m(Ti),N=5 I - ^ - - x ( y ) , N=5 1 • ••- -mcro, N=0.2| —M- x(y), N=0.2|

0.4

mcri) 0.3 0.2

d x(y) d 0.7 ^-^ '

0.1 0 ^ .

3 - - - B-- "^»^™^^-«---^:j43

-01 0 2 y(/;m)

Figure 3. The variation of the weight fraction of titanium (m^^) and of the titanium carbide composition (x in TiC^) within the graded region.

Figure 4. The variations of the thermal expansion coefficient, a and of the Young modulus, E with the titanium carbide composition within the graded region.

The variation of the thermal expansion coefficient and of the Young modulus across the FGM region is shown in Fig. 4. Residual stresses after cooling or heating FGM system involve two principal contributions, one arising from stress equilibrium due to contraction or expansion and the other coming from moment equilibrium due to asymmetric variations in the composition, and hence the elastic and expansion characteristics, across the thickness. The total residual stresses can be determined according to [3], (Eq.9), and expressions (7) and (8) for a{y) and £(>') can be substituted in the integrals (10) and numerical integration performed for various configurations of the FGM region.

f Oresh)-E{y)\a{y)

I I

£(2)l|{3'£(l)-£(2)} 1 I

£(1) £:(l)£(3)-£^(2)

(9)

401

where: A{l) = Ja{y)E{y)dy,

A{2) = Ja{y)E{y)ydy,

-c

-c

£(1) = j£(y)dy, E{2) = jE{y)ydy, -c

-c

c

m

= fE{y)fdy

and c = | y J = | j _ | .

(10)

3. RESULTS AND DISCUSSION Residual stresses were calculated for three types of graded microstructure, each with a different value of the parameter N, corresponding to the concentration profiles shown in Figs.3. The presence of the metal substrate and the titanium carbide coating was not taken into account. The calculations were performed by considering both materials parameters (E, and a ) as being independent of temperature and assuming 1300 **K for the processing (diffusion anneal) temperature. The results of the numerical integration of the integrals (10) for a relatively thick FGM layer (-c = -5;/m and c = 5/<m, 2c=10/<m) and for three values of N are tabulated in Table 1 and the resulting stress distributions shown in Fig.5. Table 1. Calculated A(i) and E(i) values for three composition profiles {2c=l0pim). N=l

N=0.2

N^5

A(l) A(2)*106

32.78

30.20

35.25

-6.36

-3.03

-4.37

E(l)*10-6

4.06

3.58

4.56

E(2)

-1.22

-0.52

-0.91

E(3)*106

33.92

30.28

36.56

Similar calculations were also performed for a thinner (-c = -Ipim and c = l/<m, 2c=2/<m ) FGM layer. The results, with the exception of the difference in scaling along the x-axis, are similar to those obtained for the 10 pim wide FGM region, indicating that in this particular case, the results are not sensitive to the width of the transition zone. Turning now to the effect of parameter N, the concentration profile corresponding to the linear variation of the titanium carbide composition , namely for N=l, results in the lowest residual stresses, as shown in Fig.5 . For values other than N=l, the residual stresses have higher values. The concentration profile for N=5 (concave up) gives rise to residual tensile stresses at both low and high carbon content and compressive stresses at intermediate concentrations. A reverse distribution of stresses is generated for the N=0.2 (concave down) concentration distribution, with both low and high carbon content regions in titanium carbide being under compression and the intermediate region in tension. Particularly noteworthy in the latter case is the elevated level of the compressive residual thermal stresses at the left edge of the FGM region, the one located in the vicinity of the high carbon content titanium carbide, or in other words, in the vicinity of the cferattlic-liKe part of the FGM zone. In general, tensile stresses are undesirable in ceramics, cchversely the generation of compressive stresses on the ceramic side of a FGM region by designing a conoentotion profile corresponding to a value N<1 and cooling from high temperature, might be aavantagedus aJid improve the mechanical strength of an FGM-based transition region.

402

2 10°

1

• • ' I

I

I

I

I

I

I

I

— — Stress N=1 - » - Stress N=0.2 • • - • Stress N=5 ^8

-3 10'

I

-4

I

i

I

li

I

I

I

I

L.

J_

-I

L

0

yOtvm)

Figure 5. The residual thermal stresses in the FGM region for various configurations of the initial Ti-TiC multilayers.

ACKNOWLEDGMENT This work was supported by the Ministry of Science and Arts grant 5842-1-95.

REFERENCES. 1. M. Ohring, The Materials Science of Thin Films, p. 403-449, Academic Press, Inc. San Diego, 1992. 2. C.C. Chiu, J. Am. Ceram. Soc, 73 (1990) 1999. 3. K. S. Ravichandran, Materials Science and Engineering A201 (1995) 269. 4. K. Wakashima, T. Hirano and M. Niino, in "Space Applications of Advanced Structural Materials", ESA SP-303 (European Space Agency, 1990) p.97. 5. A. R. Andrievski and I.I. Spivak, " Strength of Materials Based on Refractory Compounds", Metallurgy, Chelyabinsk, 1989.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

403

The effect of constituent and microstmcture of composites on the residual thermal stress in TiC-NisAl FGMs Ji-Hui Wang and Lian-Meng Zhang State Key Laboratory of Materials Synthesis and Processing, Wuhan University of Technology, Wuhan 430070, P. R. China Based on the measured thermomechanical properties and the microstmcture of the graded layer of TiC-NisAl functionally graded material, using the analysis method, the interface residual thermal stress of TiC-NisAl sphere in the sintering process was calculated. The relationship between the stress and the content of NiaAl was presented. The results show that the failure mechanics of TiC-NisAl composite may be different in different content of NisAl. l.INTRODUCTION Ceramic-metal functionally graded material is a new type of composite developed in the recent years. This kind of material is characterized by the continuous transition or ladder tr^isition of the constituent, microstmcture and properties. The influence of constituent and microstmcture on the mechanical behavior of composites is a very important problem dravm great attention from all over the world. Due to the great difference between metal and ceramic in FGMs in mechanics and thermal expansion properties, when a powder metallurgy is used to synthesize ceramic-metal FGMs, high residual thermal stress is generated inside the material. This stress varies with the constituent and microstmcture of composites. In the present work, the graded layer properties of the thermal stress relaxation type functionally gradient material TiC-NiaAl were investigated. First, the thermomechanical properties of the each layer of the composite was tested. Second, the microstmcture of the composite was examined. Next, the interface residual thermal stress of TiC-NisAl sphere in the sintering process was calculated by the analysis method. Then, the relationship between the stress and the thickness of NisAl was also presented. Finally, the effect of the residual thermal stress on the properties of TiC-NisAl composite was discussed. 2.PROPERTIES AND MICROSTRUCTURES OF TiC-NisAl COMPOSITES Figure 1-2 show the relationship between Young's modulus, Poisson's ratio and the thermal expansion coefficient of TiC-NisAl composites and the volume fraction of NisAl, respectively. For the discussion of the relationship between the properties and the material microstmctures, figure 1-2 also give the numerical results of Young's modulus, Poisson's ratio and the thermal expansion coefficient of TiC-NisAl composites calculated by a two-

404 phase micromechaical formula [1]. From these figures, the tested values are close to their numerical values, but not monotonic with NisAl content. At 40% NisAl, in particular, the two tested curves are convex. This non-monotonic phenomenon in TiC-NisAl composite was attributed to the existence of a network structure at 30--50% of NisAl content. To verify the network structure in TiC-NisAl composites, figure 3 gives the SEM images of the TiC-NisAl composites with NisAl contents of 20%, 40%, 60%[2]. From figure 3, it is clear that, at 20% and 60%) of NisAl content, the TiC and NisAl are mutually dispersive phase in the composites and the images reveal a dispersive material microstructure; at 40% NisAl, however, the TiC particles and the NisAl particles are connected by themselves, and therefore a network SEM image forms. The percolation observed from the figure 3 is an important phenomenon; from figures 1-2, Young's modulus, Poisson's ratio and thermal expansion coefficient are seen to have a significant increase at 40% NisAl content, it needs more explanations. 0.6

500 —-•—Experiment 1

^

i^

1

1 .•

14

5 J 0.5

.400 I O

. —^ - Experiment

•^300

o ^(A J 0.4

1 ^^^-o^ f

-Calculation

Itt 12

^ " Z ^ ^ ^^

(D O

o c o 0.3

f200

-^^^

X

100

13

1 r ^

0.2

f

1

1

0

20

1

g 8 lit^T^^-

1

.-r"

1—'

1

40 60 Ni3Alvol%

80

Figure I.Young's modulus and Poisson's ration related to the content of NisAl.

0.1 100

.

6

0

20

1

1

40 60 Ni3Alvol%

80

1

100

Figure 2. Thermal expansion coefficient (xlO"^) related to the content of NisAl.

.1-* >d*

'^•Vr**',iiMy 60 urn

Optics microscope photo

Figure 3. Photomicrostructure of TiC-NisAl with various mixing ratios of TiC and NisAl.

405 3.MICROMECHANICAL MODEL FOR THERMAL STRESS CALCULATION In this paper, a spherical-shaped micromechanical model is considered. The diameter of the inside sphere is 5|am, the thickness of the outside spherical shell can be changed. The calculation condition for the thermal load is sintered at 1300°C and cooled down to room temperature (25°C). 3.1 Formulation Due to the spherical symmetrical problem, an spherical symmetrical equilibrium equation can be written as:

=0

dR - + - R The strain-displacement relations are expressed as

(1)

dUo

^'~

dR

(2)

The stress-strain relationship are expressed as aET l + ju 1 - 2 / / E

M

\ + ju , 1 - 2 / /

"

1-2;/

(3)

oET -O+Sj.

1-2//

Substituting expression (2) into equation (3) yields E \

fi

_ du„\ du,^

oET \-2ju

E

(7r =

Ur M -e+1 + //U-2//

(4)

oET 1-2//

Introducing expression (4) into equation (1) and applying expression (2), we find Un=-—-a—j] TR^dR + CR^—:r ^ \-ju R^ J^ R^ 2aE 1 EC 2ED 1 TrdR + c r „ = - \-juR' 1-2// l + ^R'

(5)

f

oE C7r =

1-//

U

EC TR^dR + 1-2//

ED

(6)

1

oET

\ + juR'

\-ju

(7)

The integrating constants C and D are to be determined by the condition of compatibility of the interface of the TiC sphere and the outside spherical shell NisAl.

406 3.2 The radial displacement of TiC ceramic sphere Considering a TiC ceramic sphere as shown in figure 4, the sphere is subjected to a uniform load qi, the radius is equal to a.

Figure 4. TiC ceramic sphere

Figure 5. NisAl spherical shell

The boundary conditions are given by (8) (9) Substituting equation (8) and (9) into expression (5) and (6) yields

1+A

r ^ l z M \^(^E,T

(10)

*«i

3.3 The radial displacement of NiaAl spherical shell Considering a NiaAl spherical shell as shown in figure 5, the inside of the spherical shell issubjected to a uniform load qi, the inside radius and the outside radius of this spherical shell are equal to a and b respectively. The thickness of NisAl is equal to subtracting afromb. The boundary conditions are given by

(o-«X^, = - ^ 2

(11)

(^«L*=0

(12)

Substituting equation (11) and (12) into expression (6) and using expression (5) leads to MR2=-

4l-2/^)

la^EJ'

a

3(l-//.)^^-«

r92

(i+;^VV lE^a%^-a')

(13)

3.4 The determination of the interface stress When the temperature varies , the TiC ceramic sphere and NisAl spherical shell are deformed simultaneously. The displacement of the interface between TiC sphere and NiaAl spherical shell are equal, namely, it must satisfy the condition of compatibility of the interface displacement, and the radial stress is equal and opposite in the interface.lt can be written as "/?] = ^R2

(14)

qx=qi=q

(i5)

407 From expression (14) and (15), we can obtain 1+M 3(1-A)

^ ^ 1 - 2A

^(1-2//^) la^EjT

2«i£ir 3(1-A)

-^

a^q

3(1-/^)" (16)

4.ANALYZING RESULTS AND DISCUSSION In this study, the properties of TiC and NisAl are for TiC, ai=7.4xlO"^/K, Ei=320GPa, m=0.195 for NisAl a2=11.9xlO-^/K,E2=199GPa, 1^2=0.295 When a=2.5|im, b=3.0|im, t=0.5|Lim andT=-1275K, substituting the upon data into equation (16) and solving , we find q =— 460.5MPa. For the same reason, we can obtain the results that when t is equal to 0.4|im, 0.3|im, 0.2|im and 0.1|im, q is equal to -392.1MPa, -313.8MPa, -223.9MPa, -120.2MPa, respectively. Based on the evaluated interface stress, the radial residual thermal stress and tangential residual thermal stress of different NisAl thickness are calculated. Figure 6 shows the relationship between the radial stress and NisAl thickness. In this figure, the radial stress in the TiC ceramic sphere is uniform compressive stress and the magnitude of this stress increases with the increment of thickness of NisAl. The radial stress is equal to zero on the outside surface of NisAl spherical shell. Figure 7 shows the relationship between the tangential stress and NisAl thickness. In this figure, the tangential stress in the TiC ceramic sphere is also uniform compressive stress and the magnitude of thes stress increases with the increment of thickness of NisAl. However, the tangential stress suddenly becomes tensile stress in the NisAl spherical shell area. The magnitude of this stress is increased with the decrement of the thickness of NisAl. For the thickness of NisAl is proportional to the content of NisAl, in other words, the radial compressive stress and tangential compressive stress increase with increment of the content of NisAl; the tangential tensile stress is decreased with the increment of the content of NisAl. This is a very important phenomenon and in agreement with the previous test results in which the strength of TiC-NisAl composites is increased with the increase of the content of NisAl. The larger the radial compressive stress, the harder to fracture in the radial direction in the material; the larger the tangential tensile stress, the easier to fracture in the tangential direction in the material. So, the failure mechanics of TiC-NisAl composite may be different in different content of NisAl. 5.CONCLUSIONS Through the analysis of a spherical-shaped micromechanical model of TiC-NisAl, the following conclusions can be drawn (1) The radial and tangential residual thermal stress in TiC sphere is a uniform compressive

408

(2)

(3)

stress and increase with the increase of the thickness of NisAl; The radial stress in NisAl spherical shell is rapidly decrease and equal to zero on the surface of the sphere, the tangential tensile stress is decrease with the increase of the thickness of the thickness of NisAl; The different content of NisAl may cause different failure mechanics. 100 0 -100^1\D

—1

0.5

1.0

m

i_

1.5

2.0

2

-200

•'•'!

-300

••!!

-400 -500 -600

t=0.1 t=0.3 t=0.5

t=0.2 t=0.4

-700 Radial length,|Lim

Figure 6. The relationship between radial stress and thickness of Ni3Al(t, |Lim) 1800 c« 1400

1=0.1 t=0.2

. 1000

1=0.3 t=0.4 t=0.5

200 ^

-200o[er -TQ:^-r.WT.TX-Jr.z.t.^r. ^.5

3

-600 Radial length,|;im

Figure 7. The relationship between tangential stress and thickness of Ni3Al(t, |im)

ACKNOWLEDGEMENTS This work was supported by the National Science Foundation. REFERENCES 1. J.H.Wang, Q.J.Zhang and D.H.Wu, J.Compos. Soc.,Vol.l3,No.2, (1996),89-93. 2. L.M.Zhang, J.Liu R.Z.Yuan and T.Hirai, Materials science & engineering, A203 (1995) 272-277.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

409

New application of FGM to identification of unknown multicomponent precipitates I. Itoh% H. Yaraada', Y. Kojima', Y. Otoguro% H. Nakata' and M. Matubara' ^Department of Mechanical System Engineering, Gunma University 1-5-1, Tenjin-Chou, Kiryu-Shi, Gunma, 376, Japan ''Graduate Student, Gunma University

1. INTRODUCTION It is well known that FGM is developed to relax the thermal stress which generates near the bonding interface of metal and ceramics. Recently FGM has been actively applied to many other fields. For example, .there is a research project on the development of energy conversion materials through the formation of gradient structures(Research periodiPhase 1 Fiscal 1993 to 1995)[1]. We have joined the project and studied the diffusion bonding of AI2O3 and Mo using the Al-Sn bonding method which we developed[2]. The objective of this research is to offer the new application of FGM that is the compositionally graded material(CGM). It is difficult to identify multicomponent precipitates because of the lack of multicomponent phase diagram and X-ray diffractometer's data. Then we try to develop the method by which the precipitates can be easily identitified, that is, the method using CGM. When the diffusion couple of Cu-10 mass^Sn and Al-0.5 mass%Sn is heated at 773 K which is lower than the eutectic temperature of 821 K in Cu-Al system, 4 intermetallic compounds are formed in the diffusion layer[3]. In this research, these compounds are judged by using CGM.

2. EXPERIMENTAL METHOD Diffusion couple of Cu and Cu-10 mass%Sn is pressed with screws of a clamp and treated at 1173 K for 360 ks in a vacuum of about lO^^Pa. Each specimen size is 10 mm cube. One of the 10x10 mm surfaces is selected as the bonding surface. The bonding surfaces of Cu and Cu-lOmass^Sn are polished with a #800 polishing paper. After bonding, the couple is cut perpendicularly to the bonding interface and separated into two pieces. These pieces have the same size of 20 mm in length, 10 mm in width and 5 mm in thickness. The cross section of one piece is polished with 1 //m-diamond powder for analysis of each element near the interface using EPMA. The other piece is used for identification of the intermetallic compounds as mentioned later. Both pieces change gradually in concentration of Sn from 0 to 10 mass%, that is, CGM.

410 Next the CGM is used as one piece of diffusion couple bonded to the other piece of Al-0. 5 massJKSn whose size is 20 mm in length, 10 mm in width and 5 mm in thickness. We call this couple the compositionally graded diffusion couple(CGDC). The bonding surfaces of the couple are 10 mmx20 mm in size and polished with jJSOO polishing paper. This couple is heated at 773 K for 7.2 ks under the pressure of 2MPa in the atmosphere. After this treatment, the couple is cut perpendicularly to the bonding interface and in parallel to the direction of the length at the middle point(Ref. Figure 1). The cross section of one piece is polished for the observation of the structure using SEM and optical microscope, and for analysis of each element near the bonding interface. From now on, we will adopt the simple expression of CGDC which contains the concrete heat treatment, that is, in this case Cu:eu-10Sn(1173K-360ks)/Al-0. 5 Sn(773K-7. 2ks).

3. EXPERIMENTAL RESULTS

Figure 1, (a) Schematic diagram of the method to make the compositionally graded diffusion couple. (b) Schematic view of the intermetallic compounds formed in the couple mentioned in (a). (c) and (d) Composition images using EPMA on Cu side and Cu-lOSn side near the bonding interface, respectively.

411 Figure l-(a) shows the schematic diagram of the method to make the compositional ly graded diffusion couple. The couple of Cu/Cu-10 mass^Sn has the diffusion zone of 2 mm in length after heating. Figure l-(b) shows the schematic view of the intermetallic compounds formed in the couple mentioned above. Figures l-(c) and l-(d) show the composition images using EPMA on Cu side and Cu-lOSn side near the bonding interface, respectively. The concentration of Sn gradually increases from 0% on Cu side to IQ% on Cu-lOSn side. Such intermetallic compounds as 7 2, 772 and 6 are formed at the interface of Cu side from Cu to Al-0. 5Sn alloy[4]. 4 intermetallic compounds of a'^d are formed at Cu-lOSn side. The intermetallic compounds of b, c and d are identified as r 2, V2 and 6 according to the continuation of structure from Cu side to Cu-lOSn side, respectively. Each thickness of these intermetallic compounds almost does not change throughout the couple and the total thickness of those keeps about 20 jim in width. The intermetallic compound of a, which is produced in the concentration range of Sn over 2.5%, shows an island having 20 to 30//m in length and 5 to 10 //m in width. This compound is cleared to be d phase(Cu-32 mass%Sn) in CuSn system by a quantitative analysis of EPMA. Figure 2 shows SEM and Sn X-ray image at the interface of the couple of Cu and Cu-lOSn bonded at 773 K for 14. 4ks. Sn concentrates at the interface of Cu and r 2. Thus the concentration of Sn on Cu-lOSn side is considered to create the intermetallic compound of 6,

Figure 2. SEM and X-ray image at the interface of the couple of Cu/Cu-lOSn bonded at 773 K for 14. 4ks.

4. APPLICATION OF CGDC TO OTHER CASE Figure 3 shows the structures throughout CGDCs of Cu:Cu-8Al(1173K-360ks)/ Cu-52Zn(/S-brass, 973K-57. 6ks) and Cu:Cu-4Si(1173K-360ks)/Cu-52Zn(973K-57. 6ks). In both couple, the growth of )S-brass into Cu alloy increases with increasing the content of Al or Si, and the width of /S-brass on Al or Si rich side is about two times compared with that on Cu side. Moreover, a structure like martensite is formed in yS-brass on Al or Si rich side. Also, the layer

412 of new precipitate is recognized at the interface of Cu and )S-brass on Si rich side. Thus, many information can be obtained by only one CGDC and will be available in the field of surface-treatment. 1. Itoh also reported the effect of the moving direction of the interface on morphological stability of al fi phase interfaces in the Cu-Zn system using the CGDC of Cu:Cu-52Zn(973K-72ks)/Cu(973K-57. 6ks)[5]. The details will be omitted here.

Weld I Interlace Cu-SAl

Figure 3. CGDCs of Cu-Al alloy or Cu-Si alloy and )8-brass.

5. CONCLUSION As a new application of FGM, the identification of unknown multicomponent precipitates in the system of Cu-Al-Sn is studied. Three among four intermetallic compounds can be identified as r 2, ?7 2 and Q in the system of Cu-Al with using the compositionally graded diffusion couple(CGDC). Also, CGDC is expected to use in many fields. As one example, the surfacetreatment is proposed to fit this case.

REFERENCES 1. S. Sasaki, FGM News No. 31(1996) 2. Journal of FGM Forum. The Society of Non-Traditional Technology. 2. I. Itoh, H. Yamada, T. Ichiba, Y. Kozima and Y. Otoguro, Proc. of (1995) 37. 3. I. Itoh, H. Yamada, Y. Otoguro, J. of the Japan Copper and Brass Research Association, Vol. 35(1996) 126. 4. I. Itoh, Y. Otoguro, K. Hasegawa, J. Japan Inst. Light Metals, Vol. 41(1991), 590. 5. I. Itoh, J. Japan Inst, of Metals, Vol. 40(1976) 1117.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

413

Evaluation of Graded Thermal Barrier Coating for Gas Turbine Engine M. Kawamura, Y. Matsuzaki, H. Hino, S. Okazaki Akashi Technical Institute, Kawasaki Heavy Industries, Ltd. Kawasaki-cho 1-1, Akashi, Hyogo, 673 JAPAN

The degradation processes of the bilayer and 6-layer graded TBCs are discussed both experimentally and computationally. Transient thermal cycle tests using the hydrogen combustion gas burner showed that the surface layer of both types of coating spalled immediately after the initiation of heating; namely, the top coat was delaminated near the top coat/bond coat interface in the bilayer coating case, while the graded coating showed delamination at the boundary between the 100% stabilized zirconia layer and the adjacent 80% stabilized zirconia-20% MCrAlY composite layer. The maximum heat load to cause the top coat spallation of the graded TBC was clearly greater and the substrate temperature was 200 degree higher than that of the bilayer coatings. Finite element model analysis simulating the thermal cycle tests demonstrated that the top coat spallation of both types of coating is caused by the buckling driven by delamination due to the transient large in-plane compressive stress development immediately after initiate heating. 1. INTRODUCTION Thermal Barrier Coatings (TBCs) are widely used in some components of the commercial gas turbine engines, including the combustion chambers, the nozzles and the blades, to control the high heat flux entering from the combustion gas to the structural components. Higher heat insulation performance and reliability of the TBC, however, is demanded to realize the gas turbines with higher performance and efficiency. Graded typed TBC consisted of the top coat (typically stabilized zirconia), the bond coat (MCrAlY), and the in-between composite layer(s) has been anticipated to exhibit higher thermal shock tolerance as compared with the conventional bilayer TBCs (stabilized zirconia / MCrAlY). In this study, the degradation processes of the bilayer and 6-layer graded TBCs are discussed both experimentally and computationally. 2. EXPERIMENTAL PROCEDURE Thermal shock resistance of the bilayer and the graded TBC were evaluated using the hydrogen burner rig test system, and oxidation behavior was evaluated by thermal exposure test. Coating was carried out as follows : For bilayer coating, 120 ju m of NiCoCrAIY bond coat was applied to the Co-base superalloy substrate using a vacuum plasma spray, and 230 ju m YSZ ceramic (Zr02-8wt%Y203) applied on the coating using atmospheric plasma

414

100 « m

(a) Bilayer type TBC

(b) Graded type TBC

Fig. 1 Cross sections of bilayer type and graded type TBC. spray. For the graded type TBC, after applying of bond coat (120 // m) on the substrate, a number of mixutures of NiCoCrAIY and YSZ were further applied on the previous coating. Each time the mixing ratio was changed by 20vol%, eventually forming a 6-layer multicoat from bond coat (0% YSZ) to top coat (100% YSZ). Each layer except bond coat was applied to 100 // m , and total thickness was 620 // m, so as to produce the same thermal barrier performance as bilayer coating. All test samples underwent heat treatment to removing Fig.2 Hydrogen burner rig stress after coating. Figure 1 shows the cross section of the bilayer and the graded coating prepared following the procedures above. 3. RESULTS AND DISCUSION 3-1 Thermal shock resistance of graded type TBC The thermal shock resistance of TBC was evaluated using the hydrogen burner rig test system. Figure 2 shows the external view of the test system. While forcibly cooling the back surface of the ^ 17mm cylindrical test sample, the front surface was heated by hydrogen-oxygen combustion gas. This can cause a temperature gradient in the thickness direction of the test sample. This system also does the real-time measurement and analysis of temperature on the material's front and back surfaces, heat flux during steady state heating, and effective heat conductivity. Moreover, monitoring AE (Acoustic Emission) makes it possible to determine the crack initiations. As the method of evaluating thermal shock resistance, the thermal load incremental method[l] was used and a heat cycle comprising 360 sec of heating and 240 sec of cooling was repeated with the thermal load step-likely

415 1600 1400

o Interface temperature • Heat flux

Spall

1200 1000 800 ^ *CO

CD Q. C

^1

600 400 c-

R....0.

a.-Q

800

a

v-^

1600

M...X)

^

2400

g

3200

4000

Elapsed time (sec) (a) Bilayer type TBC

4800

Elapsed time (sec)

(b) Graded type TBC Fig.3 Comparison of thermal shock resistance between bilayer and graded TBC.

Fig.4 Appearrence and cross-section of bilayer TBC after spallation by HBR test.

416

Fig.5 Appearence and cross-section of graded TBC after spallation.

raised for each cycle. Figure 3 compares experimental results for the two types coating with regard to coating / substrate interface temperature, heat flux. These results demonstrate that the graded TBC is more heat resistant and the substrate temperature to cause the top coat spallation was 200°C higher than that of bilayer TBC.

Sample diameter: j> 17mm Diameter of heated area:
Gas temperature: 200010 Heat transfer coefficient : 1250-4250 W/m^K

Topcoat thickness : 250// m Bondcoat thickness : 150// m

Substrate thickeness: 6mm

Coolant temperature: 20*0 Heat transfer coefficient :600W/m2K

Fig.6 Analysis model for hydrogen burner rig test.

3-2 Thermal shock degradation mechanism In the hydrogen-bumer-rig test, delamination occurred immediately after the start of the heating cycle. Figures 4 and 5 show the appearance and cross-section of both types of TBC after the test. In the bilayer coating, delamination occurred on the top coat side in the vicinity of the bond coat / top coat interface, whereas in the graded coating it occurred in the vicinity of the boundary between the zirconia 100% layer (the outermost layer) on the top surface and the second layer containing metal components of 20vol%. Regarding the delamination mechanism, FEM analysis was conducted for the bilayer TBC. TBC was applied on one of the end faces of a ^ 17 mm cylindrical sample, and the sample was heat-treated at 980*0. Assuming the condition after being cooled down to room temperature from heat-treated temperature as the initial condition, the state of the sample heated with the burner was analyzed. Figure 6 shows the analytical conditions. For data of the material's physical properties, values in literature were used[2][3]. Figure 7 presents the temperature history during one thermal cycle as a result of burner heating and the in-plane stress of the hot side surface. Figure 8 indicates the in-plane stress history immediately after the start of heating. From these results, it was revealed that the surface immediately after heat treatment was already compressive stress state and that the compressive stress was increased by heating, reach the maximum within a very short time

417 namely 1 sec. Therfore, the top coat spallation of top coating is caused by the buckling driven by delamination due to the transient large in-plane compressive stress development immediately after initiate heating.

-50

k

-100 -150 -200 -250 -300 -350

t

-400 3-3 Lifetime of graded type TBC 1500 f heating cycle cooling cycle based on thermal exposure tests In the above section, it was • I demonstrated that graded TBC exceeds bilayer coating in thermal shock resistance. As shown in Figure 5, however, the observation of crosssection after hydrogen-burner rig tests implies that the surfaces of NiCoCrAIY Elapssed time (sec) particles dispersed throughout the YSZ ceramic had already been oxidized. It Fig.7 In-plane stress and temperature is generally considered that the histories at hot side surface on thermal oxidation of metal components governs cycle. the lifetime of the graded coating. Hence, high-temperature atmospheric exposure tests were necessary to evaluate damage morphology and service life related to oxidation. The -150 test samples were cylindrical with the -200 dimensions of ^ 10 X 70^ mm. 1" Exposure tests were carried out at -250 1100X3. The test samples were cooled -300 ^ 1 • to room temperature every 50h and ; • »® tI *• * -350 checked for cracking. Figure 9 presents ! I »• the test results of bilayer coating. The -400 longitudinal cracks in the coating as Elapsed time (sec) shown in Figure 9 often developed at either end of the test sample, growing Fig.8 In-plane stress histories at hot side in the axial direction of the cylinder as surface just after heating cycle started. time elapsed. Coating delamination from the substrate occurred on the top coat side on the bond coat / top coat boundary. Although similar to the delamination in the hydrogen burner rig tests, this delamination results from the oxidized of bond coat. In the graded TBC, as shown in Figure 10, turtle-shell-pattern cracks which developed throughout the coating, eventually breaking the coating into pieces. The lifetime of graded TBC was nearly half that of bilayer TBC. Therfore, to apply ceramic / metal graded coating to industrial gas turbine engines, there is a need to provide preventive measures against the oxidation of metal components.

r.

r::::::::::::::::::::i V

.'••. - 1

[. ..

Gas side heat transfer coefficient • 1250

• 42S0

,••••'•••••

418

^vCv;l?.

Fig. 9 Damage of bilayer TBC by thermal exposure test (11001), 500h).

Fig.lO Damage of graded TBC by thermal exposure test (1100"C, 200h). 4 . CONCLUSIONS As a result of evaluating the thermal shock resistance of graded type thermal barrier coating and its lifetime in a high-temperature oxidizing atmosphere, the following results were revealed. (1) The maximum heat load to cause the top coat spallation of the graded TBC was clearly greater and the substrate temperature was 200 degree higher than that of the bilayer TBC. (2) Finite element model analysis simulating the thermal cycle tests demonstrated that the top coat spallation of both types of coating is caused by the buckling driven by delamination due to the transient large in-plane compressive stress development immediately after initiate heating. (3) From the thermal exposure test, the lifetime of the bilayer TBC is longer than that of the graded one. There is a need to provide preventive measures against the oxidation of metal components.

REFERENCES 1. A. Kawasaki, A. Hibino and R. Watanabe : J. Japan Inst. Metals, 56(1992) 472. 2. R. V. Hillery, B. H. Pilsner, R. L. McKnight, T. S. Cook, and M. S. Hartle (GE): NASA CR180807, November 1988. 3. J. T. DeMasi, K. D. Sheffler, and M. Ortis (United Technologies, Pratt & Whitney): NASA CRl82230, December 1989.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

419

Mechanical and electrical properties of multilayer composites of silicon carbide J. Hojo^ F. Hongo^, K. Kishi^ and S. Umebayashi^ ^Department of Chemical Science and Technology, Faculty of Engineering, Kyushu University, 6-10-1 Hakozaki, Higashi-ku, Fukuoka-shi, 812 Japan '^Graduate School, Kyushu University, 6-10-1 Hakozaki, Higashi-ku, Fukuoka-shi, 812 Japan* ^Kyushu National Industrial Research Institute, Shuku-machi, Tosu-shi, 841 Japan

Multilayer composite disks of SiC-TiC-Ni system were fabricated by a powder sintering technique. The composition was graded stepwise between SiC and TiC, and between TiC and Ni. Cracking in the ceramic part was prevented by increase of the number of stacked layers. It was found by the Vickers indentation technique that large tensile stresses existed on the surface of SiC layer and at the disk edge. The reduction of residual stress by compositional gradient was confirmed by an analytical calculation. Furthermore, the electrical property of SiC-TiC gradient material was evaluated. The electrical conductivity measured in the direction from SiC to TiC revealed the ohmic behavior, indicating that metallic TiC is useful as electrode of semiconductive SiC. 1. INTRODUCTION Silicon carbide (SiC) has been considered as potential high-temperature material with high strength and chemical stability. The bonding of SiC to metals is useful to develop the applications. However, when SiC is bonded directly to metals, the serious reaction weakens the bond. Hojo et al. have reported that SiC can be stably bonded to Ni by using the intermediate TiC layer, and the residual stress can be reduced by grading tiie composition at the boundaries [1,2]. Semiconductive SiC has been receiving a lot of attention as high-temperature electronic device. In this application also, TiC may be useful as electrode because of its high electrical conductivity. In the present paper, the properties of gradiently-bonded SiC materials were investigated with the following viewpoints: the residual stress of SiC-TiC-Ni system and the electrical conductivity of SiC-TiC system.

* Present address: Toto Ltd., Kitakyushu-shi, Japan

420

•'-' ^'-'^^^mm^^^^^^MiM. ••;::•:'; ''^\' •"' 'i-'^^^^m^^^^W^^^

^^f

Figu^*® 1. Crack generation in SiC-TiC-Ni multilayer composite.

2. EXPERIMENTAL The multilayer composite disks of SiC-TiC-Ni system were fabricated by two-step process. Commercial powders were supplied as follows: P-SiC 0.3|J,m, TiC 2-6[im and Ni 4-7|im. AI2O3 (2wt%) was added to SiC powder as sintering aid. SiC powder, SiCTiC mixed powder and TiC powder were stepwise stacked in a graphite mold and hotpressed at 2000°C and 28-64MPa for 0.5h in Ar. The sample was annealed for Ih at 1600*'C before cooling. TiC-Ni mixed powder and Ni powder were stacked and coldpressed in a steel die. The ceramic part was put on the metallic part, and heated at 1320^C for 0.5h in Ar under a light load of 27-60kPa. The multilayer composite disk of SiC-TiC system used for electrical measurement was hot-pressed in N2. Cracks were observed on the top surface and the section surface of the disk with an optical microscope. The microstructure was observed by SEM. The surfaces were polished with diamond paste. The Vickers indentation technique was used to evaluate the direction of residual stress. The residual stresses were calculated by an analytical technique on the assumption of elastic condition [2]. For electrical measurement, Pt electrode was formed on the disk surfaces with Pt paste. The electrical resistivity was measured by two-probe method.

3. RESULTS AND DISCUSSION 3.1. Stability of SiC-TiC-Ni system The multilayer composites of SiC-TiC-Ni system were fabricated with the following structures. (a) SiC/TiC//Ni (3 layers) (b) SiC/SiC50yTiC//TiC50/Ni (5 layers) (c) SiC/SiC80/SiC60/SiC40/SiC20/TiC/mC67A'iC33/Ni (9 layers) (d)SiC/SiC90/SiC80/SiC70/SiC60/SiC50/SiC40/SiC30/SiC20/SiC10mC //TiC90/TiC80/TiC70mC60A'iC50/TiC40/TiC30/TiC20mC10/Ni (21 layers) where the compositions are shown in volume fraction and // means the bonding position of ceramic part to metallic part. The size of composite disk was 10mm in diameter and 3mm thick for samples (a), (b) and (c), and 15mm in diameter and 4.5mm thick for sample (d). Sample (a) was non-gradient material. Samples (b), (c) and (d) were gradi-

421

z(mm)

z(mm)

c •

1

•

4.5• /

1

•

T - ^ SiC

C •I

y

/V

V

/;

•\"

\\

y

.

T »

2500'a

A : a X (SiC surface) H O :CTz (SiC+TiC layer) |

•

^ 2000

1

0 1 a x(GPa)

.

Ni

I 1500

X

TiC

/ •

1

••

X

y*

-1

»

^s ^\ \v'v

/ // < // y^ // /.

t r

.

v:

\ V^ ^\ 1 \. > -1 0 1 a z(GPa) .

i

1000 1

Figure 2. Distribution of residual stress in SiC-TiC-Ni system. T : tension, C : compres.

500^ 3 5 1

1

^^-^°^^^^—^11:^ —o o 1

1

10 15 Layer number

1

20

Figure 3. Change in residual stress with the number of layers.

ent materials, in which the composition was linearly graded between SiC and TiC, and between TiC and Ni. In the present system, the thermal expansion coefficient varies in the following order: SiC(4.8xlO-^rC) < TiC(8.0xlO-V°C) < Ni(16.8xlO-^rC). When the SiC-TiC part was mechanically polished after hot-pressing, cracks were sometimes formed on the surface of TiC layer, especially in samples (a) and (b). This is caused by residual stresses generated from the gap in thermal expansion coefficient between SiC and TiC. After bonding to the TiC-Ni part, cracks were also formed in the ceramic part of SiC-TiC-Ni system of samples (a) and (b). The example is shown in Figure 1. Cracks proceeded typically in the perpendicular direction from SiC surface and in the parallel direction along the interface in SiC layer. Cracks were not observed in samples (c) and (d). 3.2. Residual stresses The stress distributions, which cause cracking in the ceramic part, are typically characterized by two stress components as follows: tensile stress parallel to interface near disk center; tensile stress perpendicular to interface near disk edge [2]. The residual stresses were calculated for cooling process after bonding of the ceramic part to the metallic part. Figure 2 shows one example of the stress distributions along gradient direction (z-axis). ax is the residual stress parallel to interface. Large tensile stress exists on the surface of SiC layer. This stress causes cracking in the perpendicular direction from SiC surface, a z is the residual stress perpendicular to interface at disk edge. Large tensile stress also exists in the range from TiC to SiC. This stress causes cracking in the direction parallel to interface. These calculations can explain the cracking behavior of layer composite of SiC-TiC-Ni system as indicated with a x and az in Figure 1. The variation of residual stresses with the number of layers was evaluated with respect to tensile stresses on the surface of SiC layer and in the middle of SiC-TiC part at disk edge in the stacking model seen in Figure 2. The results are shown in Figure 3. It is confirmed from the calculations that the residual stresses are effectively reduced by increase of the number of layers.

422

iKjc(//)

H100M,m Figure 4. Vickers indentation on section surface of SiC+TiC layer.

Figure 5. Fracture toughness of section surface of SiC-TiC part (sample (c)).

Crack propagation by the Vickers indentation was observed to prove the presence of residual stresses. When the surface of SiC layer was indented, the crack propagation was serious near disk center. This means that tensile stress expanded the cracks on SiC surface. Figure 4 shows the crack propagation on the section of the ceramic part. It was found that the cracks parallel to interface were longer than those perpendicular to interface. This means that tensile stress was large in the perpendicular direction to expand the cracks. These results are well consistent with the expectation from the calculation of stress distributions. The fracture toughness of section surface was evaluated from the crack length in each direction, that is, parallel or perpendicular direction, by using Niihara's equations [3]. The result of sample (c) is shown in Figure 5. The fracture toughness was smaller in the parallel direction than in the perpendicular direction over the range from SiC to TiC, because large tensile stress existed in the perpendicular direction. The fracture toughness exhibited the maximum around 50vol% TiC in both directions. Such an improvement of fracture toughness has been reported in the particulate composite of SiC-TiC system [4, 5]. In the present work, the cracks were formed especially at or near SiC layer although large tensile stress was imposed over the whole range of the ceramic part as seen in Figure 2. It is thought that some strengthening effects work in other layers. 3.3. Electrical properties of SiC-TiC system The multilayer composite of SiC-TiC system was sintered in N2 to make SiC semiconductive by nitrogen doping. It is known that nitrogen is incorporated into SiC and works as electron donor [6]. However, N2 atmosphere retards the sintering of SiC. Figure 6 shows typical microstructures of SiC+TiC layers sintered in N2. Many pores were observed when the content of SiC was above 70vol%. In Figure 6, black matrix is assigned to SiC and white grains to TiC. With increasing content of TiC, TiC grains connected to each other and formed the network. The V-I curve of the multilayer composite, which had the same structure as in the

423

TiC : 10vol%

TiC : 50vol%

Figure 6. Section surface of SiC-TiC multilayer composite (sintering atmosphere : N2).

i a -0.4

-0.2

0.0

0.2

0.4

0.6

Voltage (V)

Figure 7. V-I curves of SiC-TiC systems. Specimen size : d=10mm<^, t=2.8mm (TiC), 1.8mm(SiC), 3.0mm (SiC/TiC)

100 200 Temperature (°C)

300

Figure 8. Temperature dependence of electrical resistivity.

ceramic part of sample (d), was measured at room temperature by two-probe method in the direction from TiC to SiC. The result is shown in Figure 7, including the data of monolithic SiC and TiC sintered under the same condition. All curves were linear, indicating the ohmic contact between the electrode and the specimen, and between SiC and TiC in the multilayer composite. Figure 8 shows the temperature dependence of electrical resistivity. The electrical resistivity of SiC decreased with a rise of temperature because of its semiconductive property. On the other hand, the electrical resistivity of TiC was low and increased with a rise of temperature because of its metallic conductivity. The multilayer composite of SiC-TiC system exhibited semiconductive property similar to that of monolithic SiC but the electrical resistivity was low because of bonding to TiC with low resistivity. These results mean that gradiently-bonded TiC is useful as electrode of semiconductive SiC device.

424 4. CONCLUSIONS SiC can be stably bonded to Ni by using TiC intermediate layer. However, when they are directly bonded, cracks are formed in the ceramic part owing to residual stresses. According to the analytical calculation of residual stress and the crack propagation by the Vickers indentation, there are two stress components inducing the cracking. One is tensile stress parallel to interface, which is large on SiC surface. Another is tensile stress perpendicular to interface, which is large in the ceramic part near disk edge. The cracking can be prevented by grading composition between SiC and TiC, and between TiC and Ni, because of the reduction of residual stresses. Such a gradient bonding may be useful to joint SiC to Ni for structural application. TiC can be used as electrode of semiconductive SiC because TiC is metallic conductor with low electrical resistivity. In this case also, the gradient bonding is important to prevent cracking. SiC-TiC gradient material reveals the ohmic behavior and has low resistivity. Such a gradiently-bonded TiC electrode may be available for hightemperature application of semiconductive SiC device because TiC is thermally stable as well as SiC.

REFERENCES 1. J. Hojo, T. Kajiya, T. Kuga, K. Kishi and S. Umebayashi, Proc. 1993 Powder Metallurgy World Congress, Kyoto (1993) 1279. 2. J. Hojo, T. Kajiya, S. Fukuoka and K. Kunoo, J. Jpn. Soc. Powder and Powder Metall., 41 (1994) 658. 3. K. Niihara, R. Morena and D. P. H. Hasselman, J. Mater. Sci. Letters, 1 (1982) 13. 4. S. Shirasaki and A. Makishima (eds.). Handbook of Control Technique of Superfine Ceramics (Japanese), Science Forum, Tokyo (1990) pp.448-457. 5. H. Endo, K. Tanemoto and H. Kubo, Proc. of 4th Int. Symposium on Sci. and Tech. of Sintering (Sintering '87), Tokyo (1987) 1052. 6. K. Koumoto, M. Shimohigoshi, S. Takeda and H. Yanagida, J. Mater. Sci. Letters, 6 (1987) 1453.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 1997 Elsevier Science B.V.

425

The Effect of Thermal Shock on the Thermal Conductivity of a Functionally Graded Material* AJ.Slifkaa, A.Kumakawab, B.J.Fillaa, J.M.Phelpsa, N.Shimodac aNIST, 325 Broadway, Boulder, Colorado, 80303, USA. bKakuda Research Center, National Aerospace Laboratory, STA, 1 Koganezawa, Kimigaya, Miyagi 981-15 Japan. cSteel Research Laboratories, Nippon Steel Corporation, 20-1 Shintomi, Futtsu, Chiba 299-12 Japan.

Abstract We have measured the thermal conductivity of a Ni20Cr / 8% yttria-partially-stabilized-zirconia functionally graded 1.1 mm thick coating on a substrate of 403 stainless steel. We measured thermal conductivity of the as-received coated specimen, then thermally shocked the specimen and measured thermal conductivity again. The measurements were done using an absolute, steady-state technique over a temperature range from 400 K to 1200 K. The specimen was thermally shocked by heating in a furnace to 475 K, then quenching in water at 295 K. We discuss the effect of moderate thermal shock on the thermal conductivity of the coating. 1. INTRODUCTION Functionally Graded Materials (FGMs) are known to be useful for reducing thermal stress in applications where high heat fluxes and large temperature gradients exist [1]. Applications of this type are found primarily in aerospace, where nose caps, leading edges and engines will experience these types of severe thermal loading [2]. The thrust chamber Hner for rocket combustors is a target appHcation of Ni20Cr / 8% yttria-partially-stabiUzed-zirconia FGM due to the low thermal conductivity and heat resistance of 8YSZ and the oxidation resistance and matching coefficient of thermal expansion of NiCr alloys to typical substrate metals. The thrust chamber liner is actively cooled, usually by liquid hydrogen or other liquid fuel, and heated on the other side by the combustion of nitrogen tetroxide / monomethyl hydrazine or other propellants, which results in a large thermal gradient with high heat flux. We have measured the thermal conductivity of a 1.1 mm thick Ni20Cr / 8YSZ FGM over the temperature range from 400 to 1200 K to evaluate the potential thermal performance of the material system. We have also applied moderate thermal shock to the coating and re-tested thermal conductivity to

^Contribution of the National Institute of Standards and Technology and therefore not subject to copyright in the USA.

426 discern any possible short-term degradation effects on the thermal conductivity of the coating. 2. SPECIMEN AND TEST PROCEDURE The specimen was low-pressureplasma-sprayed (LPPS), which results in high bond strength and little or no oxidation of the metallic FGM component. A four-port spray gun was used, which simultaneously introduces the metalHc and ceramic powders to the plasma jet using two surrounding ports for each component. The two ports are diametrically opposed, which results in a uniform mixture of components in the plasma jet. By using a spray atmosphere of 26.6 kPa, a coating Figure 1. Optical micrograph (50x) of optimum density is made [3]. The showing the graded structure of the 1.1 coating was sprayed on a substrate of 403 stainless steel 69.85 mm in mm thick FGM coating. diameter. A linear grading of Ni20Cr to 8YSZ was sprayed and then polished to a 0.19 |im surface finish. The total FGM coating thickness was 1.1 mm. The coating was sprayed using six steps, beginning with pure Ni20Cr alloy and adding 20% 8YSZ powder for each step up to 100% 8YSZ. Each layer was approximately 200 |im thick. Approximately 100 |im was polished from the top of the coating after spraying, to achieve theflatnessand surface finish required for the measurement. Figure 1 is an optical micrograph of the coating, showing the graded structure. Figure 2 is a backscatter electron micrograph showing the splat structure typical of plasma-sprayed materials. The bright areas in the micrograph indicate microcracks, pores and splat boundaries. Figure 2. Backscatter electron The thermal conductivity was measured using an absolute, steadymicrograph (2000x) of the top layer of state technique that is similar to the the coating. ASTM-177 guarded-hot-plate specification [4]. The primary difference between the technique used here and a typical guarded-hot-plate is that one specimen

427

Radiative Heat Sink (Nickel)

^Thermocouple ^,..^
Outer Guard. Heater

Bottom Heater \ Plate

ooooooooooooooooooooood. > ^ n »^^/" f

\^^ '>-^. \^^ '>^. .>\. .y\.

/y // =//

r

// =//

f/ =/y

// =^/^ ff =/y

ff =/y

.Main/Inner-Guard Heater Plate -Control RTD Plate W

ff =

-Heat Flux Transducer

Figure 3. Schematic of the measurement stack of the one-sidedguarded-hot-plate. is used in our apparatus, rather than two. Figure 3 shows the important features of the measurement system. Heat-flow calculations show that the one-sided nature of the apparatus reduces heat losses. Details of the apparatus design are found elsewhere [5, 6]. The principle of operation is to create one-dimensional axial heat flow through the specimen so that the Fourier heat conduction equation may be used to determine thermal conductivity: q = -kA

dx

(1)

where q is steady-state heat flow, k is thermal conductivity, A is the specimen cross-sectional area and -dT/dx is the temperature gradient [7]. Heat is injected electrically into the main / inner guard heaters at a constant power density over the plate surface. The plate below the main / inner guard plate has a platinum resistance temperature detector (RTD) embedded in it to measure the absolute control temperature for the measurement stack. The outer guard cylinder is controlled to the same temperature as that of the main / inner guard plate. The inner and outer guards keep heat from the main heater from flowing radially. The bottom heater plate is also controlled to the same temperature as the main / inner guard plate so that heat from the main heater cannot flow downward. Therefore, all of the heat generated by the main heater flows upward, through the thermocouple plates and the specimen, and into the heat sink, where the heat is radiated away to the cooler furnace. Knowledge of the specimen's crosssectional area, coating thickness, temperature difference across the specimen as determined from the thermocouple plates, and steady-state heat flow as measured using the electrical power input to the main heater, yields the raw data needed to calculate the thermal conductivity.

428 This raw data, when used in equation 1, gives total I I I I I I I I I I I thermal conductivity, E which is shown in figure 4. The total thermal 5 . i....Q..O..O-.6.^.conductivity data 6 o o o p '^ represent essentially five •>

•p 23 stainless steel substrate, o we use data obtained "a previously on 410 stainless steel [8]. Figure 5 shows thermal conductivity as a function E 21 of temperature for 410 stainless steel as measured 20 in the guarded hot plate. 400 500 600 700 $00 900 1000 We extrapolated the quadratic fitting function Temperature, K to 1200 K for this work. Figure 5. Thermal conductivity of 410 The properties of 403 and stainless steel. 410 stainless steel are similar since 403 and 410

429 stainless steels have the same 0.000 28 composition, except 403 stainless steel has a narrower 0.000 26 acceptable chromium range [9]. We also use 0.000 M data measured previously on the interfacial resistance 0.000 2 Z L between stainless steel 410 and the upper thermocouple 0.000 2 plate for the interfacial resistance o between the t : 0.00018 substrate and measurement plate 0.00016 required here. 500 600 700 800 900 1000 400 Again, we extrapolate our data Temperature, K to 1200 K for use Interfacial resistance between 410 Figure 6 here. Figure 6 shows the substrate- stainless steel and the measurement plates. coating interfacial resistance function. The interfacial resistance between the coating and the lower thermocouple plate is different from the interfacial resistance between the substrate-measurement plate due to light oxidation of the steel substrate, which is not observed for the coating. We have observed that the interfacial resistance of ceramics depends primarily on the surface finish of the ceramic material, since ceramics are chemically stable. Therefore, we use an interfacial resistance that was previously measured between Pyroceram 9606, of the same surface finish as the FGM coating, and the upper thermocouple plate as our interfacial resistance function for the coating-measurement plate interface [6]. Tliermal conductivity of the FGM coating can then be extracted from the total conductance data by using these mterfacial resistances and the thermal conductivity data for 410 stainless steel. The estimated uncertainty of the measurement system is 5%. 3.

RESULTS

Figure 7 shows thermal conductivity data for 5 test cycles of the specimen. A test cycle entails a conditioning thermal ramp, where the specimen is heated from 473 K to 1173 K in steps of 100 K, and a final thermal ramp. At each temperature of interest, the system is allowed to thermally equilibrate for a minimum of 3 hours. The conditioning ramp is used to allow complete heUum diffusion between the plates, mechanical settling due to thermal expansion differences, and stable formation of oxides on metallic specimens. Since the interfaces between measurement plates and the specimen are changing during this ramp, we do not use any of the data during this ramp in the analysis. Following the conditioning ramp, the temperature is dropped to 373 K, then the final thermal ramp is done in 50 K steps, up to 1173 K. The thermal cycling due to the testing itself is a mild form of thermal shock. Thermal conductivity dropped by an average of 12% from the first test to the second test, probably due to significant microcracking from differential thermal expansion of the two components in the coating. At lower temperatures, test 1 yielded thermal conductivity data that were appreciably higher than corresponding data from any of the four subsequent tests. Since the coating had

430 not yet been subjected to high temperatures, the coating would have had a 2.5 low microcrack density o t95t Tl until sufficiently high test i\ n temperatures would give test 3l o £ enough thermal stress X test 4l from the mismatch in the test sl + thermal expansion coefficients of the two coating components to •S 1.5 o oO start generating o microcracks. After the © 3 specimen had been tested QH = C twice, each successive 2*+ + — 1 test was preceded by an -»*additiond moderate _i_ ~ thermal shock to the specimen outside the apparatus before re0.5 I testing by heating the 200 1000 400 600 800 1200 1400 specimen to 473 K and Temperature K quenching into water at 295 K to induce a small Figure 7 . Thermal conductivity of the FGIVI amount of microcracking coating, showing data from 5 tests. in the specimens. Figure 7 shows that the thermal conductivity did not change appreciably from test 2 to test 3, which would be due to this moderate thermal shock. The first external thermal shock appears to have had no effect on the thermal conductivity of the coating. The next test (test 4) showed an additional average drop in thermal conductivity of 17%, while thefinaltest (test 5) showed an additional average drop in thermal conductivity of 9%. Even a moderate thermal shock is seen to significantly decrease the thermal conductivity of this coating. Since the thermal conductivity does not increase from test to test, the upper temperature limit of 1200 K is probably too low to allow sintering of the splats.

xji+

of

An interesting feature of figure 7 is the large drop in thermal conductivity between 1050 K and 1100 K. A possible explanation for the drop in thermal conductivity could be based on differences in the thermal expansion coefficients of 8YSZ and Ni20Cr. Since the thermal expansion of zirconia is about half that of Ni20Cr, there could be a separation at the zirconiaNi20Cr splat boundaries as temperature increases and the metal expands more rapidly than the 8 YSZ. Tliis phenomenon would only occur in the metal-rich portion of the coating, though. The coating is deposited at approximately 725 K to 925 K. Since the discontinuity in thermal conductivity occurs above this temperature range, this explanation is at least plausible. A more likely explanation for the sharp drop in thermal conductivity between 1050 K and 1100 K is the occurrence of a phase change in the coating. One thing to note is the repeatability of the drop in thermal conductivity. We observed a small, -50 K temperature hysteresis in this thermal conductivity transition upon cooling. The repeatability over many thermal cycles and the small temperature hysteresis upon cooling points toward a phase transition occurring, causing the decrease in thermal conductivity. The mechanism here would be a splat-boundary opening occurring when monoclinic 8YSZ material transforms to a tetragonal crystal structure. The monoclinic-to-tetragonal phase transformation in zirconia results in a volume contraction of 3 to 4%. Even a small volume contraction in the 8 YSZ splats could result in a much more

431 efficient thermal barrier at the splat boundaries. The thermal hysteresis of the zirconia phase transformation has been observed in polycrystalline zirconia [10]. The phase transformation possibly observed here occurs at a lower temperature, which could be due to fine grain structure. Since plasma-spray is a non-equilibrium process, grain sizes of material within the splats may be very fine. The monoclinic-to-tetragonal phase transformation in zirconia is loiown to be a function of grain size [11]. We plan to powder some of the coating and do xray diffraction analysis at room temperature, to determine whether a significant amount of monoclinic material exists in the coating. As Httle as 2% monoclinic phase may be enough to cause the discontinuity in thermal conductivity because the effect would not be due to a difference in phonon transport through different crystallographic phases, but due to the opening of splat boundaries. If we observe a significant amount of monoclinic 8 YSZ at room temperature, we plan to do high-temperature x-ray diffraction analysis on the coating to determine the phase transformation temperature in our plasma-sprayed material. The increase in thermal conductivity as temperature increases from 400 K to 1050 K is probably due to the Ni20Cr, as we have observed virtually no temperature dependance of thermal conductivity for pure 8YSZ coatings from 400 K to 800 K [8]. The strong temperature dependance of thermal conductivity above 1100 K is probably due to a radiative component of thermal conductivity in the 8 YSZ. If a radiative contribution is becoming a significant contributor to thermal transport in this temperature region, these data should be viewed as apparent thermal conductivity, composed of both conductive and radiative heat transfer. 4. CONCLUSIONS We have direcdy measured the thermal conductivity of a plasma-sprayed Ni20Cr / 8 YSZ FGM. Thermal conductivity decreased about 12% after the first test, probably due to microcracking from differential thermal expansion between the two coating components. Subsequent tests showed no drop, then a 17 % drop, then a 9 % drop, on average, of the thermal conductivity of the coating. We observe a sharp drop in thermal conductivity of the coating between 1050 K and 1100 K, probably due to a phase transformation or thermal expansion effects. We plan to do future work investigating this phenomenon.

432 REFERENCES 1. R.L. Williamson and B.H. Rabin, "Numerical Modeling of Residual Stress in NiAI2O3 Gradient Materials," Ceramic Transactions, 34, American Ceramic Society, pp.55-65 (1993). 2. R. Watanabe, A. Kumakawa and M. Niino, "Fabrication of Panel Assembly of Functionally Gradient Material with Active Cooling Structures," Ceramic Transactions, 34, American Ceramic Society, pp. 181-188 (1993). 3. N. Shimoda, H. Hamatani, Y. Ichiyama, S. Kitaguchi and T. Saito, "Fabrication of PSZ / Ni-Alloy FGM by Applying Low Pressure Plasma Spray," Advanced Materials '93, III/ B: Composites, Grain Boundaries andNanophase Materials, M. Sakai et al. Ed., Trans. Mat. Res. Soc. Jpn., 16B, pp.1263-1266 (1994). 4. "Standard Test Method for Steady State Heat Rux Measurements and Thermal Transmission Properties by Means of the Guarded-Hot-Plate Apparatus," C 177-85, Annual Book ofASTM Standards, 6, ASTM, pp. 17-28 (1988). 5. B.J. Filla, "Design and Fabrication of a Miniature High-Temperature Guarded-HotPlate Apparatus," Thermal Conductivity, 21, Plenum Press, pp.67-74 (1990). 6. B.J. Filla, "A Steady-State High-Temperature Apparatus for Measuring Thermal Conductivity of Ceramics," submitted to Review of Scientific Instruments. 7.

W.H. McAdams, Heat Transmission, 2nd Ed., McGraw-Hill, p.7 (1942).

8. A.J. Slifka, B.J. Filla, J.M. Phelps, C.C. Berndt and G. Bancke, "Thermal Conductivity of a Zirconia Thermal Barrier Coating," submitted to Journal of Thermal Spray Technology. 9.

R.A. Lula, Stainless Steel, ASM, p.37 (1986).

10. G.M. Wolten, "Diffusionless Phase Transformations in Zirconia and Hafnia," Journal of the American Ceramics Society, 46:9, pp.418-422 (1963). 11. D.J. Green, R.H.J. Hannink and M.V. Swain, Transformation Toughening of Ceramics, CRC Press, p.24 (1989).

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

433

Non-Destructive Evaluation of Carbon Fibre-Reinforced Structures using High Frequency Eddy Current Methods G. Mook, O. Koser, R. Lange Institute of Materials Science and Materials Testing, Otto-von-Guericke-University Magdeburg, POB 4120, 39016 Magdeburg, Germany Eddy current (EC) inspection of carbon fibre-reinforced structures bases on their anisotropic electrical properties. Great differences in conductivity between carbon fibres, polymer matrix and integrated functional components contribute to this circumstance. The paper describes suitable eddy current probes, the fundamental idea of modelling and measurement of EC distribution and provides a short survey of application potential to characterise CFRP nondestructively.

1. EDDY CURRENT PROBES The essential idea of EC inspection of CFRP is to use the anisotropic conductivity of the material [1,2]. Along the carbon fibres much higher conductivity is found than in the opposite direction (9... 15 kS/m and 0.07... 0.7 kS/m respectively). Additionally, capacitive couplings between the fibres can be used at high test frequencies. Non-axial EC probes (Fig. i) are necessary to detect these anisotropic effects [3,4]. Differential probes (Fig. 2) are capable to detect local inhomogeneities due to their gradients in conductivity. The comparatively low conductivity of the material lets eddy currents penetrate deep into the material thus opening the opportunity of volume inspections. The range of test frequencies starts at about 500 kHz and actually ends with the limitations of commercial EC devices at about 10 MHz. For higher frequencies network anaFig. 1: Non-centric eddy current probe for lysers can be used. Fig 2: Differential probe and CFRP inspection

sensitivity profile

434 2. MODELLING AND MEASUREMENT OF EDDY CURRENT DISTRIBUTIONS

S (x,y+dy) infinitesimal current loops

y+dy

S(x,y)

S (x+dx,y)

X

x+dx

Fig. 3: Model for the relation between current distributions and the magnetic field

In order to develop eddy current measurement techniques for CFRP it is necessary to understand the effect of the anisotropic resistance on the eddy currents. A method was developed which enables the visualisation of eddy currents in CFRP. In this method the z-component of the magnetic field was measured using a receiver coil. From the two-dimensional magnetic field distribution the current distribution can be calculated. The method bases on a formulation of the current density in form of a grid consisting of elementary current loops (Fig. 3). Each loop has a current flow with the size S(x,y). The current density J is then the result of the partial derivations ofS: ^ . . d S(x,y) J,(x,y) = — dy

. J, . d S(x,y) and J,(x,y) = — d X

Each elementary current loop produces a magnetic field which is proportional to S, The magnetic field of the current distribution is the sum of the magnetic fields of the elementary loops. The relation between S and B can be formulated as a convolution because the magnetic fields of the elementary current loops are equally shaped. The relation between the zcomponent of the magnetic field measured with a coil and S can be written as: B(x,y) = b^*S . The distribution of b^ in a plane close to the plane of the current distribution has ^^ approximately the shape of a delta function. Substituting this in Eq. (2) gives: (3) B^(x,y) = C'S(x,y) with C being a proportional constant. For this reason the current density can be evaluated directly by forming the partial derivations ofB^.

435

Experimental setup receiver coil

scan area

transmitter coil Fig. 4: Magnetic field and current distribution of unidirectional CFRP specimen

The eddy currents inside the CFRP were induced using a transmitter coil. The coil had a ferrit core with a diameter of 3 mm and a length of 5 mm. The magnetic field was measured directly above the CFPR using a small receiver coil so that Eq. (3) is valid. The magnetic field was measured with and without specimen. The difference between both measurement results gives the magnetic field of the eddy current distribution. The method was applied first to investigate current distributions in unidirectional CFRP. Fig. 4 left hand side shows the z-component of the magnetic field in form of grey values. The picture size corresponds to a scan area of 2.4 cm square 7.6 cm. Next to the magnetic field the current distribution evaluated with Eq. (1) and (3) is shown. It can be seen that the currents induced in fibre direction are dominating as expected for reasons of the anisotropic conductivity. In an additional experiment four layers of CFPR having different fibre orientation were stacked. Fig. 5 shows the measured magnetic field. The picture size corresponds to a scan area of 2.8 cm square 3 cm. The fibre orientation is clearly visible.

Experimental setup scan area

transmitter coil Fig. 5: Magnetic field of multidirectional CFRP specimen

436 3. APPLICATION TO CFRP 3.1. Detection of fibre orientation

Mil LA/ Fig. 6: Eddy current polar diagrams of unidirectional (left) and multidirectional CFRP (right)

Fig. 7: Eddy current image of bidirectional CFRP (above), Hanning masked and Fourier transformed image (below)

Fig. 6 shows the EC signal in a polar system obtained by a rotating non-centric probe [3]. Obviously, the diagram on the left represents unidirectional CFRP. The signal maximum correlates with fibre orientation. Special probes have been developed to detect the accuracy of fibre orientation in fatigue specimen up to 0.5°. The diagram on the right results from multidirectional laminate with 0, +45, 90 and -45° fibre orientation. The signal amplitude decreases with increasing distance between probe and layer. These diagrams have been recorded at 10 MHz. Higher frequencies provide sharper separation of fibre directions but require special equipment like network analysers [5] which hardly can be used for in-field inspection. The diagram in Fig. 6 on the left shows another interesting effect. Perpendicular to the fibre direction side maxima can be observed. In an ideal CPRP material, all fibres are parallel and isolated each from another. In reality, the manufacturing process leads to a certain fibre deformation and redistribution. Both the fibres within a layer and the fibres of neighbouring layers can contact each other thus causing variations in the eddy current field. This effect can be used for evaluating some matrix and bonding properties. Another way of evaluating fibre orientation is to scan a certain area with stationary differential probes. Fig. 7 presents the result of bidirectional CFRP specimen (RTM system) scanned in a 15x15 mm^ area. The pattern above results from natural inhomogeneities always occurring parallel with the fibres [4]. The image below indicates the fibre orientation very clearly. It was obtained by Hanning masking the original image and subsequent Fourier transform [6].

437 3.2. Ageing Effects The complex conductivity across the fibres changes with materials ageing. A set of specimen was thermally aged using a cycle between -160 and +120 degrees centigrade. The ageing process was interrupted after defined numbers of cycles and EC signal has been recorded. A significant growth of the side maxima during the ageing process could be found (Fig. 8). 3.3. Local Inhomogeneities To locate and analyse local inhomogeneities rotating EC probes scan the material. The m e a s u r e m e n t signal is on-line visuaUsed in a Fig-8: Eddy^current inspection of thermally aged C F ^ ^ Number of thermal cycles (from inner to outer polar coordinate system on a PC screen. Two loop): 0, 10, 200 cursors allow to choose angle positions of the probe where measurement signals should be taken. These two values whether can be recorded directly or they can be processed before recording [7]. Commonly, the ratio between the maximum and the adjacent minimum is the most suitable choice. Fig. 9 shows changing fibre orientation covered by unidirectional material twice as thick as the layer of interest. Another type of local inhomogeneity is a gap remaining between two tapes during the manufacturing. Fig. 10 demonstrates the potential of EC method to evaluate these gaps nondestructively. Although the CPRP is made from pregregs both figures clearly show varying conductivity across the fibres resulting from slightly varying fibre fraction.

Fig. 9: Eddy current detection of covered changes of orientation

Fig. 10: Eddy current detection of a gap between fibre tapes

438 3.4. Impact damage The impact damage combines fibre braking and delamination. Both inhomogeneities influence the EC signal. Fibre braking interrupts the conduction current and delamination significantly reduces the interlayer connection. Fig. 11 displays an 40x40 mm^ area including an impact not visible from the front side.

3.5. Adaptive structures Recent developments in materials science and technology try to integrate active components into the structure. Among a wide variety of possible combinations piezoceramic foils laminated into CFRP are of great interest. At first stage, the manufacturing of these structures should be optimised. Non-destructive methods help to characterise the material immediately after manufacturing and after certain static and dynamic load. High resolution eddy current methods are capable to detect, for instance, inhomogeneous fibre distribution, impact damages in fibres and actuators. Fig. 12 reflects an eddy current image of a broken piezoceramic actuator (44x30 mm^ ) under CFRP covering. Eddy current spreading bases on the thin metallisation of the piezoceramic foil which is necessary for electrical contacting.

Fig. 11: Eddy current image of an impact

Fig. 12: Eddy current image of a broken actuator in an adaptive CFRP structure

ACKNOWLEDGEMENTS This work was supported from the Deutsche Forschungsgemeinschaft and the Kultusministerium of Sachsen-Anhalt. REFERENCES [1] [2] [3] [4] [5] [6] [7]

S.N. Vernon, Materials Evaluation 46(1988)11, p. 645 S.N Vernon, J.M. Liu, Materials Evaluation 50(1992)1, p. 36 R. Lange, G. Mook, NDT&E International 27 (1994) 5, p. 241 G. Mook, R. Lange, Materialprtifung 36(1994)9, p. 345 M.P. de Goeje, K.E.D. Wapenaar, Composites 23 (1992) 3, p. 147 G. Mook, H. Heyse, J. Simonin, A. Tchemov, J. Berger, R. Lange, DACH-Jahrestagung ZerstCrungsfr. Materialpr., Lindau, 13.-15.5.1996, proc, p. 99 G. Mook, R. Lange, ICCE/2,21.-24.8.95, New Orleans, proc, p. 519

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

Thermal diffusivity

439

measurement for SiC/C compositionally graded

graphite materials O J. Nakano, K. Fujii and R. Yamada Japan Atomic Energy Research Institute, Tokai-mura, Naka-gun, Ibaraki-Ken, Japan

ABSTRACT Thermal diffusivity of oxidation-resistant SiC/C compositionally graded graphite materials has been measured by using the laser flash method. In order to study the effect of the SiC/C graded layer on the diffusivity, the thickness of the graded layer and the SiC content were changed. In addition, the specific surface areas of the SiC/C materials have been measured. It is shown that the effect of the SiC/C graded layer on thermal diffusivity was small within SiC contents (0.27-8.52 mass%) used in this study. 1.

INTRODUCTION In various fields including nuclear energy systems, graphite materials have been well used although they have a drawback of high reactivity to oxygen. To improve oxidation-resistance, SiC/C compositionally graded graphite materials with a gradual change in concentration between SiC and C have been fabricated, and have shown good oxidation-resistance and thermal shock resistance! 1-4]. Since they are used at a high temperature, it is important to study the thermal properties. This report describes the results of thermal diffusivity measurement for the SiC/C compositionally graded materials with different thicknesses of the SiC/C graded layer and SiC contents. 2. EXPERIMENTAL PROCEDURE 2 . 1 . Materials Isotropic fine-grained nuclear grade graphite (IG-110 graphite, Toyo Tanso Co., Ltd., Japan) and solid SiO of 99.9% purity (Wako Pure Chemical Industries Co., Ltd., Japan) were used for the formation of the SiC/C materials. Sample dimensions were

440

1.5, 2.0, and 3.0 mm in thickness and 10 mm fixed in diameter. The SiC/C graded layer was formed by the following reaction: 2C (s) + SiO (g) -^ SiC (s) + CO (g), where the graphite heated at 1380 °C was exposed to SiO molecules that were gasified at 1300 °C and carried by high purity helium gas. Here, all the external surfaces of the sample were graded with SiC. The formed SiC had the /? structure identified by X-ray analysis (RINT-1000, Rigaku Co., Ltd., Japan). After the above reaction, sample weights increased because of the formation of SiC. The average SiC content in the SiC/C graded layer was calculated by the following equation: M, w — w X Average SiC content (mass%) ^^ after ^^ before-xioo. (1) ^Si

^C

w.

where A^^ , A^, and M^j^ are the atomic weight of silicon, that of carbon, and the molecular weight of SiC, respectively. W^gf^^^^ and W^^^^ are the sample weights before and after the SiC/C layer formation, respectively. Fig. 1 shows wavelength-dispersive analysis of X-ray (WDAX) maps for Si for samples with 1.5 mm thickness. Si atoms penetrated a depth of about 500// m from the surface in Fig. 1 (a), and reached the center of the sample in Fig. 1 (b).

Laser Beam Direction

I

SiC/C Graded Area

Observed Area Fig. 1 X-ray maps for Si for the SiC/C graphites cross-sectionally cut (thickness: 1.5 mm); (a) SiC content 1.95% and (b) 6.87%. 2.2. Measurement of thermal diffusivity Thermal diffusivity was measured by using a laser flash equipment, PS-2000 (Rigaku Co., Ltd., Japan) which has a ruby laser of maximum 20 J in power. The sample temperature was changed from room temperature to 1320 °C . Thermal

441 diffusivity was obtained by t^/g method using the following equation: (2)

a=1.388LV7r't,

where t^/g is the time required to reach half of the total temperature rise on the back surface of the specimen and L is the sample thickness. SiC contents are 0.37-7.61% for 1.5 mm sample thickness and accordingly, 0.27-6.82% for 2.0 mm, and 0.30-8.52% for 3.0 mm. Before the thermal diffusivity measurement, the specific surface area for each sample was measured (AcceSorb2100E, Shimazu Co., Ltd., Japan) using Kr as an adsorption gas. IG-110 graphite and CVD SiC (Toshiba Denko Co., Ltd., Japan) were also used for the measurements for comparison. 3.

R E S U L T S AND DISCUSSION

3 . 1 . T h e effect of SiC c o n t e n t on t h e r m a l

diffusivity

In Fig. 2, the volumetric change and the bulk density after the formation of SiC/C graded layer are shown as a function of SiC content. It is seen that the sample volume was expanded with an increase in SiC content, but the bulk density hardly changed with SiC content. No variance of the density is reasonable when the weight gain due to the formation of SiC is taken into account. Fig. 3 shows the results of measured specific surface areas. The fact that the specific surface area increased with increasing SiC content means that the porosity in the surface layer also increased with SiC content. As shown in Figs. 2 and 3, these tendencies did not depend on the sample thickness.

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SiC content(mass%) Fig. 2 Volumetric change and bulk density after the formation of SiC/C graded layer.

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1.5 2.0 3.d • • • SiC/C graded material LQJ _ AI| 2 4 6 8 10 IG-110 Graphite

lAI

SiC content(mass%) Fig. 3 Specific surface area as a function of SiC content.

442

Fig. 4 shows the thermal diffusivities of graphite and the SiC/C samples of 1.5, 2.0, and 3.0 mm thickness as a function of SiC content. The thermal diffusivities depend on neither dimension nor SiC content. There was a concern that the thermal diffusivity would deteriorate due to an increase in specific surface area since thermal properties worsened with decreasing bulk density[5] or with increasing porosity[6]. In the case of the SiC/C samples, however, the thermal diffusivity hardly changed although the specific surface area increased with increasing SiC content. Since the bulk density was almost constant, it may be reasonably to consider that the specific surface area only gives the porosity localized at or near the surface region, which has little effect on thermal diffusivity. Thermal diffusivities of the samples of }5 10 X 3.0 mm as a function of temperature are shown in Fig. 5, where three specimens were used for each type of sample. Samples of ^ lOX 1.5 mm and j> 10 X 2.0 mm were also used, and had a similar temperature dependence. The I G l l O graphite and the SiC/C sample had almost the same values. Although the values for CVD SiC were scattered, the thermal diffusivity at room temperature was higher than those of other samples, and the literature data (0.018 - 0.694 cmVs) for the sintered SiC[7]. Since the CVD SiC sample used here had larger grain size (about 1.3 mm) and higher density (3.21 g/cm^) than sintered SiC, the heat flow induced by a laser pulse easily went through the CVD SiC with less barriers. The temperature dependence of the CVD SiC generally agrees with that of graphite materials, where the thermal diffusivity decreases with increasing temperature due to reduction of the mean free path of phonon[6].

[Thickness 1.5 2.0 IG-110 Graphite SiC/C graded material | O

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> CO 13

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--^-13-110 Graphite -^ SiC/C graded material (SiC cont.:0.3-8.5mass%) ^-CVD SiC

010X3.0 £1.5 o

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SiC content(mass%) Fig. 4 Thermal diffusivities of the samples of 1.5, 2.0, and 3.0 mm thickness as a function of SiC content.

500

1000

Temperature(*C)

Fig. 5 Thermal diffusivities of the samples of ^10X3.0 mm as a function of temperature.

1

443

3.2 Comparison of measured and calculated data When two phases exist in a specimen, the resultant thermal conductivity is calculated by the following equation[6], l+2V3,c(l-K,/Ksic)/(2K,yK3ie+l) K =K„

(3)

The resultant thermal conductivity and heat capacity are given by Eqs. (4) and (5). (4)

III

pm

* Gr

^pGr

SiC

pvSiO

(5)

where K, V, Cp, a , and p are thermal conductivity, volume -!!V-measured values {SiC:4.57 vol%) (8.52 mass% = 4.57 vol%) fraction, heat capacity, thermal •HS-caluclated values {a ) E 0.8 diffusivity, and density, respectively. o ' > 0.6 The subscripts, m, Gr, and SiC CO denote mean, graphite, and SiC, io.4 respectively. The values of these CO properties were referred from the § 0.2 0 literature[8-10]. The calculation was 0 carried out for 4.57 vol%, 1500 500 1000 Temperature(°C) corresponding to 8.52 mass%, 40 and 80 vol% for V^j^., as shown in Fig. 6 Comparison of the measured thermal diffusivities with the calculated values. Fig. 6. The calculated thermal diffusivity for 4.57 vol% agreed well with the measured value. Here, 4.57 vol% corresponds to 2.71 mol% SiC. Enweani et. al.[ll] also observed that the thermal diffusivity hardly changed for the 3 at% Si-doped graphite. From the above, the effect of the formation of SiC/C graded layer on the thermal diffusivity was small for the SiC contents used in this study. The calculations suggest a distinct effect will be observed for the case of more than 40 vol% of SiC content. 4. CONCLUSIONS Thermal diffusivity for the SiC/C materials with different thicknesses of the graded layer and SiC content has been measured by the laser flash method to study the influence of the graded layer on diffusivity. In addition, the specific surface areas for the SiC/C materials have been measured. The results obtained in this study are summarized as follows. (1) The effect of the SiC/C graded layer on thermal diffusivity was slight within

444

the range of SiC contents(0.27-8.52 mass%) used. (2) The temperature dependence of thermal diffusivity and their values themselves for the SiC/C graded materials were similar to those for graphite. (3) There was no effect of sample thickness on the thermal diffusivity when the thickness ranged from 1.5 to 3.0 mm. (4) The calculated thermal diffusivity based on the two-phase model agreed with the measured one . (5) The influence of the porosity in surface region on the thermal diffusivity was small when there was no change of bulk density.

REFERENCES [I] K. Fujii, H. Imai, S. Nomura and M. Shindo; J. Nucl. Mater., 187, 204-08, (1992). [2] K. Fujii, J. Nakano and M. Shindo; J. Nucl. Mater., 203, 10-16, (1993). [3] J. Nakano, K. Fujii and M. Shindo;]. Nucl. Mater., 217, 110-117,(1994). [4] K. Fujii, J. Nakano and M. Shindo; PROCEEDINGS at 3rd Inter. Sympo. on FGM '94, 541-547,(1994). [5] H. Matsuo, K. Fujii, H. Imai and T. Kurosawa; J. Nucl. Mater., 136, 229-237, (1985). [6] W. D. Kingery, H. K. Bowen and D. R. Uhlmann; Introduction to Ceramics, p.612-643,(1981). [7] R.Taylor; J. Appl. Phys., vol.16, 509-515, (1965). [8] D. R. StuU and H. Prophet; JANAF Thermochemical Tables, (1978). [9] H. Matsuo; Netsu Sokutei, 17(1), 2-8 (1990). [10] A. T. D. Rutland and R. J. Maggison; AEEW-R-815, (1972). [II] B. N. Enweani, J. W. Davis and A. A. Haasz; J. Nucl. Mater., 224, 245-253, (1995).

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

445

High-Temperature Ductility of TiC as Evaluated by Small Punch Testing and the Effect of CraCz Additive L. M. Zhangi J.-R Li2 R. Watanabe2 T.Hiraii 1 Institute for Materials Research, Tohoku University, Sendai, Japan 2 Department of Materials Processing, Faculty of Engineering, Tohoku University. Sendai, Japan ABSTRACT To study the mechanism by which TiC/NisAl FGMs are damaged under cyclic thermal shock, pure TiC and TiC composites with 1 to 5 vol% Cr3C2 additives were prepared by HP (Hot Pressing) at 2703 K and 1573 K, respectively, the relative densities of all the composites being above 97%. The high-temperature deformation behavior was investigated by means of a modified small punch testing (MSP) method at the critical temperatures of 1208 and 1373 K. Typical brittle fractures were found to occur in all the samples at 1208 K. When the testing temperature was increased to 1373 K, all the Cr3C2 -doped TiC composites exhibited obvious ductile deformation, such deformation behavior remaining unchanged. Compared to pure TiC, the composites had lower yield strength. The yield strength decreased and the ductile deformation increased with increasing CT3C2 content, probably due to the existence of lower charged Cr compounds in the composites. The deformation transition behavior can be applied to explain the thermal shock property of TiC/NisAl FGMs. 1. INTRODUCTION It is known that the high-temperature cyclic thermal shock resistance of thermal stress relaxation type of ceramic/metal FGMs (Functionally Graded Materials) depends on two points: 1. the thermal stress relaxation design of FGMs [1]; 2. the compatibility of the ceramic/metal system (i. e. thermal expansion coefficient. Young's modulus, Poisson's ratio, etc.) and the high-temperature mechanical properties of ceramics, including high temperature strength and ductile deformation behavior [2]. In particular, high-temperature mechanical properties of ceramics are also the intrinsic characteristics for high temperature cyclic thermal shock resistance of its FGM because they are major factors causing the damage in the ceramic phase of FGMs [2]. Their experimental determination is very important, not only to explore the damage mechanism of cyclic thermal shock of ceramic/metal FGMs, but also to facilitate the design of FGMs having better performance. Recently, a modified small punch testing (MSP) method has been developed for mechanical evaluation of ceramic and ceramic/metal composite materials used in FGMs [5]. In the present study, the MSP method was applied to evaluate TiC-based ceramics, which are employed as ceramic-side materials in a TiC/NisAl FGM. Because the TiC/NisAl FGMs apparently show different damage behavior at temperatures from 1208 to 1373 K, the fracture strength and deformation of TiC-based ceramics were measured at these two critical temperatures. The differences in the behaviors of different TiC sintered bodies and their

446 TiC ceramics with similar densities OTiC tended to be lower with increasing • CrrCs grain size, that is, it was easier for larger grained TiC to deform plastically at the same temperature. In the present experiment, there was a low charged Cr compound O (CrvCs), the melting point of which was just 1873 K [8], much lower o than that of TiC, which had an effect on the high temperature mechanical behavior of TiC ceramics. Namely, o from Fig. 4 and Fig. 5, it can be concluded that the low melting point JL of Cr-jC^ at the TiC grain boundary 60 50 softened at high temperature, which 26 / degree, CuKa caused the TiC-Cr3C2 specimen to Fig. 4 Effect of CT2C2 addition on XRD pattern of be plastically deformed at a lower hot-pressed TiC. temperature (1373 K) than that at (a) 1% CT3C2; (b) 3% CrsCa; (c) 5% CrgCz. which the pure TiC specimen was deformed. This temperature was lower with increasing Cr-jC^ additions. Otherwise, the grain size had no crucial effect on the high temperature ductile behavior although the grain size of the pure TiC was almost two times larger than that of the Cr3C2 -doped TiC. Yield strength was calculated from Equation (1), and the results are shown in Table 2. For the specimens without yield deformation, the values were fracture strength. Fracture strength (no yield point) of the pure TiC sintered at 2073 K was 269 MPa at 1208 K and 229 MPa at 1373 K, which was almost the same as the high temperature strength measured by 4-point bending by Miracle et al . [6] for their TiCo.75 sintered bodies. The — Tested at 1208 K — Tested at 1373 Kl strength of TiC-Cr3C2 sintered at 1573 K was obviously lower than that of the pure TiC sintered at 2073 K, at both 1208 K and 1373 K. For the CT3C2 -doped TiC specimens, the yield strength obtained at 1373 K decreased coincidentally with increasing Cr3C2 additions. However, it is difficult to fully explain the high temperature yield strength of TiC-Cr3C2 by their relative density [9] (see Fig. 2) and 100 150 200 250 300 grain size. From the experimental Deflection, e /fim results, it seems that Cr7C3 should be the main factor causing the variation Fig. 5 MSP load vs. deflection curves of hot-pressed in yield strength. Cracks or other damage were TiC.

JIJ

±

i

1

447 relationships with sintering temperature and additives are discussed. 2. EXPERIMENTAL 2-1. SAMPLE PREPARATION Commercial TiCo.76 and CrsCi powders, both with an average size of 5 // m, were used as raw materials. To densify TiC at a temperature of 1573 K, which is lower than the melting point (1670 K) of the metal phase (Ni3Al), three different amounts of Cr3C2 were added, such that a nearly dense TiC/NisAl FGM sample could be prepared by hot pressing. Table 1 shows the compositions of the mixtures. Table 1 The composition of samples TCO TCI TC3 TC5 TO Sample TiC (vol%) 100 99 97 95 100 Cr3C2 (vol%) 0 1 3 5 0 Hot-pressing temp. (K) 1573 1573 1573 1573 2073 Relative density (%) 72 93 97 98.1 97.1 Powder mixtures were placed in a graphite mould and then sintered at the temperatures shown in Table 1 for 2 h by hot pressing at 30 MPa in an Ar atmosphere. The sintered compacts were ground to a thickness about 0.5 nmi and then diamond-polished. They were finally cut into small round disks, 10 mm in diameter and about 0.5 mm in thickness, with a ultrasonic cutting machine. 2-2. MSP TEST Figure 1 (a) is a schematic diagram of the MSP equipment employed, and Fig. 1 (b) is the model used for calculating strength, a and b in Fig. 1 (b) represent the radii of the supporting die's hole and of the upper punch, respectively, t is the thickness of the specimens. Yield strength,
(D

(1)

puncher

guiding die 2b specimen

H (§)

2a

2a=4.5mni 2b=l.8min t = ~ 0.5mm

H

supporting die

(b) deflection-detecting prot)e

Fig. 1 Schematic drawing of Modified Small Punch (MSP) test (a) and model for strength calculation (b).

448 Where, Pmax is the maximum load in an elastic range, u is the Poisson's ratio of the tested specimen {o = 0.19) [3]. The MSP test was carried out in a vacuum of less than 10-5 torr. For comparison of load-deflection curves, real load (P) was converted to an equivalent load (Peq) to remove the thickness effect of tested specimens by the following equation: Peq=(0.5/t)2P

(2)

Specimens were heated at a rate of 10 K/min, soaked for 10 min when the predetermined peak temperature was reached, and then loaded at a crosshead speed of 0.05 nmi/min. 3. RESULTS AND DISCUSSION Figure 2 shows the variation of the relative density of sintered bodies as a function of Cr3C2 contents. TiC without Cr3C2 additive was sintered to only 72% at 1573 K, and to 97.1% even at 2073 K. Only a small amount of Cr3C2 remarkably increased the density. At the hot-pressing temperature of 1573 K, it reached 92.5% with 1 vol% Cr3C2, much higher than that of pure TiC (without Cr3C2 ). When 3 vol% Cr3C2 was added, it reached 97% and was almost as high as that of the pure TiC obtained at 2073 K. If more Cr3C2 (5 vol%) was added, the increase of relative density was not so obvious (98.1%). Figure 3 shows the cross-sectional morphologies of the pure TiC and Cr3C2-doped TiC samples. As shown in Fig. 3 (a) and (b), the pure TiC (without Cr3C2 additive) was structurally densified and developed perfectly with grain growth (average grain size about 10 fi m) when hot-pressed at 2073 K, whereas the pure TiC hot-pressed at 1573 K had a loose structure with many large pores. The grain size of latter was almost the same as that 100 of the starting material (about 5 // m). o ^ Figure 3 (c), (d) and (e) show the morphologies of TiC with 1 vol%, 3 vol% 90 and 5 vol% Cr3C2 additives at 1573 K. It gradually became densified with o:2073K increasing Cr3C2 content, and TiC with 3 ¥. 80 •:1573K vol% and 5 vol% Cr3C2 were already densified. Compared to the pure TiC •S 70 sintered at 2073 K, however, their grain (D sizes were much smaller. This shows that 60 1 1 1 1 1 1 1 1 1 1 1 1 the addition of Cr3C2 not only improves 2 4 6 8 10 12 the low-temperature-densification of TiC Content of Cr3C2 , fj / vol% but also restrains the grain growth. Figure shows the XRD patterns of TiC Fig. 2 Effect of Cr3C2 addition and sintering sintered bodies containing different temperature on relative density of TiC. amounts of Cr3C2 additives. Reflection of phases other than that of TiC could not be detected when 1 vol% Cr3C2 was added (Fig. 4 (a)). A small amount of low-charged Cr compound (Cr7C3) appeared when Cr3C2 content was increased to 3 vol% (Fig. 4 (b)). The reflection intensity of Cr7C3 increased more obviously when more Cr3C2 up to 5 vol% was added. This indicates that a small amount of Cr3C2 additive is sufficient to facilitate sintering densification of TiC ceramics, whereas low-charged Cr compounds are increasingly formed with increased Cr3C2 content. •

449

^^^m^

' ' • « » • ' •

12 ju m

a: without CrsCi (2073K); b: without Cr3C2 (1573); c: l%Cr3C2; d: 3% Cr3C2; e: 5% CT2C2. Fig. 3 Effect of CrsCi addition and sintering temperature on fracture section of TiC. Figure 5 compares the MSP load-deflection curves of the pure TiC and Cr3C2-doped TiC samples (pure TiC sintered at 2073 K, TiC with 1 vol%, 3 vol% and 5 vol% Cr3C2 sintered at 1573 K) at 1208 K and 1373 K. At 1208 K, two materials (pure TiC sintered at 2073 K and TiC-3vol% Cr3C2 sintered at 1573 K) with almost the same density basically exhibited linear deformation behavior before the final fracture. In other words, it appears that no yield deformation occurred in the above two materials at 1208 K, while at 1373 K, all the Cr3C2 doped TiC samples, except the pure TiC, showed obvious signs of yield deformation. It was found that ductile deformation increased with increasing content of Cr3C2 additive, although for purposes of clarity, whole curves are not shown in the figure. After having studied the high temperature deformation behavior of nonstoichiometric TiC by the 4-point bendingmethod, Miracle et al. [6] pointed out that its ductile deformation was mainly caused by the softening of impurities at the TiC grain boundary and the slippage of TiC grains. G. Das et al. [7], who studied the high temperature deformation behavior of nonstoichiometric TiC by the compression method, indicated that the ductile-brittle transition temperature (DBTT) of

450 Table 2 Yield strength of specimens Sample Sintering temp. (K) Yield strength 1208 K (MPa) 1373 K

TiC TiC+l%Cr3C2 2073 1573 269 / 229 150

TiC+3%Cr3C2 1573 193 129

TiC+5%Cr3C2 1573 / 113

found to occur in the CrsCi -doped TiC top side of TiC/NisAl FGMs after approximately twenty heating-cooling cycles under simulated large-temperature-difference conditions [4]. The surface temperature in that case was confirmed to be near 1573 K. The ductile deformation as revealed by the present study is apparently responsible for the hightemperature damage in the TiC/Ni3Al FGMs, according to the mechanism proposed by Kawasaki ^r tz/. [2]. 4. CONCLUSIONS AND SUMMARY (1) Highly densified TiC ceramics can be obtained by hot-pressing at 1573 K with the addition of a small amount of Cr3C2 . The densification is significantly enhanced with increasing Cr3C2, whereas the compound CryCs, which has a low melting point, is also increasingly formed. (2) High-temperature mechanical evaluation by a modified small punch (MSP) testing method indicated that no ductile deformation occurred below 1373 K for the pure TiC ceramics sintered at 2073 K and below 1208 K for the Cr3C2 -doped TiC ceramics sintered at 1573 K. The Cr3C2 -doped TiC ceramics, however, exhibited ductile deformation at 1373 K, which was likely caused by the softening of Cr7C3 at high temperature. (3) High-temperature yield strength of Cr3C2 -doped TiC ceramics decreased with increasing Cr3C2, also indicating the influence of the compound Cr7C3, which has a low melting point.

REFERENCES 1. T Hirano and K. Wakashima, Technology Materials, Vol. 38, No. 12 (1990), 26-32. 2. A. Kawasaki, A. Hibino and R. Watanabe, J. Japan Inst. Metals, Vol. 56, No. 4 (1992), 472-480. 3. L. M. Zhang, J. Li, R. Z. Yuan and T Hirai, Mater. Sci. & Eng., 203 (1995), 272-277. 4. L. M. Zhang, A. Kumagawa and T Hirai, J. Compo. Eng., (1996), in press. 5. A. Kawasaki, R. Watanabe and H. Takahashi, J. Jpn. Soc. Powder and Powder Metall. 36 (1989), 217. 6. Daniel B. Miracle and Harry A. Lipsitt. J. Am. Ceram. Soc, Vol. 66, No. 8, (1983), 592599. 7. G. Das, K. S. Mazdiyasni and H. A. Lipsitt. J. Am. Ceram. Soc, Vol. 65, No. 2, (1982), 104-110. 8. H. J. Goldschmid: Interstitial Alloys, Butterworths, London, 1967. 9. W. D. Kingery, H. K. Bowen and D. R. Uhlmann, Introduction to Ceramics. Second Edition, by John Wiley & Sons, Inc. 1976,472-494.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

451

Mechanical and Thermal Properties of PSZ/Ni-Base Superalloy Composite Shinya Akama 4th Research Center, Technical Research & Development Institute Japan Defense Agency 2-9-54, Fuchinobe, Sagamihara, Kanagawa 229, Japan ABSTRACT Mechanical and thermal properties of non-graded Partially Stabilized Zirconia (PSZ)/IN100 (Ni-Base Superalloy) composite materials were investigated at elevated temperature. The specimens were made by sintering method. The following results were obtained by 4-point bending tests. The yield strength of these uniformly composite materials abruptly decreases at temperatures over 1123 K when the concentration of PSZ in weight is under 5 0 % . The specific yield strength (yield stress/density) of composites with more than 75 % PSZ is nearly equal to that of the CMSX-2 single crystal which is recently used for the turbine blades of the aero jet engines in the temperature range from 1173 K to 1273K. 1.INTRODUCTION High engine performance and compactness are strongly required to gas turbine engines; therefore it is desired to apply new materials in order to increase turbine inlet temperature. Functionally graded materials, which are composed of ceramics and metals and the concentration of each element changes gradually from one side to another, have much potential for the application in the field of aerospace industries because of their thermal stress relaxing function. PSZ is one of the most popular refractory materials and used as a thermal barrier coating material in combustion chamber and/or turbine vanes of jet engines. INIOO is also used widely as a material of turbine blades and vanes. For this reason, FGM of PSZ/INIOO was selected. In order to design the optimum structure of FGM for stress-relaxation, mechanical and thermal property data of each layer at the practical used temperature are greatly needed, but in the recent work on FGM of PSZ/IN 100, most of such data are obtained only at room temperature ^^' ^^ . The present paper describes results of the investigation on hightemperature mechanical and thermal properties of uniformly composite

452 materials of PSZ/INIOO made by a sintering method. 2. EXPERIMENTAL Uniformly composite materials are produced by a sintering method. In the sintering method, average diameters of PSZ and INIOO powders are 0. 8 //m and 25. 0//m respectively. As the sintering condition, specimens are made by hot pressing at 1423K under the pressure of 196MPa. The concentration of PSZ and INIOO was chosen as shown in Table 1. Table 1 Concentration of PSZ and INIOO (wt«)

PSZ

0

25

36

50

62

75

100

INIOO

100

75

64

50

38

25

0

Two types of specimens were used in this examination. One is a diskshaped, 10 mm in diameter and 2 mm in thickness that is used to obtain the thermal conductivity. The other one is a rectangular-shaped, 4 mm in width, 3 mm in thickness and 40 mm in length that is used to examine the thermal expansion coefficient and bending strength. 3. RESULT AND DISCUSSION 3. 1 Thermal Expansion Coefficient The linear thermal expansion per change in temperature were measured at temperature range from 293K to 1473K used dilatometer for each composite respectively. The mean coefficient of linear thermal expansion (a) is represented by

1

AL/AT

The thermal expansion coefficients from 293K to are plotted in Fig. 1.

573K and to 1173K

3. 2 Thermal Conductivity Coefficient The thermal diffusivities of each composite were measured at 673K, 1073K, 1273K, 1373K and 1473K by laser-flash method. The thermal conductivity coefficient, the product of thermal diffusivity, specific heat and density, are plotted as a function of the concentration of PSZ and test temperature in Fig. 2 and Fig. 3. In Fig. 2, the thermal conductivity coefficient at room temperature is plotted simultaneously ^^. From

453 Fig. 2 and 3, it was recognized that the thermal conductivity coefficient became stronger temperature dependence with decrease the concentration of PSZ. 3. 3 Bending strength 4-point bending tests were carried out using Instron-type testing machine in vacuum pressure of about 8 xlO "^ Torr at temperature range from 1073K to 1373K. Load-deflection curves obtained in elevated temperature are classified into four groups as shown in Fig. 4. D : deflection is large and yield point is clear. D^: deflection is very large and yield point is not clear. F : fracture behaviour occur immediately after yield. B : fracture behaviour occur before yield. According to this classification, the change of the form of loaddeflection-curve with test temperature is indicated using these symbols as shown in Fig. 5. The form of the flow curve in same PSZ/INIOO composites changed B ^ F ' = > D ^ D ^ as the temperature became higher. The form of the flow curve in same temperature changed D^ t = > D ^ F ^ B as the concentration of PSZ became higher. From the load-deflection curves, the yield strength of each composite was obtained in elevated temperature as shown in Fig.6. It was found that the yield strength in the PSZ/INIOO composites became lower as the temperature became higher but the relationships between yield strength and temperature can be distinguished into two groups. Namely, where the concentration of PSZ is under 50%, the yield strength abruptly decreases over 1123 K. Over lb% PSZ, that slowly decreases. The yield strength is one of the most important design parameters where "no plastic deformation" is design criterion. But in case of the rotating parts, the specific yield strength (yield strength/density) becomes one of the most important design parameters. Fig. 7 shows the relation between the specific yield strength and temperature. In this picture, the specific yield strength of the Single Crystal (CMSX-2) which is recently used for the turbine blades of the aero jet engines is plotted simultaneously. It is recognized from Fig.7 that the specific yield strength of composites with more than 15% PSZ is nearly equal to that of the CMSX-2 in the temperature range from 1173 K to 1273K.

4.CONCLUSIONS 1.The relationships between the thermal expansion coefficient and the concentration of PSZ and the thermal conductivity coefficient and the concentration of PSZ at elevated temperature were obtained. 2. 4-point bending tests were carried out at temperature range from 1073K to 1373K, the change of the form of Load-deflection curve was obtained.

454 3.The yield strength of composites, where the concentration of PSZ is under 50^, abruptly decreases over 1123K. 4. The specific yield strength is nearly equal to that of the Single Crystal (CMSX-2) which is recently used for the turbine blades of the aero jet engines, where the concentration of PSZ is over 75^. ACKNOWLEDGEMENT The author wish to thank Mr. Y. Tada and Dr. R. Ishikawa of National Aerospace Laboratory, Science and Technology Agency, JAPAN for their contribution to the bending test program. REFERRENCE 1.M. Yuuki, H. Nakanishi et al, Proc. of FGM'93 symposium of Functionally Graded Materials Forum. 47-50, (1993) 2. M. Yuuki, H. Nakanishi et al, Proc. of the 22th annual conference of Gas Turbine Society Japan. 185-189, (1994)

CONCENTRATION OF PSZ

wt%

Fig. 1 Thermal expansion coefficient of PSZ/INIOO composites

455 ^ 30 CONCENTRATION OF PSZ

2

CONCENTRATION OF PSZ

20

673

WT%

-—-—-j—-1-

©

1 CO Q_

^

2

1

E

®

1

50 1

©

1

—-i-©---H— • ® © ^ I

U_ CD

^

1273 1473

Fig. 3 Relation of thermal conductivity coefficient and temperature

N9

L/;

1073

TEMPERATURE

Fig. 2 Thermal conductivity coefficient of PSZ/INIOO composites

®

873

1— =2: CJ5

1

1

1

-—(i>--i1

1

1

1

CD

" 100 .

Fig. 4 Classification of loaddeflection curve

-—[—\—^®
1173 1273

1373

TEMPERATURE Fig. 5 Distribution of form of load-deflection curve

456

a 800 \

CONCENTRATION OF PSZ 0 o 25 A 36 D 50 o 75 • 100 •

600

CO CO I— CO

400 200 1

1073

1

1173 1273 TEMPERATURE

1

1373K

Fig. 6 Yield strength of PSZ/INIOO composites

CO

c:>

1073

1173 1273 TEMPERATURE

1373K

Fig. 7 Specific yield strength of PSZ/INIOO composites

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

457

Processing-working stress unified analysis model and optimum design of ceramic-metal functionally graded materials Peng-Cheng Zhai, Qing-Jie Zhang and Run-Zhang Yuan State Key Laboratory of Materials Synthesis and Processing Wuhan University of Technology, Wuhan 430070, P.R.China

The processing-working unified analysis model and optimum design of ceramic-metal functionally graded materials are studied. Emphasis is placed on the effect of the residual stress on the working stress and optimum design of the materials. Two analysis models for the residual stress effect are examined: one is a separate analysis model and the other is an unified analysis model. It is indicated that the effect of the residual stress is significant and the optimum design of the FGMs should be based on the processing-working unified analysis model when the materials response is inelastic. 1.INTRODUCTION The stress analysis of ceramic-metal FGMs constitudes a basis for the materials optimum design, either in the processing process or in the working process. There have been many studies on the residual stress and the optimum design process of the materials, for instance Ref [1-3], and some studies on the working stress and the optimum design process, such as Ref [4-5], but few studies have been reported to deal with the processing-working stress unified analysis. Therefore, the optimum design for the working process considering the residual stress effect is not well established at present. The present study concerns the working stress and optimum design of ceramic-metal FGMs considering the residual stress effect. The residual stress effect is examined with two analysis models: one is a so-called separate analysis model and the other is an unified analysis model. The difference between the two analysis models is discussed. When the response of the materials is inelastic, the separate analysis model is proved to be unsuitable and the unified analysis model is proposed for the processing-working stress analysis and the optimum design of the compositional gradation. 2.M0DEL DESCRIPTION Consider a ceramic-metal FGM which undergoes a thermal load history as shown in figure 1. The thermal loading history consists of two phases: phase I corresponds to the cooling after sintering(processing phase), the stress in this phase is called residual stress; phase II

458 corresponds to the working phase and the stress is called working stress. There are two methods that can be used t6 treat the residual stress effect on the working stress: one is so-called separate analysis model and the other is an unified analysis model. In the separate analysis model, the residual stress and working stress is analyzed separately and I he resulted stress is obtained from a direct superimposition of the two separate ones. It is obvious that this method is correct only when the two responses are linear and the material remains elastic in the two phases. For the unified analysis model, the processing phase and the working phase are treated as an unified loading process and the resulted stress is obtained through an unified analysis for the two phases. l.RESULTS AND DISCUSSION A TiC/Ni FGM is used as an example. The material properties, geometrical sizes and compositional function were given in Ref [6]. The thermal load in the working phase is thermal a shock with heat flux magnitude q = 5 AffV/m^ and duration ^^ = 2.0^. The sintering temperature in the phase I is taken as 1300K. The effect of the residual stress is considered by comparing the results of the separate iiuilysis of phase II(which abandoned the residual stress effect) and unified analysis of phase I mill II. Numerical results are indicated in figure 2. It is obvious that the effect of the residual stress is significant. For the case of not considering the residual stress effect, the stress is compressive during the heating process and decreases with the time increases. During the unloading process, the stress becomes tensile and then decreases with the time increases and becomes compressive again. The maximum tensile stress at the ceramic surface is about 7M)lVIPa. For the case of considering the residual stress effect, the stress is compressive during I lie heating process and increases with the time increases. During the unloading process, the stress becomes tensile and remains tensile. The maximum tensile stress at the ceramic surface is about 350MPa. It is suggested that the residual stress exerts significant effect on the working stress and it will reduce the maximum tensile stress at the ceramic surface obviously. Figures 3(a) and (b) compare the results from the separate analysis model and the unified analysis model. Figure 3(a) corresponds to the elastic analysis and it is seen that the two models give the same resulted stress. Figure 3(b) corresponds to the elastic-plastic analysis \\H\ it is seen that the results found from the two models are different. This is because in the eListic analysis the two responses in the phase I and II are both linear and the resulted response can be obtained from a direct superimposition of the two separate ones, but in the I'liLstic-plastic analysis, the two separate responses are both nonlinear and the separate analysis model based on the direct superimposition of the two separate responses is no longer suitable. Figure 4(a) and (b) give an important relationship for the optimum design of the FGM: the relationship between the maximum tensile stress at the ceramic surface and the compositional L'.xponent. Figure 4(a) shows the effect of the residual stress on the optimum design. From figure 4(a), it is seen that the effect of residual stress is very significant. When the effect of the residual stress is considered, the maximum tensile stress at the ceramic surface almost

459 decreases monotonously with the compositional exponent increases but if the effect of the n'sidual stress is abandoned, the maximum tensile stress has a minimum value at P = 0.6 . It is suggested that the effect of the residual stress on the optimum design of the FGMs can not l)c igonred. Figure 4(b) compares the results from the separate analysis model and unified analysis model. From figure 4(b), there exist some difference between the two analysis models.. 4.CONCLUSIONS The thermo-elastic-plastic response and optimum design of ceramic-metal FGMs under thermal shock loading are studied. An unified analysis model is used to consider the effect of residual stress on the working stress and optimum design of the FGMs. In this model, the cooling process and working process is treated as an unified thermal loading process. A TiCNl FGM is taken as a numerical example. Emphasis has been put on two aspects: the effect of (he residual stress on the working stress and the optimum design of the graded composition. The main conclusions obtained are as follows. i 11 fhe effect of the residual stress on the working stress and optimum design of FGMs is very significant. It not only changes the stress history but also changes the magnitude of the tensile stress at the ceramic surface. If this effect is abandoned, it will overestimate the tensile stress at the ceramic surface. (2) In the elastic analysis, the separate analysis model and the unified analysis model give the same result because the two responses of the cooling and working phases are both linear. In this case, the direct superimposition is suitable. I}) In the elastic-plastic analysis, the separate analysis model and the unified analysis model give the different results because the two responses of the cooling and working phases are both nonlinear. In this case, the separate analysis model is no longer suitable and the unified analysis model should be used..

\ ( KN OWLEDGEMENTS This work was supported by the National Science Foundation REFERENCES I

R.L.Williamson and B.H.Rabin, Proc. 2nd Int. Symp. on FGMs, 1992, 55-66

> ^ 4. 5. 6.

R.L.WilHamson and B.H.Rabin, Proc. 3rd Int. Symp. on FGMs, 1994, 215-222 N.Cherradi, P.Moechli and K.Dollmeier, Proc. 3rd Int. Symp. on FGMs, 1994, 253-258 J.Teraki, T.Hirano and K.Wakashima, Proc. 2nd Int. Symp. on FGMs, 1992, 67-74 T.Maruyama, M.Harada, etal., Proc. 2nd Int. Symp. on FGMs, 1992, 433-440 P.C.Zhai, Q.J.Zhang and Y.Z.Yuan, Acta Mechaica Solida Sinica, vol.9, 1996,in press

460 T(K)

time

phase I

phase n

Figure 1 : Thermal loading history of ceramic-metal FGMs phase I: cooling phase after sintering phase II: working phase o o o

• working stress • resulted stress o o

o o

o o o

Figure 2 : Influence of processing stress(residual stress) on working stress (P=1.0)

461 o o (N

1 1

t

^

^s-

time (s) \. '

2

'

'

3^*-»..^ 4

5

• o o (N

—^r- unified analysis model -•— separate analysis model

O

o

a: elastic case O O

o o

O

4

-s

time(s) 5

O O (N

O O

• unified analysis model • seperate analysis model

o o

b: elastic-plastic case Figure 3: processing-working stress obtained from the unified analysis model and separate analysis model (P=1.6)

462

0.4

0.6

0.8 1 1.2 compositional exponent

1.4

1.6

1.4

1.6

(a) effect of the residual stress • : without the effect of the residual stress A : with the effect of the residual stress o o

o o

(D

O

<4

o o

o o

0.4

0.6

0.8 1 1.2 compositional exponent (b) effect ofthe analysis models • : the separate analysis model • : the unified analysis model

Figure 4 : determination ofthe optimum compositional gradation ofthe FGMs

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

463

Evaluation test of C/C composites coated with SiC/C FGM, under simulated condition for aerospace application Y.Wakamatsu% T.Saito^ F.Ono", K.Ishida\ T.Matsuzaki^ O.Hamamura^ Y.Sohda^andY.Kude' ^Kakuda Research Center, National Aerospace Laboratory, Kakuda, Miyagi, 981-15, Japan ^National Aerospace Laboratory, Chofu, Tokyo 182, Japan '^Advanced Engineering Service Company, Ltd., Tukuba, Ibaraki, 305, Japan "^Central Technical Research Laboratory, Nippon Oil Company, Ltd., Naka-ku, Yokohama, 231, Japan The authors fabricated carbon/carbon composites coated with SiC/C functionally graded material (FGM). The fabricated specimens were hemispherical in shape with a diameter of 50 mm. They were tested in two kinds of heating devices to evaluate their durability. In hot gas flow tests, specimens were heated seven times for 30 seconds each before damage occurred. In arc-heated wind tunnel tests, a specimen was heated twice, 1,100 seconds each time, before damage appeared. Pinhole-like damage was observed on the surface of all the specimens and pores were found in the carbon under the SiC coating layer by means of X-ray CT. The results of these evaluation tests demonstrated that the FGM layer has reasonable durability although the SiC coating is subject to damage due to oxidation.

1. INTRODUCTION C/C composite (CCC) shows promise as a material for space vehicles or engines because of its low density and high strength at high temperature. The greatest problem involved in the practical use of CCC under atmospheric conditions is the lack of oxidation resistance, and thus the surface of the composite is usually coated with SiC to prevent oxidation. However, this SiC layer is apt to spall off at elevated temperatures because of the difference of the thermal expansion coefficient between SiC and CCC. In an effort to reduce this difference and the consequent spalling, in a previous

464 study, the authors fabricated two hemispherical CCC specimens with SiC coating, the interface consisting of a SiC/C FGM layer to relax thermal stress between the substrate and the coating. They conducted heating tests twice for twenty seconds each under conditions simulating those of an aerospace-plane by using a hot gas flow device. Results showed no significant changes on the specimens, and the weight loss after the test was negligible [1]. In this study, the authors fabricated similar specimens and evaluated them under conditions of repeated heating and heating of long duration. 2. SPECIMENS The specimens were of hemispherical shape with a diameter of 50 mm. Three SiC 100A6m FGM 30Mm specimens (A,B,C) were fabricated. CCC was fabricated as follows. A threeC/C 2.5mm dimensional quasi-net-shape fabric of pitch-based carbon fibers was formed. It was repeatedly penetrated with pitch and was baked at 2,000 K to increase the apparent density. SiC/C FGM was Figure 1. Configuration of specimens. formed on the substrate by the chemical vapor deposition (CVD) process before the surface was layered with a pure SiC layer by the CVD method. Details of the specimens are described in a reference 1.

3. EVALUATION TEST DEVICE AND PROCEDURE Major simulation parameters for aerospace applications include heat flux and flow dynamics. Radiation equilibrium temperature is an alternative parameter to simulate the heat flux. Parameters simulating the flow dynamics include flow velocity, dynamic pressure, etc.. Two heating devices were used to evaluate the durability of specimens. One was a hot gas flow device and the other was an arc-heated wind tunnel. In the former device, to obtain the strong effect of thermal stress, repeated brief periods of heating under conditions of high heat flux and high dynamic pressure were imposed. In the latter device, to obtain the strong oxidizing effect, chemically active heating for a long duration under conditions of high heat flux and high velocity was imposed. The maximum temperature of the specimens was observed at the top of the nose due to the fluid dynamic conditions. The target temperature was set at 1,800 K to maintain the margin at the melting point of SiOg, the oxidation resistant layer produced from SiC. 3.1. Hot Gas Flow Device This device generates a simulated hot air flow by mixing and burning hydrogen

465 gas and a mixture of oxygen and nitrogen. To simulate the air flow, the mixture ratio of gases was determined so as to keep the content of oxygen within the combustion gas at a mole fraction of about 21 %. The temperature of combustion gas could be varied by adjusting the mixture ratio of the reactants. Water vapor was always included within the gas flow because hot gas is generated by the combustion of hydrogen. Details of the combustor can be found in a reference 2. A specimen was fixed to a water cooled support just downstream of the nozzle of the combustor. When the device was Table 1 operated under the conditions shown in Typical operation of hot gas flow Table 1, the top of the specimen was 0.83 (kg/s) subjected to a temperature of 1,810 K. Mass Flow Rate of N 2 Mass Flow Rate of 0 2 0.74 (kg/s) The surface temperature of the top was 0.04 (kg/s) measured by a radiation pyrometer. The Mass Flow Rate of H 2 specimen was heated for 30 seconds in Combustion Chamber Pressure 0.7 (MPa) each test and kept at equilibrium tem- Combustion Gas Temperature 2000 (K) perature for 20 seconds. It was then mod- Ambient Pressure 0.1 (MPa) erately cooled by the nitrogen gas flow Mole Fraction of Oxygen 20.6 (%) for about 30 seconds. 3.2. Arc-Heated Wind Tunnel The air was heated by using an electric arc heater. The enthalpy of the air flow was controlled by adjusting the arc current. Molecular oxygen easily dissociates into atomic oxygen under arc discharge conditions. Therefore, the air flow includes active atomic oxygen and is highly oxidative if the flow is frozen in the nozzle. The nozzle and the specimen were set in a low pressure test cell and the flow was evacuated by a vacuum pump. The specimen and the heat flux meter Table 2 were fixed to a revolving arm. To obtain T5T)ical operation of arc w i n d t u n n e l the required temperature of the spec18.2 (g/s) imen, the air mass flow and the arc cur- Mass Flow Rate of Air 1036 (V) rent were adjusted. When the tunnel was Voltage of Arc operated under the conditions shown in Current of Arc 700 (A) Table 2, the top of the specimen was Electric Input Power 724 (kW) heated up to 1,840 K. The surface temChamber Pressure of Heater 100 (kPa) perature was measured using a thermo170 (Pa) camera. The specimen was heated for Chamber Pressure of Test Cell 1,100 seconds in each test.

4. RESULTS OF HEATING TESTS 4.1. General Results Although all specimens were eventually damaged by corrosion, no spalling of the CVD-SiC layer from the FGM layer due to thermal stress occurred in any of the specimens. Microscopic examination showed an absence of a tendency for cracks to increase on the surface. These results suggest that the FGM layer

466 achieved thermal stress relaxation of the interface between SiC and CCC and demonstrated reasonable durability within the test range of heating duration and number of times heated. A feature common to all the specimens was the appearance of spots. White spots with a diameter from about 0.5 to 1 mm appeared on the center region of the nose on the heated side. Violet spots with a similar diameter were observed on the back side. Those spots appeared in the first test and increased with subsequent tests. 4.2. Hot Gas Flow Tests Preliminary tests were conducted five times using specimen A. In the fifth test, corrosion occurred because of overheating due to operational errors. Regular tests were conducted seven times using specimen B. Measured temperature of the nose varied from 1,770 K to 1,840 K. After the seventh test, pinholelike damage was found as shown in Figure 2. Total heating time was 210 seconds.

Figure 2. Specimen B after tests.

4.3. Arc-Heated Wind Tunnel Tests Two heating tests of specimen C were conducted. The surface temperature was 1,840 K in the first test and 1,870 K in the second. At 1,000 seconds after initiation of heating in the second test, surface temperature and the brightness of the specFigure 3. Specimen C after tests, imen abruptly increased. After the second test, corrosion and pinholes were observed as shown in Figure 3. This damage was thought to have been caused when the fluctuation of temperature was observed. Heating time amounted to 2,100 seconds until occurrence of damage. 5. DISCUSSION The histories of mass change rate of Specimen B and C are shown in Figure 4. The air flow of the arc wind tunnel was expected to be chemically active. However, the mass reduction rate for the arc-heated wind tunnel test was unexpectedly smaller than that for the hot gas flow test. This can possibly be attributed to the ambient pressure, because the ambient pressure for the hot gas flow test was

467 about six hundred times greater than 0.0000 1 1 1 1 that for the arc-heated wind tunnel test. 9 0 ! Severely damaged specimens A and O Arc Wind Tunnel C had a common feature without regard . -0.0002 L i i L D Hot Gas Flow ..J i l l to the heating method. Pinhole-like dam1 9 4 i l l age was also observed in the case of spec^ -0.0004 tp \ ^ \ "1 imens A and C as in the case of specimen B. From these observations, it is speculated that corrosion was initiated with I -0.0006 r j T i i m : i 1 the pinhole-like damage on the SiC layer. In X-ray CT photographs of specimens -0.0008 K B and C, Figures 5 and 6, pores can be i i 1 i i i observed behind the pinhole-like dami i i i i i 1 age. As to the relationship between pores -0.0010 1 i 2 3 4 5 6 and pinholes, the probable cause of pinNumber of Times Heated hole damage in high temperature air flow is a discharge of SiO produced by Figure 4. Variation of mass change rate. I

Figure 5. X-ray CT of specimen B.

I

I

Figure 6. X-ray CT of specimen C.

active oxidation or a discharge of melted or vaporized SiOg produced by passive oxidation. Figure 7 shows the equilibrium composition of products in the reaction of SiC with air at 100 kPa. It shows that solid phase SiOa does not become liquid phase SiOg, nor gas phase SiOg and SiO below 2000 K. This means that solid SiOg is lost only when it is heated to above 2000 K. Therefore, one possible cause of the pinhole-like damage is local overheating of material. If a pore is generated on the surface of CCC in the process of production, it may create thermal resistance and prevent thermal diffusion from the heated surface to the cooler side, thus causing overheating of the CVD-SiC layer and generating a pinhole. If heating continues, the size of the pore increases. Oxygen diffuses into the interior carbon through the pore if there is pinhole

468 damage at the CVD-SiC layer by any 0.12 reason, even though there are no pores 0.1 after fabrication. Then the pore is P=100l<Pa I 0.08 formed due to the oxidation of carbon -e-SiO (G) and increases in size. vt 0.06 -e-Si02(S) Present evaluation tests were cono 0.04 -^-Si02(L) -•X"Si02(G) ducted from the viewpoint of applica0.02 tion of the CCC in the field of aerospace technology. It was clarified that good 0 1000 1500 2000 500 joining durability was feasible if FGM Temperature (K) technology was utilized to form the SiC coating on the CCC. At the same Figure 7. Equilibrium composition. time, the importance of understanding the behavior of the SiC layer itself under the condition of high temperature and high speed flow is suggested. These factors should be further clarified by a different approach.

6. CONCLUSION Heating evaluation tests of CCC coated with SiC/C FGM were conducted under conditions simulating those in aerospace applications. The findings obtained were as follows. 1. Reasonable durability of the FGM layer between the SiC layer and the CCC was shown within the experimental range, evaluated. 2. Pinhole-like damage and corrosion were eventually observed. The mechanism of pinhole generation and the oxidation behavior of the SiC should be further clarified. 7. ACKNOWLEDGMENT The authors would like to express sincere appreciation to the late Prof. H. Takahashi of Tohoku University, and Mr. Y. Watanabe and Mr. T. Kanda of NAL for their invaluable advice. They also appreciate to the cooperation of Mr. T. Chiyokubo of Miyagi Prefectural Industrial and Technical Center and Mr. K. Aida of Japan Philips Ltd..

REFERENCES 1. Y. Sohda, Y. Kude, T. Saitoh, Y. Wakamatsu and M. Niino, Ceramic Trans., The American Ceramic Soc, Vol.34 (1993) 125. 2. Y. Wakamasu, S. Ueda, T. Saito, M. Niino, T. Nomura, K. Kosaka, T. Saito and S. Kiyoto, Ceramic Trans., The American Ceramic Soc, Vol.34 (1993) 263.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

469

Durability and High Altitude Performance Tests of Regenaratively Cooled Thrust Engine Made of Zr02 /Ni Functionally Graded Materials Yukio Kuroda^, Makoto Tadano^, Akio Moro^, Yoshihiro Kawamata"^, and Nobuyuki Shimoda^ ^Kakuda Research Center, National Aerospace Laboratory 1, Koganezawa, Kakuda, Miyagi 98915, JAPAN ^Mitsubishi Heavy Industries, LTD. 1200, Higashi-Tanaka, Komaki-Shi, Aichi-Ken 485, JAPAN ^Nippon Steal Corporation. 20-1, Shintomi Futtsu Chiba-Ken, 299-12, JAPAN Durability and high altitude performance tests of regeneratively cooled 1200 N thrust engine composed of two kinds of Zr02/Ni FGM chambers were conducted with NTO/MMH propellant to evaluate the real engine performance of Ispv,, C* efficiency of the engine and also to obtain thermal data of the engine. It was shown that a high level of combustion efficiency can be achieved in the operating test conditions. In the high altitude performance tests, the vacuum specific impulse Ispv was 318 s at Pc=1.3 MPa. 1. INTRODUCTION Ceramic thermal barrier coating systems will be of growing importance for the reusable high performance orbiting maneuvering system (QMS) engine. High performance thrust chamber, however, coupled with extended reuse requirements impose difficult cooling requirements particularly at throat region. To meet the cooling and life enhance requirements, the inner wall of the thrust chamber are fabricated with thermal barrier, such as ceramic coating to reduce the large wall temperature difference. One promising method of improving adhesion of the ceramic coating to the metal wall in the thrust chamber is to apply the functionally graded materials (FGM) composed of Zr02 and nickel (Ni) which can withstand the lower thermal expansion of the Zr02 thermal barrier coating and is characterized by high thermal conductivity. In this study, durability and high altitude performance tests (HAPT) of the two kinds of Zr02/Ni FGM chambers were conducted with NTO/MMH propellant. 2. EXPERIMENTAL THRUST CHAMBER AND TEST FACILITY The thrust chamber assembly used in this test series is shown in Figure 1. It consists of a core injector, two kinds of combustion chambers, i.e., one is perfect Zr02/Ni FGM chamber with a high area ratio nozzle for high altitude performance tests and the other is the pertial Zr02/Ni FGM chamber for sea level durability tests.

260

O^

(Unit: mm)

Figure 1. Thrust chamber

470 2.1 Injector The basic configuration of the injector was an unlike quadlet elements without acoustic cavities. The film cooling fraction were reduced to zero percent of the total fuel flow rate to obtain the high performance. Core elements were designed so that Rupe's optimum spray mixing condition, defined by Eq. (1)=0.5, was obtained at an overall mixture ratio of 1.65 for the NTO/MMH bipropellant. RN=

(1)

l/(l+PoUo^do/pfUf^df)

p is the density, u is the velocity through injector orifice, d is the orifice diameter. Subscripts o and f represent the oxidizer MTO and fuel MMH respectively. 2.2 Regeneratively Cooled Zr02/Ni FGM Chamber for HAPT The thrust level of the experimental engine was 1200 N at a chamber pressure of 1.0 MPa. Figure 2 shows the engine set in an altitude capsul. The nozzle area ratio of 300:1 was used for HAPT. To increase coolant speed and pressure drop, the width of the channel varied from 2.0 mm at the cylindrical section to 0.9 mm at the throat section. Figure 3 shows the outer cylinder cross section of the chamber. The combustion chamber used in HAPT was composed of perfect Zr02/Ni FGM, i.e., chamber inner wall was made of Zr02 (100 vol %), coolmg side was made of pure Ni and intermediate material is made of ZrC)2/Ni FGM. The inner wall of the overall section was constant to 3.1 mm in thickness. Calculated temperature and heat absorption distribution of the chamber is shown in Figure 4 for each cross sectional area under the conditions of Pc=1.0 MPa and MR=1.65. The heat absorption qc at the throat section was 4.3 MW/m^ which was about 5 times compared to that of cylindrical section of the chamber. Twgs* Inner surface temperature of t h e Z r 0 2 10® vol % T b i : Temperature between Z r O i 100 vol and FGM Twci Cooling wall temperature 1

2400

1

^1600

Figure 2. Engine set in an altitude capsul Electro-formed Nickel Ooseout Wall

SUS304LCloseoutWaIl

^

0.7 mm

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400

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0 1 0 L (cm) I Throat Chamber profile

0 L (cm)

30 Coolant Channels p-^^^^ 4 Calculated tempemture and heat absorption distribution of the chamber Figure 3. Outer cylinder cross section of the chamber

471 3. EXPERIMENTAL RESULTS 3.1 Estimated Precision of Performance Parameters The primary performance parameter is the vacuum specific impulse Ispv which calculated using equation (2), I spv = Fv /nit C * = ( Pc , inj fp) (A t fdis ) / mt

was (2) (3)

TlC* = C*/C*oDE (4) where Fy is a vacuum thrust, mt is a total propellant flow rate. The characteristic velocity C* was calculated using measured chamber pressure, defied by Pc,inj, at the injector face, m^ and measured geometrical throat area A^. In equation (3), fp is the correction coefficient to convert Pc,inj to nozzle stagnation pressure, and f^^jg is the discharge coefficient. C efficiency (r]C ) is also used to indicate relative measure of combustion efficiency. In equation (4), C ODE is the theoretical characteristic velocity for one-dimensional equilibrium flow. The precision of measured Fy can be estimated to be within ±0.12 %. The precision of the calibration factors for the turbine-flow meters were generally within ±0.25 % at a nominal flow rate, and the long term variation of K-factors were within ±0.2 %. The overall precision of the measured I^py , is calculated to be within±0.4 %. 3.2 Thermal Characteristics Figure 5 shows the timewise variation of throat temperatures T T H 1 , 2 and that of coolant outlet temperature TcOF. Thermocouples were embeded at near the cooling channel of the chamber. It is seen that the temperatures at the throat section reached steady values at 100 s, whereas those at the outlet of the coolant increased sligtly even after 100 s of firing. This may be due to the heat conduction from the non activity-cooled regions near the upstream of the cooling outlet manifold. The coolant outlet temperature difference between 100 s and 150 s was less than 1.5 K. Figure 6 summarizes the temperature increase through the cooling channels and the total heat load Q^. The outlet temperature increased with mixture ratio as shown in Figure 6. The increase of outlet temperature may be attributed to the decrease of the cooling fuel flow rate. The temperature increase through the cooling clannels with respect to P^ is very small. The smaller temperature increase is very important for prevension of possible fuel decomposition on the

1.2 20

40

60

80

100

1.4

1.6

1.8

2

2.2

120

Figure 5. t (s) Timewise variation of throat temperatures Figure 6. Temperature increase through the cooling and that of coolant outlet temperature channels and the total heat load Qc

472 nozzle throat as well as to obtain a high injector performance precluding reaction stream separation which will degrade the engine performance. 3.3 Durability Tests In this test series, partial ZrCb/Ni FGM chamber with an expansion area ratio of 11.3:1 and the perfect Zr02/Ni FGM chamber were used.(^) For the pertial Zr02/Ni FGM chamber, the inner wall of the cylindrical section was 4.0 mm in thickness and that of the throat section was 2.0 mm in thickness. FGM section in this chamber is composed of ZrCb (24.5 vol %)/Ni (75.5 vol %). The inner surface of the chamber was covered with Zr02 (100 vol %). The heat absorption at the throat section was calculated to be about 8 times compared to that of cylindrical section of the chamber under the conditions of Pc=1.0 MPa and MR=1.65. The thruster underwent 260 cycle tests at sea level, a total of 2780 s of firing time. Figure 7 (a) shows the C and r]C vs test cycles for the engine operating at the design point chamber pressure of 1.0 MPa and MR=1.65 where the C had a peak. Slight increase in C and r\C can be seen to about 100 cycles with respect to an increase in the test cycles. This may be due to an increase in the fuel injector manifold temperature caused by the spalling of the Zr02 (100 vol %) at inner wall surface of the chamber. After 100 cycles, a small amount of decrease in r]C can be seen tol50 cycles and r\C was settled to 97 % after 150 cycles. This is presumably due to the decrease in injector performance in the cycle tests. In the 100 s durability test, the initial throat temperature of the thruster remain constant at 650 K. It can be concluded that a high level of combustion efficiency can be achieved at the test conditions compatible with the engine life requirements of 130 cycles in this test conditions. On the other hand, the r\C for the durability tests of the perfect ZrC)2/Ni FGM chamber (as shown in Figure 7 (b)) was almost constant at during 50 cycle tests. 100

1800

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(a) Pertial ZrCh/Ni FGM chamber

(b) Perfect Zr02/Ni FGM chamber Figure 7. C and C* efficiency vs test cycles

3.4 Specific Impulse Performance Figure 8 shows the effect of mixture ratio on performance at a chamber pressure of about 1.2 MPa. It was shown the maximum vacuum specific impulse at near the mixture ratio of 1.60. From Figure 8, it may be seen that the mixture ratios of peak specific impulse coincides with those corresponding to the designed Rupe number of 0.5. This is perhaps due to the nonreactive effects of the core flow with the film coolant as shown in the film cooled thrust engines.^ ) The mixing factor M at the peak vacuum specific impulse was 0.85, which was the small value compared to the designed point of 1.0. The vacuum specific impulse Igpy decreased as mixture ratio MR increased from around 1.60 to 1.87. C* efficency decreased from 97 3 % to 95.5 % as the mixture ratio increased from 1.42 to 1.87. Figure 9 shows the effect of chamber

473

pressure on performance at a mixture ratio where the specific impulse peaked. The C* efficency and Igpv attained almost steady values at chamber pressure as low as 1.3 MPa and continued steady values of Ispv=318 s after Pc=l-3 MPa, which may be attributed to the steady state combustion (no separation at high chamber pressures ) with respect to the high performance unlike quadlet elements design injector used. The chamber pressure at which the low frequency combustion instability, namely chagging, occurs was observed at below the chamber pressure of 1.02 MPa due to the decrease of the pressure drop in the injector. The definite decrease at chamber pressure of around 1.0 MPa may be attributed to the combustion instability which was observed in the injector performance tests before the HAPT. 100

325

3

320

99

315

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96

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1.2

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1.6

1.8

Pc (MPa)

Figure 8. Effect of MR on performance

Figure 9. Effect of Pc on performance

4. COMPARISON WITH PREDICTION Figure 10 shows the flow subdivisions used for the computation.(^) In Figure 10, ODE stands for one dimensional equilibrium flow, ODK for one dimensional kinetic flow, TDK for two dimensional kinetic flow, and TBL for turbulent boundary layer flow. The value of "(TDKTBL)" which is equals to I^py (energy release efficiency; riE;R=1.0) is defined as I spvCnER = 1.0) = loDE - KL - TDL - BLL

(5)

were KL, TDL and BLL is a kinetic loss, divergence loss and boundary layer loss respectively. A comparison of calculated delivered specific impulses and experimental ones is made in Figure Boundary layer friction and heat transfer loss ( T B L ) Injection w i t h reduced enthalpy 3/

,J>*>^"'

r d ^ ^ a>

*Flow

I X Supersonic Combust 1on\-\1 1 1 ^ chamber Transonic TM

region

region ODE Incomplete energy release (Core J?^*)

ODK 1

TDK

Kinetic and TD loss

Figure 10. Flow subdivisions used for the computation

Figure 11. Comparison of calculated delivered specific impulses and experimental ones

474 11 under the conditions of P^ 1.2 MPa and the nozzle area ratio of 300:1. In this calculation, the effect of chamber energy release loss and that of mixing between core propellant was not taken into account in the calculation for simplicity. As shown in Figure 11, the kinetic loss was smaller in the lower mixture ratio. Experimental performance tends to be higher in the lower mixture ratio and lower in the higher mixture ratio compared to the calculated performance. 5. CROSS SECTIONAL PROFILES Figure 12 shows the cross sectional profiles of the perfect Zr02/Ni FGM chamber after the durability tests. As shown in Figure 12 (a), the electro-forming (E.F) FGM layer at cirindrical section was gradually changing from one side of the Zr02 18 vol %/Ni 82 vol % to the other side of the pure Ni layer. At the throat section ( shown in Figure 12 (b)), however, the width of the E.F FGM layer was not the same width with the designed value of 2 mm. After the cycle tests, the large degradation of the inner surface coating at the throat section was observed. It may be attributed to this defection of FGM layer.

(%)

(a) Cylindrical section (b) Throat section Figure 12. Cross sectional profiles of the perfect Zr02/Ni FGM chamber

6. CONCLUDING REMARKS 1) A high level of combustion efficiency can be achieved by using Zr02/Ni FGM chamber and the delivered maximum specific impulse was 318 s at Pc=1.3 MPa near the design mixture ratio of 1.65. 2) Long duration (150 s) firing tests at a 1.05MPa chamber pressure were conducted without any explosion due to the heat soak-back. 3) It was shown that FGM coating layer remained in tact except the spalling near the inner surface of the chamber composed of ZrOo (100vol%). REFERENCES 1) Kuroda, Y : Evaluation Tests of Zr02/Ni FGM for Regeneratively Cooled Thrust Engine Applications. Ceramic Transactions, FGM, Vol.34,1993,pp.289-296. 2) Ueda, S : Performance of the N2H4/MMH Mixed Fuel Regeneratively Cooled Engine. NAL TR-1082, 1990.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

475

Research on Enhancement of Thermoelectric Figure of Merit through Functionally Graded Material Processing Technology in Japan

Takenobu Kajikawa Shonan Institute of Technology Fujisawa, Kanagawa 251 Japan

Abstract For usefiil thermoelectiic elements, the figure of merit should be high over wide temperature range. In order to realize such a thermoelectric performance, Functionally Graded Material (FGM) approach has attracted engineers' attention. The researdi activities on the enhancement of thermoelectric figure of merit through FGM processing technology in Japan are reviewed in the paper. After the fundamental design guideline of FGM thermoelectric element is presented, the experimental results for several kinds of FGM approach are introduced mainly based on the FGM projects named 'T)evelopment of Energy Conversion Material though Formation of Graded Structures" promoted by Science and Technology Agency and Thysics and Chemistry of Functionally Graded Materials" supported by Ministry of Education and Culture. The technological problems and future prospects are also discussed. 1. Introduction The worth of a thermoelectric material is expressed by the figure of merit Z, which depends upon the electrical and thermal properties of material , expressed as a^ 6/K, where (7 is the electrical conductivity (S/m) and K is the thermal conductivity (W/mK). The energy convention efficiency V is given by (AT / Th) (M-1) / (M + TcATh), where M is (1+Z Tav)o ^ Tav is mean temperature = (Th + Tc) / 2, Th and Tc are hot and cold junction temperature respectively. Therefore, a good thermoelectric material is required to have a large Z over the whole temperature range. However, the figure of merit for usually used materials is represented as a function of temperature and has a peak value at the pecuhar temperature. Hence, the concept of Functionally Graded Material thermoelectiric element is proposed in order to get the efficient thermoelectric element over a wide temperature range. The objective of this paper is to review the research on an enhancement of thermoelectric performance through fimctionaUy graded material processiag technology

476 in Japan. The experimental results are introduced mainly based on the FGM projects named 'Development of Energy Conversion Materials through Formation of Graded Structures" promoted by Science and Technology Agency since FY1993 and Thysics and Chemistry of Functionally Graded Materials" supported by Ministry of Education and Culture since FY1996. The technological problems on thermoelectric device and future prospects are also discussed. 2. FGM i^proach for Thermoelectric Material The general characteristics of thermoelectric material are summarized as follows: 1) The figure of merit Z consists of three parameters; Seebeck coefficient, electrical conductivity and thermal conductivity. 2) The value of Z is sensitively dependent upon the material species, composition, dopant level, structure and temperature. 3) The peak value of Z exists at the pecuhar temperature in the impurity conduction range in general. By the usage of these characteristics FGM approaches for thermoelectric material are proposed to control the temperature dependence of the figure of merit as follows: 1) Material species graded 2) Composition graded 3) Dopant concentration graded 4) Structure graded xlOFor a material species graded P777771 Dopant Graded FGM approach, several kinds of (AO^b^ Composition Graded FGM thermoelectric material of which the optimum operating temperature is different firom each other are laid in order. Concerning composition graded and dopant Integrated FGM Element concentration graded approach FGM thermoelectric element is made based on the fact that the sHghtly different parameters such as the ratio different species se^ lentation of material composition or hBiiTei i—PbTt. J 1 i dopant concentration can 600 800 1 000 1 200 1 400 200 400 control the temperature Temperature T (K) dependence of the performance, because the Fig.l Integrated FGM thermoelectric element using the Seebeck coefficient and conventional materia]

477 electrical conductivity are a ftmction of such parameters. The structure graded FGM approach is classified with macro-, mezo and micro-scale approaches. From the point of the shape of the gradation, multi-staged and continuously graded type are proposed. The optimal FGM approach should be selected or integrated in consideration of operating condition and material characteristics. By the usage of the conventional thermoelectric material; BisTea for low temperature range, PbTe for mediimi temperature range, Sio.TGeoa for high temperature range and LaTei.4 for the highest temperature range, an integrated FGM thermoelectiric element appHed with various kinds of FGM approach is designed as shown in Fig.l {1}. The performance of the FGM thermoelectric element represents nearly ZT = 1 for the temperature range form 300K to 1400K. The energy convention efficiency is calculated 23.3%, twice of that for single thermoelectric material such as Si-Ge {1}. 3. Experimental Research on FGM Thermoelectric Elements 3.1. Dopant concentration graded FGM The dopant concentration graded FGM thermoelectric element based on PbTe has been experimentally researched by National Research Institute for Metals, STA {2}. The PbTe material was prepared using a vertical gradient freezing method, for which temperature gradient was 8K/cm, in a quarz ^ass ampoule. P b t was used as a dopant. First, three kinds of PbTe material of different dopant level were made separately. Each ingot was pulverized and was filled in a graphite die in order of the magnitude of carrier concentration. Three-staged graded PbTe element was hot-pressed. At each boimdary the mixture of neibouring material was bound as a relaxation of thermal stress and a reduction of Peltier effect. The carrier concentration is distributed 3.51 X lO^^-^ for high temperature range, 2.6 XIO^^^ for medium temperature range and 2.26 XIO^^-^ for low temperature range respectively. The thickness of each layer is 2.0mm and the diameter is 10mm. The overall Seebeck coefficient and overall resistivity were 129.2 /i V/K and 7.91 X lO-^ Q m at room temperature. The overall resistivity is higher than that of each layer. The temperature dependences of the power factor are shown ia Fig.2 for the temperature range from 300K to 770K. The performance of FGM element was lower than that of each layer except b layer for low temperature range while it was higher than that of each layer for high temperature range. It can not say an5rthing on the performance of FGM as composed with a homogeneous thermoelectric element, because the performance of FGM element should be determined with overall temperature difference at the designed temperature distribution. Then, the relationships between maximum power output per volxmae/length and the temperature difference at the constant cold junction temperature (300K) are shown in Fig.3. It is apparent that FGM element is superior to non-FGM element. At 486K in temperature difference, it is higher than that of (C), the best of non FGM element by 11% in this case. The experiments on dopant concentration graded FGM using n-type BisTea have

478 0.005 been carried out by Kogakuin University {3}. Two samples doped with J 0.06 and 0.15 wt% of HgBr2 were § 0.004 soldered to form an FGM element. It was found that the temperature of b maximum power factor shifted to a Q 0.003 higher temperature with increasing ^ dopant concentration. "o 3.2. Composition graded FGM ^ 0.002 Basic research on the composition U graded FGM thermoelectric element ^ based on p-type Bi-Sb-Te has been 0.001 carried out by Electrotechnical 250 350 450 550 650 750 850 Laboratory, AIST, MIT I Tenv^erature T (K) Fig.2 Temperature dependence of power factor {4}. In the experiment pfor dopant graded FGM element

type

(Bi2Te3)i.Y(Sb2Te3)Y

elements were prepared by PIES (Pulverized and Intermixing Elements Sintering)method {5}. The characteristics of their thermoelectric properties are shown in 3 Fig.4. The FGM effect to the power output due to the three staged gradation was detected with the experiments varying the direction of 0 100 200 300 400 500 600 Temperature Difference A T (K) the heat flow; that is, the direction available to Fig.3 Characteristics of power output v.s. temperature FGM effect and the difference for dopant graded FGM element reverse direction. The

I

i

resultant power output was 5.4% larger than that for the reverse direction. As the temperature range was narrow, the FGM effect was obtained to be small. It is necessary to select the proper thermoelectric materials and the proper temperature width for composition graded FGM element. 3.3. Continuously Graded FGM A modified sputtering method controlling the power supphed to the target combined

479

(Bi2Te3)i-y(Sb2Te3)y

Measument point CASE!

CO

o

(J CO

o Q.

0 F' I 300

I

I—I

I

1 I

350

I

I

I — I I I

400

Temperature /

i_-i

u-x.

450 K

Fig.4 Thermoelectric performance for (Bi2Te3)i.Y(Sb2Te3)Y elements

10

15

Location ( m )

Fig.5 Profiles of dopant concentration for continuously graded FGM

with shding shutter has been developed to make a thin film of the linearly graded of the dopant concentration on the substrate by Hitachi Co.,Ltd {6}. Target material was Si8oGe2o+ B(1.2atm%) and SisoGe2o respectively. The supphed power to each target was controlled: for example, the power to target A was changedfi-omOW to 500W, while the power to target B was changed from 500W to OW simultaneously . The controllable parameters are the rate of power change, the changiag pattern (linear or quadratic) and the shding speed of sht of 1mm in width. In the experiment three cases were carried out: easel; OW to 500W for target A and 500W to OW for target B, case2: 250W to 500W for target A and 500W to 250W for target B, and case 3: quadratic changeficomOW to 500W for target A andfi-om500W to OW for target B. The resultant distribution of the dopant concentration is shown in Fig.5. It suggests that it is possible to make a continuously graded FGM thermoelectric element changing not only the dopant concentration and also the composition of the element due to this method for various kinds of thermoelectric material . 4. Technological Problems and Future Prospects At the present research phase, it can be said that the approach on FGM for thermoelectricity has been classified and the experiments on multi-staged graded FGM thermoelectric elements have been carried out to verify the FGM concept by using the conventional material . It is the most important to develop processing methods for more efficient FGM thermoelectric element over a wide temperature range including the processing methods for a continuously graded FGM element. One of the important issues on FGM thermoelectric technology is concerned with the stabihty and durabihty of FGM performance for a long term. The diffiision phenomenon due to thermal process is inevitable for aU graded material or inhomogeneous material .

480 Hence, it probably causes the deterioration of thermoelectric performance of FGM element at high temperature range. However, such a situation is usual for a thermoelectric device which has the junction between thermoelectric element and electrode or the junction between n type thermoelectric element and p type one. The durability for Si-Ge device and Pb-Te one has been already established at high temperature range {7}. Such an experience can be appHed to the constitution of FGM thermoelectric element. The technological approaches to keep the thermoelectric performance of FGM element for a long term can be classified as foUows: 1) To make a difiRision barrier for each layer to multi-stage graded FGM element. 2) To apply micro level diERision barrier mechanism to a continuous type FGM element. 3-staged graded FGM thermoelectric element made of Si-Ge has been made and tested by Shinku-Yakin Co.,Ltd.. {8}. The carrier concentration was graded in the element. There was no additional material at the interface. The distribution of the electrical potential shows that the additional contact resistance occurs at the interface. The increment of interface resistance of FGM element was estimated about 7-8% at 8mm and 7.5 5.0 2.5 0 0 2.5 5.0 7.5 5mm in element thickness.Therefore, distance from the interface (1000 A ) it is necessary to select a proper substi^ate n-type diffusion barrier material for each p-type Ni FGM thermoelectric element. The research on diffusion barrier 1 ; Ni! using Ni and Al has been carried out 1 Te 1 A : for the p-n junction between Bio.5Sb1.5Te3 and Bi2Seo.6Te {9}. In the ^^ 1 case of nickel film as a diffiision 7.S S.O 2.5 0 0 2.5 5.0 7.5 lO.O barrier, after 1 hour at 200°C nickel distance from the interface (1000 A ) film of 1 // m in thickness was diflftised Bi, Te, Se etc. On the other Fig.6 AES depth profile of annealed p/n hand, in the case of alumiaum fiJm, junction containing Al- and Ni- layer its layer of 1 /i m in thickness could

481 protect the invasion of Bi, Te, Se etc. as shown in Fig.6 {9} .The reason is considered that Al2Se3 at the junction between n type and aluminum layer and that AbTea at the junction between p type and aluminum layer have been formed. The free energy of formation of these compounds is very large as compared with Ni compounds. The largeness of formation energy can make an important role to protect the diBRision. This experiment can suggest the possibihty of durabihty of multi-staged graded FGM thermoelectric element due to the proper diEfusion barrier layer. That is, the interface layer should be made of certaia kinds of compound of which the free energy of formation is high and the resistivity is low. Item 2 is concerned with the surface modification technique of the powder before sintering. It is sirmlar to the microstructure-controUed thermoelectric element, which is composed of FeSi2 grains coated with undoped SiGe by means of rf-plasma processing developed by Yamaguchi University {10}. The schematical constitution is shown in Fig .7 {10}. It can be seen that FeSi2 cluster is connected with each other through homogeneous (thickness; lO^lOOA) imdoped SiGe surface. SiGe From the view point of keeping FGM structure, this technique can be appHed to aU macro-graded ionizfld donors FGM thermoelectric elements. The mother •lectrons material clusters of Fig. 7 Schematic diagram of the structure of FeSi^/SiGe diBferent carrier concentration or of different composition, or different microstructure are coated with a certain neutral or positive material by means of nonequilibrium high energy processing such as rf-plasma processing. The coated mother material clusters make up macrograded FGM structure. FGM clusters don't contact each other but through homogeneous coating material. At the interface there is no diffusion phenomenon. The combination of macro-graded FGM and microscopic structure control is needed to keep the durability of continuously graded FGM thermoelectric element.The above-mentioned approaches are igeneous schematically shown in Fig.8. vDiffusion ing material jD i f f u s i a barriers Another technological problem on FGM thermoelectric element is concerned with the operation of thermoelectric Continuously graded Multi-staged graded devices. That is, the performance at the offFig.8 Schematical approaches on durable design operating conditions FGM element should be considered,

482 because FGM thennoelectric element is made under the disigned temperature distribution at fixed heat flux. It is necessary to estimate the performance at off'-design operating conditions as the temperature level of the heat source varies with time or operating conditions in general. The definition of inhomogeneous thermoelectric properties and the measurement techniques of FGM element also include several technological problems to be solved. 5. Conclusion 1) Approach to the formation of Functionally Graded Material (FGM) thermoelectric element is classified as the gradation of material species, composition, dopant concentration and structure and the shape of gradation such as multi-staged and continuously graded. 2) The FGM effiect on the thermoelectric energy conversion efficiency was estimated to get 23.3% for the temperature rangefirom300K to 1400K, twice of single thermoelectric element such as Si-Ge. 3) The experiments on several kinds of FGM approaches were introduced to verify the usefulness of FGM thermoelectric element. 4) The several technological problems on FGM thermoelectric element are discussed. In particular, the approaches for the durabiUty of FGM structure are proposed to suggest the direction of basic research on FGM technology.

Acknowledgement Author expresses his appreciation to Dr. I. Nishida, National Research Institute for Metals, Prof. I. Shiota, Kogakuin University, Prof. T. Hirai, Tohoku University and Mr. T. Ohta, ETL for their support and useful suggestions. References l.I.Nishida,J.of Energy and Resources, 16,4(1995)416 2.I.Nishida, Proceedings of Symposium of the Japan Institute of MetaLs,(1996)9 3.H.Kohri, KSato, N.Abe, T.Suzaki, I.Shiota, LANishida, Proceedings of FGM95, The Society of Non-Traditional Technology(1996)95 4.T.Ohta,AYamamoto,J. of Advanced Science,7,3&4(1995)141 5.T.0hta,T.Kajikawa,CRC Handbook on Thermoelectrics,CRC Press(1995)109 6.M.Hayashibara,private communication 7.D.A.OT?iodan,Proceedings of 4th International Conference on Thermoelectric Energy Conversion,The University of Texas(1982)15 8.K.Takahashi,T.Masuda,T.Mochimaru,T.Noguchi, Proceedings of FGM'95(1996)123 9.I.H.Eim,D.H.lee,AIP Confewrence Proceedings 316(1995)254 10.K.Kishimoto,K.Nagao,T.Koyanagi,K.Mastsubara,FGM News 29(1995)14

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 1997 Elsevier Science B.V.

483

A Design Procedure of Functionally Graded Thermoelectric Materials J.Teraki^ and T. Hirano^ ^Daikin Industries, Ltd., MEC Laboratory, 3 Banchi, Miyukigaoka, Tsukuba-shi, Ibaraki, 305, JAPAN.

Abstract A computational design procedure of a thermoelectric power device using Functionally Graded Materials (FGM) is presented. A model of thermoelectric materials is presented for transport properties of heavily doped semiconductors, electron and phonon transport coefficients are calculated using band theory. And, a procedure of an elastic thermal stress analysis is presented on a functionally graded thermoelectric device by two-dimensional finite element technique. First, temperature distributions are calculated by two-dimensional non-linear finite element method based on expressions of thermoelectric phenomenon. Next, using temperature distributions, thermal stress distributions are computed by twodimensional elastic finite element analysis.

1. Introduction As part of the Japanese National Project on Functionally Graded Materials (FGMs), we have been studying design methodology for the application to be used in future space programs. The purpose of this study is to control the compositional and microstructural distribution in the materials so as to minimize thermal stresses imposed on the materials. This study can be understood as the tailoring of mechanical properties of materials. On the other hand, it is well known that electric properties of materials are also dependent on composition, microstructure and crystal structure. Therefore, the electric function of a material can be improved by application of the FGM concept to the optimization of the carrier conduction mechanism[l]. In Japan the new project for FGM research will focus on energy conversion materials. For the design of FGM thermoelectric device, it is necessary to get optimum energy conversion efficiencies, and to decrease thermal stresses by the high temperature difference within the device. In this paper, we describe about the electric property calculation using the band theory, and device efficiency calculation and thermal stress calculation using two dimensional finite element method.

484

2. Analysis Method .Energy

2.1 Transport Properties For the design of thermoelectric materials, it is necessary to obtain analytical transport properties, such as Seebeck coefficient, electrical conductivity and thermal conductivity. For this work, the band theory using the relaxation time approximation of Boltzmann's equation based upon simple semiconductor band model (Figure 1) has been adopted [2]. The concentration of carries, «±, is given by n- = 4d

Im^kgf

Conduction Band ]

^DiE)

0

Ed Dd -Eg Balance Band

SP(E)

Figure 1 Semiconductor Band Model

Fu.{r)

where + and - indicate holes and electrons respectively, ffi+ is the effective mass, T is the temperature, kg is Boltzmann's constant, h is the Planck's constant. And 4- is the reduced Fermi level, Fn(^ is Fermi integral given by "

r

v"

"^^^ J l + exp(x-|)

kj

kJ

where ^ ^ i s the Fermi level. Eg is the band gap, ^ is the reduced energy. level is obtained by the charge neutrality condition N,=n--n^

, N,

ew{^d)^exp{^-)

The Fermi

kJ

where Nd is the net number of ionized impurities, Dd is total Dopant concentration. Using the relaxation time approximation of Boltzmann's equation, the expression for the properties (electrical conductivity, Seebeck coefficient, and Lorenz number) of the holes and electrons are given by %m CF~

+k.

- •

3/w

r-=\^ where e is the magnitude of charge on an electron, G^n(^ represents an average of

485 the relaxation time given by

0 {l-\-exp{x-<^)f where T ± is the appropriate scattering rate. Two scattering mechanisms (by acoustic lattice vibration and by ionized impurities) for the carriers are considered. The total electrical conductivity , Seebeck coefficient, and thermal conductivity by carriers are given by cr = a^+cr , a^ia'^a^

-\-a a )/(J,

K^ =

VG^

+L-G~ +

a G

{a*-a-)\T

For the lattice thermal conductivity, the model due to Steigmeier and Abels is adopted, here [3]. The lattice thermal conductivity is given by ' '^R-'

i^B

'•4

where

':-h.

0/T c'* exp(x) x"^ exp(^x) dx, /3 = J -dx, 7 2 = 1 0 {exp{x)-\) T^N {exp(x)-\y

,

/"

^N V

x"^ exp(x) r^y

{exp{x)-lf

dx

V is the speed of sound, 0 is the Debye temperature, TC is the total phonon-scattering rate, TN is the phonon-scattering rate due to three phonon normal processes. In this model, two additional scattering mechanisms of phonon (by point defects and by charge carriers) are considered. 2.2 Thermoelectric Analysis Figure 2 shows a simplistic model of a FGM thermoelectric power device. we proceed to calculate the temperature and electric potential distributions within each arm of the device. The internal physics is a coupled phenomena of heat flux and electric current, and the expressions of thermoelectric phenomenon in isotropic body are given by G

Ga

TaG

K + TGa^

-vr

and, for steady-state, we obtain

Figure 2 Schematic of Device Model

486

-V-^ + y ( - V ^ ) = V(;fVr)-7Vay + / / a = 0 wherey is current density, q is heat flux, ^ is electric potential, Tis temperature, a is electrical conductivity, a is Seebeck coefficient and K is thermal conductivity. For two-dimensional finite element analysis, we employ a weighted residual method, then we obtain JVw(oV^ + aaV7)c/v + j w y „ ^ = 0 V

S

V

V

S

where w is weighted fimction, >, qn are the current density and the heat flux in the direction normal to the boundary, respectively. These equations are nonlinear with respect to the electric potential and temperature. Therefore, Newton method is employed to solve the equations. Two-dimensional fore-node elements are used, the FGM properties within the element are interpolated with the nodal point values. 2.3 Thermal Stress Analysis We calculate thermal stress in thermoelectric device by means of conventional finite element method using calculated temperature distributions. Twodimensional four-node plane strain elements are used in calculation. 3. Calculation Results 3.1 Graded Dopant Concentration Design for FGM Here, we calculate the conversion efficiency of FGM having the optimized graded dopant concentration. Figure 3 shows the temperature dependent figure of merit (ZT) calculated for Bi2Te3 with dopant concentration as a parameter. The ZT curve for the uniform material of dopant concentration is indicated by the solid line. It is

1

1 y^^''

A.., 1 1

Dd :Dopant concentration

•'

i ^^^^-^'^"'^

geoo 1 1 41 li

Dd! 0

100

;'.-'!',.. 2 7 ' i \ 25.^-

y ^

j/"'^

1

Temperature 1

J

f^-'/y 1 n - ~ ~i

log!(Dd) = ,25 1—

Dopant conpentration 200

300

400

1 500

1

|--

""i

600

700

800

Temperature ( K )

Figure 3 Dimensionless Figure of Merit Calculated for n-type Bi2Te3

0

0.2

1

1

0.4

0.6

1

0.8

1

X (Location in thicl(ness)

Figure 4 Graded Distribution of Dopant and Temperature in Optimized FGM

487 shown that the higher dopant concentration in the hot side is increase the total ZT. Figure 4 shows the optimized gradient distribution of dopant concentration for FGM and the temperature distribution in the FGM. Table 1 shows the conversion efficiencies of power generator using three t5^ical thermoelectric materials. In the case of Bi2Te3, the conversion efficiency of the FGM is calculated as 12.25%, which corresponds to 17% increase in the conversion efficiency. Table 1 Conversion efficiency of Optimized FGM and n-FGM Bi2Te3 Material PbTe SiGe 300-'700 700-1000 1000-1300 Temperature (K) 5.20 12.25 4.51 FGM Efficiency (%) 4.96 4.49 10.46 n-FGM Efficiency (%)

Total 300-1300 20.56 18.72

3.2 Thermoelectric Analysis and Thermal Stress Analysis of FGM Device Thermoelectric analysis and thermal stress analysis are performed on single-stage Electrical Insulator thermoelectric generator by the finite , Electroad I element method. An analysis model is shown in Figure 5. The device is constructed by ceramic electrical insulator, p-type n-type 4.6mm copper electrode and Bi2Te3 thermoelectric Bi2Te3 8i2Te3 arms. 300~500K is loaded as temperature JL difference, and a symmetry condition is set on the right face as a mechanical boundary 1^ 2mm <^ condition. 8mm Heat flux and temperature distributions are shown in Figure 6, and current density Figure 5 Analysis IVIodel and electrical potential distributions are shown in Figure 7. It is shown that the electric current is generated by temperature difference in the device. With the calculated temperature distribution, we can obtain the thermal stress distribution. Displacements and maximum principal stress distributions are shown in Figure 8. By the material expansion in *JM*JW

~»MX73

Figure 6 Temperature and Heat Flux Distribution for Biales Device

Figure 7 Electric Potential and Current Density Distribution for BiaTea Device

488 the hot side, the module is deformed, and the thermal stresses are concentrated at the interface between the electrode and ceramic insulator. Figure 9 shows the maximum principal stress distributions in the copper electrode and the copper/ceramic FGM electrode. It is found that the 30% decrease can be obtained for the thermal stress in the interface between the electrode and the ceramic insulator. S«rfa**

Figure 8 Deformation and Maximum Principal Stress Distribution for BiaTes Device

4. Conclusions Figure

9

Maximum

Principal

Stress

Computational approach to the design of Distribution in copper Electrode (upper) and thermoelectric FGM device was presented, copper/ceramic FGM Electrode (lower) which includes theoretical calculation of transport properties of thermoelectric materials, and two-dimensional thermoelectric and thermal stress analysis. Graded dopant concentration design for Bi2Te3 was studied and it was found that the 17% improvement can be obtained for the conversion efficiency. From the thermal stress analysis results for Bi2Te3 thermoelectric device, it was shown that the maximum thermal stresses were concentrated at the interface between the electrode and the ceramic insulator. And, the FGM electrode can be effect the 30% decrease for thermal stress. Acknowledgment This study was performed through Special Coordination of the Science and Technology Agency of the Japanese Government.

REFERENCES 1. T.Hirano, L.W.Whitlow, J. Teraki and M.Miyajima, Froc. of 3rd Int. Symp. on Structural and FGM, (1994), 633. 2. C.B. Vining, J. Appl. Phys. 69(1), 1 January, 331, (1991). 3. Steigmeier and B. Abeles, Phys, Rev. 136, A1149 (1964).

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 1997 Elsevier Science B.V.

489

Transport Properties in Multi-Barrier Systems Y. Nishio and T. Hirano MEC Laboratory, Daikin Industries Limited, 3 Miyukigaoka, Tsukuba, Ibaraki 305, Japan. Thermoelectric properties in a multi-barrier system have been investigated theoretically to develop a high-Z materials by eliminating the lower energetic carriers by barriers. It has been shown that the elimination of lower energetic carriers by barriers always brings an enhancement of thermoelectric figure of merit. The optimal height of potential barrier is given analytically. By using the Kane model, role of minority carriers are clarified. The enhancement of figure of merit due to the elimination of minority carriers are greater than that due to the elimination of lower energetic majority carriers. 1. INTRODUCTION Moizhes and Nemchinsky[l] proposed a new method to improve the thermoelectric figure of merit by eliminating the lower energetic carriers using potential barriers as an energy filter. Subsequently Whitlow and Hirano[2] suggests the applicability of this method(energy filtering method) to SiGe system, and the physical interpretation has been given by Rowe and Min[3]. Recently authors have also considered the energy filtering method theoretically[4] and the optimal height of barrier and the scattering dependence of improvement of figure of merit has been clarified within the parabolic single band model. However most of thermoelectric materials known presendy are described by two band non-parabolic model and harmful effect of minority carriers cannot be neglected. In this article we have employed the Kane model and clarify the role of minority carriers. 2. DETERMINATION OF OPTIMAL BARRIER HEIGHT Dimensionless thermoelectric figure of merit Z = zT can be considered as SL functional of differential conductivity a(e). The existence of barrier changes a(e) as o(e) + 5o(e). The resulting change in Z can be expressed as[4]: 8ZHZ[CT + 5CT]-Z[o] = J _ " ^ d e ^ 5 a ( e )

(1)

490 The functional derivative of Z with respect to a(e) can be calculated easily and we have obtained following results:

=smh

8Z _ - z f ^ f " ( e ) l l . . .^f + ^ ( i + z ) ( e - n ) - H £ ^ ( l + Z ) 8a(e)

(2)

Here o is electrical conductivity, a is thermopower, K is thermal conductivity, e is energy of carrier, ji is chemical potential, e is bare charge of electron, and f^(e) is Fermi-Dirac distribution function. In deriving eq.(2) we treat the lattice thermal conductivity KL as a constant. Following we consider the n-type semiconductors, then the change of differential conductivity can be given by: 5a(e) = -a(e)0(e-eB)

(3)

i.e. carriers below potential barrier do not contribute to transport processes. In order to keep 6Z>0, the functional derivative should be negative. Therefore (e -^i)' + ^ ( 1 +Z)(e -^i) + 2 ^ ( 1 +Z) > 0.

(4)

Equation (4) determines the height of potential barrier Eg and the solution of eq.(4) is

Therefore if we make the potential barrier below E^_ or above £3+, the change of figure of merit 6Z becomes always positive. However the effect of the barrier above 63+ is neglected in following analysis, since the expression of 6Z contains the derivative of Fermi distribution function (eq.(2)) and therefore the contribution of carriers above £3+ is considered to be small. It is easy to see that eq.(5) gives the generalization of height of the optimal barrier derived earlier[4] 3. ENHANCEMENT OF Z IN TWO BAND SYSTEM In this section we estimate the enhancement of Z and characteristic feature of this method by using Kane model. In Kane model, the transport coefficient, when the particular type of scattering mechanism is pre-dominant, is given by[5]: a = Z cJi, a = i 5^ Gitti, K = LaT+ K^^, + KL, i=e,h

i=e,h

(6)

491

o. = A 0^3 ,j(Ti,p),

a,a. = - ^ A [ G ; . 3 /,(Ti,p) - I]GI, /^(Ti.p)], [GJ.3,,(TI,P)

(7)

(8)

G;.3/2(^.P)

(9) (10)

(11) Here the suffix e denotes electron and the corresponding expression for holes is obtained by substituting T] -> - T| - x^ in above equations(the expression of hole contribution for thermopower is required to change the sign in addition to the substitution). We have used the relaxation time approximation in above expressions and assumed the energy dependence of relaxation time is given by[5]: T(X) = X?

x^(l + px)^ (1 + 2px) '

(12)

and r depends on the scattering process, i.e. r=-l/2 corresponds to acoustic phonon scattering, r=l/2 to optical phonon scattering, and r=3/2 to ionized impurity scattering[5]. To perform the numerical calculation, we use the semi-empirical formula of energy gap forBii_ySby [6] with y=0.1 as a typical example: E^ =270y - 15 + 4.0 x 0.032T (meV), and set the temperature at 100K(therefore x^ = 2.878). We also fix the ratio K^ / KL = 1.

(13)

492 Figure 1 shows the ratio of the contribution between electrons and holes to thermopower as a function of reduced chemical potential. We have also shown the results in the case of parabolic band(|5 set to be zero while x^ kept fixed value corresponds to the parabolic band). KambJK total, 1

10^ -•-Acoustic Phonon -'-Optical Phonon -•-Ionized Impurity -^Parabolic Band

10^

'

'

!

'

•

•

'

1

1

.

.

T-

r-1

-•-Acoustic Phonon ] i - ' - Optical Phonon 1 j -•-Ionized Impurity f 1 -©-Parabolic Band | ]

0.8 ?!?i8^ 0.6

H^^^^H'^'"'^ 1

0.4

10' 0.2

10°

-1

01 -1 -0.5

0

0.5

1

-0.5

1

I

1

i

0

1

>

>

•

i

1

0.5

•

i

•

i

1

>

1

•

>

i

I

1

1

1

1.5

1.5

Figure 1. Ratio of contributions between electrons and holes to thermopower as a function of reduced chemical potential.

Figure 2. Ambipolar contribution to thermal conductivity devided by total thermal conductivity as a function of reduced chemical potential.

Z From Fig.l the contribution from minority -Acoustic PhononI 1.5 -Optical Phonon carriers to thermopower is comparable -Ionized Inipurity with that of majority carriers near T] = - 1. -Parabolic Band 1.25 This reflects the increase of the ambipolar contribution toward r| = - 1 as seen from Fig.2. Therefore the rapid decrease in Z 0.75 toward r| = - 1 is due to the minority carriers in this model. We also see the properties near r| = 2 correspond to be single band. 0.25 Next we estimate the enhanced figure of merit as a function of reduced chemical potential simply by setting the lower limit of integral in eq.(lO) by corresponding Figure 3. Thermoelectric figure of merit as barrier height using eq.(5). Figure 4-7 a function of reduced chemical potential for various scattering processes. shows the effect of energy filter. We see the influence of elimination of minority carriers has considerable effects on thermopower and thermal conductivity.

493 Since, from Figs.4-6, the effect of increase in thermopower and of decrease in thermal conductivity is superior than that of decrease in electrical conductivity, the figure of merit increases toward T] = - 1. Comparing Fig.3 and Fig.7 we see that the enhancement of figure of merit comes from two different origin. The enhancement near Ti = - 1 is due to the elimination of minority carriers and that near r| = 2 is due to the elimination of harmful majority carriers. a Filtei ^Klteicd/^O

J%

1 -Acoustic PhononI -Optical Phonon - Ionized Impurity I -Parabolic Band

0.9 I 0.8 [•• 0.7 -•-Acoustic Phonon -^-Optical Phonon 0.6 H -•-Ionized Impurity o-Parabolic Band 0.5

-1

-0.5

0

0.5

1.5 -0.5

Figure 4. Electrical conductivity with barrier divided by that without barrier as a function of reduced chemical potential.

0

r| Figure 5. Thermopower with barrier divided by that without barrier as a function of reduced chemical potential.

^Filterec/S

11

• . . . , . . . . , , . . . , . .

.

.

j • .

-•-Acoustic Phonon i 0.8 11 ^ ^ Optical Phonon [ 1 -^•-Ionized Impurity \ |-«-Parabolic Band y\ 0.6

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. 1 j .

• • • 1

1

j

A———^"^

T

3.5 3

•

•

•

'

j

:

i

bs^

j

'

'

•

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i"*

\ 2.5 \ ^ ^

-0.5

'

1

•

•

•

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I

'

'

' •'

j

'

- ^

^ ^

kj^^^^i—-•—T^*—~t^——-i

^!tite}^

1.5 1

• i .

'

-•-Acoustic Phonon ...J - ^ - Optical Phonon -•-Ionized Impurity j... ^^ Parabolic Band

2

t

. .

•

'

•

J

0.2 Ol.. ,1 -1

1

.

. •. . i . . . . i . . . . i . . •. 1

• .

0

0.5

1.5

Figure 6. Thermal conductivity with barrier divided by that without barrier as a function of reduced chemical potential.

0.5

'

-1

1

1

•

•

i

•

-0.5

•

•

•

i

0

•

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.

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.

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•

•

I

i

1

I

•...•

i

1.1

1

•

1

1.5

Figure 7. Thermoelectric figure of merit with barrier as a function of reduced chemical potential.

494 Although our results of calculation are rather qualitative one, Efimova et.al.[7] estimated the influence of the minority carriers on Z for n-type PbTe in more realistic fashion and indicated the minority carriers are responsible for a decrease of Z by 40-50% around T « 500-600°C. 4. DISCUSSION AND CONCLUSION It has been proved that the elimination of lower energetic carriers brings the enhancement of Z and the optimal height of potential barrier within the two band model is given analytically. Such elimination (confinement) of minority carriers has already studied in GaAs/AlAs systems and it has been known that the use of graded alloy composition at interfaces of hetero-junction are important to confine the minority carriers effectively[8] due to decrese in the defect density near interfacial region. Finally we note that the possibility of destroying the effect of enhancement of Z by internal electric field, which is produced by blocked carriers located near the barrier interface. Such internal electric field is inevitable and may degrade the effect of enhancement by this method. We expect the cancellation of positive and negative blocked carriers located at the interface of barrier makes this internal electric field weaker, to some extent. However the effect of internal electric field is left as a future problem. ACKNOWLEDGMENTS One of the author (Y.N.) would like to thank Professor J. Yoshino of Tokyo Institute of Technology and Professor T. Koyanagi of Yamaguti University for helpful discussions. This work is partly supported by a grant from the Science and Technology Agency of Japan for the Development of Functionally Graded Materials for Energy Conversion. REFERENCES 1. B. Y. Moizhes and V. A. Nemchinsky, 11th Int. Conf, on Thermoelectrics, Texas, 1992, (University of Texas Press, Texas, 1993), p.232. 2. L. W. Whitlow and T. Hirano, J. Appl. Phys. 78(1995)5460. 3. D. M. Rowe and G. Min, 13th Int, Conf. on Thermoelectrics, Kansas, 1994, (AIP Press, New York, 1995), p.339. 4. Y. Nishio and T. Hirano, Jpn. J. Appl. Phys., submitted. 5. A. Anselm, Introduction to Semiconductor Theory, Prentice-Hall Inc., New Jersey, 1981, Chapter 9. 6. B. G. Martin and L. S. Lemer, Phys. Rev. 86(1972)3032. 7. B. A. Efimova, L. A. Kolomoets, Yu. I. Ravich and T. S. Stavitskaya, Sov. Phys. Semicond. 4(1971)1653. 8. L. W. James, J. Appl. Phys. 45(1974)1326.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

495

Theoretical estimation of thermoelectric figure of merit in sintered materials and proposal of grain-size-graded structures. Junji Yoshino Department of Physics, Tokyo Institute of Technology, Meguro-ku, Tokyo 152, Japan

Thermoelctric figure of merit for n-type sintered materials of SiGe and PbTe has been calculated based on Boltzmann equation and a heterostructure grain boundary model. The results reveals that thermoelectric figure of merit are expected to be improved in both materials reducing the grain size. However, maximum thermoelectric figure of merit is expected to achieve at different grain sizes for SiGe and PbTe, owing to difference of their sound velocities.

1. INTRODUCTION Recently, thermoelectric energy conversion has paid a great attention in terms of energy saving once again. However, improvement of its energy conversion efficiency is still a essential issue. Since thermoelectric quantities, which determine thermoelectric figure of merit, Z, strongly depend on temperature, functionally graded (FGM) structures, where impurity concentration or/and alloy concentrations are changed along temperature gradient in a thermoelectric element to achieve maximum dimensionless thermoelectric figure of merit, ZT, at each point, are expected to improve Z, effectively. However, it is probably inevitable to adopt sintered materials to prepare such practical FGM thermoelectric elements. Therefore it is important to examine the energy conversion efficiency of sintered materials theoretically. Although extensive theoretical studies have been made to clarify the grain boundary effect on thermal conductivity, the effects on electric properties have not been well examined. In this paper, thermoelectric figure of merit for n-type sintered materials of SiGe and PbTe has theoretically been estimated considering effects of grain boundaries on carrier mobility and Seebeck coefficient based on a model.

2. THEORETICAL MODEL In this paper thermoelectric figure of merit, Z, for sintered materials are evaluated based on theoretical calculation of bulk properties. The framework employed in present study to calculate bulk properties is similar to Vining's model [1], except some small modification mentioned in the following. Scattering processes included in Boltzmann equation are ionized impurity scattering, acoustic phonon scattering, alloy scattering and intervalley scattering for

496 calculations of electronic properties, and three phonon scattering, carrier scattering and alloy scattering for calculation of thermal conductivity. Since the effect of non-parabolic band is essential especially in narrow gap semiconductors, following non-parabolic effective mass is introduce in the calculation of PbTe properties. m :\ + In present calculation, it is assumed that the effects of grain boundaries for Z are brought only through electric conductivity, Seebeck coefficient and thermal conductivity. Namely, grain boundaries effects on these three important properties are taking account independently, and then thermoelectric figure of merit, Z, are obtained by using the equation, Z = cra^/fc, where cr, a and K are electric conductivity, Seebeck coefficient and thermal conductivity, respectively. The effects of grain boundaries for thermal conductivity is accounted by a simple relaxation time formula, r = L/v^, where v^ and L are velocity of phonons and grain size, respectively [2]. On the other hand, mean free path of charge carriers is fairly smaller than grain size, as typical grain size of sintered thermoelectric materials is about 1-100 |nm. Therefore, carrier scattering at grain boundaries is not essential. However, grain boundaries contain a lot of defects and sometimes it is consisted of amorphous materials. Then carrier mobility in grain boundaries is expected to be smaller than that of grain inside. Furthermore, localized energy levels which are created by dangling bonds, expected to pin Fermi level and form band bending. Several models have been proposed to account for temperature dependence of mobility in polycrystalline silicon. In this paper, a crystalline-amorphous-crystalline heterostructure model for grain boundaries presented by Kim et al. [3] has been accepted to take account of grain boundary effects on electric properties of sintered thermoelectric materials. Figure 1 presents the grain boundary model used in this study. The carrier transport in grain boundaries are assumed to take place by Brownian motion and the carrier mobility is inversely proportional to temperature. Several interface parameters, which characterize its band profile are also indicated in fig. 1. The meanings of these parameters are the same as that presented in Kim's paper. Figure 1. Energy band profile of sintered Since Seebeck coefficient is depend on Fermi materials based on heterostructure grain energy, it is position dependent, if one assumed a boundary model. band profile as shown in fig. 1. Overall Seebeck coefficient is calculated as follows. Total output voltage of single grain is estimated integrating Seebeck coefficient over a single grain. Position dependent Seebeck coefficient is obtained as a function of Fermi energy based on calculation for bulk, Then overall Seebeck coefficient is obtained by dividing total output voltage by a grain size. Major parameters used in

497 this calculations are summarized in table 1.

Table 1. Important parameters used in present calculation. SiGe anhrmonicity parameter of alloy scattering for phonon: y rate constant ratio related to intervalley scattering: W2/W1 band discontinuity at grain boundary: A thickness of grain boundary: S localized energy level around grain boundary: Ei

PbTe

0.91

1.5

1.0

0.5

OeV

OeV

lxIO"'m

IxIO'^m

- 0.5 eV

- 0.5 eV

3. RESULTS Electric and thermal properties of bulk materials have been calculated to examine reliability of our calculation. Figures 2(a) and (b) indicate temperature dependence of calculated electron mobility for bulk Sio.yGeo.s and PbTe, respectively. The mobility of SiGe is almost inversely proportional to temperature, while that of PbTe is almost proportional to T~^. It is notable that strong temperature dependence of PbTe, which is experimentally obtained, is successfiilly achieved. Although the effect of grain boundaries on Seebeck coefficient have been examined, it is found that the effect is only to become significant at a grain size less than 0.1 fim and it is negligible in common sintered materials. Therefore, the effect on Seebeck coefficient is neglected in the following calculation. Figures 3(a) and (b) indicate grain size dependence of Z r for sintered SiGe and PbTe, respectively. We can find significant difference in the grain sizes,

1000 Temperature (K)

1000 Temperature (K)

Figure 2. Temperature dependence of calculated carrier mobility for (a) SiGe and (b) PbTe.

498

10-6

10-5

Grain size (m)

Figure 3. Grain size dependence of dimensionless thermoelectric figure of merit, ZT ^ for (a) SiGe and (b) PbTe

which give maximum ZT. Namely, maximum ZT is obtained at a grain size around 0.1 jim for SiGe, while that of PbTe is achieved at a smaller grain size. Origin of the notable difference can be easily understood, if one examine grain size dependence of electric conductivity and lattice thermal conductivity show in fig. 4. Lattice

"T

100000

'"1

I

400 K

(b) PbTe

600 K 50000 800 K

•H«j—I I IHIH]

1 I Mllll|—t I l l l l l l l — I I I mil

400 K

600 K 800 K

10-8

10-7

10-6

10-5

Grain size (m)

10-4

10-3

10-8

10-7

10-6

10-5

10-4

10-3

Grain size (m)

Figure 4. Grain size dependence of electric conductivity a and lattice thermal conductivity KL for (a) SiGe and (b) PbTe.

499 thermal conductivity of sintered SiGe and Nd= 1 X 10 26 PbTe decrease with decreasing grain size as Si Ge 0.7 0.3 Nt= 3.2 X 10 28 shown in fig. 4. However, it is notable that J-400K grain size, where suppression of thermal 1.0 1200 K conductivity is take place, is not identical 1000 K for SiGe and PbTe. Lattice thermal 0.5 conductivity decrease gradually below a 800 K grain size of 10 \\.vci in SiGe, while it begin 600K below a grain size of 1 |Lim in PbTe. The nn significant diflTerence can be attributed to 10-6 10-5 large difference of sound velocity. Grain size (m) So far, extensive experimental studies on grain size effects have been made [4]. Figure 5. Grain size dependence of dimensionless thermoelectric figure of merit ZT for cases of high Although reduction of thermal conductivity locaHzed level density. Density of localized level is due to reduced grain size has been 3.2xl0^^m•^ confirmed by many authors experimentally, results on ZT are still controversial [4,5]. The calculated grain size dependence of thermal conductivity shown in fig.4(a) is in good agreement with experimental results obtained by Vining et al. [4]. On the other hand, their results indicate that degradation of electric conductivity takes place below the same grain size, while present result predicts that it takes place below one order of magnitude smaller grain size. This large discrepancy is due to the difference of electric conductivity in grain boundaries. The grain size dependence of electric conductivity is strongly affected by mobility ratio between grain inside and grain boundary, deference of their temperature dependence, and band profile. Therefor the grain size itself, which gives 2 7 maximum, shown in fig. 3 are not meaningful. The inconsistency of experimental results is also attributable to difference of grain boundary quality, since it strongly depends on preparation procedures. Therefor a great effort should be given not only to clarify the structure of grain boundaries, but also to control the quality of grain boundaries, in terms of thickness and defect concentration. Furthermore, interesting behavior on grain size dependence is expected. Figure 5 shows grain size dependence of ZT for a case of high localized level density. The grain sizes, which give maximum ZT, shift toward smaller grain size as increasing temperature. This results offer a new concept of a FGM structure, namely a grain size graded structure. However, band bending at grain boundaries reduces the peak value of ZT, and control of defect density at grain boundaries is essential problem.

4. SUMMARY Thermoelectric figure of merit for sintered materials has been investigated theoretically based on classical transport equation and a grain boundary model. The calculations reveals that enhancement of ZT by phonon scattering at grain boundaries are expected to be dominant at grain sizes of below 10 |im for SiGe and 1 ^im for PbTe. Since such enhancement is strongly affected by quality of grain boundaries, fijrther effort to control the quality of grain boundaries is expected to result in enhancement of Z r due to grain size reduction.

500 ACKNOWLEDGMENTS This work was performed through Special Coordination Funds of the Science and Technology Agency of the Japanese Government.

REFERENCES 1. 2. 3. 4. 5.

C.B. Vining, J. Appl. Phys., 69 (1991) 331. J.E. Parrott, J. Phys. C2, (1969) 147. D.M. Kim et at., TEEE Trans. Electron Device, ED-31 (1984) 480. C.B. Vining et al., J. Appl. Phys., 69 (1991) 4333, and references therein. D.M. Rowe et al, J.Appl. Phys., 73 (1993) 4683.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

501

Computer design of thennoelectric fimctionally graded materials L.I.Anatychuk, L.N.Vikhor Institute of Thermoelectricity, General Post-Office, Box 86, 274000, Chemivtsi, Ukraine Abstract The results of the optimal control theory use have been presented for program creation of functionally graded material computer modeling. Modeling results such as limiting values for FGM generators efficiency, fimctionally graded material advantages resulting from the computer modeling have been. Problem solution Achievements in thermoelectricity during the recent 30-40 years are first of all due to the progress in the creation of a thermoelectric material with high figure of merit values Z = ^

(1)

The maximum figure of merit is achieved by way of material optimization. It is well known that in semiconductors Seebeck coefficient a, conductivity a and heatconductivity K are the fiinctions of current carrier concentration «, p which in their turn are the fimctions of impurity concentrations Ne, Np a = a(n,p)= a(NeNp) a=a(n,p)=c7(NeNp) K= K(n,p)^

(2)

K(NeNp)

To reach the maximum of Z, one must find the optimal concentration fi-om the extremum condition —

=0

(3)

Thus, the result of optimization is a number, namely the value of the optimal concentration of donor or acceptor impurities. Naturally, the theory cannot indicate the exact figure of doping impurity concentration. Therefore the optimal concentration is experimentally determined by way of creating materials with different impurity concentrations and determination of their Z. This method served as the basis for highly efficient thermoelectric material technologies and guaranteed the progress of thermoelectricity on the whole. However, by the present time these ways of improving the thermoelectric figure of merit have practically exhausted themselves. Despite the numerous efforts, the increase in the figure of merit is insignificant. Therefore, it is only natural to ask whether this situation is a casual one or there are vaUd reasons for these restrictions.

502 To answer this question, special investigations were carried out. The value of Z was calculated for the model thermoelectric materials from which the best possible values of microscopic constants were taken in crystalline structures. The calculated values of ZT lie between 2 and 3. In reality this value will be considerably lower, therefore one can assume that this way of improving the thermoelectric figure of merit of material seems to have exhausted itself What are the further possible ways of increasing the efficiency of materials for the thermoelectric energy conversion? At present one of the most challenging trends of increasing the efficiency of thermoelectric materials is the transition from a classical thermopile model where legs are made of the homogeneous material, to the thermopiles where legs are made of materials whose properties are coordinate functions. The prospect of using the inhomogeneous legs was already indicated by academician loffe in his book [1]. A large series of investigations on the use of the inhomogeneus materials in cooling batteries has been conducted in Ukraine. A theory of computer-aided design of materials with a programmable inhomogeneity has been developed. Fig. 1 shows an example of the optimal inhomogeneity functions for the materials based on Bi-Te. These investigations have been described in the monograph by Anatychuk L.I. and Semenyuk V.A. [2]. Technologies have been developed for obtaining these materials by pressing, extrusion, zone melting and Czochralski methods. These materials have been used for the manufacturing of thermopiles whose properties are shown in Fig. 2. It is seen from the figure that when materials with programmable inhomogeneity are used, the maximum temperature drop in cooling thermopiles increases from 70 to 90 degrees, i.e. by 25%, and the coefficient of performance can increase several times. To reach such parameters with the homogeneous materials, one must create a material with a thermoelectric figure of merit Z=4,510"^ K"^ which for the moment is impossible. This example shows that the use of thermoelectric materials with programmable inhomogeneity is a new efficient trend of improving the quality of thermoelectACxlO",cm^

Number of ATxlO" cm'' sanipJc sections '

'

TV

FGM

«-type

{CdCL) L

Fig. 2. J - the increase in ^Tmax with Fig.l. Optimalfunctions of thermoelectric approximation to the optimalfunction; material inhomogeneity.n-type Bi2Te2jSeo.3 + 2 - the increase in the coefficient of per(0,09,,. O.OSJCdCh; formance zfor the optimal inhomogeneity; P'typeBiosSbi^sTes + 4% Te 8o is coefficient ofperformance of the homogeneous leg.

503 lie devices. Materials with programmable inhomogeneity can be also used to improve the efficiency of a generator. This idea is extremely popular with the Japanese investigators. Thus, in the paper by Nishida [3] it is suggested that a material with programmable inhomogeneity should be created on the basis ofPb-Te. The generally accepted name for these materials in Japan is functionally graded materials. In conformity with the paper [3], FGM is defined as the envelope by the maximums of Z at various temperatures and, accordingly, at various concentrations. Equivalently, for wider temperature ranges FGM is formed as the envelope by the maximums of Z for various materials suitable for each temperature level. In reality, to design FGM, use must be made of more accurate, hence more complicated methods. Let us consider them in more detail. The purpose of material design is to determine the distribution of impurity concentration along the leg where at the given temperature of hot and cold sides Th and Tc the maximum efficiency or maximum power is reached. Thus, we have a basically new approach to the optimization of thermoelectric material. Really, if before optimization served the purpose of number, today it results in the optimal function of impurity concentration as coordinate function. To determine this optimal function is a complicated task. Unlike finding the optimal concentration in the homogeneous material, experimental determination of the optimal function is a complicated and expensive task with a result which is apriori not very reliable. Really, for this experimental selection one must make a leg of individual source materials with various concentration values. The greater the number of such source materials and sections, the greater the chance of more accurate determination of the desired function. Naturally, here we must accept a limitation that we have to change a continuous concentration function for a step one. Combinations of these constituent samples are extremely numerous. Under these complicated experimental conditions it is advisable to pass fi*om a real experiment to a computer-made experiment, which in our opinion is more accurate and involves considerably less expenditures. Fig. 3 gives a physical model for the construction of a computer-made FGM design. This model includes an inhomogeneous leg of nand p-type of conduction, ^iW' 'AT;"" K gr^TriT, connected to electric and thermal circuits, external optimal electric load Ropt A L where electric power W is
MAAAA/VV

504 different Seebeck coefficient values inside the leg. The model allows for the existence of contact resistances ro^'^^\ as well as losses of heat ATi!''^^^ and electricity in commutation plates with the electric resistance i?/. The source system of equations describing thermal and electric processes in the infinitesimal leg part dx is given by expressions dT = ai T rr i q -—dx K K (4) dq _ a ^i rp ai i + —q + — dx K K a with the boundary conditions T„(0) = Tp(0) = T,,

TJl)

= T^(l)

= T,

(5)

where / is the density of generated current. The equations are based on the differential Ohm's law for the circuit comprising thermoelectric sources and on the law of energy conservation. From the solution of the system of equations one can find temperature distribution in the legs T„,p=Tn,p(x); the distribution of thermal flow qn,p=qn.p(x), as well as the thermal flow entering and leaving n- andp-type legs, the integral electric power on the external load and, accordingly, the generator efficiency. At the predetermined functions of the leg material inhomogeneity by varying the value of the external resistance R one can determine the maximum efficiency and the maximum circuit power, and here, as is well known, R^pf ^^^ ^ R^pf jy. To solve this direct task that comes to finding efficiency and power, a program has been developed based on the combination of Euler method and the method of shooting for the system of differential equations (4). However, our object in this investigation is to find the ways of solving the inverse task, namely, the search for the optimal function ofp- and n- type leg material inhomogeneity at which the maximum efficiency of thermoelectric generator is reached. This task was solved by way of development of a special-purpose computer progralm based on the use of Pontryagin mathematical theory of optimal control [4]. The essence of the method lies in the fact that a zero approximation of the desired function is entered into the computer in the form of material parameters Ne and Np that are coordinate independent or linearly changing. Pontryagin method comes to the search for the optimal function by calculating the efficiency for various functions that are different from the assigned zero ones. In this way a computer experiment is actually carried out instead of a real one. In so doing, the method and computer speed allow to search for as many functions as necessary for obtaining a predetermined accuracy. Thus, usually for the exact finding of functions the leg is divided into 50-^100 sections. Use here is made of 30-5-50 concentration intervals of Ni^\ We get an enormous number of combinations which cannot be sought for even by computer method. Pontryagin method allows to make a purposeful search for function and thus considerably reduce the number of computer experiments. The essence of this method lies in the determination of the optimal fiinctions N„,p(x), characterizing the inhomogeneity of thermoelectric materials of«- and/?-type from the condition H^Jil/(x)J(x),q(x),N(x),i)

= mcDC H„^pMx)J(x),q(x),N,i)

where Hamiltonian function H has the form

(6)

505

0

0.2

0.4

0.6 0.8

Fig. 4. Optimalfunctions of Ne, Np impurity concentrationfor n- andp-type alloys based on Bi-Te, accordingly Hn,p=(¥lfl-^¥2f2)n,p

0.2

0.4

0.6 0.

0.6 0.8 1 x/l Fig. 5. Optimal functions of Ne, Np impurity percentage Fig. 6. Optimalfunctions of Ne, for n- andp-type alloys Np impurity concentration for nbased on Pb-Te. andp-type alloys based on Si-Ge. (7)

(fly h )n,p are right sides of equations (4), i/f=(y/i, y/^n.p is pulse vector which is conjugate to the vector of phase variables >'=(^r,^„,;,. To use this method, the system of equations (4) must be complemented by the functions relate material parameters a, a and K to the impurity concentration N„,p and temperature. The more accurate the definition of these relations, the higher the accuracy of the Special investigations have been conducted based on our empirical data and the data from the world literature that allowed to approximate the fiinctions of an,p, an,p and Kn,p by the method. These relations are dictated by the energy spectrum of material, microscopic constants of substance and the character of current carrier scattering, polynomials. The system of equations (4)-(7) with polynomials has served as the basis for the program of search for the optimal inhomogeneity functions of the materials of interest for us. Examples of the optimal inhomogeneity functions obtained by computer-aided design method are shown below. Fig. 4 shows an example of the optimal concentration fiinctions for Bi-Te alloys, Fig. 5 - for Pb-Te alloys. Fig. 6 - for Si-Ge alloys. One of the dependencies can be obtained within 10-15 minutes of work of even not very high speed computer of the type IBM-486, which proves the efficiency of the developed program. Naturally, each specific case of FGM use requires individual function of material inhomogeneity, otherwise the use of FGM will be inefficient. The influence of deviation from the optimal functions on the thermogenerator efficiency has been analyzed and it has been found that deviation fi-om the motion of the optimal function leads to considerably fast reduction of FGM method efficiency. Therefore, for each specific case of FGM use it is important to find its individual optimal function of inhomogeneity, hence to make the inhomogeneous thermoelectric material for this case. The next important step of computer-aided design is the development of programs for the cascade thermoelectric generators. The program is complicated by the necessity to reach mutual match of the optimal inhomogeneities in each cascade at simultaneous finding of the best working temperature ranges for each of the cascades. In terms of the optimal control theory it means the introduction to the previous equations of additional relationships for the intercascade temperatures and coordinates of the intercascade junction points which are as follows:

506

2Wi{^>c)- SV^i{x,), n,p

k=

l,...,N-l

(8)

n,p

(9) '**.r/^0''»cni3 p«.^xio;.cm

«^«^10"", ;jsMv^lO", On the basis of equations (4)-(9) a program has cml cm" ^^^^ developed for the design of cascade thermoelectric generators with any number of cascades. Design according to this program results in the determination of the optimal matched inhomogeneities for each cascade, as well as in the determination of the optimal working temperature ranges for the materials of each cascade. It is quite obvious that no direct experimental methods seem to give these results. Fig. 7 shows an example of a three-cascade thermopile design made of FGM. In the cold cascade of the thermopile the alloys of 5 / Te have been used, in the medium-temperature 300 K 1300 K cascade - the alloys of Pb-Te, in the high Fig. 7. Optimal FGMfor a three-stage temperature cascade - the alloys of Ge-Si. It is generator. seen from the graphs that computer-aided experiment provides us with rather unexpected changes in the concentration: in addition to the monotonous concentration increase which could be predicted qualitatively, one can also observe functional bends caused by the optimization of simultaneous contradictory processes occurring in the inhomogeneous legs. The methods of computer simulation have been used to design 50 generators optimized for different hot temperature values in the range of 300^1300 K. The results of these investigations are shown in Fig. 8. In the figure one can see the existence of rational temperature ranges for which 0 200 400 600 800 Ar,K it is advisable to use a one -, two- or three-cascade Fig. 8. The efficiency ofN-cascade gegenerator: The results of these investigations also nerators with FGM material as function determine the efficiency values of the generators of temperature drop. Tc=300K that can be made of FGM. It is seen that at cold temperature of 300°C and the temperature drop of 1000 degrees one can expect the efficiency of 19%. It should be borne in mind that these efficiency values are realistic enough, since the program took into account thermal intercascade losses, as well as the losses in the thermopile contact resistances. The value of contact resistance ro =10"^ Ohmcm^ was accepted for calculation. The results obtained testify in favour of FGM use both for thermoelectric cooling and generation. One can expect that further progress in thermoelectricity will be closely associated with the use of the inhomogeneous thermoelectric structures. ^

507 In conclusion one must note that the search for further qualitative improvement of thermoelectric material is not restricted to creation of FGM structures formed by the monotonous macroscopic material inhomogeneities. Of exceptional interest is creation of the microscopic inhomogeneities with quantum well formation, as well as the inhomogeneous structures allowing to combine on the microscopic level both thermoelectric and emission ways of thermoelectricity generation in a solid body. These possibilities were discussed in detail at the VII International School on Thermoelectricity and published in the "Journal of Thermoelectricity" [5,6]. References 1. loffe A.F. Semiconductor thermal elements, Moscow, 1960. 2. Anatychuk L.I., Semenyuk V.A. Optimal control of semiconductor material and device properties, Chemivtsy, 1992. 3. Nishida LA. Highly efficient thermoelectric materials in FGM program. In Proc. of JapanRussia-Ukraine Int.Workshop on energy conversion materials, Japan (1995) 1. 4. Pontryagin L.S., Boltyanski V.G., Gamkrelidze R.V., Mishchenko E.F. Mathematical theory of optimal processes, Moscow, 1976. 5. Melnichuk S.V., Kosyachenko S.V. Thermoelectric properties of superlattices (Review), J. of Thermoelectricity No. 3 (1996) 15. 6. Ravich Yu.Is. Exceptional thermoelectric properties of cluster superlattices based on opal, J. of Thermoelectricity No. 3 (1996)43.

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I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

509

Anisotropic carrier scattering in n-type Bi2Te2.85Seo.15 single crystal doped with HgBr2 LJ. Ohsugia, T. Kojimaa, H.T, Kaibe^, M. Sakata^ and LA. Nishida^ «Salesian Polytechnic, Igusa 2-35-11, Suginami-ku, Tokyo 167, Japan ^Tokyo Metropolitan University, Minami-osawa 1-1, Hachioji-shi, Tokyo 192-03, Japan. <^Kyorin University, Shinkawa 6-20-2, Mitaka-shi, Tokyo 181, Japan <*National Research Institute for Metals, Sengen 1-2-1, Tshukuba-shi, Ibaraki 305, Japan Anisotropic thermoelectricity of n-type Bi2Te2.85Seo.15 single crystal doped with HgBr2 is discussed on the basis of a multivalley model and an anisotropic scattering mechanism. Temperature dependence of the anisotropy parameters of the six-valley model are calculated from observed galvanomagnetic data. It is confirmed that below 200 K, a weakly anisotropic band ( the first conduction band) is dominant, while above 200 K, a strongly anisotropic band (the second conduction band) also contributes the conduction. Anisotropic scattering parameters are determined from observed thermoelectric powers. It is concluded that conduction electrons are scattered by not only acoustic phonons but also optical phonons, and that the effect of optical phonons in the c-direction is greater than that in the a-direction. 1. INTRODUCTION Bismuth telluride is known as a thermoelectric material particularly useful for low temperature segments of the FGM thermoelectric module[l]. In single crystals of n-type Bi2Te3, anisotropy is observed in thermoelectric powers as well as galvanomagnetic parameters, because of its uniaxial crystal symmetry R3m[2]. It is known that the anisotropic galvanomagnetism can be analyzed on the basis of the six-valley model proposed by Drabble et al.[3-5]. The anisotropy in thermoelectric powers, however, cannot be deduced from the six valley model. On this compound, attempts have been made to analyze the transport properties

510 by introducing a mixed carrier scattering mechanism involving the optical phonon scattering[6]. Kutasov and LukVanova have reported the carrier concentration dependence of the "effective" scattering parameter [7]. The purpose of this work is to discuss the anisotropic thermoelectricity of ntype Bi2Te2.85Seo.15 single crystal on the basis of anisotropic carrier scattering.

2. EXPERIMENTAL PROCEDURE AND RESULTS The starting material was prepared by mixing weighed amounts of 99.999% pure Bi, Te and Se and doped with 0.06% HgBr2. The mixture was encapsulated into an evacuated Pyrex tube, and melted in a rocking furnace to form a homogeneous ingot free from segregation. The ingot was crushed into flakes of about 5 mm by 5 mm, before encapsulated into a Pyrex crucible with 1.33 X10^ Pa hydrogen. A boule of about 16 mm diam and 50 mm length was grown by the Bridgman method with a growing rate of 0.6 mm/h and a temperature gradient of 26 K/mm. The boule included several large crystal grains growing in the direction of soUdification. A single crystal sample was cut out of the largest grain. The crystal was confirmed to be a single crystal by the back Laue method, measurement of the resistivity distribution, and polarized microscope observation. In Bi2Te3 single crystals, the general relationship between the electric field E, the current density J, and the magnetic field H is expressed, in terms of galvanomagnetic tensor components, by Ei = PijJj-PijkJjH, ^p^^JjH.Ht +• • •,

(1)

where p^, /O^^, and p^^ are the resistivity, the Hall coefficient and the the magnetoresistivity tensor components, respectively. Since the crystal structure of Bi2Te3 belongs to the symmetry group R3m, the Unearly independent components are AI> Pas* Piu^ Aan Ani* Am* Am* Assa* Asm A132' A211 and P2323 (^1^® subscript "3" refers to the c-axis). The galvanomagnetic tensor components were measured by the high-speed d.c. method[l] to prevent the error caused by the Peltier effect. The Hall coefficients and magnetoresistivity tensor components were measured in a magnetic field of 0.5 T. Strong anisotoropies were observed in the resistivities and the Hall coefficients as shown in Fig.l. The resistivity ratio P33 //?„ and the Hall coefficient ratio A321 / Ai3 were approximately constant under 200 K, while above 200 K p^2\ / A13

511

tended to increase slightly with temperature.

10'

-J^Q

300 290

Ca__^^06

£ 10^

n

:

P321

O

:

P213

a u

u

"e

^

10-^

10',6,

390 29Q

19Q •

_asi_

: 10"^

-. Q.

Q.

>^ 'a 10^

10^ '^

CO

HH

(D 0

'SCo D

U

0

)

CD

CO - J — I — I

10^

I

I

5

Q

I

10

C

00

'^

10"^

:

-ri—'

5

Temperature lOV T (K-^)

^

1

1 J_

10

Temperature 10^/ T

Fig. 1

I

I

I

(K~M

Fig. 2

Temperature dependence of the Hall coefficient and resistivity tensor components.

^300

_Q

1

'

1

f—

Temperature dependence of the magnetoresistivity tensor components.

I

Q^

QL

^^^ftM

"2001

0

o

-X^^

100

••

-^^

: lOiil : IQ33I

(T3

a.

CD O

e 0

H

CD CD

1

100

,

1

200

Temperature T

i_

1

300 (K)

100 200 300 Temperature T (K)

Fig. 3

Fig. 4

Temperature dependence of the thermoelectric power tensor components.

Temperature dependence of the parameters u,v,w and /? of the Six-Valley Model determined from the observed galvanomagnetic tensor components.

512 Figiire 2 shows the temperature dependence of the observed magnetoresistivity tensor components. The ratios among the observed components were approximately constant below 200 K, while a sUght change was observed in the ratio of p^^^^ to the other components above 200 K. The observed temperature dependence of the anisotropic thermoelectric power is shown in Fig.3. The anisotropy was negUgible below 200 K, while above 200 K the anisotropy increased with temperature. DISCUSSION In the six-valley model proposed by Drabble et al.[3-5], the equi-energy surface of a valley is expressed by

E = £"o +

(ankn^ +ank2i^ +^33^:33^ -{-lankiki)

(2)

,

2/wo

where E^ is the minimum energy of the valley, mo the rest mass of a free electron, k^ (/ = 1,2,3) the wave number vector components, and « „ , a^^, a^^ and a^ the configuration

factors of the equi-energy

surface.

In this model,

the

galvanomagnetic tensor components should satisfy the following relationships: A3 ^ Ai

(3a)

2v

1

—1

1

I

r—

P213

(w + z/v)(l + i/) 4z/v

(3b)

e u

(3c)

(l + «)' - 1

-0.5

bo C u(U

AiAiii _ (>«'-5«M' + 3tn' + t<^vXl + ")

AiAi

!

-

0 A21

;

}-<

CtJ U CO

(3d)

•

: ri

•

: r3

j

-^

~ Hf J » -

j

^

-1 1

100

_l

1

l _

1

200

300 Temperature T (K)

Fig. 5 A1A311 A213

(l + t<)^(*«' + »v) 16/9W/^

1

(3e)

Temperature dependence of the anisotropic scattering parameters determined from the observed thermoelectric tensor components.

513 PnPu22 _ (3w--mv-mv4-3i/^v)(l-m) pj

^^'

2v a'(l + M)

(30

where

a=

4wv (w + in^Xl + w)

(4)

The non-dimensional parameters w, r and vr are defined by

U-GC\\I(^22^

^ = ^33/^22' ^^^ ^^^ = (^22^33 •" ^23^)7^22^ • The factor j8 is the function of the scattering parameter r and the reduced chemical potential r? expressed by

p-

(2r + 3/2)'F(2r + l/2,7)^ (r + 3/2X3r + 3/2)F(r +1/2, ;/)F(3r +1/2,7) '

(5)

where F{r,rf) is the Fermi-Dirac integral defined by

F(r,ri) = ]

-dx . J exp(x-;7)+l

(6)

The values of u, v, w and J3 were determined from the observed galvanomagnetic tensor components by the aid of the least square method as shown in Fig. 4. All the parameters, especially u, showed an inflection at 200 K, which suggest change in the configuration of the eqvii-energy ellipsoid. The thermoelectric power of n-type BiaTea is expressed by Q{r,r})=-

(2r + 5)F(r + 3/2,/7) (2r + 3)F(r + l/2,;7)

-n

(7)

Assuming the anisotropy in r i s very small, the value of n can be estimated from the determined j3 value and the mean value (2Q„ +0^^)/^ by using Eqs.(5) and (7).

Each component of the anisotropic thermoelectric power is expressed by

Q(r>,t?) = -

(2r,+5)F(r,+3/2,»7) -fj l(2r,+3)Fir, + l/2,t})

where r, is the component of an anisotropic scattering parameter.

(8)

The value of

514 r^ can be calculated from the observed thermoelectric powers, substituting the estimated r? value into Eq.(8), The calculation results are shown in Fig.5. At low temperatures below 200 K, the anisotropy in the scattering parameter is negUgible and r^^r^^w -0.7, which is slightly smaller than that in the pure acoustic phonon scattering case (r = - l / 2 ) . It is concluded that the sUght difference is due to the optical phonon scattering, because the observed temperatiu'e range includes the Debye temperature of Bi2Te3 (165 K)[8] and ionic bonding exists between Bi and Te atoms[9]. In the high temperature range above 200 K, r, approaches t o - 1 / 2 , while r^ decreases with increasing temperature. It is known that the second conduction band exists 30 meV above the first conduction band[10]. From these facts, it is concluded that the effect of the optical mode is greater in the c-direction than that in the a-direction, and that the second band is more anisotropic than the fij*st band, which is consistent to the report that the anisotropy in thermoelectric power increases with increasing amount of dopant[2]. ACKNOWLEDGMENTS Thanks are due to Prof. Ohta of Keio University for his helpful discussions. REFERENCES [1] K. Uemura and LA. Nishida, Thermoelectric Semiconductor and Its AppUcations, Nikkan-Kogyou Shinbun , 1988 (in Japanese). [2] H.T. Kaibe, T. Okumura, M. Sakata, Y. Isoda and LA. Nishda, Proc. the 10th Int. Conf. Thermoelectrics, Cardiff, U.K.,p.35, 1991. [3] J.R. Drabble and R. Wolfe, Proc. Phys. Soc. London, 69 (1956) 1101. [4] J.R. Drabble, R.D. Groves and R. Wolfe, Proc. Phys. Soc. London, 71 (1958) 380. [5] J.R. Drabble, Proc. Phys. Soc. London, 72 (1958) 380. [6] M. Stordeur, Phys. Stat. Sol, (b), 98 (1980) 199. [7] V.A. Kutasov and L.N. Luk^anova, Sov. Phys. SoUd State, 28 (1986) 502. [8] G.E. Shoemake, J.A. Rayne and R.W. Ure, Jr., Phys. Rev., 185 (1969) 1046. [9] D.R. Lovett, Semimetals & Narrow-Bandgap Semiconductors, Pion, London, p. 182, 1977. [10] R.B. MaUison and J.A. Rayne, Phys. Rev., 175 (1968) 1049.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

515

Percolation design of graded composite of powder metallurgically prepared SiGe and PbTe R. Watanabe^, M. Miyajima", A. Kawasaki^ and H.Okamura^ a Department of Materials Processing, Faculty of Engineering, Tohoku University, Sendai 980-77, Japan b Graduate Student, Faculty of Engineering, Tohoku University, Sendai 980-77, Japan

The design and the powder metallurgical fabrication of functionally graded thermoelectric material have been studied with specific interest to enlarge the working temperature range and increase the figure of merit. The percolation control of the electrical resistivity of the composite phases has been taken into account in the design scheme of the graded composite. The materials combination of SiGe and PbTe was selected as a model system for the verification of the percolation concept. The measurements were carried out on the thermal conductivity, electrical resistivity and Seebeck coefficient, which are involved in the figure of merit, of non-graded composites, as well as of the monolithic SiGe and PbTe. 1. Introduction To enhance the performance of the thermoelectric generating device, it is important to enlarge the working temperature range of the device. Toward this goal, functionally graded thermoelectric material [1,2] has a great potential. As the figure of merit of thermoelectric materials has remarkable temperature dependence, we have to use several materials along the temperature gradient to achieve the best performance of thermoelectric conversion. We have two large problems on preparing such sturucture i.e. segment bonded sturucture [3]. The first problem is the residual thermal stress at the material interface. And the second problem is the electrical properties of the material interface. For the furst problem, recent progress in the field of stress relaxation type functionally graded material helps us very well, and the principle is utilizing multi phase composite material. On the other hand, for the second problem, it is not desireble to introduce complicated interface geometry, for example muti phase composite structure Unfortunately, there are not very many works in the field of composite thermoelectric material. One of the rare clear theoretical works by Bergman and Levy [4] tells us that any multi phase composite in which unclusions are dispersed in matrix phase shows smaller figure of merit. But, as they modeled, this calculation does not contain cotinuous microstructure networks, which appears very often in functionally graded materials. We would like to see the continuous network structure effect, i.e., percolation property [5-11] of composite thermoelectric material. The following section 2 describes basic numerical modeling of percolative mateials, and the section 3 describes on experimental study of SiGe / PbTe graded composite.

516 2. Percolation Model 2-1 Percolation on thermoelectric material Percolation phenomena is characterized by large scale transition of physical property as a function of composition. The characteristic composition at which the percolation transition occurs is called percolation point. Generally, high electrical resistivity material has rather small thermal conductivity. So we may hope in thermoelectric material composite, preparing composite with high and low electrical resistivity material, large thermal conductivity reduction, whereas keeping rather low electrical resistivity, as shown in Figure 1. These are the key concept of introducing percolation microstructure. 2-2 Microstructure and equivalent electric circuit model To model the microstructure and evaluate the thermoelectric properties, we used following simple equivalent electric circuit model shown in Figure 2. We considered the two phase composite as a cluster pararrel network circuit. Setting for each cluster the characteristic single phase physical property, and settle the material composition to the cluster number ratio, we can simulate the total thermopower of the system by Millman's theorem of d.c. circuit. Numerical results of this model at fixed temperature are shown in Figure 3. The figure of merit are shown as a function of SiGe / PbTe composition. Physical properties for each materials are from ref [2]. As shown in Figure 3, such microstructure that contains SiGe property, which is modeled as high thermal conductivity / low electrical resistivity phase, for large composition range, shows better thermoelectric property. For such materials, only the thermal conductivity is effectively reduced by containing PbTe phase. On the contrary, such microstructure that is dominated by PbTe property, which is modeled as low thermal conductivity / high electrical resistivity phase, which has rather small percolation point, shows smaller figure of mertis compared to the former material model.

!5- «

\

^ K

\^^

•fi| ^

x^

3

]

T3

a8 D

13 'B

p

N.

'0>. O'^

V3 *oo

i^

quality enhancement available ....

l / / j

^ Phase A percolation point, Pc Phase B Compositon of phase A / B composite

_ £ai/pi

a "" Zl/pi

Figure L Schematic diagram of percolation Figure 2. Schematic diagram of equivalent transition of transport property and its use on electrical circuit model of phase A/ B (suffix a andb) thermoelectric composite material. thermoelectric composite material.

517 2-3 Modeling graded two phase thermoelectric composite In the next step, we would like to consider using such composites to prepare graded thermoelectric material. In Fig.4, results of simple numerical evaluation are presented. These curves present the basic features of percolation designed graded thermoelectric composite, that is, we can bridge the valley of figure of merit of different thermoelectric materials, by controling the percolation point and the transition width ( here we used the form of well-known Fermi-Dirac distribution function to model the gradual percolation to let the transition complete one half at ther percolation point) via microstructure control. The model parameters in Figure 4, however, do not include the composite phase interface geometries, explicitly. This is not enough for practical designing of graded structure and to be considered in further step.

l.U

^ ^ 0.9 ' ^ ^

2l.50

1

h = 0.2 (a) Pc=70 (b) Pc=50 V (c) Pc=30

P c = 3 0 \ p c = 5 0 \ p c == 7 o \ 0.7

~ Z(PbTe) \

n^

1

0

\ 1

1

25 50 75 100 PbTe Phase Fraction, Vf / vol.%

Figure 3. Thermoelectric figure of merit of SiGe/PbTe composite numerical model.

A = 0.2 A = 2.0 (a) Pc=70 (d) Pc=70 (b) Pc=50 (e) Pc=50 (c) Pc=30 (f) Pc=30

I

o

4.50

r^^

1

Z(PbTe)

m \^ \ ycc) 1

11.00 h

^^^y

|Z(SiGe)

tL,

Z(SiGe)

2.00 h=1.0 (d) Pc=70 /^ (e) Pc=50 / 1 (f) Pc=30 /

fe^^^W

N •a

T = 900K

\

0.8

2.00

I

numerical model with percolation point Pc [vol.%] / ' ^ r \

4g

^ ^

, \ Z(PbTe)

0.50

1000 800 900 Temperature, 7 / K (1) thermal conductivity reduction and figure of merit of numerical model with percolation point Pc (vol.% PbTe), where :

700

/c=ph

PbTe + (l-ph)/CsiGe

0.50

700

800 900 1000 Temperature, 7 / K (2) percolation transition width and figure of merit of numerical model with percolation point Pc (vol.% PbTe), where : P =f' P PbTe + ( ! - " / ) • P SiGe / =1.0/(exp(Pc-p)/A+1.0)

Figure 4. Parametric consideration for the figure of merit of the numerical model of SiGe/ PbTe graded thermoelectric composite.

518 3. Fabricating SiGe/PbTe graded composite 3-1 Aim of experiment Altiiough die numerical evaluation shows die possible thermoelectric property enhancement, there are not any appropriate experimental data on thermoelectric properties of composite. The aim of the following experimental procedure and some results is to collect basic data for composite thermoelectric materials design. 3-2 Experimental procedure The experimental procedure so far we have tried is described in Figure 5. The starting non-dope powders were prepared by High Purity Chemicals Co. Ltd. The impurities for conduction type adjustment were added via mechanical alloying, using steel tumbler mill, in argon atmosphere.

Starting powder Si-20at%Ge, PbTe ( 3N) Prealloyed / non-doped

I

Doping (via mechanical alloying) SiGe: B(p-type), P(n-type) PbTe: Ag2Te(p-type), Pbl2(n-type) Powder blending SiGe-30,50,70vol%PbTe Die compaction 100 MPa CIP 200 MPa I HIP in vacuum glass capsule 1123K / Ih / 200 MPa (Argon)

3-3 Results The thermoelectric properties of doped and Evaluation sintered samples are presented in Table 1. The p \ room temp.( Van der Pauw) electrical resistivity of the composite made by AC : 2 9 3 -- 1073 K ( laser flash) those doped powders become very high, XRD, EPMA a \ room temp. compared to monolithic phases. So far, we consider the results to have been caused by Ge diffusion between SiGe/ PbTe interface. By EPMA analysis, 0.9-3.5at% of Ge were probed Figure 5 Possible procedure of fabricating in PbTe phase. After the experimental work monolithic SiGe, PbTe and also the graded by T.Abakumowa [12] such Ge atoms could thermoelectric composite. work as donor site in PbTe lattice. The thermal conductivity of the composites showed modest monotonous reduction with increase of second phase (PbTe) composition.

I

Table 1 Mechanically doped impurities and the thermoelectric properties of sintered monolithic SiGe and PbTe test samples used to prepare SiGe / PbTe composite.

thermoelectric properties

Seebeck coefficient ( fi VK'^) Electrical resistivity ( 0 cm) Carrier density (cm'^)

SiGe p-type

n-type

B 0.5at%

P 0.5at%

171 135 5.417E-3 3.407E-3 6.514E+19 1.372E+20

PbTe p-type Ag2Te 2.0 mol%

n-type Pbl2 2.0 mol%

59 156 3.136E-4 9.217E-4 1.504E+19 5.333E+19

519 4. Summary Application of percolation conduction model on thermoelectric composite material has been considered using simple numerical model. The possible thermoelectric property enhancement were illustrated on SiGe / PbTe composite and graded structure model. The model, however, does not include composite microstructure geometry explicitly, so far. The possible experimental procedure and some results were also described. With simple powder metallurgical experiments on preparing composite SiGe/PbTe mixed phase material, the following results were obtained : (1) The electrical resistivity of the composite made by those doped powders become very high, compared to monolithic phases. We consider the results to have been caused by Ge diffusion between SiGe/ PbTe interface. (2) The thermal conductivity of the SiGe/PbTe composite showed modest monotonous reduction with increase of second phase (PbTe) composition. Acknovpledgement This work is partly supported by a grant from the research project in the Science and Technology Agency of Japan on the Development of Functionally Graded Materials for Energy Conversion. Reference 1. G.D.Mahan, J. Appl. Phys.,70 (1991) 4551. 2. T.Hirano, L.W.Whitlow and M.Miyajima, Ceramic Transactions, Functionally Gradient Materials, Ed. by J.B.Holt, M.Koizumi, T.Hirai and Z.Munir, The American Ceramic Soc, 34 (1993) 23. 3. F.D.Rosi, J.P.Dismukes and E.F.Hockings, Elec. Eng., 79 (1961) 450. 4. D.J.Bergman and O.Levy, J. Appl. Phys., 70 (1991) 6821. 5. S.Kirkpatrick, Rev. Mod. Phys., 45 (1973) 574. 6. R.Landauer, AIP conference proc., No.40 (1978) 2. 7. J.P.Straley, J. Phys. D, 14 (1981) 2101. 8. D.Stauffer, Introduction to Percolation Theory, Taylor & Francis, (1985). 9. O.Levy and D.J.Bergman, J. Phys., A25 (1992) 1875. 10. F.Lux, J. Mat. Sci., 28 (1993) 285. 11. R.Watanabe, J.Takahashi and A.Kawasaki, 3rd Ind. Symp. on Structural and Functionally Gradient Materials, Lausanne, Ed.by B.Ilshner and N.Cherradi, Pressis polutechniques et universitaires romandes, (1995) 3. 12. T.A.Abakumowa, et al, Neorg Mater, 30 (1994) 1121.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

521

Design of Multi-Functionally Graded Structure of Cylindrical RI Heat Source for Thermoelectric Conversion System S.Ainada% J.Terauchi^ and T.Senda^ ^Dep. of Mech. Eng., Gunma University, 1-5-1, Tenjin, Kiryu, Gunma, 376, JAPAN ^Ship Research Institute, 6-38-1, Arakawa, Mitaka, Tokyo, 181, JAPAN The graded structure with the multiple functions is required to construct a radioisotope (RI) heat source for a thermoelectric conversion system. The first function is to get the surface temperature of the heat source as possible as high and the other must have a radiation shielding function. This structure is a cylinder, composed of RI-SrTiOg as a heat source and BN as a radiation shielding material. The composite Rl-cylinder must be designed to maxmize the surface temperature and to minimize the radiation intensity. It is presented that an optimum graded distribution of RI-SrTiOg exists to satisfy two distinct requirements. 1. INTRODUCTION Substitute energy developments are indispensable problem in our generation. As one of those candidates, RI thermoelectric conversion system is noticed [1,2]. This system is based on Seebeck effect which transforms temperature difference into voltage difference directly. In order to increase the efficiency of RI thermoelectric conversion system, not only a development of high efficient thermoelectric conversion materials but also a Rl-heat source structure must be done. The Rl-heat source must satisfy two requirements, those are an effective radiation shielding property and keeping the operating temperature below the melting point of heat-source material. Under these conditions, the surface temperature of Rl-heat source must be increased as high as possible. As Rl-heat source, we selected SrTiOg, a chemical compound of Sr-90 extrated from the waste nuclear fuel, and boron notride (BN) which controls to generate bremsstrahlung r -ray. Assuming that the RI heat source

522 is constructed from a cylinder mixed SrTiOg and BN, its optimum multi-functionally graded structure is studied to satisfy two distinct requirements, high surface temperature of the heat-source and low dose^, which leads to an efficient thermoelectric conversion system. 2. ANALYTIC METHOD 2.1. Temperature analysis The steady state heat conduction equation with heat source in cylindrical coordinate system is given by 1 d>. dT

1 dT d^T g(r) ^ + —5- + -2l2 = o (1) A, dr dr r dr dr X where T is temperature, r is radius, A is the temperature-dependent thermal conductivity and g is the heat generation rate. The estimated formula of the thermal conductivity of the composite material is given by Maxwell-Eucken formula [3]. +

2.2. r -ray shielding analysis Sr-90 is /3 -decay and its daughter nucleus Y-90 is also /3 -decay. Since jS -ray interacts with the coulomb field of the atomic nucleus, it radiates the secondary radiation, called bremsstrahlung [4,5]. Radiated r -ray has a continuous energy spectrum (0.11 '^ 2.09 MeV). The radiation shielding analysis is focused on the bremsstrahlung r -ray in this study. Assuming that the radiation source is an area radiation one accumulating of point isotropic sources, the distrbution of radiation energy is approximated by mono-energy (1 MeV). It is also assuming that there is no buildup (B=l). Putting the gamma-ray evaluating point P on the distance 30 cm from the central axis of the Rl-cylinder, the r -ray dose is calculated by •SB 2 exp(-i;|iiti)rdrde

(2)

where (f) is the r -ray flux at the point P, r is the distance from the point of radiation source to the point P, S is the r - r a y strength, B is the buildup factor, p. is the r -ray attenuation coefficient, t is the thickness of the shield. 2.3. Analytic model We selected a hollow cylindrical model. The base structure consists of two-layers of SrTiO, 100% and BN 100% as shown in Fig.l. This structure has a power 50 kW

523

with the heat generation rate of RI-SrTiOg 2.3 W/cm^ Let us look for the graded structure to optimize two functions high surface temperature and low dose based on the base structure. The composition distribution is represented by a modified probability density function based on un-symmetric T distribution function [6].

f{r) = r(a + l)Z?^ ^ { - r + 0.0l(r2-ri)ci}'exp

-r + 0.0l(r2-ri)q (3)

where F is the F function, r is radius, r^ the inner sufiface and r^ the outer surface radius. l{f(r)>l,f(r) equals to 1. Eq.(3) is defined as the graded composition distribution function. This function is not only symmetrical, but also r e p r e s e n t s various distributions by changing parameters (a,b,c^c^) as shown in Fig.2. In this study, focusing the parameter a of four parameters, and we call it the gradient parameter. We analyze the temperature and the radiation dose by changing the graFig.l Base structure (two-layer dient parameter. hollow cylindrical structure). 2.4. Evaluation method Let the temperature performance index r) ^ be defined by eq.(4) to estimate an improvement of the temperature performance. "^^

T^-T,

(4)

where T^and T^ are the surface temperature of the graded and the base structure, T^ is the melting point of Rl-cylinder (2313 K). Next, the r -ray shielding performance index 77 is defined by eq.(5) to estimate an improvement of the radiation shielding performance. Tly=l-*/K

(5)

where (t> and (f) ^ are the r -ray flux of the graded and the base structure. Let us introduce the integrated performance index 77 . defined by eq.(6) to estimate by combined i] ^ and 77 ^ . r|i = a-r|T+(l-a)-r|^ where a is the weight coeficient.

(6)

524 3. RESULTS AND DISCUSSIONS Fig.3 shows the rerationship between the temperature and the r -ray performance index r? ^, 7? ^ and the gradient parameter a. Generally, moving the distribution of SrTiOg to outside, the temperature performance increases. Oppositely, moving the distribution of SrTiOg to inside, the r -ray shielding performance increases. There is a trade-off relation between them. Accordingly, it is not enough only one-functional design of the graded structure for the Rl-heat souce. The graded structure of the Rl-heat souce must be designed by integrating two distinct functions. Assuming that the optimum design demands n =0, the optimum gradient parameter lead to a=2.2, and the temperature performance index n ^=0.39. It shows that the temperature performance is raised 39 % as compared with the base structure to make the graded structure, which means that the surface temperature increment increases 500 K. Its optimum graded structure and temperature distribution are shown in Fig.4 and Fig.5. According to these figures, the optimized structure leads to the gentle temperature gradient and higher surface temperature. In case of changing the weight coefficient a which specifies a weight of temperature and dose in the design, the integrated performance index curves are shown in Fig.6. The gradient parameter a which gives the maximum value of each curve corresponds to the optimum gradient parameter a^^^. Fig.6 shows that there is the optimum gradient parameter which obeys the specification of the design, and it exists the optimum graded structure. The relationship between the weight coefficient a and the optimum gradient parameter a^^^ is shown in Fig.7. It was shown that a^ ^ is smaller as a is larger which corresponds to the design weighted on the

Temperature performance index TJ

0.5 0.0

r ''i

0.01(r^-rj)c^ ^2

Radius, r Fig.2 Graded Structures (a=0.1, 1.0, 2.0, 3.0).

1

2 3 4 5 Gradient parametera, a

Fig.3 Variations of temperature and r - r a y shielding performance index (b=4, c^: =86).

525 2500 |Meltingpoint(2313K)

^ 100^ ^ c o

Base structure

80-

1 60Co

E

40

S

20

CO

structure

^ 2000 4 I— 1500

Optimum graded structure

Base structure

I 1000 E" ^

500 4

OH

04 10

25

15

0

Radius, r / cm

10

15

20

25

Radius, r / cm

Fig.5 Temperature distributions of base structure and optimum graded structure (a=2.2, b=4, Cj=86).

Fig.4 SrTiOg volume fractins of base s t r u c t u r e a n d o p t i m u m graded structure (a=2.2, b=4, c^=86).

surface temperature. That is to say, the design puts emphasis on the temperature performance, as the distribution of RI-SrTiOg moves to outside. The temperature gradient becomes smaller because the SrTiOg layer with low thermal conductivity can be thin to keep a constant volume as the layer moves to outside. At the same time, the r -ray shielding performance deteriorates because the radiation source moves to outside and the shielding layer is thinner. Deciding the weight coefficient a to satisfy the requirement, we can get the optimum graded structure for each a as shown in Fig. 8. Compared with the temperature performance, the radiation shielding performance was not improved in this study whose analytic model is two-ingredient system of SrTiOg and BN. Its maximum value is + 7 ^ 8 % . The reason is why BN can control producing the bremsstrahlung r -ray, because it is light material, but cannot shield ?r"~ 1 . 0

§• 4

0=0.9

I 0.5 03

S 0.04 c

//x\

CO

1-o.s

rr/\\

2.-1.0

j

^^^"^

2 4T

^Q3

0)

\ a=0.1"; E iE o. O

? -1.5 \ = -2.04 0

O

i 1

i 2

1

i

3

4

5

Gradient parameter, a

Fig.6 Integrated performance index curves ( a =0.1, 0.3, 0.5, 0.7, 0.9).

0

0.2 0.4 0.6 0.8 1.0 Weight coefficient, a Fig.7 Relationship between optimum gradient parameter aopt and weight coefficient Of . 0.0

526 radiated r -ray. It is expected that the radiation shielding performance may be considerably improved by two phase shielding. The first shielding must be done by a light material like BN located at the radiation souce in order to restraint producing the b r e m s s t r a h l u n g r -ray, and t h e second shielding by a heavy material (ex:Pb,W,etc.) located at the outer layer in order to shield radiated r -ray. The design of multi-ingredient graded structure is a future problem.

uu-

"•'^^oc=0.1

80-

(x=0.5"'"->i pa=0.3

600=0.7

4020-

cx=0.9

oJ

— 1 —

5

""'

— r"^*^*****!^— 10 15 20 Radius, r / cm

'"i"

Fig.8 Optimum graded structures ( a = 0 . 1 , 0.3, 0.5, 0.7, 0.9).

4. CONCLUSIONS We got the following conclusions. 1) Since there is a trade-off relation between the temperature performance and the r -ray shielding one, the optimum graded design of only one function is not enough. 2) Assuming that the optimum design demands 7? =0, the surface temperature increment increases 500 K, the temperature peraformance could improve 39%. 3) Moving the distribution of RI-SrTiOg to outside, the temperature performance increases. Oppositely moving, the r -ray shielding performance increases. 4) The graded distribution of the Rl-cylinder is effective for a high performance of the heat source. REFERENCES 1. K.Uemura and I. Nishida, Thermoelectric Semiconductor and Application, NIKKAN KOGYO SHIBUN, (1988), 13, (in Japanese). 2. The Science and Technology Agency, The Development Project of High Efficient Thermoelectric Conversion System by Using Functionally Graded Materials: Survey Report, The Society of Non-traditional Technology, (1993),(in Japanese). 3. A.Maezono, Ceramics, 29, (1994), 421, (in Japanese). 4. J.R.Lamrash, Atomic Nucleus Engineering Guide, SACHI-SHOBOU, (1982), 42, (in Japanese). 5. A.Ohashi et al., Procd. Functionally Graded Material , The Society of Non-traditional Technology, (1994), 143, (in Japanese). 6. T.Indou, Probability and Statistics, CORONA-SHA, (1957), 80, (in Japanese).

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

527

Fabrication of N-type Poiycrystallinc Bi-Sb and Their Thermoelectric Properties M.Miyajima^, G.G.Lee", A.Kawasaki^ and R.Watanabe^ a Daikin Industries Ltd. MEC Laboratory, Graduate Student, Faculty of Engineering, Tohoku University, Sendai 980-77, Japan b Korea Institute of Machinery and Metals, Changwon Industrial Complex, 66, Sangnam-dong, Changwon, Kyungsang-namdo, Korea c Department of Materials Processing, Faculty of Engineering, Tohoku University, Sendai 980-77, Japan We report on the powder metallurgical fabrication of bismuth-antimony solid solution and the thermoelectric properties of the fabricated composites. The solid solution powders were prepared by mechanical alloying (MA) aiming at large reduction of the thermal conductivity with the very fine microstructures obtained through MA process. The prepared bismuth-antimony powders (Bi-7.5at%Sb) have been sintered by hot pressing. The thermoelectric properties of the sintered Bi-Sb compacts were influeneced by the MA conditions. With the development of very fine grain microstructure, the thermal conductivity was reduced to the level of 2.0 W/mK which is about the 40% of the sintered compacts made from melt-solidified (MS) Bi-7.5at%Sb powder. 1. Introduction The mechanical alloying ( MA ) method has been applied in thermoelectric materials research [1-3] as this process has such practical merits that : (a) a very wide range of alloy compositions is available (b) multi phase composites can be fabricated and especially (c) because of the building up process characteristics of mechanical alloying, sintered compacts from MA powders, generally have very fine crystal grains and very large total sum of grain boundary. The long wave length phonons which carry the heat energy in semiconductor are mainly scattered at grain boundaries, so, we can expect thermal conductivity reduction and hence the thermoelectric figure of merit enhancement, from preparing materials with mechanical alloyed powders. The present work reports on mechanical alloying process of the n-type Bi-Sb [4-10] solid solution which is known to be the best material in low temperature region. The effect of milling conditions on the microstructure development and the thermoelectrical properties were studied. 2. Experimental procedure The experimental procedure is summarized in Figure L (1) mechanical alloying The composition of Bi-Sb solid solution was fixed to 7.5 at% Sb which is the lower limit of

528 the n-type semiconducting transsport region [7,8]. The blended Bi and Sb powders whose purities are 5N were mechanical alloyed using stainless tumbler mill and carbon steel milling ball. The weight ratio of filled powder to the milling balls was 1 / 62. (2) sintering of sample powders The milled powders were compacted in a metal die at a pressure of 100 MPa, then cold isostatic pressed at 200MPa. The size of the green compacts were 15 mm in diameter and 4 mm in height. The green compacts were sintered by hot pressing in a alumina die under a pressure of 70 MPa at the temperature of 543K, for sintering time of 3600 s. Ar gas flow containing and 5% hydrogen was used as the sintering atmosphere. (3) phase and microstructure evaluation The alloying process was observed through the changes of the X-ray diffraction patterns of milled powders. And the microstructures of the sintered compacts were observed by TEM. The samples for TEM observation were prepared by Ar ion milling. (4) measiu*ement of the thermoelectric properties The thermal conductivity of the sintered Bi-Sb compacts were measured by laser flash method at room temperature. The direction of the heat flow was parallel to the compacting direction during hot pressing. The sample size was 10mm in diameter and 1mm thick. The electrical resistivity and the Hall coefficient were measured using Van der Pauw method at 77K and 300 K. Both of the properties were meassured along the plains perpendicular to the compacting direction during hot pressing. The sample size was 3mm X 4mm and 0.5 mm thick rectangular shape. 3. Results and Discussion 3-1 Phase analysis and microstructures The x-ray diffraction patterns of mechanical alloyed Bi-7.5at%Sb powders are shown in Figure 2. The diffraction peaks of Sb almost vanished afrer 86.4ks and the profile showed formation of Bi-Sb solid solution. The diffracdon patterns did not show large peak broadening with longer milling time. The sum of the impurities amount from the milling vessels such as Fe, Ni, Cr were less than 300ppm for each milUng time sample powders. In Fugure 3, the TEM micrographs of the sintered

STARTING MATERIAL Bi powder 5N, d < 150|im Sb powder 5N, d<150|im

MECHANICAL ALLOYING Composition Bi-7.5at.%Sb Milling: Tumbler MiU(sus304) Ball diameter 12.7mm Vessel size 70mm(t)x 135mmL Powder / Ball weight ratio 1 / 62 Atmosphere Ar Rotate speed llOrpm Milling time max 720ks

COMPACTION DIE press 100 MPa, CIP 200 MPa Compact size: 15mm(t) X 4mm height

HOT PRESSING Compacting pressure 70MPa Sintering atmosphere Ar+5%H2 Sintering temperture 543K Sintering time 3600s

EVALUATION Phase analysis : X-ray diffraction Composition : ICP Microstructure : TEM Electrical properties : Resistivity Hall coefficient (Van der Pauw : 77K,300K) Thermal properties: Thermal conductivity (Laser flash : 300K) Figure 1. Experimental procedure

529 compacts are presented. The microstructures of the sintered Bi-7.5at%Sb compacts show the difference of mechanical alloying conditions, that is, the milling time of powders. With increase of the milling time length from 86ks to 172ks, the grain size are kept in the range of 0.1 to 0.5 jxm. With further increase of the milling time, to 360 ks, the grain size become finer to less tiian 50 nm. Further increase in milling time, however, resulted in larger grain size, of several |im. As the sintering temperature and sintering time were fixed for all the samples, these differences in microstructure formation might be related to the recrystalization process during sintering. As the amont of the strain energy differs with milling time,the driving force of the recrystallization process for each powder might have been varied, though the difference of the strained energy should not be so large, as the peak broadening of the X-ray patterns in Figure 2 are not so evident.

Figure 3 TEM micrographs of mechanical alloyed and hotpressed Bi-7.5at%Sb. Milling time: (a) 86ks (b)172ks (c) 360ks Figure 2. Change of X-ray diffraction pattern (d)720ks. Sintered at 543 K for 3600 s of Bi-7.5at%Sb powder during ball milling. in Ar+5%H2. Sintering pressure 70 MPa. 30 40 50 60 70 Diffraction angle , 28 / degree

530 3-2 Thermoelectric properties and milling condition The Hall mobilities of the sintered Bi-7.5at%Sb compacts from MA powders at 300K and at 77K are shown as a function of milling time. The values of the sintered compact from melt solidified (MS) powders are also presented. The mobilities were very much reduced by nicclianical alloying compared with the compacts from MS powders. But with increase in milling time, the mobility monotonously increased. This results might be related to the amount of lattice defects, such as lattice vacancies, interlattice atoms, dislocations, which affects the scattering of conduction electrons. As the mobility showed increase, these lattice defects should have been decreased with longer milling time. One of the possible interpretation for these results, is that enhaced recrystallization during sintering effectively removed the lattice defects in MA synthesized Bi-Sb solid solution . The thermal conductivities are presented in Figure 5 as a function of milling time. 10000 • The thermal conductivity decreased rapidly o Bi.7.5%SbT=300K with with increasing milling time first, • Bi-7.5%SbT=77K L however, with further increase in milling Milling time t =0 : 7500 time, the thermal conductivity increased. compact from melt solidified As the main carrier of heat at BOOK is considered to be long wave length phonon, It 5000 L powder which is scattered mainly by grain [ boundary, the milling time dependence of y^ • thermal conductivity is considered to be in o 2500 good coincidence with the microstructure B fonnalion shown in Figure 3. \^^^t/f^^^^\ 1 The figure of merits at 300K are shown 600 800 200 400 in Figure 6. With reduction of the thermal Mining time, //lO^s conductivity and recovery of the mobility by long time milling, the compacts from Figure 4 Hall mobility of Bi-7.5at%Sb MA powders have higher figure of merits compacts as a function of milling time. than the compacts from MS powders. 1

T

5.0 c^

1

101-3 Bi-7.5%Sb T=300K Milling time t =0 :

compact from melt solidified powder

N •J"

I

Bi-7.5%SbT=300K Milling time t =0 : compact from melt solidified powder

\/

101-4 200 400 600 Milling time, //lO^s

800

Figure 5 Thermal conductivity of Bi-7.5at%Sb compacts as a function of milling time.

,

1

i

i__

1

1

200 400 600 Milling time, //lO^s

800

Figure 6 Fgure of merit of Bi-7.5at%Sb compacts as a function of milling time.

531 4. Summary Bi-7.5at%Sb thermoelectric semiconductor powders were prepared by mechanical alloying. The following results were obtained. (1) Bi-7.5at%Sb solid solution powders were successfully prepared by mechanical alloying. (2) The sintered Bi-7.5at%Sb samples showed submicron size microstructure. The grain size varied with the change of milling time lengh in mechanical alloying process. And these difference in microsturucture resulted in different thermoelectric properties. The figure of merit of the compacts from MA powders were better than that of the compacts of melt solidified and pulverized powder samples. References 1. B.Cook, B.Beaudry and J.Harmga, Proc. 9th ICT,(1990),28. 2. K.Hasezaki, M.Nishimura, M.Umata, H.Tsukuda and M. Araoka, Proc. 12th ICT,(1994),307. 3. K.Pixius, W.Wunderlich, J.Schilz and W.Amend, Phys. stat. sol. (a), 146(1994),109. 4. A.L.Jain, Phys. Rev., 114(1959),1518. 5. S.Tanuma, J. Phis. Soc. Jpn., 14(1959), 1246. 6. G.Smith and R.Wolfe, J. Appl. Phys., 33(1962),841. 7. H.J.Goldsmid , phys. stat. sol. (a), 1(1970),7. 8. W.Yim and A.Amith , Sol.-St.Electron.,15(1972),1141. 9. R.B.Horst and L.R.Williams, Proc. 3rd ICT, (1980),139. 10. Y.Suse,Y.H.Lee JI.Morimoto,T.Koyanagi,K.Matsubara and A.Kawamoto, Proc. 12th ICT, (1994),248.

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I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 1997 Elsevier Science B.V.

533

Development of Functionally Graded Thermoelectric Materials by PIES Method A.Yamamoto and T.Ohta Energy Fundamentals Division, Electrotechnical Laboratory, AIST, MITI, Umezono 1-1-4 Tsukuba Ibaraki 305, Japan An attempt was made to develop compositionally graded p-type Bi-Sb-Te thermoelectric materials by PIES (Pulverized and Intermixed Elements Sintering) - Hot Press method. The sample developed here consisted of three segments of different alloy composition. The electrical output of the sample under temperature difference of 200K showed an asymmetry with forward and reversed temperature gradient. 1.INTRODUCTION Thermoelectric energy conversion (TEC) is one of the potential candidate for waste heat recovery path in energy system designing. Its high reliability and maintenance-free operation are big advantages over conventional energy conversion technologies. Bismuth telluide based alloy is the most typical material for thermoelectric application in temperature range from 300K to 500K. Over the past thirty years a considerable number of studies have been made on bismuth telluride based alloy and up today the highest figure of merit ZT= 0.96 have been achieved for p-type single crystal[l]. From the view point of application thermoelectric material is usually used under large temperature difference. The figure of merit of the material should be large in the whole temperature range so that the average figure of merit value is large. The figure of merit value of bismuth telluride based alloy , however , strongly depend on temperature and have a maximum point at a certain temperature. Hence there is a great possibility to increase the conversion efficiency by flattening the temperature dependence of the figure of merit ZT. Since the temperature giving the maximum figure of merit is controled by varying a doping level and / or composition of the alloy , an enhancement of the average figure of meritZJis expected with a sample where such parameters are gradually changing. Here we can adopt a concept of functionally graded marterials (FGM). In this study an attempt was made to prove such positive effects on thermoelectric properties by preparing and characterizing compositionally graded thermoelectric materials.

534

2.EXPERIMENTAL 2.1 PIES method for sample preparation PIES method (Pulverized and Intermixed Elements Sintering method) is one of the most promising sample preparation techniques for mass production of thermoelectric materials. Figure 1 shows the flow diagram of sample preparation by PIES method. The method is featured with two simple processes. The first one is ball-milling process where raw powder elements are intermixed and pulverized by mechanical energy input. In this process, the powder forms a solid solution without any temperature increment. The second one is sequent sintering process where the powder is compacted in a temperature much below its melting point by using hot-pressing (HP), hot isostatic pressing (HIP) or plasma activated sintering (PAS). Samples made by PIES method usually have very fine grains which favor a reduction in the thermal conductivity and an enhancement in the mechanical strength. PIES method does not require an intricate process such as sealing in a quartz ampoule, melting in Rocking furnace and grinding and sieving which all are used in conventional melt techniques. P-type bismuth telluride based materials prepared by PIES method have already proved its potential as same level as conventional melt samples and the process has been employed in many other material systems[2-7]. In this study 99.999% purity bismuth, antimony and tellurium powdered elements were used as starting materials. Hot pressing method was chosen as a sintering process. Hot pressing conditions were 753K, 400kgf/cm^, for one hour. Typical sizes of hot-pressed specimen were 10mm(() X 5mm and sliced into 1mm thick disk specimen. The density of specimen measured by Archimedes' method were more than 98% of the ideal value. 2.2 Measurement The Seebeck coefficient were calculated from measurement of electromotive force with temperature difference of lOK. The electrical resistivity and Hall measurement were performed by van der Pauw method. The thermal conductivity were calculated from the thermal diffusivity, the specific heat and the density. The thermal diffusivity and the specific heat were measured by laser flash method and differential scanning calorimeter (DSC), respectively.

Bi

c;b

T

Te

T Pulverizing and intermixing

Hot Pressing

Raw materials (elements)

Agate Ball-mill 20 hrs., Aratomsphere

Graphite dies 753 K, Aratomsphere Ann Ifnflnm ^

1 hr. TTieimoelectric specimen 1 {B\^Te^),_^ (Sb2Te3)y

Figure 1. Procedure of sample preparation by PIES method

535

2.3 Preparation and characterization of FGM sample In this study FGM samples were prepared to prove the enhancement in output characteristics by introducing compositionally graded structure. Through the study on a compositional dependence of thermoelectric properties , segmentation structure of y = 0.80 / 0.825 / 0.9 was selected where y indicated the antimony telluride fraction in (BiJ^e^)^ (Sb^Te^) system. In the experiment samples which have compositionally graded structure was prepared by piling several powder materials of different composition at a charge process into graphite dies before hot-pressng. This technique was performed quite well since no cracking was observed at the interfaces inside the sample and there were not any unfavorable effects on mechanical strength. Nickel electrodes were soldered at both sides of the sample and then it was cut into 2mm X 5mm x 5mm. Energy Dispersive X-ray spectroscopy (EDX) observation was performed to confirm the compositional distribution of the sample. The sample was fixed between a heater and a heat sink which apply a temperature difference of 200K to the sample. Current lead and voltage lead were attached to nickel electrodes and the V -1 characteristic was measured by varying an external load. 3.RESULTS 3.1 Thermoelectric properties Figure 2 shows the compositional dependence of carrier density. As shown in the figure, carrier density strongly depended on composition of the samples. The samples of composition y > 0.7 showed p-type conduction while those of y < 0.7 showed n-type conduction. This result is quite different from the results of typical bismuth telluride based materials which is made by melt process[8]. All n-type samples had 1.5 low Hall mobilities around 100 cmWs r) I (Bi2Te3Vy(Sb2T%)^ therefore only p-type compositions are "E 1.0 in selected to further study on thermoelectric 0 9>\ 0.5 properties. = P-type ^ ^ 0 ^ Figure 3 , 4 and 5 are temperature ^W 0 -_ c : n-type © dependence of Seebeck coefficient , the -0 0.5 1 @ i resistivity and the thermal conductivity as a S CO 1.0 — 0 M function of composition. All the samples 0 lo @ here are of p-type composition. The lines in • 11 1.5 0 0.2 0.4 0.6 0.8 1 Figure 3 and 4 vary in order of composition Composition y y and clearly each peak of line shift to higher temperature range as composition y ^-^^^^ 2. Carrier density - composition 1

1

1

'

1

•

1

1

1

I

1

1

1

1

1

1

1

1

L

536 4.0

300 y=0.8

, o o •-' u o o ,

> =1 250 c0 .2 200

y=o.8o

3.0 • ««• •

0.825

• #o

o

^nDnn°Bi|B! h

o0 ^ 150

0.875 ^ Q O

.

• •

ii"

I I I I I t I I I I I I I I I I I I I I

350

400

450

350

Seebeck coefficient

1

D C

•D

o o

1.3

/

[ L

0.875 \

r

0.9

L

/

1.2

L

- •- '

o 0

.-•''

«

K

t

r °"^ 8 2 5 r

f

0.8 h /

0.6

^•^ /

^^^

500

/

^^ 0.875

^

„

\— 0.7

/•/

^

Ft 0 . 8 2 5 " ^ • ^ - / - . -

1.1 k

^

" ° ^Q _ - c r / ^

h

[-

0.9

//°

"^

450

Figure 4. Temperature dependence of resistivity

1.5

^

400

Temperature /

Figure 3. Temperature dependence of

>

i l " " 0.9

I I I I I I I I I I I I I I I I I I I I I

300

500

Temperature / K

1.4 ^

iB iB

0

100 300 I

Bi

0.875g|

•I 10

0.9

0.825»**

2.0

E

- t •

o°

r *

y = 0.8

\

^^

0.5

L

0.4

ri ^,.1,,. , 1 , , , , 1 , , , , 1 , , , ^ 1

y=0.8

1 r 1 , , , . 1,„i t ^ l i 300 350 400

i x-iJ„ 1 i - o ^ l l

450

500

Temperature / K

Figure 5. Temperature dependence of thermal conductivity.

•^

300

350

400

450

500

Temperature / K

Figure 6. Temperature dependence of figure of merit.

increase. This peak shift may account for change of carrier density which is strongly depended on composition y. Figure 6 shows the temperature dependence of dimensionless figure of merit ZT value which were calculated from data shown in Figure 3 , 4 and 5. The highest figure of merit value 0.88 was obtained at 340K in sample of y = 0.825. There are overlaps between each temperature dependence. This means a possibility of an enhancement in conversion efficiency through FGM designing. For example , as seen in Figure 6 a compensative combination of curves of

537

y= 0.825 and y=0.9 will produce larger average ZJvalue in the temperature range from BOOK to 500K. 3.2 FGM sample Compositionally graded thermoelectric sample was obtained by using PIES / hot-press method. The sample consisted of three segments , y = 0.8 , 0.825 and 0.9. Figure 7 shows the bismuth and antimony contents measured by EDX along the direction of temperature difference to be applied. Bismuth to antimony ratio changed gradually tracking the target composition. Figure 8 shows power output characteristics of FGM thermoelectric sample shown in Figure 7. Voltage-Current plot and Power-Current plot show the difference in the direction of temperature gradient applied to the sample. Output power with forward temperature gradient is 6% larger than that with reversed temperature gradient. The most likely explanation of this asymmetry can be found in a graded structure of the sample. But there is room for argument on this result because there seems a problem on reproducibility of both the material properties and the soldering technique. In this study the expected enhancement on conversion efficiency is rather small, it seems necessary to solve the problem mentioned above.

40

>

.1 I I I I I I I I I I I I I I I I I I I I I I I I I I I I

35

"

E

v30

-"^ CD

25

> "D

•• 1

CO

"o

• • •

20

. °—» •

15

y=0.825

y=0.9

y=0.8

Figure 7. Bithmus and antimony content measured by EDX.

15 10

o

Upper

20

E %

o Q.

•+-•

O

30

D forward : • reversed 25

10

5

5

0

Q.

"3 O

I I I I I I I I I I I I I I I I I I I I I Igfif^.

0.0 0.5 1.0 1.5 2.0 2.5 3.0

Current / A Figure 8. V-I and P-I characteristic of FGM sample.

538 4.CONCLUSIONS In this paper preparation and characterization of compositionally graded thermoelectric element were discussed. The sample preparation was successfully performed by using PIES / hot-press method and the figure of merit showed a maximum value of ZT= 0.88 for y = 0.825 at 340K. The figure of merit Z J showed a strong dependence on its composition and design and preparation of compositionally graded sample was successfully performed. Output characteristic of the FGM sample was measured under practical operating condition , AT=200K. The electrical output for forward temperature gradient was 6% higher than that for reverse one. This asymmetry for forward and reverse temperature gradient seems to stem from functionally graded structure of the sample. REFERENCES l.T.Caillat et.al., Mat. Res. Soc. Symp. Proc. Vol. 234 (1991), 189. 2.T.0hta et.al., The 8th Proc. Int. Conf on Thermoelectrics , Nancy , (1989), 7. 3.T.0hta et.al., The 13th Proc. Int. Conf on Thermoelectrics , Kansas City , (1994), 267. 4.T.0hta et.al., The 14th Proc. Int. Conf on Thermoelectrics , St. Petersburg , (1995), 24. 5.T. Caillat et.al., The 11th Proc. Int. Conf on Thermoelectrics , Arlington , (1992), 240. 6.T.0hta et.al., The 11th Proc. Int. Conf on Thermoelectrics , Arlington , (1992), 74. 7.B.A. Cook et.al., The 9th Proc. Int. Conf on Thermoelectrics , Pasadena , (1990), 234. 8.J. Shim et.al., The 9th Proc. Int. Conf on Thermoelectrics , Pasadena, (1990), 27.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

539

Microstructure and thermoelectric properties of p-type Bio.5Sb1.5Te3 fabricated by hot pressing Doo-Myun Lee, Jun-Ho Seo, Kyeongsoon Park^, Ichiro Shiota^, Chi-Hwan Lee Dept. of Metallurgical Engineering, Inha University, Inchon 402-751, Korea ^Dept. of Materials Engineering, Chung-ju National Univ., Chungbuk 380-702, Korea Dept. of Chemical Engineering, Kogakuin National University, Tokyo 192, Japan The p-type Bio.5Sb1.5Te3 compound doped with 4.0 wt% Te was fabricated by the hot pressing at the temperature range 380 to 440 °C under 200 MPa in Ar. The microstructure and thermoelectric properties of the compound were investigated. Optical microscopy, scanning electron microscopy, and X-ray diffraction were used to examine the microstructure. The microstructure was relatively dense. The density was increased with increasing the pressing temperature. The grains were preferentially oriented through the hot pressing and also the degree of preferred orientation was increased with the pressing temperature. It was also found that the figure of merit was increased with increasing the hot pressing temperature. The highest figure of merit (2.69 x 10"'^/K) was obtained at 420 °C. 1. INTRODUCTION Bi2Te3-based compounds are well known to be good thermoelectric materials for the applications near room temperature. The crystal structure of Bi2Te3 at room temperature is rhombohedral (a=0.438 nm and c=3.049 nm) [1]. This crystal is composed of atomic layers in the order of Te/iSi/Te/Te/ Bi/Te/Bi/Te/Te/ along the c-axis. The Te/Te layers are considered to be weakly bound with van der Waals forces [2]. The crystal has distinct cleavage planes perpendicular to the c-axis. Owing to the cleavage feature, the crystal has low mechanical properties and poor ability in micro-processing for fabricating the miniature thermoelectric modules and is inappropriate for mass production of thermoelectric modules. Many attempts were made by sintering to fabricate miniature modules without cleavage. However, sintering technique is not effective because the figure of merit of sintered compounds is lower than that of single crystals. In this work, we fabricated the p-type Te-doped Bio.5Sb1.5Te3 compound by the hot pressing and then investigated the microstructure and thermoelectric properties of the compound. 2. EXPERIMENTAL PROCEDURE To fabricate the p-type Bio.5Sb1.5Te3 compound doped with excess 4.0 wt% Te, the starting powders with >99.99 % purity were mixed. The powders mixture was placed into Si02 tube with 25 mm diameter and 330 mm length. Then, the tube was evacuated below 10" torr and sealed. The powders

540

mixture was heated at 700 °C. The melt in the tube was stirred under a frequency of 5 times/min at 700 °C for 6 hours using a rocking furnace to make a homogeneous melt without segregation. The tube containing the melt was cooled in furnace. The solidified ingot was crushed into fine flakes using AI2O3 bowl. The resulting flakes were ball milled for 12 hours and then sieved to prepare powders with 45-74 fm size. To remove the oxygen developed during the crushing and ball milling, the resulting powders were reduced in hydrogen atmosphere at 380 °C for 4 hours. The powders were compacted by the hot pressing at the temperature range 380-440 °C at steps of 20 "C under 200 MPa in Ar to produce the billets with 30 mm diameter and 6 mm length. The density of the hot-pressed compound was measured by pycnometer (Micrometric Co.). The compound for optical microscopy was etched with a solution of HN03*H20=1:1. The preferred orientation of grains was investigated by X-ray diffraction (XRD). The thermoelectric properties were measured at room temperature along the direction perpendicular to the pressing direction. The samples with dimensions of 2 x 2 x 1 5 min and of 4 x 4 x 4 mm were cut out of the compound for the measurements of Seebeck coefficient a and thermal conductivity /c and of the electrical resistivity p, respectively. Then, their surfaces were polished with a series of SiC polishing papers of up to #2000 and further polished on a polishing cloth impregnated with AI2O3 powders of 0.3 //m size. To measure the Seebeck coefficient a, heat was applied to the sample which was placed between the two Cu discs. The thermoelectric electromotive force (E) was measured upon applying small temperature difference ( J T < 2 "O between the both ends of the sample. The Seebeck coefficient a of the compound was determined from the E / J T . The electrical resistivity p of the compound was measured by the four-probe technique. The repeat measurement was made rapidly with a duration smaller than one second to prevent errors due to the Peltier effect [3]. The thermal conductivity K was measured by the static comparative method [3] using a transparent Si02 ( K =1.36 W/Km at room temperature) as a standard sample in 5 x 1 0 torr. 3. RESULTS AND DISCUSSION 3.1 Microstructure It was found that the p-type compound was relatively dense. The density was increased with increasing the pressing temperature because of the porosity decrease. The decrease results from an improvement in the bonding between the powders. We could not fabricate successfully the compound at 440 °C because of the local melting of the powders. The melt was identified as Te used as a dopant. Fig. 1 shows the optical microstructures along the longitudinal and transverse directions for the compound hot pressed at 420 °C. The dark areas shown in Fig. 1 correspond to the pores. The porosity present in the compound was decreased with the pressing temperature. To investigate the orientational change of grains depending on the pressing temperature, XRD analyses from the perpendicular and parallel sections to the hot pressing direction were made. Fig. 2 (a) and (b) show the XRD patterns obtained from the perpendicular and parallel sections, respectively, for the compounds hot pressed at 380, 400, and 420 °C. The intensity of the (0 0 15) and (0 0 18) planes from the perpendicular section is much stronger than that from the parallel section and is increased with increasing the pressing temperature. Also, the intensity of (0 0 6) plane is only observed at the perpendicular section. The (0 0 6), (0 0 15), and (0 0 18) planes are perpendicular to the c-axis. This indicates that the grains are preferentially oriented through the hot pressing and also the degree of preferred orientation

541

Fig. 1. Optical micro structures along the (a) longitudinal and (b) transverse directions for the compound hot pressed at 420°C

liT 380 X)

o o!2 OO

N

lii*iiiiiiiA

400*0 c

I LuJjJLaXjJ 42013

JJ\.J„ 20

30

XiiiU^^ 40

2e (DegrM)

20

30

40

26 (D«grM)

Fig. 2. XRD patterns obtained from the (a) perpendicular and (b) parallel sections to the hot pressing direction for the compounds hot pressed at 380, 400, and 420 °C is increased with the pressing temperature. It is thus expected that the thermoelectric properties will be improved with increasing the temperature owing to the increase in density and preferred orientation. It has previously been reported that the preferred orientation of grains in unidirectionally solidified materials is observed and the growing direction is perpendicular to the c-axis.

542 3.2 Thermoelectric properties Fig. 3 shows the carrier concentration nc and mobiUty ju as a function of the pressing temperature. With increasing the pressing temperature, the carrier concentration and mobility of the compound are decreased and increased, respectively. The increase in mobility results from the porosity decrease. The variation of Seebeck coefficient a with hot pressing temperature is shown in Fig. 4. As the temperature is increased, the Seebeck coefficient is slightly increased because of the

> CM

E b

Hot Pressing Temperature (ic)

Fig. 3. Carrier concentration nc and mobility // as a function of the hot pressing temperature.

decrease in carrier concentration. The relationship between the a and nc can be expressed as follows- a ~ r-ln nc, where r is the scattering factor [4]. The variation of electrical resistivity

300

250

>

200 h

150

380

400

420

Hot Pressing Temperature ("C)

Fig. 4. Variation of Seebeck coefficient a with hot pressing temperature.

380

400

420

Hot Pressing Temperature Cc)

Fig. 5. Variation of electrical resistivity p with hot pressing temperature.

p with hot pressing temperature is shown in Fig. 5. As the hot pressing temperature is increased, the electrical resistivity of the compound is decreased. The electrical resistivity can be expressed as the following relationship: p=l/nce/jt. As a consequence, two competing factors, carrier concentration and mobility, determine the electrical resistivity. Therefore, it seems that the decrease in electrical resistivity with increasing the pressing temperature would result from a significant increase in mobility and a slight decrease in carrier concentration. Fig. 6 shows a plot of thermal conductivity K VS. hot pressing temperature. The thermal conductivity is increased with the pressing temperature probably because of the density increase. Fig. 7 shows the figure of merit Z of the compound hot pressed at various hot pressing temperatures. The figure of

543

merit is increased with the pressing temperature because of the decrease m porosity and increase in preferred orientation. The compound hot pressed at 420 °C shows the highest figure of merit (Z-2.69X10 /K).

380

400

420

Hot Pressing Temperature ("c)

Fig. 6. Relationship between the thermal conductivity K and hot pressing temperature.

380

400

420

Hot Pressing Temperature (x;)

Fig. 7. Figure of merit Z of the compounds hot pressed at various hot pressing temperatures.

4. CONCLUSIONS The hot pressed p-type doped with 4.0 wt% Te Bio.5Sb1.5Te3 compound was found to be densified up to 99.0 % of theoretical density. The hot pressmg gave rise to a preferred orientation of grains and also the degree of preferred orientation was increased with increasing the pressmg temperature. It was also found that the figure of merit was increased with mcreasmg the pressing temperature due to the porosity decrease and the preferred orientation increase. The compound fabricated at 420 °C showed the highest figure of merit (2.69x10 7K).

REFERENCES 1. R. W. G. Wyckoff, Crystal Structure Vol. 2, Interscience Publishers, New York 1964 2. J. R. Weise and L. Muller, J. Phys. Chem. Solid, 15 (I960) 13 ^^^.^ .^ 3 A Goudot, P. M. ScWicklin, and J. G. Stockholn, 5th ICTEC (1984) 49. A K Uemura and I. Nishida, Thermoelectric Semiconductors and Their Application, Nikkan-Kogyo Shinbun Press, Tokyo, Japan, 1988.

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I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

545

Microstructural and thermoelectric properties of hot-extruded p - t y p e Bio.5Sb1.5Te3 Jun-Ho Seo, Doo-Myun Lee, Kyeongsoon Park^, Jong-Hoon Kim^, Isao A. Nishida"^, Chi-Hwan Lee Dept. of Metallurgical Engineering, Inha University, Inchon 402-751, Korea ^Dept. of Materials Engineering, Chung-ju National Univ., Chungbuk 380-702, Korea Korea Academy of Industrial Technology, Shiheung 429-450, Korea ^National Research Institute for Metals, Tsukuba 305, Japan The p-type Bio.5Sb1.5Te3 compound doped with 4.0 wt% Te was fabricated by the hot extrusion at the temperature range of 300-510 °C under an extrusion ratio of 201 and a ram speed of 5 cm/min. The microstructure of the compound was investigated by scanning electron microscopy and X-ray diffraction. The microstructure was highly dense and fine-grained (~1.0 fim). The hot extrusion gave rise to a preferred orientation of grains. With increasing the extrusion temperature, the bending strength and figure of merit were increased due to the porosity decrease. The highest bending strength (92 MPa) and figure of merit (2.94xiO'7K) were obtained at 440 °C. It is proposed that the hot extrusion reduced a grain size and increased a density, resulting in an improvement in bending strength and figure of merit. 1. INTRODUCTION Bismuth telluride (Bi2Te3) compound has been used as thermoelectric cooling and heating materials, since it has a high figure of merit (2.5xlO~/K 3.0 X10' /K) at room temperature and can be fabricated easily and chiefly. Many workers have reported on the fabrication and thermoelectric properties for the compound. The compound has a rhombohedral structure (a=0.438 nm and c=3.049 nm) and belongs to space group R3m. The electrical and mechanical properties are higher along the two equilivalent directions parallel to the (001) cleavage planes than the c-axis [1, 2]. Since the compound has the easy cleavage planes, it has difficulty for the mass production of small thermoelectric modules. The grain refinement and mechanical properties improvement may avoid the easy cleavage features. In this work, we attempted to refine the grains and to improve the mechanical properties by means of the hot extrusion. 2. EXPERIMENTAL To fabricate the p-type Bio.5Sb1.5Te3 doped with excess 4.0 wt% Te, the starting powders with >99.99 % purity were mixed. The powders were

546 compacted by the hot pressing at 420 "C and 200 MPa in Ar to produce the billets with 30 mm diameter and 60 mm length. Subsequently, the compacted billets were hot extruded at the temperature range 300-510 °C at steps of 70 "C under an extrusion ratio of 20:1 and a ram speed of 5 cm/min. The density of the compound was measured by pycnometer (Micrometric Co.). The preferred orientation of grains for the compound was investigated by X-ray diffraction (XRD). The mechanical properties were measured at room temperature under a crosshead speed of 0.5 mm/min by three-point bending in accordance with ASTM D790 using a universial testing machine. The thermoelectric properties were measured at room temperature along the direction parallel to the extrusion direction. The samples with dimensions of 2 X2X15 mm and of 4 x 4 x 4 mm were cut out of the compound for the measurements of Seebeck coefficient a and thermal conductivity K and of the electrical resistivity p , respectively. Then, their surfaces were polished with a series of SiC polishing papers of up to #2000 and further polished on a polishing cloth impregnated with AI2O3 powders of 0.3 pm size. 3. RESUTLS AND DISCUSSION 3.1 Microstructure and mechanical properties At the hot extrusion temperatures of 300-440 °C, we obtained good extruded bars without any defects such as tearing, orange peel, and blister. However, hot cracks were developed during the hot extrusion at 510 "C. This would result from the local melting due to the heat formed by the friction between the billets and die. The relative density of the compound was increased with increasing the extrusion temperature. This increase occurs due to the decrease in the porosity because of an improvement in the bonding between the powders. The highest relative density was obtained at 440 °C and its value was 99.6 % of theoretical density. The XRD patterns obtained from the compound hot extruded at 440 °C are shown in Fig. 1. Fig. 1(a) and (b) in (a) show the XRD patterns obtained from a. 1 the perpendicular and parallel sections to the hot extrusion direction, respectively. The intensity of (0 0 6), (0 0 15), and (0 0 18) planes, which are perpendicular to the c-axis, is only observed at the parallel section. (b) This indicates that the hot extrusion gave rise to a preferred orientation of s 0 grains. The bending strength was 0 ^ & increased with increasing the extrusion 0 e temperature. The increase in bending strength results from the porosity 50 60 decrease. The bending strength of the 2 e (Degree) compound hot extruded at 440 °C was 92 MPa. The fractograph of the compound hot extruded at 440 °C is shown in Fig. 2. The fractograph Fig. 1. XRD patterns obtained from transgranular cleavage the (a) perpendicular and (b) parallel represents sections to the hot extrusion. features. The fracture path follows transgranular cleavage planes. The

UJiii

547

,^>f\c-',*.

•yft.

iGjLOB

orientation change from grain to grain was also found. The grain size estimated from this fractograph is ~ 1.0 fM. We believe that the hot extruded compound with high strength and small grain size leads to an improvement in the bonding strength between the thermoelectric materials and metal electrode during soldering for the fabrication of thermoelectric modules and leads to a good ability in micro-processing for fabricating the miniature thermocouples for the semiconductor devices.

Fig. 2. Fractograph of the compound hot extruded at 440 °C.

3.2 Thermoelectric properties Fig. 3 shows the carrier concentration Uc and mobility // as a function of the hot extrusion temperature. As the hot extrusion temperature is increased, the charge carrier concentration is decreased and the mobility is significantly increased. The significant increase in mobility occurs due to the porosity decrease. The variation of Seebeck coefficient a with hot extrusion temperature is shown in Fig. 4. This figure represents that the Seebeck coefficient is increased with increasing the extrusion temperature because of the decrease in 300 370 440 Hot Extrusion Temperature (Tc) carrier concentration. The relationship between the a and nc can be expressed as follows^ a ~ r-ln nc, where r is the scattering factor [4]. Fig. 3. Carrier concentration HC and The values of a for the compound hot mobility /u as a function of the extruded at 300 and 440 t are 145.8 hot extrusion temperature. and 231.1 //V/K, respectively. The relationship between the electrical resistivity p and hot extrusion temperature is shown in Fig. 5. As the hot extrusion temperature is increased, the electrical resistivity was decreased. This is because with increasing temperature, the scattering of carriers is decreased due to the decrease in porosity and thus the mobility is increased. The value of p for the compound extruded at 440 °C is 1.85x10"^ Qm.

548

250

^% 6h

^"^^^

200

^ >=1 Q

150

•

X

100 1 300

370

440

300

Hot Extrusion Temperature ("C)

Fig. 4 Variation of Seebeck coefficient a witii hot extrusion temperature.

370

440

Hot Extrusion Temperature (ic)

Fig. 5. Relationship between the electrical resistivity p and hot extrusion temperature.

Fig. 6 shows plots of thermal conductivity K VS. hot extrusion temperature. The thermal conductivity is increased with increasing the temperature. The increase in thermal conductivity would be strongly affected by the decrease in porosity. Fig. 7 shows the figure of merit Z of the compound hot extruded at various hot extrusion temperatures. The figure of merit is increased with the extrusion temperature due to the decrease in porosity. The compound hot extruded at 440 "C shows the highest value of Z (Z=2.94xl0"^/K).

1.2

•

•

1.0

?

0.8

i^

0.6

Jff

0.4

•

0.2 0.0

1

300

370

1

440

Hot Extrusion Temperature ("C)

Fig. 6. Relationship between the thermal conductivity K and hot extrusion temperature.

300

370

Hot Extrusion Temperature ("C)

Fig. 7. Figure of merit Z of the compound hot extruded at various hot extrusion temperatures.

549 4. CONCLUSIONS It was found that the microstructure of p-type Bio.5Sb1.5Te3 compound doped with 4.0 wt% Te was highly dense and fine-grained ( — 1.0 jm). The grains were preferentially oriented through the hot extrusion. The bending strength and figure of merit were increased with increasing the extrusion temperature because of the porosity decrease. The bending strength and figure of merit of the compound hot extruded at 440 °C are 92 MPa and 2.94xiO'^/K, respectively. We believe that the hot extrusion provides enhanced bending strength and figure of merit and is thus a very useful technique for the fabrication of thermoelectric materials. REFERENCES 1. 2. 3. 4.

Y. M. Yim and F. D. Rosi, Solid State Electronics, 15 (1972) 1121. D. M. Rowe, Applied Energy, 24 (1986) 139. A. Goudot, P. M. Schlicklin, and J. G. Stockholn, 5th ICTEC, (1984) 49. K. Uemura and I. Nishida, Thermoelectric Semiconductors and Their Application, Nikkan-Kogyo Shinbun Press, Tokyo, Japan, 1988.

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I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

551

Effect of Dopants on Thermoelectric Properties and Anisotropics for Unidirectionally Solidified n-BiiTca N.Abe', H.Kohri\ I.Shiota', and I.A.Nishida' 'Department of Chemical Engineering, Kogakuin University, 2665-1 Nakano-machi, Hachioji-city, Tokyo, 192, Japan ''Institute for Materials Research, Tohoku University, 2-1-1 Katahira, Aoba-ku, Sendai-city, Miyagi, 980-77, Japan TSfational Research Institute for Metals, 1-2-1 Sengen, Tsukuba-city, Ibaraki, 305, Japan

It is wellknown that thermoelectric properties of n-type Bi2Te3 compounds are affected by dopants. In this work, we used HgBr2, Hgl2, SbBrs and Sbis as the dopants. As a resuh, electrical resistivity p of specimens doped with Sb compounds showed a drop of about 18% as compared with Hg compounds. Thermal conductivity K of specimens doped with bromide showed a drop of about 6% as compared with iodine. The figure of merit Z of specimen doped with Sbis showed the greatest value of 3.88 X lO'^K"^ in these specimens. On the other hands, Bi2Te3 compounds have large anisotropy in the thermoelectric properties because of its crystal asymmetry of R3m.

In this work, anisotropics in hall

coefficient Ru and p of specimens doped with Hg compounds showed about 1.3 times as compared with Sb compounds.

1. INTRODUCTION The Bi2Te3 compound is the most excellent material for thermoelectric cooling materials at around room temperature, and has been widely used in the precise temperature control machines etc. The compound has a layered rhombohedral structure. Then the compound has strong anisotropics on mechanical and thermoelectric properties. The properties vertical to the direction of c-axis are better than parallel[l]. Anisotropy of

552

galvanomagnetism has been investigated by Drabble et al. to deduce expressions for anisotropic galvanomagnetic parameters on the basis of the six-valley model[2][3]. The resistivity along the c-axis is greater than vertical direction to the c-axis, and the hall coefficient measured with a magnetic field in the c-axis is less than that with a magnetic field vertical to the c-axis. HgBr2 has been empirically doped in the practically used n-type Bi2Te3 compounds. Recently, Kaibe et al. reported that Sbis is more effective than HgBr2 as the dopant in sintered n-type Bi2Te2.85Seo.15 [4]. But it is not confirmed which element was effective to improve the figure of merit. In this work, HgBr2, Hgl2, SbBrs and Sbis were used as the dopants. These dopants consist of combination Hg or Sb and Br or I. Each dopant was added in the unidirectionally solidified n-type Bi2Te3 to form a n-type compound. The effects of each element were determined by the combination of the observed values of four specimens. Anisotropics in the unidirectionally solidified n-type Bi2Te3 was also investigated.

2. EXPERIMENTAL PROCEDURE 2.1. Preparation of the specimens Bi and Te of normal purity 99.999% were used as the starting materials. Firstly, an ingot of Bi2Te3 without dopant was fabricated. Proper ratio of Bi and Te were weighed and the mixture of the powder was sealed in a quartz glass tube with 460mmHg Ar. The tubes was stirred sufficiently by using a rocking furnace at 923K, and the melt was unidirectionally solidified by using the Bridgman method with a cooling rate of 2mmh'^ under a temperature gradient of 5Kmm'\ The Rn of the ingot was measured to determine the carrier concentration. Dopant is required to form an n-type Bi2Te3 as Bi2Te3 is p-type. HgBr2, Hgt, SbBr3 or Sbl3 was mixed in the powder which was obtained by pulverizing the ingot. Amount of dopants were determined to adjust the electron concentration of 1.0 X lO^^m"^ at room temperature under the hypothesis that a halogen atom and a Sb atom give an electron and a hole, respectively. The n-type Bi2Te3 doped with each dopant was prepared by same process described above. The obtained n-type Bi2Te3 consisted of a few large crystals, which is very like a single crystal. The ingot consisted of several large crystal grains growing in the direction of soUdification. 2.2. Measurements of thermoelectric properties A sample of 1 X 2 X 7mm\ as shown infig.1 (a) (b), was taken out of the largest grain

553 which consists of a single crystal to measure p and Ru. Temperature dependence of p and Ru were also measured over the temperature range from 80 to 500K where the highest performance can be expected. A specimen of 4 X 4 X 4mm^, as shown infig.1 (c) (d) lower, was also cut out, which is also single crystal. Then the Seebeck coefficient a and K of the specimen were also measured at room temperature by the static comparative method, as shown in fig.2. Afiasedquartz block of the same size as the specimen was used for the standard material of the K. /?l,i^H

1

P//Ifn // Pressure (lOg/cm^)

Heater Thermo couple Specimen 1

Specimen 2 Copper Plate (c)

(d)

Fig. 1 Measurements of thermoelectric properties Subscript of p, Ru, a, and K : direction of measurements to c-axis

Fig. 2 Measurements of K and a

3. RESULTS AND DISCUSSION 3.1.Thermoelectric properties The carrier concentration n^ vertical to the c-axis of each specimen at room temperature were shown in table 1. It was observed that the n^ of each specimen were about 1 X lO^^m"^ equal to expected n^. The temperature dependence of p and Ru in the direction vertical to c-axis of each specimen are shown in fig. 3. The Ru of each specimen shows a same value over the observed temperature range because of the same n^. The Ru of each specimen decreased suddenly at the temperature range more than about 300K. Therefore it is obviously that it

554 Table 1 Thermoelectric properties (at R.T) Ru a [IQ-^Q m] [IQ-VC^I riQ"'mV^s-^] [lO^^m'^] [ptVK'^] [WK'^m'^]

Dopant

[IQ-'K'^]

HgBr2

7.13

2.74

3.84

0.98

-210

1.78

3.44

Hgl2

6.79

2.49

3.67

1.07

-216

1.90

3.61

SbBrs

5.92

2.64

4.46

1.01

-205

1.99

3.57

Sbis

5.46

2.68

4.91

1.00

-212

2.12

3.88

begins to be influenced by intrinsic range around 300K. The p of each specimen increased with increasing temperature up to about 400K, and decreased suddenly at the temperature range over 400K. It is found that the p of specimens with iodide are less than those of specimens with bromide. It is also found that the p of specimens with Sb compounds are less than those with Hg compounds. The temperature dependence of hall mobility \i vertical to c-axis of each specimen is shown in fig.4. The \x of each specimen decreased with increasing temperature over the observed temperature. The [x of specimens with Sb compounds are larger than those of specimens with Hg compounds over the observed temperature range. This phenomena can be considered as follows; Hg ion traps more electrons than Sb ion as reported that Bi2Te3 compounds are the same crystal structure as Bi2Se3 compounds, and Hg atom acts as acceptor in Bi2Se3 compounds[5]. Moreover, Sb atom influences Bi2Te3 directly, and ionic bond become weaker by doping with Sb compounds[6].

t

1

1

1

1

1

1

^m^

1

1

^is

I

1

%i

1

1

.

10-

i '• 1 1 1

0

i 0 0:HgBr2 • :Hgl2 0:SbBr3 • :Sbl3

[ ^

^>*

4 \

IQ-^h

r

4> 1

1

1

1

1

110

; -

CM

E

•k«1 1

0:HgBr2

a:

• •:Hgl2

2:* €>';^

1 1 10" 10 103/T[K-^] Fig.3 Temperature dependence of Resistivity and Hall coefficient

5

-

> 10-5 Q-

^*

1 91

1

1

• 2*

IQ-^b f O

10

1

i

l

l

O o • O

^

0:SbBr3 10" - • : S b l 3 '

•

1

1

1

9 _ • 1

200 400 T[K] Fig.4 Temperature dependence of Hall mobility

555

The a, K and Z vertical to c-axis of each specimens at room temperature are shown in table 1. The a becomes smaller by doping with Sb compounds. The K of specimens with bromide are smaller than those with iodine. It is considered as follows; the ion radius of bromine is smaller than that of iodine and shows a tendency to exist substitutionally in the Bi2Te3 lattice, and the Kiattice which is the lattice component of the K becomes smaller. The K of specimens with Hg compounds are also smaller than those with Sb compounds. It might be caused by weak ionic bond due to doping with Sb compounds as described above. The Keiectron wMch is the Carrier component of the K become larger. Therefore the Z of each specimen doped with iodine or Sb compounds are larger than those of specimens doped with bromide or Hg compounds. The Z of specimen doped with Sbis is the largest, 3.88 X 10'^K'\ because of very high \x. It is obviously shown that Sbis is the most excellent as the dopant.

3.2.Anisotropies The temperature dependence of p and Ru vertical and parallel to c-axis of each specimen are shown in Fig. 5. Anisotropic coefficients of p and Ru of specimens with Hg compounds are about 2.5 and 2.7 at room temperature, respectively. Anisotropic coefficients of p and Ru of specimens with Sb compounds are about 2.1 and 1.9 at room temperature, respectively. From these results, it is found that anisotropics of specimens with Hg compounds are larger than those with Sb compounds. The anisotropic coefficients of a was only 1.1 in both cases of the specimens doped with Hg or Sb dopants,as shown in table 2.

—

1

i—

1

1

1

1

1

10"

'*

1

• ;o'

1 J

10-

8 " -4 E G

iio-

1

1

1

1

5

t

1

1

• . 1

1

, ,10 10^/T[K-^] (a)HgBr2 and SbBr3

1

1

1

1

1

1

i^B" • • • • • • • • A ^ ^ ^ i ^ ^ 412^^

U m

•

6

a

^

1

10-

•^H

: ^ '

o

' ^ ^

CJ

10-

—

1

1

•

1

•

A

A

1

-

10-

• A

^ ^ ^ ' -ho -4 E a D Hgi2 J-: Hgl2// •A Sbl3 ±•

D

^m^ - 10X :

•• ••

1

. ^ A A A A A A AA A

0:HgBr2 -L; • :HgBr2// 0:SbBr3 -L • :SbBr3// _

o ^ 10-

1

10-

- 10

- o

1

* *

RH

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10-

1

• ^ e 8 6 o a ^8 5 e. ^

•

u

1

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;

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• • : « « a -iio1 1 1 1 ^

1

103/T[K-^j

10

(b)Hgl2 and Sbl3

Fig.5 Temperature dependence of Resistivity and Hall coefficient

556 Table 2 Anisotropies of a (at R.T) Dopant HgBr2 Hgl2 SbBrs Sbl3

-210 -216 -205 -212

a//

a i /a//

-191 -200 -183 -194

1.10 1.08 1.12 1.09

4.CONCLUTION n-type Bi2Te3 were doped with HgBr2, Hgl2, SbBrs or Sbls. The carrier concentration n^ of each specimen was highly controlled, and was 1 X lO'^^m"^ at room temperature. The p of a specimen doped with Sbis showed the minimum value, 5.46 X 10'^ Q m. The a became larger by doping with Hg compounds. The K of specimen doped with HgBr2 was the minimum value, 1.78WK"^m'\ The anisotropies of specimens with Hg compounds were larger than those with Sb compounds. The Z of specimen doped with Sbis is the largest, 3.88 X 10"^K'\ because of very high \i. It is obviously shown that Sbis is the most excellent as the dopant.

ACKNOWLEDGMENT We express our appreciation to supporting by Science and Technology Agency.

REFERENCES 1. K.Uemura and I.A.Nishida, Thermoelectric Semiconductor and its appliciations, Nikkan Kogyo Shinbun, LTD., 1983 2. Teledyne Energy System, Gas Fueled Thermoelectric Generators -Engineering and application Manual, TELAN/DECAP, August, 1983 3. S.W.Petric, Opt. Eng. 26, 965(1987) 4. H.Kaibe, "The Studies on the Thermoelectric Properties for Semiconducting Bi2Te3 Compounds", Thesis, 1989 5. Wessenstein. J., Horak. J., Tiehy. L., Vasko. A., :Cryst. Lattice Defects 8 (1980) 223 6. J.Sugihara,Electric Structure and Thermoelectric Properties of Bi2Te3, the Collected Papers of TEC'96,1996

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

557

Thermoelectric Properties of Arc-melted Silicon Borides Lidong Chen, Takashi Goto, Toshio Hirai Institute for Materials Research, Tohoku University, Katahira 2-1-1, Aoba-ku, Sendai 980-77, Japan

ABSTRACT Silicon borides were prepared by arc melting in a boron content range from 80 to 94mol%. As-melted specimens consisted of SiBn and free silicon. After heat treatment at 1500-1673K, SiB4 formed near the SiBn-Si boundary due to the solid reaction between free silicon and SiBn, and as the result SiBn-SiB4 composites were obtained. The SiBn-SiB4 composites showed larger electrical conductivity and smaller thermal conductivity than the as-melted silicon borides, which leads to an improvement of thermoelectric figure of merit.

1. INTRODUCTION The efficiency of thermoelectric power generation is in proportion to hot junction temperature and the temperature drop between hot- and cold-junctions [1]. Recently, in order to achieve a high conversion efficiency, a new concept called functionally graded thermoelectric material was proposed, in which several types of materials will be joined by graded structure so a broader range of temperature is expected to be covered practically than that is when a homogeneous material is used [2-3]. At present, Bi2Te3, PbTe and SiGe are generally considered practical materials for thermoelctric power generation in the low (about 300 to 500K), medium (about 500 to 800K) and high (about 800 to HOOK) temperature ranges, respectively. Unfortunately, no applicable material has been discovered for the ultrahigh temperature range over HOOK, though some candidate materials have been studied for the past three decades [4-9]. Boron-rich silicon boride is one of the candidate materials for ultra-high temperature thermoelectric conversion because of its moderate Seebeck coefficient (a) and small thermal conductivity (K) at high temperatures over lOOOK [1, 4-6]. In the Si-B binary system, there are many types of compounds such as SiB4 (rhombohedral), SiB6 (orthorhombic) and SiBn (hexagonal, n=15-49) [10-12]. Among them, SiB4 has a low thermal conductivity and high electrical conductivity (a) but a low Seebeck coefficient [13]. On the other hand, SiB6 and SiBn have a large Seebeck coefficient and a low thermal conductivity but a moderately low electrical conductivity. The authors have reported the synthesis and thermoelectric properties of arc-melted silicon borides [14]. As-melted sihcon borides in the boron content range of 80 to 94mol% consisted

558 of SiBn and free silicon, in which the free silicon dispersed in a network structure when boron content is below 90mol%. The existence of free silicon made both the electrical conductivity and thermal conductivity increase, so it did not improve the thermoelectric figure of merits {Z=a^c/K). If the composite of SiBn-SiB4 with SiB4 dispersed in network can be obtained, the decrease of the thermal conductivity accompanied with increase of the electrical conductivity will be expected, and so the thermoelctric figure of merit would be increased. Silicon borides such as SiB6 and SiBn can be synthesized by hot pressing [4], plasma activated sintering [14], arc melting [14] and chemical vapour deposition (CVD) [15]. In the cases of sintering including hot pressing and plasma activated sintering, a high sintering temperature over 1800K is needed. Unfortunately, SiB4 decomposes in 1500-1673K due to its metastability [10, 12, 16-17]. Therefore, SiBn-SiB4 composite is difficult to be synthesized by a sintering method. In the present research, we tried to prepare SiBn-SiB4 composites using arc melting combined with sequent heat-treatment process. In this paper, we report the microstructure changes after heat-treatment and their effects on the thermal conductivity and electrical conductivity.

2. EXPERIMENTAL The mixtures of silicon and boron powder in a boron content range of 80 to 94mol% were pressed into disk-shaped pellets (10 mm thickness and 20 mm diameter) and then arc-melted in an argon atmosphere. The arc-melted samples were then heat-treated in argon atmosphere at temperatures of 1400 to 1700K. The phase composition of the resulted specimens was identified by X-ray diffraction (XRD). Rod-like pieces (3x3x15mm) and disk-shaped pieces (2mm thickness and 10mm diameter) were cut out for the electrical conductivity measurement and the thermal conductivity measurement, respectively. Microstructure and phase distribution were observed by a scanning electron microscopy equipped with EPMA (JEOL: JXA-8621MX). Electrical conductivity was measured using a D.C. four-probe method. Thermal conductivity was measured using a laser-flash technique. All the measurements were performed in the temperature range of 300 to 1200 K.

3. RESULTS AND DISCUSSION Figure 1 shows a typical X-ray diffraction pattern of as-melted silicon boride containing 90mol% boron. In the present boron content range (80 to 94mol%), all the arc-melted specimens consisted of SiBn and free silicon. Figure 2 shows the changes of free silicon content with the boron content in the raw material. The content of free silicon decreased from about 30 to 3vol% as the boron content in raw material increased from 80 to 94mol%. The lattice parameter (a) of

20

25

30

35

26, CuKa / degree Figure 1. X-ray diffraction pattern of an arc-melted silicon boride (B=90mol%).

559 the free silicon was 0.5415nm, which was smaller than the JCPD value (0.543 Inm) [18]. It was reported that the lattice constant of silicon changes from 0.5431 to 5412nm as boron dissolves in silicon up to 3mol%, and the solubility limit of boron in silicon is generally considered about 3mol% [11, 19-20]. Therefore, the free silicon in the melted silicon boride could be almost saturated with boron. We reported that the free silicon phase is inter-connected to form a network structure when B=80-90mol%, while it dispersed isolatedly when B>90mol% [14]. And the Seebeck coefficients is almost independent of phase composition and changes from 100 to 300 jlVK' as the measuring temperature up to HOOK [14]. Figure 3 shows the changes of X-ray diffraction pattern of the melted sample containing 90mol% boron after heat-treatment. Only trace amount of SiB4 formed after annealing at 1538K for 5hr (Fig.3(a)). When the annealing time was 40hr at the same temperature, SiB6 phase appeared but without much change in SiB4 amount (Fig.3(b)). After annealing at 1673K for 0.5hr, moderate amount of SiB4 formed without formation of SiB6 phase (Fig.3(c)). However, after annealing at the same temperature for 2hr, SiB6 appeared and the amount of SiB4 decreased (Fig.3(d)). In the annealing temperatures between 1500 to :3 1673K, SiB4 formed at first, and then SiB6 appeared while the SiB4 amount got decreasing. When the annealing temperature was below 1473K, no new phase appeared even if annealing time is 40hr. And when (c) annealing temperature was over 1680K,

'o

(d) r

I

15

0 75

80

85

Boron content

90

•

14-

20

•

•

.

.

I

25

.

.

•

•

I

I

30

I

I

I

35

29, CuKa / degree 95

/ mol%

Figure 2. Relationship between free silicon content in arc-melted silicon borides and boron content in raw materials.

Figure 3. X-ray diffraction patterns of arcmelted silicon boride (B=90mol%) after heat-treatment. The heat-treatment conditions are: (a) 1538K, 5hr; (b) 1538K, 40hr; (c) 1673K, 0.5hr; (d) 1673K, 2hr.

560

Figure 4. Microstructures of arc-melted silicon boride (B=90mol%). (a) as-melted; (b) after heat-treatment at 1673K for 0.5hr. The A, B and C parts are SiBn, free-Si and SiB4, respectively. only SiB6 formed as the new phase when the annealing time is even lOmin. In general, SiB4 can be prepared in the form of mixtures with SiB6 and/or silicon by melting silicon and boron powder mixtures or heating the mixture powder at temperatures between 1500 and 1673K [12, 16-17, 21], though there is one report that the pure SiB4 was synthesized by heating the stoichiometric mixture at 1473K for about three weeks [16]. It is generally known that, when the mixture powder of silicon and boron are heated at 1500 to 1673K, SiB4 forms at first, and then it slowly decomposes into SiB6 and silicon [10, 12, 16-17]. SiB4 is stable below 1473K. The present experimental results agreed with those reports. The free silicon in the arc-melted silicon borides can be partially changed into SiB4 due to the solid reaction between silicon and SiBn through heat-treatment process by a proper annealing temperature and annealing time. Figure 4 shows the microstructures of the specimen containing 90mol% boron before and after annealing at 1673K for 0.5hr. In Fig.4(a) of the as-melted sample, the EPMA

104

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S 103 0.5

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as-melted heat-treated as-melted heat-treated

1.5

X-1

/

2.0

2.5

10-3 K-1

Figure 5. Temperature dependence of electrical conductivity of arc-melted silicon borides before and after heattreatment at 1673K for 0.5hr.

561 study revealed that, the black area (A) is SiBn and bright area (B) is free silicon. In Fig.4(b) of the annealed sample, the black area (A), bright area (B) and the gray area (C) are SiBn, free Si and SiB4, respectively. The free silicon phase inter-connected to form a network structure in the as-melted specimen. After annealing, the free silicon and SiBn reacted near the boundary to form SiB4, and as the result, the silicon network changed into SiB4 network though some silicon still remained within SiB4 mainly at the cross points of the network. Figure 5 shows the temperature dependence of electrical conductivity (a) of arc-melted silicon borides before and after annealing. The larger the free silicon amount, the greater the a values. After annealing at 1673K for 0.5hr, the a values increased about 30-60% for all the specimens. Particularly, the a values of the specimen containing 88mol% boron almost agreed with those of CVD-SiB4 [13]. Figure 6 shows the temperature dependence of thermal conductivity (K) of arc-melted silicon borides before and after annealing. The larger the free silicon amount, the greater the K values. After annealing at 1673K for 0.5hr, K values decreased about 20-30% for all the specimens. They changed from 8 to 6WK-im-^ when measuring temperature increased from room temperature to HOOK, which almost agreed with those of the SiB4 [13, 17] and SiBn [6, 14]. The composite of SiBn-SiB4 showed the greater electrical conductivity and smaller thermal conductivity than the as-melted silicon borides. These changes contribute to the improvement of the thermoelectric property.

4. CONCLUSION 12.0

Silicon borides were prepared by arc melting in argon atmosphere using silicon and boron powders in a boron content range from 80 to 94mol%. The as-melted specimens consisted of SiBn and free silicon. The contents of free silicon decreased from 30 to 3vol% as the boron content in raw material increased from 80 to 94mol%. The free silicon phase interconnected to form a network structure when B=80-90mol%. The as-melted specimens were heattreated in argon at temperatures of 1400 to 1700K. Below 1473K, no phase changes occurred by annealing for 40hr. In the anneahng temperatures of 1500 to 1673K, SiB4 formed at first due to the solid reaction between free silicon and SiBn near the S i B n - S i boundary, and then SiB6 appeared after a longer annealing time due to the d e c o m p o s i n g reaction of SiB4. Through annealing at 1673K for 0.5hr,

•: ii.oH1 r

^ ^

10.0

•

A :

•••••

•---..

' ^^^x ****• N

7.0

as-melted heat-treated as-melted heat-treated

\

••.. 9.0

B=88mol%, B=88mol%, B=90mol%, B=90mol%,

°: A:

6.0 H 5.0 '. 400

i

1

1

600

1

800

Temperature

1

1

1000

1

1200

/ K

Figure 6. Temperature dependence of thermal conductivity of arc-melted silicon borides before and after heat-treatment at 1673Kfor0.5hr.

562 SiBn-SiB4 composites, in which SiB4 is inter-connected into network, were obtained. The SiBn-SiB4 composites showed larger electrical conductivity and smaller thermal conductivity than the as-melted specimens, which leads to an improvement of thermoelectric figure of merit.

ACKNOWLEDGEMENTS We thank Mr. Y. Murakami of IMR, Tohoku University for helping EPMA analysis. This research was supported in part by the Grant-in-Aid for Scientific Research from the Ministry of Education, Science and Culture, under contact nos. NP0701 and (B)06453081, also supported by the Special Coordination Funds for Promoting Science and Technology from the Science and Technology Agency of Japan.

REFERENCES 1. C. Wood, Materials Research Society Symposia Proceedings, Vol 97, Ed. by D. Emin, T. L. Aselage and C. Wood (Materials Research Society, Pittsburgh, 1987), p.335. 2. T. Hirano, L. W. Whitlow, M. Miyajima, Ceramic Transactions, Vol.34, Ed. by J. B. Holt, M. Koizumi, T. Hirai and Z. A. Munir (Amer. Ceram. Soc, Westerville, Ohio, 1993), p.23. 3.1. A. Nishida, Material Technology, 14 (1996) 9. 4. C. Wood, D. Emin, R. S. Frigelson and I. D. R. Mackinnon, Materials Research Society Symposia Proceedings, Vol 97, Ed. by D. Emin, T. L. Aselage and C. Wood (Materials Research Society, Pittsburgh, 1987), p.33. 5. B. Armas and C. Combescure, J. Less-Conmion Met., 47 (1976) 135. 6. J. M. Darolles, T. Lepetre and J. M. Dusseau, Phys. Stat. Sol. (a), 58 (1980) K71. 7. O. A. Golikova and I. M. Rudnik, USSR Neorg. Mater., 14(1978)17. 8. T. L. Aselage: Modern Perspectives on Thermoelectrics and Related Materials, Ed. by D. D. Allred, S. B. Vining and G. A. Slack (Materials Research Society, Pittsburgh, 1991), p.l45. 9. T. Goto, E. Ito, M. Mukaida and T. Hirai, J. Jpn Soc. Powder & Powder Metallurgy, 43(1996)311. 10. R. W. Olesinski and G. J. Abbaschian, Bull. Alloy Phase Diagrams, 5 (1984) 478. 11. B. Armas, G. Male and D. Salanoubat, J. Less-Common Met., 82 (1981) 245. 12. H. R Rizzo, B. C. Weber and M. A. Schwarz, J. Amer. Ceram. Soc, 43 (1960) 497. 13. M. Mukaida, T. Goto and T. Hirai, Mater. & Manufacturing Processes, 7 (1992) 625. 14. L. Chen, T. Goto, J. Li, M. Niino and T. Hirai, Trans. lEE Jpn, 116-A (1996) 248. 15. M. Mukaida, T. Goto and T. Hirai, J. Mater. Sci. 27 (1992) 255. 16. C. Brosset, B. Magnusson, Nature, 187 (1960) 54. 17. R. S. Feigelson and W. D. Kingery, Ceram. Bull., 42 (1963) 688. 18. JCPD card. File No. 27-1402 (1977). 19. G. V. Samsonov and V M. Sleptsov, Russ. J. Inorg. Chem., 8 (1963) 1047. 20. J. Hesse, Z. Metallkd., 59 (1968) 499. 21. H. Moissan and A. Stock, Compt. Rend., 131 (1900) 139.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

563

Graded thermoelectric materials by plasma spray forming Jiirgen Schilz"*, Eckhard Miiller", Wolfgang A. Kaysser'', Gregor Langer^, Erich Lugscheider^, Giinter Schiller^, Rudolf Henne^ "German Aerospace Research Establishment (DLR), Institute of Materials Research, D-51140 Koln, Germany ^Aachen, University of Technology, Materials Science Institute Jiilicher Str. 342, D-52056 Aachen, Germany ^German Aerospace Research Establishment (DLR), Inst, for Technical Thermodynamics, D-70503 Stuttgart, Germany Plasma spraying is a consolidation process for powders with the additional capability of a composition control of the spray formed structures. The paper reports on the first steps to adapt this method to the production of functionally graded thermoelectric materials with a locally maximized figure of merit. Iron disilicide (FeSi2) was used to test the performance of the technique on thermoelectric material. It was found that plasma spray forming is applicable to produce dense materials with thermoelectric properties comparable to hot pressed ones. Problems were however found with the thermal stability of the microstructure. Final goal is the employment of plasma spraying to form compositionally graded materials of the Mg2(Si,Ge,Sn) system. Here we report on the preparation and thermoelectric transport properties of Si-rich quasibinary Mg2(Si,Ge) and Mg2(Si,Sn) mixed crystals by mechanical alloying. 1. I N T R O D U C T I O N Thermoelectric (TE) generators have shown their reliability in space applications and there are already a number of efforts to bring TE generation of electricity into terrestrial commercial use. Requirement, however, are lower investment costs for the modules and simultaneously higher conversion efficiencies. In terms of costs it is indispensible to develop manufacturing processes for TE devices, which not only consolidate the powders in automized procedures, but also have the potential of lateral shape forming and composition control of the material. This would not only give the possibility to form complete elements or modules in a single manufacturing step, but also to prepare functionally adapted TE devices which show a higher figure of merit [1]. Here, we introduce plasma spraying as a forming process with the required capabilities. As first T E material to be bulk solidified by plasma spraying, we have chosen iron disilicide (FeSi2) [2]. Due to its well known properties, the processing of FeSi2 allows a valuation of

564

,

•

Al-doped PS FeSij

Co-doped PS FeSi,

i 2

I 1 ^ 20

40

60

Tempertime at 800°C in h

Figure 1. Scanning electron micrograph of SPS formed Al-doped FeSi2 perpendicular to the growth direction. 1: FeSi2 + 0.7at% Al, 2: FeSi2 + 1.4at% Al, 3,4: FeSia + 1.8at% Al, 5: FeSi, 6: Oxide.

Figure 2. Room temperature thermal conductivity values of SPS formed FeSi2 as a function of heat treatment time at 800° C.

the novel method's applicability. The results of the experiments are presented in the next section. Additionally, the first outcomes of the production of graded electrical contact junctions are reported. Final goal, however, is the preparation of compositionally graded materials made of Mg2(Si,Ge,Sn) crystals. The preparation of the powders by mechanical alloying and their properties after hot pressing is topic of the third section. The high reactivity of the Mg is the main difficulty which has to be overcome. For the preparational route the problems have been solved. The next step will be the transfer to the plasma spray process. 2. P L A S M A S P R A Y F O R M I N G Plasma spraying (PS) is a technology that has been widely used for coating both metals and ceramics [3]. In the present case, however, the spraying process is applied with the aim to form bulk material. Therefore we name the method plasma spray forming. 2.1. P l a s m a spray experiments on iron disilicide As first material to test PS forming of thermoelectric materials, we used gas atomized, Co- and Al-doped, FeSi2 powders [4]. The powder batch was sieved into two fractions. The powders with the larger mean particle diameter of 70 /im were successfully deposited by plasma spraying in ambient air (air plasma spraying, APS) or with an argon shroud which aggravates oxidation (shrouded plasma spraying, SPS). The targets were made of unalloyed conventional steel or of Ti-6A1-4V. PS powers between 20 and 30 kW with pure argon as plasma gas, or with hydrogen or helium additions were employed. Spraying distances were chosen from 70 down to 30 mm. The smaller distance made finer lateral structures of about 10 mm possible, but also led to a lower deposition rate. The scan velocity was varied by one order of magnitude between 16 and 160 m m per minute. With

565 a powder feeding rate of about 30 g per minute the typical layer thickness per scan reached 0.1 mm. When employing so-called vacuum plasma spraying (VPS), i.e. spraying in a closed chamber operated below atmospheric pressure filled with 200 mbar inert argon gas, only the use of the finer powder fraction with mean particle diameters of 20 //m resulted in a suffient melting and deposition rate in the high velocity plasma jet. However, though the spraying distance was larger in the present case, the layer thickness per scan did not exceed 0.02 mm. 2.2. Metallographic characterization For APS and SPS, the chosen spray parameters were appropriate to produce FeSi2 bulk material up to a thickness of 10 mm with a relative density of at least 92%. The mechanical strength was almost identical to that of hot pressed material and thus there were no problems in cutting and machining samples to be tested. In the case of VPS, however, the maximum sample thickness achieved did not reach much above 1 mm, but its relative density approached the high value of 98%, which is comparable to that of hot pressed FeSi2. A scanning electron microscopy picture of SPS FeSi2 (cf. Fig. 1), which is taken perpendicular to the growth direction, shows the main features of an as sprayed structure. Striking are the microcacks which develop in the splat bond regions. A similar microstructure has been observed in PS mullite [6]. These regions also contain oxides, which were found in all types of PS FeSi2. Up to now not explainable is the result that a chemical analysis of oxides in VPS FeSi2 gives much higher values than those found in SPS and hot pressed material (cf. Table 1). The PS FeSi2 in its as-sprayed state was mostly in its metallic a-phase, and had therefore to undergo a temper process of about 1 h between 700 and SOO^'C to obtain the desired semiconducting /?-phase [5]. 2.3. T h e r m o e l e c t r i c transport properties The thermoelectric properties of the plasma spray formed iron disilicide might be influenced by a possible change in material composition due to the high temperatures, and the highly disordered structure. From the three thermoelectric parameters, Seebeck coefficient S, electrical conductivity a, and thermal conductivity /c, the thermal conductivity is most sensitive to the microstructure. From Fig. 2, which shows a plot of room temperature thermal conductivity of SPS FeSi2 vs. heat treatment time at lOOO^C, it can be seen that there is obviously a change in microstructure due to the thermal treatment. At the beginning, the thermal conductivity is extremely low which is probably due to the large number of microcracks in the bond regions of the flats. Tempering induces an annealing process which lets K rise eventually reaching the value of hot pressed FeSi2. VPS FeSi2 does not exhibit such a large change in thermal conductivity. From the beginning, the material is more dense and /c quite large. Table 1 lists AC-values of SPS and VPS FeSi2 which were measured after a 2 h heat treatment at 800°C (i.e. the time required to induce the a to ^ transition) and after a certain temper time at 800°C. It can be seen that in both cases the thermal conductivity of hot pressed FeSi2 is approached. APS FeSi2 shows an extremely low electrical conductivity of only 15 (Ocm)"^ (Co-

566

30

^x . 0 *,

0.5 h 1.5 h 2.5 h 72 h

E

^ to

20

300

400

500

600

TIKI

Figure 3. Electrical conductivity of SPS formed Co-doped n-type FeSi2 after different heat treatment durations at 800° C.

Figure 4. Graded junction between FeSi2 (bottom) and an iron based alloy (top). The height of the picture represents 2 mm.

doped). The a-values of SPS material are somewhat higher (cf. Fig. 3), but still a factor 5 lower compared to hot pressed material. Only VPS FeSi2 exhibits acceptable values around 65 (Ocm)"^. Interesting is the fact, that the electrical conductivity does not show any significant depence on the aging process by which the thermal conductivity is influenced so dramatically. Fig. 3 shows, that even after 72 h of heat treatment at 800°C on an Co-doped SPS sample, the cr-values are almost constant. The same behaviour is valid for the thermopower S (cf. Table 1). 5-values for PS FeSi2 are higher than those of hot pressed material - especially for the VPS material - which is obviously due to a dopant loss during the spraying process.

Table 1 Comparison of hot pressed and plasma spray formed Co-doped FeSi2. Hot pressed SPS VPS "as sprayed" 50h at 800°C "as sprayed" 50h 22 a (1/ncm) Too 22 65 180 S (/.V/K) 165 190 180 1 • 10^1 n (cm~^) 2 • 10^1 ; • lO^*" 4.9 1.5 K (W/Km) 3.4 4.9 5.5 4.7 6.8 Z (10-VK) 1.6 0.17 0.7 0.15 O2 (wt%) Room temperature values.

at 800°C 190 4.5 5.2

2.4. G r a d e d j u n c t i o n Fig 4 shows a graded junction between the semiconducting FeSia and an iron-based alloy. Purpose is the formation of a metaUic surface on which electrical contacts can

567

100-

Mg2Sio,,Ge„3

80v>

J

Hall-mobility — o — Preparation as usual — • — Oxide-free preparation

> rTeoE o £40-

MgjSiogSrioi

20-

MgzSioeSnoz Mg;SySno3

/ . ^

J J _^ooo

/

"

•

()

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o ^

\

050

100

Mg2Sio6Sno4

Figure 5. Layered Mg2(Si,Ge,Sn) ingot prepared by hot pressing.

/

/ // /

=L

]

/ / /

150

200

250

300

TinK

Figure 6. Hall mobility of two mechanically alloyed, hot pressed Mg2Si samples.

be soldered with a low melting point, ductile solder. Such a contact can release stresses. Gradient formation was achieved by simultaneously feeding FeSi2 and Fe into the plasma gas jet, and stepwise varying their relative fraction. From the however, it can be seen that the mixing of the powders was not complete which in a quasi-layer structure. In spite of this, the contact is mechanically stable.

thermal powders picture, resulted

3. S Y N T H E S I S A N D P R O P E R T I E S OF Mg2(Si,Ge,Sn) The Mg2(Si,Ge,Sn) mixed crystal system shows wide ranges of solid solubility and the compounds have been proposed to be good candidates for high-Z TE-materials [8,9]. Additionally, the material system offers the possibility to form compositionally graded elements. In the present project, in order to obtain the materials in the required powder form, the compound is prepared by a mechanical alloying process in a planetary ball mill Retsch PM4000DLR [10,11]. Up to now the method succeeded in synthesizing Si-rich quasibinary Mg2(Si,Ge) and Mg2(Si,Sn). Consolidation occured by hot uniaxial pressing. Layered structures were prepared by pressing stacks of powders having different compositions (cf. Fig. 5 for example). During the preparation process it became evident that a paramount parameter governing the thermoelectric properties is the oxide contents in the material. Figs. 6 shows mobility curves of two mechanically alloyed Mg2Si. The powders were prepared in the same way with n-hexane as oxygen free milhng fluid, but in the "oxide-free" preparation route, the powder was not dried in an argon atmosphere, but, still wet, filled into the pressing die. The room temperature mobiUty of such prepared material appeared to be a factor of three higher. Thus it is important to avoid any oxidation during the whole milling and consolidation procedure — a fact which has to be considered when transferring the consolidation into the plasma spray technology.

568 4.

CONCLUSIONS

Our experiments found that the employed Co- and Al-doped gas atomized iron disilicide powders exhibit a good spray ability in case of APS and SPS deposition. For VPS, the deposition rate has to be improved. Plasma spray formed layers have to undergo a postdeposition annealing process in order to obtain the semiconducting ^-phase. Regarding thermoelectric properties, the Seebeck-coefficient is comparable to hot pressed FeSi2 and the thermal conductivity is found to be lowered. However, the electrical conductivity was lowered as well, which in the case of APS and SPS material led to a lower Z-value. On the other hand, VPS FeSi2 showed an electrical conductivity high enough to compete with conventionally hot pressed material. The PS FeSi2 is not thermally stable. An aging process occurs, which may be due to recrystalHsation and/or microcrack annealing. Interesting is the fact, that the aging process seems to affect only the thermal diffusivity (conductivity), whereas the electrical conductivity and the thermopower remain unchanged. Quasibinary solid solutions of Mg2(Si,Ge) and Mg2(Si,Sn) could be prepared by mechanical alloying. By consequently avoiding the formation of oxides, the room temperature electron mobility values could be raised to 100 cm^/Vs. The obtained values indicate that the material is indeed a good candidate for future high efficient TE-material. However, due to its high reactivity, the maximum working temperature may be limited to 500° C. This work was supported by the Deutsche Forschungsgemeinschaft (DFG), Schwerpunktprogramm Gradientenwerkstoffe. REFERENCES 1. T. Hirano, L.W. Whitlow, M. Miyajima, in: Ceramic Transactions, Functionally Gradient Materials, ed. by J.B. Holt, M. Koizumi, T. Hirai, Z.A. Munir, The Amsterdam Ceramic Soc. 34 (1993) 23. 2. R.M. Ware, D.J. McNeill, Proc. lEE 111 (1964) 178. 3. R.W. Smith, R. Novak, pmi 23 (1991). 4. U. Stohrer, U. Taibon, E. Gross, U. Birkholz, Proc. IX Int. Conf. on Thermoelectrics, Pasadena CA, (1990) 242. 5. J. Schilz, M. Riffel, R. Mathesius, G. Schiller, R. Henne, R.W. Smith, Proc. of the XV Int. Conf. on Thermoelectrics (ICT'96), Pasadena, USA (1996). 6. W. Braue, G. Paul, R. Pleger, H. Schneider, J. Decker, J. Eur. Ceram. Soc. 16 (1995) 85. 7. H.E. Eaton, J.R. Linsey, R.B. Dinwiddle, In: Thermal conductivity 22, Ed. T.W. Wong, Technomic (1994) p. 289. 8. R.J. LaBotz, D.R. Mason, D.F. O'Kane, J. Electrochem. Soc. 110 (1963) 127. 9. R.J. LaBotz, D.R. Mason, D.F. O'Kane, Proc. of the XII Int. Conf. on Thermoelectrics (ICT'93), Yokohama, Japan (1993). 10. M. Riffel, J. Schilz, Proc. of the XV Int. Conf. on Thermoelectrics (ICT'96), Pasadena, USA (1996). 11. J. Schilz, K. Pixius, W. Amend, M. Plate, H.-J. Meyer, Proc. of the XIII Int. Conf. on Thermoelectrics (ICT'94), Kansas City MO, USA, AIP Conf. Proc. 316 (1995) 71.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

569

Preparation of PbTe-FGM by Joining Melt-grown Materials M. Orihashi ^, Y. Noda ^, L. -D. Chen ^, Y. -S. Kang ^, A. Moro ^ and T. Hirai ^ ^ Institute for Materials Research, Tohoku University, Katahira, 2-1-1, Aoba-ku, Sendai, Miyagi 980-77, Japan " Faculty of Engineering, Tohoku University, Aramaki Aoba, Aoba-ku, Sendai, Miyagi 980-77, Japan ^ National Aerospace Laboratory, Kakuda Research Center, Koganezawa 1, Kimigaya Kakuda, Miyagi 981-15, Japan The 2-stage carrier concentration FGM of PbTe were prepared by plasma activated sintering (PAS) using the discs cut from melt-grown PbTe ingots. The component materials of PbTe ingots were prepared by the Bridgman method after the direct melting of the constituent elements (Pb and Te) and 2000 or 4000 molppm Pbl2 as a n-type dopant. The PAS conditions were as follows; in vacuum, pressure of 30 MPa, temperature at 1050-1100 K, and current of 600 A and voltage of 25 V. The powder of either of the component materials was adopted as joining reagent of the discs. The thermoelectric characterization of the FGM was made in the temperature range from 300 to 700 K. The electrical conductivity of the FGM at 300 K was almost intermediate between those of the components, while the thermoelectric power corresponded to that of the component with high carrier concentration at high temperature side. The thermoelectric power for the FGM was almost intermediate between those for the components. At T>650 K, the electrical figure of merit for the FGM was larger than that for the components. 1. INTRODUCTION PbTe is among the best materials used in construction of thermoelectric generators working at intermediate temperature region (450-840 K). Since the maximum figure of merit (Z) shifts in wide temperature range depending upon carrier concentration, the carrier concentration FGM of PbTe is expected to attain high efficiency of thermoelectric energy conversion.

570 Plasma-activated sintering (abbreviated as PAS) has been developed in the preparation of a variety of functionally graded materials (FGM)[1], where the powder particle surface are activated by plasma leading to an accelerated reaction or sintering. The PAS has lately been applied to the thermoelectric materials such as SiGe and PbTe[2,3]. In the present study, we prepared the stepwise carrier concentration FGM of n-type PbTe by using PAS. 2. EXPERIMENTAL 2.1 Preparation of PbTe-FGM The PbTe ingots as the source material were prepared by the Bridgman method after the direct melting of constituent elements of Pb and Te (nominal purity of 99.9999%). The weighed amount of the elements in the stoichiometric composition (Pb/Te=l) was vacuum sealed in a quartz ampule (10 mm ID and 120 mm length) with 2000 or 4000 molppm Pbl2 as the source of n-type dopant of iodine. The growth condition was as follows; the maximum heating temperature of 1223 K, the temperature gradient of about 1200 K/m at melting point, and growth rate of 4 mm/hr. The PAS was performed in a carbon die with 11 mm ID. The n-type 2-stage FGM was prepared by direct PAS of the two discs (^2 mm thickness X 10 mm diameter) cut from the 2000 and 4000 molppm Pbl2 doped ingots, where the powder from the 4000 molppm Pbl2 doped ingots was placed between the discs as the binder. The PAS condition was; in a vacuum of about 10 Pa under the condition of load of 30 MPa, and pulse current of 600 A at 25 V with 6 Hz for 90 s, followed by the heating at 810 K for 540 s. 2.2 Measurement of thermoelectric properties The thermoelectric properties were measured at 300 K for the FGM and its component layers separated from the FGM. The electrical conductivity (cx) and Hall coefficient (JR^) were measured by the 6-probe method for the FGM and by the van der Pauw configuration for the components cut from the FGM using Pt-wire electrodes. The carrier concentration (n) and Hall mobility ( A H ) ^^^^ calculated using the equation n=lleRY{ (e: electric charge) and /^jj=/?H cr, respectively. The thermoelectric power ( a ) at 300 K was estimated from the linear relationship between thermoelectromotive force (EMF) and temperature difference within 5 K. The thermoelectric properties of the FGM and its components were measured in Ar atmosphere in the temperature region from 300 to 700 K. The sample size was ^^3 X 3 X 8 mm^. On the thermoelectric power measurement, the components of high and low carrier

571 concentration were arranged to high and low temperatures, respectively, and thermoelectromotive force was measured at the temperature difference within 20 K between the hot and cold ends. The temperature was monitored by using Pt-13%Rh thermocouples and the additional Pt electrodes were adopted for the EMF measurement. For the conductivity measurement, the Pt wires of the thermocouples were served as the current lead and the additional electrodes as the potential lead. 3. RESULTS AND DISCUSSION Table 1 lists the thermoelectric properties of two stage n-type FGM and its components at 300 K. The data for the components (1) and (2) of the FGM corresponded to those for the melt-grown PbTe single crystals doped with 2000 and 4000 molppm Pbl2 respectively[4]. This indicates that PAS can produce a high quality material. The a value for the FGM was almost intermediate between those for the components and never became lower than those of the components[5], while thermoelectrical power, carrier concentration and Hall mobility seem to represent those for either component. Fig. 1 shows the potential profile near the joint boundary in the FGM. The abrupt potential change was found within a width of about 0.5 mm, which is also reported for the SiGeFGM[2]. Since the potential gradient corresponds to an increase of resistivity at the joint interface, the interface resistivity must be minimized by optimizing the sintering in the further studies in order to attain high efficiency of energy conversion. Fig. 2 shows the temperature dependence of CT for the FGM and the components. The (7 value of all the samples monotonously decreased with an increase of temperature, indicating that the samples are the typical degenerated semiconductors. The <J value for the FGM was almost intermediate between those for the components and never become lower than those of the components in spite of the existing interface resistivity. At Jb>500 K, the O value for FGM was larger than that for the components. Fig. 3 shows the temperature dependence of a for the FGM and the components. The a Table 1 Thermoelectric properties of the two stage n-type FGM of PbTe at 300K layer n-FGM(l+2) Component (1) (2)

Pbl2 (molppm) 2000 4000

a

n

(Q-V^)

(mV^s'l)

4.1X1024*

1.0X10^

1.5X10-1*

(V-K-1) -1.0X10-4

4.2X10^4 4.9X10^5

7.6 XIO^ 1.9X10^

1.1X10-1 2.4X10-2

-2.4X10-4 -1.1X10-4

Data obtained by 6-probe Hall measurement.

572

T

^S ^^ a

0.08 0.06 boOOmolppm Pblj

|

1

T"•~T

^.lO'

'>

"^^'^^Qj..

Z

0.04

'^^^^^^fej^

s -8 |10^ r

0 0.02

:

0 1 1 • A 2 1 ^ FGM|

\

4000molppm Pblj

1

^A'^'^^AAA^VtoiA^^S^ [Hjgiijr'|fc^

^**i»M^^H|

"""^^

S r

at300K I=20mV

^

1 2 3 4 Position ,x /mm Fig.l Plots of electrical potential versus position near at the boundary joined between 2000 and 4000 molppm Pbl2doped PbTe in the two-stage FGM

u

103

- 1

1

1

1

1

200

300 400 500 600700600 Temperature, T/K Fig.2 Temperature dependence of electrical conductivity for 2-stage FGM of PbTe shown with those for component (1) and (2) with different electron concentration listed in Table 1 ^

10-^

1—I—]—1—r

r\w 2FGM ^ 10-2 b

^°

I 10-^1

I

300

400

500

600

700

Temperature, T/K Fig.3 Temperature dependence of thermoelectric power for 2-stage FGM of PbTe shown with those for component (1) and (2) with different electron concentration listed in Table 1

10-^

300

400 500 600 700 Temperature, T/K Fig.4 Temperature dependence of electrical figure of merit a ^ a for 2-stage FGM of PbTe with those for component (1) and (2) with different electron concentration listed in Table 1

value for the component (1) linearly increased up to near 600K and then took a maximum of 3.2 X lO""^ V-K"^ while the maximum for (2) was 2.1 X 10""^ V-K"^ at about 630 K. The curve for the FGM was almost intermediate between those for the components. Fig. 4 shows the temperature dependence of the electric figure of merit (or power factor, O) for the FGM and the components. With an increase of temperature, the a^ a value for the component (1) monotonously decreased with an increase of temperature, while that of (2) increased. In the measured temperature range, the a^ cr value for the component (2) was

573 lower-lying than that for (1). At J>650 K, the a^ a value for the FGM was found to exceed that for the components. However, the increase of efficiency for the FGM in the present study did not enhance the power factor in low temperature region. The large difference in carrier concentration between components (1) and (2) may degrade the thermoelectric properties of the FGM under the temperature difference within 20 K. Therefore, it is concluded that the difference in carrier concentration must be kept small between the adjacent component materials in a stepwise FGM in order to attain high efficiency, which leads to a FGM with well-designed continuous carrier concentration profile. 4. Acknowledgment The authors wish to thank Mr. K.Horasawa (Tohoku University) for manufacturing the glass equipments. REFERENCES 1. T. Hirai: Materials Science and Technology, A Comprehensive Treatment, ed. R. W. Cahn, P. Hassen and E. J. Kramer, Chapter 20 Functional Gradient Materials, VCH Verlag. (Weinheim), (1996), 293 2. K. Takahashi, T. Masuda, T. Mochimaru and T. Noguchi: Froc, FGM symp. (FGM95), (1995), 123. 3. M. Miyajima, K. Fujii and T. Hirano: Froc XII Int. Conf. Thermoelectrics (X II-ICT), Yokohama, (1993), 272. 4. Y. Noda, M. Orihashi and I. A. Nishida: Trans. lEE of Japan, 116-A (1996), 242. 5. Y. Noda, M. Orihashi H. T. Kaibe, Y. Imai, I. Shiota and I. A. Nishida: Froc. Int. Conf. Thermoelectrics (ICT96), Pasadena, (1996), in print. 6. Yu. I. Ravich, B. A. Efimova and I. A. Smimov: Semiconducting Lead Chalcogenides, Prenum Press, New York, (1970).

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I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

575

Improvement and thermal stability of thermoelectric properties for n-type segmented PbTe S. Yoneda^, H.T.Kaibe^, T. Okumura^, Y. Shinohaxa^, Y. Imai^, LA. Nishida^, T. Mochimaru^ K. Takahashi^ T. Noguchi^ and I. Shiota^ ^Department of Electronics and Information Engineering, Tokyo Metropolitan Universisty, 1-1, Minami-Ohsawa, Hachioji-shi, Tokyo, 192-03, Japan ^National Research Institute for Metals, 1-2-1, Sengen, Tsukuba-shi, Ibaraki, 305, Japan ^Vacuum Metallurgical Co., Ltd., 516 Yokota, Yamatake-cho, Yamatake-gun, Chiba, 289-12, Japan ^Kogakuin University, 2665-1 Nakano-cho, Hachioji-shi, Tokyo, 192, Japan The resistivities p of the n-type melt-grown and plasma activated sintered PbTe were measured as a function of temperature during heating and cooling cycles to progressively higher temperatures until 903 K in Ar atmosphere. The hysterisis of the temperature dependence of p appeared for both specimens after heated at 703 K, which indicates that the change of thermoelectric properties occurs during the operation as a thermoelectric generator. This phenomena was more remarkable for the sintered specimen than that for the melt-grown one. However, X-ray analyses did not detected the evidence such as the apperance of the secondary phase or oxidation. The each specimen prepared by meltgrown and plasma activated sintering was also annealed in an evacuated quartz tube at 793 K for 8.64x10^ s (24 h). p and the carrier concentrations for both specimens were almost unchanged after annealing. Then, it is pointed out that a PbTe thermoelectric generator is required to be used in a closed evacuated container in order to prevent the change of the thermoelectric properties during operation. 1. INTRODUCTION The n-type PbTe as a thermoelctric material whose carrier concentration n is controlled to have around 3.0 x 10^^ m~^ is commonly used in the temperature range of a hot side electrodes between 600 and 950 K.[l] The thermoelectric figure of merit Z has a maximum value of 1.4 X 10"^ K~^ at 700 K and the dimensionless thermoelectric figure of merit ZT attains to almost unity. Then, PbTe and its solid solutions are utilized such as a solar thermoelectric generator (STG) and a radioisotope thermoelectric generator (RTG).[1,2] Z of the n-type PbTe with a single and homogeneous carrier concentration has a maximum value Zmax at an optimum temperature T^pt and decreases rapidly with temperature higher than T(ypt. Topt and Zmax are the functions of n. Then, Topt shifts to the higher temperature side and Zmax decreases with increasing in n. When n varies from 5 x 10^^

576 to 7 X 10^^ m"^, Zjnax holds its value between 1.7 x 10"^ and 1.2 x 10"^ K"^ in the wide temperature range between 400 and 800 K. Then, the average figure of merit in an operation temperature range is expected to be improved to 1.5 times larger than that of a conventional homogeneous material with a single carrier concentration. [1] Imai et al. showed the evidence of the improvement in thermoelectric performance by using the hot-pressed n-type PbTe with a 3-stage carrier concentrations. [3] However, in order to form the FGM structure and to attach the metal electrodes such as Fe and other alloys the PbTe materials are needed to be heated up to around 1000 K during the diffusion bonding and hot pressing. [4] Then, it is possible to be caused the change of the thermoelectric properties and it becomes important and serious problems for designing the thermoelements used for a generator and for reliablity of the PbTe thermomodules.[5,6] The purposes of this study are to clarify the mechanism of the change of the thermoelectric properties for n-type melt-grown and sintered PbTe and to establish the methods and techniques for prevention of the change of the thermoelectrc properties. Then, the resistivities were measured during a series of heating and cooling cycles for the n-type melt-grown and sintered PbTe. X-ray analyses were carried out to examine the identification of the mother phase, precipitation of the second phase and oxidation. The annealing effects for n-type sintered and melt grown PbTe evacuated in quartz tubes were also examined. 2. EXPERIMENTAL P R O C E D U R E 2.1. Preparation of n-type melt-grown PbTe Pb and Te with purity of 99.999 % were individually weighed out corresponding to the stoichiometry. They were loaded into an evacuated quartz tube whose size was 50 mm in diameter and 300 mm in length together with 0.5 wt% Pbl2 as n-type dopant. The tube was sealed off under a pressure of I x 10"^ Pa. The melt was stirred sufficiently for 7.2 x 10^ s (2 h) at 1273 K and then it was cooled down along 2.8x10"^ K/s (10 K/h) with holding a temperature gradient of 0.8 K/mm. The obtained n-type melt-grown PbTe ingot was 50 mm in diameter and 30 mm in length. Hereafter, this ingot is named as Ingot 1. Parallelepiped specimens for resistivity p and Hall coefficient Ra were cut out of the top and bottom portions of Ingot 1. The measurements of p and Ru were carried out by d.c. method. Ru was measured with an applied magnetic field of 0.33 T. The p were 2.16x10"^ fi m at the top portion of Ingotl and 2.41x10"^ fi m at bottom one, which indicates that Ingotl was fairly homogeneous in respect of chemical composition and carrier concentration. Another ingot was prepared by the same process with the equal amount of Pb, Te and Pbl2 to those of Ingot 1. It was sealed off again under a pressure of 1x10"^ Pa in a quartz tube consisting of the two cylindrical parts with 16 and 50mm in diameter respectively and with a conical shaped end whose tip angle was 60^ as shown in Fig. 2 of reference 3. The melting and solidification were carried out the same process as the case of Ingot 1. The obtained ingot was 16 mm in diameter and 130 mm in length. The ingot is named as Ingot 2. The conduction type of whole region of Ingot 2 was n-type, while n was decreased along the growth direction from 2.1x10^^ to 2.8x10^^ m~^. X-ray powder diffraction analyses using Cu-Ko; (40kV-50mA) confirmed that both Ingot 1 and 2 consist of a single phase of PbTe.

577 2.2. Preparation of n-type sintered PbTe Each portion of 20, 60 and 96 mm from a growing tip was cut out of Ingot 2. They corresponded to those with p of 2.84x10"^, 5.68x10"^, 9.46x10"^ fim and with n of 2.0x10^^, 1.0x10^^, 0.6x10^^ m~^, respectively They were used as starting materials for the plasma activated sintered PbTe. Each starting material was ground into a fine powder with an average particle size of 42.5/im under 106 /xm. The plasma activated sintering was carried out using a carbon dice in an atomosphere of Ar+5 %H2. Then sintering pressure was 4x10^ Pa and the temperature of the dice was controlled to be kept constant between 684 and 709 K during each sintering. The obtained sintered specimens were the pellets with 15 mm in diameter and 3 mm in thichness.The apparent density of each specimen was larger than 99 %. Each pellet prepared from the portions of 20, 60 and 96 mm of Ingot 2 is named as p-1, p-2 and p-3, respectively. The p-1 was ground into a fine powder again and then sinter process under the same condition as mentioned above was carried out, since the cracks were observed in it. X-ray powder diffraction analyses confirmed that all the pellets were in single phase state of PbTe. Each pellet was cut into the parallelepiped specimens with a dimensions of 1x2x5 mm^ for the measurements of the resistivity p and the Hall coefficient RH2.3. Measurement of hysteretic temperature dependence of resistivity p during a series of heating cind cooling cycles to progressively higher temperatures Each parallelepiped specimen that was cut out of Ingot 1 and p-2 was set on an AI2O3 plate binding with Pt wire in a stainless stem with a dimension of 26 mm in diameter and 650 mm in length. Pt wire with 50 /xm in diameter was attached to the specimen as electrodes. The stainless stem was evacuated under a pressure of lxlO~^ Pa and then it was filled with pressure of 1x10^ Pa of Ar gas. The resistivities p were measured as a function of temperature during heating and cooling cycles. The heating was carried out with a rate of 2.5x10"^ K/s (90 K/h) and holding was for 3.6x10^ s (Ih) and cooling was with rate of 4.2x10"^ K/s (150 K/h). The holding temperatures were raised to progressively higher temperatures at intervals of lOOK from 423 to 903 K. The temperature of specimes were measured using the R-type thermocouples with 76 /xm in diameter which were put under an AI2O3 plate. The specimens cut out of the Ingot 1 and p-2 are named as h-m and h-p, respectively. After h-m and h-p were heated up to 903 K, they were identified by the means of the X-ray powder diffraction. 2.4. Annealing of the n-type melt-grown and the sintered PbTe Each parallelepiped specimen cut out of Ingot 1 and p-2 was sealed off under a pressure of lxlO~^ Pa in an evacuated quartz tube with a dimension of 11 mm in diameter and 70 mm in length. They were annealed at 793 K for 8.64 x 10^ s (24 h). The annealed specimens of Ingot 1 and p-2 are named as a-m and a-p, respectively, p and RH of a-m and a-p were measured at room temperature. 3. RESULTS AND DISCUSSION The resistivities piw at room temperature for p-1, p-2 and p-3 were 1.42x10"^, 1.03x10"^ and 1.04x10"^ Qm, respsctively.The carrier concentrations npT estimated from the Hall coefficients for p-1, p-2 and p-3 were 3.22x10^^, 1.83x10^^ and 1.69x10^^ m-^ respectively

578

2 2.5 XQ^IT (1/K) Figure 1. Temperature dependence of resistivity p measured during a series of heating and cooling cycles to progressively higher temperatures for the n-type melt-grown PbTe in Ar atmosphere (h-m). Max. Temperature of each cycle from 1 to 6 are 423, 503, 603, 703, 803 and 903 K, respectively.

2

2.5

3

3.5

10^/r (i/K) Figure 2. Temperature dependence of resistivity p for the n-type plasma activated sintered PbTe in Ar atmosphere (h-p) under the same heating and cooling cycles as ones for h-m. The measurement could not be carried out during coohng from 903 K because of the crack occurrence.

PUT and URT were higer and lower respectively comparing with those for the corresponding portions of Ingot 2 as the starting material. The Hall mobility ^H calculated from PRT and UJIT were less than 10~^m^/(Vs) which is two orders less than those of Ingot 2. This is mainly due to the evaporation of Pbl2 as dopant when the electric current flew directly through the powder materials and when temperature of the portion at which the particles contact each other was raised. It was also reported that p after being powdered and pressed for the n-type PbTe doped with PbBr2 considerably increased. [7] The Hall mobihty pin for the n-type hot pressed Pbo.95Sno.05Te doped with Pbl2 was decreased and the temperature dependence of fin changed drastically comparing with those of the melt grown one. [8] Then, it can be considered that powdering and sintering process forms the potential barrier at the grain boundaries and that the additional scattering mechanism besides due to the lattice vibration, ionized impurity and interaction of poin defect is introduced. [7,9,10] Fig.l is the temperature dependences of the resistivity p during a series of heating and cooling cycles to progressively higher temperature for h-m. The resistivity PRT of the as-grown one at room temperature is 2.35x10"^ Qia and was almost unchanged after heated up to 603 K. However, after heated at 703 K, PRT decreased to 2.26x10"^ Hm after heated at 803 K and finally decreased to 1.91 x 10~^ Hm after 903 K. The carrier concentration URT of the as-grown one estimated from the Hall coefficient was 1.82 x 10^^ m~^ and increased to 2.44 x 10^^ m~^ after heated up to 903 K. It is considered that those are due to the deviation of the stoichiometry caused by the evaporation of the Te with a comparatively high vapor pressure from the surface of the specimen. [11-13] Fig.2 is the temperature dependence of the resistivity p during a series of heating and

579 Table 1 PUT and nj^ for the melt-grown (a-m) and plasma activated sintered (a-p) specimens annealed in the evaquated quartz at 793 K for 8.64x10^ s. Pj^To (Hm)

a-m a-p

2.11 X 10-^ 9.36 X 10-^

PRT (^m)

URTQ (m~^)

TIBT

(m~^)

2.06 x 10"^ 8.85 x 10"^

2.92 x 10^^ 2.06 x 10^^

3.21 x 10^^ 2.19 x 10^^

PRTO and rijirpQ are the values before annealig.

cooling cycles to progressively higher temperature for h-p. ppj^ of the as-pressed one was 1.03x10"^ d m . p had a hysteretic temperature dependence after heated up to 703 K similar to that in a case of h-m. PRT decreased to 1.89x10"^ Dm after heated to 803 K. Both h-m and h-p had hysteretic temperature dependence of the resistivity after heated up to 703 K. However, the decreasing ratio was larger for the h-p comparing with the case of h-m, which indicates that the volume of the grain boundaries in the specimen relates closely to the decreasing of p. However, the mass decreasing was negligible small and X-ray analysis detected no evidences of the precipitation of the secondary phase nor oxidation. Then, it can be pointed out that the more precise and accurate structural and compositional analysis are required in order to clarify the mechanism of the decreasing of the resistivity. Table 1 summarizes the results for a-m and a-p. The change after annealing were less slight comparing with those for h-m and h-p. Then, It is found that using lead telluride as thermoelectric material in an encapsulated container is promisingly useful to prevent the changes of the thermoelectric properties. However, since it was also reported that p for the n-type sintered PbTe decreased one order of magnitude after annealed at 1073 K for 1.12x10^ s (310 h), the more detailed investigations concerning the other factors such as operation temperature, container volume, species and pressure of ambient gas are needed. [7] 4. C O N C L U S I O N The resistivities p for the n-type melt-grown and the plasma activated sintered PbTe were measured as functions of temperature in the heating and cooling cycles to progressively higher temperature up to 903 K. The irreversible phenomena in the temperature dependence of p occured and p decreased for the both specimens after heated up to 703 K. The phenomena was more remarkable for the sintered specimen than for the melt grown one, which indicates the correlation of the volume of the grain boundaries. However, the evidences of the precipitation of the second phase nor oxidation could not be detected. The changes of the thermoelectric properties for both specimens evacuated in the quartz tubes were less slight comparing with those in the cases of heating and cooling cycle. Then, it can be concluded that to use PbTe as thermoelectric material in an encapsulated container is promisingly useful to prevent the changes of the thermoelectric properties.

580 5. ACKNOWLEDGMENT The authors wish to thank Prof. T. Kojima and Dr. I.J. Ohsugi of Salesian Polytechnic for their collaboration in X-ray analyses. This work was supprted by a Grant-in Aid for Physics and Chemistry of the Functionally Graded Materials from the Education, Science and Culture. REFERENCES 1. LA. Nishida, Materia Japan, 35 (1996) 943, in Japanese. 2. K. Uemura and LA. Nishida, Thermoelectric Semiconductors and their applications (in Japanese), Nikkan-Kogyo Shinbun-sya, 1988. 3. Y. Imai, Y. Shinohara, LA. Nishida, H.T. Kaibe, K. Sato, H. Kohri and L Shiota, in Proceedings of FGM'95, Tokyo, 1995, edited by L Shiota, pp.101-106. 4. M. Weinstein and A.L Mlavsky, Rev.Sci.Instr., 33 (1962) 1119. 5. E.H. Putley, Proc.Phys.Soc.London, 68 (1954) 22. 6. J. RxDsenzweig, J. Zhang and U. Birkholz, phys. stat. sol. (a)83 (1984) 357. 7. R. Breschi, A. Olivi, A. Camanzi and V. Fano, J. Mater. Sci., 15 (1980) 918. 8. H.T. Kaibe, S. Yoneda, Y. Shimazaki, T. Okumura, Y. Imai, LA. Nishida and I. Shiota, Journal of Advanced Science, 7 (1995) 157. 9. J. Yoshino, in Proceedings of FGM'94, Tokyo, 1994, edited by I.Shiota, pp.223-228. 10. J. Yoshino, in Proceedings of FGM'95, Tokyo, 1995, edited by I.Shiota, pp.66-64. 11. LB. Cadoff and E. Miller, Thermoelectric Materials and Devices, (Reinhold Publishing Corporation, New York, 1960), pp.149. 12. R.F. Brebrick and R.S. AUgaier, J. Chem. Phys., 32 (1960) 1826. 13. Vacuum Handbook, (ULVAC Co., Ltd.,1989 edited by S. Yoshikawa) pp.133.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

581

Preparation and thermoelectric properties of IrSba M. Koshigoe', I. Shiota^, Y. Shinohara^ , Y. Imai'', I. A .Nishida*' ^ Department of Chemical Engineering Kogakuin Univ., 2665-1 Nakano-cho , Hachioji-city , Tokyo 192, Japan ^ National Research Institute for Metals ,1-2-1 Sengen , Tsukuba-city , Ibaraki 305 , Japan An attempt is made to fabricate IrSbs compound and Iri.xCoxSb3(X=0.12) solid solution by the liquid-solid phase sintering and hot-pressing techniques. The compound and solid solution were found to be single phase of IrSbs and Iro.88Coo.i2Sb3 by X-ray diffractometry, respectively. The electrical resistivity and Hall coefficient were measured in the temperature range from 80K to room temperature. The hot-pressed IrSbs and Iro.gsCoo.nSbs were p-type degenerated semiconductors with hole concentration of 9.53 and 6.80 X lO^'^/m^, respectively. The cobalt atoms act as donors in the solid solution. The hole mobilities for both of the hotpressed materials (IrSbs and Iro.88Coo.i2Sb3) were larger than for the liquid-solid phase sintered and hot-pressed materials. A high thermal stability was obtained in hot-pressed IrSbs. 1. INTRODUCTION The maximum figure of merit of thermoelectric materials is obtained at a characteristic temperature Tc. Tc of BiaTes, PbTe and SiGe alloy are around room temperature, 500K and lOOOK, respectively. Tc can be controlled by the carrier concentrations, compositions and textures in a thermoelectric material. Thermoelectric materials with a functionally graded structure have been investigated in a project supported by Science and Technology Agency of Japanese Government since 1993. In the project, high efficient thermoelectric materials in the low, intermediate and high temperature ranges are produced by forming graded carrier concentrations in Bi2Te3, PbTe and SiGe, respectively. Besides forming graded carrier concentrations, attempts have also been made to join these materials to form a stepwise graded structure parallel to the heat flow. Combining these two methods it should be possible to expand the utilization temperature range from 300K to 1300K [1] which is much larger

582 than monolithic thermoelectric materials. PbTe is a typical thermoelectric material for the intermediate temperatures and shows a high figure of merit for the temperature rangefrom400K up to approximately 800K according to the carrier concentration. However, PbTe is unstable above TOOK because of a compositional change due to the high vapor pressure of Te. IrSbs compound is an interesting thermoelectric material in the intermediate temperature range. This material is expected to have better thermal stability at high temperatures than that of PbTe, because of its higher melting point and lower vapor pressure of Sb. The IrSbs compound and the solid solution of the isostructural compound have high carrier mobilities, low electrical resistivities, moderate Seebeck coefficients and moderate thermal conductivities [2]. IrSbs belongs to a large family of compounds with the skutterudite crystal structure. This structure is composed of a cubic lattice and the unit cell contains 8 of the AB3 group lattices[3].

t

2000

20

10

Weight percent Sb 30 40 50 60 70 80 90 100 — ^ - ^ — 1

\'

i<

i

• '

•

1800

|-

;

1600

y

[

2720K M.P.

K 1400

L[

§ 1200 ^-(Ir)

1 1000

888K \ |904K

1 800' 600

(a Sb)—4

400" 0

10

20

30

40 50 60 70 Atomic percent Sb Fig. 1 Ir-Sb Phase Diagram

80

90

100

IrSb3 is formed by peritectic reaction at a temperature of 1373K as shown in Fig.l[3], so it is difficult to obtain the stoichometric IrSbs directlyfromthe molten state. Some researchers attempted to crystallize IrSbs from a Sb-rich melt[3]. In this method, the IrSba ingot contains many Sb inclusions. On the other hand, an attempt was made to form IrSba via liquid-solid phase sintering (LSPS) with an Ir and Sb powder mixture. However, the relationship between the fabrication process and the IrSbs thermoelectric properties was not clear.

583 In this work, an attempt was made to prepare IrSbs by liquid-solid phase sintering (LSPS) and hot pressing of pulverized LSPS powders. The LSPS powder were characterized by using an X-ray diffractometer. These procedures were also applied to form an IrSbs-CoSbs solid solution. The electrical conduction parameters of these materials were examined on the electrical measurement. The thermal stability of the hot-pressed IrSba was also investigated. 2. EXPERIMENTAL PROCEDURE Ir powder and a Sb ingot 99.9% and 99.9999% pure, respectively were used as the raw materials. The Ir and Sb powder mixture was encapsulated in a 10^ Torr vacuum. The powder mixture was heated at 1373K for 12 hr in an electric furnace to promote the reaction by solid state interdiffusion (hereafter referred to as the reacted material). Two small samples were removed from the upper and lower sections of the reacted material. They were pulverized and characterized by a X-ray diffractometer. The residual reacted material was also pulverized to form a powder under 32 // m in size, and the powder was hot-pressed under a pressure of 120MPa at 1330K for 15 min. in a 250kPa Ar atmosphere( hereafter referred to as the hotpressed material). The composition of Co in IrSba-CoSbs solid solution was 3at%, since it has been reported that the maximum solid solution limit of Co in Iri.xCoxSbs system is for X=0.12 [2]. A "rocking" furnace was used to homogenize the solid solution. The purity of Co powder used was better than 99.5%. The reacted and hot-pressed Iro.ssCoo.iaSbs materials were prepared by the techniques mentioned above. The reacting and hot-pressing temperatures were 1230K and 1073K, respectively. Powders from two regions of the reacted material were characterized using a X-ray diffractometer via the same procedure described above. Samples were cut (1mm X 1mm X 6mm) from the reacted and hot-pressed materials. The electrical resistivity of each specimen was measured by DC method. A magnetic field of 0.5T was applied to determined the Hall coefficient. 3. RESULTS AND DISCUSSION 3.1. Characterization by X-ray diffractometry The diffraction patterns of the IrSbs reacted material taken from the two different regions were identical to each other as shown in Fig.2, and they agreed with the pattern of a skutterudite structure. It is confirmed that a homogenous skutterudite IrSbs structure was formed by the liquid-solid phase sintering of Ir and Sb powder mixture.

584

< UUJl

jUJJl^iJLJLij^

Upper prat Lower part

LU 50

2 e

60

70

80

90

[Cu-K a 1

Fig.2 X-ray diffraction pro3fll fileoflrSb.

H^

1

3

C

M^-MUJULJJUJ

Upper prat

OF

LuJUJ UJajLj

Lower part

T

80 90 50 60 70 2 0 [Cu-K a ] Fig.3 X-ray diffraction profile of Iro.ggCoo.nSbs

20

30

40

The two parts of the ingot of the reacted solid solution also possessed the same X-ray profile as shown in Fig.3. For the case of fabrication without using the "rocking" furnace, the ingot was inhomogeneous. Rocking the ampoule during reaction was an effective way to homogenize the ingot because the molten Sb penetrated the Ir powder. 3.2. Electrical properties The electrical properties of IrSbs and Iro.88Coo.i2Sb3 at the room temperature were shown in Table 1. The reacted materials were more porous than the hot-pressed materials. It was found that the hot-pressing was effective to improve the apparent density. The carrier concentrations of the hot-pressed materials were lower than for the reacted materials for both IrSbs and Iro.88Coo.i2Sb3. This may be caused by cancellation of the holes by electrons from the oxide, which was formed during pulverizing. The electrical resistivity and Hall coefficient values of IrSbs and Iro.88Coo.i2Sb3 are shown in Table 1.

585 Table 1 The apparent density and electric properties of IrSba and Iro.ggCoo.nSbs at room temperature apparent

resistivity

d

p

RH

fJ-V

np

material

Hall coefficient hole mobility hole concentration

[%J

I.><.l(3ln.ml

IXigrV/Cl

InAVs]

I.X.1()^V^

IrSbs

Hot-pressed

97.2

5.69

6.55

0.115

9.53

Reacted

69.4

6.36

5.04

0.0792

12.4

Iro.ggCoo.iiSbs

Hot-pressed

95.4

13.9

9.18

0.0661

6.80

Reacted

75.4

13J

6,61

0.0487

9.45

loV

.10-^

10<^

p

a

•(>

•

o

•

• : Hot pressed o : Reacted

>> 10*

^ 10-^1

10' oo o o

10-^ 0

o

u8

5

10-^ 15

10

,10-^

Temp. 1000/T[1/T]

A : Hot pressed ^ Reacted Jio-^

S. 10- iffi^AAV 10-

Jio-^

110' 5 10 15 Temp. 1000/T [1/T]

Fig.4 Temperature dependence of resistivity

Fig.5 Temperature dependence of resistivity

and Hall coefficient of IrSbg

and Hall coefficient of Iro.g8Coo.i2Sb3

The temperature dependence of electrical resistivity and Hall coefficient for IrSbs and IroggCoonSba are shown in Fig. 4 and 5, respectively. The resistive behaviors of the hotpressed material is the same as that of the reacted material. However, the Hall coefficient RH for the hot-pressed material is larger than the values for reacted material. RHS are constant over the observed temperature range and carriers are in the degenerated state. u.z

•^^..^^^^^rp-O.S °cO^^Soo^^ IrSb3<;^ ^

> ^ 0.1 ^0.08 ;:§0.06

A AA ^ A

-

/J>^

Q Heating •Cooling ^

AZV^^

^ ^ ^ ^

A

-

^

| 0 . 0 4 Aro.ggCoo.iiSbs o A Hot pressed • ^Reacted

S X 1

1 1 1 1

100

1

1

1

1

1

1 1 1

1000 Temp. [K] Fig.6 Temperature dependence of hole mobility for IrSb3 and Iro.88Coo.i2Sb3

Fig.6 shows the hole mobilities of the hot-pressed material and reacted material as a function of temperature. The hole mobility ju p of each material decreased with increasing t e m p e r a t u r e . The hole m o b i l i t i e s of Iro.88Coo.i2Sb3 were lower values than IrSbs. As shown in Table 1, the Co atoms in IrSbs reduce hole concentrations and act as donors.

586 The measurement of resistivity and Hall coefficient for the hot-pressed IrSba was carried out during heating to TOOK and cooling to room temperature. The results were also shown in Fig.6 plots of fi pS during heating (O) and cooling (O). The fi p plots of heating and cooling agreed with each other and it was confirmed the high thermal stability of IrSbs. 4. CONCLUSION IrSbs compound and Iro.88Coo.i2Sb3 solid solution were prepared by two techniques, liquid-solid phase sintering and hot-pressing, and the electrical properties were investigated. It was found that the rocking process during LSPS was useful for forming a homogeneous IrSbsCoSbs solid solution. The sintering technique was effective to synthesize the IrSbs compound and IrSbs-CoSbs solid solution and to improve the apparent density. The high thermal stability of hot-pressed IrSbs was confirmed. These methods suggest the feasibility of developing high performance hot-pressed IrSbs compounds and IrSbs-CoSbs solid solutions with a FGM structure. ACKNOWLEDGMENTS A part of this research was supported by the Special Coordination Funds of the Science and Technology Agency of Japan. The authors are grateful to STA for their financial support. We would also like to thank Dr. H.Kaibe of Tokyo Metropolitan University and Mr. J.F.Atkinson of the USA Science and Technology Center for their useful discussion. REFERENCES 1) I. A. Nishida, Highly Efficient Thermoelectric Materials in FGM program, proc. JapanRussia-Ukuraine. Int'l, workshop on Energy Conv. Mater.(ENECOM '95), Sendai, (1995) 15. 2) A. Borshchevsky, J.-P. Fleurial, E. Allevato, T. Caillat, CoSbs-IrSba Solid Solutions: Preparation and Characterization, Proc. 13th ICT, New York (1995) 3-6. 3) T. Cillat, A.Borshchevsky, and J.-P. Fleurial, Thermoelectric Properties of a New Semiconductors: IrSba, Proc. 11th ICT, Arlington, Texas,(1993) 98-101.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

587

p-n Joining of Melt-Grown and Sintered PbTe by Plasma Activated Sintering Y.-S. Kanga, Y. Nodab, L. -D. Chens K Kisara^ and M. Niino^ a National Aerospace Laboratory, Kakuda Research Center, Koganezawa 1, Kimigaya, Kakuda-City, Miyagi 981-15, Japan. b Faculty of Engineering, Tohoku University, Aoba-Aramaki, Aoba-ku, Sendai, Miyagi 980-77, Japan, c Institute for Materials Research, Tohoku University, Katahira 2-1-1, Sendai, Miyagi 980-77, J a p a n ABSTRACT A p-n junction of PbTe thermoelectric material was prepared by plasmaactivated sintering(PAS) in order to investigate the junction interface of a stepwise carrier concentration FGM. The component materials were undoped ptype and 2000 or 4000 molppm Pbfc-doped n-type, respectively. The n-type component was made from a mixture of PbTe and Pbfc powder or a powder which was obtained by grinding a Pbl2 doped melt-growth crystal. The microscopic observation was carried out for the component materials, which indicated the grain growth during PAS. The electric properties were almost consistent with those for the single crystals. The characterization of the p-n junction was carried out by measuring thermoelectromotive force, voltage profile and current(l)voltage(V) relationship across the joint boundary. An abrupt change of conduction type for the thermoelectromotive force and a continuous distribution for the voltage were observed at the joint boundary. The voltage profile was different from that of the stepwise FGMs of SiGe and PbTe with abrupt change at the joint boundary. The I-V relationship in forward and reverse direction bias differs each other which was limited in low voltage region. These results indicate that the junction was successfully performed by the PAS process. 1. INTRODUCTION Functionally graded material(FGM) in thermoelectric energy conversion is expected to attain higher conversion efficiency than those of constituent homogeneous thermoelectric materials [1-4]. However, a lot of difficulties are associated with the processing of a graded structure, which are an optimized design of the carrier concentration profile, unnecessary diffusion of constituent and dopant elements, an increased resistivity at the interface of stepwise FGM

588 joint, and so on. Especially, an important research concerning the conversion efficiency of the graded thermoelectric materials is to investigate the jimction interface of a stepwise carrier concentration FGM. This problem must be solved in relation to the process of manufacture of functionally graded thermoelectric materials. The preparation of the thermoelectric materials has mainly been based on the melt-growth, hot pressing(HP) and hot isostatic pressing(HIP). On the other hand, plasma-activated sintering(PAS) or spark plasma sintering(SPS) has been developed as a new convenient process in the preparation of a variety of FGM [5]. The PAS has been lately applied to prepare thermoelectric materials such as SiGe[6], PbTe, Bi2Te3[7] and semiconductor-electrode joint[8]. PbTe is one of the best materials to construct thermoelectric generators working at intermediate temperature region. The 3-stage stepwise carrier concentration FGM of PbTe was reported to have higher resistivity than those of the constituent homogeneous materials[9], which was attributable to the effect of interface of the adjacent component, as was the case of the SiGe FGM[6]. In this research, the PAS is an attempt to prepare sintered p and n-type PbTe and construct a p-n jimction of PbTe in order to investigate the junction interface of a stepwise carrier concentration FGM. The junction was characterized by the thermoelectromotive force, voltage profile and current(I)-voltage(V) relationship in order to explore the joint quality. 2. EXPERIMENTAL The powder materials (mean particle size ~ l p m ) of undoped p-type and 2000 or 4000 molppm Pbl2-doped n-type PbTe were used as the source materials. The PAS was performed in a carbon die with 11mm in internal diameter in a vacuum of about 10 Pa under an axial load of 45 MPa, and pulse ciu-rent of 400 A at 25 V with 60 ms for 80 s, followed by heating at about 1073 K for 540 s. A p-n junction was prepared by joining undoped p-type and Pbfc-doped n-type PbTe powder. Two kinds of n-type PbTe powders were used; one is a mixture of undoped PbTe and Pbl2powders, and the other is the powder obtained by grinding a Pbfc-doped meltgrown crystal. The sintered p and n-type PbTe materials were characterized by SEM observation, after polishing and chemical etching for 4 min using the etchant(HN03: glycerine=l:l). The electrical conductivity(c^, carrier concentration(i2) and Hall mobility(jUH) were measured at 300 Kfor the p and n-type sintered PbTe. The characterization of the p-n junction was conducted at 300 K by measuring thermoelectromotive force within temperature difference of 5 K, and voltage (electric potential) distribution using 4-probe method and current(I)-voltage(V) relationship in forward and reverse bias.

589 3. RESULTS AND D I S C U S S I O N Both of the p and n-type PbTe materials obtained by PAS indicated relative density higher than 99 %. Figure 1 shows a SEM micrograph of the p-type PbTe material. The grain boundary was distinct and neither voids nor cracks were observed in the microstructure. The grain size are 20 to 70 times Fig. 1 SEM micrograph of the p-type PbTe larger than the powder particle of the source sin tered m a terial. material. This fact indicates that the grain growth occurred during the PAS process, although not in the preferred orientation. Figure 2 shows dendrites in the sintered material, indicating a trace of local melting during PAS. These complex microstructure may reduce thermal conductivity and enhance figure of merit of the sintered materials, which must be investigated in further studies. The properties of electrical conductivity((^, carrier concentration(i2) and Hall mobility(/iH) for the sintered materials of p and 4000 molppm Pbl2-doped n-type PbTe are listed in Table 1. The electric properties of the p-type material correspond to those for the meltgrown undoped p-type single crystal[10]. The result indicates that the good quality in the sintered material can be obtained by PAS in spite of the short sintering time, which may be due to the accelerated sintering through the plasmaactivated surface of the powder particle. Table 1 Electric properties of undoped p-type and 4000 molppm Pbh'doped n-type PbTe. sample

a

(Q-im-i)

n (m-3)

/XH(m2V-is-i)

undoped p-t5T)e*

2.64X104

3.87X1024

4.26X10-2

Pbl2-doped n-type*

5.20X103

3.98X1023

8.17X10-2

Pbl2-doped n-type**

1.90X105

4.90X1025

2.40X10-2

* obtained by PAS.

**obtained by melt-growth.

590 In case of the n-type material obtained using a powder mixture source material, the value n were much lower than that of the material obtained from metal-growth crystal. This fact indicates that the iodine atom was not dissolved enough to act as donor in the PbTe matrix during the PAS process. Therefore, the n-type source material was used Fig.2 SEM micrograph ofp-type PbTe sintered hereafter by grinding the Pbl2-doped melt-growth material indicating dendrite structure PbTe crystal. due to local melting. Figure 3 shows plots of thermoelectromotive force 1 —r —T 1 1 1 1 against position across the o ^ n tvnf* ^ o s^ II l y p ^ ^ ^ joint interface in the p-n o o o o o junction. The conduction type at the boundary of p-n o S IIjimction was found to change abruptly within 1 O ^ 1 mm. This result shows ° o o o o that a well-defined joint 1 1 1 1 1 1 S .11-1 boimdary can be prepared 2 3 4 5 6 by PAS process without a o Position, x/mm long distance of diffusion o UJ of the dopant in matrix material. Figure 4 shows Fig. 3 Plots of thermoelectromotive force across the I-V characteristic of the joint interface of the p-n junction between the p-n junction. The undoped p-type and 2000 molppm Pbh doped ncurrent is found to depend upon the forward and type PbTe. reverse bias and the curves evidently deviate each other in the low voltage region. Since the measurements were performed within 1 s, the deviation may not caused from the voltage of Pertier effect. Thus, the observed I-V relationship may be a characteristic feature of the p-n jimction for the sintered material, though not so distinct as is the case in a rectifying p-n junction of single crystals.

591 The measurement seems to be helpful to estimate the joining process for preparation ofFGM. Figure 5 shows the voltage distribution measured across the joint interface of the p-n junction. A continuous voltage profile measured at the joint boundary means a decreased resistivity at the interface. This result is different from that of the stepwise FGM of SiGe with abrupt changes at the joint boundary [6]. Although the origin of the interface resistivity is not clear so far, an inspection into the low resistive interface of the p-n junction will give us a solution for preparing high performance of interface in stepwise carrier concentration FGMs. 4. C O N C L U S I O N S

< E

c

o

0-^

10-^ Voltage, V /mV

Fig. 4 I-V relationship of the p-n junction between undoped p-type and 2000 molppm Phh-doped n-type PbTe sintered material.

> E 1.5

p-type

n-type OQ

0) CD (D

OO

*^ O

0.5

> The thermoelectric materials of undoped p-type 0 and Pbl2-doped n-type PbTe 79 80 81 82 78 84 83 were successfully prepared by Position, X /mm PAS with high density. The grain growth diu*ing PAS Fig. 5 Voltage profile of the p-n junction on leads to the electric forward bias between undoped p -type properties consistent with and 2000 molppm Pbh-doped n-type those for the single crystals. PbTe. The sintered n-type material using the powder mixture of PbTe with P b t was low in carrier concentration compared with the one using the Pbl2-doped melt-growth crystal. A p-n junction of undoped p-type and Pbl2-doped n-type materials was prepared by PAS process. The characterization of the junction was carried out by measuring thermoelectromotive force, I-V relationship and voltage distribution. The thermoelectromotive force measured across the joint boundary of junction

592 showed an abrupt change of conduction type, indicating a well-defined joint boundary. The observed continuous voltage profile was different from that of the stepwise FGM of SiGe and PbTe with a abrupt change at the joint boundary, which indicates that the junction in a stepwise FGM could be obtained with low interface resistivity by the inspection into the interface of the p-n junction. The I-V relationships of the junction measured by applying a pulse voltage differ each other depending upon forward and reverse bias at the low voltage region. Considering the p-n junction of the sinterted material, the bias dependence indicates the good quality of the junction. From the results of characterization of the p-n junction, we consider the junction was successfully performed by the PAS process and the sintering process should be inspected in further studies use the same method of characterization. ACKNOWLEDGMENT Part of this research was supported by the Special Coordination Funds of the Science and Technology Agency of Japan. The authors are grateful to the STA for financial support. REFERENCES [1] T. Hirano, L.W. Whitlow, J. Teraki and M. Miyajima: 3rd Lit. Symp. on Structural and Functional Gradient Materials, Lausanne, 633, (1994). [2] I. A. Nishida: Proc. of Japan-Russia-Ukraine Int. Workshop on Energy Conversion Materials, Sendai, 1, (1995). [3] M. Niino and L. Chen: Proc. XII int. Conf. Thermoelectrics (XII-ICT), Yokohama, 527, (1993). [4] Y. Noda, M. Orihashi, H. Kaibe, Y. Lnai, I. Shiota and I. A. Nishida: Proc. Xm int. Conf. Thermoelectrics (XIII-ICT), in press, (1996). [5] T.Hirai: Materials Science and Technology, A Comprehensive Treatment, ed. R.W.Cahn, P.Haasen and E.J.Kramer, Chapter 20 Functional Gradient Materials, VCH Verlag. (Weinheim), 293, (1996). [6] K. Takahashi, T. Masuda, T. Mochimaru and T. Noguchi: Proc. FGM Symp.(FGM'95), 123, (1995). [7] M.Miyajima, KFujii and T.Hirano: Proc. XH Int. Conf. Thermoelectrics(XHICT), Yokohama, 272, (1993). [8] Y.-S.Kang, Y.Noda, L.-D.Chen, S.Moriya and M.Niino: Proc. Thermoelectric Energy Conversion(TEC'96), 22, (1996). [9] Y.Noda, M.Orihashi, H.T.Kaibe, Y.hnai, LShiota and LA.Nishida: Proc. hit. Conf. Thermoelectrics(ICT96), Pasadena, in press, (1996). [10] Y.Noda, M.Orihashi and LA.Nishida: Trans. DEE of Japan, 116-A, 242, (1996).

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

593

Trial Manufacture of Functionally Graded Si-Ge Thermoelectric Material T. Noguchi, K. Takahashi, and T. Masuda Vacuum Metallurgical Co., Ltd. 516 Yokota, Sambu-machi, Sambu-gun, Chiba-ken, 289-12 Japan

A functionally graded material of SiGe was manufactured by using spark-plasmasintering(SPS) in a trial. The studied material was mainly a p-type 80Si-20Ge alloy in which doping concentration of B were graded stepwise in three levels of 0.027, 0.2, and 0.58 at %. Preliminary evaluation was made for the manufactured material samples in terms of their Seebeck coefficients together with electrical resistivities. Measurement of resistivity distribution around the FGM boundary was carried out to localize the boundary.

1. INTRODUCTION An alloy made of Si and Ge has been known as a thermoelectric material to be most useful at high temperatures. We have studied the SiGe alloy for several years[l]. To enrich the thermoelectric properties of the SiGe alloy, we have been studying the functionally graded structure of the material. A trial manufacture of the SiGe material graded in doping content was conducted by using the method of Spark-Plasma-Sintering (SPS). The alloy composition mainly studied was Si 80% and Ge 20% in atomic mass ratio, and p-, and n-type doping elements were B and P, respectively. Concentration of the dopant was varied from approximately 0.03% to 0.5% for both P and B. Many kinds of graded structures could be deigned for the thermoelectric material. However, we chose the following design; a step-like structure where 3 layers of p-, or ntype materials having different concentrations of dopants were piled up stepwise. SiGe powders of three levels in doping concentration were compacted in a Carbon mold as a pre-form, and sintered simultaneously by the SPS process. Then the sintered specimen was cut into an appropriate size by using a discharging wire cutter for measurement of the thermoelectric properties.

594 The electrical resistivity and the Seebeck coefficient of the specimens were measured up to 800 ""C. To localize the boundary between the portions of graded doping, the distribution of electrical resistivities around the boundary was investigated by the fourterminal method in which one end of voltage probes was scanned automatedly. 2. PROCESSING The functionally graded material,abbreviated as FGM, of SiGe alloy was processed in the following manner. The first step was an induction melting of mixed lumps of Si and Ge doped with P or B in vacuum or Ar atmosphere of the air pressure. The doping element B was mixed with SiGe lumps prior to the vacuum melting, however, P was dropped down into SiGe alloy being melted in the Ar atmosphere to minimize volatilization during melting. Concentration of the doping element was controlled through the above process of mixing and melting. The second step was crushing, pulverizing and milling the melted SiGe alloy doped with P or B into powder of about 2 /i m in diameter. The crushing and pulverizing were made manually with a hammer and a mortar, however, the milling was made by using a planetary ball-mill. Sintering was the third step. The SiGe powder was molded into a carbon mold of about 40 mm in diameter prior to sintering. To obtain a FGM structure, the powders with different doping concentration were pressed to form pre-formed disks. Levels of the doping concentration in atomic percent were 0.027B, 0.2B, and 0.58B for p-type 80Si-20Ge, and 0.027P, 0.3P, and 0.5P for n-type one. Three disks of the pre-formed SiGe alloys were (n^-Type} placed by turns into the mold to form a stepSi Ge Si Ge like FGM monolith structure by sintering in I n d u c t i on I n d u c t i on one process. Melting Melting The block diagram of the above mentioned P u I v e r i zat i on P u I v e r i zat i on process is shown in Fig.l. Figure 2 shows & Sieving & Sieving schematically a carbon mold filled with preT formed disks of SiGe material having different I M o l d i ng doping concentration. Typical sintering I conditions are 25'^30MPa in pressure, S i n t e r i ng around 1,285''C in temperature, and l'^2 I Cutting minutes in time. In the last step, the sintered material was cut into a size having a cross-section of about 3mm x 3mm by a Fig. 1 Block diagram of processing

595 discharging wire cutter for measurement of its thermoelectric properties. It should be noted here that no optimiza- tion was made for combination of the graded concentration in doping, or for the dimension of each portion of different doping.

SiGe coipacted Powder for FGM

/^p-type >, ijOSi-20Ge/

Carbon Punch and Die

Fig. 2 SiGe powders compacted in a Carbon mold

3.THERM0ELECTRIC MEASUREMENT As the thermoelectric characteristics of FGM samples, the electrical resistivities ( p ) and Seebeck coefficients (a) were measured in a method shown schematically in Fig. 3. It should be noted in the method that the U- or IT-shaped specimen is never p-n device but a single p- or n- type material. This was aimed to remove possible errors due to Seebeck and Peltier effects in the resistivity measurement where a large temperature gradient is given to a specimen. An apparent Seebeck coefficients were calculated as the temperature derivatives from the measured electromotive forces. .3iiiiThe resistivity, that could be also regarded Hot End as an apparent quantity, was also calculated in the manner that the electrical resistance I %^M ^ I (V/I) was multiplied by the cross-sectional area and divided by the length of the current EMF pass. Such a method as the above mentioned iMeasurement may not be always verified right for thermoelectric measurements, but would be a simpler and easier manner to make a Cold End Resistivity preliminary evaluation of FGM samples to Measurement be placed in large temperature gradients. As for the thermal conductivities, however, no Fig. 3 Thermoelectric measurement data were obtained because of difficulty system originated in our equipment of thermal conductivity measurement (the Laser Flash method) for a thick specimen such as our FGM material. In the measurement of p and a of the FGM specimens, the portion of the heavily doped was always located at the hotter side during heating.

J

596 4.THERM0ELECTRIC PROPERTIES Figure 4 shows results of Seebeck coefficient measurements for p-type materials. At the same time, for comparison, it shows results of the non-FGM samples that were measured by the same method with the FGM's. Here the notation H-M-L indicates the FGM sample.

^

Apparent Seebeck Coefficient! I

Doping FGM p -Type

8 0 S i -2 0 Qe

^

Apparent Seebeck C o e f f i c i e n t

\^

Dopinf FGM n-Type

80S

i-20Ge

\ °-^ r >e

>

>-

0.4

HLL(0.027B)

Uk((0.2B) LH(0.58B) • H-M-L

0

Hot End Temp, (^t)

^H(0.50P) 1 •H-M-L 1

1 0-2

SOS

100

200

300

400

500

600

700

800

Fig.5 Seebeck Coefficient of n-type SiGe

1 Apparent Power Factor

Resistivity

1 Doping FGM p -Type

•L(0.027P)|

Hot End Temp. CC)

Fig. 4 Seebeck coefficient of p-type SiGe

Apparent

1 0.3

a> 1 w 1 <^ 0 1 1 0

100 200 300 400 500 600 700 800

1

g o

Doping FGM p-Type

i-20Ge

SOS

i-20Ge

*- 12 o • 10 O! E

^

8

30

1 LL(0.027B)

LL(0.027B) WM(0.2B) UH(0.68B)

5 20

>• 6

1 UH(0.58B)

LH-M-L

:E 4 •M

T****

Z 2 °^ 0

Jt+tf

Si

4.«»« • • • • f l i t K ^ i t i ' - l i i l l i^kAAA A A A *

100 200

300

' '500 ''

400

'

600

1

'

'

^

700

800

Hot End Temp. CX))

Fig. 6 Resistivity of p-type SiGe

1 111 11 1111 ± J - U 0 100 200 300 400

lilt

500

Hot End Temp.

600

700

800

CO

Fig. 7 Power factor of p-type SiGe

As for n-type materials, though the measurement was made not for 11-shaped but for stick like specimens, Seebeck coefficients of the FGM as well as non-FGM were shown in Fig.5. It seems likely in Fig.5 that the n-type FGM can exceed a non-FGM in Seebeck coefficient, whereas the p-type FGM is almost competitive with the non-FGM of middle doping. The apparent electrical resistivity is shown for the p-type material of the FGM together with non-FGM's in Fig.6. The apparent resistivity of the FGM is higher than those of the non -FGM of the middle

597 and heavily doped. This is never strange because the FGM sample includes a portion of the lightly doped which has a very high resistivity. Within this work, because neither the doping content nor the length of each divided portion was optimized, the FGM sample does not always perform better compared to the non-FGM middle or heavily doped material. The disadvantage of no optimization is definitely shown in Fig.7 showing the calculated power factors. 5. RESISTIVITY DISTRIBUTION AT BOUNDARY Investigation of the boundary between graded doping is important to make the boundary location clear, and also to see how much carriers were diffused during sintering. The distribution of electrical resistivity in the vicinity of the boundary will give us interesting information on such a subject. Using the method shown in Fig.8, the electrical resistivity distribution was measured automatedly at the room temperature. To make boundary effect prominent, we prepared a SiGe specimen consisting of two phases of a large difference in doping concentration;0.027B and 0.58B. Figure 9 shows the resistivity distribution around the FGM boundary. It should be noted that the boundary layer is not sharp but is broad as wide as 3'^4mm in contrast to the case of compositional FGM sample presented elsewhere[2]. This would not be due to mixing of powders during molding, but be due to diffusion of carriers during sintering. The sintering, even in SPS process for several minutes, could have a large thermal and mechanical potentials to make B atoms diffuse easily each other in the boundary layer. The large gradient in B concentration could also have enhanced diffusion reaction. This result is very important not only to design FGM thermoelectric materials but also to carry out further development in manufacturing

Fig.8 Resistivity measurement around FGM boundary

Resistivity Distribution around Boundary 80SI-20Ge-0.027B/80S i-20Ge-0.58B

'••§-•••»«• *tt1

Position (Scale Unit:0.5nin)

Fig. 9 Resistivity distribution around FGM boundary

598 methods. 6. SUMMARY AND CONCLUSION A functionally graded material of SiGe was manufactured by using spark-plasmasintering(SPS) process in a trial The studied material was mainly an 80Si-20Ge alloy in which doping concentrations of B were graded stepwise in three levels of 0.027, 0.2, and 0.58 at%. Some preliminary evaluation was made for the manufactured material samples in terms of their Seebeck coefficients as well as electrical resistivities. Thermoelectric properties were unsatisfactory so far because of no optimization in material design such as doping content, dimensions, and geometrical structures. The boundary layer between the graded doping was investigated in the manner of electrical resistivity measurement. It was suggested that diffusion of dopant occurred during sintering and resulted in a broad boundary layer as wide as 3-4mm.

REFERENCES 1. T. Mochimaru, K.Takahashi, M. Masuda, T.Ikeno, Y.Higashiguchi, T.Noguchi and I.A.Nishida, Proc. 12th International Conference on Thermoelectrics, Nov., 1993, Yokohama 2. K.Takahashi, T.Masuda, M.Mochimaru, and T.Noguchi, Proc.FGM'95(in Japanese)

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

599

Microstructure and property of (Si-MoSi2)/SiGe thermoelectric convertor unit J. S. Lin", K. Tanihata", Y. Miyamoto' and H. Kido^ "The Institute of Scientific & Industrial Research, Osaka University, Ibaraki 567, Osaka, Japan ^ Osaka Municipal Technical Research Institute, Morinomiya, Jhoto-Ku 536, Osaka, Japan This research suggested a one-step process to fabricate the thermoelectric conversion cell with its electrode. SioygGco22 alloy and (Si-MoSi2) composite were used as a thermoelectric conversion material and an electrode material, respectively. The HIP process was employed to prepare samples. The experimental and theoretical analyses on the thermal stress and electrical resistivity identified that this process has good prospects to substitute for jointing of SiGe with electrode. 1. INTRODUCTION The Figure of Merit of thermoelectric conversion (TEC) Heating units is depending on the behavior of thermoelectric Electrode < > materials and their service temperature. Figure 1 is a schematic diagram of the TEC unit. For the application at vy/y/ ZZZZZZ2^h high temperature about 1000 °C, SiGe alloys are believed to P n be suitable for TEC cell. The /'-type of SiGe cell is Cell Cell r. connected to the «-type of SiGe cell with electrode [1]. r-EZZZZa EZZZZ3-1 Obviously, the materials used as an electrode should satisfy Cooling the requirements on the low electrical resistivity, high heat \ Load h conductivity as well as the good corrosion and oxidation resistance at high temperature. Our former research work Fig. 1 A schematic diagram of [2,3] has developed an electrode material in MoSi2/Al203/Ni the TEC unit. system, which not only exhibited good electrical properties but also excellent mechanical properties because of the strong surface compressive stress introduced by the symmetrically graded structure. However, the electrode was failed to be jointed with SiGe TEC cell due to the considerable mismatch of the thermal expansion between the MoSi2 and SiGe alloys. Further development of electrode materials is required for the efficient and reliable operation of thermoelectric power generators. The researches in this paper try to suggest a new process to fabricate the electrode of TEC unit with TEC cell in one step based upon the concept of functionally graded materials (FGMs). For this purpose, MoSi2, undoped Si and SiGe were selected as raw materials to design the (Si-MoSi2)/SiGe thermoelectric conversion unit. Preliminary experimental and theoretical analyses were carried out on the material design and evaluation of mechanical and electrical properties.

600 2. EXPERIMENTAL The raw materials used in this work were MoSi2 powder, undoped Si and SioygGeoii powders. Their nominal sizes were 1.09 jxm, 20 |Lim and 30 jim respectively. The powders in pre-determined volume fraction were mixed together by AI2O3 ball milling for over 48 hours, then dried in a vacuum furnace. The green compacts were pre-pressed under 200 MPa using CIP before HIP sintering. X-ray residual stress determination was performed on the surface of the samples prepared by HIP sintering. The measured residual stress was compared with the results calculated by the finite element method (FEM). The electrical resistivity was measured by the four probes method on the slices cut from the cylinder samples. In order to inspect the thermal stability, the samples were annealed at 900 °C for 24 hour in vacuum. The microstructure on the section was observed by scanning electron microscope. 3. DESIGN AND FABRICATION OF (Si-MoSi2)/SiGe TEC UNITS In this work, the TEC unit was imagined as a structure as shown in Figure 2. SiGe alloys possessing excellent thermoelectric properties were used as the TEC material at high temperature. The silicon has good thermal conductivity (1.5 Wcm'^deg'^ at 0 °C) and low thermal expansion coefficient (2.5x10"^ deg'^ at 0 °C); MoSi2 has low electrical resistivity (2x10"^ Q c m at room temperature). The composites of the both were expected to be an excellent electrode material with high thermal and electrical conductivity at high temperature. The electrical conductivity of (Si-MoSi2) composites would be dependent on the volume of MoSi2 in the Si matrix. Increasing the ratio of MoSi2 to Si can effectively reduce the electrical resistivity of (Si-MoSi2), however, the addition of MoSi2 will enhance the mismatch of

al E S F G M e l e c t r o d e ^ 0(2

06

SiGe TEC cell

^SiOe

E F G M electrode strode Z 3 a2 //A al w/ / / / / / / Fig.2 Structure of the sample.

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L

•

n=0.6

^

r

n=1.0

Z^

150

L

0

100 1 1 50

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0

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< -100 r

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(a)

0.0

FGM 1 1 1 1 1

2.5 5.0 7.5 Distance away from center (mm)

(b)

Fig. 3 Residual stress distribution with different n values (a) radial stress, (b) axial stress.

601 thermal expansion between the SiGe and (Si-MoSi2), which may lead to crack of the sample during cooling from sintering temperature to room temperature. In order to satisfy the requirement on the fabrication process, it is necessary to understand the behavior of the residual stress produced in the process of sample preparation. Figure 3 shows the preanalysis results on the thermal residual stress with FEM method assuming the composition distributing function, f{x) = (x/ dy, where x is the distance from the end of the sample, and d is the thickness of the FGM layer. It can be seen that the maximum radial stress, which is in the surface layer, decreases while the maximum axial stress, which is on the cylinder surface, increases with the n value rising. Therefore, in the case of the sample D, the n value of 1.0 was taken to design the graded structure, but in order to introduce a compressive stress into the surface layer, relatively lower thermal expansion material was employed for the surface layer. The samples with four kinds of structures were listed in the Table 1. Table 1 The structures and composition of the samples used in the present study (vol%) Number

B

C

D

First layer

Si

Si+15MoSi2 Si+20MoSi2

Second layer

Si+30SiGe

/

Si+20MoSi2+40SiGe Si+30MoSi2

Third layer

Si+70SiGe

/

/

Si+20MoSi2+40SiGe

Central layer

SiGe

SiGe

SiGe

SiGe

Si+15MoSi2

The samples were fabricated by HIPing method. The green compact was sealed into a borosilicate glass container with BN powder bed in vacuum, then placed into a graphite crucible [4]. HIP sintering was performed at 1250 °C which is lower than the melting temperature, 1268 °C while applying an Argon pressure of 100 MPa. The SiGe sample sintered in the same process gave a density of 3.06 g/cm^ which is 99.6% of the theoretical density, 3.07 g/cm^. Figure 4 illustrates the microstructures in the vicinity of interfaces for the sample D. The interfaces between every two neighboring layers were well bonded, and the microstructures were distributed uniformly along the radial direction.

Fig. 4 Microstructures in the vicinity of every interfaces for sample D.

602 4. THERMAL RESffiUAL STRESS The residual stress often leads to the failure of HIP sintering. It should be controlled to be lower than the strength of the material. In this study, FEM method was used to analyze the stress distribution of every kind of samples supposing a temperature drop of 1250 °C. Four-node elastic axisymmetric element was employed. Material parameters used in the calculation are listed in the Table 2. In the case of sample D, the coefficients of thermal expansion were controlled as an arrangement of «/< aj> as >asiGe, the radial stress would be in a state of compression/tension/tension/ compression, as shown in Figure 6(a). The maximum tensile stress, about 150 MPa, existed in the second layer. The stress near the interface of electrode/ SiGe was relaxed effectively. The compressive stresses on the surfaces for all values of x-ray measurement.

(a)

Table 2 Physical properties used in FEM analysis of residual stress a(xlO-^) ^ ( G P a )

Materials (vol%)

V

199 0.22 4.18 210 0.22 4.55 234 0.23 5.31 205 0.22 4.75 150 0.22 3.27 Sio.78Greo.22 a. thermal expansion; E\ Young's modules; v: Poisson's ratio Si+15MoSi2 Si+20MoSi2 Si+30MoSi2 Si+20MoSi2+40SiGe

Table 3 Residual stress on the surface of samples Sample A B C D Residual stress measured (MPa)

-40

Residual stress calculated (MPa)

17

-33

-130

35

-47

-144

the samples are listed in the Table 3 with the

(b)

Fig. 6 The contour map of the residual stress in the sample D; (a) radial stress, (b) axial stress

603 The axial stress in the sample is showed in Figure 6(b). It can be seen that the maximum tensile stress in the direction of axis occurred on the cylinder surface near the electrode/SiGe interface. It often caused samples to crack in the SiGe alloy, and should be controlled lower than the strength of SiGe (about 170 MPa [5]). 5. ELECTRICAL RESISTIVITY AND THERMAL STABILITY Figure 7 shows the experimental results of the electrical resistivity before and after annealing for four kinds of samples. Obviously, the electrical resistivity of the electrode was strongly dependant on the volume of MoSi2 in Si matrix. In the case of MoSi2 free, the electrical resistivity of electrode was about 1.5x10'^ Q c m as shown in Figure 7(a). However it decreased sharply from 1.7x10"^, 4.0x10'^ to 1.0x10'^ Qcm while the volume fraction of MoSi2 increased in the order of 15 (Figure 7(b)), 20 (Figure 7(c)) to 30% (Figure 7(d)). This 0.40

0.10

I

[ (b)i

I 11'111111 I

I'

Samj)le B

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S 0.06 i a

I 0.04

i I

%

0.02

5 10 Distance (mm)

15

0.00

I I I I I I I I I M I I I I t I I I I I I I I I I I I I

5 10 Distance (mm)

15

5 10 Distance (mm)

15

0.50 I

0.20

5 10 Distance (mm)

15

Fig. 7 Electrical resistivity of the samples, • before annealing; O after annealing at 900 °C for 24 hours.

604 phenomenon can be attributed to the compositional and microstructural changes of the SiMoSi2 composites. The microstructure observation indicated that it was not enough for MoSi2 to be adjacent each other in (Si+15vol%MoSi2) composite, as shown in Figure 8(a). For 30 vol% MoSi2, however, the MoSi2 was completely connected each other to form good conductor as seen in Figure 8(b).

Fig. 8 Microstructures of Si+15vol%MoSi2 (a) and Si+30vol%MoSi2 (b). Thermal stability is highly important to ensure high efficiency and long life of the TEC unit with graded structure. The measurements of electrical resistivity after heat treatment were illustrated in Figure 7 as well. These results identified that there were no evidently effects on electrical resistivity after a 1-day thermal exposure at 900 °C. In this study, the electrical resistivity of undoped Si and SiGe was much lower than standard undoped Si and SiGe, the diffusion of B might be responsible for it because the green compact was surrounded by the BN powder in the process of sintering. CONCLUSION It is expected that one-step sintering process of SiGe and electrode can substitute for jointing of the both, and the mechanical properties can be significantly improved by the design using the concept of FGMs. The addition of MoSi2 into the Si matrix can reduce the resistivity and control the thermal expansion mismatch between Si and SiGe. The resistivity measurement is a useftil way to examine the thermal stability of the graded composition and microstructure. REFERENCES 1.1. Nishida, Ceramics Japan, 21(1986) 516 2. Y.S. Kang, Y.Miyamoto, Y.Muraoka and O.Yamaguchi, J.Soc.Mat.Sci.Japan, 44(1995) 705 3. Y. Miyamoto et al., Proc. Int. Symp. on FGMs at the Annual Meeting of Am. Ceram. Soc, Indianapolis, April, 1996, in press 4. Y. Miyamoto, J. S. Lin and K. Tanihata, Proc. of Composite & Advanced Ceamics Materials and Structures, Jan. 7-11, 1996, Flolida. 5. R. A. Lefever, G. L. McVay and R. J. Baughman, Mat. Res. Bull., 9(1974) 863

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

605

Temperature Dependence of the Porosity Controlled SiC/B4C+PSS Thermoelectric Properties K. Kato, A. Aruga, Y. Okamoto, J. Morimoto and T. Miyakawa Department of Materials Science and Engineering, National Defense Academy, 1-10-20, Hashirimizu, Yokosuka, Kanagawa, Japan, 239

In the previous paper, it was reported that the thermoelectric properties of SiC/B^C system could be controlled by the addition of PSS at room temperature. In this report, the porous structure of these samples were confirmed by using SEM, and the measurements were carried out on the temperature dependence of the thermoelectric properties from room temperature up to 600 V (the thermal conductivity up to 300 °C ). The figure of merit of the sample (B^C: 2.0 wt.% + PSS: 5.0 wt. %) is estimated about 2 X10'^ K^^ around 600 V.

1. INTRODUCTION The thermoelectric semiconductor converts directly thermal energy into electric one. It can generate electric power even when the temperature difference between the heat source and atmosphere is much smaller compared with the one needed in conventional thermal power generation. The figure of merit Z, which is one of the measure for the thermoelectric energy conversion, is defined by the following equation: Z=aV/0 /c, where a , p and K is Seebeck coefficient, electrical resistivity and thermal conductivity, respectively. One of the criterion for the practical thermoelectric materials is that the value of figure of merit exceeds 10'^ K"\ Usually, it is difficult to increase a and at the same time decrease p and /c , because these parameters depend not only on the carrier concentration, carrier mobility, scattering mechanism in the crystal grains, but also on many factors such as grain size, microstructure of grain boundaries. At room temperature, Bi^-TCj is one of the practical thermoelectric semiconductors with its

606 large figure of merit. However, thefigureof merit of this system decreases with temperature [1]. SiC is chemically stable and it has the high mechanical strength even in high temperature region. These characteristics are essential features for the high temperature applications. At the present stage of our study, the figure of merit of the sample (SiC + 20 wt. % B^C) is 8 X10"^ K"^ at 600 V [2]. This is still 1000 times smaller than the value of practical materials, and leaves a room for improvement. PSS (polysilastylene) is one of the sintering additives in this system. Excess PSS evaporates during the sintering process and pores remain in the sample [3]. We have reported that the thermal conductivity at room temperature could be controlled by the amount of pores in this system [4]. In this paper, we report on the results of the SEM(Scanning Electron Micrograph) observation for the cross sections and also of the measurements for the temperature dependence of the thermoelectric characteristics of porosity controlled SiC + B^C 2.0wt.% samples. We also discuss temperature dependence of these parameters in the light of microscopic structure of the system. 2. EXPERIMENT and RESULT 2 . 1 . Sample preparation Sample preparation procedure is the same as described in our previous papers [2,4]. yS -SiC (average particle size 0.15 // m and BET surface area 19.5m^/g, Mitsui Toatsu Co., Ltd.), B^C (average particle size 0.7 // m, Denki Kagaku Kogyo Co., Ltd.) and PSS were the starting materials. Slurries were made from mixed powder of SiC, 2.0wt.% B^C and 0.5'-"^25 wt.% PSS in polyethylene jars with nylon coated iron balls. The mixing agent is xylene. After mixing for 20 hours and passing through 75 // m mesh sieve, these slurries were dried. The dried mixture were granulated using 500 ju m mesh sieve and pressed into 20 mm 0 x 4 mm pellet at 2 X 10^ N/m^ Then the pellets were sealed into an evacuated rubber tube and then pressed isostatically at 2 X 10^ N/m^. Each pellet was covered with the same compositional powder to prevent the change in composition and it was placed into a carbon crucible. The samples are pre-sintered to evaporate PSS at 1000 X^ for 60 miu. in an atmosphere of 1.0 atm. Ar gas flow. After furnace cooling, the sintering procedure was carried out in a furnace with a RF induction heater. First, the furnace was heated up to 1000 °C ata rate of 20 degree/min. in vacuum, then Ar gas was introduced up to 1.0 atm. Then the temperature was raised to sintering temperature (2100 *C) at a rate of 10 degree/min., the sample was kept at this temperature for 2 hours and cooled naturally down to room temperature. Sintered materials were cut into rectangular shaped specimens of 3 X 4 X S'^IO mm^ in dimensions for measurement of the themioelectric properties.

607 2.2 Sample structure The large X-ray diffraction peaks could be assigned to SiC and B^C. These peaks show that dominant structure in our SiC is 6H-SiC. The small additional peaks were found which could not be identified to the composite materials of Si, B and /or C. Figure 1 shows the PSS concentration dependence of the density. The sample with the PSS concentration of 0.5wt.% is the most dense one in all samples in this work. The sample density then decreases when PSS concentration exceeds 0.5 wt.% . The thermal conductivity of the sample also decreases with sample density.

,90 •35 80 c u •o

.S 701

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20

30

PSS concentration (wt.%) Figure 1. Initially added PSS concentration dependence of sample density and thermal conductivity. The right side ordinate (closed circle) indicates thermal conductivity measured at room temperature. The left side ordinate (open circle) indicates sample density and packing density.

SEM observations were made on the surface and cross section of the samples. One can see pores in the sample which seems to be induced by the evaporation of PSS during the sintering process. Figure 2-a is the SEM observation of the surface of the sample with PSS concentration of 0.5 wt.% sample. Pores with diameters around 1 nm can be seen. Figure 2-b is the SEM

20 /im Figure 2-a. SEM observation of the surface of 0.5 wt.% PSS concentration sample.

20 /im Figure 2-b. SEM observation of the surface of 25 wt.% PSS concentration sample.

608 observation of the surface of 25 wt.% PSS concentration sample. The size of pores increases. 2.3 Temperature dependence of the Thermoelectric properties Figure 3 shows the temperature dependence of the electrical resistivity with the PSS concentration as a parameter. The current-voltage characteristics of the sample obeys Ohm's law. The values of the electrical resistivity around 600 V is reduced by a factor of 100--1000 times compared to the values at room temperature. Figure 4 shows the temperature dependence of tiie tiiermal conductivity of die samples also with PSS concentration as a parameter. They are measured by the PPE method, proposed in the previous reports [2,4,5]. The unstable tiiermal contact at higher temperatures of tiie sample to thermocouple in our experimental system limits the range of the thermal conductivity measurements under 300 "C. The thermal conductivities of tiie samples witii 0.5, 5.0 and 10 wt.% PSS concentration decrease with increasing temperature. Hie thermal conductivities of PSS concentration 15 and 25 wt. % samples slowly increase with temperature above 150 *C. Figure 5 shows the temperature dependence of the Seebeck coefficient. The data are obtained by the conventional DC metiiod. The Seebeck coefficient of all the samples increases with temperature monotonously. The rate of increase, however, is not large. Figure 6 shows the temperature dependence of the figure of merit. In this calculation of tiie figure of merit, the interpolated values of the Seebeck coefficient a and the electrical resistivity

120 r

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A D

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D

°

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200 400 Temperature (°C)

D

600

Figure 3. Temperature dependence of electrical resistivity.

100 200 Temperature (°C)

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[-

1

300

Figure 4. Temperature dependence of thermal conductivity.

609

9 > 150

PSS 0.5 wt.% PSS 5.0 wt.% PSS 10 wt.% PSS 15 wt.% PSS 25 wt.%

o A • + 1 X

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200 400 Temperature (°C)

600

Figure 6. Temperature dependence of figure of merit.

p are used from room temperature to around 600 °C. The values of the thermal conductivity /c are also interpolated ones using polynomials under 300 °C, Above 300 "C, we have assumed that the thermal conductivity stays constant. The values of the figure of merit of all samples increase with temperature monotonously. Around 600X), 5.0 wt.% PSS initially added PSS concentration sample represents the largest figure of merit (2 X10"^ K"^) in this work. 3. DISCUSSION We discuss on the miaoscopic structure of our samples, in reference to the result of thermal conductivity measurements. A K. Collins et al. have reported that grain size of their poly crystal line sample could be estimated from the temperature dependence of low temperature thermal conductivity [6] with peak at around 200 K. They obtained values from several to 10 fi m. The grain sizes in our samples are smaller than these value from SEM observations. Although the temperature of our measurement is higher the phonon scattering on the grain and pore boundaries may still be important. Our samples can be classified into two groups according to the temperature dependence of thermal conductivity /c, Group 1: K decreases with increasing temperature. Group 2 : /c increases slightly with increasing temperature above 150 "C.

610 Samples with 0.5, 5.0 and 10 wt.% PSS and almost all samples in our previous paper [2] belong to group 1. Samples with 15 and 25 wt. % PSS belong to group 2. Acx:ording to Litovsky et a/.'s classification, group 1 and group 2 samples seems to correspond to structure with "dense grain boundaries" or "microcracks", and "porous grain boundaries with microcracks", respectively. It seems that excess PSS causes changes in the microstructure of samples in addition to the reduction in sample density. For the high temperature applications group 1 samples are preferable. 4. CONCLUSION We found that the addition of PSS has three effects. First (--0.5 wt.%) it acts as the sintering additive and increases density of the sample. Then the density and also the thermal conductivity K decreases ^parently because of pores introduced by evaporating PSS. (0.5--10 wt.%). In these concentration range, K decreases with temperature. At still higher PSS concentrations, they seem to cause some change in microstructure of grain boundaries and the temperature dependence of the sample becomes unfavorable for high temperature applications. From these results, one may infer that the optimum PSS concentration (with 2.0% B^C) is about 5 wt.% under conditions of our present work. Ai this concentration we could find Z=2 X10"^ K"^ at 600 °C. Tliis is somewhat lower but very close our best record in which the improvement of the electrical resistivity was the main theme. Acknowledgement The authors are largely indebted to Mr. M. Furuta for assistance and Prof. S. Fujimoto for discussion. References 1. K. Uemura and I. Nishida, in Netsudenhandotai to sono oyo (Thermoelectric Semiconductor and its Application) (Nikkan Kogyo, Tokyo, 1988), 33-38. 2. Y. Okamoto, A Aruga, K. Shioi, J. Morimoto, T. Miyakawa, and S. Fujimoto, Proc. 12th. Int. Conf. on Thermoelectr., Yokohama, edited by K. Matuura, 184, (1993). 3. K. Sugamuma, G. Sasaki, T. Fujita, M. Okumura, A Nakazawa and K. Niihara, J. Jpn. Soc. Powder and Powder Metallurgy, 38, 62 (1991) . 4. Y Okamoto, K. Tanaka, A Aruga, F. Furuta, J. Morimoto, T. Miyakawa,and S. Fujimoto, AIP Conf. Proc, 316, 62 (1994). 5. H. Wada, M. Watanabe, J. Morimoto, and T. Nfiyakawa, J. Mater. Res. 6, 1711 (1991). 6. A K. Collins, M. A Pickering and R. L. Taylor, J. i^pl. Phys. 68, 6510 (1990). 7. Efim Ya. Litovsky and Michael Shapiro, J. Am. Ceram. Soc, 75, 3425 (1992).

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

611

Preparation of B^C-B System Composites adding PSS and their Thermoelectric Properties A. Aruga, K Tsuneyoshi, Y. Okamoto and J. Morimoto Department of Materials Science and Engineering, National Defense Academy, 1-10-20 Hashirimizu, Yokosuka, Kanagawa 239, Japan B4C-B quasi-binary composites were prepared at 1950 °, 2100 ° and 2250 °C by pressxireless sintering in argon atmosphere using commercially manufactured B4C and amorphous B powders adding 0.5 wt.% polysilastyrene (PSS). As B9C possesses higher thermoelectric property as p-type semiconductors applied at high temperature, in this study we focused upon the lattice constant, density, electric resistivity, thermal conductivity, and Seebeck coefficient. These values as a fiinction of carbon concentration added on preparation over range 7'^20 at.% at room temperature were addressed. Also figure of merit (Z) of specimens were calculated. X-ray powder diffraction (XRD) revealed that the phase of specimens were mainly B]^2(C' Si, B)3 and a little carbon. Si was detected by energy disparsion Xray spectrometer (EDX), but SiC did not appear on XRD, though Si and C elements were due to a pyrolysis of PSS. Namely, isostructure substitution occurred at a site of C in B4C. Sintering density is less than 70 % that of B4C, especially lower in the range of 7'^9 at.% C owing to excess B. SEM observation showed a lot of opening pores, but particles were binded and grain growth occurred at higher temperature. The values of thermal diflfusivity were from 3 to 13 m^/s. Specific heat of samples was almost all 1.0 J/gK Then thermal conductivities of these composites were at a range of 4'^23 W/mK Electric resistivity (7 X lO'^'^e X10'^ Qm) tended to increase with increasing C content, while Seebeck coefficient (0.30^0.38 mV/K) showed opposite tendency. Consequently, the largest Z value of these B4C-B composites was about 2.4 X10"^ K'^ at room temperature. So it is estimated that this value of Z is between BgC and B4C, although near to B4C, and the value of ZT at 1000 °C will be predicted over 0.06.

1. INTRODUCTION Boron carbide is very hard, refi-actory solids (m.p.>2400 °C). And its thermoelectric properties are unconventional in high temperature range above

612 700 °C; that is, low electric resistivity, high Seebeck coefficient, and low thermal conductivity [1]. So it is useful material as a p-type semiconductor applied for thermoelectric power conversion at high temperature. Nuclear properties such as high cross-section and resistance to irradiation are also grateful for neutron absorption (^^B). These are attractive especially for radioisotope fueled thermoelectric generator in space. Nonetheless, boron carbide does solemnly not utilize as thermoelectric devices except for thermocouple of B^C/C. Boron carbide is easily oxidized in air above 500° C [2], and react with rare metal such as Pt and Pd [1], which is used as an electrode. However, these are only technical problems. It would use for thermoelectric devices if it were highly efficient, a low cost one and easy to make it. Therefore, we first focused on a low cost and an easy making. Starting materials selected are commercially manufactured B^C and amorphous B powders. And selected process is usual pressureless sintering. Generally, boron carbides were prepared from boron and carbon atoms by hot pressing (HP), because it is said that pressureless sintering is very difficult against covalent material such as B4C. But our selections are cheaper than HP, chemical vapor deposition (CVD) [3], and other processes [4], We did not get B9C powder commercially manufactured. So our trial is to react B4C with B so as to obtain BgC ceramics in consideration of following; the single phase regime of boron carbides extendsfi:-omnear 9-20 at.% [5], and sintering is promoted at lower temperatures by introducing free carbon or other impurities [6]. In addition, one of our attention is to make machinable ceramics [7]. Polysilastyrene (PSS) is adopted as a binder, which acts to bind B4C particles, and produces porous ceramics so that it improve the electric properties and thermal conductivity affected by the density while maintaining the physical properties of B^C.

2. EXPERIMENTAL The starting materials used in this experiment were B4C (average particle size of 0.7 Jim, Grade: #1200, Denki Kagaku Kogyo, Japan), amorphous B powder (average particle size of 0.1 jim, Rare Metallic, Japan), and polysilastyrene (Grade: PSS-100, Nippon Soda, Japan) as a binder. After mixing above B4C and O-'S.S mol B with 0.5 wt.% PSS in xylene solution, the green bodies were prepared by same manner described in a previous report [8]. First, the green bodies were slowly treated up to 1000 ° C in a flowing argon atmosphere for pyrolysis of PSS. Then each preheated body covered with each powder of similar composition in each graphite crucible were sintered at 1950 °, 2100 °, or 2250 °C for 2 hrs in an argon atmosphere. The specimens with dimensions of 3.0 X 4.0 X10 mm^ for the electrical transport measurements were cut out from the sintered bodies. Samples were cleaned in ultrasonically agitated baths of acetone and ethanol.

613 The microstructure of the specimens were examined by a scanning electron microscope (SEM; model S-2100, Hitachi, Japan) and simultaneously elements were analysed by using the energy dispersion X-ray spectrometer (EDX; model 3700-2000S, Horiba, Japan). Also B4C-B composites were identified using an Xray diffractometer (XRD; Rint 2500, Rigaku, Japan) equipped a graphite monochromator and precise lattice parameters were corrected with internal standard of Si (NBS, 640B). Electrical resistivity measurement adopted conventional four probes method. Seebeck coefficient was measured by the standard DC method. Thermal conductivity K, was calculated fi-om density, specific heat, and thermal diffusivity. Specific heat measurement was carried out by use of a differential scanning calorimeter (DSC; model 8230, Rigaku, Japan) compared with a standard material of a -AI2O3. The values of thermal diffusivity obtainedfi^oma differential phase analysis of photo-pyroelectric signal (AL-A 6 analysis) [9]. All measurements were done at room temperature.

3. RESULTS AND DISCUSSION Figure 1 shows the XRD profiles for B4C+6.5B (B^Q 5C) composites adding 0.5 wt.% PSS sintered at various temperature. And hexagonal lattice parameters of a and c for B4C-B sintered bodies added 0.5 wt.% PSS are given in Fig. 2, compared with a starting material of B4C. XRD analysis shows that some sintered bodies included a little amount of amorphous or crystallized carbon, but there was no SiC. Then all lattice parameters were slightly greater than these of B4C as raw material. 2250°C These revealed that a little amoxuit of added Si, which arisesfi:-omp3a'olysis of ,XJ PSS and which in all specimens was detected by use of EDX, was incorporated 2100°C into the B4C lattice. Furthermore, the lattice parameters slightly increased with increase in content of added B, though extension of the lattice upon incorporation 1950°C of Si and B suggested thatSi and B atoms -li vL JUJ substituted for C site. Fig. 3 shows SEM 20 30 40 50 photographs of cutting surfaces for Cu Ka 2 /9 ( deg.) B4C+5B ceramics sintered at various

Li

temperature. And density of B4C-B ceramics was plotted out in Fig. 4. Relative density is less than 70 % theoretical density

u

Fig. 1. XRD profiles of B4C+6.5B composites sintered at various temperature.

614

[•

• • 1 "

• 1 • ' '

1 1 1 1 1 1 1 1 1 1 1 1 1 1 1.

I r

, 5.615J-L a>

t fit^^fi

i

, . . . 1

- - • — 1950°C —e—2100°C --5^2250t • raw mat.

•1

j

i

5.611

.^^-.^ -> 5.605

6

8

10

12

14

16

18

20

C content initially added (at.%)

22

C content initially added (at.9

Fig. 2. Hexagonal lattice parameters (a and c) of sintered B4C-B composites as a -. function of initial C content. 1 Jim

(a) (b) (c) Fig. 3. Scanning electron micrographs of the cutting surfaces of the B4C+5B composites sintered at (a) 1950 °C, (b) 2100 °C, and (c) 2250 °C. 1.8 of B4C, especially lower (about 50 %) in the range ofl'^d at.% C owing to excess ^ 1.7 B, which prevented sintering. SEM observation showed a lot of opening pores, 1.5 h i T /r^^^^^~^^^ but particles were binded and grain growth 'JA 1.4 r // --•—1950°C occurred at higher temperature. That is, C j —e—2100t ^^-^i* Q 1.3 7*^|p* —^^ 2250°C these composites are porous ones and ,_^2 1.2 component is mainly B^g^^, Si, B)3 [10], a 6 8 10 12 14 16 18 20 22 little amount of carbon and excess Concentration (at.%) amorphous B. Roughly speaking, these are Fig. 4. Density of B4C-B ceramics close to B4C ceramics. sintered at various temperature. Figure 5 shows electrical resistivity p measured at room temperature as a function of C concentration initially added. The p of B4C-B ceramics tended to increase with increasing carbon content. This tendency is opposite with other reports [1]. For almost all specimens prepared in this study are similar to B4C except for a little substitution of C to B and Si. And . . 1

1 .

615

B 40 LI I I I I I I I I I I I I I I j I I I I I I I I I I I I I I 35

-•—1950^: -e—2100°C ^5r- 2250°C

30 25 20 15 10 OS

8

10

12

14

16

18

20

C content initially added (at.%)

Fig. 5. Electric resistivity as a function of initially added C content.

22

B u H

5 0

' ' • • • ' ' • ' ' • • * ' • • • ' • • • ' • • •

8

10

12

14

16

18

20

22

C concentration (at.%)

Fig. 6. Thermal conductivity as a function of initially added C content.

in a view of sintering temperature, it is insufficient for sintering at 1950 °C, especially for B4C. But its values sintered at 2250 °C are about 1.4X lO"^ Qm, which is good agreement with other reports [1], and minimum value of 6 X10"^ Q m for B4C+8B ceramics sintered at 2250 °C is lower than that of the recent report [11]. This proves that B4C particles were successfully binded with each other by use of B and PSS (see Fig. 3), and yet the physical properties of sintered bodies remain such as low electric resistivity even though porous ones. Thermal conductivity K of samples, is easily deducible by simple arithmetic from data of the density (see Fig. 4), specific heat and thermal diflfusivity, graphs out in Fig. 6. Among them, specific heat of all samples was 1.06 ±0.07 J/gK, which is good agreement with general values, on the groiinds that its value of B^C and B is similar to each other. The values of thermal diffiisivity were from 3 X10"^ to 1.3 X10"^ m^/s. Higher sintering temperature rises, lower the values of thermal diffusivity tend. The K of these composites were at a range of 4 ^ 2 3 W/mK The tendency was similar to that of thermal diffusivity. And all samples sintered at 2250 °C were with low thermal conductivity. The Seebeck coefficient a and figure of merit Z for B4C-B ceramics as a function of C content are given in Fig. 7 and Fig. 8, respectively. The a was always positive, and its absolute value increases with increasing carbon content except for B4C+5B. a (0.30-^0.38 mV/K), whose maxima was observed at 20 at.% C (B4C) sintered at 2250 °C, showed opposite tendency of electric resistivity (7 X lO'^^-^ex 10"^ Qm), whose minimum was at B4C+8B sample fired at 2250°C. Though the figure of merit Z is evaluated from electrical resistivity, thermal conductivity and Seebeck coefficient, the Z values showed maximimi of 2.4 X 10'^ K'l at B4C+8B composite fired at 2250 °C. Therefore, the electric resistivity affects more than the Seebeck coefficient. This material in analogy with SiC is good in high temperature range but not in low one. So in practical use it needs to combine other effective materials in low and middle temperature range for a high performance device.

616

8

^

10

12

14

16

18

20

22

C content initially added (at.%)

Fig. 7. Seebeck coefficient as a function of initially added C content.

6

8

10

12

14

16

18

20

22

C content initially added (at.%)

Fig. 8. Figure ofmerit as a function of initially added C content.

4. CONCLUSION The largest Z value of these B4C-B composites added PSS was about 2.4 X10"^ K"^ at room temperature. So it is estimated that this value of Z is between BgC and B4C, although near to B^C, and the value of ZT at 1000 ''C will be predicted over 0.06. It will be obtained a good result if we have get B9C powder with a little B, adding more PSS (about 2 wt.%), and some other additives will be able to improve the electric resistivity.

REFERENCES T. L. Aselage and D. Emin, "CRC Handbook of Thermoelectrics", Ed. by D. M. Rowe, CRC Press, New York (1995) 373. D.-H. Riu, R. Choi and H.-E Kim, J. Mat. Sci., 30 (1995) 3897. K Koumoto, T. Seki, C. H. Pai and H. Yanagida, J. Coram. Soc. Jpn., 100 (1992) 853. T. L. Aselage, D. Emin, G. A. Samara, D. R. Tllant, S. B. Van Deusen, M. O. Eatough and S. M. Johnson, Phys. Rev., B48 (1993) 11759. M.Bouchacourt and F. Thevevot, J. Less-Common M e t , 82 (1981) 219. F. Thevenot, J. Eur. Coram. Soc, 6 (1990) 205. K Suganuma, G. Sasaki, T. Fujita, M. Okumura, A. Nakazawa and K Niihawa, Funtai oyobi Funmatsu Yakin, 38 (1991) 62. Y. Okamoto, A. Aruga, H. Tashiro, J. Morimoto, T. Myakawa and S. Fujimoto, Proc. 14th Inter. Conf. Thermoelec, Ed. by A. F. loffe, Phys.-Tech. Institute, St. Petersburg (1995) 269. H. Wada, Y. Okamoto and T. Miyakawa, J. Mat. Res., 6 (1991) 1711. 9. 10. File of X-ray powder diffraction standards of American Society for Test and Materials. Inorganic Compounds; BJL2(C, Si, B)3, 19-178. 11. T. Goto,J. Li and T. Hirai, Funtai oyobi Fxmmatsu Yakin, 43 (1996) 306.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

617

Joint of n-type PbTe with Different Carrier Concentration and its Thermoelectric Properties Y.Imai, Y.Shinohara, I.A.Nishida, M.Okamoto, Y.Isoda, T.Ohkoshi, T.Fujii, I.Shiota' and H.T.Kaibe^ National Research Institute for Metals, STA, Ibaraki 305, Japan '"Kogakuin University, Tokyo 192, Japan ^'Tokyo Metropolitan University, Tokyo 192-03, Japan ABSTRACT The experimental examination and investigation have been carried out on the jointed ntype PbTe ingot and the sintered n-type PbTe compact with 3 layers of graded electron concentrations n^s to develope a higher efficient thermoelectric material. The jointed ingot composed of materials with /^7e=0.18 and 4.5Xio^Vm'' was prerared by difBasion jointing apparatus. No jointing boundary was confirmed by the opitical micro structure observation. Both the electrical resistivity p and Seebeck coefficient a of the jointed ingot were strongly affected by the lower part of ^Ze^O. 18 X 10^Vm' at lower temperature and by the higher part of nc=4.5 X lO^Vm^ at higher temperature. The sintered compact with 3 layers of ne= 3.51, 2.60 and 2.26 X lO^Vm^ was prepared by the hot-pressing technique. It was found that the graded structure of rie can perform a 7% higher effective maximum power than the homogeneous structure at a temperature difference of 280K. 1. INTRODUCTION The lead telluride compound systems are typical thermoelectric materials in an intermediate temperature region, and had been developed as space power generating materials in the 1960's[l,2]. Currently, the thermoelectric generators with the combustion heat source of LNG and LPG come into the market as self-supporting electric sources on the earth[l,2]. The figure of merit Z for thermoelectric materials with a homogeneous composition generally shows a maximum value at a certain temperature Topt and rapidly decreases above and below ^opt [1-3]. As shown in Figure 1, the Topt can be changed higher or lower by higher or lower carrier concentration in the material, respectively[4]. Theftmctionallygraded material(FGM), which have a slope of carrier concentration in the heat flow direction, can have a higher Z in a wider temperature range(a broken line in Figure 1). As a part of the national project in Japan,

618 highly efficient thermoelectric materials have been recently studied in Bi2Te3, PbTe and Si-Ge alloy systems with graded structures of carrier concentration, crystal grain size, composition, etc[4,5]. The present study was to examine the thermoelectric properties of jointed ntype PbTe with different electron concentrations n^s. The 3-layered FGM of sintered PbTe with graded n^ was also prepared by hot pressing technique, and the performance of the FGM was discussed to attain high energy conversion efficiency. 2. EXPERIMENTAL PROCEDURE

Temperature Distribution Low ^ — High I

1

11

JJ[

IV

V

Electron concentration FGM

300 400 500 600 700 800 Temperature(K)

900 1000

Figure 1 Temperature dependences of figure of merits

The mixture of weighed amounts of for n-type PbTe with various electron concentrations. Pb and Te in a stoichiometric composition of Pb/Te=l was vacuum sealed in quartz tubes with dopant of Pbt. The purity of raw materials was 99.999%, and the «e was controlled by the amount of Pbt in the range from 0.2 to 1.2wt%. The mixture was mehed at 1300K and solidified at the cooling rate of 30K/h by the rocking furnace to obtain homogeneous ingots. The rocking cycle was IHz, and the slope of temperature given to quarts tubes was 0.5K/mm. The PbTe ingots with different n^ of 0.18 and 4.5 X lO^Vm^ were cut and subsequently jointed by the difiRision jointing apparatus to form the jointed PbTe specimens. The jointing was under 1.4Pa at 1073K for 30min in an argon atmosphere. PbTe ingots with n,=3.5, 2.5 and 2.0X lO^Vm' were crashed to pieces and shifted to obtain the starting powders of 74-124 A^ m in particle size. The powders with n,=3.5, 2.5 and 2.0 X 10^Vm'' were stacked one by one in a carbon die and subsequently hot pressed to form the 3-layered FGM specimen of sintered PbTe with graded n^. The hot pressing was at 1 lOOK for 30min under 3.3MPa in an argon atmosphere. The dimension of FGM was 10 ^ X 6^mm' and each layer was 2mm in thickness. The measurement of thermoelectric properties employed the DC method with high speed and high resolution[2] to remove fully errors occurring by Peltier effect. The thermoelectromotive force EQ was measured as a function of temperature difference A 7 at both ends of a specimen. The thermoelectric power a was obtained from a slope of Eo-A Jcurve and expressed as an absolute value.

619 Tabic 1

Conduction paramcrtcrs of solidified n-lypc PbTc and jointed electron concentrations.

doped Pbl2 wtSli

0.2 1.2 0.2/1.2

Rfl

ne

u-

lO-'^xDiVC 33.88 1.40

lO^Vm^ 0.184 4.459

mVV-S 0.190 0.064

P

lO'^xfim 17.74 2.20 7.33

material with two kinds of

3. RESULTS AND DISCUSSIONS 3.1. Joint of Solidified Ingot The electrical conduction parameters at room temperature for solidified ingots with «e=0.18 and 4.5 X 10^Vm^ and their jointed material are shown in Table 1. for ingots with «e=0.18 and 4.5XlO^Vm^ were 0.2 and 1.2wt%.

The amount of P b t

The jointed material is

refered to as 0.2/1.2wt%Pbl2 in tables and figures afterwards. No jointing boundary was confirmed by the opitical micro structure observation, and the joint was satisfactory.

The

resistivity p of the jointed material is smaller than an average value(14.4jiQm) of ingots, although the length of 0.2wt%Pbl2 part from the jointed face is 3.7 times larger than that of 1.2wt%Pbl2 part.

Yoneda e/.a/.[5] have reported that p of the solidified n-type PbTe is

decreased to 15 and 50% by annealing at 703 and 803K for Ih in an argon atmosphere, respectively.

The ingots were jointed at

1073K for 30min in an argon atmosphere, 10'

resulting that the p of the jointed material is 50%

lower

than

soHdified ingots.

an average value

of

Figure 2 shows the

CJ X

temperature dependence of p for the ingots and the jointed material.

The p of the ingot

1.2wt%)Pbl2 increases

with

increasing temperature monotonously.

This

doped

with

10^

o:0.2wi%Pbl2 A:1.2wt%Pbl2 o:0.2/1.2wt%Pbl2

tendency reveals that the heavily doped ingot is a degenerated semiconductor.

On the

contrary, the ingot doped with 0.2wt%Pbl2 and the jointed material show maximum values at high temperatures, and the jointed material has the maximum at an higher temperature.

As resuhs, it was found that p

of the jointed material is affected by the part

10^

_i

I

1.5 2 2.5 3 3.5 Reciprocal of temperature (IO'VK)

Figure 2 Temperature dependences of resistiviteis for solidified n-typc PbTe doped with 0.2 and l.2wt%Pbl2 and jointed material.

620 with lower n^ at lower temperature and by 0| ~1 I I I I I I I I T" the part with higher n^ at higher 1.2wt%Pbl2 temperature. Figure 3 shows the temperature dependence of a obtained from Eo-/^T curves for the ingots and the jointed material. In the measurement of the jointed material, the length of 0.2wt%Pbl2 part is 1.5 times larger than that of 1.2wt% Pbl2 part. The 0.2 and -350 _J I I I I I I I I L_ 300 400 500 600 700 800 1.2wt%Pbl2 parts were set up at a heat Temperature (K) sink of 300K and a heater to obtain AT", Figure 3 Temperature dependences of Ihermoelectric respectively. The |a| of the ingot doped properties for solidified n-type PbTe doped with 0.2 with 0.2wt%Pbl2 increases with increasing and 1.2\M%Pbl2 and jointed material. temperature and shows a maximum value near 600K, while the |a| of the ingot doped with 1.2wt%Pbl2 and the jointed material increase with increasing temperature lineally and also the latter is larger than the former. It is difficult to judge whether the joint of thermoelectric materials may improve the maximum output power from the above results, because the ratios of (the length of 0.2wt%Pbl2 part) to (the length of 1.2wt%Pbl2 part) for the jointed material are different on the measurements of a and p. However, it is clarified that both the temperature dependences of a and p for the jointed material are affected by the higher and lower n^ parts at higher and lower temperatures, respectively. 3.2. Sintered PbTe FGM The thermoelectric properties at room temperature of the sintered PbTe FGM with different n^s and of the high(a), intermediate(Z?) and low n^ (c) layers composed FGM were shown in Table 2. The high n^ (a)- layer gives an optimum Z to the homogeneous PbTe as reported Stavitskaya[5], and the n^ is not changed from the starting solidified ingot. However, «eS of (by and (c)-layers are higher than those of the starting ingots, and the lower n^ layer has a higher increasing ratio of «e by sintering. The a and p of FGM are -85.0^V/K and 3.95piQm, respectively. The a is close to the average a value of (a)-, (b)- and (c)-layers, while the p is even larger. Takahashi et.al{6'\ have reported that the Si-Ge alloy FGM with 3 steps of different n^ has higher electrical resistance at jointed faces. The phenomenon of increasing p for the PbTe FGM is not clear yet, but it is likely that the jointed faces of the FGM cause increase of p.

621 Tabic 2

layer

(a) (b) (c) FGM

Thermoelectric properties of 3-layered FGM of PbTc with different electron concentrations and of layers composed FGM at room temperature.

a(/zV/K) -67.1 -72.3 -88.6 -85.0

p(Qin)

RHCIDVC)

/7e(1025/Di3)

2.54x10-6 2.81x10-6 3.34x10-6 3.95x10-6

1.78x10-'' 2.30x10-'' 2.75x10-7

3.51 2.60 2.26

XieCmVVs) 7.00x10-2 8.18x10-2 8.22x10-2

a2a(ff/K2m) 1.77x10-3 1.86x10-3 2.34x10-3

Figure 4 shows the temperature dependence of p for the sintered PbTe FGM and for (a)-, (by and (c)-layers shown in Table 2. All ps increase with increasing temperature, and they have typical variations for the degenerated semiconductor. However, the increasing ratio of p for the FGM is smaller than those of (a)-, (b)- and (c)-layers. p of the FGM is strongly affected by the lower n^ layer at lower temperature and by the higher n^ layer at higher temperature as same as the jointed material in § 3.1. Figure 5 and Figure 6 also show the temperature dependences of |a| and electrical figure of merit a^a, respectively. All measured specimes have maximum values for a^a. The temperature corresponding to a maximum a}a for each layer increases with increasing n^. The a^o of the FGM is lower than those of all layers below 500K because of the increased p, while is higher above 63OK. This result confirms the graded structure of A?e is effective to higher thermoelectric performance. Figure 7 shows the relationship between the effective maximum power Pmax and AT for the sintered PbTe FGM and for (a)-, (b)- and (c)-layers at cold edge temperature T^ of 500K. 10^ • :FGM 0:a-layer A:b-layer D:c-layer

X

• :FGM 0:a-layer A:b-Iayer D:c-layer

b^

>a

^ 10^

10'

1.0

1.5

2.0

2.5

3.0

3.5

Reciprocal of temperature (10"VK) Figure 4 Temperature dependences of resistivities for n-type PbTe FGM and for high(a)-, intermediate(/))and low electron concentration(c)-layers in FGM.

300

400 500 600 700 800 Temperature (K)

Figure 5 Temperature dependences of thermoelectric powers for n-type PbTe FGM and for high(«)-, intermediate(Z))- and low electron concentration(c)-layeri: in FGM,

622 The Pmax of FGM at Ar=280K is 150Wm/m^ and is about 7% larger than that of (a)-layer whose Pmax IS largest in all layers. This value shows the electric power reaches 25kW/m if the operation temperature range from 500 to 780K is given to the FGM with 6mm of thickness. From these results, it is concluded that the highly efficient n-type PbTe FGM can be performed by the optimum control of electron concentration graded structure.

300 400 500 600

700

800

Temperature (K) Figure 6 Temperature dependences of electrical figure of merits for n-type PbTe FGM and for high(o)-, intermediate(6)- and low electron concentrations(c)-layers in FGM.

50

100 150 200 250 300

Temperature difference AT (K) Figure 7 Effective maximum powers for n-type PbTe FGM and for high (a)-, intermediate (b)- and low electron concentrations(c)-layers in FGM at cold edge temperature Tc of 500K.

REFERENCES 1. D.M.Rowe and C.M.Bhandari, Modern Thermoelectrics, Holt, Rinehart and Winston, London, 1983. 2. K.Uemura and I. A.Nishida, Thermoelectric Semiconductors and Their Applications, Nikkan-Kogyou Shinbun-Sha, Tokyo, 1988. (in Japanese) 3. LB.Cardoff and E.Miler, Thermoelectric Materials and Devices, Reinhold Publishing Co., New York, 1960. 4. I. A.Nishida, Proc. Japan-Russia-Ukraine Int'l. Workshop on Energy Conv. Mater.(RNCOM'95), (1985)1. 5. S.Yoneda, H.T.Kaibe, Y.Imai, Y.Shinohara, I.A.Nishida, T.Mochimaru, K.Takahashi and T.Noguchi, J. Advan. Sci., (1996). to be pubiUshed. 6. K.Takahashi, T.masuda, T.Mochimaru and T.Noguchi, Proc. FGM Domestic Sympo. (FGM'95), Tokyo, (1995)123. (in Japanese)

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

623

Effects of plasma treatment on thermoelectric properties of Si8oGe2o sintered alloys K. Kishimoto^, Y. Nagamoto^, T. KoyanagP, and K. Matsubara^ ^Department of Electrical and Electronic Engineering, Yamaguchi University, 2557, Tokiwadai, Ube 755, Japan. •^Department of Electronics and Computer Science, Science University of Tokyo in Yamaguchi, 1-1-1, Daigakudori, Onoda 756, Japan.

A microstructure control at grain boundaries of SiGe sintered alloys has been attempted to improve their thermoelectric figure-of-merit. SiGe raw micrograins were exposed to GeH4-plasmas to be coated with Ge layers, and then were sintered. Results of the microstructure analysis suggested that the plasma-treated alloys had Ge-rich layers around grain boundaries. In these alloys, a reduction of the thermal conductivity and increases in the electrical conductivity and the Seebeck coefficient were observed. This favorable change of the thermoelectric properties may be associated with the microstructure modification around grain boundaries. 1. I N T R O D U C T I O N Silicon germanium is one of the most well-known and efficient thermoelectric materials [1]. So far, many efforts on silicon germanium[2-4] have been devoted to improve its thermoelectric figure-of-merit Z = S^GJK^ where *S, c, and K is the Seebeck coefficient, the electrical conductivity, and the thermal conductivity, respectively. In previous papers [5,6], we reported that the thermoelectric properties of sintered iron disilicides were improved by the microstructure control around grain boundaries using the plasma processing. Iron disilicide micrograins were exposed to rf plasmas of reactive gases, such as O2, SiH4, or GeH4, and then were sintered. It was found that the obtained sintered alloys had a larger electrical conductivity and a lower thermal conductivity than untreated alloys due to modification of grain boundaries. In this paper, we investigate effects of GeH4-plasma treatment on the microstructure and the thermoelectric properties of Si8oGe2o sintered alloys. 2. E X P E R I M E N T A L Si8oGe2o micrograins of 1 to 10 /xm in size were obtained by milling arc-melted SiGe

624 102 E

0.18

-—> O) (D •o

0.16

JZ •o

^

*r

CO

X

0.14

50 100 150 Treatment time (min)

200

Fig.l. Dependence of the increase in Ge concentration by the plasma treatment and half width of SiGe(lll) XRD peak on the treatment time.

o C/)

>, >

10^

r

•C^

o3 T3 C

\^ %. \ %. ^^.

treatment time • : 0 min O: 20 min 0\ 60 min A: 180 min

10^

o o

.^

a

•c 10o CD LU

10-'

SisoGego 1

,

^^ 1

,

1

iooo/r(K-^) Fig.2. Temperature dependence of the electrical conductivity of GeH4-plasmartreated SiGe sintered alloys.

ingots. The micrograins were exposed to an rf-plasma. The processing apparatus was reported previously[5]. The treatment was made under the following conditions: processing gas 5 % GeH4 diluted with Ar of 132 ml/min; rf-power 200 W; total pressure 0.4 Torr; processing time 0 to 3 h; micrograin amount 15 g. The treated micrograins were sintered by the hot-pressing under the following conditions: vacuum pressure 10~^ Torr; sintering temperature 1473 K; pressing pressure 32 MPa; pressing time 3 h. For comparison, untreated Si8oGe2o sintered alloys were prepared without treating the micrograins in an rf-plasma. The sintering conditions were same as those of the plasma-treated alloy. The microstructure of the samples was examined by the x-ray diffraction (XRD), the scanning electron microscope (SEM), and the electron probe microanalysis (EPMA). The electrical conductivity and Seebeck coefficient were measured from 300 to 1200 K. The thermal conductivity was measured by the laser-flash method at room temperature. 3. R E S U L T S A N D D I S C U S S I O N The Ge concentration of SiGe sintered alloys measured by EPMA increased with increasing the treatment time. This result indicated that a thickness of Ge coating layer on micrograins by the GeH4-plasma treatment increased with the treatment time. In XRD patterns of the GeH4-plasma-treated sample, Ge crystalline peaks were hardly observed, indicating that a part of Ge might diffuse into the grains while sintering. Ge concentration of the sintered alloys was obtained by XRD peak position of SiGe(lll)[7]. Figure 1 shows increases in the Ge concentration by the plasma treatment measured by EPMA and XRD. The half widths of XRD peaks of S i G e ( l l l ) are also plotted in the figure. The increase in Ge concentration obtained by EPMA are larger than that obtained by XRD. A difference between the two values increases with increasing the

625 1.0

0.036

0.6 —

0.8

^

> E "c

0.4

A A A A A A ~" A A _ A A

jt:

o -0.4

o c 0.030h o o

-0.8 1 ^ ^ ^ ^y O

1

400

1

0.032h

•a

-0.6

-i.o'

0.034

\>

0.0 —

0

C/)

A

0.? —

o o -0.2

^0

treatment time • : 0 min O: 20 min O: 60 min A : 180 min

J 600

Si8oGe2o LJ

^_}

^_±_

800 1000 1200

Temperature (K)

Fig.3. Temperature dependence of the Seebeck coefficient of GeH4-plasmartreated SiGe sintered alloys.

CD

0.028h

0.026

50

100

150

200

Treatment time (min)

Fig.4. Thermal conductivity versus the plasma treatment time.

treatment time as well as the half width. The Ge concentration obtained from XRD corresponds to that of only grains which are crystallized, whereas that obtained from EPMA is an average value of the sintered alloy. Therefore, it seems that a part of the Ge coating layer is still left at the grain boundary even after sintering. From these analysis and consideration on the microstructure, it seems that the plasma-treated alloy was composed of SiGe grains with Ge-rich grain boundaries and that the Ge-rich region increases with the treatment time. Figures 2 and 3 plot the temperature dependence of the electrical conductivity and the Seebeck coefficient, respectively, on the plasma treatment time for the SiGe sintered alloys. As the treatment time increases, the conductivities below 700 K increase first, and then decrease when the treatment time is 180 min. However, the electrical conductivities above 700 K are not changed by the treatment. The Seebeck coefficient below 700 K increases first with increasing the treatment time, and it converts from negative to positive when the treatment time is 180 min. The Seebeck coefficient above 700 K is hardly changed by the treatment. The increases in the electrical conductivity by the plasma treatment have been observed in the y5-FeSi2 sintered alloys[5,6]. In the case of sintered FeSi2, the plasma treatment reduced some of defects at the grain boundaries [8]. Usually defects at the grain boundaries form a potential barrier, so that the carrier mobility is lowered. Therefore, it is considered that the reduction of the defects by the plasma treatment is related to the increase in the carrier mobility. In this experiment, it is also considered that Ge coating layers may act as a sintering aid because of its low melting temperature, so that sintering may be promoted and that such defects may be reduced. Figure 4 shows the room temperature thermal conductivity of the SiGe sintered alloys as a function of the treatment time. The thermal conductivity is reduced by the plasma treatment monotonously with increasing the treatment time. The reduction of the thermal conductivity may be attributed to presence of Ge-rich layers around grain

626 boundaries, which act as a scattering center for phonons. From these results, it was found that the GeH4-plasma treatment increased the electrical conductivity and Seebeck coefficient and reduced the thermal conductivity for SiGe sintered alloys. As a result, the figure-of-merit of the GeH4-plasma-treated SiGe sintered alloys was improved about three times at room temperature as large as that of untreated ones. 4. CONCLUSIONS We investigated the effects of the GeH4-plasma treatment on the thermoelectric properties and the microstructure of Si8oGe2o sintered alloys to improve their figure-ofmerit. The plasma treatment modified the microstructure of the alloys. The GeH4plasma-treated alloys seem to be composed of SiGe grains with Ge-rich grain boundaries. For these alloys, the thermal conductivity was reduced, and the Seebeck coefficient and electrical conductivity were increased, so that the figure-of-merit was improved. The improvement of the thermoelectric properties may be associated with the microstructure modification, especially around the grain boundaries. REFERENCES 1. J. P. Dismukes, L. Ekstrom, E. F. Steigmeier, I. Kudman, and D. S. Beers, J. Appl. Phys., 35(1964)2899. 2. N. Savvides and H. J. Goldsmid, J. Phys. C: Solid St. Phys., 13(1980)4657. 3. D. M. Rowe, V. S. Shukla, and N. Savvides, Nature, 290(1981)765. 4. C. B. Vining, W. Laskow, J. O. Hanson, R. R . Van der Beck, and P. D. Gorsuch, J. Appl. Phys., 69(1991)4333. 5. K. Matsubara, N. Minemura, J. Miyata, K. Kishimoto, K. Kawamura, T. Koyanagi, and T. Miki, Proc. 10th Int. Conf. on Thermoelectrics, Cardiff, UK, (1991)40. 6. K. Kishimoto, Y. Nagamoto, K. Nagao, T. Miki, T. Koyanagi, and K. Matsubara, AIP Conf. Proc. 316 (13th Int. Conf. on Thermoelectrics), Kansas, USA, (1994)123. 7. K. Ando, Bussei, (1962)381 (in Japanese). 8. T. Miki, Y. Matsui, Y. Teraoka, Y. Ebina, K. Matsubara, and K. Kishimoto, J. Appl. Phys., 76(1994)2097.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

627

Control of Temperature Dependence of Thermoelectric Properties of Manganese SUicide by FGM Approach

Takenobu Kajikawa, Seiya Suzuki, Keisuke Shida, and Sunao Sugihara Shonan Institute of Technology Fujisawa, Kanagawa 251 JAPAN

ABSTRACT The manganese sihcide is one of interesting materials for power generation appHcation utilizing high temperature resources. In the paper the temperature dependence of thermoelectric properties is described for the sintered manganese sihcide with two kinds of dopant concentrations from 350K up to 1 lOOK. In the experiments the characteristic of the temperature dependence is possibly controlled by the combination of carrier concentration with macro-structure aiming a broad temperature dependence of the figure of merit by FGM approach. The experimental results suggest that the combined control of two kinds of parameters is effective to get a good FGM thermoelectric element due to the sintered manganese sihcide for a broad temperature rangefi-om350 to 1 lOOK. 1.INTR0DUCTI0N The manganese sihcide is a kind of 3-d transition metal sihcides. The composition is formed as MnSi^-x, for which x=0.25-0.273 had been determined {1}. The microscopic structure is hke a chimney ladder, and the unit cell of the crystal consists of a lot of combinations of Mn ans Si. For example, it is reported that the unit ceU is composed of 11 sub-lattices with 44 Mn and 76 Si for MniiSii9 {2}. This complex structure can possibly lead to reduce the lattice thermal conductivity to enhance the thermoelectric performance. Besides, it can be used up to 1200K and is safe, abundant and potentially inexpensive. Hence, the manganese sihcide is one of interesting thermoelectric materials for power generation apphcations utilized high temperature resources such as heat recoveryfi:omsohd waste incinerator system {3}. The electric properties of MnSi^ x crystal is reported to be strongly anisotropic {2}. The

628 thermoelectric properties of the polycrystalline HMS( Higher Manganese Sihcide) were investigated by E.Gro B et al {4}. The temperature dependence of the dimensionless figure of merit ZT showed that Ge-doped HMS element had a peak value of ZT around 800K Therefore, the objective of our research is to control the temperature dependence of the enhanced thermoelectric performance for the sintered manganese sihcide by Functionally Graded Material (FGM) approach aiming the efficient thermoelectric performance over a broad temperature rangefi:om350K to lOOOK. 2.EXPERMENT 2.1 Sample preparation and characterization The p-type manganese sihcide siatered element used ia this investigation were formed in the shape of sintered (7mm in diameter) cylindrical pellets by means of the processing as shown in Fig.l. The raw materials of Mn(4N) and Si(4N) were mixed in the ratio of (Mn : Si = 1:1.73) and directly melted by Radio-Frequency induction heating system in a graphite crucible under vacuimi. The small amount of Ge was added in the raw material as a dopant. In order to control the temperature dependence of Mn Si thermoelectric properties by FGM 1 : 1.73 approach, sample with various Ge Dopant Gel dopant concentrations were prepared. Intermixing The ingot was pulverized by using I Melting planetary ball null. The powder of less i Pulverizing than 70 x/m was cold-pressed and sintered at 1400-1470K in argon Cold Pressing atmosphere by a conventional I sintering method. The sample Green characterization were analyzed with Sintering SEM, X-ray di£fractometry(XRD) etc. iThermoelectric Element] The density of the sample was

Z

I

measured by Archimedes' prindple.The realtive densities of samples were inthe Fig. 1 The process of the sample preparation range of 85 to 93%. SEM photographs show the sintered samples with small grains, several large agglomerates and pores. The grain size is several micron meters to several tens micron meters. The XRD patterns were somewhat complex and identified as Mn27Si47 or Mni5Si26 mainly. According to the chemical analysis the overall ratio of Mn to Si was nearly 1.1-1.2. And the dopant was

629 contained 0.25atm.% and 1.0atm% of Ge. The former is represented as GE25 in the figure or table ,and the latter is represented as GEIO. 2.2 Measurement of thermoelectric properties The temperature dependence of Seebeck coef&dent and electrical resistivity of the sintered Mn-Si element were measured simultaneously by the power factor measurement device {5}. The temperature difference was kept at constant lOK in the temperature range up to HOOK. Hall coef&dent and electrical resistivity were measured with van der Pauw method up to 500K. The each sample is spot-welded a 50 /i m platinum wire as the electrode. The current and magnetic field were O.IA and 0.356T respectively in the Hall measurement. In the measurement of electrical resistivity the current was 3mA. 2.3 Evaluation of FGM effect By the usage of two staged graded FGM element the FGM effect on thermoelectric performance such as Seebeck coef&dent and power factor was investigated. Moreover, the thermoelectromotive force for overall temperature difference was also evaluated. 3.RESULTS AND DISCUSSIONS Figure 2 shows the temperature dependence of the Seebeck coef&dent for the p-type sintered Mn-Si elements with two Ge-dopant concentraions (0.25 and 1.0 atm%) measured fi:om 350K to HOOK. Each value of Seebeck ^iJ\J coef&dent has maximum 200 and positioned in the near••rgP intrinsic temperature range. The results were explained 150 • quahtatively on the basis of conventional scheme, for which 100 D GE10C2 at the low temperature range A GE2501 • FGM the Seebeck coef&dent xR-FGM 50 increases with temperature, and decreases over the nearn 300 400 500 600 700 800 900 1000 1100 intrinsic temperature range Temperature [K] where both electron and hole contribute to thermoelectromotive force. Such Fig.2 Temperature dependence of the Seebeck coef&dent

630 characteristics can be also said to be dominated with the scattering mechanism from structural defects and impurity as the values increase with temperature and are rather insensitive to the dopant concentrations. Figure 3 shows the temperature dependence of the electrical resistivity for the samples from 350K to HOOK in logarithmic scale . The start of intrinsic region can be obtained in t h e temperature range from 970K to lOOOK. The tendency which h a s the iacrement

of

resistivity

the

with

Q

electrical temperature

E -10.0

iadicates

that

the

scattering

mechanism is domiaant similar to the

behavior

of

the

Seebeck

coefficient. In the iatrinsic region the characteristic for each sample is not so dififerent from each other.

2

1

3

4

.9.9

^

A

D GE10C2 A GE2SC1

-10.1

.^ -10.2 .2

\

S

10

•S -10.3

.8

\

^

« -10.4 UJ

From the slope (Ego/2kB) involving

%

-10.5

the energy gap and Boltzmann o

constant respectively, the energy

Temperature 10*/T [1/K]

gap is calculated to be about 0.45

Fig. 3 Temperature dependence of the eV, whereas the energy gap of the electrical resistivity single crystal of MnSii.73 is 0.67Table 1 Semiconducting propeties of the sample 0.72eV{l}. Table

1

shows

the

6E25 6E10

Ref.4

semiconductive properties of the sintered sample with two kinds of

a

(/^V/K)

110

dopant concentrations (0.25 and

P

(10-5 Qm)

1.0

atm%)

obtained

from

120

112

2.72

2.55

1.67

0.045

0.003

the

^

(10-W/C)

0.055

measvirement of Hall parameter

n

ClO^'in-')

1.12

a n d thennoelectric power etc. at

u

(10¥/Vs)

room temperature in comparison

d

(lO'kg/m' )

20.2 4.72

1.40 17.5 4.45

22.0

1.6 5.10

with the value for Reference 4.. The values of a were similar to t h a t in the reference, but the resistivities were higher t h a n t h a t in the reference, because each sample showed low density and complex structure. The values of HaU

coefficient were larger t h a n t h a t in the reference by one

order of magnitude. At present the reason is uncertain. As the results, the dopant concentration sHghtly corresponds to carrier concentration and is around

lO^^ m-^.

Figure 4 shows the temperature dependence of the power factor from room

631 temperature to HOOK. The maxiinum value for each sample is almost equal to each other. It is about 1.05-l.lxlO-3(W/mK2), which is similar to that obtained in the hotpressed HMS doped 0.5% of Ge at 750K {4}.

1^ I.

V

a QE10C2 A GE25C1 • FQM KR-FGM

300

400

500

600

700

800

900

1000 1100

Temperature [K]

Fig.4 Temperature dependence of the power factor

300

400

500

600

700

aOO

900

1000 1100

Hot junction temperature [K]

Fig.5 FGM effect on the total voltage over the temperature difference

In order to understand the FGM effect on such thermoelectric material, the temperature dependence of the Seebeck coef&dent and power factor for a two staged graded thermoelectric element is measured shown in Figs. 2 and 4. The key "FGIVF refers to the case that the direction of temperature difference is normal. The key HFGIVP means the case that the direction is up side down. The Seebeck coefficient of FGM element is superior to the other non-FGM elements in the temperature range up to 850K. The temperature dependence of power factor for FGM element seems to be broad from 500K to 900K as compared with non-FGM elements, but the absolute value is quite low because of the high contact resistance due to no adhesion material at present. The cumulative thermoelectromotive force V= I a dT in the case of the fixed cold junction temperature Tc=344K is shown in Fig.5 to evaluate the FGM effect on total voltage over a wide temperature difference range. 4.C0NCLUSI0NS Aiming the efficient thermoelectric performance over a broad temperature range, the temperature dependence of the sintered Ge-doped Mn-Si thermoelectric element was

632 experimentally investigated to control the characteriistics ,and the preliminary experiment on FGM was carried out. The results obtained in these experiments are summarized as follows: 1) The temperature dependence of thermoelectric properties such as the Seebeck coefficient, electrical resistivity and power factor for two kinds of Gre dopant levels were clarified to shift the temperature of the maximimi power factor due to the dopant levels and sintered condition. The power factor at the maximum for the sintered Mn-Si element was obtained 1.05-1. IxlO-^ (W/mK^) as the promising thermoelectric element for high and middle temperature range. 2) The preliminary experiment on the combined FGM with the dopant concentration gradation and the macro-structure gradation showed that the temperature dependence of the power factor became broadfi-om500K to 900K as compared with non-FGM elements. 3) As the fiiture problems, the evaluation method of FGM thermoelectric performance, the processing method for the formation of FGM element and the durabihty should be established.

ACKNOWLEDGMENT The authors express our appreciation to Dr.LNishida, National Research Institute for Metals, Prof. T.Hirai, Tohoku University, Prof. I.Shiota and Ms M..Koshigoe,Kogakuin University for the fi^tfid discussion and suggestions. This research is carried out in the research project Thysics and Chemistry of Functionally Graded Materials" under the leadership of Prof.T.Hirai supported by Ministry of Education and Culture. REFERENCES 1.1.Kawasumi,M.Sakata,LNishida,K.Masumoto,J/of Material Science 16 (1981)355 2.1.Nishida,Material Science 15,2(1978)72 3. T.Kajikawa,M.Ito,I.Katsube,E.Shibuya, AIP Conference Proceedings 316(1995)314 4. E.Gro 5, V.Neu, U.Stohrer, Proceedings of 11th International Conference on Thermoelectrics (1992)107 5. T.KaJLkawa,I.Katsube,Y.Takada,K.Kuwabara,M.Inobe,T.Ohta,Proceedings of 12th International Conference on Thermoelectrics, IEEE(1994)44

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

633

Heat Sensing Device with Thermoelectric Film Laid on Insulated Metal Sheet T. Amano, N. Kamiya and S. Tokita IMRA MATERIAL R&D CO., LTD. 5-50, Hachiken-cho, Kariya city, Aichi pref., 448, Japan We investigated a thermoelectric thermal sensor which responded to heat in a short time and generated electricity to drive electric parts such as a light emitting diode or a piezoelectric buzzer. The sensor was composed of the insulated metal sheet and thermoelectric thick films on it. The device contained 8 pairs of iron disilicide thermocouple connected in series, generated 1.5 V within 5 seconds and 2.8 V within 25 seconds reacting to a gas flame. The compositions of the thermoelectric iron disilicide thick films were Feo92Si2.5Mno.08 for p-type and Feo.98Si2 5Cooo2 for n-type, respectively. The iron disilicide thick films are melted to reduce their porosity for good heat resisting property. Their peeling sensibility and residual stress were also reduced by melting. Adding Si in the film more than stoichiometric composition of the iron disilicide promoted recrystallization to a semiconducting )^FeSi2 phase and improved thermoelectric property of the film. 1. INTRODUCTION One of the most interesting application for the thermoelectric material is a heat sensing device which needs no power supply. It requires good sensibilty and appropriate output power. When the heat sensing device drives a light emitting diode (LED) or a piezoelectric buzzer to inform an existence of a heat source, the power ranging from lO'"^ to 10"^ watts is needed generally. But it is rarely that thermoelectric devices generate the above range of the output power [1-2]. Therefore we have developed a new type of thermoelectric device which can drive above electric parts immediately by heat of a gas flame. Reducing a thermal capacitance of a device by decreasing thickness is effective for good response. It is well known that functionally graded material, a composite of ceramic and metal, can reinforce the brittleness of the ceramic decreasing the amount of weight. An iron disilicide material is suitable for the sensing device because it has a good heat resisting property when a porosity is lower than 8 vol. % [3]. It is also well known that the thick film hardly shrink during sintering because of restriction of a substrate. We investigated a processing of the dense iron disilicide film laid on an insulated metal sheet. The iron disilicide systems were newly developed for the thermoelectric films having thermoelectric property as well as an ordinary sintered iron disilicide material. We also report a character of the sensing device. 2. EXPERIMENTAL PROCEDURE 2.1. Substrate The substrate was composed of a metal sheet and insulating layers which was coated on the both side of the metal sheet to be prevented from warping. Al-doped ferritic stainless steel (0.2 mm in thickness) was selected for the metal sheet. Because the value of liner thermal expansion coefficient ^ of it (1.0 X10"^ K"^) was in good agreement with that of iron silicide

634 and the heat resisting property of the metal sheet in air was approximately 1443 K. The solution of sodium-silicate and silica powder was starting material to form the insulating layer. After the solution was coated by spray printing, the substrate was dried and subsequently baked. Finally the layer was changed to porous insulating silica (O.lmm-0.2 mm in thickness), where the value of ^ was 1.3 X10"^K\ 2. 2. Iron disilicide thick film Compositions of the iron disilicide material were Feo92Si2.oMnoo8 and Feo.98Si2oCooo2- The purity of both material were 99 %. Each ingot was milled to powder of 4.0 /jm and 3.6 fj,m in median diameter, respectively. For improvement of the thermoelectric property of the iron disilicide film. Si powder (99.999 % purity, 32 ^m in median diameter) was blended with the raw iron disilicide powder in the range of 2.0^;c^2.6, 2.0^^:^2.6, where x and y are amount of Si for Feo.92Si^Mnoo8 and Feo.98SiyCooo2» respectively. These powder were kneaded with terpineol and turned into paste. The green line was drawn by squeezing the paste through a needle on the moving substrate. The width of lines were 1-2 nmi. The binding terpineol was burned out at 673 K for 30 minutes and subsequently sintering was performed in the range of 1463 -1488 K for 15 minutes under vacuum (<1.33 X 10'^ Pa). Annealing for semiconducting )3-FeSi2 phase was carried out at 1033 K for 10 hours. 2. 3. Measuring procedure The cross section of the films were observed and their porosity were calculated by the point counting method. Cracking of the films were also observed by optical microscope. The phase structure of the film was analyzed by XRD analysis using Cu-Ka radiation. Chemical dispersion of a cross section was evaluated by EPMA. An area of a cross section of the thick film was also calculated with a outline including porosity. Adhesive strength was qualitatively analyzed whether thefilmcan be peeled with a knife or not. Thermoelectric power a and electric resistivity p of the specimen were measured with a 4prove method. Thermo electromotive force E and resistance R of the sensing device were measured by d. c. method heating hot junction with a gas flame 2. 4. Thermoelectric sensing device The lines of Feo92Si^Mnoo8 and the Feo98SiyCooo2 paste were drawn alternately on the substrate. The length of lines was 40 mm and the area of cross section was approximately 2 mm^. The spacing distance between them was 4.6 mm. The two lines were connected directly for hot junction. The device contains 8 pairs of thermocouple connnected electrically in series. 3. RESULTS AND DISCUSSION 3.1. Iron disilicide thick films Figure 1 shows SEM photographs of cross sectional fracture of the Feo.98Si2.oMnoo8 thick films. The melting temperature seems to be approximately 1483 K. Table 1 shows the porosity of the films. The porosity is larger than 21 vol. % at 1478 K, while porosity decreases to 0 vol. % when the film was melted at 1483 K. These results show that the substrate restricts shrinkage of thefilm.The porosity can be decreased down to 8 vol. % when thefilmis melted. With increasing sintering temperature, the adhesive strength was decreased. However, the molten film could not be peeled off. The cracks were also increased with increasing sintering temperature and the very few cracks were found in the moltenfilms.This result shows that the residual stress caused by shrinkage of the film increases with increasing the sintering

635 temperature and released when thefilmwas melted.

a) 1463 K b)1473K c) 1478 K d) 1483 K 25 fmi Figure 1. Broken surface of the Feog8Si2.o Mno.os thickfilmssintered in various temperatures. Table 1. The porosity of the thick films sintered in various temperatures Sintering temperature (K)

1463

Porosity Feo.92Si2.oMnoo8 (vol. %) Feo.98Si2.oCoo.o2 *sintered film, ** molten film.

53* 41*

1473 31* 35*

1478 21*

1483

1488

Q**

13*

2**

The Feo.92Si20Coo.02filmwas also investigated by the same manners described above. It was found that both peeling and cracking characteristic for the film was similar to these of the Feo.98Si20Mno.08 film. The melting temperature of the Feo.92Si2.oCoo.02 film was approximately 1488 K. In addition, the adhesive strength of the Feo.92Si2.oMno.08filmseemed lager than that of the Feo.98Si2.oCoo 02film.Figure 2 shows a cross sectional microphotograph of molten films. It shows that a contact angle of The Fco 92Si2 oMno.o8filmis larger than that of The Feo.92Si2.oMno.08 film. It is suggested that the diffusion of Mn into the insulating layer effects on a wettability of the film and the adhesive strength was improved. Besides the shrinkage of the film during the sintering and difference in the linear thermal expansion coefficients between the layers, the residual stress is caused by both the shrinkage of the film by solidification of the melting iron disilicide material and an expansion of the film by recrystallizing (measured values were 4 - 5 %). These factors caused warping of the device or peeling and cracking of the film.

Feo.92Si2.oMno.08 b) Fe0.98Si2.0Co0.02 Figure 2. Cross sectional microphotograph of molten films. 1; metal sheet, 2; insulator, 3; thermoelectric film.

0.2 mm

Figure 3 shows the thermoelectric properties of the Feo.92Si2.oMno.08 film. Figure 4 also shows the thermoelectric properties of the Feo92Si2oCooo2film. These values are compared with that of ordinary sintered specimen[4]. It is found that the values of a and p of the film are smaller than those of ordinary sintered materials.

636 Figure 5 shows the XRD patterns of the Fe098Si2.0Co0.02 films. It is found that these microstructures consist of the ^FeSi2 phase and the residual e-FeSi phase. The recrystallization was caused by the decomposition of a-Fe2Si5 -^ ^FeSi2 + Si and the successive reaction of e-FeSi + Si -> )S-FeSi2 [5]. Therefore the amount of the ^FeSi2 phase will be unable to increase any more. It is expected that the amount of the j^FeSi2 phase increase by adding Si more than stoichiometric composition of the j^FeSij phase.

600

^ \ 0)

^ x=2.6 ' 2.4'^ 400 . 2.5^

500

vv

I

300 -

o

200

s E

10-2

100 0 200

2.2'

10-3

\\\

~~~^\N

2.0^

CL 10-4

10-5

400 600 800 1000 1200 Temperature 7/K

1.5 2.0 2.5 10007-1/K-1 a) Thermoelectric power b) Resistivity Figure 3. Thermoelectric properties of the Feo.92Si,Mno.o8 molten film.

10-3 E

to E

10-4

0)

400

600

800

1000 1200

Temperature T/K

1.5

2.0

10007-1/K-1

a) Thermoelectric power b) Resistivity Figure 4. Thermoelectric properties of the FcoggSiyCoo02 molten film. Figure 3-4 also shows the thermoelectric properties of these new materials. The value of a increased with increasing Si content, while the value of p increased constantly. The maximun value of a was obtained at x=y=2.5. These results were also confirmed by phase analysis shown in Figure 5. Si content increased with decreasing e-FeSi phase. The e-FeSi phase almost disappeared at x=y=2,5. On the contrary, the Si phase was appeared in the range of jc>2.5 and y>2,5 where the value of a was saturated and the value of p increased further. Not only the value of a of the Fe098Si2.5CO0.02 film but also the value of p were higher than those of the

637 ordinary sintered Feo.98Si2.oCoo.02, while the Feo.92Si2.5Mno.08 film corresponded to the ordinary sintered Feo.97Si2.oMno.03 material [4] because of evaporation of Mn. A lack of Si is supposed to be caused by selective oxidation of Si [6] and evaporation rate of Si is larger than that of Fe. But these causes are not sufficiently related to above results. Molar ratio of Mn and Co in each film were evaluated by the EPMA analysis and found to be 0.03 and 0.02, respectively. The dispersion in the direction of thickness and diffusion in the insulator were observed in the case of Mn as shown in figure 6.

0

^

Cu - Ka

e) (0

'c 13

d)

_^ ^__^^____

i^Jv.^-^^

0

j8 - FeSi

A

a - Fe Si

•

e - FeSi Si

T

^

2

2

5

c) .

CO

c

b)

_c

AJ

_J/

A

.

*

a) 1

20

1

30

1

_iit___

1

40

50

60

2(9/deg Figure 5. Result of XRD analysis of the Feo98SiyCoo.o2film. a) y=2.0; as melted, b) y=2.0; as annealed, c) y=2.2; as annealed, d) y=2.4; as annealed, e) y=2.6; as annealed.

a) SEM image b) the image of Mn -Ka 0.1 mm Figure 6. Cross sectional photos of SEM image and image of Mn-Ka of the Feo.92Si25Mno.08 films. 1; metal sheet, 2; insulator, 3; thermoelectric film. 3.2. Thermoelectric sensing device The device was shown in figure 7. The size was 40 (W) x 60 (L)x 0.9 (T) and the weight was about 5 g. The electrical output of the device loaded an LED (Stanley Co., Inc., BR5334S) was evaluated with a gas flame as shown with figure 8. The value of E increased steeply reacting to a gas flame and decreased gradually after the temperature of the device was saturated by the gas flame within 30 seconds. The value of i^ decreased steeply accompanying the change of the value oiE and gradually after the value oiE was saturated. The temperatures of hot junction and cold junction were saturated at approximately 890 K and 350 K, respectively. As the value oiE reached 1.5 V within 5 seconds, the device lighted the LED. The maximum output was nearly 3 mW.

638 The resistivity of the hot junction was extremely high because of the difference in the melting temperature between Feo.92Si2.5Mno.08 material and Feo.98Si25Coo.02 one. As the former material melt earlier than the latter, the former material penetrated into the latter in the sintering process through the hot junction by the capillary effect. Fortunately the resistivity of the hot junction decreased by the hot junction being heated as the device was operated. For the cold junction, it was better to connect them electrically by soldering or using electrical conductive adhesive to avoid the defect described above. The thick film shrunk during the sintering process and expanded during annealing process. Though these factors arose the warping of the device, the degree of warping was adjustable by dimensions of the film and thickness of the substrate. Therefore the device could be finally flattened. > —.

E: open circuit

UJ

o E

/ Figure 7. The view of the heat sensing device. 4.

Figure 8. The thermoelectric properties of the heat sensing device. The LED begins lighting up above £"=1.5 V.

CONCLUSIONS

The heat sensing device with the thermoelectric thick film laid on insulated metal sheet was investigated. The device well responds to a heat source and generated enough power to drive an LED. The thick films were made from the thermoelectric iron disilicide materials whose compositions were Feo.92Si25Mno.08 and Feo98Si25Coo.o2 for p type and n type, respectively. It was found that melting was effective to decrease the porosity of the film and to increase mechanical strength of the device. It was also found that, by adding Si more than stoichiometric composition of the j^FeSi2, the residual f-FeSi phase could be minimized and thermoelectric property of the film could be maximized as much as that of ordinary sintered iron disilicide material. REFERENCES 1 2 3 4 5

D. M. Rowe, Proc. 12th ICT (1993) 429. H. lizuka, T. Yamazaki, M. Komabayashi, Proc. 12th ICT (1993) 295. T. Kojima, N. Hiroyama, M. Sakata, J. Materials Science Soc. Japan 28 (1991) 252. T. Kojima, M. Okamoto, I. A. Nishida, Proc. ICTEC, Arlington, March, (1984) 56. I. Isoda, K. Masumoto, T. Kojima, I. A. Nishida, K. Tanaka, Abstracts of Japan Institute of Metals, April, (1985) 175. 6 K. Herz, M. Powalla, Appl. Surf. Sci. 91 (1995) 87.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

639

Recent Developments in Oxygenated Thermionic Converters *. J.-L. Desplat Rasor Assoc. Inc., 5670 Stewart Ave., Fremont, CA 94358, USA Addition and control of oxygen in thermionic converters is the key to achieving high efficiencies and power densities. Converters with collectors containing cesium oxide, and with cesium oxide vapor circulation have shown barrier indices near 1.9eV. A new method, based on sputtering and readsorption for measuring the vapor pressures of both cesium and oxygen-bearing species in research converters was developed, hi addition, the vapor phase in research devices operating under similar conditions was sampled by time-of-flight mass spectrometry: vapor pressures of CS2O around 1x10"^ Torr were detected in Cs pressures around 1 Torr. These CS2O pressures were consistent with those obtained using the readsorption technique.

1. INTRODUCTION Parametric studies have shown that, in order to achieve high efficiencies and power densities at realistic spacings and emitter temperatures, it is necessary to have a high emitter bare work function O^Q at or above 5.2eV and a barrier index Vg at or below 2 eV. One approach successfully implemented in the past was to incorporate oxygen in the collector layer by subliming molybdenum in an oxygen atmosphere [l]. Extensive performance mapping of such converters has shown that the release of oxygen is increased as the collector temperature is raised, resulting in higher O^Q [2]. This linkage however restricts the range of collector temperatures available to design thermionic nuclear reactors by establishing a constraint on the rejection temperature. Rasor Associates therefore initiated a research program with the DOE to test other approaches and in the process gain some insights about controlling the oxygen addition. This paper is mostly a summary of a report available from the DOE [3l

2. NEW APPROACHES IN OXYGENATED CONVERTERS 2.1.Cesium oxide loading in collectors The mechanism by which the sublimed oxygen-containing Mo collector leads to better performances is not well understood. It is commonly assumed that Cs reduces the molybdenum oxide formed during the sublimation, the resulting volatile CS2O then dissociating on the emitter surface. To bypass the molybdenum oxide reduction step, the first approach tried was to load commercial "cesium oxide" (composition close to that of CSO2) in hollow collectors. * This research was sponsored by the U.S. Department of Energy

640 Figure 1 shows two versions of such converters assembled around a planar variable-spacing research converter body built by ThermoTrex: one was fitted with a collector whose cavity contained - 0.6g of CsO^ , the other had a porous W disc in which CsQ had soaked in after melting, hi both cases cesium vapor was supplied by an outside liquid reservoir.

COLLECTOR BASE-

Figure 1. Two versions of planar converters with CsO^ - loaded collectors. After extensive activation both designs achieved good performance: OgQ-^ 5.5eV and Vg at 5A/cm^ < 2eV. (see Figure 3 below for some JV curves). However given the complexity and expense associated with variable-spacing research converters, a new test vehicle was designed to allow faster iterations: Builtfi"omdemountable standard UHV components, it featured a 7cm stainless steel cube to which variousflange-mountedelectrodes and instrumented ports could be attached, hence its name - the "cube" converter. 2.2. Development of a vapor circulation source Early tests in the cube consisted of investigating the performance of fixed-spacing planar converters in which a pool of liquid Cs was gradually oxidized. However cesium oxide migration out of the reservoir was observed, against the temperature gradient, leading to a segregation of the oxide in a hotter region of the converter. To confine the oxide to the reservoir a technique already used with molten alkali salts was adopted [4]: a porous platinum disc, resulting fi'om partial sintering of platinum black, was placed in the reservoir. Figure 2 shows the cube with its OXYGEN DETECTOR electrodes, the reservoir with the Pt pellet, and the newly developed oxygen detector (see section 3.). The emitter was a W layer CVDed fi-om WClg, the collector polycrystalline Mo, the spacing ~0.6mm. Cesium POROUS Pt PELLET could be distilled in measurable amounts into the reservoir. Two such converters with similar Pt pellets exhibited very good performance without activation Figure 2. "Cube" converter. : Both achieved barrier indices below

641 2 eV and O^Q above 5.5eV. Figure 3 shows that the performance of the first one was nearly identical to that of the converter with the CsO^ - containing collector. Figure 4 shows that the second one had a Vg of 1.91eV at the beginning which later settled at 1.97eV, during a limited life test. \ \ VB = 2.0« V I

>vM8k

l

"

^

TcyBE=710K Tcup'VAR.

\ \ \ \ \ \ \

N

\ ^

Tc=670K

I

V s^9k fe 15

Tg« 1800K

'"••

516k

— PL E-4 HE T-4R

" ^

XAV'^ 02

0.4

0.6

0.8

::.225^^

1.0

0.4

ELECTRODE VOLTAGE(V)

Figure 3. JV curves for two converters: HET4R with CsO^ -collector, PUE4 with Pt pellet reservoir.

0.6

0.8

VOLTAGE. V

Figure 4. JV curves during a short life test of PUE5 with Pt pellet reservoir (no further change from 600h to lOOOh).

EXTERNAL VAPOR FtOW

Figure 5. Processes assumed to be occurring in DECOR. These data together with visual observations of the reservoir and SEM examinations of the sintered Pt lead to a description of the cesium-loaded Pt pellet as a "dynamic equilibrium cesium oxygen reservoir" (acronym: DECOR) capable of generating pressures of Cs and oxygen-bearing

642 species and circulating them throughout the converter. A brief summary of this model follows: 1. Due to its large specific area the sintered Pt contains a significant amount of adsorbed oxygen. Cesium addition results in the formation of a dilute Cs/0 solution. 2. The internal structure of the Pt pellet is bi-porous, consisting of a continuous web of small pores,filledwith a Cs/0 liquid, around which large interconnected pores carry a Cs/0 vapor. 3. This bi-porous structure under a heat flux (e.g. from the electrodes) generates internal liquid and vapor flows [5], and extemal flows as well. Figure 5 shows the resulting concept. 4. The thermal gradient within the pellet generates composition gradients in the liquid and vapor. 5. Chemical equilibrium establishes the local Cs/0 liquid and vapor composition. 6. The oxygen concentration in the extemal vapor is determined by the temperature and composition of the liquid at the hot end of DECOR. 7. Thermochemical calculations suggest that the material holding the Pt pellet is involved in determining the steady state oxygen content of the liquid film at the cold end of the pellet.

3. OXYGEN PRESSURE MEASUREMENT 3.1 Basic Method The technique consists first of removing the adsorbed oxygen layer through in-situ sputtering by Cs"*" ions generated during a voltage sweep of the emitter current, and then to monitor the rate of oxygen readsorption through its effect on the emission. As oxygen is readsorbed on the emitter, the Cs coverage also rises concurrently, resulting in a rise in emitted current. The relationship between the current rise and the "equivalent" oxygen pressure was derived by Ned Rasor [3] under the following assumptions: •small oxygen coverage, GQ <0.05, so that the bare work function rise A^g stays proportional to OQ , the correlation between them being derivedfi'omgaseous O2 adsorption on W [4]. •sticking coefficient ~ 1, which also results from the small coverage, •the arriving oxygen species is molecular oxygen and the adsorbed species atomic oxygen (as if O2 gas were impinging on a bare emitter), hence the term "equivalent" oxygen pressure ^^ • Under these approximations: Po, (Torr) - UxlO-'" (7^ /t) Ln(IA,),

(1)

where IQ is the emitted current at the end of the sputtering (t=0), I is the current at time t and Tg the emitter temperature in K. Thefilamentis operated near the maximum of the S-curve so that the emission current is then proportional to the cesium pressure: Given the filament and collector dimensions, and excluding the current collected by a guard ring surrounding the colder filament ends, this translates into Pcs(Torr)~4]o (A).

(2)

3.2. Implementation As seen in section 2.3, a special "oxygen detector" was added to the research converters . It consisted of a simple W filament (length ~ 1cm, diameter ~ 0.15mm) directly heated, acting as the emitter, surrounded by a cylindrical collector and guard ring . hi order to minimize the effect

643 of filament temperature variations (such as resulting from electron cooling and ion heating) on the emitted current, the filament was operated near the maximum of the "S-curve", in our case a t - 1300 K. The sputtering of this filament was achieved by simply raising the amplitude of the 60 Hz voltage sweep used to measure the filament emission. The sputtering threshold was very sharp and reproducible at 7.5 volts, but the actual voltage used was 15-20 volts to remove the oxygen in 1 to 4 half cycles . A computer program was later designed to handle the acquisition of the filament I-V curve as afimctionof time and calculate the oxygen and cesium pressures [6]. The same method can be applied to an actual converter (with considerable caution because of the magnitude of the currents involved): Figure 6 shows the evolution of a JV characteristic as oxygen was being readsorbed.

.4

.6 .8 VOLTAGE IV)

1.0

1.2

Figure 6. Successive JV curves after sputtering a converter emitter: PQX ~ 2x10"^ Torr.

4. SAMPLING OF THE GAS PHASE BY TIME-OF FLIGHT MASS SPECTROMETRY In order to identify the oxygen-carrying species responsible for the enhanced performance, it was decided to sample the converter gas phase by mass spectrometry. Since cesium can form suboxides with amu around 1,500 (CS11O3) and because quadrupole mass spectrometers suffer from a transmission loss as the amu rises, a time-of flight mass spectrometer (TOFMS) was chosen. A special, short line-of-sight valve was designed to interface between the TOFMS and a small orifice inside a cube converter. Between the valve and the TOFMS itself extra features were added to reject the main ions (Cs"^) and improve the sensitivity.

644 4.1. Sampling bakeable line-of-sight valve To minimize the distance between the cube converter and the TOFMS, a custom bakeable valve was assembledfromcomponents of a commercial high temperature valve and regular UHV hardware. Figure 7 shows a crosssection ofthe sampling valve (overall length: 7.1 cm): • a removable aperture with an orifice 0.25mm in diameter is mounted on a 7cm flange attached to the cube. • on the exit side of the valve a STELLITE TIP cooled skimmer defines a narrow molecular beam. Since the entrance side is at -- 300 C, there is a large temperature gradient along the valve body. • the valve handle is actuated by a computer-controlled torquesensing tool, which allows fast and reproducible opening and SKIMMER GAS INLET closing. After exiting the valve the molecular beam is ionized in a conventional Figure 7. Cube sampling valve ionizer. Then before entering the actual flight tube of the TOFMS, the ions are subjected to mass filtering to reject the Cs^ ions which would otherwise overwhelm all others, and then stored to boost the detection sensitivity. 4.2. Ion mass filtering and ion storage To prevent Cs^ ions from reaching the TOFMS a quadrupole lens acts as a high-pass filter : To that effect the low mass setting is at - 140 amu and there is no mass sweep. HELIUM INLET ' QUADRUPOLE LENS

TIME OF FLIGHT MASS SPECTROMETER * p=

Figure 8. A quadrupole lens filters out unwanted ions and an ion trap enhances the sensitivity.

645 To increase the sensitivity the remaining ions are sent into a quadrupole ion trap where they accumulate and are slowed down through collisions with rare gas atoms: Helium is injected into the trap thrpugh a leak valve to keep the pressure in the trap at ~ 1x10"^ Torr. Figure 8 shows the quadrupole and ion trap assembly. The loading time of the trap and the time during which the ions are stored in the trap are both adjustable to increase the sensitivity: As seen later sensitivities of 30 parts per billion were easily obtained. 4.3. Sampling of the gas phase generated with a circulation vapor source A circulation vapor source similar to the DECOR source previously studied was placed inside a cube converter attached to the TOFMS as explained above. This cube also featured an oxygen detector to get independent oxygen pressure measurements. To reduce spurious signals showing as cesium oxides peaks with the sampling valve closed, it was necessary to remove all grids in the ionizer where these oxides would condense and be released due to radiant heating from the filament. Some spurious signals still remained at a much reduced level, attributed to desorption from the side of the valve tip facing the TOFMS. Nevertheless as seen in Figure 9, spectra with the valve opened and closed showed clearly that the only significant oxygen-carrying gas was CS2O. These data were obtained with ion trap

CsjO

Scan it 0825951S VALVE OPEN Atten. Setting = Odb Gain Setting = 2 Electron Current = 2mA Pc. = 1.15 Torr P„ =8x10-'Torr Pc.2=5xlO-»Torr

Csj

y \AlA}iit\MM^yKM^ 54.885

58.725

82.585

68.405

CsjOi or

j (CsOH), I

70.24S

74.885

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|iSeC

78.245

74.885

77.825 M^SCC

Scan #08259516 VALVE CLOSED Same Settings

54.885

58.725

1.585

88.405

Figure 9. Two TOFMS spectra, with the sampling valve open (top) and closed (bottom), showing that the oxygen-carrying species in the cube was CS2O.

646 loading and storage times of 0.6s while averaging 16 spectra, in effect leaving the valve open for 19s, thereby losing at that P^g about 0.33mg of Cs. The presence of the Cs2 dimer peak in Figure 9 provided a convenient way of estimating the CS2O pressure by assuming that the ionization cross-sections of CS2 and CS2O in the ionizer and the secondary electron yields for the corresponding ions at the electron multiplier were the same: Then the relative signals are proportional to the peak heights. The absolute value of the dimer pressure is then computed from the equilibrium 2Cs(g) 1= Cs2(g). The C§ O pressure thus obtained was 1x10"^ Torr, in good agreement with PQ^ determined by the readsorption method.

5. CONCLUSIONS The converters with cesium-oxide-containing collectors exhibited good performance, with emitter bare workfimctionsaround 5.60eV and barrier indices between 1.9 and 2.0eV, but they were intrinsically two-reservoirs systems. Furthermore migration of cesium oxide away from its initial location had negative effects on the stability of these performance levels. The identification of cesium oxide migration as the cause of the instabilities and irreproducibilities observed in previous studies lead to the development of the DECOR source. DECOR provided a stable way of dispensing cesium and cesium oxide at the electrodes and reaching good performance without special electrode materials, even though DECOR was not a strictly isothermal reservoir. The development of the oxygen pressure measurement method allowed for the first time to measure the vapor pressure of active species inside an operating converter. Finally the sampling of the gas phase generated by a DECOR source, combined with oxygen pressure measurements, provided a detailed picture of thegas phase of a converter with oxygen-enhanced performance.

ACKNOWLEDGMENTS The participation of the following scientists was essential to the success of the program and is gratefiilly acknowledged: K. Greek, L. Hansen, L. Hatch, J. McVey and N. Rasor. hi addition J. Chang contributed his vast skills to the assembly of the various experiments.

REFERENCES 1. D. Lieb et al.. Proceedings of the Thermionic Conversion Specialist Meeting, Eindhoven (Netherlands), 11(1975) 2. G. L. Hatch, Thermionic Technology Program: Thermionic Converter Performance Final Report DOE contract DE-AC03-86SF15954 (1988) 3. Final Report of the High Efficiency Thermionics (HET FV) and Converter Advancement (CAP) Programs, DOE contract DE-ACl 1-93PN38195, order # DE96010173 (1996) 4. K. H. Lau et al., J. Electrochem. Soc, 132,3041 (1985) 5. Ned S. Rasor and J.-L. Desplat, Proc. 24th hitersociety Energy Conversion Engineering Conf, Washington DC, 2847 (1989) 6. K. Greek, Rasor Assoc, internal report, unpublished

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

647

Development of Refractory Metal Oxide Collector Materials and Their Thermionic Converter Performance R.Fukuda, Y.Kasuga and K.Katoh Energy Division, Electrotechnical Laboratory, Umezono, Tsukuba-shi, Ibaraki 305 Japan

ABSTRACT Refactory metal oxides of NbO x ,W0 x ,TaO x, and Silver oxide (AgO x) have been studied for a high performance collector. The metal oxide materials were deposited on metal substrates by RF sputtering in the Ar/O 2 gas mixture, in which the partial pressure of O 2 was deliberately set at the lower values in order to sputter in the stoichiometrically oxygen gas short conditions. Work function of the metal oxides was measured by cesium plasma immersion technique. As results, minimum work function values of each oxide materials were obtained as follows; AgO x =L25eV, NbO x =1.38eV, WO x =1.42eV, TaO X =1.43eV. NbO x and AgO x are considered most promising for a collector. A thermionic converter with a plane parallel type of a polycrystalline W emitter and an AgO X collector, an interelectrode spacing at room temperature 0.1mm, was set up and the power generation experiments were conducted. The maximum power, 3.9W/cm ^ , 0.6V, 6.5A/cm , was obtained under the unignited mode operation at T E = 1 5 8 3 K . The barrier index was V B = 1 . 5 V at T E = 1 5 7 8 K . Based on the experimental results, a new type of a PGM collector was proposed for a micro-gap thermionic converter. 1. INTRODUCTION Development of collector electrode materials with lower work function { — 1.0 eV) is the most effective methods to improve energy conversion efficiency of a thermionic converter. At present refractory metals such as Mo, W, Nb with cesium adsorption are used as a collector with the work function values of about 1.7eV. It is suggested that metal oxide collectors can adsorb cesium more strongly and show lower values of work function. ^ ^ In the study, refractory metal oxides and AgO x were experimentally examined concerning the work function values and high temperature durability. A research thermionic converter of a W emitter and an AgO x collector was fabricated and power generation tested to examine the effectiveness of the AgO x collector. A new type of a FGM collector which integrates a The study was supported by the project "Devek)pment of Energy Conversion Materials through Functionally Graded Structures", funded by Science and Technology Agency, Japanese Government.

648 metal oxide layer and an electrical insulation spacer in a compositionally graded structure was designed and proposed for a micro-gap thermionic converter. 2. EXPERIMENTAL 2.1 Fabrication of metal oxides collectors Metal oxides of W, Mo, Ta, Nb, and Ag were fabricated on the metal substrates by RF sputtering method. (The RF sputtering apparatus is ANELVA,SPF-210H.) The sputtering was done in Ar and O 2 gas mixture. The sputtering conditions in case of the dispersed, partially oxidization condition are shown in Table 1. Table 1 Sputter conditions of dispersed, partial oxidization of metals Target

Gas composition (Ar/0 ,0

Pressure (Torr)

W Mo Nb Ta Ag

90/10 90/10 95/05 95/05 90/10

5.7x10"' 6.2x10"' 6.0x10"' 6.6x10"' 5.6x10"'

Sputter power time (W) (min.) 51 65 75 70 40

25 60 26 17 38

Layer thickness elect, resis. color (Q) ( /^ m) 0.63 1.29 0.39 0.34 2.74

43 32 67 21 0.3

metal ic metal ic metal ic metal ic white gray

SEM photos of AgO x of the dispersed, partially oxidized case are shown in Fig.l AgO X makes a glanular white-gray color layer. In all the case the sputtering was done in the stoichiometrically O 2 short conditions in order to prevent the generation of fully-oxidized layer. The physical characters of fully-oxidized metals are generally higher electrical resistivity, higher thermal radiation emissivity, and higher vapor pressure of sublimation, i.e. volatile. These characters are unfavorable for a thermionic converter

^•2L.l&^fli Ago

(a) Surface Fig.l SEMphoto of AgO

(b) Side layer of sputtered electrode

649 operation. On the other hand,when metal oxide are sputter-coated in the stoichometrically O 2 short conditions, the characters of the layer keep basically the same character to the original metals. Bcause oxygen quantity required to reduce the work function of the electrode is very minute, so the dispersed and partially oxidized metal oxide layers can be expected to have lower work function values for a collector. 2.2 Power generation tests of the W-AgO x thermionic converter AgO X was chosen as a collector because of its lowest work function in our experiments. The research thermionic converter with a poly-W emitter and an AgO x collector was fabricated and power generation tested. The converter, as shown in Fig.2, is a plane-parallel type with a collector guard, 16 0 mm of the emitter diameter, 0.1mm of the interelectrode spacing between the collector and the emitter(C - E), 0.3mm between the collector guard and the emitter (CG - E) at room temperature. Six small plates of YSZ (Yttria stabilized zirconia), 1.5 X 1.5 X 0.1 ^ mm, were set on the emitter for spacers between C - E. The emitter electrode was heated by the electron bombardment of D.C. high voltage, the collector was cooled by a copper radiator. The temperature of the emitter was measured by a pyrometer and W-Re thermocouples, the collector by Chromel-Alumel thermocouples. As shown in Fig.3, output current-voltage characteristics (J-V curve) were measured by applying the 50Hz A.C. voltage to the converter through a transformer. Voltage and current signals were recorded by the Digital Oscilloscope at the phase of 50Hz sinusoidal half wave changing from minus peak to plus peak, i.e. from unignited toward ignited mode opration. radiation fin

AI2O3 spacer collector guard collector YSZ spacer W emitter V(ch-l)

^ - h e a t choke •+0 o4—1

W filament

Fig.2 Cross sectional view of thermionic converter 3. RESULTS AND DISCUSSION

im'

'

^

l(ch-2)

Digital I Oscillol... scope!

I

"H

PLOTTER

To CH-2

Fig.3 J-V curve measurement

3.1 Work function measurement of metal oxide collectors The sample electrodes for a thermionic converter collector, which were fabricated by RF

650 sputtering in Ar and O 2 gas mixture, consist of functionally graded materials (FGM) of the composition gradually changing from the metal to the metal oxide. The sputter-coated layers of the sample electrodes were all sound, no peeling off, and showed strong adhesion. A plasma anode technique was used to measure work function values of sample electrodes in cesium vapor . The results of work function measurements are shown in Fig.4 (a) ^ (d). In the Figures, Tc is the collector sample electrode temperature, T R is the cesium reservoir temperature. The work function values were measured for Tc ranging 723K — 923K, T R ranging 425K -- 513K. The lowest work function values of four kinds of metal oxides were as follows. AgO x :1.25eV, NbO x :1.38eV, WO x :1.42eV, TaO x :1.43eV. The maximum temperature of the AgO x electrode was 536 °C , and the NbO x was 620 °C , in both cases the change of work function values was not observed in the 2.2

—T

1

1

r — I

1

1

1

AgOx

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1

2.2

1

TRO433K A473K 0 493K D 513K

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1

r—

TR 0 433K A456K 0493K D513K

1.8

Mo ]

_ 1.6

;i.6 y

14

14

.^

1.2

1.2 1.0 1.0

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1.6

1.8

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_i

I

I

i_

1.2

1.4

1.6

2.0

1.8

2.2

2.4

TC/TR

TC/TR

(b) N b O X

(a) A g o X 2.2

2.2 TR 0 433K A473K 0493K a513K

2.0 1.8

TR 0 453K A473K 0 493K D513K

2.0 1.8 ;i.6

;i.6

_> 14

14

1.2

1.2 I

1.01—I—i—t1.0 1.2 1.4

1.6

1.8

TC/TR

(c) W O X

2.0

2.2

2.4

1.0 1.0

—I

1.2

1

1

1.4

I

L_

1.6

1.8

TC/TR

(d) TaO X

Fig.4 Work function of metal oxide in Cs vapour

2.0

2.2

2.4

651 temperature range. Then the AgO x was intentionally heated up to 620 °C for 3hrs,and the NbO X up to 745 °C for 3hrs respectively. After heating, 0 c (mi n) of AgO x increased by O.leV, i.e. 1.25 -^ 1.35eV. The 0 c cmi n) of NbO x did not change at all. 3.2 Power generation tests of the W-AgO x thermionc converter Output J-V curves of the thermionic converter with T R below 515K are shown in Fig.5 and Fig.6. The open curcuit voltage of Fig.5 was lower than the one of Fig.6. The reason was guessed to be a very slight short-circuit between the collector and the emitter . Therefore, after that , both J-V curves of the collector - emitter (C-E) and the collector guard - emitter (CG-E) were taken. A noticeable point of Fig.5 and Fig.6 is that the forward saturation current J f o r in the unignited mode was very large , being inconsistent with the theory. (Refer an appendix.) For instance, J for of electron rich ( /3 < 1, /3 means " ion richness ratio"), diffusion conditions (d/ A e-n > 1, A e n means " electron -Cs atom mean free path") under the unignited mode operation, is estimated to be not greater than 0.4A/cm ^ as shown J f o r in Table 2. The value of J f o r of the experiment greatly exceeded 0.4A/cm ^ . J-V curves at T E = 1300 °C are shown in Fig.7, Fig.8, and those at T E = 1350 °C in Fig.9, Fig.10, respectively. The J-V curve of T R =553K in Fig.7 should be noteworthy because the curve has barrier index V B = 1.5V. Although it is not clearly shown in Fig.7, all curves in Fig.7 ignited in 2nd quadrant. They were measured at the phase of half sinusoidal wave of applied voltage from the minus peak toward the plus peak, i.e. from unignited toward ignited mode. Therefore all curves in Fig.7 were in unignited mode opration. Again it is noteworthy J f o r in unignited mode operation is veiy large. In Table 2 are shown the converter parameters calculated on the curve of T R = 5 5 3 K in Fig.7, when the Richardson current J R from the emitter supposed to be J R =10A/cm ^ . The value of ion richness ratio /5 =4.44 X 10 ~ means the electron rich condition and nevertheless the very large forward current was obtained in Fig.7. This is not explainable . Except this point there was not any contradiction in Fig.7. The value of 0 c in Table 2 was calulated by Fig.4 (a), and the 0 c value was far from the minimum value of 0 c because of higher T c . Therefore if the collector temperature could be lowered to the value of minimum 0 c , the output J-V curve could be greatly improved. In fact there are high levels of back emission current from the collector as clearly shown in Fig.7 ^ Fig. 10. The J-V curves in Fig.8 ^ Fig. 10 can be explained as the same way as in Fig.7, and barrier index V B = 1.7V was obtained in these Figures. In Fig. 11 is shown the maximum output power with regard to the emitter temperature. The maximum power of 3.9W/cm " , 0.6 V, 6.5Aycm ' were obtained on the J-V curve of CG-E, at T E = 1 5 8 3 K . 3.3 Future study An image of a FGM collector for a thermionic converter is shown in Fig. 12. The distinctive feature of the FGM collector is that the metal oxide materials with the compositionally graded structure is coated on the metal substrate by sputtering method. At the same time the electric insulation materials with the small columnar shape are placed uniformly dispersed on the collector to hold the gap of 10 jLL m between the emitter and the

652 lOOOi

Em'W.Col'AgOx TE = 1580K Tc = 978K d = 0.3mm (CG)

-0.5

Fig.5

0.0

0.5

1.0

-0.5

1.5

Output Voltage, V (V) J-V curve of T E = 1 5 6 5 K

0.0

0.5

1.0

1.5

Output Voltage, V (V) Fig.6 J-V curve of T E = 1 5 8 0 K Em'W.Col'AgOx TE = 1583K Tc = 1075K d = 0.3mm (CG)

20V, 7y

-0.5

0.0

0.5

1.0

1.5

-0.5

2.0

0.0

0.5

1.0

1.5

Output Voltage, V (V)

Output Voltage, V (V) Fig.8

Fig.7 J-V curve of T E = 1 5 7 8 K

J-V curve of T E =1583K

24 ...

;r20

20 ^ o 16

\

\ (

523K-\^n

' ' ' \

1

:.i2

Em'W.Col'AgOx 1 TE » 1625K Tc = 1098 K d = 0.1mm (C)

2.0V 1.7V

543K \533K \

\

\

L 553K^M| ,

r 533l<\

\

5 16 f 523K\ -3 I •^ 12

r

rTR = 5 1 3 K \

h

1 1 1 1 1 1 1 t 1 1 1 1 1 1 l-t 1 1 1 1 1 1 iM 1 1 rs 1

-0.5

0.0

0.5

1.0

1.5

Output Voltage, V (V) Fig.9 J-V curve of T E = 1 6 2 5 K

2.0

\

\ VB=1.5V

\

^~^~--^>^.

0

-4

\

O^^ ^

t^ n o

Em'W.Col'AgOx 1 TE = 1623 K Tf = l ? | 6 K d » 0.3mm ICG)

M

\

8h

VVB=I.5V

\

^

20V1-7V : T ; •

\

£

^^"•~~-^-oL [ T R = 513K

553K\

1 1 1 1 1 1 1 1 1 1 1 i_i 1 1 1 1 1 1 1 ' ^ ' ' 1 ' ' ' '

-0.5

0.0

0.5

1.0

1.5

Output Voltage, V (V) Fig. 10 J-V curve of T E = 1623K

2.0

653 Table 2 Estimated parameters of J-V curve of Fig.7 T E

Tc

1578K

1074K

01

TR

553K

2.35eV

l.OSTorr 1.40x10 ' N / m

d/ A 4.44.X10

3.5x10

^ mm

1.48eV

2.8

lOA/cm '

9.02x10 " ' A / c m '

Tc/T

0 E- 0 (

1.94

0.87eV

0.37A/cm

0 E is estimated, supposing J R =10A/cm -

D:C

—

o:CG

Q_ 4

0^ Q

\h 3h h

0

D

0 D D

80

o a.

I

collector. A higher efficiency of the thermionic converter operating in unignited mode can be expected by developing the FGM collector.

0 0

L_

L1 1 1 1 1 1400 1500 1600

1

1 . 1 . 1700 1800

1900

Emitter Temp. , TE ( K ) Fig. 11 Maximun output power density of the converter Insulator Oxide graded layer Metal substrate-^

- F G M layer 1^^Metal substrato Fig. 12

Image of FGM collector

4. Conclusion 1) Fabrication of metal oxide collectors The sample electrodes of metal oxides for a thermionic converter collector were fabricated by RF sputtering method in Ar + O 2 gas mixture. Refractory metals of W, Mo, Nb, Ta and also Ag were chosen as target materials of the RF sputtering apparatus. The oxygen partial pressure was intentionally set at stoichiometrically oxygen short levels so as to make the metal partially oxydized and dispersed into the metal matrix of the sputter coated layer. The sample electrodes consist of the functionally graded materials (FGM), the composition changing from the metal to the metal oxide by controlling the oxygen partial pressure of sputtering. The sputter -coated layer showed sound, no peeling off and good adhesion. 2) Work function in cesium vapour The work function values of these sample electrodes were measured by immersing them into the cesium plasma.

654 The lowest work function values 0 c (mi n) of four kinds of metal oxides were as follows. AgOx:1.25eV, NbO x :1.38eV, WO x :1.42eV, TaO x :1.43eV. The AgO x was intentionally heated up to 620 °C for 3hrs, and the NbO x up to 745 °C for 3hrs respectively, in order to obtain data on the high temperature endurance capability. After heating them, 0 c (mi n) of AgO x increased by O.leV, i.e. 1.25 ^ 1.35eV. The 0 c (mi n) of NbO X did not change at all. The materials of AgO x and NbO x made by sputtering were judged to be promising for a thermionic collector. 3) Power generation test The research thermionic converter, with the poly-W emitter, the AgO x collector, interelectrode spacing 0.1mm, was fabricated and power generation tested. The maximun output power was obtained in the unignited mode operation. The maximum power, 3.9W/cm ^ , 0.6V, 6.5A/cm ^ was obtained between the collector guard and the emitter at T E = 1 5 8 3 K . The barrier index V B = 1 . 5 V was obtained between the collector and the emitter at T E =1578K, T C =1074K, d=0.1mm, under the unignited mode operation. 4) Subjects to be solved In the experiment the forward saturation current density J f o r under the electron rich and diffusion conditions of the unignited mode, was 20 times larger than the value expected by the theory. This inexplicable results should be examined by another experiment. Also, what extent the V B could be reduced by T c optimization and what the maximum endurance temperature of the AgO x collector could be, should be examined further. 5) Future study The FGM metal oxide collector which has the integrated spacers and can hold 10 ILL m interelectrode spacing is proposed for the realization of the micro-gap thermionic converter. APPENDIX; The forward saturation current J f o r

under the electron-rich, diffusion conditions of unignited mode

(5 )

Operation is expressed by the equation J for

=

(2 A e - n

n 0 V 1 )/(3d)

(1)

n o = 4(J R J c /v 1 V 2 )

equihbriun density determined by emission currents

V 1 =(8kT E / ;r m)

electron average thermal velocity

V 2 =(8kT E / TT M)

Cs ion average thermal velocity

J R : electron emission current of Richardson eq. J c : ion emission currents of Saha-Langmuir eq.

/I e - n : electron-Cs atom mean free path d : interelectrode spacing

REFERENCES: 1) I.Langmuir, K.H.Kingdom: Phys. Rev. 23, p. 112 (1924). 2) R.Fukuda et al :7th Functionally Graded Materials Symposium(FGM'93), pubhshed by the Fomm of FGM Society, p.211,Tokyo, Nov. 1-2,(1993). 3) R.Fukuda et al : 29th Inter. Energy Conv. Eng. Conf.,Monterey, CA. U.S.A., p. 1041,(1994). 4) R.Fukuda et al :8th Functionally Graded Materials Symposium(FGM'95), pubhshed by the Fomm of FGM Society, p.l61,Tokyo, Oct. 12-13,(1995). 5) P.Stakhanov, et al: Plasma Thermal Emission Energy Conversion, cliapter IE , pl57, machine aided trans, version (1969).

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

655

Thermionic Properties and Thermal StabiHty of Emitter with a (0001) Oriented Rhenium Layer and Graded Structure M. Katoh", R. Fukuda*" and T. Igarashi' "Tokyo Tungsten Co., Ltd., 2 Iwasekoshi-machi, Toyama-shi, Toyama 931, Japan ^lectrotechnical Laboratory, 1-1-4 Umezono, Tsukuba-shi, Ibaraki 305, Japan Rhenium exhibits excellent thermionic electron emission characteristics in a cesium plasma converter. Chemical-vapor-deposited rhenium with a preferred (0001) orientation should have advantages over polycrystalline rhenium with respect to thermionic emission. Rhenium layers with (OOOl)-oriented three-dimensional surfaces have been successfully deposited on molybdenum substrates. The work function of rhenium emitters was estimated to be 4.9 eV from the results of power generation tests. The maximum power of a thermionic energy converter incorporating a rhenium emitter is 2.0 W/cm^ when the temperature of the emitter is 1800 K, that of the molybdenum collector is 990 K and that of the cesium reservoir is 561 K. The composition and surface morphology of the rhenium layer may change as molybdenum from the substrate diffuses into the rhenium layer at high operating temperatures. In order to improve the thermal stabiUty of the rhenium layers at elevated temperatures, tungsten layers were inserted between the rhenium layers and the molybdenum substrates to form graded composition interfaces. The tungsten layers prevented excessive diffusion of molybdenum into the rhenium layers. 1 . Introduction Thermionic energy conversion is a power generation method which can be used to convert thermal energy into electric energy directly. A thermionic energy converter is a nonmechanical device that has high reliability. Research has been carried out on thermionic energy converters for use in space and ocean environments. On thermionic energy conversion, electrons are emitted from a hot electrode and collected using an electrode at a lower temperature. These electrodes are usually separated by a distance of 0.5 mm or less. The electrons condense at the collector and return to the emitter via the electrical load, thereby delivering electrical work. The limitation of the output current density due to the space charge is prevented by enclosing cesium vapor between the electrodes, because of the low ionization energy of cesium. Due to the adsorption of cesium ions, the work function of the electrodes decreases with decreasing electrode temperature (T^ ) and increasing cesium reservoir temperature (Tj^) [1]. The work function of the electrodes is maintained at the Optimum value by controUing the operating conditions.

656 The emitter for the thermionic energy converter is required to have not only excellent thermionic emission characteristics but also hot strength, good workability and bondability with adjacent component materials. Surface-coated emitters are unique in that they possess the thermionic emission characteristics of the overlayer and the mechanical characteristics of the substrate. On the other hand, it is necessary to fabricate an overlayer, which has good thermal stability, for a surface-coated emitter, because the overlayer easily generates damage at the interface between the overlayer and the substrate during heating and cooling due to the thermal stress which results from the difference between the thermal expansion coefficients of the overlayer and the substrate, and the diffusion of substrate constituents to the emitter surface under high temperature operation. Therefore, it is necessary to fabricate a diffusion prevention layer, which suppresses chemical diffusion between the overlayer and the substrate, and results in a composition gradient at the interface, which reduces the thermal stress. Due to the composition gradient at the interface, the formation of a concentration gradient is suppressed, and the thermal stabiHty is improved. In this work, the work function of a (0001) preferred orientation rhenium layer, the thermionic power generation characteristics under conditions in which this layer operates as an emitter and the thermal stability of the interface between the rhenium layer and the substrate are described. 2 . Experimental 2.1. The formation of a rhenium thermionic emission layer and a tungsten diffusion prevention layer Rhenium layers with (0001) preferred orientation were fabricated on molybdenum substrates by chemical vapor deposition (CVD). Molybdenum has light weight, high thermal conductivity and good workability compared to other refractory metals. A schematic drawing of the CVD apparatus is shown in Figure 1. Rhenium base powder was reacted with the chlorine gas at 1070-^ 1170 K, resulting in the formation of ReCls gas. This gas decomposed on the substrate, which was thermally heated to 1400-^ 1520 K, and thus rhenium was deposited. The ReCls generation reaction and the thermal decomposition reaction are represented as 5

Re

— a.

ReCL

(1)

CL

(2)

2

ReCL

Re

heater

It was found that tungsten is an effective diffusion prevention layer, which improves the thermal stability of the rhenium layer by reducing the amount of chemical diffusion between the rhenium layer and the molybdenum substrate. The tungsten layer was deposited by reduction of WFg using hydrogen, at a substrate temperature of 900^1100 K, and a reaction pressure of 1.3 kPa.

induction coil

CZH Figure 1. Schematic diagram of the rhenium CVD apparatus.

657 2 . 2 . Observation of the rhenium layer surface morphology The surface morphology of the rhenium layers was observed using a scanning electron microscope (SEM), and the surface area was measured using an atomic force microscope (AFM). Using X-ray diffraction (XRD), the macroscopic crystal orientation was examined, and the (0001) crystal plane was identified from electron channeling patterns (ECP) [2]. 2 . 3 . Formation and observation of the composition gradient Diffusion was promoted by heating the samples in a vacuum of about 6.5 X 10" Pa at 2000 -^2300 K in order to form a composition gradient between the molybdenum substrate and the rhenium layer. The composition gradient was investigated using an electron probe microanalyzer (EPMA), 2 . 4 . Measurement of the work function and thermionic power generation characteristics A schematic diagram of the power generation test apparatus is shown in Figure 2. The emitter, which was a disc with a diameter of 16 mm and a thickness of 5 mm was joined using ruthenium-molybdenum braze to a support made of tantalum, and was placed facing the molybdenum collector at a distance of 400/^m. The emitter was heated to 1400-^ 1900 K by electron bombardment. Then the collector was cooled by radiative cooling to about 1000 K. By maintaining T^ at 400^570 K, the cesium ^Radiation fin ( C u ) vapor pressure (P^) was kept at 0.35~2.05 Pa. ^ Thermocouple hole An AC power source (100 V, 50 Hz) and X Collector guard ring ( M o ) a load were connected to the power generator, and the current-voltage characteristics were y / O ^ " ^ Spacer (AI2O3) measured using a digital oscilloscope. The Collector ( M o ) effective work functions were evaluated under Emitter (CYD-Re) conditions of P^^ = 0.8~ 13 Pa, for which there Temp, measurement hole is no volume ionization (unignited mode). The power generation characteristics were measured Emitter holder ( T a ) under conditions of P^^ = 1.1 X 10^ --" 2.6 X Heat choke ( T a ) 10^ Pa , i.e., the cesium was in the plasma state Filament guard ( Mo ) (ignited mode), in which an output current Filament ( W ) density more than ten times that for the unignited Figure 2. Schematic diagram of the mode is obtained. thermionic power generation test apparatus. 3 . Results and Discussion 3 . 1 . Evaluation o f the rhenium layer In Figure 3, an SEM image (a) and an AFM image (b) of the rhenium layer surface formed at a substrate temperature of 1520 K are shown. On the surface, (0001) facets, with the characteristic hexagonal structure of the most dense plane, are obseived. The surface area with (0001) orientation was estimated as 42%, and the surface area was increased by 13% due to the surface roughness. 3 . 2 . Fabrication of a gradient structure emitter Figure 4(a) shows the composition of a rhenium/molybdenum interface that was heated for 10 hours at 2300 K. At the rhenium/molybdenum interface, the thermal stabiUty is low, because a rhenium-molybdenum alloy layer over 200/i m thick is formed.

658 (b)

Figure 3. Surface morphology of an as-deposited CVD rhenium layer : (a) SEM image and (b)AFM image. The composition gradients of the rhenium/tungsten and tungsten/molybdenum interfaces formed by heating for 10 hours at 2300 K are shown in Figure 4(b). The samples have an intermediate tungsten layer between the rhenium layer and the molybdenum substrate. The rhenium-tungsten alloy layer is 40// m thick, i.e., about 1/5 of the thickness of the layer formed in the rhenium/molybdenum system. The tungsten-molybdenum alloy layer is 65 // m thick and the rhenium/tungsten/molybdenum structure is stable at high temperatures. Chemical diffusion is suppressed by the composition gradient at the interface at temperatures which are sufficiently higher than the operating temperature, and the thermal stability of the rhenium layer is improved. In Figure 5, the composition dependences of the average thermal expansion coefficients measured at 303-1073 K for the rhenium-tungsten alloy and at 293-1273 K for the tungstenmolybdenum alloy are shown. For the rhenium-tungsten system, the dependence is almost linear, including that for the ^-phase intermetallic compound, formed in the 30-^ 55 mass% tungsten composition range. The results indicate that the stress concentration can be reduced by the composition gradient, even if the (T-phase, which has low ductility, is formed. In the tungsten-molybdenum system, the thermal expansion coefficients change by about 13% in the composition range 0^^ 15 mass% molybdenum. However, for molybdenum compositions greater than 15 mass%, thermal stress concentration is unlikely to occur, since the change in the theimal expansion coefficients is small.

I

I

I

-I

r~

303-1073 K

o *

1

1

293-1273 K

\ \ . \. _l

I

I

I

I

L_

0 20 40 60 80 100 80 60 40 20 0 (Re) (W) (Mo) mass % W distance

Figure 4. Composition gradient at (a) Re/Mo and (b) Re/W/Mo interfaces.

Figure 5. Mean thermal expansion coefficients of W-Mo and Re-W systems.

659 Based on these results, a (0001) oriented rhenium thermionic emission layer ( 100/^ m ), and a tungsten diffusion prevention layer (500jum) were formed on a molybdenum substrate, and emitters with a composition gradient at each interface were produced. 3 . 3 . Evaluation of the power generation characteristics 3 . 3 . 1 . Evaluation of the work function The relationships between ^ g and 7^ / 7]^ for polycrystalline rhenium and CVD rhenium are shown in Figure 6. It has been reported that the work o A functions of (0001) single-crystal rhenium (^ ^Q^^^ ) 3.5 D and polycrystalline rhenium (<> f ^^^y^) in vacuum are 5.5eV [3] and 4.6eV [4], respectively. Using the measured values of ^ ^ ^^^ the relationship between 7^/ TR and ^ ^ ^^^ various bare work function ((/> Q ) , reported by N. S. Rasor, ^ o ^^ ^^^ ^ ^ ^ rhenium emitter was evaluated as 4.9 eV.

1450-1500K 1500-1600K 1600-1700K

^o = 4.9eV

3.3.2. Evaluation of the power generation w characteristics 2.5 When P^ and the interelectrode gap (d) have 4.0 3.5 3.0 TE/TR values greater than a specified value (P^^ • d>20 mil • Torr), the cesium ionizes and forms a plasma, referred Figure 6. Comparison between the to as the ignited mode. In this ignited mode, an output work functions of polycrystalline current density more than ten times that for the and CVD rhenium emitters. unignited mode is obtained. The output currentvoltage characteristics obtained at Tg = 1800 K, T^ = 1000 K and T^ = 497 ~ 561 K are shown in Figure 7. The maximum output increases as the reservoir temperature increases. A maximum output of 2.0 W/cm^ was obtained at 7]^=561 K. The current-voltage characteristics for the above conditions were calculated using the standard computer model "TECMDL" [5] for the ignited mode, developed by J. B. MacVey et al. The current-voltage characteristics were calculated using the measured values of Tg, T^, T^, <j> E^ i> c and d. Although this result does not take into account time degradation, the experimental value for the output exceeded the calculated value. For a cesium plasma thermionic energy converter, there are six important parameters, T^ , ^ c ^R' ^ E ' i^ c ^^^ ^' ^s mentioned above. To improve the output of a thermionic energy converter, it is necessary to determine how these parameters affect the power output. The adjustment of Tg , T^ and T^ is necessary in order to produce appropriate power generation conditions. ^ ^ and j5 ^ of the materials have a significant effect. The interelectrode gap, d, is also an important parameter, since the output depends on it exponentially. E ^^ ^^^o an important factor, since it strongly affects the maximum output. The results calculated using the relationship between the work function and the maximum output of the emitter are shown in Figure 8. In this figure, T^ is optimized for the work function of each emitter. When the work function of the emitter is increased, the maximum output increases rapidly. If the area ratio of rhenium (0001) increases, then (f> ^ increases, and an improvement in the output is expected.

660

0.5 1.0 Output Voltage, V/ V

4.0 4.2 4.4 4.6 4.8 5.0 5.2 Emitter Work Function, ^ / eV

Figure 7. Current-voltage characteristics Figure 8. Calculated variation of maximum of CVD rhenium emitter. (Dotted Hne output with the work function of the emitter. indicates calculated value for T^ = 561 K.) 4 . Conclusion A gradient structure emitter, in which the molybdenum substrate is covered by a (0001) rhenium layer with a three-dimensional surface and an intermediate tungsten layer was fabricated, and the work function and thermionic power generation characteristics were evaluated. The following results were obtained. (1) The thermionic emission area was about 13% greater than that of a flat plane, and a (0001) preferred orientation rhenium layer with a three-dimensional surface and a work function of 4.9 eV in vacuum was obtained. (2) An intermediate tungsten layer was formed, which prevented thermal diffusion of the surface layer into the substrate, and improved the thermal stability of the surface layer. (3) A maximum output of 2.0 W/cm^ was obtained at T^ = 1800 K, T^ = 1000 K, T^ = 567 K a n d d = 400 //m. Acknowledgement This work was carried out as part of the second phase of the FGM R & D program for the development of direct energy conversion technology. This work was supported by the Science and Technology Agency of Japan. References 1. 2. 3. 4.

N. S. Rasor and C. Waner, J. Appl. Phys., 35 (1964) 2589. D. C. Joy, D. E. Newbury, D. L. Davidson, J. Appl. Phys., 53 (1982) R81. V. S. Fomenko, Zmissionne Svojstva Materialov, (1991) 59. Thermo Electron Corporation ed.: Thermionic Research Computation Aids, TE320-1-77, July (1976) 17. 5. J. B. McVey and N. S. Lasor : in Proceedings of the 27th Intersociety Energy Conversion Engineering Conference, SAE/P-92/259, Sandiego, CA, August 3-7, 1992.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

661

Development of Efficient Thermionic Energy Converter T.Kato% K.Morimoto% K.Isogai^ M.Kato\ T.Fukushima^ R.Fukuda^ ^ Mitsui Engineering & Shipbuilding Co.,Ltd., Ichihara-shi, Chiba 290, Japan ^ Tokyo Tungsten Co.,Ltd., Toyama-shi, Toyama 931, Japan ^ National Research Institute for Metals, Tsukuba-shi, Ibaraki 305, Japan ^ Electrotechnical Laboratory, Tsukuba-shi, Ibaraki 305, Japan In the near future, thermionic energy converter will be a more promising device for space power generation. To this end, a solar heated thermionic energy converter has been developed by applying functionally graded materials (FGM) to the electrode. It is majorly composed of a W/Re-FGM emitter, a NbOx collector, and a TiC/Mo-FGM solar receiver. 1. DESIGN OF FGM-BASED CONVERTER 1.1. High Temperature Electrode High temperature electrode, called emitter, is made up of a refractory metal block with three circular parts, as shown in Fig. 1. The surface of upper disc is coated with a layer of TiC/Mo materials by a plasma spray method^^l The TiC surface has emissivity of approximately 0.9 and can be heated up to a temperature range of 1800 to 2000 [K] by solar concentration. Furthermore, such a FGM layer reduces the stress caused by difference of thermal expansion between TiC surface and Mo base. The surface of bottom disc is covered with a layer of ReAV materiales, which is formed through heat treatment after two step CVD coatings^^l The Re surface, having preferred orientations with (0001) plane, shows higher work function than a surface of conventional type. The FGM layer pre-coated with W prevents the diffusion at high temperature. Then both Mo discs are brazed to a middle block of Ta. 1.2. Low Temperature Electrode The structure of collector electrode at lower temperature is similar to the emitter one, but with two circular parts, as shown in Fig. 1. The surface of Mo disc opposite to the emitter is covered with NbOx (Niobium oxide) layer coated by a sputtering method^^l NbOx surface, which shows low work function in Cs vapor, can reduce the energy loss of electron condensing in to collector. This coated Mo disc is brazed to base block of Ta, similar to emitter.

662 1.3. Connections and Other Components Fig. 1 also shows that the emitter and the collector are connected by several units. To reduce the heat conduction loss, Ta tubes called heat-chokes are joined to the both electrodes. Ceramic seal rings are used to insulate the electrodes electrically. Two rings are arranged for Cs vapor tube inlet and electrical leads. Each unit is connected by a electron beam welding or a vacuum brazing method. Insulator pins, called ceramic spacers, are attachet on the collector to keep a constant interelectrode gap. Small crown parts are assembled inside the converter in order to block the radiative heat. 1.4. Converter for Laboratory Tests We have to prepare another type of converter for laboratory testing, shown in Fig. l(a dotted line). Such a converter has a long emitter block with three holes for thermocouples, instead of Mo disc. It can be heated by electron bombardment. The surface temperature of emitter, TE, is estimated from the temperature distribution of the emitter block.

1®J ?ctron-bombardment heat i ng

radiation 18 00

TiC/Mo graded Re/W graded emitter ceramic spacer radiation seal NbOx col lector ceramic seal

m

M

-^ electrical lead Cs vapor inlet Fig. 1 Cross-sectional view of eflBcient thermionic energy converter.

663

radiation seal vacuum chamber

ai r cool ing fin

fi lament

thermionic converter Cs vapor pipe

Fig. 2 Schematic of experimental apparatus.

2. EXPERIMENTAL APPARATUS An experimental measuring apparatus as shown in Fig. 2 consists of an electron bombardment heating system, an air cooling system, a vacuum system, and an electrical measuring system. By adding assistant alternating voltage to the electrodes, the corresponding output current density is measured, differently from a passive method. The way to estimate the surface temperature of collector, Tc, is the same as that of the emitter. The collector codling system has three holes to insert thermocouples along conductive heat flow. The pressure of Cs vapor filled in the interelectrode gap is controlled by the temperature of liquid Cs reservoir, TR. 3. EXPERIMENTAL RESULT AND DISCUSSION 3.1. J-V C h a r a c t e r i s t i c s The electron current density, JR, emitted from the electrode is given by a relationship^'*^

664

J R = AT^-exp

kTJ

[A/cm2

(1)

where A = 1 2 0 [A/cm^-K^] T = electrode temperature [K] (f) = electrode work function [eV] k = 8.62 X 10"^ [eV/K] : Boltzmann's constant Eq. (1) is called Richardson-Dushmann equation. JR never indicate output current density in itself, although it helps account for the converter performance. Large JR from Eq. (1) suggests that the converter generates a great deal of output power. Since the work function of the electrodes immersed in Cs vapor can be expressed as a ratio of TR to T, following three temperatures, TE, TC, TR, are important parameters, which will be presented later. A typical curve of current-voltage (J-V) characteristics and maximum output power (Pmax) are shown in Fig. 3. For three given temperatures, increasing the voltage up to 0.8 [V] results in a gradient decrease of the output current. In this region the converter is operated in an ignited mode, and a large number of ions a r e g e n e r a t e d by i n e l a s t i c collisions in the interelectrode gap, this leading to a plasma state. Pmax is usually obtained at large voltage of the ignited mode. Then there is a rapid current drop because of transition of the operating mode. In a voltage range from 0.9 to 1.0 [V], the converter is operated in the unignited mode without plasma. In this region the output current density is limited by electron space charge. J-V characteristics helps define the interelectrode phenomena such as electron and ion emission from electrodes, sheath height at the electrode surface, voltage loss in the plasma, and so on. But it Fig. 3 Typical J-V characteristics of the is difficult to discuss those converter with an electrode surface area phenomena in detail. In following of 3 [cm^], # depicts a Pmax point for, s e c t i o n , we d i s c u s s the V= 0.7519 [V] correlation between Pmax and J=10.4159 [A] some temperatures, apart from P= 7.9069 [W/cm^]. J-V characteristics.

665 3.2. Tc Effect on Pmax In the experimental converter of Fig. 1, collector surface covered by a NbOx layer serves as a reservoir of electronegative gas (oxygen-contained molecule), which reduces the emitter work function. Therefore Tc has a large effect on the converter performance. The dependence of Pmax on Tc is shown in Fig. 4. This indicates that an optimum value of Tc is found to be between 925 and 950 [K]. As Tc increases from 750 to 950 [K], Pmax increases monotonously because emitter work function is reduced by oxygen from the collector. But further increasing of Tc results in a reduction of Pmax because, broadly speaking, the electron emission from the collector can not be negligible. 3.3. TR Effect on Pmax The dependence of Pmax on TR is shown in Fig. 5. It is seen from this figure that higher TR seems to achieve larger Pmax. If TR was in excess of 610 [K], Pmax would reach a peak at a given TR. But a great deal of current density is not suitable for practical use. Table 1 shows experimental values of electrode voltage (Vmax), current density (Jmax) corresponding to Pmax, and controlled temperatures, which present the plotted data of Fig. 5. It is found that Vmax shows a large decrease with an increase of Jmax when TR rises ten degrees from 593 [K]. When Jmax is limited below 10 [A/cm^] to avoid a large voltage drop across the actual leads, Pmax has a value of about 8 [W/cm^]. 4, CONCLUSION To date an efficient thermionic energy converter has been developed and operated at various test conditions. It is concluded that the converter design, including a technical method of coating, welding, and brazing, has been successfully done because there observed no trouble through operating at high temperature. Typical output power density is 7.9 [W/cm^] at the emitter temperature of about 1800 [K]. It is pointed out that higher TE of 2000 [K], may be desirableto produce large Pmax. REFERENCES 1. T. Fukushima, S. Kuroda, and S. Kitahara, Proc. 8th Symp. on FGM, pp.167-170. The FGM Forum (1995). (in Japanese) 2. R. Igarashi, et al., ibid., pp.155-160. (in Japanese) 3. R. Fukuda, Y. Kasuga, and K. Kato, ibid., pp. 161-166. (in Japanese) 4. G. N. Hatsopoulos and E. P. Gyftopoulos, "Thermionic Energy Conversion, vol.2," The MIT Press (1973).

666 Table 1 Testdata-operating temperatures, output voltages and current densities.

h

1 Exp No

SI

h

Vmax [V]

[fe

Pmax [W/cn?]

3.0768 3.1142 3.2003 3.7666 4.8332 7.1485 15.0007 3.9472 4.5133 5.5754 7.0158 9.3982 I 11.3566 15.5721

1.6667 1.6437 1.6667 2.0140 2.6853 3.7730 6.3543

1 2 3 4 5 6 7

1851.2 1846.2 1849.8 1847.8 1846.2 1840.8 1818.6

1061.4 1056.7 1056.1 1062.2 1073.9 1088.5 1122.6

543.3 553.3 562.8 573.3 583.5 593.9 603.4

0.5417 0.5278 0.5208 0.5347 0.5556 0.5278 0.4236

8 9 10 11 12 13 14

1844.9 1844.5 1840.6 1835.2 1831.4 1826.0 1817.4

970.8 969.9 977.6 986.4 996.9 1000.7 1015.0

543.4 552.5 563.4 572.9 583.1 593.2 604.0

0.7917 0.7873 0.7847 0.7792 0.7129 0.7062 0.6197

TE ! Emitter temperature Tc ! Collector temperature TR ! Reservoir temperature 3. 5

1

•

.

1

,

•

Vmax ! Voltage at the maximum output power Jmax ! Current at the maximum output power Pmax ! Maximum output power 1

1

o •2.5

-

T A

1. 5

0

J

A

o j T ^ oH

750

I

1

5

1

Fig. 4 Dependence of output power on collector temperature. O: T E = 1 6 2 0 ± 1 0 , TR=573±1[K] • : TE=1620±10, TR=533±1[K] A: T E = 1 6 2 0 ± 1 0 , T R = 5 2 3 ± 1 [ K ]

-

12

~ -

S 4 _

11

A

• 10

J

"6

•9

-j

0,5

°?

J

°4

S 1

540

0, 7 J

•_

a>

1

800 850 900 950 1000 C o l l e c t o r T e m p e r a t u r e [K]

14

J

-•e f^ 2 _ =3 °i

l—

^T"

• „

tr>

=3

L.

' •13

^>^ 1

•^ 1 J

r •

e ^ " y - 8

*> 3

A

1'^,"

1

^' ^

V.

T

1 0.5

o

]

AA ^ -^ A •

r A

-

^

o

0<^ O O

3.1250 3.5533 4.3750 5.4667 6.7000 8.0200 9.6500

-J

J

J

550 560 570 580 590 600 610 Cs R e s e r v o i r T e m p e r a t u r e [K]

Fig. 5 Dependence of output power on reservoir temperature, protting data of Table 1. • : TE=1835±15, TC=1090±30[K] O : T E = 1 8 3 0 ± 1 5 , TC= 9 9 5 ± 2 5 [ K ]

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

667

Radiation dose reduction by graded structures in the heat source of a ^^Sr radioisotope battery A. Ohashi, K. Ueki and T. Senda Ship Research Institute, MOT, 6-38-1, Shinkawa, Mitaka, Tokyo 181, Japan

We studied possibility of absorbed dose reduction for a cylindrical hollow heat source model of a ^^'Sr radioisotope battery by adding graded structure. The simulated result for addition of three mixed layers conserving total materials showed 22 % reduction to the dose of non-FGM structure. For compensation of a dose limitation of unmanned mission in SP-100 project, tungsten shield was needed about 5.7 cm thickness with the graded structure which indicated the minimum dose.

1. INTRODUCTION A radioisotope battery is one of the choice for energy source of meteorological observation and development of undersea and space[l]. We have considered a strontium-90 (half-life: 28.8y) heat-source model of a radioisotope battery and improved it in two aspects—radiation dose reduction and improvement of thermal conductivity—adding graded structure to the model[2]. The present study reports the dose reduction of bremsstrahlung photons from yS -ray of ^*'Sr and its daughter nuclide yttrium-90. The calculation was carried out by a continuous energy Monte Carlo code, MCNP 4A[3].

2. MODEL FOR ANALYSIS Figure 1 shows the two models which were calculated. The base structure was considered with another research group for improvement of the thermal conductivity, while three layered structure was adopted on the view point of shielding. The base structure was assumed to be a concentric cylinder of 60 cm height which was formed void, strontium titanate C^*'SrTi03, p :5.12g/cm3) and boron nitride (BN, p :2.26g/cm3) from the inside. The weight of these two materials was made to be identical each other. Three layered structure includes a new layer (mixed layer) between ^^^SrTiOa and BN of the base structure conserving each total weight of the two materials.

668 In the mixture, average atomic number decreases with the ratio of BN. Since radiative stopping power of /3 -ray is proportional to the square of average atomic number, the probability of production of bremsstrahlung photons can be reduced in the new layer, and also the spectrum of the photons becomes softer than that of the ^^^SrTiOa layer. These are factors for radiation dose reduction. On the other hand, there is another factor for radiation dose increase since generated position of bremsstrahlung photons approaches the detection position. Consequently, radiation dose on the surface of the heat source can be minimized by arranging those two factors. Four layered structure includes another new layer at the inner boundary or outer boundary of the three layered structure which showed the minimum absorbed dose. In the case of the inner boundary, the new mixed layer was made from 100% ^oSrTiOs and the mixture of the three layered structure. The new inner radius was changed in 100% ^"SrTiOa and the new outer radius in the mixture. The effect of addition of the layer was tested depending of the two radius, and also at the outer boundary between the mixture of the three layered structure and 100% BN. The structure which obtained the better dose reduction in the two tests was adopted as the minimized four layered structure. For five layered structure, three tests were performed to the three boundary which located between 100% ^oSrTiOa and 100% BN of the minimized four layered structure. The structure which derived the best dose reduction was chosen as the minimized five layered structure. In the calculation, the inner and outer radiuses of mixed layer were changed by 0.5 cm step to simulate the effect of dose reduction. The density of mixture was calculated theoretically from the inner and outer radiuses. Although strontium has many isotopes, we treated ^"Sr only. From this condition the model included 2.9 x lO^^ Bq of ^'^Sr at the outset and its thermal output became 52 kW. 90

SrTi

BN NEW LAYER (MIXTURE)

NQUTER BOUNDARY INNER BOUNDARY

BASE STRUCTERE Figure 1. Calculation model.

THREE LAYERED STRUCTURE

669 3. CONDITION FOR MONTE CARLO SIMULATION We performed two kinds of calculation using MCNP 4A code. One was simplified calculation. Its source was set to 2.245 MeV as monoenergy and cutoff energy for electron set to 2.24 MeV. The calculation could save time and search for the tendency of the dose variation to the inner and outer radiuses of a mixed layer. On the other hand, normal calculation was performed to the geometry which showed the calculated value within one standard deviation from the minimum value of simplified calculation. The source spectrum for the calculation was compounded from those of ^^^Sr and ^'T which were cited from Ref. 4 as a function of 0.15 MeV step. Source bias was employed in the calculation. Cutoff energy for electron was set to 0.2 MeV testing the effect of the energy to the result [5]. Both calculations were carried out to obtain average absorbed dose on the cylindrical side of the model in Figure 1 by means of surface crossing estimator in MCNP 4A. As the setting which is common to two, cell importance for electron and weight window for photon were used for variance reduction. The response function for the absorbed dose was quoted from Ref. 6. A card BBREM in MCNP was adopted to bias the probability of production of bremsstrahlung photons at higher energies. The minimum absorbed dose was searched up to five layered structure by means of the same procedure. With the structure of the minimum value, thickness of tungsten shield was calculated to compensate the dose criteria of SP-100 project for unmanned mission(5 x 10^ rad for 7 years)[7]. Workstations of Hewlett Packard Company (9000/735,755) were used for calculations. It took one hour per case for simplified calculation, five hours per case for normal calculation and 2 days for the calculation with tungsten shield.

4. RESULTS AND CONSIDERATION Figure 2 shows the results of simplified calculation for three layered structure. The axis of abscissa indicates the inner radius of mixed layer for the three layered structure. Eight graphs correspond to the radius of the outer boundary in the figure. Each vertical line of the gi'aphs reveals relative intensity of absorbed dose in linear scale and gives the same range each other. Error bar expresses standard deviation from the Monte Carlo simulation. The lower values of absorbed dose are obtained in the range from 11.5 to 12 cm of the inner radius and from 15.5 to 18.0 cm of the outer radius. It is considered that the minimum absorbed dose locates within these range. In the same figure, filled square indicates the position where showed the values within one standard deviation from the minimum value of the simplified calculations. These geometry was recalculated by normal calculations with the 0.2 MeV cutoff energy. The minimum value of the normal calculation was shown in it by filled

670 circle (inner radius: 12.0 cm, outer radius: 17.5 cm). The minimum value resulted in 17 ± 1 % reduction to that of the base structure. Using the same procedure, the minimum absorbed dose was searched up to five layered structure. Table 1 shows the results of the calculation. For four layered structure, the effect of addition of a new mixed layer was tested to two boundary of the three layered structure which showed the minimum absorbed dose. From the simulations, the better reduction in the two tests was obtained when a new layer was added to the inner boundary. The minimum value showed 19 % reduction to that of the base structure. Also for five layered structure, the tests were performed to three boundary of the minimized four layered structure. From the results, the improvement of the most inner

OUTER BOUNDARY 100% SrTiOa

ri^ /MIXTURE

m

M 100% BN

INNER BOUNDARY

10

11

12

INNER BOUNDARY OF MIXTURE (cm) • •

WITHIN 1 S.D. FROM MINIMUM VALUE OF SIMPLIFIED CALCULATION MINMUM VALUE BY NORMAL CALCULATIONS ERROR BAR: STANDARD DEVIATION

Figure 2. Results of simplified calculation for three layered structure.

671 Table 1 Results of calculation. Boundary Weight % Weight Density Structure radius(cm) ofsoSrTiOa %ofBN (g/cm3) 3.14 50. 50. Homogeneous 6 . 4 - - 2 L 0 6.4-12.785 5.12 100. 0. Base 12.785-21.0 2.26 0. 100. 6.4 - 12.0 5.12 100. 0. 3 layers 12.0 - 17.5 2.60 76.4 23.6 17.5-21.0 2.26 0. 100. 6 . 4 11.5 100. 5.12 0. 4 layers 11.5 - 14.5 43.4 56.6 2.98 14.5 - 17.5 23.6 76.4 2.60 17.5-21.0 100. 0. 2.26 6 . 4 9.5 100. 5.12 0. 5 layers 9.5 - 12.0 92.1 7.9 4.65 12.0 - 14.5 43.4 56.6 2.98 14.5 - 17.5 23.6 76.4 2.60 17.5 - 2 1 . 0 0. 100. 2.26

Dose Standard ratio deviation ±0.03 2.38 1.00

±0.008

0.83

±0.01

0.81

±0.01

0.78

±0.01

boundary showed the minimum dose in the three tests. The minimum absorbed dose resulted in 22 % reduction to that of the base structure. The two-layer addition caused 5 % reduction from the absorbed dose of three layered structure. As these data include about 1 % error from the precision of Monte Carlo simulation, it is difficult to show clearly the effect of more multi-layered structure. The result of homogeneous structure is also shown in the same table, which is one layer structure of 50 %(weight) ^"SrTiOa and 50 %(weight) BN. It indicates a factor of 2.38 increase to the dose of base structure. A conceptional design of a radioisotope battery in Ref. 1 adopted a homogeneous mixture of 20 %(volume) '^^>SrF2 and 80% (volume) BN in the heat source. Hence it is possible to reduce the radiation dose of the design by the application of gi*aded structure. Since the radiuses of mixed layers were changed by 0.5 cm step, it is possible that the ratio would become lower if the step became smaller. But it is also very difficult to indicate the difference between the minimum value and the new result because the difference is within one standard deviation of Monte Carlo simulation. Figure 3 shows the results of calculation for thickness of tungsten shield with the five layered structure which showed the minimum result. Horizontal dotted line indicates a converted value of dose limitation of unmanned mission in SP100 project. Tungsten shield is needed about 5.7 cm to reduce the dose to the limitation.

672

2 10

-13 f

'

T-

~i

'

r

14

i 10 -'' b'a

10

-15

10

-16

10

-17

10

-18

O Q

Q PQ CQ P<

O

c/3

D.OSE UMITATIQN LEVEL.

0

1 2 3 4 5 6 THICKNESS OF TUNGSTEN SHIELD (cm)

Figure 3. Results of calculation for thickness of tungsten shield with the five layered structure in Table 1. The surface of the structure corresponds to 0 cm of the axis of abscissa 5. CONCLUSION The possibility of absorbed dose reduction for the heat source model of a ^^Sr radioisotope battery was studied by adding graded structure to the model using MCNP 4A code. The results showed 17, 19 and 22 % reduction to the dose of the base structure for three, four and five layers, respectively. A calculation was performed to evaluate a thickness of tungsten shield for compensation of a dose limitation of unmanned mission in SP-100 project. From the result, tungsten shield was needed about 5.7 cm thickness with the five layered structure which indicated the minimum dose. In the future, the model have to be integrated with the results of improvement of thermal conductivity. After that the improvement of a conceptional design in Ref. 1 should be gone ahead. REFERENCES 1. N. Ito, M. Fujiwara, M. Watabe, M. Toyoda and Y. Kamishima, MITSUBISHI JUKO GIHO, 29, 531-536(1992), in Japanese. 2. A. Ohashi, K. Ueki and T. Senda, FGM'94, 143-146(1994). 3. J. F. Briesmeister, Ed., LA-12625(1993). 4. A. B. Brodsky, Ed., CRC handbook of radiation measurement and protection, CRC Press, 1978. 5. A. Ohashi, K. Ueki and T. Senda, Proceedings of the Sixth EGS4 Users' Meeting in Japan, KEK, Tsukuba, Japan, 1996. 6. ICRP Publication 21, Pergamon Press, 1971. 7. V. Keshishian, L. Gay and R. D. Meyer, Intersoc. Energy Convers. Eng. Conf. 23rd[3], 219-223(1988).

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

673

Output increase of thermionic energy converter due to the illumination of Xenon short arc lamp Y.Shibahara and M.Kando Department of Electrical and Electronic Engineering, Shizuoka University, Johoku 3-5-1, Hamamatsu 432,Japan

ABSTRACT The influence of illumination of Xenon short arc lamp radiation on the thermionic energy converter output has been investigated. Here, the thermionic energy converter used in the experiment is designed to keep the space between both electrodes enough widely for the efficient illumination. Through the open end of the cylindrical collector, the light from the Xenon lamp with the nominal output power of 3kW at the maximum illuminates the emitter after being focused by lenses. The cesium gas temperature Tg can be controlled up to 450 K. The weak magnetic field with cusp geometry is applied to guide the electrons to the collector. The main results obtained by the experiment are summarized as follows: 1) The output current operated by the unignited mode increases by 10^^50 times due to the illumination for the emitter temperature TE less than 900 K and Tg from 300 to 450 K. But the difference of the output currents with and without illumination decreases as TE increases. 2) The ignited mode appears when Tg is higher than 400 K and the collector is positively biased with respect to the emitter. But, the bias collector voltage for the ignited mode operation largely decreases by the illumination. Especially, the ignited mode operation can be accomplished at TE around 1300 K by the illumination, even if the collector voltage is lower than the emitter voltage.

1 .INTRODUCTION As is well known, the thermionic electron emission from the tungsten hot plate in the cesium gas has an excellent nature that the thermionic electron current from the tungsten hot plate becomes large at the lower hot plate temperature than 1000 K, although it depends on the hot plate temperature and cesium gas pressure. This is because the hot plate is covered by cesium thin layer, which reduces the work function of the hot plate to around 2.0 eV at low hot plate temperature. This nature seems to be very attractive from the view point of the development of thermionic energy converter because it enables to the operation of thermionic energy converter at the low emitter temperature. However, the operation of low emitter temperature causes the large negative space potential near the emitter since the number of cesium ion near the emitter produced by

674 the contact ionization at the surface of the hot plate is insufficient to neutralize the space charge, compared with the number of the electron emitted from the hot plate. Therefore, the large number of the positive ions should be supplied in the space between emitter and collector in order for the whole electrons emitted from the emitter to be able to reach the collector, which leads to an extraction of enough large output current even in the low emitter temperature operation. The photoionization of cesium atom or molecule due to the sunbeam is adopted as the way of the positive charge supply in the present work. The structure and arrangement of electrodes in the thermionic energy converter used here have been improved in order for the space between both electrodes to be well illuminated by the sunbeam. By the way, the effect of the electrodes gap length on the space potential distribution has been investigated by Otto[l]. According to his numenical analysis, the space potential distribution is closely related to the space charge neutrality a which is defined by ni/ue^ where rii and rie are the ion and electron density around the emitter. The space potential distribution objectionable to the electric power generation is created in the case of a < 1, because the minimum space potential becomes lower than the emitter potential. However, the improper effect of the minimum space potential on the characteristics of the thermionic energy converter is remarkably improved according as a becomes larger than unity. In the present paper, the increase of output current of thermionic energy converter due to the Xenon short arc lamp illumination will be reported and discussed.

2.EXPERIMENTAL APPARATUS 2.1. Thermionic The thermionic the space between As shown in Fig.l,

energy converter energy converter has been designed to satisfy with the condition that both electrodes is effectively illuminated by the Xenon short arc lamp. the cyfindrical ring with 38 mm in inner diameter and 20 mm in height.

Heat shielding plate 5 [cm]

-"0 Emitter ( T u n g s t e n spiral) L 0

Collector ( C y l i n d r i c a l ring) JL 5

J10

Figure. 1 Schematic diagram of thermionic energy converter.

Heater

15 [cm]

675 made of SUS 304, is set near the end of the pyrex glass tube with an inner diameter of 50 nmi and 150 mm long and used as the collector. The end glass near the collector is optically flat and works as the window. Emitter ^^^ emitter is the tungsten spiral with a di(Tungsten spiral) amter of 18 mm. It is made by winding the tungsten wire with a diamter of 0.5 mm by 8 turns and is set at one end of collector ring. The external magnetic field with the cusp Collector (Cylindrical ring) configuration is applied to the thermionic energy converter as shown in Fig.2. It should be noted that the lines of magnetic field intersect the surfaces of the emitter and collector so that the diffusion loss of electron position [cm] emitted from the emitter and created by photoionization of cesium atoms or molecules Figure 2. The external magnetic field will be suppressed. with cusp configuration 2.2. T h e operation of thermionic energy converter The emitter is directly heated up to 1500 K by a half rectified current of 60 Hz. The whole measurements of the characteristics are carried out during the period without the heating current to avoid the disturbance caused by the heating voltage which appears at both ends of the emitter spiral. The radiation from the Xenon short arc lamp with the nominal output power of 3 kW is focused by several optical lenses to illuminate the emitter through the glass window and the another open end of the collector ring. Therefore, the emitter, plasma and cesium gas in the space between the window and emitter can be sufficiently illuminated. The cesium gas pressure in the thermionic energy converter can be controlled by changing the temperature of glass wall so that the termionic energy converter is set in the electric furnace to heat up around 450 K at the maximum. 2.3. M e a s u r e m e n t of t h e characteristics of t h e thermionic energy converter The measurement of the output characteristics is carried out during the period that the heating current is switched off. The electronic circuits for the measurement are consisted of the comparator circuit, delay pulse generator, sampling and hold circuit and triangular wave generator, as shown in Fig.3. The radiation from Xenon short arc lamp is so intense that the emitter and the wall of converter will be heated and the operating condition will be significantly modified when the illumination continues for long time. To investigate the photo-effect on the characteristics of the thermionic energy converter, the illumination time should be short enough for the heat-effect not to be actualized. Therefore, the whole experimental data have been taken within 1 second just after the illumination The external magnetic field with cusp configuration is produced by two solenoid coils. The magnetic field intensity on the central surface of the emitter is 20 Gauss, while it is

676

around 180 Gauss on the middle plane of the collector for the coil current of 5 A. Electric furnace

H^

"pu fJl'' U Samp 1 i ng and Hold c i r c u i t

generator

X axis

7^

p-

Y axis

Ext. trigger Figure 3. The schematic diagram of the experimental apparatus. 3. E X P E R I M E N T A L R E S U L T S The effect of the illumination on the output current of the thermionic energy converter are examined as a function of the emitter temperature and cesium gas temperature. In general, the thermionic energy converter is operated by the unignited mode at the lower cesium gas temperature. Figures 4(a) to 4(d) show the output characteristics of the thermionic energy converter operated by unignited mode. It is clearly indicated that the output current remarkably increases by the illumination of xenon short arc lamp. Figures 5(a) and 5(b) summarize the measured short circuit current of the thermionic energy converter as a function of the emitter temperature, together with the electron saturation current given by Richardson-Dushmann equation and the space charge neutrahty a calculated under the consideration of the emitter and cesium gas temperatures. It is clearly found that the remarkable increase of short circuit current due to the illumination is obtained at the lower emitter temperature. The characteristics operated by the ignited mode are observed at the higher emitter and cesium gas temperature, as shown in Fig.6(a) to 6(d). The output voltage needed to induce the ignited mode operation depends on the illumination, in addition to the emitter and cesium gas temperatures. It should be noted that the ignited mode appears even in the positive output voltage at the emitter temperature of 1280 K and cesium gas temperature of 443 K as shown in Fig.7. It is found that the output current becomes sensitive to the external magnetic field under this condition.

677 ^oi^^^cnf]

(a)

(b)

^^ g^g temp.Tg:350[K]

'o[^,^'^']

Emitter temp.TE:1070[K]

0.3

0.3

O.2I

0.21 With irradiation

Without irradiation

no!

0.1

-3

-2

^1

^Vo[V]

6|

-3

-2

-1

-i-^^5:^v°[^

0| -0.1

-0.1

lo[mA/cm^] 0.4r

lo[mA/cm^] 0.4r

(d)

Cs gas temp.Tg:350[K] Emitter temp.TE:1200lK]

0.3

Without irradiation

With irradiation

"o^ 0.1 Vo[V]

-3

-2

-1

-^Vo[V] 2^^^S3

0| -0.1

With magnetic field(Coil currenf.1,3,5[A])

- Without magnetic field,

Figure 4. The output characteristics of the thermionic energy converter operated by the unignited mode. The experimental conditions are written in the figure. 10- r

rio^10-

r

2 /^ / y L -^-^

-""t--^

^ ^ ^ ' ^ ^ 1

.^10- F;^^" o

**

E10-

r

irradiation X

•

0

A

X

A

0

/. <^^'<^3^'miv

/s^'/^/^ / ' lii^'; /o 2 ^ ^

clQ- r •^ dio-

\* \ •

/

^ » •

// /

•

o10-^

(fi-—^

/liiK

•a

0

|_10^°

Cs gas temp. T,:350[K]

-rio-

lioi

10°

Q)

•o clO 10-^° t5

•\

03 Q.

\i)__^

. |-

^Q-20

500 1000 1500 Emitter temp.TE[K] (i)T.=350[K] magnetic field X (ii)t=400[K] X 0 (iii)Tg=TJK](Wall temp.) 0

(a)

Cs gas temp, Tg:400[K]

Q)

0IO

.lio-^^> o XL

^loi irradiation

•

0

A

X

A

0

0

500 1000 1500 Emitter temp.TE[K] magnetic field (i)T,=400[K] (ii)t=450[K] X 0 (iii)Tg=TJK](Wall temp.) 0

(b)

Figure 5. The relation between the emitter temperature and the short circuit current with and without illumination and external magnetic field.

678 lo[A/cm^ 2r

Cs gas temp.Tg:400[K]

lo[A/cm^ 2\

(u)

Emitter temp.TE:1070[K]

Without irradiation

-3

2~^

-2 " -

loIA/cm"] 3

(c)

3^°^

With irradiation

-3 • -2 • -V"Ol

2

3Vo[V]

lo[A/cm^ 3r

(d)

Cs gas temp.Tg:443[K]

1

Emitter tempTg: 1280[K]

Without irradiation

T

0'

^VoM

1

• Without magnetic field,

With irradiation

-3

-2

-1

0"

;Vo[V]

-With magnetic field(Coii current:1,3,5[A])

Figure 6. The output characteristics of the themionic energy converter operated by the ignited mode. The experimental conditions are written in the figure. lo[A/cm2]

4. D I S C U S S I O N

With irradiation

0.3r

Cs gas temp.Tg:443[K] Emitter temp.TE:1280[K] 0.2

\ Coil current: '^\5[A] 0.1 .3[A]

Im. 0.1

0.2

yoM

0.3

Figure 7. The output characteristics of the thermoionic energy converter. The data shown in Fig.6(d) is enlarged.

The increase of output current in the unignited mode operation due to the illumination can be explained by the impovement of the space charge neutrahty a. As mentioned above, a is usually smaller than unity at the low emitter temperature so that most electrons emitted from the emitter cannot reach the collector. However, the illumination creates so many electrons and ions by photoionization that a approaches to unity [2]. Though the remarkable effect of illumination on the output current is observed at the operation of the low emitter temperature, the increase of a does not contribute to the increase of the output current at the higher emitter tem-

679 perature where the number of the positive ion produced by the contact ionization becomes larger than that of the electron. These features are clearly shown in Fig.5. As for the ignited mode operation, the increased electrons and ions in the space between both electrodes contibute to induce the breakdown because the probabiHty of collisional ionization is proportional to the electron and cesium atom density. Therefore, the larger output current will be obtained when the ignited mode operation will take place by increasing the cesium gas pressure and emitter temperature and by illuminating the intense light on the thermionic energy converter. REFERENCES l.W.Otto,Z.Naturf,22a(1967)1057. 2.M.Kando,H.Furukawa,M.Ichikawa and S.Yokoi,Proc.of 29th IECEC,Vol.2(1994)1067.

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681

Hybrid Mode Concept of a Thermionic Converter with a FGM Structured Collector Mitsuo Iwase

Yoshihiko Hirai

Department of Electrical Engineering, Tokai University Kitakaname 1117, Hiratsuka-shi, Kanagawa 259-12, Japan

ABSTRACT A new hybrid mode concept of the thermionic energy converter was proposed. A number of papers deaUng with the hybrid mode have been published, but they mainly deal with the mechanism of the arc at the emitter. On the contrary, this paper proposes that the ball of fire at the collector can result in an hybrid mode. To that effect a FGM structured collector is considered to generate the new hybrid mode. In this study, the FGM structured collector was produced by means of the ion beam sputtering for mosaical target. The material distribution deposited on the substrate was detected by the X-ray photoelectron spectroscopy(XPS). Results of the XPS were recognized to be FGM structured collector.

INTRODUCTION 1989, a new idea was proposed by the study group resident in Sendai: to alter gradually the internal components or the microscopic organization of the materials with the objective of achieving arbitrary material properties. This is a epochal concept of material design named Functionally Graded Material (FGM).l) At the start, the techniques based on the concept of the FGM for realization of compositoin and structure gradients have succeeded in development of the wall materials for the space-plane that is needed to achieve relaxation of thermal stress. Subsequently to this success, FGM technology has been applied to many fields, because FGM technology, in its original sense, means providing a gradient in materials properties such.as electronic, chemical, optical, nuclear, and biological applications. 2) In this study, we have as its objective the improvement of the efficiency of thermionic energy converters. The electrical output characteristic of the converter shows two mode of operation, arc mode and unarc mode. Each mode have some merits and demerits. Therefore, hybrid mode is an ideal mode which makes use of the each mode merits. So, we applied the FGM technology to the development of the electrode materials needed for the hybrid mode. Usually, the hybrid mode has been realized by using the arc generation on the "emitter", such as a grooved type emitter. However, FGM technology makes it possible to realize the hybrid mode by using the ball of fire generated at the collector. We report on the new hybrid mode and on a FGM collector which could be used to generate it.

682

THEORY In order to discuss the I-V characteristic of thermionic converter, a schematic of such a characteristic curve(pc? < 20, /? < 1) is shown in Fig.l. The positions on the I-V curve have been useful in determining a inner potential distribution which can be a effective guide in developing the theory of the hybrid mode. During the negative resistance region ((3)-(4))on the I-V curve, the discharge has been observed to look as a "ball of fire".3) Motive diagrams for the negative resistance region on the I-V curve are shown in Fig.2. At (4) on the I-V curve, the ball has been observed to be positioned near the collector.

b a l l of

/"VT::L-Z

V\

ICA3

arc mode

' negative resistance region unaicmode^ v:

V,

nv]

Fig. 1 Schematics of the critical I-V curve (pd < 20, ^ < 1)

fir*

ball of f i r *

^4>z\

/

v/

4>z

FL( i )

V4=o Vi Vt

0

XQ

(b)

Fig. 2 Motive diagram for the negative resistance region of I-V curve. (a) corresponding to (4) of Fig. 1, (b) corresponding to (3) of Fig. 1 As soon as the current increases the ball moves toward the inter electrode space, as shown in Fig.2(a). At the point of (3) on the I-V curve, the ball has been observed to positioned near the emitter, as shown in Fig.2(b). Probe measurements indicate the existence of an accelerating electric field for electrons near the collector which leads to the formation of the ball. The mechanisms of the ball formation depend on the collector work function. If we prepare two kinds of collector work functions, then the lower forms the strong accelerating field, and the higher forms the weak field, as shown in Fig.3. The former operates as an arc mode, and the latter operates as an unarc mode.There are two modes together in the converter, so called hybrid mode.

683 This new hybrid mode model, however, can be realized very rarely. If we have various collector work functions the probability of the realization of the model increases. Thermionic converters usually use cesiated, refractory metal electrodes. The work functions of these surfaces are dependent on the electrode materials. Accordingly, the collector materials produced by FGM technology can have a spatially dependent work function. In other words, FGM technology makes it possible to realize the new hybrid mode.

I

pr*-are nodt

7^

^

Fig. 3 Motive diagram for the hybrid mode corresponding to the two kinds of work function of the collector.

EXPERIMENTAL The ion beam sputtering (IBS) system is used in the present experiment as shown in Fig.4. The chamber was pumped down to 8 X 10~^ Torr prior to the deposition using the cryopump. The ion source is the hotcathode type with an effective beam diameter of 30 mm. Argon gas was supplied into the ion source and the pressure in the sputtering chamber was set 1.6 x 10"'^ Torr during operation. The target was designed for FGM technology. Other operation conditions of IBS system are a,s follow : Ion energy 1.2 keV and current density 0.4 mA/cvn? . The deposition rate of Nb/Mo under those condition was about 17.0 A / min.

^

1). Beam Shutter ft Qineot Deast^ Kiisoc. 2).Neutiillz»r.

3).T«i9ct. 4). Sutatntc.

Fig.4 Schematic diagram of ion beam sputtering system.

RESULTS AND DISCUSSION By the IBS system, about 500 Athin film deposition, (Mo/Nb) was formed on the substrate (Inconel : 100 x 100mm)during about 30 min. As another, we picked up the test samples (Inconel : 10 x 10) which placed along the substrate at intervales of 10 mm. The detection of the MQ/NI deposited on the Inconel was executed by XPS. As seen on Fig.5, the presence of Mo is recognized by the peaks at 227.7 eV for 3^5/2 and 230.9 eV for 3^3/2. In the same way, the peaks at 202.7 eV for 3^5/2 and 205.4 eV for 3^3/2, identify A^^, as seen on Fig.6. Since the ratio of the peaks in the spectra at each point corresponds to the ratio of the concentrations of each material, we ploted them as a function of distance. As seen on Fig.7, the distribution of the components is fiat, when the target materials are placed like a symetrical mosaic. On the contrary, for the unsymetrical mosaic target.

684 as seen on Fig. 8, the distribution of the components has a gradient. The slope for Mo is about -0.6 %/cm, that of Ni is about +0.6 %/cm. Those values are very small, but they are FGM structured collectors. Mo 3d

Nb3d 1

10000

t'

3000 'w*

j\\iftVA/ \

8-

III

C

1000 230 Binding Energy (eV)

•

•

'

1

lA V /

.

'

240

1

,

200

i\ ]\

I

U

•u

^AWNM/^^NV^

1

i

i.„ -1

i

1

H

1, , J

210 Binding Energy (eV)

1—J

220

100

uu Mo 0

0 b

"

0 0

0

^

Nb

X X

X

n —1

1

0 ^-

^

^^"iO X(cnn)

Target

Fig. 7 Distribution of the deposited atoms from the symmetrical target.

.Nb ou

30 o \ .

0

MoV ^ r* 0 ^ 700"

Target

1 J

Fig. 6 XPS Nb spectra at 3 locations on the substrate.

Fig. 5 XPS Mo spectra at 3 locations on the substrate.

•Mo-

1

V\,^i,Mh\-f^*f*^ xi A^M\

'**-V*' 1 \

32000

220

'

2 0-

~^'' Nb," ^

v^ X t

X,

1

X.

1

1 20 1

X'lO

X(cm)

Fig. 8 Distribution of the deposited atoms from the unsymmetrical target.

685

SUMMARY In this study, we proposed a new hybrid mode model of the thermionic energy converter with a FGM structured collector. Based on the model, we have experimented to deposit graded Unearly molybdenum (Mo)and niobium {Nb) on the substrate considered as the collector. The results can be summarized as follows: 1. "The ball of fire" yielded on the collector is available for the hybrid mode. 2. This is not to say that the model realize usually. However, FGM technology makes it possoble to realize the hybrid mode easily. 3. In the present experiment, Ion Beam Sputtering system was used, in which we designed unsymetrical target for the FGM technology. 4. The existence of Mo on the substrate was recognized by the XPS peaks, at 227.7 eV for 3 (^5/2 and 230.9 eV for 3 ^3/2, and for, A^^ at 202.7 eV for 3 ^5/2 and 205.4 eV for 3 ^3/2-

5. Results of XPS for the samples versus distance showed concentration gradients of -0.6 %/cm for Mo and +0.6 %/cm for A^^,. Thus, we have obtained a FGM collector, whose cesiated effective work function can be predicted by the SIMCON code. 4),5)

ACKNOWLEDGMENT This paper is a part of results prepared as an account of the Government sponsored work; the Ministry of scientific and Technical Administration of Japan. The authors would like to thank Mr.Watanabe Noboru of the department of general study, and Mr.Azuma Shiro of the department of material development for their efforts in support of this study. Many thanks are also due to the chairman of the FGM - II project. Dr. Niino Masayuki, to the chief of the working group. Dr. Fukuda Ryuzo, and all of the member of FGM - II project for the useful advices.

REFERENCES 1) R. Watanabe : "Functionally Graded Material" committee of the FGM.(1993) 2) M. Koizumi and M. Niino : "Overview of FGM Research in Japan" MRS BULLETIN/JANUARY (1995) 19 3) J. M. Houston, etal : J. Appl. Phys. vol.38, (1967) 3425 4) D. R. Wilkins : AEC Research and Development Report, GESR - 2109 (1968) 5) N. S. Rasor and C. Warner : J, Appl. Phys. vol.35 (1964) 2589

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I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

687

THERMOELECTRICALLY MODULATEDE/NANOSCALE MULTILAYERED GRADIENT MATERIALS FOR APPLICATION IN THE ELECTROMAGNETIC GUN SYSTEMS M.A.OTOONI^, John F. ATKINSON^, and I.G.BROWN^, US Army Armament Research, Development and Engineering Center, Picatinny Arsenal, N.J. USA Science and Thechnology Center, Far—East Lawrence Berkeley Laboratory, University of California, Berkeley, CA ABSTRACT Analysis of fired rails from electromagnetic railguns indicates severe surface damage occurs due to high current arcing and tribological mismatch. We have explored the behavior of several nanoscale multilayered materials as possible routes to improve the thermomechanical properties of the rail and armature materials Structures investigated include (i) Ti—Co alloy on Ta—Cu alloy on die (diamond—like carbon) on stainless steel; (ii) Ti—Co alloy on Ta—Cu alloy on die on Cu, (iii) Ti—Co alloy on Ta—Cu on Cu; and (iv) Ti—Co on Ta—Cu alloy on Al. The alloys were all 50:50 at% and film thicknesses were in the range 400-1000 A. The films were formed using a repetitively pulsed vacuum arc plasma deposition method with substrate biasing- and IB AD—like techniques. The surfaces were characterized by scanning electron microscopy, transmission electron microscopy, Rutherford backscattering spectroscopy, optical microscopy, microhardness measurements, arc erosion resistance and scratch resistance tests. Preliminary results show improvement in the microhardness, arc erosion resistance and scratch resistance, most especially for the die—coated surfaces. This kind of multilayered approach to the fabrication of electromagnetic railgun and armature surfaces could be important for future advanced Electromagnetic EM Gun systems. l.INTRODUCnON Copper and aluminum have been used in the design of the electromagnetic (EM) rails and armatures. The primary reasons for selecting these materials are based on the high electrical and high heat conductivities of both metals. Recent data indicate that in addition to these two important properties, the creep behavior also plays an important role. This is particularly so when the EM gun is designed to function in the hypervelocity regime with repeated firing capabihty. Research on the materials behavior of copper railguns and aluminum armatures was initiated in the early 1980s

688 by several universities and centers where most of the EM gun research thus far has been conducted. Based on these efforts, the thermomechanical behavior of most elements and alloys materials used in EM system, such as Cu, Mo, W, and Al alloys, has been known for some time. Thermoionic and thermoelectric properties have also been investigated.

In addition, extensive data on the behavior of these materials have now become available through simulation, structural design studies, laboratory studies of materials commonly employed in various phases of the EM gun, theoretical results derived from application of the imposed electrodynamic conditions, fluid materials. These investigations have shown that the electromagnetic dynamics, and the characteristics of field tested launch package and electrodynamic properties of the railgun vary with time and space throughout the length of the gun barrel. From a structural design standpoint these variations, if true, give rise to severe electromagnetic and electromechanical responses which may be unattainable from materials currently used in EM Gun design. This is particularly true when singleelement materials are used, i.e., Cu for the radl or Al for the armature or sabot designs. In the light of these analyses, which are supported by many field demonstration results, there appears to be no adequate justification for using these materials in electromagnetic guns other than for their superior electrical and heat conductivities. These requirements are indeed necessary but insufficient for modem EM Gun systems. In view of the transient thermomechanical and electrodynamical functions of the rail and armature two avenues for improvement may be possible: (a) appUcation of advanced coating materials on railgun and armature where simulation results indicate application of coatings may prevent localized degradation due to localized melting, wear and spark erosions; and (b) a new design of the rail and launch package (armature and sabot) from a new class of synthesized materials whose properties are made to correspond to the transient fluid dynamics of the system. These new classes of materials are collectively known as gradient materials[l-4]. It is envisoned that future designs of improved EM systems will incorporate combinations of these two strategies for a significant enhancement in EM gun performance. We have synthesized and examined the behavior of several nanoscale multilayered materials that could provide possible routes to improved thermomechnical behavior of rail and armature materials. 2.EXPERIMENTAL PROCEDURE 2.1. Metallurgical Procedure Several nanoscale multilayered structures were prepared. These included (i) Ti—Co alloy on Ta—Cu alloy on dlc(diamond—like carbon) on stainless steel; (ii) Ti-Co alloy on Ta~Cu on die on Cu; (iii) Ti-Co aUoy on Ta-Cu on Cu; (iv) Ti-Co alloy on Ta-Cu alloy on Al. The alloys were all 50:50 at% and film thicknesses were in the range of 400—lOOOA. The specimens were plasma processed, and then evaluated by scratch testing, spark erosion testing, Rutherford Backscattering spectroscopy (RBS), scanning electron microscopy (SEM) , transmission electron microscopy(TEM), Reflection High Energy Electron Di&action, optical microscopy, and microhardness measurements. The plasma processing technique and other treatments employed in the preparation of the test specimens are briefly described below.

689 2.2.Plasma Processing Vacuum arc plasma discharges are intense sources of dense metal plasma, and can be used to deposit metal alloy thin films of various kind including both conventional alloys as weU as non—equilibrium alloys. In our approach, the basic plasma deposition process is combined with the ion bombardment; the method is environmentally friendly, highly efficient, can be scaled up to large size, and can synthesize films of a wide range of materials[5-9]. A metal plasma of the required species is formed by a vacuum arc plasma gun and directed towards the substrate with a moderate streaming energy, typically in order of 100 eV. At the same time, the substrate is repetitively pulsed biased to moderate negative voltage (typically a few hundred volts to a few tens of kilovolts), thereby accelerating a fraction of the incident ion flux and energetically bombarding the ions into the substrate and the previously—deposited film. This technique provides a means for precise control of the energy of the depositing plasma ions. At early—times high ion energy will be used as to atomically mix the film into the substrate, and a later—time, when the film is in the growth stage, lower energy ions bombardment is used to add an " ion assist" to the deposition - a process which is similar to an ion beam assisted deposition or DBAD technique. In this way the film is ion stitched to the substrate and has very strong adhesion, high density (void—free), good microstructur and excellent morphology (close to being atomically smooth). Using this plasma materials synthesis techniques four nanoscale multilayered specimens with different substrate materials have been prepared. Two of the specimens, one stainless steel and one copper, were coated with die (diamond-like carbon) prior to the plasma processing. Two other specimens, one 7075—T6 aluminum and one copper, were directly subjected to the plasma processing. Table 1 describes characteristic features of the samples. Table 1 The four nanoscale multilayered structures were produced as indicated. A. B. C. D.

Multilayered Multilayered Multilayered Multilayered

SS/dlc/Ta-Cu(50:50)/Ti-Co(50:50) Cu/dlc-/Ta-Cu(50:50)/Ti-Co(50:50) Al(7075-T6)/Ta-Cu(50:50)/Ti-Co(50:50) Cu/Ta-Cu(50:50)/Ti-Co(50:50)

To form the Ta—Co and Ti—Co films, vacuum arc cathodes were first fabricated by using pressed powders of 50:50 at% composition ratios. The plasma formed by the plasma gun has a composition that approximately reflects the composition of the cathode[10-13]. 2.3.Test Procedure Following surface processing, the test coupons were tested for relative improvement in their surface characteristics, and depth profiled. The characterization techniques included scratch testing, spark erosion testing, Rutherford Backscattering, scanning electron microscopy, optical microscopy, high resolution transmission electron microscopy, and reflection high energy electron diffraction microscopy. Scratch testing was done using a simple but accurate device in which a diamond probe was drawn across the surface at a

690 constant rate and constant loading. The scratch mark was of a length of a few mm. A region of the coupon was masked during the plasma processing of the surface so as to provide an unprocessed reference region. Spark erosion testing was performed using a specially made instrument in

which

a

single, highly—reproducible

spark

was formed

by

the

discharge

of

a 10 F capacitor at 3.5 kV though a precise gap in which the test surface was one of the electrodes. The spark craters so formed were then examined under an optical microscope. The measurements were carried out several times for both the scratch testing and the spark to insure consistency of the results. Elemental depth profiling was done with RBS using 1.8 MeV He ions. Scanning electron microscopy and cross—sectional transmission electron microscopy were employed in an effort to examine the nature of the several boundary layers of the multilayered structure. Reflection high energy electron diffraction technique was also employed to study the nature of the coated surfaces. 3.RESULTS AND

DISCUSSION

S.I.Scratch Tests Scratch test measurements indicated that die is a far more scratchresistant surface (i.e.,it is harder) than any of the metal surfaces used in the present work, as expected. It is possible that a more extensive testing progranni, perhaps scanning in diamond stylus loading, could reveal further differences between the different metal surfaces. This procedure wiU be attempted in the future and will be the presented in another paper at a later date. Several specimens were subjected to microhardness measurements. The microhardness values varied from 3365 to 3942±66 in the Hv scale. 3.2.Spark Erosion Tests Erosion craters from single discharges of the spark tester were formed on surfaces of all specimens of the multilayered materials. Figures 1—3 show spark tests and the photomicrographs of the damaged craters of the four specimens. The geometry of the craters on the die surface is quite different from that of the other surfaces. This effect could be attributed to different electrical conductivities of the surfaces, particularly because of the addition of Mo as dopant to make the die somewhat conductive. The depth of the craters on the die surface are also much shallower, indicating a much harder surface.

MO

\

^

MkV >

UW

^ ^

lOiloekO

rM 1 -^

1 T

^?

I v = ^

Figure 1 Schematic representation of the spark erosion apparatus

691

Figure 3 Optical photographs of spark erosion craters of the nanoscale Multilayer A after coating. Mag.lOOx

Figure 2 Optical photographs of spark erosion craters of the nanoscale Multilayer A before coating. Mag.lOOx

3.3.Rutherford Backscattering Spectroscopy RBS was used to measure the concentration and depth profiles of the implanted Ta in SS/dlc/Ta-Cu(50:50)/-Ti-Co, Cu/dlc/Ta-Cu(50:50)/Ti-Co, Al-Ta Cu (50:50)-Ti-Co, and Cu/Ta-Cu-(50:50)/Co-Ti surfaces. Depth profiles of several nanoscale multilayered specimens are shown in Figures 4 - 5 . As can be seen from these results, the Ta is implanted to a depth of about 150-170A from the surface. The composition is shown in atomic percent. While the presence of the implanted Ta ion has the effect of increasing the microhardness of the near-surface layer, the optimum values for the Ta implantation dose and energy for maximum hardening will need to be tested at a later stage. These specimens were Ta implanted to a dose of 4-4.5x10*^ cm ^ and energy range of 100 keV. Note that for these high doses, the applied implantation dose is not the same as the retained implantation dose because of the effect of surface sputtering including sputtering away of the previously implanted Ta. The deposited concentration of Ta ions varies in the range of 0.3—0.4 atomic fractions. 3.4.Electron Microscopy and Microanalysis A scanning electron micrograph showing the TaN on die on SS is shown in Figures 6(a,b). The micrograph indicates the nature of the interface to be r m 1 1 1 1 1 1 1 1 1 1 1 1 1 1 1 1 1 I M 1 1 1 1 1 1 t I t T T !T I1 1 1 1 1I I I1! 1-nq 2 #2 Tl/Co, Ta/Cu on Cu

: =- • •• ••• I

2 Li- 0.4

E o

oOoO^% o o

-• t c

• o • o

r.

o

oooooo

ho

* o

O O

..

O

••••A

•Oo^ • *ooO o • o • . o • .

Co Ti Ta Cu

i 1

-\ -. ij

• o •f

: ;

[h *^¥> p. , M i l l i » i » i # t T i 1111111111111111111 1 1 1 iTi^PtH^i *M 200

(o)

400

600

800

1000

Depth (e15/cm2)

Figure 4 RBS graph from Multilayer B showing atomic fractions of the elements as a function of Depth

0.8

1 1 I I I 1 1 1 I I 11 1 1 1 1 1 1 M 1 r I i-rr111.. j i i ini -ir m i i- r

#2 Tl/Co, Ta/Cu on A17075

F 0.6 tr • L *

c .2

-•

I 0.4,P° 3

o E o < 0.2

• o

••**•. * °°o

• o

-

o

o

•

d

*•

•

• •

o

Co

q

Ti

H

Ta

d

Cu A J

•

q

L

•

J ^< ° •. % d <^ ° O* • . . "^ • O ^ : =,., l i i i i ^ U ^ i i i i i i l M r i . ^ i f t n l t , ' LJ " " .t>l^^l**fi.«ivi

0

^lo)

200

400

600

Depth (e15/cm2)

Figure 5 RBS graph from Multilayer C showing atomic fractions of the elements as a function of Depth

1000

692

fairly smooth. EDAX from the same region of the fihn shows concentrations of the various elements in the muMayer fihn. A high resolution transmission electron micrograph of the interface between die and Cu in the multilayer Cu-dlc-Ta-TaN-Ti/Co is shown in Figure 7. This micrograph is especially interesting because it shows the die has recrystallized to a diamond structure particulate embedded in the copper. The heating of the copper substrate either during deposition of the die layer or during the implantation process may have contributed to the transformation process. The presence of diamond nuclei in the copper surface may indeed provide an excellent explanation for the scratch-resistant properties of the die surface. Figure 8 shows a RHEED electron micrograph of the same region of the Cu-dlc interface, and here again there appears to be an indication of an epitaxial relationship between the multilayer and the substrate. These data are preliminary and will be further investigated and reported on in general detail later.

ENERGY

(KEU>

Figure 6 (a)Scanning electron micrograph showing armagements of the coated layers in nanoscale Multilayer B(Mag.lOk) (b)The associated EDAX

Figure 7 Transmission electron micrograph showing the nature of microstructure in nanosale Multilayer B and its associated electron diffiracti on (Mag. 100k)

Figure 8 High resolution electron micrograph from a region of Cu/dlc interface indicating evidence of orientational relatinship

693 4.CONCLUSIONS Several nanoscale multilayered materials have been prepared. Techniques of Rutherford backscattering, electron microscopy and microanalysis and other metallurgical tools have been used to investigate wear resistant, scratch resistant, microhardness, and spark erosion properties of these nanoscale multilayered materials. Preliminary results indicate that nanoscale multilayered materials with improved thermomechanical, properties can be synthesized for application in the EM gun system. AppHcation of ion beam technology for the synthesis of gradient materials appears to have great potential for design of new materials with improved properties to be used in fabrication of many armament materials. ACKNOWLEDGMENTS The financial support for this research has been provided by the Armament Engineering Directorate of the US Army ARDEC. The bulk of the electron microscopy and microanalysis work were performed at the University of Pennsylvania and the Phillips Electronic Laboratory at Mahwah, New Jersey. The author is indeed indebted to Professor David Luzzi and Dr. Shue Chen Wang of the University of Pennsylvania whose efforts made this work possible. REFERENCES 1. T.Hirano, T.Yamada, M.Nino, and A.Kumakawa, Proc.l6th Int.Symp. on Space Technology and Sciences, Sapporo, Japan, May (1988)23. 2. F.Erdogan and M.Bakioglu, Int.J.Fractured Mechanics, 13(1977)739-747. 3. M.Otooni, I.Brown and S.Foner, Mat.Res.Soc.Symp.Proc.,316(1994)569-575. 4. M.Yamanouchi, M.Koizumi, T.Hirai, and I.Shiota(eds.). Proc. of 1st Int. Symp. on Functionally Gradient Materials, FGM Forum,Sendai,Japan(1990). 5. D.M.Ruk, H.Emig, B.Wolf, I.G.Brown, and Bo Torp, Vacuum, 39(1989)1191. 6. H.Zhang, X.Zhang, S.Zhang and Z.Han, Rev. Sd. lustrum, 63(1992)2431. 7. I.G.Brown and H.Shiraishi, EEE, Trans.Plasma Sci. PS-18(1990) 170. 8. H.Shiraisi and I.G.Brown, Rev.Sd. lustrum., 61(1990)3775. 9. I.G.Brown, M.R.Dickinson, P.B.Fojas, and R.A.MacGill, Rev.Sd.Instrum., 65(1994)1441. 10. G.Brown, Wiley, " Arc Beam Technique" , New York, 371,(1989). 11. A.W.Swanson and A.E.Bell.in " The Physics and Technology of Ion Sources" , edited by I.G.Brown, New York, Wiley, 313, 1989. 12. R.G.Wilson and G.R.Brewer, " Ion Beams with Applications to Ion Implantation" , Wiley, New York, 1973. 13. P.Spsdke, H.Emig, J.Klabunde, D.M.Ruke, B.H.Wolf and I.G.Brown, Nucl. lustrum. Meth., A.(1989)278,643.

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I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

695

Synthesis of In - Sb alloys by directional solidification in microgravity and normal gravity condition Hideki MINAGAWA, Yoshikazu SUZUKI, Katsuyoshi SHIMOKAWA, Yoshinobu UEDA, Jiro NAGAO, and Junichi KAWABATA Materials Division, Hokkaido National Industrial Research Institute, Agency of Industry and Science Technology, Ministry of International Trade and Industry. Sapporo 062, JAPAN, Buoyancy convection and gravitational segregation caused by density differences can be suppressed under microgravity conditions. Therefore, it is expected that high quality materials unobtainable on the ground can be produced by homogeneous mixing materials with different densities. On the other hand, compositional gradient which had once formed under normal gravity condition is maintained under microgravity condition. This compositional gradient will be available for the production of functionally graded material. In this study, as a fundamental research for production of functionally graded material, the synthesis of indium antimonide with several compositional ratio has been carried out under a short time microgravity conditions. Microgravity experiment has been performed by 1.2 sec drop tower facility in Hokkaido National Research Institute ( HNIRI). The structure of indium antimonide alloy concerning with composition has been investigated.

1. INTRODUCTION Indium antimonide have been studied by Hall element of electronic device and photo elective conversion device. Especially, Indium antimonide cell is famous as photo conductive device and photo magnetic device for infrared radiations^ Moreover, lattice constant of indium antimonide is expected to be varied by doping of other element or compositional gradation. Heterogeneous junction at electrode by the change of lattice constant can be applied for the reduction of the potential barrier of conductive band. Therefore, compositional graded indium antimonide with a view to produce functionally graded material seems to be available for the manufacturing of new functionally electronical devices. In this study, in order to produce the indium antimonide as a functionally graded material, compositional change due to powder mixture ratio of indium and antimonide has been examined. Generally, it is difficult to produce high quality InSb alloy because of contamination from the atmosphere ( water vapor, oxygen, and etc. ) and gravitational segregation due to difference of crystal density in the melting matrix. Therefore, multiple zone refining ^ or use of the floating zone method ^ is used for the production and the concentration of InSb alloys. Method of melting under ultra high vacuum is not applicable, because of the high vapor

696 pressure of In an Sb. Moreover, exclusive melting in the closed tube is difficult, because a little amount of residual gas in the ampoule would be dissolved in the InSb molten metal as contamination. Hereby, a semi-closed tube (effusive glass ampoule) arrangement "^ combined with ultra high vacuum system has been employed in the production of high quality InSb alloys. The semi - closed tube allows continuous removal of impurities. Furthermore, production of the homogeneous InSb alloy has been performed under microgravity conditions. In this paper, the difference of the structure of indium antimonide alloy produced under microgravity conditions is described. 2. EXPERIMENTAL The mixture powder of In (-10 jim) and Sb (-10 fim) were stoichiometrically mixed. The mixture powder of about 0.2 g was pressed to form a disk-shaped pellet ( 4.0 mm^ x 2.0 mm') with pressing pressure of 100 MPa which corresponds to the packing density of 80 %. The mixture powder of atomic ratio of indium and antimony is 0:100, 10:90, 20:80, 30:70, 40:60, 50:50, 60:40, 75:25, and 100:0, respectively. The samples were reduced under hydrogen atmosphere with 1 atm at the temperature of 973 K in 30 minutes in the quartz glass crucible. The microgavity in the time duration of 1.2 seconds is obtained on the 10 m Drop Tower of HNIRI. The Ultra high vacuum process chamber was attached to the drop capsule of HNIRI. The UHV process chamber was evacuated by a turbo molecular pump to an ultimate pressure of 2 x 10'^ Pa. The reduced pellet was set into the quartz ampoule with hole of 0.1 mm**, the so-called, semi-closed tube. The tube was set into the furnace of the UHV chamber. When the system was evacuated, the tube was evacuated simultaneously. The temperature of the sample was monitored by chromel - alumel thermocouples. Directional solidification experiments have been performed under microgravity and normal gravity. At first, the sample was heated from room temperature ( 296 K ) to 973 K. On reaching temperature of 973 K, the system was dropped with the inner capsule and the microgravity condition. After 0.2 seconds heating at 973 K, the sample was cooled by directional exposure for liquid nitrogen. The sample was cooled from 973 K to 823 K during the 1.0 seconds of microgravity. After dropping, the sample was cooled to 573 K within 2 minutes. The pressure of the system increased from 5x10"^ Pa to 2 x 10"^ Pa due to the exposure to the liquid nitrogen. The same procedures have been carried out under normal gravity conditions. After the solidification experiments, each sample was cleaved and analyzed on the cross section of the cleaved surface. The characterization of the samples which were produced under microgravity and normal gravity have been performed by following methods. Scanning Electron Microscopy ( SEM ) for micro structure. X-ray diffraction and Laue X-ray diffraction for crystal structure, and Auger Electron Spectroscopy (AES ) for compositional analysis, respectively.

697 3. RESULTS AND DISCUSSIONS Figure 1 shows phase diagram of indium and antimony. Indium and antimony formed indium antimonide alloy at compositional ratio of 50 at.% : 50 at.%, and formed eutectic structure of Indium antimonide and antimony at compositional ratio of 70 at.% : 30 at.%. In case of the mixture powder ratio of 0:100, 100:0, and 50:50, the homogeneous structure and composition have been formed by directional solidification under microgravity. Especially, the sample with pellet mixture ratio of 50:50, the X-ray Laue diffraction of the sample similar with single crystal InSb, which suggests that highly crystallized sample have been produced under microgaviy condition. The segregation of excess element was observed as a domains except of ratio of 0:100, 100:0, and 50:50. Figure 2 shows the cross sectional Auger electron image of the sample with mixture ratio of the other composition. The light area corresponds to the indium antimonide alloy (In:Sb : 50:50 ), dark area corresponds o the compositionally excess constituent. For example, dark area corresponds to antimony and light area corresponds to indium antimonide in case of In:Sb = 10:90, and dark area corresponds to indium and light area corresponds to indium animonide in case of In:Sb = 75:25. Figures 3(a) and (b) show the concentration of constituent and domain size of constituent against atomic composition of mixture powder, respectively. Figure 3(a) consists with phase diagram of figure 1. The domain size of excess constituent was seems to change with excess constituent composition as shown in figure 3(b). The domain size increases with the difference of compositional ratio of In:Sb = 50:50. This result indicates that compositional difference from 50:50 causes the segregation with domain size of excess constituent. Furthermore, it is observed that the orientation of the domains were arranged toward one direction. This direction of the domains arrangement seems to be related with the direction of cooling gas injection. This specific solidification causes the highly crystallized structure formation under microgravity solidification. Under normal gravity solidification, the cavitation due to rapid heating, preferential sublimation of indium, and heterogeneous structure formation due to primary crystalline formation prevented from forming of highly crystal materials. It suggests, however, that 1.2 second microgravity experiment makes one dimensional solidification and formation of highly crystallized materials possible. 4. CONCLUSION The structure of the sample which were solidified under microgravity, consist with phase diagram. The compositional deviation from 50:50 of indium and antimony causes the segregation with domain size of excess constituent. The direction of domains arrangement suggested that the solidification under microgravity would be related with the direction of cooling gas injection. Furthermore, the cavitation due to rapid heating, preferential sublimation of constituent, and heterogeneous structure formation due to primary crystalline formation seems to be suppressed under microgravity conditions.

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REFERENCES 1. R.C. Bouke, S.E. Miller and W.P. AUred, J. Electrochem Soc, 106(1959)61. 2. H.C. Gatos, P.L. Moody and M.C. Lavine, J. Appl. Phys., 31(1960)212. 3. I. Nakatani, Summary Report on Science Results of Fuwatto '92, Space Experiment. Tokyo, Japan, Dec. 6-7 (1993)68. 4. B. Carswell, M. Zugrav, and F. Rosenberger, American Institute of Aeronautics and Astronautics, Inc. 90-3546-CP, (1990)1-4. 5. M. Mori and Y. Sakai, Proceedings of JASMA (Japan Society of Microgravity Application), No. 28, Nov. 14-15, (1994) Osaka, Japan. 6. Sven Fries, Proc. Int. Symp. of Space Utilization " IN SPACE", Sapporo, Japan, Nov. 10-11(1992)1-29.

699

10 Jim

100 M-m (b) In : Sb = 25 : 75

(a) In : Sb = 10 : 90

10 Jim

10|im (c)In:Sb = 40:60

(d)In:Sb = 75:25

Figure 2 : The cross sectonal Auger electron image of the sampe produced from powder mixure of indium and animonide with several ratio

700 100 Composition of / Composition of •^

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Sb concentration of powder mixture ( at. % ) (a) Concentration of constuent of indium antimonide alloy against animonide concentration of powder mixure

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i

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10 20 30 40 50 60 70 80 90 100 ^^ Sb Sb concentration of powder mixture ( at. % )

(b) Domain size of constuent of indium antimonide alloy against antimonide concentration of powder mixure Figure 3 : Concentration of constituent and domain size of constituent versus atomic composition of mixture powder

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

701

Full-colored Zinc Gallate Phosphor with Graded Composition Tadashi Endo, Kyota Uheda and Hirotsugu Takizawa

Department of Molecular Chemistry and Engineering, Faculty of Engineering, Tohoku University, Aoba, Aoba-ku, Sendai, Miyagi 980-77, JAPAN (Zni.^M,)GaP4(M=Cd^", Mn^") and Zn(Ga i.^M^)204(M=Al^", Cr^") were prepared under the conditions of 1000°C tol400°C for 3 to 24 h in air and flowing nitrogen or argon gas. The solid solutions were identified as a single phase with spinel structure. On doping Al^^ and Cd^^ ions, absorption edges were continuously changed with the concentration of dopant. In particular, the self-activated emission peak sifted to the positions in the range of 450nm to 570nm with increasing the x of (Znj.^Cd^)Ga204. The doped Mn^^ and Cr^^ ions functioned only as color centers, corresponding to the 3d-3d transitions. As a result, it was found that the intense emissions of ZnGap^ substitutes were essentially induced by the donor-acceptor recombination due to the graded and non-stoichiometric compositions. 1. Introduction The mineral spinel, MgAl204, is only one of a large class of isostructural AB2O4 compounds that have become known as spinels. The unit cell contains eight AB^O^ molecules. This gives rise to 96 interstices per unit cell of which 32 are octahedral (B) sites and 64 tetrahedral (A) sites. Normally only 16 of the B and 8 of the A sites are filled with every cations. A complicating factor in some other compounds is that the cation distribution may vary. Two extreme types of behavior may be distinguished. ZnGa204, as being self-activated phosphors, is normal spinel in which the Zn ions are in tetrahedral sites and the Ga ions in octahedral sites. In the inverse spinel half of the B ions are in tetrahedral sites, leaving the remaining B ions and all the A ions in octahedral sites. As well as the two extreme types of behavior exhibited by the normal and inverse spinels, the complete range of intermediate cation distributions is possible and in some cases, the distribution changes with temperature [1]. It is characteristics of nearly all phosphor synthesis that the spectroscopic characteristics of most phosphors depend heavily on details of the temperature and the method of processing. The present study focuses on ZnGa204 which is ^ self-activated phosphor that is blue-emitting under excitation by both ultraviolet light and low voltage electrons [2,3].

702

This study attempts to identify the experimental parameters required to achieve the full-colored emission under accelerated electrons (cathode ray tube) and ultraviolet light excitations for ZnGa204 phosphor with graded and non-stoichiometric composition. 2. Experimental (Zni.,M^)GaP4 (M=Cd^\ Mn^^) and Zn(Gai.,M,)204(M=Al'\ Cr'^) were prepared by a conventional ceramic processing. Commercial oxides with 99.99% purity were weighted in the desired proportion and mixed intimately with ethanol for 24 h by ball-milling. Successively the mixture was molded in the shape of disc under the pressure of 100 MPa and calcined at 1000°C to 1400°C for 3 h to 24 h in flowing N2 or Ar gas. In order to design the texture and structure with graded composition, the following two methods were performed; one was the crystal growth of the solid solution in the PbO-B203 flux under controlling the heating and cooling rates, and the other was the stratification and sintering of several green compacts with different compositions. A X-ray powder diffraction analysis was performed on a Rigaku X-ray diffractometer with monochromatic CuKa radiation. The scanning rate of 0.5 min"^ was chosen and silicon was used as a internal material to evaluated precisely the lattice parameters. The chemical composition of the product was determined by using Shimazu-EPMA V6, because all the products were hardly dissolved in the acid and base solutions. Luminescence measurements were performed using a Shimazu RF-540 spectrofluorometer system equipped with a liquidnitrogen cryostal. Diffuse-reflectance spectra were measured on a Shimazu UV-265 spectrometer with an integrating sphere attachment. 3. Results and Discussion According to the X-ray powder diffraction patterns, all the diffractions were faithfully indexed as cubic symmetry of spinel structure for every doped-ZnGa204. Figure 1 shows various lattice a parameters as a function of the amount of dopants in ZnGa204. The change in the lattice a parameter are monotonously and faithfully brought by the size of dopants. Accordingly, the substitutional solid solution was considered to form in whole compositional range. In the normal spinel structure, the substitutions of Mn^^ and Cd^^ ions for Zn^^ ion at tetrahedeal site, and the Al^^ and Cr^^ ions for Ga^^ ion at octahedral site were partially performed as following the nominal composition. In figs. 2-(a), (b) and (c), the diffuse reflection spectra of every solid solution are illustrated. All the optical adsorption edges are successively shifted with the amount of dopants. The corresponding peaks to the 3d-3d transition were observed in the ZnGa204

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704

respectively. Figure 3(a)and(b)show the excitation and emission spectra of ZnGa204 doped with Cd^^ mdM^\ The figure Em. Exc. demonstrates that X-0.3 X-0.2A-0-4 both exciting and x=D.1 (b) (a) x^o 1 11 X=:0^ emission spectra shift Exc. Em. \, 1 ' finl 11 'i/ulU with the amount of X=0.2 ! tvjH III JC=^0.1 XB0.3> dopants. The ^—x«o /' / similarity with the JM|x«0.6.^ shift in the diffusion ,.^r?r?^ KT F111 >»1 >>iin reflection spectra is 200 300 400 500 200 300 400 500 Waval6iH|Ui
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705 luminescence center) to the corresponding 3d-3d transition sites. This possibility was examined by decay measurements of the broad band luminescence. C.B.

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Fig.5 Configurational coordinate diagram of ZnGaP4 substitute. Figures 7(a) and (b) show the the variation of the emission intensity in the series Cr^"^ 400 GOO Wavelength(nm) and Mn^^ substitutes at room temperature The most intense luminescence are Fig.6 Excitation and emission spectra of demonstrated near x=0.01 in Zn(Gaj ZnGaP4 Substitute. ^Crjp^ and x::i0.004 m (Zni_,Mn,)Gap„ diminishing markedly with increasing x. This behavior are probably due to the concentration quenching of the substituting emission. Appropriate energy levels of the dopant ion are not devoted to the cross-relaxation, but suitable non-radiative , superexchange process via the intervening oxide ions. 140

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X.exc.=302nm for the Mn^^ dopedZnGa204. It is clear that the profile consists of two lines and crossed each other at x=0.004, whose value corresponds to the intensity maximum of emission. Consequently, the concentration quenching seems to be taken place by the formation of Mn^^-Mn^^ clusters. Further study is required to understand the energy migrations, i.e., a random walk of the excitation energy over the sublattice where the coordination polyhedra of Mn^^ ions were connected by common edge and comer in the spinel structure. Furthermore, the design of shape in is expected to realize by the

REFERENCES 1. Z. Wen-Chen, Solid State Commun., 80, 213 (1991). 2. S. Itoh, H. Toki, Y. Sato, K. Morimoto and T. Kishino, J. Electrochem. Soc, 138,1509 (1991) 3. CW.W. Hoffman and J.J. Brown, Jr., J. Inorg. Nucl. Chem., 30, 63 (1968). 4. Y. Tanabe and S.Sugano, J. Phys. Soc. Jpn., 9, 753 (1954) 5. Y. Tanabe and S.Sugano, J. Phys. Soc. Jpn., 9, 766 (1954)

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 1997 Elsevier Science B.V.

707

Synthesis and characterization of a model CuO/Sn02 oxygen sensor Preben. J. Moller*, Zheshen Li^ and Qinlin Guo* Department of Chemistry, University of Copenhagen, Universitetsparken 5, DK-2100 Copenhagen, Denmark

1, INTRODUCTION Tin dioxide (Sn02) is an insulator in its stoichiometric form, but in practice the nonstoichiometric form is always encountered and behaves as a semiconductor due to the presence of defects or deliberate doping. It is a very promising material for gas sensor application because of its very good chemical stability and good surface conductivity. It is generally agreed that the gas absorption on the material surface modifies the density of free electrons close to the surface that changes the resistance. But relatively little is known about the mechanism of the surface reactions. So it is quite important to study the microstructure and electronic properties of the surface. There are many works on polycrystalline Sn02, and it has achieved great success in gas sensor applications [1-3]. But only a few works have been carried out on single crystals. This may be because of the difficulty of growing large and perfect crystals. Shen's work [4] shows the surface structure and the sheet conductivity change with temperature of the Sn02 (110) surface. Most work related to the single crystal Sn02 are concentrated to the (110) surface while very few consider the other low index surfaces. In practice, the polycrystalline grains may have other low-index oriented surfaces, such as (101), which may behave differently as compared to (110) [5]. The properties of metal overlayers have also attracted much attention. The properties of submonolayer metal atoms on the crystal surface are quite sensitive to gas exposure and are very important in understanding the catalytic properties of the surface. In this paper, we will discuss the structural and electronic properties of interfaces formed by Cu deposits on the (110) and the (101) surfaces of Sn02, and the subsequent formation of a CuO overlayer by oxidation. Due to the lattice mismatch, epitaxial growth of CuO on the Sn02 single crystal is not expected. We will therefore deposit nanoscale small particles of Cu in small steps, and oxidize these at each step to CuO, gradually increasing the size of Corresponding author. Fax: -f 45 3532 0299; e-mail: [email protected] ^Permanent address: Surface Physics Laboratory, Fudan University, Shanghai 200433, China ^Present address: Department of Chemistry, Texas A&M University, College Station, TX 77843-3255, USA This work is sponsored by the Danish Science Research Councils through the Electroceramic FGM Materials Science II program.

708 the particles of this material in order to obtain a stronger bonding to the Sn02 surface. With this in mind, we have studied two low-index Sn02 surfaces in order to see the influence of Sn02 crystal structure in the interfacial region on sensor properties. To understand the micro-properties of that structure used as an oxygen sensor, the changes in electrical conductivity of the ultrathin sandwich-layer system was at the same time followed in situ through measurements of the changes of the current passing through this n-type ( Sn02 substrate ) and p-type (the fully oxidized CuO overlayer ) heteroj unction system as a function of oxygen exposure, thus enabling us to characterize and optimize its behaviour as an oxygen sensor (and actuator). 2. EXPERIMENTAL The two Sn02 samples were cleaved along the (110) and (101) orientations, respectively, from a natural single crystal. The sizes were about 3mmx8mm with a thickness of 0.7 mm for the (110) orientated sample and 2mmx6mm with a thickness of 0.4 mm for the (101) one. The samples were pre-washed in acetone and ethanol solutions and then in deionized water. After that, they were blown dry with dry N2 gas and mounted side by side onto the same manipulator for low-energy electron diffraction (LEED) and photoemission measurements. This will give the same experimental conditions for the two samples, enabeling us to compare their properties. The vacuum system was baked to 450K for 24 hours. The samples were cleaned by repeated annealing cycles at BOOK with and without exposure to 1x10"^ mbar O2 for 20min. All the measurements were taken at a base pressure 1x10"^^ mbar. Synchrotron-radiation-induced photoemission spectroscopy (SRPES) measurements were carried out on beamline 5 at the ASTRID synchrotron-radiation storage ring in Arhus. A Zeiss SX700 plane grating monochromator was used to give desired photon energy. The photon energies used were 75.5eV for valence-band structure measurements. SRPES data were obtained by use of a VG CLAM spectrometer running at 20 eV pass energy. Cu deposition was performed by physical vapor deposition (PVD) from a K-cell evaporator at a coppersource temperature of 1300K. The temperature of the copper source was controlled to within ± IK. The conductivity measurements were obtained in high vacuum in a different chamber which is facilitated with LEED, Auger electron spectroscopy (AES) and highresolution electron energy-loss spectroscopy (HREELS) at a base pressure of 3x10"^^ mbar. The preparation of the samples were the same as that of the SRPES chamber, but the mounting method was different as described below. 3. RESULTS AND DISCUSSION 3.1 Growth of Cu The sample surface had some carbon contamination after system bake-out. To remove the carbon, the SnO2(110) was annealed for one hour in 1x10"^ mbar O2 at 875K. After this treatment, the carbon on the SnO2(110) surface was removed and the SRPES spectra then showed only Sn and O elements. Cu was deposited onto the (110) and (101) samples simultaneously to ensure the same evaporation rate for both of them. The growth mode of Cu on the (110) surface and the corresponding LEED patterns during the evaporation are discussed

709 elsewhere [6]. Briefly described, the growth mode of Cu on (110) surface is a threedimensional crystalline growth of Cu islands (Volmer-Weber growth mode). The LEED patterns changes are as follows. Initially it is 1x1 for the clean (110) surface, but during Cu deposition the 1x1 pattern becomes weaker and weaker until an overlayer Cu (111) pattern appears. At this time, the 1x1 weak pattern is still visible. Finally, when the Cu layer is thick enough, only the Cu(lll) LEED pattern can be seen, and it became more clear and sharp. For the (101) surface the growth mode is similar to that of the (110) surface, but the LEED patterns, as shown in Fig. 1, are different. The clean SnO2(101) surface shows a 1x1 reconstruction ( Fig. la ). This surface is very stable and maintain its 1x1 structure even when the sample temperature was raised to lOOOK. This behaviour is different from that of the (110) surface [7]. The LEED pattern of the Cu-deposited (101) surface exhibits patches of polycrystalline Cu ( Fig. lb ) while the (110) surface exhibits epitaxial Cu(lll) singlecrystal film growth as for TiO2(110) [8].

(a)

(b)

Fig. 1 LEED pattern of Sn02 (101): (a) clean surface, incident electron energy £'p=41 eV, and (b) Cu-covered surface, £:p = 124 eV.

3.2 Oxidation of the Cu/Sn02 surfaces Fig.2 shows the SRPES spectra from the Cu/Sn02 system before and after oxidation. All binding energies are given with respect to the Fermi energy. For the SnO2(110) sample there is a shift of 0.7 eV to higher binding energy in all the four spectra in Fig. 2(1). We have already moved them back to zero shift in Fig. 2(1) since it is due to a charging effect. This sample is rather thick and has been preheated in air at 900K which causes the sample to have fewer oxygen vacancies. Curves (a) ( Fig. 2 ) for both surfaces ( I and II) are from the clean surfaces of Sn02 while curves (b) are from the Cu-covered Sn02 surfaces. The amount of Cu is about 0.7 monolayer. The characteristic peak at around 14.5 eV, with a photon energy of 75.5 eV, clearly indicates metal phase Cu [9]. When the samples are heated at 500K in O2 ( 1x10"^ mbar partial pressure) ( Fig. 2(c)), the peaks at binding energy 15.3 eV show that the Cu at the surfaces is oxidized into CU2O. When the temperature

710 increases to 600K ( fig. 2(d) ), very strong CuO-characteristic peaks appear. The peaks at 15.3 eV are near the detection limit. The given temperatures are only accurate to within lOOK which are less than that of the real sample surface because the thermocouple was located on the holder 3 mm from the sample. We have chosen the photon energy of 75.5 eV since it is quite near the strong resonant photoemission regime of Cu 3d->4s (the resonant peak positions are 75.5, 76.5 and 74 eV for Cu, CU2O and CuO, respectively [9-10] ), where we can easily distinguish the different Cu chemical states. SnO2(110)

SnO2(101)

Binding Energy ( eV )

Binding Energy ( eV )

Fig. 2 SRPES spectra from Sn02 surfaces from the (110) surface (I) and the (101) surface (II). (a) clean surface, (b) Cu-covered surface (0.7 monolayer), (c) the sample annealed to 500K in 1x10"^ mbar O2 for 20 min., and (d) the sample annealed to 600K in 1x10'^ mbar O9 for 20 min.

3,3 Conductivity measurements of CuO/Sn02 system Sn02 is practically an n-type semiconductor and CuO is a p-type semiconductor. They can be used to form a p-n junction to make a sensitive gas sensor [1]. To measure the conductivity of this system, gold was evaporated to the rear side of the Sn02 single crystal to achieve better ohmic contact. Then several monolayers of Cu was deposited in submonolayer steps to the front side. The cleanness of the surface was checked by AES until the impurity level was beyond the detection limit. The copper-covered Sn02 surface was at each step annealed in O2 at a pressure of 1x10"^ mbar. The temperature was increased in steps of 50 K, staying 20 min at each step, reaching finally 700K. After the first CuO layer was

711 formed, the deposition of Cu and oxidation was repeated until the CuO layer thickness finally reached 30 A. No LEED pattern appeared during the oxidation process of the copper side view overlayers and the substrate LEED was greatly attenuated even if the Cu coverage was only 0.2A. From an AES spectrum we know that the substrate Sn signal becomes smaller and smaller top view as the CuO layer becomes thicker. It is quite important also to increase the temperature in ^ steps since by increasing the temperature too fast the deposited Cu will shrink to form islands or evaporate before it has been oxidized, and that Fig.3 Schematic of sample with dewill make it difficult to form uniform layers of posited contacts CuO. After the growth of a 30A layer of CuO on SnO2(110) at which the Sn (MNN) intensity was greatly attenuated, a thick layer (thicker than aprox. 1 jam ) of gold islands was deposited through a mask onto the CuO surface. The size of the islands were 0.2mmx0.2mm in area and 0.2 mm separated from each other ( Fig.3 ). Then one electrode was moved to touch the gold in the front side of the sample while another electrode was permanent in contact to the rear side. The resistance measurement was through the p-n junction to which the oxygen diffused. All of the above work was carried out in ultra high vacuum. The resistance change of the CuO/SnO2(110) system vs O2 partial pressure at room temperature (RT) shows a resistance change of 3% as the pressure changes from 10"^ to 10'^ mbar. When the sample temperature is kept at 450K, the resistance changes about 50% when the O2 pressure is increased from 5x10'^ to 5x10"^ mbar. To make sure that this change is not due to the Sn02 substrate itself, the resistance change of the clean substrate Sn02 was measured at same temperature and same O2 pressure range as above. The change was now only 6%. It should be mentioned that for our model the sample electrode area was 0.04mm^ while the CuO layer was only 30A thick ( 3xlO"^mm). Most of the current originated directly from the part that was not exposed to O2. In a rough estimation, the area exposed to O2 is 0.2mmx4x30A=2.4x10"^ mm^, so the ratio of the area of exposed to and not 6

2

^

exposed to O2 is 2.4x10' :4xl0" =6x10' . Hence the quite large change in conductivity originates from only a very small exposed area. The conductivity change may be due, however, to a combined effect of oxygen interface diffusion and the presence of the heterojunction. 4. CONCLUSION Ultrathin films of Cu were PVD-deposited at RT onto model SnO2(110) and (101) crystal surfaces in ultrahigh vacuum. SRPES spectra show CU2O epitaxial layers were formed by using 1x10" mbar O2 pressure at the temperature of 500K to oxidize the Cu-covered Sn02 (110) and (101) surfaces. A CuO layer can be obtained at the temperature of 600K. The changes in electrical conductivity of the ultrathin sandwich-layer system was followed through in situ measurements of the changes of the current passing through this n-type and

712 p-type heteroj unction system as a function of the exposure of oxygen diffuses to the interface. The relatively large conductivity change was compared to the ratio of perimeter to area, and may be due to a combined interdiffusion and heteroj unction effect. By using an atomic-scale gradual way of synthesizing a CuO layer on top of a Sn02 crystal surface, we have joined two materials to a sensor with better functional properties. ACKNOWLEDGEMENT We are very grateful to G. Thornton and P. L. Wincott for providing the samples and to S. V. Christensen for assistance during the experiment.

REFERENCES 1. G. Sarala Devi, S. Manorama, and V. J. Rao, J. Eletrochem. Soc, 142 (1995) 2754. 2. S. Lenaerts, M. Honore, G. Huyberechts, J. Roggen, and G. Maes, Sensors and Actuators B, 18-19(1994)478. 3. G. Williams and G. S. V. Coles, Sensors and Actuators B, 15-16 (1993) 349. 4. G. L. Shen, R. Casanova, G. Thornton, and I. Colera, J. Phys: Condens. Matter, 3 (1991) S291. 5. E. De Fresart, J. Darville and J. M. Gilles, Surf. Sci., 126 (1983) 518. 6. Z. S. Li, Q. Guo, and P. J. Moller, Z. Phys. D, in press. 7. D. F. Cox, T. B. Fryberger and S. Semancik, Surf. Sci. 224 (1989) 121. 8. P. J. M0ller and M. C. Wu, Surf. Sci. 224 (1989) 265. 9. M. R. Thuler, R. L. Benbow, and Z. Hurych, Phys. Rev. B, 26 (1982) 669. 10. Z.-X. Shen, R. S. List, D. S. Dessau, F. Parmigiani, A. J. Arko, R. Bartlett, B. O. Wells, L Lindau and W. E. Spicer, Phys. Rev. B, 42 (1990) 8081.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

713

Fabrication of Magnetic Functionally Graded Material by Martensitic Transformation Technique Yoshimi Watanabe % Yuzo Nakamura ^ and Yasuyoshi Fukui ^ ^ Department of Functional Machinery and Mechanics, Shinshu University, Ueda 386, Japan '' Department of Mechanical Engineering, Kagoshima University, Kagoshima 890, Japan Abstract It is known that the paramagnetic phase in an austenitic stainless steel, such as Fe-18Cr-8Ni, transforms into a ferromagnetic a '-martensite phase by plastic deformation at low temperature. The amount of saturation magnetization due to the martensitic transformation increases with increasing plastic strain. Thus, a manufacture of magnetically graded materials based on the concept may require the inhomogeneity of plastic deformation. In the present study, it is aimed to obtain a suitable gradient of the magnetization by introducing inhomogeneous deformation and to examine the relationship between the magnetization and applied plastic strain using Fe-18Cr-8Ni. A simple model to evaluate the distributions of strain and saturation magnetization is obtained in order to clarify the results mentioned here.

1. INTRODUCTION Functionally Graded Material (FGM) is known as a new material whose compositions and microstructures are varied continuously from place to place [1-3]. FGMs can be classified into composite and monolithic materials. FGMs made of composite materials are, however, not considered to be preferable from the viewpoint of material recycling, since it is difficult to separate the dispersion phases from the matrices. In contrast, monolithic FGMs have a possibility of the recycling or the recovering of their functions using very easy methods like melting or other heat treatments. Unfortunately, very few works have concerned with the monolithic FGMs. FGMs are applicable not only to the mechanical field but also to the electronic, chemical, optical, nuclear, biomedical and other fields [1-2]. However, most of the previous studies on FGMs have deah with those mechanical function but little attention have been focused on other functions and those applications. It is well known that paramagnetic phase in austenitic stainless steels, such as an Fe-18Cr-8Ni, transforms into ferromagnetic a ' martensite phase by plastic deformation at low temperature [4,5]. Since the amount of the deformation-induced martensite increases at a larger

714 strain, the saturation magnetization of the deformed austenitic stainless steel increases with increasing strain. Therefore gradually inhomogeneous deformation is considered to bring about the change of the saturation magnetization depending on local strain [6]. The idea is the origin to manufacture a magnetic FGM using an austenitic stainless steel. The method is named a martensitic transformation technique which controls the degree of the inhomogeneity of plastic deformation [7]. However, in the previous study [6], the saturation magnetization of the austenitic stainless steel before the deformation is not zero. This is because that the specimen was used without heat treatment. In the present study, therefore, it is aimed to obtain a suitable gradient of the magnetization using an annealed Fe-18Cr-8Ni alloy. If the relation between the amount of the deformation-induced martensite and the amount of plastic deformation be known, it would be easy to design the profile of the saturation magnetization by changing the local strain. A simple model to evaluate the distributions of strain and saturation magnetization is obtained in order to clarify the results mentioned in the present study. 2. EXPERIMENTAL PROCEDURE Three kinds of tensile specimens, hereafter called Type (A), (B) and (C) respectively, were machined from a 1 mm thick plate of SUS304 stainless steel. The chemical composition of the steel was: Cr; 18.06, Ni;8.47, C;0.04, Si;0.52 Mn;1.27, P;0.026, S;0.005, Cu;0.06, Mo;0.06, N;0.049, O;0.0039, all in mass pet. The shape and dimensions of the specimens are shown in Fig. 1. Type (A) specimen is a standard test piece with a uniform cross-sectional area in its reduced gauge section (Fig. la). In this type of test pieces, a uniform plastic strain is expected to be introduced, resulting in the corresponding saturation magnetization constantly distributed along tensile axis. In Type (B) and (C) specimens, the cross-sectional areas of their reduced gauge sections decrease linearly from the left shoulders of Fig. lb and Ic (designated "O") in a direction of tensile axis. The inclination angles, defined as the angles of specimen edges from tensile axis, are 1 degree for Type (B) and 3 degrees for Type (C), respectively. All specimens were solution -annealed at 1283 K for 3.6 ks in evacuated quartz capsules and then quenched in air. In this study, it is necessary to observe the distribution of plastic strain along tensile axis precisely.

JL

:^^i'

40

Type (A) .Point 0

-I

3.P2 1

-i^

Type (B) .Point O

2.02

Type (C)

Figure 1 The dimensions of three types of specimens used for experiment. Specimen thickness is 1 mm. Type (A) specimen is a standard test piece which has an uniform cross-sectional area. Type (B) and Type (C) specimens have non -uniform cross-sectional areas. The inclination angles are 1 degree in the Type (B) specimen and 3 degrees in the Type (C) specimen, respectively.

715 The plastic strains are measured by means of point marking method. The markings with 1 mm interval in a direction of tensile axis were made by a micro-vickers hardness tester before tensile tests. Tensile tests were conducted at room temperature at a cross-head rate of 0.5 mm/min with an Instron-type testing machine. After the tensile deformation, the relative displacements between markings were measured under a precision machinery microscope, and the corresponding local strains in the direction of tensile axis were calculated. Then deformed specimens were cut into pieces of 1 mm width by a low speed cutter perpendicular to the tensile axis. The saturation magnetization of each piece was measured by magnetic balance at room temperature. 3. EXPERIMENTAL RESULTS AND DISCUSSION 3.1. Uniform Deformation in Type (A) Specimens A typical true stress-true strain curve of Type (A) specimens is shown in Fig. 2. The true strain is calculated from the displacement of cross-head. As can be seen, the fracture of the specimens occurs at a strain of about 0.5. Relation between true stress, a , and true strain, E , in uniformly deformed Type (A) specimen is given by the following Ludwik law, = a, . Ke"

(1)

where a o = 235 MPa, K = 1820 MPa and n = 0.917. Figure 3 shows the change in the saturation magnetization with strain in Type (A) specimens. It is clear that the saturation magnetization increases with increasing deformation (strain). This is because that the paramagnetic phase in austenitic stainless steels transforms into ferromagnetic a ' martensite phase by the plastic deformation. It is also found that the saturation

0.2 0.3 0.4 True Strain Figure 2 A typical true stress-true strain curve obtained by tensile deformation test of the Type (A) specimen. The presence of a strain of 0.5 causes the specimen to fracture.

0.2 0.3 0.4

0.6

True Strain Figure 3 The values of the saturation magnetization of Type (A) specimens as a function of tensile strain.

716 magnetization of a non-deformation specimen is zero. Relation between saturation magnetization, L (tesla), and true strain, E , is given by following equation. /

(2)

= 5.948^ - l.OOe^ + 0.498

3.2. Inhomogeneity of Plastic Deformation in Type (B) and (C) Specimens As shown above, the fracture of the specimens occurs at a strain of about 0.5. Therefore, the maximum local strain was limited to 0.35 in both Type (B) and (C) specimens. Figure 4 shows the change of the local strain in Type (B) and (C) specimens with distance from the point O of Fig. 1 in a direction of tensile axis. It is seen in the figure that the inhomogeneity of plastic deformation in Type (C) specimen is larger than that in Type (B) specimen. Moreover, as shown in the figure, the linear relationship between distance and the local strain is not observed, although the cross-sectional areas of their reduced gauge sections decrease linearly. The strain distributions of these specimens will be discussed later. 3.3. Magnetic Gradient in Type (B) and (C) Specimens Figure 5 shows the distributions of the saturation magnetization in Type (B) and (C) specimens. In Type (C) specimen, although the gradient of the saturation magnetization is produced between the distances of 30 mm and 50 mm, magnetic gradient is not observed at distances less than 30 mm. This implies that the inclination angle of 3 degrees is so large that the degree of inhomogeneity of plastic deformation is unsuitable for our purpose. In contrast to Type (C) specimen, the saturation magnetization is gradually distributed over the whole gauge length in Type (B) specimen. In this way, we can obtain a magnetic graded material with suitable gradient of the magnetization by the martensitic transformation technique.

0.5

0.5 0.4

1 0.3

• : Specimen (B) A : Specimen (C)

0.4

g tsa

0.2

1 • : Specimen

1 • : Specimen (C)

(B)

JT J/

T * ^ *

,

0.3

a 0.2

0.1 0

10 20 30 40 50 60 Distance / mm

Figure 4 The change of the local strain in Type (B) and (C) specimens with distance from the point O of Fig. 1 in a direction of tensile axis.

.2

0.1 0

0

A,

,

,

1

10 20 30 40 50 60 Distance / mm

Figure 5 The distributions of the saturation magnetization in Type (B) and (C) specimens.

717 3.4. Evaluation of the Distributions of Strain and Saturation Magnetization If the relation between the amount of the deformation-induced martensite and the amount of plastic deformation be known, it would be easy to design the profile of the saturation magnetization by changing the local strain. For this purpose, the following model is used Point 0 to evaluate the distributions of the plastic strain and saturation magnetization in wedge-shaped plates of SUS304. In iohomogeneously deformed specimens, the X-axis is taken in a direction of tensile axis as shown in Fig. 6. When the specimens are separated into 50 elements Fig. 6 The schematic representation of the along tensile axis, the cross-sectional area wedge-shaped specimen and reference axis, X-axis. a(x) is given as a function of jc.

a(x) = a, - (a,-a3^x(_)

(3)

where ao and aso are the areas at x=0 and x=50, respectively. If the area of the ith element changes from a, to a,' by tension, the plastic strain of the ith element is given by the following equation. (4) " t

At the same time, the plastic strain of the ith element follows the Ludwik law mentioned above. P/a. e, = [-

_0-|l/«

(5)

where F is a tensile load. The value of e is calculated iteratively with increasing the tensile load, P, by 1 MPa. The elongation of the ith element Ax, is given by the following equation. Ax, = (X, - x,.^)[cxp(8,) - 1]

(6)

Therefore, the distance from the point O to ith element after the tensile test, dt, is given by «/, = / - E Ax,

(7)

when we assume the gage length is 50 mm. Figure 7 shows distributions of the plastic strain calculated for Type (B) and Type (C) specimens. As seen in Fig. 7, the theoretical distributions fit in with the measured distributions for low strain levels. However, the calculated values are not in agreement with the experimental values at high strain levels. The lack of such coincidence at high strain levels is considered to be attributed to the contribution of strain associated with stress-induced martensite.

718 T

1

1

r-

Specimen (B) Specimen (C) //%^

a 0.3

1 ^

I

CO

: Specimen (B) : Specimen (C)

.0.4 O

0.3

0.2 "^0.1 o

10 20 30 40 50 60 70 Distance / mm Figure 7 Distributions of the plastic strain calculated for Type (B) and Type (C) specimens.

2 0 en

10

20 30 40 50 Distance / mm

60

70

Figure 8 The theoretical distributions of saturation magnetization in the wedge-shaped plates.

When Eq. (2) is substituted in the theoretical distributions of the plastic strain (Fig. 7), we obtain the theoretical distributions of saturation magnetization in the wedge-shaped plates as shown in Fig. 8. Although the calculated values differ from the measured ones at high strain levels, they are in good agreement with the experimental values at low strain levels. 3.5. Advantages of Martensitic Transformation Technique Several fabrication methods, which can produce the tailor-made distribution, are proposed. These methods are applications of the relatively new technology, and the fabrication facilities are expensive. In addition, it is difficult to obtain relatively large materials by these methods except for the centrifugal method [8-10]. On the contrary, as shown in this study, no special facility is required in the martensitic transfomiation technique. It is easy to obtain large materials by this technique. Moreover, since the deformation-induced martensite transforms into austenite when the deformed stainless steels are heated to temperatures of austenitic region, it is also easy to obtain virgin materials for FGMs. This may be one of the simplest methods to obtain virgin materials for FGMs. The magnetic FGM may be available as a position measuring device by combining it with a magnetic sensor. For instance, it can be used as a device that determines the focus point in an automatic focusing camera. REFERENCES 1. M. Koizumi and Y. T2ida,Kinzoku, 58,No4 (1988) 2. 2. M. Koizumi and K. Urabc Jetsu to Hagane, 75 (1989) 887. 3. T. Hirai and M. SsiS^ki, JSME InternationalJournal Series 1,34 (1991) 123. 4. R. Lagnemorg,^crfl Metall, 12 (1964) 823. 5. P.L. Mangonon, Jr and G. Thomas,Mem//. Trans., 1 (1970) 1587. 6. Y. Watanabe, Y. Nakamura, Y. Fukui and K. Nakanishi/. Mater, ScL Letters, 12 (1993) 326. 7. T. Hirai, "Functional Gradient Materials" (Materials Science and Technology, 17B), VCH, Germmany, (1996). 8. Y. Fukui JSME Int. J. Series lU 34, (1991) 144. 9. Y. Fukui, N. Yamanaka, Y. Watanabe and K. Nakanishi, J. Jpn. Soc. Heat Treat, 35 (1995) 11. 10. Y. Watanabe and Y. Fukui,/ Jpn. Inst. Light Metals, 46 (1996) 395.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

719

Characterization of single-crystalline Cu/Nb multilayer films by ion beam analysis S. Yamamoto, H. Naramoto, B. Tsuchiya and Y. Aoki Dept. of Materials Development, Japan Atomic Energy Research Institute, 1233 Watanuki, Takasaki, Gunma, 370-12 Japan

Successful growth of Cu/Nb single crystal multilayer films on sapphire substrates with different orientations is reported based on the crystal symmetry characterization using the planar channeling technique. The difference in lattice structure and the big lattice mismatching was overcome by adjusting the substrate temperatures. The layered structure was prepared by using electron-beam evaporation under UHV condition. The typical thicknesses of Cu and Nb layers were in the range of 50 nm to 250 nm. The exact orientation relationship among Cu, Nb layers and sapphire substrate was determined through the mapping of planar channeled points at each layer of Nb, Cu and sapphire. The growth habit depends on the substrate orientations: Nb(110)/Cu(lll)/Nb(110) on sapphire (1120) and Nb (lOOyCu (100)/Nb (100) on sapphire (0112).

1. INTRODUCTION In modern technologies, materials are exposed to the severe conditions like high temperature, high pressure, low temperature and chemically reactive environment to exert their abilities effectively. The trials to improve the materials properties in bulky phases have been performed but for further improvements the surface modification with the graded nature will become important. This concept of the graded nature is also useful to design or to synthesize new materials with special function. The multi-layered structure with a certain periodicity, so called superlattice, can be assumed to be one of the graded materials controlled in nm scale. The superlattice with one dimensional crystal nature is useful to improve the reflectivity for soft X-rays but there appear another kind of problems to be solved. For the applications to soft X-ray mirrors it is critical to prepare multilayered structure with sharp interface in an atomic scale[l]. For that purpose, two kinds of layer unit, single-crystalhne or amorphous layers, are expected typically. In the present study, the efforts are described to prepare the single-crystalline layered structure with the sharp interface. In this context, Nb/Cu system which are immiscible even at higher temperature[2], is suitable for exploring a condition of multilayered structure growth with the sharp interface but might not be so for siagle crystal growth. It is

720 demonstrated that the single-crystalline multilayered structure composed of Nb and Cu on aAI2O3 is formed with the excellent crystal quality by controlling the substrate temperatures. The crystal quality was assessed by RBS/channeling technique[3]. 2. EXPERIMENTAL Nb and Nb/Cu multilayers were deposited on three kinds of major crystallographic planes, (1120) , (0001) and (0112) sapphire (a-AlaOa) substrates using the electron beam evaporation technique under UHV condition. All of sapphire substrates were pre-heated at 1500°C for 24 hours in air to eliminate the induced strain during polishing. This process is very important also for preparing the crystallographically stepped faces. During evaporation, the vacuum in a growth chamber was maintained around 5 x 10'^ torr after a long term evaporation to trap the residual gases with deposited Nb on the chamber wall. The thickness of each layer was monitored with quartz oscillators which were calibrated with Rutherford backscattering spectroscopy (RBS) measurements. The Nb and Cu fihns were deposited at the rate of about 0.2 nm/s onto the sapphire substrate kept at TSO'^C for Nb and less than 600°C for Cu. The typical thicknesses of Cu and Nb layers were in the range of 50 nm to 250 nm, respectively. At each step of evaporation, the surface structure was examined with low energy electron diffraction (LEED) technique. The surface structure of the top Cu layer was observed by SEM to check a possibility of columnar growth. The layered samples were analyzed with RBS/channehng method using 3 MV single stage accelerator at TIARA, JAERI/Takasaki. The analyzing beams of "^He ions with energy of 1.5 to 2.7 MeV were incident on samples. The size of the beam was about 1 mm in diameter and beam current was about 10 nA typically. Backscattered particles were detected by standard surface barrier detectors at 160® and 110° to the incident beam. 1500

n—I—I—r

1—I—I—r

T—I—I—r

2.0MeV '^He'^ RBS-C 6=110° Nb(110)/sapphire(1120) Thickness: 100 nm 1000

Random

>^ 500 <110>

O Al I

I

0.5

I

hmi^\

I

1

\

I

l\ I

I

I

L

1.5 Energy (MeV) Fig. 1. 2.0 MeV ^He RBS/channeling spectra from Nb(llO) epitaxial fihn on (1120) sapphire substrate. Thickness of Nb fOm is about 100 nm.

721 3. RESULTS AND DISCUSSION Fig. 1 illustrates the <110> axial channeling results of 2.0 MeV "^He ions in single crystal Nb(llO) layer with 100 nm thickness on (1120) sapphire substrate. Xmin, the ratio between the random and the aligned yield at the fixed thin layer is an important parameter to characterize the crystal perfection. Judging from Xmin specified at the depth (~10 nm) just behind the surface peak at Nb layer, the crystal quality is ahnost the same as in a bulky Nb crystal at the corresponding depth. A comparison among the channeling data from Nb layers prepared under several different conditions shows that the quality of single crystal Nb films is dependent on the substrate temperature during evaporation. For the growth of high quality Nb single crystal layer it is needed to employ the high temperature condition. In this figure it is recognized that some amounts of disorder exist around the interface region between Nb layer and sapphire substrate, however, the interface is not mixed with each other under the present condition. At higher temperature there is a possibility to form a compound with Nb. The crystallographic analysis was examined by setting the energy windows of RBS spectra at Nb and Al component of sapphire. The orientation relationships obtained from planer channeling measurements were as follows: Nb(110)/sapphire(1120) Nb(lll)/sapphire(0001) Nb(100)/sapphire(0112). The same crystallographic relationship was observed based on TEM analysis in small area but this is the first time to examine the whole area of grown layer under the channeling condition. This kind of thin Nb layer is expected to show the different chemical and physical behaviors because clamping the atoms at surface and/or interface might induce the anisotropic nature even if Nb has the cubic nature originally. In Cu deposition on sapphire with the thickness of 100 nm, the epitaxial growth was not realized at the substrate temperature ranging from 200°C to 700°C. The structure of Cu layer tends to be highly textured perpendicular to the substrate with the increase of substrate temperature.

Fig. 2. SEM observation of Cu (50nm) deposited on Nb(110)/sapphire(1120) at different temperatures, .(a) 200°C,(b) 400°C, (c) 500°C.

722 Fig. 2 shows three kinds of SEM photographs on Cu layer deposited on Nb single crystal fikns on sapphire. An additional deposition of Cu on Nb layer has spent so much time to find out a suitable condition for the single crystal growth of Cu. The substrate temperature was changed from RT to 600 °C. Different from simple imagination, the best condition for single crystal growth with the smooth surface was around 200 °C. The SEM photographs here show that the higher substrate temperature induces the island growth, however, each island is connected coherently judging from the studies of channeling and LEED analyses. Fig. 3 is the LEED pattern from the top Cu(lll) layer on Nb(110)/sapphire(1120) taken with a conventional LEED/Auger spectrometer. As the first layer, Nb was deposited on sapphire substrate at 750 °C, and then the sample temperature was lowered down to 200 °C to assure the hetero-epitaxial growth of Cu layer. Here in this photograph one can see the formation of Cu single-crystalline layer with good quality but the present information is not good enough about the three dimensional packing of atoms. Fig. 4 illustrates RBS/channehng results of the same sample as in Fig. 3. Nb and Cu layers were deposited with the same thickness of 50 nm. By employing rather higher energy 2.7 MeV "^He"*" ions, the mass-resolution in the spectra has become good enough to judge the interface sharpness. The peak at the 2.2 MeV corresponds to the Nb layer, and the peak at 2.1 MeV to the top surface Cu layer. The orientation relationship among Cu layer, Nb layer and sapphire substrate was determined through the angular mapping of planer channeling. Curiously fee Cu(lll) matches with bcc Nb(llO) with the following relationship: Cu(lll)/Nb(110)/sapphire (1120) 1000

T — I — I — I — I — r — I — I — I — I — I — I — I — I — I — I — | -

2.7 MeV ^He^ RBS-C 6=160° 800 h

Nb

Cu(50nm)/Nb(71nm) on sapphire(1120) Cu

600

Random

.SJ 400

200

Aligned

1

Fig. 3. LEED pattern from the top Cu(lll) layer on Nb(110)/sapphire(1120) substrate.

1.5 2 Energy (MeV)

u 2.5

Fig. 4. 2.7 MeV^He RBS/channeling spectra from Cu(lll)/Nb(110) epitaxial fihn on (1120) sapphire substrate. Thickness of Cu and Nb layers are 50 nm and 71 nm, respectively.

723 The single crystal growth of Cu on Nb(100)/sapphire (0112) was also successful around 200 °C but at higher temperatures the columnar structure appeared. As a third choice, the Cu deposition on Nb(lll)/sapphire(0001) was made in the temperature range of 200 °C to 800 °C. In this case any single crystal growth was not found. Further deposition of Nb layer was examined on the top Cu layer at 200°C. The Nb deposition on sapphire at 200 °C resulted in the polycrystalline growth but Nb(llO) single crystal film was grown under the same condition. The crystallographic relationship between Cu and Nb layers can be explained by stacking the closest-packed planes of bcc Nb and fee Cu.

4. CONCLUSION As one of synthesizing techniques of functional materials with the graded nature, the effectiveness of molecular beam epitaxial growth technique was demonstrated in Cu/Nb single crystal multilayer films on sapphire substrate. The sharpness at the interface and the crystal quahty were assessed with RBS/channeling analysis. The orientation relationship among Nb layer, Cu layer and sapphire substrates was determined by mapping the planar points on the angular coordinate. The results obtained are Nb(110)/Cu(lll)/Nb(110)/aAl2O3(1120) andNb(100)/Cu(100)/Nb(100)/a-Al2O3(0li2) .

REFERENCES 1. T. W. Barbee Jr., Materials Research Society Bulletin/February(1990)37. 2. Ed. By T. B. Massalski(chief), H. Okamoto, P. R. Subramanian and L. Kacprzak, Binary Alloy Phase Diagrams 2nd edition vol. 2(ASM International, 1990). 3. For example; Ed. by J. R. Bird and J. S. Williams, Ion Beams for Materials Analysis (Academic Press, 1989). 4. D. M. Tricker and W. M. Stobbs, Phil. Mag. 71, 1037(1995). And Phil. Mag. 71, 1051(1995).

This Page Intentionally Left Blank

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

725

Enrichment of ^^Si by Infrared Laser Irradiation T.Tanaka", I.Shiota", H.Suzukib, and T.Noda^ ^Kogakuin University, 2665-1, Nakano-cho, Hachioji, Tokyo 192, Japan ^National Research Institute for Metals, 1-2-1, Sengen, Tsukuba, Ibaraki 305, Japan The enrichment of -^Si has been studied using isotope selective decomposition of Si2F6 by infrared pulse laser. ^^Si was enriched in the residual Si2F6 by the laser irradiation at 952-956 cm^ where Si2F6 containing ^sgi and ^"Si preferentially decomposed to SiF4. The ^^Si content increases with lowering the flow rate of Si2F6 and increasing the laser power. Especially the addition of inert gases such as He and Ar to the source gas accelerates the enrichment of ^^Si. Under the optimum condition, Si2F6 gas with a maximum ^^Si content of 99.72% could be continuously produced.

1. INTRODUCTION Materials of which compositions are isotopically controlled are expected to show improved physical and nuclear properties which can not be attained by usual combinations of elements 11]. Natural silicon is composed of three stable isotopes such as -^Si, -'^Si and ^"Si. If the purified -^Si is obtained, high thermal conductivity of silicon and its compounds is achievable because of suppressing isotope scattering against phonon conduction. In the present paper, effects of wavenumber and energy fluence of a laser, and gas pressure and flow rate of a reactant gas on the enrichment of ^^Si were examined using isotopic selective decomposition of Si2F6 under irradiation of a pulse CO2 laser. The effect of inert gas addition to the reactant on the concentration of '^"Si was also studied.

2. EXPERIMENTAL Hexafluorodisilane(Si2FG) with a purity better than 99%, which was prepared by fluorination of Si2CL> with ZnF2, was used as a source gas for silicon isotope separation. Figure 1 shows the schematic diagram of the experimental apparatus. The transversely excited atmospheric pressure CO2 pulse laser (LUMONICS TEA-841) was used as a light source. The laser beam was slightly converged by a ZnSe lens with a focal length of 1.5 m and introduced into a reaction cell. The reaction cell was a cylindrical stainless steel tube, 2000 mm long x 54.6 mm inner diameter, equipped with NaCl windows at both ends. The wavenumber of the laser was set

726 at 929.023-983.248 cm-i by adjusting the rear grating mirror. The laser pulse with 104 ns full width at half maximum was set at the fluence of 1.62-3.2 J. The beam size is 30 x 30mm in the front of the reaction cell and 6 x 7mm at the focal position which is the midpoint of the tube length. The flow rate and pressure of Si2FG were kept at 3.33-83.5 mm%(at standard state) and 3.34-798.0 Pa, respectively. The irradiation was performed with a repetition of 10 Hz at room temperature. Inert gases such as He, Ar and Kr were also introduced to the reaction tube with a flow rate of 1.67-167mm%. Si2F6 is decomposed by the laser irradiation to form SiFj. Both SiFj and residual Si2F6 were captured with a liquid nitrogen cold trap followed by dividing into the respective component with low temperature distillation. The isotope ratios of silicon were determined by using a quadrupole mass spectrometer (ANELVA AQA-360) from the relative ion intensities of isotope species.

NaCI Window

He.Ar.Kr Power j\ Detctor

To Vacuum

Cold Trap

Col l e c t i o n Cylinder

Fig. 1 Schematic drawing of the apparatus.

1050 1000 950 Wavenunnber (cm"^)

0^ 900

Fig.2 Infrared spectrum of Si2F6 and the laser emission lines

3. RESULTS AND DISCUSSION 3.1 Effect of wavenumber Figure 2 shows infrared spectrum of Si2FG and CO2 laser Hues in the wavenumber region of 900-1100 cm^. Si2F(; has a strong absorption peak at 990 cm^ due to antisymmetric stretching vibration of-^Si-F bond. It is considered that the absorption peaks of-'^Si and ^"Si exist at the lower wavenumber region[2]. CO2 laser has four branches of lOP, lOR, 9P and 9R in the infrared region. Since the emission lines in lOR and lOP branches of CO2 laser appear in the absorption region of Si2F(i, infrared multiple photon decomposition[3] of Si2F(} molecules including ^-^Si and ^*'Si occurs effectively by selecting appropriate wavenumber. The decomposition reaction[3] producing SiFi by the laser irradiation is assumed as Si2FG(gas) + nhv -> SiF,,(gas) + SiF2(solid) (1) A mixture of SiFi and residual Si2FG is easily separated into each component by

727 vacuum distillation utilizing a large difference in boiling point between these gases. Figure 3 shows the concentrations of 2«Si, ^^Si and ^"Si in the Si2F6 after the laser irradiation as a function of wavenumber. Other parameters such as energy fluence and repetition rate of the laser, and flow rate and pressure of Si2F6 were fixed. ^«Si, ^-^i and ^"Si abundance in the natural Si are 3.1, 4.67 and 92.23 %, respectively. When the Si2F6 was irradiated at 945-955 cm-i, SiF4 containing large amounts of -^Si and '^"Si was produced. As a result, ^sSi was concentrate in the residual Si2FG. Especially at 956.205 cm-i, ^sSi content increased to 97% from the natural abundance of 92.23% On the other hand, at 975-980 cm-i which is close to the absoii^tion spectrum of ^^Si-F bond, Si2F6 with ^sSi is preferentially decomposed to SiFj. -^Si and ^'*Si concentrations increase in the Si2FG at this wavenumber range and reached 6 and 4.3% , respectively, at 983.305 cm^ while 28Si concentration was increased slightly in the SiF4. Energy Si2F6 flow rate Pressure

1.67J 16.7mm^/s 133Pa

Wavenumber Energy pressure

100j

28si

952.925cm~^ 3.14J 133Pa

"^^--^.^^^

'\-

^

./

- 2o

95! /

29Si

1° 'o

o

> ^^^^ 1

990

980 970 960 950 940 Wave number (cm"^)

930

Fig. 3 Si isotope concentrations in the residual Si2F6 as a function of w a v e n u m b e r .

iC>0-Oi—r

1

1

50

1

1

1

1

~

0

100

Si2F6 Flow Rate (mm^/s) Fig.4 Relation between Si isotope content as a function of t h e flow rate.

3.2 Effect of flow rate and pressure of Si2Fr> Figure 4 shows the relation between concentrations of ^^Si, ^^Si and ^''Si in the residual Si2F6 and the flow rate after irradiation at 952.925 cm-i where Si2FG with -•^Si and ^"Si is preferentially decomposed. The decomposition of Si2FG especially with --^Si proceeded with lowering the flow rate of the source gas. That is, -«Si content in the residual Si2FG increased with decreasing the flow rate and attained 99.65% at 3.34 mm3/s. The concentration of 28Si in the residual Si2F6 increased with the pressure of Si2FB and exceeded 99% at around 100-600Pa. Figure 5 shows the energy fluence dependence of isotope concentrations in the residual Si2FG. -«Si content increased with increasing energy. In the present study, higher energy than 1.62 J at 956.206 cm-i, which is the most suitable wavenumber as shown in fig.3, could not be obtained because of the limitation of the laser power as shown in figure 6.

728 3.3 Effect of inert gas addition to Si2FG Figure 7 shows the concentration change in silicon isotopes in the Si2F(5 with He gas flow rate after the irradiation at 956.206 cm-^ at 1.62 J. A di-astic increase in 28Si content was observed by the addition of He and the concentration reached 99% even though the flow rate of Si2F6 was 16.7 mm^/s. It became constant almost independently on the He flow rate above 33.4 mm%. As have seen in fig. 6 and 7, the increase in laser energy and the addition of He to the source gas are effective to increase the concentration of 2«Si in the residual Si2F6. Since the spectrum of the laser shows the maximum in the energy at around 940 cm-i of lOP branch as shown in fig.6. the laser beam with higher energies than 1.62 J can be emitted if the smaller wavenumber than 956 cm-i is selected. Wavenumber pressure Si2F6 flow rate

956.206cm~^ 133Pa 16.7mnn^/s

5.0

COj laser line

4.0 LIT

3.0

S 2.0 Q)

1.0

0.5

1 Energy (J)

0.0 990

1.5

980 970 960 950 940 Wavenumber, i^/cm"^

Fig.6 Relation between laser energy and wavenumber.

Fig.5 E n e r g y dependence of Si isotope concentration in t h e residual Si2F6. Si2F6 flow rate 16.7mm^/s , Wavenumber 956.206cm"^ Energy 1.62 J

100|

930

He flow rate 33.4mm^/s pressure 133Pa SigFs flow rate 16.7mm7s

95 c

42g

(D O C

C O

29Si

o O

O

>

\ 3°Si

>•

1

1

1

1

1

1

1 1 1 1 1 1 1 1

. . . . < ^do -

100 . He Flow Rate (mm7s)

Fig. 7 Relation between Si isotope concentrations a n d He flow r a t e .

956

954 . 952 Wave number (cm )

Fig.8 Si isotope concentrations a t a m a x i m u m laser energy.

Figure 8 shows the concentration of silicon isotopes in the residual Si2F(; after

729 the irradiation with a maximum power under the He flow as a function of wavenumber. Maximum concentration of ^^Si was obtained at 952.925 cm-i with a energy of 2.8IJ which is about 2 times higher than at 956.206 cm i. Under the constant energy, the concentration of ^^Si h a s a tendency to increase with increasing the wavenumber as seen in fig.3. However, the maximum power increases with decreasing the wavenumber. Then the irradiation condition showing the maximum concentration is the optimum for ^sSi with respect to the laser power and the wavenumber. Wavenumber Si2F6 flow rate 16.7mm^s"^ Inert gas flow rate 33.4mm^s"^

100

100|

R 98

o O

o O

954 , 952 Wave number (cm ') Fig.9 Effect of i n e r t gas addition on t h e Si isotope c o n c e n t r a t i o n s .

100 Ar Flow Rate (mm^/s)

956

Fig. 10 Si isotope c o n c e n t r a t i o n as a function of Ar flow r a t e .

Table 1 S e p a r a t i o n condition for m a x i m u m 28Si c o n c e n t r a t i o n . 952.925cm"' 7.47kJ/m' 931 Pa 8.35mm^s''' 83.5mmV^

Wavenumber Energy Pressure SigPeflow r a t e A r f l o w rate Natural SizFg

Si2F6

Product SiF4

2«Si=92.23% 2^Si=4.67% 3°Si=3.10%

''Si=99.72% '^Si=0.21% '°Si=0.07%

2«Si=81.63% 2^81= 5.3% ^°Si=13.07%

lO.Ocm^

0.15cm'

7.21cm'

Residual

93

9 4 9 5 9 6 9 7 9 8 99 1 0 0 28Si C o n c e n t r a t i o n i(%)

Fig. 11 Relation between yiel a n d 2«Si c o n c e n t r a t i o n in t h e r e s i d u a l Si2FG.

Figure 9 shows the comparison between inert gases on the -"Si concentration. He and Ar gases are more effective than Ki* to concentrate -"Si. The role of inert gas in the isotope selective decomposition of Si2F(; has not been clarified. One of the explanations is made by a mixing effect of the inert gas. The present reaction tube has a large inner diameter compared to the laser beam size

730 in the front position. Furthermore, the beam is focused to the smaller diameter. That means that the laser does not irradiate the whole volume inside the reaction tube. The addition of the inert gas is therefore considered to act as a mixer so that the Si2F6 in the reaction tube is efficiently irradiated. Among inert gases, Kx does not so effectively increase the concentration of ^^Si because the mean free path of Kr is about one ninth of He though the mass is about twenty times larger.

3.4 Optimum separation condition of ^sSi In order to obtain the high concentration of ^^Si, it is necessary to increase the laser power, to decrease the flow rate of Si2F6 and to add the inert gas. Figure 10 shows the ^^Si concentration change with Ar flow rate under the maximum laser energy at 952.925 cm-^ at the Si2F6 flow rate of 8.355 mm%. The concentration of ^^Si higher than 99.5% is constantly produced up to the Ar flow rate of around 100 mm^/s. Table 1 summarizes the optimum condition to continuously produce Si2F6 with a highest concentration of ^sSi in the present study. ^^Si was concentrated from 92.23 to 99.72% though the yield efficiency is only 1.5%. The yield decreases with increasing the concentration as seen in figure 11. If the ^^Si concentration of 99% is enough, the yield of above 10%o is easily attained.

4. CONCLUSION The isotope separation for ^^Si was made using the isotope selective decomposition of Si2F6 by infrared pulse laser. The followings are concluded: 1. High enrichment of ^sSi in the residual Si2F6 was observed by the laser irradiation at 952-956 cm^. 2. The concentration of ^sSi increases with lowering the flow rate of Si2F6 and increasing the laser power. 3. The addition of inert gases, especially He and Ar, to the source gas increases the concentration of ^^Si. 4. The Si2F6 gas with a maximum ^sgi concentration of 99.72% could be continuously produced. 5. The present result indicates the practical isotope separation for -"Si will be realized.

REFERENCES 1. T.Noda, Kinzoku 7(1993)32. 2. L.Halonen, J.Mol. Spectrosc. 120(1986)175. 3. H.Suzuki, H.Ai'aki and T.Noda, SiUcon Carbide and Related Materials, l O P P u b . Ltd, 1996, p i 103.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

731

Adaptive and Functionally Graded Structure of Bamboo Shigeyasu Amada* and Naoyuki Shimizu** * Department of Mechanical Engineering, Gunma University, 1-5-1, Tenjin, Kiryu, Gunma, Japan 376 ** Mitsuba Electric Go. Ltd. 1-2681, Hirosawa, Kiryu, Gunma, Japan 376

Bamboo is a natural composite material reinforced by fibers called bundle sheaths, and also a functionally graded and hierarchical structural material. Bamboo grown on the slanted ground cannot grow straight up because it is subjected to a large bending moment due to its own weight. To avoid such a state, it will be bent toward the gravitational direction. This bent part must be reinforced in some way to support the upper portion of the bamboo, for example reinforced by reaction wood for the most of trees. This paper presents how bamboo reacts to such loading conditions. Typically, the cross-section of the bamboo culm changes from circular to elliptical shape which leads to a reduction of the stress generated by the bending moment. It is concluded that the shape of the reaction wood for bamboo is closely related to the loaded bending moment. 1. INTRODUCTION

Plants including bamboo hold a fundamental structure in common which constructs their organs on the ground. Especially, cells have a hard wall and cannot move by themselves. So, the possibility is to stack cells on top of each other like bricks for growing their body. Therefore, plants have adopted a system that segmentation is located on their top portion. They also have acquired an adaptability to their environments because they cannot move from their rooted position [1,2]. Plants have phototropism to grow toward the sun light and geotropism to support themselves for catching the sun light. Their shape and structure must be designed by a compromise of both these characteristics. Trees grown on the slanted ground develop reaction wood[3,4,5] at the bend part as shown in Fig. 1 to reinforce their trunk. There are two kinds of reaction wood, tensile for Gymnosperm and compression wood for Angiosperm. Bamboo is a natural fiber reinforced composite and functionally graded material[6,7,8]. It has

732

TENSION WOOD COMPRESSION WOOD

Fig. 1 Reaction wood been recognized that fibers in the culm cross-section distribute optimaUy under bending moment[6,9]. The graded structure of diameter, thickness and intemodal change of bamboo along the height such that surface stress becomes constant[10]. These two kinds of the graded structure are hierarchical[ll]. This paper presents the structure of the reaction wood of the bamboo grown on the slanted ground. The shape of the reaction wood is related to loading condition. Finally this plant shows an optimum structure to adapt to their environments. Z BAMBOO FOR MEASUREMENTS

Culm No.14

CulmNo.syi

^'''^ sNo.a XJf

Node No.14

Node No.O

(a) Bamboo forest

(b) Coordinate system Fig.2 Curved Bamboo

The studied bamboo is Mosou-bamboo (Phyllostachys eduhs Riv.) and 2 years old. A

733

sample bamboo in the bamboo forest is shown in Fig.2(a). Fig.2 (b) shows the curved, near root part of the bamboo. This curved bamboo is subjected to bending moment due to its own weight besides the environmental loads. The slope angle of the forest ground is about 22° from the horizontal plane and the straight part above the curved cuhn part is tilted by 11.5' from the vertical axis. Internodal is numbered in terms of the cuhn number n from the ground. Judging from Fifi:.2(b), the bend part finishes around the culm number n=14. E E^

n X TO

X

o o

Hb O

o (u) Culm No. 0

(b) Culm No. 1

' •

U-

z

0

2

4

4 • 4 # S

6

CULM NUMBER: n

Fig. 4 Profile of reaction growth ^

i

o ta TENSILE SIDE

o

(e) Culm No. 4

+ \

? ^f

CO

A

20

tc

D

m15 0) LU •z.

^

COMPRESSIVE SIDE

^10 X h-

_ 1

10 (g) Culm No. 6

i 15

(h) Culm No. 7

Fig. 3 Cross section of culm

CULM NUMBER : n Fig.5 Thickness of cross section

Fig. 3 shows the cross-sectional shapes of the bamboo cuhn from n=0 to 8. The upper side is in tensile and lower in compressive side under bending moment. The cross section of tiie bamboo cuhn grow to an elliptical shape in compressive side as well as tensile one. Assuming tiiat tiie straight bamboo culm has a circular cross section, tiie distance from its circular surface to tiie elliptical one is led by the adaptive growth and noted by H, in the tensile, H^ in the compressive side, respeaively. Fig.4 shows tiie changes of H^ and W^ witii respect to tiie cuhn number n. It is clearly

734

seen that the distribution of Hu is lager than Hg. This graph concludes that bamboo has both the properties of Gymnosperm and Angiosperm of trees although the larger reaction wood is developed in the compressive side. Fig. 5 shows the thickness change of the elliptical cross section with respect to culm number. Thickness \^ and t^ become larger than 1^, and \^ to reinforce the culm subjected to bending moment. 3. STRESS ANALYSIS OF CURVED CULM

#

: NON-ADAPTIVE BENDING

O

: ADAPTIVE

D

: TENSILE STRENGTH

STRESS

3ENDING STRESS

8[

D

]• n

D

D D D

D

D

D

D

Q.

150

6

•

• • • • • • •

• •

i> 100

4

O

o Z Q Z LU CQ

CQ

CO LU

cn 1.5

b

CO UJ

on 00

2.0

^CO ^-^

c]

b 00

ib

ZOO

10

2<

) oo

oo

O

o

O O

o ()

-1

h-

z

LU

a: h-

o 50

0

2

4

L

6

CULM NUMBER :

L

8

1 10

0

z

LU

LU _J

2 0.5 Q

z i_

1.0<

00

00

CO

0

< z o

LU

12

n

Fij?. 6 Bending? stress and tensile strenjq^th

z o z

0 0.2

0.4

0.6

O.i

NON-DIMENSIONAL HEIGHT: Y

Fifif. 7 Surface stress of culm

The weight per unit length of the bamboo culm and branches is measured. It is known that weight of the branches changes parabolically with height[12]. This weight loads on the curved bamboo as a bending moment. Using this and several assumptions, the stress at surface of the bamboo culm is evaluated and is shown in Fig. 6 by white circles. Above n=4 the surface stress is about constant and below n=4 it decreases. This result shows that the adaptive growth sets a large safety margin. Under the assumption that the adaptive growth would not occur, the bamboo cukn may keep a circular cross section whose diameter is the short diameter of the ellipse. Based on this assumption, the calculated surface stress is plotted by solid circles in Fig. 6. It is almost constant around 5.8 MPa with respect to culm numbers. The adaptive growth surpasses the surface stress about 1/2.4 at the root and 1/1.5 at the finishing position of the bend part. White squares indicate the measured strength distributions whose value is 162 MPa at root and 176 MPa at n=12. Therefore, the stress generated by the bending moment is very low value compared with the cubn strength.

735 4. SHAPE CHANGE OF CULM :MACROSCOPIC GRADED STRUCTURE

Assuming that bamboo is modeled by a cantilever beam with a circular cross section and is subjected to a uniform distributed wind load in the horizontal direction, the nondimensional surface stress of the bamboo culm is calculated and shown in Fig. 7. The surface stress maintains constant along the height except the top region, which shows that bamboo is constructed based on an optimum design. The stress at n=0 has a lower value since it must have grown to adapt to the overloading. 5. FIBER DISTRIBUTION: MICROSCOPIC GRADIENT STRUCTURE

The volume fraction V,. of fiber with respect to radius is evaluated by image analysis. The obtained results are plotted in Fig. 8 for n=l with respect to non-dimensional radius in both the larger and shorter axis of the ellipse. V, is about 40% near the outer region and 15 % in the inner region. The volume fraction in both the longer axial and shorter axial directions has almost the same distribution. The total volume fraction V^ is 23.5 % in larger axis and 25.7% in shorter axis, which is approximately the same. These facts suggest that the fiber distribution and fraction are not influenced by the adapting growth. Tensile tests and a mixed principle lead to each component property , those are, CJ^ =391.1 MPa, E, =29.3 GPa for fiber and (7^^=85.1 MPa, E^=1.5 GPa for matrix, respectively, which are very close to values[l 1] reported previously. Assuming that the bamboo culm consists of multi-elliptical cylinders with different Young's modulus and strength and that is subjected to the uniform wind load, the nondimensional stress in the larger axial direction of the culm cross section is calculated and 1 f •

1

+

DIRECTION a (Total 23.5%)

•

DIRECTION b (Total 25.7%)

1

CO

1.0 AJ-

(/)

D TENSILE STRENGTH 1 0 BENDING STRESS |

UJ

h

+»

0

z g

-J

\

h

1

,•

1 •

u <

+ f

•. •+!

f

LL LU

I

D J 0

>

0

0.2

0.4

0.6

0.8

NON-DIMENSIONAL RADIUS :R=(Di/2-r)/ti

1.0

< z g z

0.5

LLI

0

1 z o z Q

n

U U

00

\

0

D 0

D

0

D

° D

0

0

0

D

D

0

0

1

0

0.2

0.4

0.6

0.8

NON-DIMENSIONAL RADIUS: R=(Di/2 - r)/ti ( i = a, or b )

(i = a, or b)

Fig. 8 Volume fraaion of fiber

Fig. 9 Bending stress and tensile strength

1.0

736 given by white circles in Fig. 9. White squares show the non-dimensional strength derived from strength at each point divided by strength at most outer layer. Both the data show approximately the same distribution. It was concluded that the fiber distribution is formed to adapt bending stress generated by environmental loads like wind. 6. CONCLUSIONS

Studying the structure of the curved bamboo grown on the slanted ground, the following conclusions were obtained. (1) Due to bending moment exerted by the own weight, the culm cross section grows from circular to elliptic, at the same time thickness in large axial direction increases. However, the fiber distribution does not change. (2) Bamboo has the property of both Gymnosperm and Angiosperm, that is, it forms compression wood as well as tension wood. (3) The bending stress due to own weight is almost constant over the curved part of the culm. (4) The surface stress remarks approximately constant along the height. (5) Fiber distribution in radial direction is the same in both the major axial and minor axial directions of ellipse. Moreover, fibers distribute to adapt to bending moment due to the environmental loads. This study was supported partially by National Science Foundation, a research project" Physics and Chemistry of Functionally Graded Materials", Ministry of Education, Japanese Government. REFERENCES

1. Y. Tsukatani, Kagaku (in Japanese), 63 (1993) 255. 2. M. Katou, Kagaku (in Japanese), 63(1993) 263. 3. M.H. Zimmermann and C.L.Brown,"Trees Structure and Functions",SpringerVerlag(1971) 4. C. Mattheck,"Trees: The mechanical Design", Springer-Verlag(1991) 5. J. Oda, Mecha-Life (in Japanese),35(1994) 12. 6. J. Oda, Trans. Japan Soc. Mech. Engr.,46(1981) 997. 7. S. Chuujo, et al., Zairyou (in Japanese), Japanese Soc. of Materials, 39(1990) 847. 8. F. Nogata and K. Seo, FGM News(in Japanese), 8(1990) 1. 9. J. Oda, J. Mech., Transmissions, and Auto, in Design, Trans. ASME, 107 (1985) 88. 10. S.Amada, et al., J. Composite Materials, 30(1996) 800. 11. S. Amada, MRS Bulletin, 20(1995) 35. 12. S. Amada and Y. Nagase, Trans. Japan Soc. Mech. Engr., Ser. A., 62(1996) 1672 13. K.Metzger, Mundener Forstliche Hefte, 3.Hefte(1893) 188.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

737

Learning about design of FGMs from intelligent modelling system in natural composites Fumio Nogata Himeji Institute of Technology, Dept. of Mechanical Engineering, Himeji,671-22,Japan

It was found that plants have a cell-based sensing system for external mechanical stimuli, which is similar to the role of piezoelectric effects in bone, i.e., mechanical stress will change the resting electrical potential of plants therefore influencing the growth activity of their load carriers. The shape and ingenious construction observed in natural composites are a continuos process of the intelligent optimization based on the adaptive modelling with a cell-based mechanosensor. It seems that this process is the best choice to survive in hard competition for energy and external mechanical condition with minimum materials that they get in the limited living space. 1. INTORODUCTION Since the shape and ingenious construction of biological hard tissues are the result of a continuous process of optimization over thousands of years, their basic characteristic such as microstructures, functions, and modelling systems fascinate the designers of engineering structures. Examining some biological load carriers, in fact, we see that their geometry of animal bones, plant and tree stems, and other biological hard tissues, under loading, change to match mainly stress- or strain-dependent requirements[1,2]. Since the time of Wolf[3] and Meyer[4], about 130-year ago, the adaptability of bone under impressed mechanical forces and the trabecular architecture organized along principal stress trajectories have been known. A possible control mechanism for the process became apparent with the discovery of the piezoelectric effect in bone[5]. Since then, many investigators have suggested that the piezoelectric properties of bone play an important role in the remodelling of the skeletal system[6-12]. In theory this effect could translate an environmental stimulus into a biologically recognizable signal controlling growth or resorptive processes [8]. In plant, we also see that tree stems are highly optimized in terms of mechanical strength. This implies that load carriers in plants and trees are also managed by a self-optimizing system with sensing mechanisms that detect external mechanical stimuli to control the growth. We present here an adaptive modelling mechanisms in plants to gain an understanding of the principles of design and processes and to apply these findings toward developing new superior material/structure concepts; such as composites in multi-phased and functionally graded materials. 2. GENERAL MORPHOLOGICAL CHARACTERISTICS IN NATURAL COMPOSITES A structure has been defined as any assemblage of materials which is intended to sustain loads [13]. In biological structural systems, however, it is difficuh to distinguish between structures

738 and materials due to their integrated and interrelated roles. This is easily understood by considering the three ways that biological systems can accommodate or sustain external loads from the macroscopic viewpoint[14], i) by changing the microstructure(thickness/shape tailoring); e.g. bamboo, moUusk shell, ii) by changing the external size and/or shape of a body; e.g. stem of a tree and plant, shaft of a leaf and feather. iii) by combining of type (i) and (ii); e.g.animal bone, tooth, spicule of sea urchin, and rat tooth. Graded structures of type (i) are found more often than other types. Herein we show some examples of the ingenious construction found in biological materials and derive ideas and guidance for the developement of new man-made functionally graded materials. 3. EXPERIMENTAL PROCEDURES Materials: Three kinds of plants, a bamboo(Moso bamboo, Phyllostachys pubescens Maxel), a Japanese cedar(Cryptmeria japonica D. Don), and a toneriko(Fraxinux japonica Blume) were used to examine the sensing ability of their cells of stem and branch stem for external mechanical stimuli. Macroscopic voltage measurement ofplant cells: On the basis of the measurement system of an electrocardiograph(EKG) machine for human body, the changes in an electrical voltage of plant's cells when stressed mechanically were measured. A half size of diagnostic EKG electrode with adhesive paste was used, which was bought from a dealer. Electrodes were pasted on the both tension and compression sides of the stem stressed by bending. Electrical signals from the electrode were recorded on the basis of an unipolar and bipolar lead system. Base potential was used as the ground voltage around in the root system of the plant. 4. RESULTS AND DISCUSSIONS 4.1, Mechanosensingability of plant cell and adaptive modelling: a) Bamboo Authors indicated[14-17] that the placement of fibre bundles is reflected to a stress situation in the bamboo stem. Also Mattheck et al.[18,19] showed that the contour shape of biological structures such as tree stems, red deer antlers, human tibia, and tiger claws are highly optimized in terms of mechanical strength and minimum weight. This implies that biological structures may have mechanical sensing devices. Then, in order to gather information and examine the sensing ability of bamboo cells, when stress is induced by external mechanical stimuli, we tried to detect a biological signal which may be induced. Figure 1 shows an example of the voltage signal curves which were obtained from a bamboo stem subjected to an external bending moment. This curve shows the presence of a spike at loading and a spike at unloading, which is very similar to the characteristic feature of the bone signal, see the box in Fig. 1[7,8]. The higher vohage signal was recorded on the compression side rather than the tension side of the bamboo stem. In bone, it is shown that the voltage reverses on release of the stress and has a sign opposite to the original signal when the stress is reversed. Thus these characterized signals in bamboo may be used as trigger to organize adaptive growth related to the stress directions, as a role of the piezoelectric properties in bone[5,6,19]. We have confirmed that there was no voltage signal induced from a specimen boiled in a hot water bath for one hour or from a dried specimen with a weight loss of one half. It is clear that the vohage signals recorded were produced from live cells in stressed materials

739 because boiling or drying of specimens means the death of the plant cells. This shows that only live plant cells have the ability to sense some information induced by external mechanical simuli. Moso bamboo (Pbyllostacbys pubescens Mazel)

size of electrode t=0.1nim U

=iB ijnam-

stressed duration

10

^ .

adhesive paste

M=8.1 Nm Z=1.9xl0'' m'

ri) bipolar lead system «>

a = M/Z=42.6 MPa stressed duratioa 20min;

(3)=®.© AgCI electrode

:^^.

*> «0

m

>

^v Tim. -

K

1 r Y

Typical bone signal[7,8] voltmeter I

t ® Uz® W+= 1 9 ^ N

stressed duration ®(2):unipolar lead system

Figure 1 An example of the voltage signals induced by a bending moment for the stem of a bamboo. The box shows a typical bone signal reported in the literature[7,8], b) Leaf shaft of a palm tree The usual concept of a palm is a tall tree with a single stem and large feather- or fan-shaped leaves. Furthermore, palms, which are chiefly tropical and subtropical trees, stand up to very strong winds that might blow from almost any direction. Figure 2(a) shows the shape change in the cross-section, perpendicular to the leaf shaft, and also shows torsional rigidity(GIp; Gishear modulus, Ipipolar moment of inertia of area) along the shaft. This data indicates that the plant is protected from strong winds by having weak torsional rigidity(X/L=more than 0.7) which has minimum resistance, allowing the free movement of the leaf. Figure 2(b)shows a typical voltage signals induced by the tension and compression stresses for a leaf shaft of a palm tree(another specimen with that in Fig. 2(a)), which shows the detection ability of the cell. It also seems that the electrical signal controls the growth activity of a leaf shaft to mach mechanical environment. This is a typical design example for a biological structural system of type (ii). On the other hand, Fukada et al.[5] found piezoelectricity properties in bone which was stressed. There are several reports [11,20,21] which are based on evidence that bone demonstrates a piezoelectric effect. This is used to explain the concept of stress- or strain-induced bone remodelling which is often refered to as Wolfs law[3]. Thus, bone converts mechanical stress to an electrical potential that influences the activity of osteoclasts and osteoblasts [1]. It is also known that the interior structure of bone(trabecular architecture) is arranged in compressive and tensile systems corresponding to the principal stress directions [4]. The role of the voltage signals induced in bamboo and palm we found may also be similar to the piezoelectric effect in bone.

740 1350

fCl electrode

Q'(Z)'(i) 2^3.3x10*" M-53 m ®®' «o»Pol" !«*<* system 0-=M/z=i6.i HPa ©bipolar l e d system stressed duration -T—1—1—1—1—I—I—I—r

1 stressed duration

Figure 2 An example of the voltage signals induced for the leaf shaft of a palm

Softwood AgQ

Hardwood

^ 9,8

Diunipolar lead system '"^

b)

a tree growing on steep ground

Srar^

a)

Figure 3 a)structure of reaction wood [ 22] b) voltage signals induced by bending moment for two different woods 4.2. Modelling mechanism of reaction wood : In bone remodelling, it is known[3] that the adaptation of bone to mechanical stress is governed as that a bone bent by a mechanical load adapts by depositing new bone in the concave side(compression side) and resorbing the bone on the convex side(tension side). We can also see similar adaptation(growing) mechanism in plants, i.e. the compression side of a softwood tree grow up faster than tension side, called as compression wood, and the tension side of a hardwood tree is grow up faster, than compression side, called as tension wood. Then we examined the mechanizm of modelling system. Figure 3(a) shows the structure of reaction wood[22] for a Japanese cedar(softwood) and a

741 toneriko(hardwood), and those voltage signals induced by bending moment are shown in Fig.3(b). These data show clearly that the cell's ability to detect external mechanical stress isdifferent, i.e.a Japanese cedar detect mainly compression stress rather than tension and itssignal may assist to grow faster the compression side of stem of tree. On the contrary, a toneriko detect mainly tension stress rather than compression and the signal may assist to grow faster the tension side of stem. Therefore, above evidence shows that the electrical potential controls the growth activity of load carriers. Also the characteristic stress/strain states may lead to the modelling of hard tissue and ingeniously customized microstructures in tree and bamboostem. We believe that the electrical properties of plant's cell play an important role in the modelling of the skeletal system in biological hard tisssues. Dc8lgii(cnvlroiuiiciitaI) Conditions External Force, Thermal, Magnetic, Corrosion, Electric, Optical, etc

7"

Artificial materials

Biological materials

^

analysis of stress, strain

iri'Situ mechanosensing system by piezoelectricity, cell's mechanical sensor (SA Activated channel)

/sei self-repair system

by computer or experiment

z:

Modelling of graded structure by 1 «placement of fiber • microstnicture porosity • density •fibresize (baamboo) (shell) (bone)

Modelling of graded structure and fiinction

i

Biological processing (biomineralization)

Processing for FGM

Figure 4 Comparison between biological and man-made materials from the viewpoint of designing procedures Shape & placement of element

fibril

fine particle (homogeneous)

needle

lamellar

porous, cellular

°oO
Example

bamboo node

shell

shell, bamboo bone, shell tooth narwhar, antler tooth

Figure 5 Various structures in biological composites

bone spicule wood

742 5. CONCLUDING REMARKS An outline of adaptive modelling mechanisms in plnats was presented. It is clear that these mechanisms are based on an in-situ stress/strain detection system. It is the author's blief that greater understanding of these mechanisms will lead to greater advances in the modelling and processing techniques aimed at developing new materials with superior properties, such as multi-phased and functionally graded composites. Figure 4 shows a comparison between biological and man-made functionally graded materials from the viewpoint of designing procedure. It is considered that this process is the best choice to survive in hard competition for energy and external mechanical condition with minimum materials that they get in the limited living space. Through these findings, one of answers for what can be learned, the author believe, that it would be better to spend more time and money on developing functionally graded materials governed by uniform strength; for example, structure using the optimal placement of fibres, various microstructures, porous or cellular structures, etc., rather than developing new materials with high-stiffness(Fig.5). ACKNOWLEDGEMENT This research is supported by Grant-in-Aid for Scientific Research on priority Areas of Japan under contract No.274(Physcics & Chemistry of PGM). REFERENCES 1 .W.C.Hayes,B.Snyder,B.M.Levine and S. Ramaswamy, In Finite elements in biomechanics(edsGallagher, R.H.,et al.),John Wiley & Sons,(1982)223-268. 2.J.C.Koch,Am.J.Anat.,21(1917)177-198 3. J.Wolff,Uirchows Arch.,50,(1870)389-450 4. Meyer,H. Arch. Anat.Phys. 47(1867)615-628 5. E.Fukada and I.Yasuda,J.Phs. Soc. Jap.,12,No.l0,(1957)1158-1162 6. A.A.Marino and R.O.Becker,Nature, 228,(1970)473-474 7. CA.L.Bassett, and R.O.Becker,Science,137(1962)1063-1064 8. W.S.Williams and L.Breger,J.Biomechanics,8(1975)407-413 9 .C.A.L.Bassett,Calc. Tiss. Res., 1,(1968)252-272 10. CA.L.Bassett,Chap.l, in Biochemistry and Physiology of Bone(2nd edn.) Academic Press,N.Y.,3(1971) 11. A.Gjelsvik,J. Biomechanics,6( 1973)69-77 12. C.Eriksson,Electrical Properties of bone , in Biochemistry and Physiology of Bone, Academic Press,N.Y.,4(1976) 13. J.E.Gordon, Structures, Pelican Book(1978)p.l7. 14.F.Nogata,Second Inter, conf. on composites engineering,ICCE/2(1995)551-552 15.F.Nogata and H.Takahashi, J.Comp. Eng. 5-7(1995)743-751 16.F.Nogata,A.Murakami and R. Hodskinson, Proc. 4th Japan Inter.SAMPESymp.(1995)609-614, 17. F.Nogat,Cell-based mechanosensing system and adaptive self-optimum modellin in plant tissues,Materia Japan(in Japanese) 35-8(1996)886-892 18.C.Mattheck and Burkhardts„Int.JoumalFatigue,12-3(1990)185-190 19.C. Matthec,Mat.-wiss.ll., Werkstofftech.21(1990)143-168 20..R.B.Martin,J.Biomechanics ,5(1972)447-455 21.S.C.Cowin and D.H.Hegedus, J. of Elasticity(1976)313.326 22.Japan Wood Resarch Society, Wonderful world in wood,Kaiseisha,(1995)p25

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

743

Development of the fire door with functionally graded wood H.Getto ' and SJshihara ' ' R&D Division, Aica Kogyo Co.,Ltd., Japan 24, Fukami, Kamikayazu, Jimokuji-cho, Ama-gun, Aichi, 490-11, Japan ^' Wood Research Institute, Kyoto University, Japan Gokasho, Uji-shi, Kyoto, 611, Japan ABSTRACT Inflammable materials like a wood can be use for the fire door by the Ministry of Constniction Notice No.ll25 in 1990. Wood has good fire endurance when it was made with a large thickness dimension. A door, however, cannot be made with such a large thickness dimension for building purpose, especially in thin panel-stile constructed doors flames the thin panels on fire. In order to improve this behavior, fire retardant wood treated with phosphorus and nitrogen compounds was tested for use in fire doors. Here two treatment methods are compared, which are heat-dried treatment and heat-pressed one. It was recognized that the latter method is more useful than the former method, because density of wood and concentration of chemicals are raised at the surface by the former method and they decrease gradually fi-om the surface to the center. In this system it was recognized that gradient of fire resistance with chemical reactions such as phosphorylation, esterification and polymerization among chemicals and wood occurred on its surface in presence of water. Then laminated boards with surface layers treated this way were used for panels of a wooden fire door and it has been authorized by the Ministry of Construction. 1. INTRODUCTION A new test method for fire door has been standardized by the Ministry of Construction Notice No.1125 in 1990 ' " .By this notice inflammable the use of materials like wood for fire door was allowed. Wood is essentially a good fire endurance material because it forms a carbonized layer on the surface during combustion which prevent flame development. The fire endurance is proportional to the cross section dimension of it ^ ^'. A glued laminated wood for structural members called gluelam for building purposes make good use of this characteristics. It is difficult to make a door, especially a panel-stile constructed one which has a large thickness dimension. Such a door is usually made thin with a nice profile. Therefore fire retardant wood has been tried for such style wooden fire door.

744

It's well known that materials made of wood like a veneer, a chip board, a fiber board etc. are treated with phosphorus and nitrogen such as ammonium phosphate, urea-phosphoric acid mixture, dicyandiamide-phosphoric acid mixture and other phosphoric amines or amides to improve their fire resistance. They form a carbonized layer acceleratively on cellulose materials which gives them superior fire endurance. Ishihara and Kobayashi improved the fire endurance of wood by coating ^ ^ \ Here an insoluble chemicals' layer has been tried to make on a surface of wood to improve the fire endurance and the treat wood has been applied to the development of a wooden fire door. 2. EXPERIMENTAL 2.1 Fire Retardant Treatment Of Wood 2.1.1 Preparation of test pieces Chemicals; Mixture of dicyandiamide (Imol), formaldehyde (Imol) and phosphoric acid (O.Smol) [DFP (N/P=8.0) ]. Mixture of melamine (0.25mol), dicyandiamide (0.75), formaldehyde (4mol) and phosphoric acid (Imol) ^ ^ ^ [MDFP (N/P=4.5) ]. Mixture of dicyandiamide (Imol), formaldehyde (2.4mol) and di-ammonium hydrogenphosphate (O.Smol) [DFAP (N/P=7.3) ] Here N/P means nitrogen and phosphorus molar ratio in chemicals. Specimens ; 3-20mm thick Basswood (Euphorbiaceae, Endospermum meduUosum L.S.Smith) were used. Basswood was straight-grained sawn timber or rotary race veneer with a cross grain and the specific gravity was 0.42-0.47 and its moisture content was 9-11%. Sample Preparation ; Wood was soaked or impregnated with chemicals. After that they were set in a hot press and pressed at 160"C/0.5-2MPa for 5-30min. [Heat-Pressed Treatment]. Comparison specimens were set in a hot oven and dried at 160 °C for 2-6hrs. [Heat-Dried Treatment]. Heat pressed material was put on top of each other or untreated material with emulsion polymer-isocyanate adhesives to make a laminated board. 2.1.2 Test Method Fire retardance and endurance ; The fire retardance of treated wood was tested by JIS A 1322 procedure that the specimen was clamped at an angle of 45"" in the holder and the yellow flame 65mm high was withdrawn for 3min. at 50mm position from the bottom edge of it. It was valuated by a char length from this position. The fire endurance was tested by JIS A 1304 procedure and evaluated by the time until theflamecame through the specimen. Measurement of functional gradation ; Treated wood was divided into 5 layersfi:*omthe top to the bottom symmetrically. The contents of nitrogen and phosphorus of each layer were measured by Kelder's method and spectrophotometric determination. Thermogravimetric analysis (TG) and differential thermal analysis (DTA) were performed over the temperature range room temperature- 600 "C in an air of lOOml/min., using a heating rate of 10 °C/min.

745

2.2 Development Of Wooden Fire Door 2.2.1 Fireproof Ability Of Panel-Stile Constructed Wooden Door The door the size of 812W X 2000H X 45T was made of Nyatoh (Sapotaceae Palaquium spp.) with specific gravity 0.55-0.65 and moisture content 8-12%. The thickness of the panel inserted portion was 21mm. It was heated according to the standard heating curve of JIS A 1304. Fireproof abiHty was measured by time until flames came through it and at the same time temperature of unheated side were measured at various position. 2.2.2 Certificated Test Of Fire Door The door the size of 890W x 2350H X 45T was made of Nyatoh mainly. The panels consisted of Nyatoh and above mentioned fire resistant laminated board and the thickness of the panel inserted portion was 20mm. It was heated for 20min. along the test method of the notice in 1990. 3. RESULTS AND DISCUSSION 3.1 Fire Retardance And Endurance Of Wood Fire retardance of treated Basswood was almost equal to that of paper as shown in Fig.l and 2. Fire retardance appears clearly at 3-5% of chemicals loading and at more than 30% of chemicals loading it didn't improve further. Fire retardance had a tendency to be better for heat- pressed treatment than heat- dried treatment and slightly better for MDFP than DFP. At more than 10% of chemicals loading heat-pressed material showed a higher char length decrease than heat- dried one in proportion to increase of chemicals loading. These chemicals react with themselves or others ti form an insoluble reactant because fire retardant ability of treated woods didn't decrease after wash by water.

BE -(m20

DFP MDFP c) A Heat-Pressed • • Heat-Dried

20

DFP MDFP O A Before • A After

m ^10

0

10 20 30 40 Chemicals Loading (%) Fig.1 Fire Retardance Of Treated Wood Note, BE : Burned TO End

10

0

10 20 30 40 Chemicals Loading {%) Fig.2 Fire Retardance Of Heat-Pressed Wood Washed By Water

Furthermore by JIS A 1304 test method the difference between the two treatment methods was more obvious. In 20mm thick of DFP treated wood, heat-pressed

746 material showed 1.5-2 times more in fire endurance than heat-dried material and fulfilled the 100 "C condition on the unheated side as shown in Fig.3. It is believed that at high concentration of chemicals reactions such as phosphorylation, esterificarion and/or condensation occurs on the surface with some destruction of wood by heat-pressed treatment, then dehydration and carbonization occur acceleratively by combustion ' ^ \ As a result a carbonized layer is formed at the surface which prevents fire development. To carry out the reaction at the temperature required for phosphorylation or esterification a medium such as a water is necessary. The medium should be a solvent for the acid component and a swelling agent for wood . That is, it should distend the wood surface and transport the phosphorylation into cell walls. The medium should be fluid in the anhydrous state at temperatures above 130 "C and as fer as possible should be chemically inert towards a wood ' ^ ^ The chemicals content in a treated wood was found to be higher at the surface and gradually decrease as it approached the center regardless of the thickness of treated wood as shown in Fig.4 The fire-regardant function would be gradated in accordance with chemicals content. To gain sufficient fire retardance it was necessary to contain more than 1.5% of phosphorus and 5% of nitrogen content in the surface layer of treated wood. 1st (Surface)

Standard Tiffle-Tenp. Curve

(J

Ay

^10

*10 (Y

1st (Surface)

c o c= o

•4-'

/"

o Heat-Pressed •Heat-Dried

•''

10

20 30 40 Time (min.) Fig. 3 Fire Retardance Of DFP Treated food And Surface Temperature In A Fire Test

& 5 +-»

Heat-Dr ied,

4ea ieat-Pressed

C3 1=

D^--m'': /

^...,i,;;::-;::iV-

•• • C )

10 20 30 Chemicals Loading {%)

\ \

2nd 3rd (Center)

0

1 2 Layer

3

F i g . 4 Gradient Of Chemicals Contents In Each Layers Of DFP Treated Wood

It was recognized that he thermal decomposition of treated woo occurred at lower and the weight loss of it was less in proportion to increase of temperature chemicals loading compared with untreated wood by thermal analysis. This should occur successively from the surface to the center. At the surface the combustion of wood is delayed because of its increased density. The remainder becomes carbonize as a result of dehydration and carbonization by heat which prevents the fire spreading effectively. From the above mentioned, it is concluded that it is useful to increase the density of wood and the chemicals content at the surface of the fire door which provided the time to block off the fire. Fm-thermore in order to increase the chemicals content and density still a compressing treatment was tried. Like Tab.l the fire endurance of treated wood was improved by compressing treatment, even the untreated wood. Regarding to the fire endurance, compressed wood with 20% of chemicals loading performed better than uncompressed wood with 30% of it. It is concluded that the above-mentioned

747 phenomena should occur highly at the surface of wood and the fire retardance in treated wood was graded in proportion to decrease of chemicals content from the surface to the center by compression. Because the fire endurance of compressed wood was as same as that of laminated board consisting of treated wood in the surface portion and untreated one in core portion. Tab.l Fire Endurance Of 15mm Thick DFP Treated Wood H e a t - Dried Compressed Structure (Chemicals Loading : 30%) (Chemicals Loading : 30%) 18'00" 22'30" (Untreated 15mm) (Untreated20mm -> 15mm) 25'00" 3roo" (Treated 15mm) (Treated 20mm -^ 15mm) sroo" 3roo" '^<-'JrAV^<^<,<'<;; (Treated 5mm 3ply) (Treated 10mm -> 5mm) 26'30" 30'30" (Treated 3mm (Treated 5mm -> 2.5mm) on both sides) Note ; Treated Untreated,

r-:::v;i

To improve the fire endurance of wood it should be more effective and important to enhance the ability of fire retardance at the surface of wood. Then rotary raced veneers treated with DFAP chemicals were used instead of surface compressed wood. They have a lot of small cracks inside from their production and chemicals can penetrate easily through these cracks by capillarity and high chemicals loading can be gained. It was tried to use for the core of fire door's panels. 3.2 Wooden Fire Door By the fire test it was recognized that the ability of fire endurance of a wooden door depended on the thickness of materials and the flame came out from panel inserted portion which was thinnest. The thickness of the carbonized layer was about 20mm after testing. From this point it was judged it was unnecessary to use fire retardant wood for thick door parts like a stile. Meanwhile it was necessary to improve the panel inserted portion to prevent flame penetration, because panels as door parts are constructed independently and the gaps of the panel inserted portions are extend more by shrinkage with wood combustion. So for a certificated test of the fire door they were made of Nyatoh mainly by panel-stile constructed style. Their panels were consisted of Nyatoh and above mentioned fire resistant laminated board and the thickness of the panel inserted portion was 20mm. It was heated for 20min. in accordance with the test method of the notice in 1990. As a result no obvious faults were found during test, they were judged as the certificated fire door. On the heating side appearance almost all of Nyatoh surface layer of panel

748 parts dropped and the surface layer of fire retardant laminated board was only carbonized. Fire seal expanded well to prevent flame from penetrating. If the test was continued until the flame came out, they would endure for more than 40min. against fire. 4. CONCLUSION Regarding to the fire retardance or endurance of wood, there was much difference between treatment methods. That is, heat-pressed treatment improves them more than heat-dried treatment, because by the former method density of wood and chemicals content become higher than with the latter method, especially at the surface of wood. Chemicals content gradually decreased as it approached the center. Furthermore wood compressed by high pressure in heat-pressed treatment showed better fire endurance, even for untreated wood. As a result the functional gradient in wood would be more pronounced. The method is, however, not practical because compressed wood swelled in water. It is easy for an entrance fire door to be affected by moisture and water Laminated board with raised fire retardance of the surface layer showed high fire endurance as well as compressed wood. In fire tests it was recognized that the ability of fire endurance of a wooden door depended on the thickness of materials. In order to improve fire endurance of a panel-stile constructed door, it was proper to improve thinner portion of door, named panels. So laminated boards with improved fire retardance at the surface were used for a wooden fire door and it has been authorized by the Ministry of Construction. It was recognized that it was more effective to raise the fire endurance at the surface of wood.

REFERENCES 1) The Japanese Official Gazette, 379 (1990) 9 2) M. Yamada, Wood Preservation, 13 (1987) 67 3) N. Kobayashi, S. Ishihara, Proc. 18th lUFRO World Cong., (1986) 31 4) S.C. Juneja, Forest Prod. J., 22 (1972) 17 5) K. Kishore, K. Mohandas, Fire and Materials, 6 (1982) 54 6) S. Ishihara, J. Ikeda, H. Getto, Preprints of the 41st annual meeting of the Japan Wood Research Society, (1991) 395 7) J. E. Hendrix, J. E. Bostic, JR., E. S. Olsen, R. H. Barker, J. Appl. Polym. Sci., 14 (1970) 1701 8) T. Ohuchi, Y. Kumagai, M. Ono, Mokuzai Gakkaishi, 18 (1972) 557

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

749

Elemental mapping of functionally graded dental implant in biocompatibility test Fumio WATARI, Atsuro YOKOYAMA, Fuminori SASO, Motohiro UO, Shoji OHKAWA, and Takao KAWASAKI School of Dentistry, Hokkaido University, Sapporo 060, JAPAN IJNTRODUCnON The dental implant composed of biocompatible metal, H, and cenimics, hydroxyapatite(HAP), with the structure of functionally graded materials(FGM) has been fabricated to satisfy both mechanical properties and biocompatibi]ity[l,2]. The miniature FGM implant was made and implanted in the femora of rats to evaluate biocompatibility [3]. The tissue including the implant was sliced to thin section and each section was observed by optical microscopy after staining[l]. In animal test biocompatibilty is evaluatedfinomhistological observation to check the inflammation, formation of new bone, contact state of new bone at the interface with inq)lant material. Elemental mapping method has rarely been used in the field of histological investigation. Electron probe microanalysisCEPMA) is one of the most common apparatus for elemental mapping. X ray scanning analytical micn>scopy(XSAM) is a new method for the fluorescenct X lay analysis. EPMA and XSAM detect the characteristic X my emitted from the specimen by irradiation of electron and X ray, respectively. In this study elemental mapping and other imaging methods by EPMA and XSAM as well as optical microscopy were applied for the evaluation of new bone fonnation around the implant which was inserted in the mandible of rabbits. The results on the Ti implant and Ti/20%HAP FGM implant were then derived and compared each other. 2JeXPERIMENTAL PROCEDURE 2.1.Specimen preparation The FGM implant wasfebricatedwith the graded structure from pure Ti at one end, mcreasing the HAP content, to 20%HAP at the other in the longitudinal direction. FGM of this type is expressed as Tiy20HAP in the followings. Both pure Ti and FGM implants of the miniature cylindrical shape 20x7mm were made through the process of packing, compacting by CIP Mp to 400MPa and sintering at ISOCC in vacuum 10" ^Pa. The Ti and T]/20HAP FGM implants were inserted respectively in the ri^t and left mandibles of New Zealand White rabbits. The age was 13 month and the weight was approximately 3kg. After 1, 4 and 8 weeks of implantationrabbitswere sacrificed and tissue blocks including the implant material were taken immediately. After fixation

750 in a 10% bufieiBdformalinsolution, the tissue blocks were stained with Villanueva's method and then embedded in methylmethacrylate resin. They were cut vertically to the longitudinal direction of implant into the sections of the thickness about 100 //m by diamcHid blade saw and then polished for the histological observation 2^.0bservation and analysis The thin section was observed by optical miooscopy. The scanning transmission X ray image and ehnental mapping for each element forming bone and implant were obtained by XSAM(HORIBAXGT2000) using the 100//m0 probe size. The secondary electron image, reflection electron image and elmental mapping were taken by EPMA (JEOLJXA8800). Quantitative microanalysis was done by both XSAM and EPMA. The images taken by various methods from the same specimen were compared each other. The results on new boneformationusing the Ti implant and FGM implant were compared. The sliced section corresponding to the part of approximately 15%HAP in Ti/20%HAP FGM was mainly shown in the results. 3JIESULTS Fig.l is the conventional X ray transmission image(so-called Roentogen image) of the implant inserted in the mandible of rabbit The implant of the size 20x7mm is observed between indsor and molar teeth. The implant was inserted to locate the Ti rich part in the upper and HAP rich part in the bottom. Fig.2 shows the growth of new boneformationaround FGM implant These are the images of Ca elemental mappiiig obtained by XSAM after 1(a), 4(b), 8 week implantation(c) respectively. Ti matrix of implant is not seen. There are some weak spots originatedfrt>mHAP particles distributed in Ti matrix in the center area of the figure where the in^)lant should be, but they are generally very weak and often negligible compared to bone area. The thick white area in Fig.2a and b is the cortical bone of mandible. The newly formed immature bone is seen in the right side and lefi: to bottom along cortical bone in Fig.2a. It is formed only partially around the implant and not contact to the implant After 4 week implantation the new bone is surrounding neariy the whole circumference of the implant The new bone becomes more mature than in 1 week. Afl^r 8 weeks the implant is completely surrounded by new bone. Newly formed bone becomes more dense and matured.

Fig.l Roentogen image of the in^)lant inserted in the mandible of rabbit

751

¥ig3 shows the histological observation after 1(a), 4(b), 8 weeks(c) by optical microscopy. The area close to FGM implant in Fig.2 is enlarged to higher magnification. In 1 week new bone is immature and rather difficult to recognize. After 4 weeks the newly fonned waven bone is clearly observed close to the implant but not yet in direct contact After 8 weeks the implant is fiilly surrounded and directly contact with the new bone. The new bone is enough mature and some parts were remodeled to the lameUar bone. Bone marrow is also seen with white contrast Fig.4 shows comparatively the results of observation of new bone formation around FGM implant after 4 weeks by various methods. Fig.4a is the reflection electron image and Figs.4b-e are elemental mappin^p of Ti(b), Ca(c), P(d), 0(e). Figs.4a-e were obtained by EPMA. Fig.4f is the scanning transmission X ray image and Fig.4g is P mapping, both of whidi were taken by XSAM. Fig.^ is the histological image observed by li^t mkroscopy. Reflected electn)ns(Fig.4a) show the compositional image where the r ^ o n composed of relatively heavier atoms exhibits the brighter contrast The Ti matrix of the implant looks the brightest The black dots scattered in the implant represent the HAP particles inside Ti matrix. The upper area in Fig.4a composed of cortical bone shows the second brightest contrast Then the new bone area is recognized with the lower brightness. They were formed along the cortical bone and around the implant with some seperation.

Fig.2 Ca mapping around FGM implant by XSAM after 1(a). 4(b), 8 weeks(c).

200M m

Fig.3 New bonefiMrmationin the region dose to FGM implant after 1(a), 4(b), 8 weeks(c), observed by optical microscopy.

752 The mapping of Ti(Fig.4b) shows the Ti matrix only. The mapping of Ca(c) shows the cortical bone with the brightest contrast and the newly formed bone. The HAP particles in H matrix were not recognized. The mappings of P(d) and 0(e) are mostly the same as Ca mappiag(c) since Ca, P and 0 are the main constituents of HAP. HAP is common to the synthetic HAP and bone, teeth as their principal component The HAP particles in Ti matrix are more sensitively observed in P and 0 mappings than in Ca moping. In 0 mapping(e) the lest of area has gray contrast suggesting the existence of oxygea This is mainly originated from acrylic resin used for embedding the specimen. Scanning transmission X ray image(f) shows the contrast inveise to the reflected electrons(a). The range of contrast in Fig.4f is too wide to show the area with low level of signals and new bone area disappeared in the ^diite contrast The P mapping image by XSAM(g) is essentially the same as that obtained by £PMA(d). The main difference is the poorer resolution in XSAM. Ihe stained specimen observed by transmission light microscopy(h) shows the new bone area with clear contrast Fig.5 shows the comparison of Ca mapping images obtained by £PMA(a) and XSAM(b) for the region around FGM implant after 1 week Because of the poor resolution of 100 /^m used in this study instead of 10//m to shorten the working time, XSAM shows the obscure image compared to EPMA. Fig.6. shows the P mappings of new boneformationafter 8 week implantation for Ti(a) and Ti/HAP FGM implant(b) in comparison. In the case of Ti implant the new boneformationis still partial in the archformaround the implant, wheras in FGM imlant the new bone area is neariy circular surrounding the implant The distribution of HAP particles inside the implant is recognized in FGM(b), which is not in pure Ti(a).

Fig.4 ObservaticHi of new boneformationaround FGM implant after 4 weeks by various methods: reflection electronsCa), mappings of Ti(b), Ca(c), P(d), 0(e), scanning transmission X ray(Of P mapping by XSAM(g), lijg^t microscopy(h).

753 4J)ISCUSSI0N 4.1.0bservation The images of leflection electrons, elemental mappings by EPMA and XSAM, scanning transmission X ray and optical microscopy were applied and compared as shown in Fig.4. Histological study generally needs the good preparation of thinly sectioned specimen and considerable experience to inteipret In mapping method the existence, distribution and quantity of component elements are clearly recognized It is simple and easy to evaluate the growth of new boneformation,forexample, as shown in Fig^. The method is useful for observation of the general view on boneformationand acquisitfon of further information to judge the problem which is difficult to interpret in optical mkroscopy. Quantitative analysis is also possible. Generally elemental mapping has, however, relatively low signal/ooise ratio and limited resolution. Reflection electron image possesses the conqilementary characteristics in this sense. As seen in Fig.4a it has hi^er signal/hoise ratio and the resolution is approximately the same as secondary electrons, about 5nm. The contrast depends on the average atomic number and the compositional image can be obtained. In Fig.4a the dififeience of reflectivity of elections by atomic number elfect gave Ti matrix the brightest contrast and the cortical bone enough bri^tQess. The newly formed bone which is rich in Ca was deariy recognized, surrounding the implant in the dari^ contrast rqspion, which is mainly composed of C, H, 0, N of protain. Scanning transmission X ray image shows the contrast formed by absorption of X ray. The region of the heavier

Fig.5 Comparison of EPMA(a) and XSAM(b) for Ca mapping around FGM implant afler 1 week

Fig.6. Comparison of P mappiqg in Ti(a) and Ti/HAP FGM implant (b) after 8 weeks.

754 atoms exhibits daiker contrast The dependence on atomic number is similar but the contrast is inverse to reflected electrons as seen in Fig.4a and t For both imaging methods of elemental mapping and ieflectk>n dectrons, specimen preparation is much easier since they are non-transparent methods yMch allows the thick specimen ai^licable. The contrast arises from the intrinsic properties of specimen, therefore there needs no stainini^ wheias ultra tbm section is necessity for good work in the conventional method by optical microscopy. In EPMA observation of non- conductive specimen, coating is preferable to avoid electron charging effect but it is not an absolute necessity. In XSAM there is no need of coating. It may enable us to observe the tissue without fixatioa The methods used in this study thus would be useful to derive an additional and conq)lementary information when used together with histological observation by optical microscopy. When the elemental mappings of Ca, P, 0, principal components of HAP and bone, were compared, signalyhoise ratio was the highest in Ca as seen in Fig.4. However the HAP particles in Ti matrix could not be recognized in Ca mapping(Fig.4c). The HAP particles are more sensitively observed in P(Figs.4d, g) and 0 mappings(Fig.4e). The mapping by P is in this sense relatively more sensitive to the low concentration in HAP particle against higher concentration in bone. One reason for this may be that the range of intensity distribution is narrower in P than in Ca and it is possible to show the contrast with the obsa:'vable gray level When the mappings of the same element taken by EPMA and XSAM were compared, the image by XSAM showed the inferior resolution as seen in Fig.5. The resolution of elemental mapping is about 1 //m in EPMA, 10 and 100//m in XSAM. In this study 100 // m probe size was used and the difference of resolution became larger. There may be another contribution other than inferior resolution for the fact that the trabeculae new bone area looks wider in XSAMCower right region in Fig.5b). The suitsice thickness to emit the characteristic X my is sevend tens fim in XSAM, which is much larger than about 1 //m in EPMA. It is also probable that the deeper sur&ce part contributes to widen the mapped area. 4.2.Biocompatibility Various observation used in this study showed that there is observed no inflammation throughout the implants &bricated and inserted in the mandible oi rabbits. Both Ti and Ti/HAP FGM implants have enough biocompatibility. When both implants, inserted respectively in right and left mandibles of the same rabbit, were compared, FGM showed the greater growth of nevAyformedbone area(Fig.6). Ti/HAP FGM implant may have further good biocompatibility.

REFERENCES 1. F.WATARI, A.YOKOYAMA, F.SASO, M.UO, TJCAWASAH, ProcSrd IntSymp.FGM, BJLSCHNER & N.CHERRADI(eds.), Pres.UnivJlonL, Lausanne, 703-708,1995 2. F.WATARI, A.YOKOYAMA, F.SASO, M.UO, TJCAWASAH, 2nd IntConf. Composite Eng., DJIUI(ed), Univ.New Orieans, 801-802,1995 3. F.WATARI, A.YOKOYAMA, F.SASO, M.UO, TJCAWASAH, Composites B6, in press.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

Characteristics of Epoxy-Modified Produced by an Infiltration Process

755

Zirconium

Phosphate

Materials

LM. Lowl, S.Yamaguchi2, A. Nakahira^ and K. Niihara^ iDepartment of Applied Physics, Curtin University of Technology, GPO Box U1987, Perth, 6001 Australia 2The Institute of Scientific and Industrial Research (ISIR), Osaka University, 8-1 Mihogaoka, Ibaraki, Osaka 567, Japan 3Department of Materials Science, Kyoto Institute of Technology, Kyoto, Japan

ABSTRACT An epoxy infiltration process has been used to fabricate zirconium phosphate with much improved physical and mechanical properties for robust applications in various electrochemical devices. Results show that phase composition, microstructure and properties vary gradually along the depth profile, confirming thefimctionally-gradedcharacter of these materials.

1. mTRODUCTION Zirconium phosphates (ZrP) are layered materials [1] with regularly spaced acid groups. The protons of these acid groups are able to diffuse through the interlayer region which enable these materials to be used as ion-exchangers and protonic conductors. In particular, proton conductors with conductivity higher than 10'^ S/cm at the working temperature, can be used in a number of electrochemical devices [2] such asfiielcells, gas sensors, and steam electrolyzers. The electrochemical properties of ZrP are optimal between 20-220°C [3]. These properties are profoundly reduced at temperatures above 200°C due to the loss of water molecule as a result of condensation of monohydrogen phosphate groups to form pyrophosphate [4]. Hence, it is virtually impossible to densify zirconium phosphate with sufficient mechanical integrity for robust applications in various electrochemical devices. In this paper, we report the preliminary results on the synthesis and properties of epoxymodified ZrP. These materials show classical graded characteristics in composition, microstructure, and properties. The microstructure-property relationships in these novel FGMs are discussed.

756 2. EXPERIMENTAL PROCEDURE The materials used for the synthesis of epoxy-modified ZrP were epoxy resin N.N.N'N' tetra glycidylmethaxy diamine (TETRAD-X) and hardener 1,2-cyclo-hexanedicarboxylic anhydride (HHPA) supplied respectively by the Mitsubishi Gas Chemical Company, Japan and the Wako Junyaku Co., Japan. Alpha zirconium phosphate was produced by Dai-ichi Kigenso Kagaku Kogyo Co. Ltd. Green ZrP bars of dimensions 40 x 6 x 6 mm^ were initially pressed in a steel die at a pressure of approximately 75 MPa, followed by isostatic pressing at 100 MPa. Infiltration of pressed ZrP bars was done by immersing them for various times (0-30 mins) in a beaker of mixed epoxy resin and hardener. The ratio of hardener to resin used was 32.3 parts per hundred resin (phr) by weight. Samples were placed in a vacuum chamber (10"^ Torr) for one hour prior to infiltration in order to maximise the kinetics of epoxy diffusion. Infikrated samples were placed in a teflon greased aluminium plate and cured in a ventilated oven at 90°C for 2 hr, followed by 2 hr at 180°C. The particle size and distribution of as received ZrP powder was analysed using a centrifugal laser particle size analyser (Shimadzu Co.). An evacuation method was used to measure the apparent density amd apparent porosity of samples. Phase analysis of ZrP powder and epoxy-modified ZrP samples was performed with a computer-controlled Rigaku diffi*actomter,using the Ni-filtered Cu Ka radiation. The thermal expansion characterisitcs of samples were analysed using a TMA Rigaku Thermoflex 8140 thermoanalyser in air. Alumina was used as reference and the heating rate used was 3*'Cmin"\ Depth profiling of thermal expansion coefficient (a) was evaluated by gradually polishing away the sample surface. Reflection optical microscopy was used to examine the microstructure. The graded composition was observed by coating the surface in gold and viewing in Normarski interference contrast. Three-point bend bars were tested in a Universal Testing Machine (Autograph, Model AglOTC, Shimadzu Co. Ltd) at a cross-head speed of 1 mm/min to determine the flexural strength (a). The elastic modulus (E) of samples was measured using a flexural vibration resonance method at ambient conditions. Depth profiling of a and E was evaluated by gradually polishing away the surface of the sample. Hardness and toughness values of polished samples were measured using a Zwick microhardness tester.

3. RESULTS AND DISCUSSION The as-received ZrP powder showed the wide range of particle size and distribution. The particle size rangedfi-omsubmicron (0.05 |Lim) to 50 |im with a mean value of 1.8 [xm and had surface area of 15.2 m^/g. This wide range of particle size distribution has also been observed by several workers. The x-ray diffraction spectra of pure ZrP powder and the epoxy-modified ZrP samples infiltrated for various times showed [5] the presence of high degree of crystallinity in as-received ZrP by virtue of the sharp reflection lines. The intensity of these lines decreased as the infiltration time or epoxy content increased, resulting in the broadening of peaks and the appearance of a broad but diffuse peak at 17°. The density and porosity (see Table 1) results show that green ZrP samples before infihration were highly porous (ie. 38% porosity) which allowed the infiltration of epoxy resin to fill-up the open pores as time progressed. The denser FGM samples were the result of this

757 pore-filling process which led to a reduction in porosity. An identical pore-filling process was observed for the epoxy-modified YBCO and BSSCO superconductors [6,7]. The kinetics of epoxy infiltration are diffusion controlled and the infiltration depth (d) is expected to vary as the square root of time (t) [8] such that d = (2/7i)(/^.coseoo/2Ti)'^ (ty^

(1)

where R, 9QO, and r\ are respectively average pore radius, equilibrium Contact angle, and viscosity of infiltrant.

Table 1: Time dependent charactersitics of various epoxy-infiltrated ZrP samples Sample EZPO EZPl EZP5 EZP15 EZP30

Infiltration Time (min) 0 1 5 15 30

Density (g/cm^) 1.79 1.84 1.88 2.05 2.07

Porosity (%) 38 30 28 19 17

Epoxy Content (wt%)

Infiltrated Depth (mm)

0 8.7 9.3 14.9 19.6

0 0.50 0.67 0.72 0.97

The presence of infihrated epoxy has imparted significant increases in strength, hardness, elastic modulus, and toughness. The improvement in these properties can be attributed to the pore-filling effect and the concomitant reduction in the overall residual porosity. Porosity has been established by various workers to have a profound influence on hardness, strength, and elastic modulus [10]. This identical strengthening mechanism has also been observed for epoxy-modified HTSC materials [7].

Table 2: Mechanical properties of various epoxy-infiltrated ZrP samples Sample Elastic Modulus Strength Hardness Toughness ^(GPa) g (MPa) H (GPa) Kjc (MPa. Vm) EZPO 3 5 25 0.80 1.05 EZPl 12 15 31 0.80 1.17 EZP5 37 0.83 1.21 EZP15 17 1.26 EZP30 19 40 0.85

The continuous change in concentration of epoxy content across the sample thickness is evident fi-om the XRD results [5] on samples before and after repeated polishing to remove the surface stepwise. As expected, the epoxy content was maximum on the sample surface and decreased gradually towards the centre. This graded composition profile is evident fi*om the microstructure viewed under Normaski interference contrast (Fig. 1) where the bright phase is

758

Fig. 1: Photomicrograph of sample EZP30 viewed under the Nomarski interference contrast. The bright phase is ZrP and the light-dark phase is epoxy resin. The pure epoxy layer is on the far right and the direction of epoxy infiltration isfi"omright to left.

UJ

0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

Sample Depth (mm)

Fig. 2: Depth profile of expansivity for sample EZP30.

0

0.1

0.2

0.3

0.4

0.5

Sample Depth (mm)

Fig. 3: Depth profile of elastic modulus for sample EZP30.

0.6

759 ZrP and the dark phase is epoxy resin. It follows that the graded composition should impart a gradual change in both the physical and mechanical properties. The thermal expansion results (Fig. 2) show that there is a profound influence of graded composition on the overall coefficient of thermal expansion coefficient (a). The decreasing value of a with an increase in sample depth can be attributed to the gradual removal of epoxy resin which has a higher a than that of ZrP. As more and more epoxy was polished away, the value of a approached that of pure ZrP. The depth of penetration can be estimated from the sharp discontinuity in the gradual change of a with depth. Similar observations were obtained for the depth profiling of elastic modulus (Fig. 3), hardness and strength [5].

4. CONCLUSIONS Functionally-graded ZrP materials with improved physical and mechanical properties of ZrP have been produced by epoxy infiltration. These materials show classical fiinctionllygraded characteristics with continuous change in composition, physical and mechanical properties.

ACKNOWLEDGEMENTS IML is very grateful to the Japanese Ministry of Education for funding a Visiting Fellowship. The authors are also grateful to Dai-ichi ICigenso Kagaku Kogyo Co. Ltd for supplying the ZrP powder, and to Dr M Hussain, Mr Suzuki, and Mr Ogawa for technical assistance.

REFERENCES 1. Clearfield, A., Inorganic Ion-Exchange Materials: Clearfield, A., Ed.: CRC Press: Boca Raton, FL. 1982, ppl-74. 2. Alberti, G., Casciola, M., Costantino, U. and Vivani, R., Adv. Mater. 8 (1996) 291. 3. Atik, M., Pawlicka, A., Aegerter, M.A., J. Mater. Sci. Lett. 14 (1995) 1486. 4. Clearfield, A. and Stynes, J.A, J. Inorg Nucl. Chem. 26 (1964) 117. 5. Low, I.M., Yamaguchi, S., Nakahira, A. and Niihara, K. to be submitted to J. Mater. Sci. 6. Low, I.M., Skala, R.D. and Mohazzab, G., J. Mater. Sci. Lett. 13 (1994)1340. 7. Low, IM, Wang, H, and Skala, R.D., J. Mater. Sci. Lett. 14 (1995) 384. 8. Semlak, K.A and Rhines, F.N., Trans. Met. Soc. AIME 212 (1958) 325. 9. Low, I.M. and Lim, F.W., J. Mater. Sci. Lett. 11 (1991) 1119. 10. Dutta, S.K., Mukhopadhyay, A.K. and Chakraborty, D., J. Am. Ceram. Soc. 71 (1988) 942.

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I. Shiota and M.Y Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

761

Preparation and Properties of PVC/Polymethacrylate Graded Blends by a Dissolution - Diffusion method Y. Agari% M. Shimada% A. Ueda% T. Anan\ R. Nomu^a^ Y. Kawasaki^ ' Osaka Municipal Technical Research Institute, 1-6-50, Morinomiya, Joto-ku, Osaka 536 Japan *^ Osaka Institute of Technology, Faculty of Engineering, 5-16-1, Omiya, Asahi-ku, Osaka 535 Japan We prepared and characterized PVC/polymethacrylate graded miscible blends by a dissolution-diffusion method, and found that those have high thermal shock resistance and our prediction model can be adopted to those graded structures. 1. INTRODUCTION Many reports have been published on functionally graded materials which are made of metals and ceramics. These graded materials show characteristics of improved strength against thermal stress, electromagnetic and optical properties. However, there have been few reports on functionally graded polymeric materials^"^ Especially, reports on widely graded and miscible blends are quite few^ although there are literature infomation on homogeneous and miscible blends. In this study, we discussed the graded and miscible blend of polyvinyl chloride(PVC)/ polymethacrylate(polymethyl methacrylate(PMMA) or polyhexyl methacrylate(PHMA)) by a dissolution-diffusion method, and characterized Table 1 Molecular weight of polymers graded structures of the blends by measuring FTIR Polymer Mn Mw spectra and Raman microscopic spectra, and thermal PVC 18200 33500 behaviors around the glass transition 35600 60400 52700 83100 temperature(rg) by DSC method, or by SEM-EDX 98500 55200 61900 124900 observation. Finally, we measured several types of PMMA 93100 59000 mechanical properties and thermal shock resistance PHMA 341100 202800 of the graded polymer blends.

762 2. EXPERIMENTAL 2.1. Materials and sample preparation Atactic PVC, PMMA and PHMA were supplied by Shin-Etsu Chemical Co. Ltd., Mitsubishi Rayon Co. Ltd. and Aldrich Chemical Co. Ltd., respectively. Their average molecular weights were determined from GPC analysis by the multi-angle light scattering method, using tetrahydrofuran(THF) as a solvent, as shown in Table 1. Sample films were prepared by the dissolution-diffusion method, described in the following: Casting of PVC solutions was made on a glass petri dish at r. t. Then, PMA solution containing different concentrations was poured on the cast PVC film in the dish at various temperatures. The weight of PMA was equal to that of PVC. PVC dissolved and diffused in PMA solution, until all solvent was evaporated. Thus, PVC/PMA blend films (about 200 |im) with several types of graded concentration for thickness direction were obtained. 2.2 Measurement Thermal behavior of the blends around their Tg were measured by DSC (Rigaku DSC 8230). The heating rate during the measurement was 10 K/min. Fourier transform infrared spectroscopy (FTIR) of the blend films was performed with a JEOL FTIR spectrometer JIRAQS20 attached to an attenuated total reflection(ATR) apparatus, using a KRS5 prism with an incident angle of 45°. Then, PMMA content was estimated from the ratios of absorbance intensities at 1728 cm"^ (stretching of carbonyl group in PMA) and 615 cm'^ (stretching of C-Cl bond in PVC). The change in PMMA content for a direction was estimated by measuring FTIRATR spectra on a sliced layer of the blend film. Measurement of Raman microscopic spectra was performed by measuring Raman spectra at the focused point, which was shifted stepwisely, 10 by 10 jim, from one surface area to the other one. Tensile properties were measured by using Shimadzu mechanical tester DSS5000. The cross head speed was 2 mm/min. Viscoelastic properties (tensile storage modulus and tan S) were measured by a tensile type dynamic mechanical analysis(SEIKO Instrument DMS 200). The frequency and heating rate of measurement were, respectively 0.2 Hz and 1 K/min. The thermal shock resistance test was performed by moving the specimens from onebox to another (kept at 253 K and 373 K) repeatedly (5 times) 30 min. The specimens were then evaluated on the thermal shock resistance by measuring a maximum angle of warp and adhesive strength in shear by tension loading.

Fig. l Schematic model of dissolution and diffusion system..

763 3. MECHANISM The mechanism of formation of a gradient strucGraded structure 1 ture is described below. After a PMA solution is poured on PVC film in a glass petri dish, PVC begins to dissolve and diffuse Graded structure 2 in the solution to the air side(Fig. 1), but the diffusion is interrupted, when all the solvent evaporates away. Thus, a blend film is produced which consists of a concentration gradient of Graded structure 3 PMA/PVC for the thickness direction. By the steps of dissolution and diffusion of PVC, the graded structures can be classified into 3 types Fig. 2 Schematic models of various (Fig. 2). types of graded structures. 1st type: PVC dissolve initially and then diffuses, but does not yet reach the air sidesurface of PMA solution. The blend has 3 phases(PVC, PMA and thin graded structure). 2nd tvpe: When all PVC have justfinisheddissolving, the diffusion frontier reaches up to the air side surface of PMA solution. The blend has one phase changing the composition gradedly from the surface to the other one, while those surfaces are composed of either PVC or PMA only. 3rd tvpe: After the dissolution and diffusion of PVC have reached up to the air side surface of PMA solution, the PVC and PMA molecules begin to mix with each other to misciblize. The concentration gradient begins to disappear. Formation of concentration gradient should depend on (a) dissolution rate of PVC in PMA solution, (b) diffusion rate of PVC in PMA solution, and (c) interruption time of the diffusion due to completion of solvent evaporation. Factors for controlling the above phenomena are: (1) type of solvent, (2) casting temperature, (3) molecular weight of PVC, and (4) amount of PMMA solution. Until PVC completely dissolves or reaches the surface of PMA solution, i. e., in formation of the first and second types of structure, the diffusion of PVC in the PMA solution is considered to obey Fick's second law(eq. 1), by assuming that the evaporation of the solvent in PMA solution during the diffusion can be neglected. (1) where, CA is concentration of PVC, f is the elapsed time, x is the distance from the surface of PVC sheet, and DAB is the apparent diffusion coefficient.

764 The point where CA becomes 1, shifts to the petri glass side, with carrying forward the dissolution of PVC. Thus, by considering this effect and rearranging mathematically eq. 2 is obtained from eq. 1. (

CA

= ^rfc\

ix-h)

[l^^DAB^t)

1.0 (h

o

jr

O FTIR method # Raman method — Predicted values

0.9 0.8 0.7

(2)

0.6

o u U >

L

i^

0.5 0.4 0.3

where, b is the distance between the petri glass side surface and the other side of remainder of PVC, which has not dissolved yet. Therefore, the gradient profile in the blend at t can be estimated by eq. 2.

0.2 0.1 0.0

1

1

1

1

1

1

0 20 40 60 80 100 120 140 160 180 DISTANCE FROM PETRI GLASS SIDE (^im) Fig. 3 The change in PVC content in thichness direction of PVC/PMMA blend.

4. RESULTS AND DISCUSSION 4.1. Optimum conditions In the case of PVC/PMMA system, we prepared samples by changing the above 4 controllable conditions, and examined the graded structures of those samples by FTIR-ATR, Raman

Graded structure 1

Graded structure 2

300

330 360 390 420 450 TEMPERATURE (K) Fig. 4 DSC curves of several types of PVC/PMMA graded blends.

U I THICKNESS DIRECTION Fig. 5 Chlorine content along thethickness of PVC/ PHMA graded blend( X 750. — ;20^m).

765 microscopic spectroscopy and DSC methods. Fig. 3 shows the graded structure of the samples in the direction of thickness, measured by FTIR-ATR and Raman microscopic spectroscopy methods. It can be confirmed that the blend had a comparatively thick graded structure phase. Then, the graded structure was examined by DSC method. The DSC curve of the blend having widely graded stmciuTt(graded structure 2), shows a more gradual slope around Tg than the others(Fig. 4). Similarly, the structures of the samples in the several types of the conditions were investigated, and then it was found that the typical optimum condition (molecular weight of PVC: Mn=35600, Mw=60400, type of solvent:THF/toluene(5/l), volume of solvent:0.23 ml/ cm^, temperature: 333K) were obtained. In the case of PVC/PHMA system, the graded structures of the samples could not be examined by FTIR-ATR and DSC methods, because PHMA was very soft at r.t. and that DSC curve did not indicate clear change around Tg, Thus, the graded structure was measured by SEM-EDX method(Fig. 5). The chlorine content in the sample increased gradually towards the petri glass side, and then it is considered to have wide graded structure. Further, the structures of the samples in the several types of the conditions wereinvestigated, and then the typical optimum condition (molecular weight of PVC: Mn=356(X), Mw=604(X), type of solvent:MEK, volume of solvent:0.37 ml/cm^ temperature: 313K) could be obtained. 4.2. Application of predicted equation to experimental data Adaptability of eq. 3 to the experimental data was examined. In the blend, which was shown in Fig. 3, experimental data agree approximately with the ones predicted by eq. 2. DAB and b were Table 2 Properties of PVC/PMMA blends PVC/PMMA blend

PVC PMMA

Type 2 D Type 1 P.M.T.2) Tensile properties (kgf/mm2) 6.4 Tensile strength Elongation at break (%) 4.5 Tensile modulus of (kgf/mm2) 200 elasticity DMA properties (Tensile mode) Tg width of storage modulus 20 (K) Half temperature width Tg in tan 8 16 (K) Thermal shock resistance 9 Maximum waip angle (•) Adhesive strength in shear by tension 98 loading (kgO

4.5 3) 2.8 3)

7.2 5.2

5.7 3.9

6.1 3.1

190 3)

220

230

230

8.6,1 P)

11

—

—

—

10

—

—

1704)

—

—

—

714)

—

—

—

1) Blend containing graded structure 2,2) Perfectly miscible blend, 3) Prepared by the hot press method, 4) Blend containing graded structure 1.

766 obtained as 6.38 jirnVsec and 57 iim, respectively. The DAB was greatly larger than Graded structure 1 that of melt diffusion. And it means that this dissolution-diffusion method is very useful. 4.3. Physical properties Physical properties of PVC/PMMA blend containing graded structure 2 (an extremely wide graded concentration), in comparison with those of blend containing graded structure 7(similar to a laminate system), perfectly miscible blend(5/5), PVC only and PMMA only, were summarized in Table 2. The blend containing graded structure 2 was superior to the others in all of the physical properties. DMA properties of PVC/PHMA 200 240 280 320 360 400 blendcontaining graded structure 2, in TEMPERATURE (K) comparison with those of blend containing Fig. 6 Tensile storage modulus and dissipation factor tan 6 several types graded structure 7, perfectly miscible of PVC/PHMA graded blends. blend(5/5) are shown in Fig. 6. Tan 5 of PVC/PHMA blend containing graded structure 2 showed a peak in a greatly wide temperature range. By this phenomenon, the blend is expected to be useful as a damping material in large temperature range. 5. CONCLUSION We prepared PVC/polymethacrylate graded miscible blends by a dissolution-diffusion method, and then found the typical optimum condition for obtaining a wide graded structur, by FTIR-ATR, Raman, DSC and SEM-EDX methods. Further, those blends have high thermal shock resistance and our prediction model can be adopted to those graded structures. REFERENCES 1. M. Kryszewski and G. Czeremuszkin, Plaste u kautchuk, 11 (1980) 605 2. Y. Koike, H. Hidaka and Y. Ohtsuka, Appl. Opt., 22(1983) 413 3. S. Asai, the 6th Symposium of Graded Functionally Materials (1993) 61 4. E. Jabbari and N. A. Peppas, Macromolecules, 26 (1993) 2175 5. K. C. Farinas, L. Doh, S. Venkatraman and R. O. Potts, Macromolecules, 27(1994) 5220 6. Y. Agari, M. Shimada, A. Ueda, S. Nagai, Macromol. Chem. Phys., 197(1996) 2017

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

767

Preparation and properties of polyimide/Cu functionally graded material M. Omori, A. Okubo, G. H. Kang and T. Hirai Institute for Materials Research, Tohoku University, 2-1-1 Katahira, Aoba-ku, Sendai 980-77, Japan A polymer/Cu functionally graded material (FGM) was prepared from thermosetting polyimide and Cu powders using a spark plasma system. These powders were consolidated at different temperatures. A FGM of Cu and polyimide was prepared under a graded temperature using a graphite die, the diameters of which were different at the top and bottom. A dense polyimide/Cu FGM was formed under the graded temperature.There were no cracks in the FGM prepared by stacking three graded layers. Thermal properties of the thin FGM were measured. 1.

INTRODUCTION

Research on a functionally graded material (FGM) has been focussed on the combination of ceramics and metals, because thermal barriers are required for rockets, space shuttles, and space planes. Ceramics/metal FGMs have been produced by various kinds of methods, e.g., CVD, PVD, electrodeposition, plasma spraying, and powder metallurgy [1]. Plasma spraying is a good method for forming ceramics/metal FGMs for practical use [2, 3], but there are substantial pores inside the sprayed layers. A dense ceramics/metal FGM has not been prepared by normal powder metallurgy, because the sintering materials generally have different sintering temperatures. A spark plasma system consists of combination of a hot-press and a plasma generator [4]. This system enables formation of a dense FGM [5, 6]. The residual stress is not completely relaxed by the graded layers, and causes cracks in the dense FGM under conditions of shock. The residual stress of a dense polyimide /Al FGM can be relaxed so as not to cause cracking [8]. Polymer/metal FGMs are practical and valuable as industrial materials. Polyimide, the thermal conductivity of which is low, is used for circuit boards. However, circuit boards require high thermal conductivity, recently. The combination of polyimide and Cu must improve thermal conductivity. The polyimide of the polyimide/AI FGM is thermoplastic. The sintering temperature of Cu is higher than that of Al. A thermoplastic cannot be used for the formation of a polyimide/Cu FGM, as a Cu layer could not be sintered densely. A thermosetting polyimide was selected for that FGM.

768 2. EXPERIMENTAL PROCEDURES The materials were Cu (Nippon Atomized Powders Co. Ltd., 5 |Lim) and polyimide (Kaneka Co. Ltd.) powders. Polyimide and Cu powders in a graphite die were first consolidated independently using spark plasma 40 system (Sumitomo Coal Mining Co. Ltd., SPS•^ •• 1050) in a vacuum. Polyimide and Cu powders A were consolidated by the spark plasma system 20 1 under pressures of 9.8 and 39.8 MPa using an ordinary die. Then those powders were mixed using an agate mortar with a pestle for 20 min. / The mixed powder was layered in a graphite die 20 as shown in Flg.1. The die was set in the system. The sintering temperature was controlled with a 1 1 20 1 thermocouple installed in a hole with a large r^ H diameter. 70 IM — • The consolidated sample was polished and its properties were measured. Density was Fig. 1 Graphite die (mm). determined by Archimedes' principle using water Immersion. Elastic properties of the bulk, 5 - 1 0 mm in length and 20 mm in diameter, were determined with a pulse-echo overlap ultrasonic technique using an ultrasonic detector ( Hitachi Kenki Co. Ltd., ATS-100) and a storage oscilloscope (Iwasaki Tsushinki Co. Ltd., DS6411). The thermal expansion of specimens ( 3 x 4 x 1 5 mm3) was measured by a differencial dilatometer (Mac Science, TD5200) using an alumina bar as a standard from 20 to 400°C under a heating and cooling rate of 20°C/min in N2 gas. The polished surface was observed using an optical microscope (Olympus, SZH and VANOX).

vt H

3. RESULTS AND DISCUSSION The spark plasma system is advantageous to form FGMs from various kinds of powders, because the graphite die is heated directly and powders are activated by plasma. The direct heating affords graded sintering temperatures using the die in Fig. 1. The plasma of this system removes adsorbed gas on the surface of powders. Metals which are easily oxidized like Al and Mg are activated by the plasma to be sintered. Effects of the plasma on polymers have not been studied, since that system was thought to sinter metal and ceramics powders. First, the consolidation of polyimide and Cu was investigated to determine consolidation temperatures and pressures for the FGM formation. The most suitable consolidation temperatures are shown in Table 1. Densities, Young's modulus and thermal expansion coefficients are also listed in Table 1. Suitable consolidation temperatures depended on the pressure. Sintering was conducted and shrinkage of the samples were observed. The sintering temperature of Cu powders was at the point where shrinkage stopped. The density and modulus of Cu bodies were similar

769 from 9.8 to 39.2 MPa, even though the sintering temperature was not the same. The sintering temperature of Cu at 39.2 MPa was 520°C. The same sintered Cu can be obtained at 630°C and Table 1 Consolidation of polyimide and Cu powders 39.2 MPa. On the other Cu polyimide hand, consolidation of the polyimide differed with 39.2 Pressure (MPa) 9.8 39.2 9.8 pressures. The polyimide, Consolidation temp. (^C) 520 630 215 380 which was consolidated at 7.74 7.68 1.34 1.06 39.2 MPa and 380^C, was Bulk density (g/cm^) Poisson's ratio 0.356 0.353 0.359 0.305 partially carbonized and 100 101 3.17 1.45 brittle. This carbonization Young's modulus (GPa) Thermal expansion is thought to be due to the plasma. Polyimide is heatcoefficient (10"%) 52 51 resistant over 400°C, and cannot be carbonized at that temperature The consolidation at 39.2 MPa was 215°C. The energy of 60 500 plasma may be 50 greater at high CO CL pressures than at low .4 40 pressures, and the Q_ consolidation at A 30 380°C partially 20 (0 deteriorated the 0 W Temperature polyimide at high 10 O Pressure pressures. The I I i I I I I density and moduli of 0 10 15 20 25 30 the polyimide Time, t/min consolidated at 380°C and 9.8 MPa Fig. 2 Temperature and pressure of FGM formation. were lower than those of the polyimide consolidated at 39.2 MPa and 215X. It is impossible to obtain a dense polyimide and copper simultaneously at 380°C and 39.2 Table 2 Consolidation of polyimide (temp.: MPa, and at 215^C and 9.8 MPa. 390^0, pressure: 9.8 - 39.2 MPa) Figure 1 shows the die that was Property Value designed for the polyimide/Cu FGM. The temperature was measured at a Bulk density (g/cm^) 1.31 hole with a large diameter. Pressure Poisson's ratio 0.301 was oriented from an upper graphite Young's modulus (GPa) 3.64 rod to a lower rod. The temperature of Thermal expansion the upper rod was the highest, because the electrical resistivity there coefficient (IQ-^/K) 50 was the highest. Polyimide was

770 stacked below and copper was stacked above. The temperature difference, which was taken by thermocouples, was more 130°C between sintering locations of polyimide and Cu. As the temperature difference was as much as 250^C at 9.8 MPa and as much as 305X at 39.8 MPa, it was not adequate to form a dense FGM. The

Fig. 3 Joined polyimide - Cu.

Fig. 5 FGM of two layers.

Fig. 4 FGM of one layer. ^^'S- ^ ^^^ ^f three layers, stacked polyimide was sintered at first near 390°C and 9.8 MPa. Then the pressure was increased from 9.8 to 39.2 MPa, with the temperature being held constant to sinter Cu. The consolidation temperature and pressure schemes of the FGM formation are shown in Fig. 2. The polyimide consolidated by these schemes was not carbonized, and its Fig. 7 Thin FGM. properties were similar to those of 39.2 MPa and 215°C, as shown in Table 2. The polyimide and Cu powders were put in the die, and joined without graded layers. The thickness of the polyimide and Cu layers was 5 mm. A cross section is shown in Fig. 3. The interface of polyimide and Cu was cracked by the residual stress coming from the thermal expansion mismatch of polyimide and Cu. A graded layer (0.5 mm) of 50 vol% polyimide and 50 vol% Cu was inserted between the polyimide layer ( 3.5 mm) and the Cu layer (3.5 mm). The photograph of a FGM with

771 a graded layer is shown in Fig. 4. A crack was observed at the boundary of the Cu and graded layers. A FGM consisting of two graded layers, which were prepared from two mixed powders of 25% polyimide - 75% Cu and 75% Cu - 25% polyimide, is shown in Fig. 5. There was a crack between the Cu and 25% polyimide - 75% Cu layers. Three graded layers, namely, 25% polyimide - 75% Cu, 50% polyimide - 50% Cu and 75% ployimide - 25% Cu, did not cause cracking, as shown in Fig. 6. The residual stress was relaxed by three graded layers. The FGM with three layers was stable and able to be cut into small fragments. A thin FGM with three graded layers was prepared as shown in Fig. 7. The thickness of each layer was 0.5 mm. The thermal expansion difference between polyimide and Cu (17x10"^/K) is about 30x10'^/K, and one graded layer must be charged with the difference of 10x10"^/K. The graded layer cannot relax such a large difference [5, 6, 7]. The stress is dispersed over the graded layers. The dispersed stress is relaxed as follows: if each graded layer has sufficient strength to resist cracking by the stress, and stresses are relaxed by plastic deformation of metals and polymers. A ceramics/metal FGM is cracked by cutting [5, 6]. Plastic deformation does not occur in ceramics. As the formation temperature of ceramics/metal FGMs Is high, a high level of stress remains. This stress is too great to be relaxed by the plastic deformation of metal only. Thermal properties of the thin FGM are shown in Table 3. The thermal conductivity of the FGM of 2.1 mm thickness was 5.69 W/m-K at room temperature, larger than that of polyimide Table 3 Thermal properties of thin FGM (0.128 W/m-K [8]). Temperature Thermal diffusivity Specific heat Thermal conductivity This FGM (cm^/s) (W/m-K) (J/g-K) has greater 1.47 25 0.886 5.69 thermal 1.22 0.913 4.87 143 conductivity 1.12 255 1.05 5.11 than the 0.737 2.94 original 355 0.915 polyimide. The thermal conductivities of alumina and Cu are respectively 21 and 406 W/m-K. That value of the FGM is not as high as that of alumina. Aluminum nitride is more conductive ( 169 W/m-K) than alumina. Thus, a composite of alumina and aluminum nitride would probably improve the thermal conductivity.

eo

4. CONCLUSION A polyimide/Cu FGM was prepared by stacking different compositions of Cu and thermosetting polyimide powders using a spark plasma system. The consolidation temperature was about 520^C for the Cu layer and 390°C for the polyimide layer. One graded layer or two graded layers resulted in cracks. There were no cracks in the FGM prepared by stacking three graded layers. The thermal conductivity of thin

772 FGM was 5.69 W/m-K. ACKNOWLEDGMENT We acknowledge the support of the Ministry of Education, Science and Culture under Grant-in Aid for Scientific Research on Priority Areas (No. 08243102, "Physics and Chemistry of Functionally Graded Materials"). We appreciate being permitted to use a spark plasma system the Laboratory for Developmental Research of Advanced Materials, Institute for Materials Research, Tohoku University.

REFERENCES 1. T. Hirai, Materials Science and Technology, Processing of Ceramics, Part 2, VHC Verlagssellschaft mblH, Germany, 1996, pp. 293. 2. H. D. Steffens, M. Dvorak and M. Wewel, Proc. 1st Int. Symp. on Functionally Gradient Materials, Sendai (1990), pp. 139. 3. T. Fukushima, S. Kuroda and S. Kitahara, Proc. 1st Int. Symp. on Functionally Gradient Materials, Sendai (1990), pp. 145. 4. M. Ishiyama, Proc. 1993 Powder Metall. World Congress, Kyoto (1993), pp. 931. 5. M. Omori, H. Sakai, A. Okubo and T. Hirai, Proc. 3rd Int. Symp. on Structural and Functional Gradient Materials, Lausanne (1994), pp. 65. 6. M. Omori, H. Sakai, T. Hirai, M. Kawahara and M. Tokita, Proc. 3rd Int. Symp. on Structural and Functional Gradient Materials, Lausanne (1994), pp. 71. 7. M. Omori, H. Sakai, A. Okubo and T. Hirai, Proc. 3rd Int. Symp. on Structural and Functional Gradient Materials, Lausanne (1994), pp. 667. 8. Lin Li and D. D. L. Chung, J. Electro. Mater., 23 (1994) 557.

I. Shiota and M.Y. Miyamoto (Editors) Functionally Graded Materials 1996 ® 1997 Elsevier Science B.V. All rights reserved.

773

Smart Functionally Graded Material without Bending Deformation J. Qiu^), J. Tani^) and T. Soga^)

1) Institute of Fluid Science, Tohoku University, Japan 2) Graduate School, Tohoku University, Japan

This paper describes a method to reduce the discontinuity of thermal stress in a smart FGM for high temperature. The original smart FGM consists of two layers: one layer of FGM base material and one layer of piezoelectric material (PZT) used as an actuator to reduce the thermal bending deformation. Large stress discontinuity is induced on the interface of the two layers due to the discontinuity of material properties when high voltage is apphed to piezoelectric layer. In this study, the stress discontinuity is reduced by introducing a graded piezoelectric layer between the FGM layer and PZT layer. Furthermore, an additional layer is introduced on the low temperature side to measure the bending deformation of the composite FGM so that the deformation induced by time-varying temperature can be actively cancelled by using feedback control.

INTRODUCTION In recent years, many functionally graded materials (FGM) were developed to bear high temperature, which have potential use on the surface of spacecraft and the first wall of fusion reactors[1,2]. Due to the large temperature difference between the two sides of a FGM plate, the thermal bending deformation is not negligible and hoped to be cancelled. On the other hand, in precision optical elements and machine parts such as reflectors and guidelines, even the small bending deformation due to slight variation of temperature or external forces may have influence on performances and some methods have been developed to suppress this kind of deformation[3,4]. As most structures in engineering are distributed parameter systems, the idea of distributed sensors and actuators has been used in vibration control of structural systems [5]. In this paper, the idea of distributed actuator is used in the control of the static bending deformation. In order to reduce or cancel the bending deformation, a two-layer material configuration was adopted for the smart FGM in the former study[6]: one layer of FGM main material and one layer of piezoelectric material(PZT) used as an actuator. Due to the discontinuity of material properties large stress discontinuity is induced on the interface of the two layers when large voltage is applied to the piezoelectric layer. The stress discontinuity is the direct cause of delamination and fracture of materials. The purpose of this paper is to design a new configuration of FGM which will can reduce the stress discontinuity.

774 Main material(FunctionaIly Gradient Material) The gradient layer of PZT and metal

L=100min /ij=10imn PZT(Piezoelectric Material)

PZT(Piezoelectric Material)

(a) Model 1

(b) Model 2 . Y

PZT(Piezoelectric Mateiial)

Main material(Functionally Gradient Materials) The gradient layer of PZT and metal

(c) Model 3 Fig.l. Configuration of smart FGM When temperature varies with time, active control is needed to cancel the thermal bending deformation. A sensor layer is introduced on the low temperature side and the simulation results show that the thermal bending deformation can be cancelled even if the temperature varies. MODAL AND ANALYSIS Model of Smart FGM The main purpose of the smart FGM discussed in the study is to cancel the thermal bending deformation induced by temperature graded in the material. However the stress distribution resulted from temperature and piezoelectric effect is also very important to ensure that the maximum stress is under the material strength and no plastic deformation will appear within the material. In the former study [6], a piezoelectric actuator was pasted on the low temperature side and used to cancel the thermal bending deformation as shown in Fig. 1(a). Due to the discontinuity of material parameters, large stress discontinuity which is the direct cause of delamination is induced on the interface of the two layers. In order to reduce the stress discontinuity, a three-layer configuration is adopted in the new smart FGM: a main FGM layer, a graded piezoelectric layer and a piezoelectric layer as shown in Fig. 1(b). Due to the graded piezoelectric layer introduced between the main FGM layer and the piezoelectric layer, the material properties of the new smart FGM varies continuously in the thickness direction. In this paper, the finite element method is used to analyze the variation of deformation and stress distribution when different voltage is applied to the actuator and to verify the effectiveness of these models in the reduction of thermal bending deformation and stress discontinuity. In order to actively control the thermal bending deformation induced by time-varying temperature Ti, the bending deformation needs to be measured. In model 3, a sensor layer with two strain gauges at position A and B is bounded to the low temperature side

775 of model 2 as shown in Fig. 1(c). Temperature Distribution First the temperature distribution is solved. Steady-state heat conduction satisfies the following equation:

Because the symmetry of the model relative to y axis, only half of it should be considered. Therefore the boundary conditions are: T = Ti on y = -hp (2) T = T2 dT ^— = 0 ox

on y = hb

(3)

on X = 0 or X =: //2

(4)

After discretization by using finite element method, the following system of algebraic equations can be obtained:

IG]{T} ^ {g}

(5)

When temperature difference between the two sides is large, heat conductivity A is a function of temperature. Therefore stiffness matrix [G] is a function of {T}, and Eq.(5) is nonlinear. Stress Analysis The stress distribution is a three dimensional problem. However for simplicity, only the two dimensional case is considered. The stress-strain relation can be written in the following form: W

= [Z)]({e}-{eo})

(6)

where

I

^x

I

I

'^x

I

^Xll

I

I

IXV

I

771

1 V V I 0

0

0 0 l-v

By using virtual work method, we can obtain a system of algebraic equations:

m{u)+{/}

=0

(7)

where [K] — stiffness matrix {u} — nodal displacement vector {/} — equivalent nodal load vector and in this case:

{f} = -Y.j i

}BY\D]{^^W

(8)

i

where [B] is a matrix consisting of shape functions and the initial strain eo can be expressed in the following form:

776

Uo}

aT + dssEx aT + d^iEx 0

(9)

The above expression is valid for all the layers. For the FGM layer, d^i and d^z equals to zero; for the PZT layer, they are constants; and for the graded PZT layer, they are functions of coordinates. Here it is assumed that electric field is applied in x direction and its intensity is E^MATERIAL PROPERTIES OF DIFFERENT LAYERS The material constitution and properties of the main FGM layer, graded piezoelectric layer and the piezoelectric layer in the smart FGM are discussed in this section. Main FGM Layer The main FGM layer considered in this study is Zr20-(Ti-6Al-4V) FGM[1]. Properties of Zr02 in the temperature range of 0°C < T < 1100°C are: f A = 1.71 + 0.21 X 10-^r + 0.116 X lO-^T^ ^ = 132.2 - 50.3 X 10-^T - 8.1 X 10-^ a = 13.31 X 10-^ - 18.9 x 10"^ + 12.7 x lO'^^ i/ = l/3 GB = 148.1 + 1.184 X lO-^T - 31.4 x lO'^T^

(10)

The properties of Ti-6A1-4V in the temperature 0°C < T < 1100°C are: A = 1.1 X 0.017T E = 122.7 - 0.0565r r 7.43 X 10-6 _^ 5 55 ^ ^Q-gj^ _ 2 59 a=l X 10-122^2 (o°C < T < 830°C) [ 11.4 X 10-6 (830°C
(11)

where A(W/niK), E{GPa), a(K~'), u and aB{MPa) are heat conductivity, Young's modulus, thermal expansion factor, Poisson's ration and allowable stress, respectively. It is supposed that these two components are mixed in an ideal way, and volume ratios of metal Ti-6A1-4V and pore are: F„ = (1 -

(12) (13)

yr

P^Ay{l-yr Table 1 Piezoelectric material properties Piezoelectric constant dsi Piezoelectric constant d^s Young's modulus E Poisson's ratio 1/ Curie point Tc

-190 X 10-1^ m/v 390 X 10-12 ^/^ 62.0 GPa 0.3 350°C

777 In the calculation, m, n and A are 2.0, 1.5, 3.0, respectively. The distribution of V and P are shown in Fig. 2. The material properties of the FGM made from the mixture of the two components can be calculated in the way as shown in the reference[l]. o

Piezoelectric Layer The material properties of piezoelectric layer are shown in Table 1. Since the temperature variation in the piezoelectric layer is not so large, it is assumed that the properties are independent of temperature.

>

Normalized thickness y/h Graded Piezoelectric Layer Fig.2. Volume ratio of metal and pore The graded PZT layer is composed of two components: metal Ti-6A1-4V which is identical to that of the adjacent FGM layer and PZT which is the same as the adjacent PZT actuator layer. It is assumed that the two components are mixed linearly with metal on the upper surface and PZT on the lower surface. The material properties of the graded FGM layer are calculated in the same way as those of FGM layer[1]. The piezoelectric constants dsi and ^33 are supposed to vary linearly with the volume ratio of PZT and equal to 0 when the volume ratio of PZT is less than 50%. The variation of piezoelectric constants d^i is the thickness direction is shown in Fig.3 N U M E R I C A L RESULTS A N D DISCUSSION R e d u c t i o n of Stress Discontinuity In the numerical simulation, it is assumed that the surface y = hb is given a uniform temperature of 1000°C and on PZT side temperature is kept zero, that is Ti = 1000°C

Oh

modell 1000

u

modell

-0.5

1

y X

model2

^

800

2

600

/

model2 o X

-1.5 -2 1 -1.5

1

-1

1

1

-0.5 0 y [mm]

.

0.5

Fig.3. Piezoelectric constant d^i

/

B H

400

y^

200 0

A

- 4 - 2 0 2 4 6 8 y [mm] Fig.4. Temperature distribution

10

778

model 1

^

1

>. 6

\

> o 3

TD

™™~»™»^ mociel2

o

model 1

\ ,

4

model2

\

c o U

\

\ ^

M 2

y

X

n -

4

-

2

0

2

4

6

8

10

0

2

y [mm]

4

6

8

10

y [mm]

Fig. 5. Heat conductivity

Fig. 6. Thermal expansion factor 0.34

0.33 O 0.32

I

0.31 h

^ -

4

-

2

0

2

4

0.29

6

0

2

4

6

y [mm] Fig.8. Poisson's ratio

y [mm]

Fig.7. Young's modulus 1000 I 800

I \

600

-1 ^

E=2.1X10-'[V/cm]

0

E=1.5X10^[V/cm]

0

E=0.0X10^[V/cm]

-200

0.2 0.4 0.6 Normalized length x/L

E=2.1X10^[V/cm] E=1.5X10^[V/cm] E=0.0X10^[V/cm]

L

I 200

6

J -3

400

allowable stress

— — \

s

—

0.8

Fig.9. Deformation of model 1

-400 '

. . . . . . ^ - • • • • • • • • • • • • • • • • ' • "

2

,/ ^^^"^--s^ 0

2 4 y [mm]

6

10

Fig. 10. Stress distribution of model 1

779 lUUW

— allowable stress 800

k

600 'o?

•

g

—E=2.2X10^[V/cm] \

—E=1.5X10^[V/cm] \

E=0.0X10'[V/cm]

400

1 200 E=2.2X10-'[V/cm] E=1.5X10^[V/cm]

0

E=0.0X10^[V/cm]

-200

1v s X ^

•

=rJ

X''

Ar\f\

0

0.2 0.4 0.6 0.8 Normalized length x/L

Fig.ll. Deformation model 2

-

4

-

2

0

2 4 6 y [mm]

10

Fig. 12. Stress distribution of model 2

and T2 = 0°C. The distributions of temperature, heat conductivity and thermal expansion factor of model 1 and 2 m y direction are shown in Fig. 4, 5 and 6, respectively. The distributions of Young's modulus and Poisson's ratio are shown in Fig. 7 and 8, respectively. In order to reduce the bending deformation, electric field is applied to PZT actuator. Figure 9 shows the variation of bending deformation when different electric fields are applied. When no electric field is applied, the maximum displacement at the end of plate is about 3.3mm. The maximum displacement is reduced to 1.0mm when Ey = 1.5 X 10^(V/cm) and it is almost completely canceled when Ey = 2.2 x 10^(V/cm). Figure 10 illustrates the variation of stress distribution of model 1 along y axis. The stress increases when electric field increases. The tensile stress in FGM is under the material strength when electric field is as high as 2.2 x 10^(V/cm). The stress discontinuity on the interface of the two layers is about 437MPa. In model 2, the bending deformation and maximum stresses are almost the same as those of model 1, but the stress discontinuity is reduced to 325MPa as shown in Fig.ll and 12. Sensor and A c t i v e Control In model 3, a sensor layer with two strain gauges are bounded to measure the bending deformation. The difference of strains between the two strain gauges varies which temperature as shown in Fig. 13. A bridge of resistances is constructed with the two strain gauges as two of the arms and the output voltage can be used as the feedback signal which is used to calculate control voltage of the piezoelectric actuator. Figure 13 also shows that with active control the bending deformation can be cancelled successfully even if the temperature varies.

0.9 0.8 „

0.7

6

0.6

^

0.5

. before control - after control

•§ 0.4

2 ^

0.3 0.2 0.1 0 -0.1 I

200

400

600

800

1000

Temperature (T, ) Fig. 13. Strain difference

1200

780 CONCLUSION From the above numerical results the following conclusion can be obtained: (1) The thermal bending deformation can be cancelled by controlling the piezoelectric actuator pasted on the low temperature side. The stress discontinuity induced by the discontinuity of material parameters on the interface can be reduced by adding a graded piezoelectric layer between the main material and piezoelectric actuator. (2) By adding a sensing layer on the low temperature side, the bending deformation can be measured. The bending deformation can be actively cancelled by using feedback control. REFERENCES [1] N. Noda and T. Tsuji, Steady Thermal Stresses in a Plate of Functionally Graded Material with Temperature-Dependent Properties, Transactions of the Japan Society of Mechanical Engineers, Vol. 57A(535), 625-631, 1991. [2] J. Tani, Present State in Research of Smart Composite Structures, Proceeding of 4th Symposium on Dynamics Related to Electromagnetic Force, Kanazawa, Japan, 309-312, 1992. [3] P.E. Barbone, and A.M.M. Braga, Influence of Electrode Size on the Active Suppression of Sound Reflection from Submerged Plates Using Distributed Piezoelectric Actuators, Proceeding of the First European Conference on Intelligent Materials, ed. B.Culshaw et al, Glasgow, England, 325-328, 1992. [4] W. Charon, Smart Structures with Piezopolymers for Space AppUcation, ibid, 329332. [5] H.S. Tzou, and M. Gadre, Theoretical Analysis of a Multi-Layered Thin Shell Coupled with Piezoelectric Shell Actuators for Distributed Vibration Controls, Journal of Sound and Vibration, Vol. 132(3), 433-450, 1989. [6] J. Qiu, J. Tani and T. Takagi, An inteUigent Piezoelectric Composite Material without Bending Deformation, Journal of Technical Physics, 35(1-2), 99-107, 1994.

781

AUTHOR INDEX Abe,N. Aboudi, J. Admon, U. Agari, Y. Akama, S. Altenburg, H. Amada, S. Amano, T. An,G. Anan,T. Anatychuk, L. I. Aoki, T. Aoki, Y. Arnold, S. M. Aruga, A. Atkinson, J. F.

551 113 397 761 451 331 521, 731 633 53, 179 761 501 257 719 113 605,611 687

Fu, Z. Y. Fujii, K. Fujii, T. Fukuda, H. Fukuda, R. Fukui, Y. Fukushima, T. Furuhashi, H.

295 439 617 257 647, 655, 661 713 59, 245, 661 337

Gasik, M. M. Ge,C. C. Getto, H. Glaeser, A. M. Goto, T. Greil,P. Guo, Q. L.

21,313 35, 65, 301 743 325 557 173 707

Bahl, 0. P. Banovic, S. W. Barmak, K. Bhatt, D. P. Bilotsky, Y. D. Birth, U. Borchert, R. Borovinskaya, I. P. Brown, I. G. Bruck, H. A.

93 227 227 93 21 167 349 283, 289 687 387

Hamagami, J. Hamamura, 0 . Harmer, M. P. Hatta, H. Hayashi, N. Heikinheimo, L. Henne, R. Higa, S. Hilmas, G. E. Hino, H. Hirai,T. Hirai, Y. Hirano, T. Hoffman, R. A. Hojo, J. Hong, C. W. Hongo, F. Hosomi, S. Hua,J. S.

Cao, W. B. Chan, H. M. Chen, L. D. Cheng, H. Cherradi, N. Chiba, A. Choi, S. C. Choules, B. D. Clarke, D. R. Dariel, M. P. Desplat, J-L. Delfosse, D. Dykes, D.

65, 301 227 557, 569, 587 53, 179 379 191 185 149 387 397 639 379 343, 373

Endo, T. Erdogan, F.

701 105

Filla, B. J. Foumie, J. F. Frage, N. Friedersdorf, L. E.

425 331 397 227

Igarashi, T. Ikegaya, A. Ilschner, B. Imai, Y. Imamura, K. Imura, T. Ishida, K. Ishihara, S. Ishikura, T. Ishizuka, T. Isoda, Y. Isogai, K. Itoh, I.

337, 343,

337, 343,

445, 557, 569, 483,

29, 283,

575, 581,

245, 343, 123,

221 463 227 257 373 313 563 373 319 413 767 681 489 319 419 173 419 289 307 655 361 15 617 191 343 463 743 373 131 617 661 409

782 Iwasaki, K. Iwase, M.

251 681

Jedamzik, R. Jin, M. J. Joensson, M. Jung, Y. G.

233 295 167 185

Kaibe,H. T. Kajikawa, T. Kamiya, N. Kando, M. Kaneta, H. Kang, C. S. Kang, G. H. Kang,Y. S. Kanzawa,T. Kasuga, Y. Kathuria,Y. P. Kato, K. Kato, M. Kato, T. Katoh, K. Katoh, M. Kawabata, J. Kawamata, Y. Kawamura, H. Kawamura, M. Kawasaki, A. Kawasaki, T. Kawasaki, Y. Kaysser, W. A. Kido, H. Kieback, B. Kim, J. H. Kim, Y. W. Kimura, H. Kirihara, S. Kisara, K. Kishi, K. Kishimoto, K. Kitahara, S. Kitayama, M. Kobayashi, H. Kohri, H. Koizumi, M. Kojima, T. Kojima, Y. Kokini, K. Koser, O. Koshigoe, M. Koyanagi, T. Krell,T.

509, 575, 617 475, 627 633 673 361 131 767 99, 569, 587 409 647 337 605 661 661 647 655 695 469 215 413 143, 515, 527 749 761 263, 563 599 167 545 373 355 197 99, 587 419 623 245 325 209 551 1, 283, 289 509 409 149 433 581 623 263

Kude, Y. Kumakawa, A. Kuroda, S. Kuroda, Y.

239, 463 425 59, 245 469

Lai, W. Lai, Z. H. Lange, R. Langer, G. Lee, C. H. Lee, D. M. Lee, G. G. Lee, S. Y. Leushake, U. Levashov, E. A Li, J. F. Li, Z. S. Lin, J. S. Liu, T. Q. Lombardi, J. L. Low, I. M. Lugscheider, E.

275 203, 269 433 563 539, 545 539, 545 527 203 263 283, 289 445 707 599 35 319 367, 755 563

Marder, A. R. Marple, B. R. Masuda, T. Matsubara, K. Matsuzaki, T. Matsuzaki, Y. Matubara, M. McCoy, B. J. Merzhanov, A. G. Meyer, H. Minagawa, H. Miura, T. Miyagawa, T. Miyajima, M. Miyakawa,T. Miyamoto, Y. Mizuguchi, F. Mochimam, T. Moeckli, P. Moller, P. J. Mook, G. Morimoto, J. Morimoto, K. Morizono, Y. Moro, A. Mu, Z. C. Muller, E. Muller, F. Munir, Z. A. Munz, D.

227 159 593 623 463 413 409 275 295 93 695 155 361 515, 527 605 1, 599 81 575 379 707 433 605,611 661 191 99, 469, 569 65 563 173 275 41

783

Murakami, Y.

343

Nagamoto, Y. Nagao, J. Nakahira, A. Nakamura, Y. Nakano, J. Nakata, H. Naramoto, H. Neubrand, A. Niihara, K. Niino, M. Niiyama, M. Nishida, I. A.

623 695 755 713 439 409 719 9,233 755 1, 99, 587 251 509, 545, 551 575, 581, 617 215 191 489 725 569, 587 737 575, 593 761 137

Nishida, K. Nishida, M. Nishio, Y. Noda,T. Noda, Y. Nogata, F. Noguchi, T. Nomura, R. Nomura, S. O'Comior, B. H. Ogura, H. Ohashi, A. Ohkawa,S. Ohkoshi, T. Ohnabe, H. Ohsugi, I. J. Ohta,T. Ohta, Y. Ohyanagi, M. Okamoto, M. Okamoto, Y. Okamura, H. Okazaki, S. Okubo, A. Okumura, T. Omori, M. Ono, F. Orihashi, M. Otoguro, Y. Otooni, M, A. Otsuka, A.

367 191 667 749 617 81 509 533 123 283, 289 617 605,611 515 413 767 575 767 463 569 409 687 251

Park, K. Peters, M. Petronis, C. M. Pettermami, H. E. Phelps, J. M.

539, 545 263 227 75 425

Pindera, M. J. Pityulin, A. N. Powers, J. D. Pratapa, S. Puerta, D. G.

113 295 325 367 227

Qiu, J.

773

Rabin, B. H. Raveh, A. Risbud, S. H. Rodel, J. Saito, S. Saito, T. Sakaguchi, S. Sakamoto, N. Sakata, M. Sano, M. Sarkar, P. Saso, F. Sato, N. Satoh,T. Schiller, G. Schilz, J. Schulz, U. Schumacher, R. Seega, C. Senda, T. Seo, J. H. Sheahen, D. M. Shen, Q. Shibahara, Y. Shida, K. Shimada, M. Shimizu, N. Shimoda, N. Shimokawa, K. Shinagawa, K. Shinohara, Y. Shiota, I. Siren, M. Slifka, A. J. Soga, T. Sogabe, K. Sohda, Y. Sugihara, S. Suresh, S. Susan, D. F. Suzuki, H. Suzuki, S. Suzuki, Y.

387 397 275 9,,233 215 463 221 215 509 245 221 749 257 355 563 563 263 93 331 521 ,667 539 ,545 137 47 , 307 673 627 761 731 425 ,469 695 69 575, 581 ,617 257, 539, 551 , 575 581, 617 ,725 313 425 773 155 239 ,463 627 75 227 725 627 695

784

Tadano, M. Taka, K. Takahashi, K. Takeuchi, Y. R. Takizawa, H. Tanaka, M. Tanaka,T. Tang, X. F. Tani, J. Tanihata, K. Tanizaki, H. Tashiro, Y. Teraki, J. Terauchi, J. Tobioka, M. Tokita, S. Tokumura, A. Tomota, Y. Touchard, G. Tsuchiya, B. Tsujikami, T. Tsujimoto, T. Tsuneyoshi, K. Tu, R. Tiiffe,S. Uchida, Yoshihisa Uchida, Yoshiyuki Uchino, K. Ueda, A. Ueda, Y. Ueki, K. Ueltzen, M. Uheda, K. Umebayashi, S. Umegaki, T. Uo,M. Vikhor, L. N. Wakamatsu, Y. Wakashima, K. Wang, J. H. Wang, X. L. Wang, Z. X. Watanabe, R. Watanabe, S. Watanabe, Y. Watari, F. Watkins, T. R. Weissenbek, E. Willert-Porada, M. A. Williamson, R. L.

469 343 575,,593 149 701 155 725 47 773 599 251 59 483 521 155 633 373 197 373 719 289 197 611 47,,307 159 337 337, 343,,373 361 761 695 667 331 701 419 221 749 501 463 123 , 131 403 387 65 ,301 143, 445, 515 ,527 337, 343 ,373 713 749 387 75 349 387

Xiang, X.

269

Yamada, H. Yamada, J. Yamada, R. Yamaguchi, S. Yamamoto, A. Yamamoto, S. Yamashita, K. Yang, Y. Y. Yatsenko, A. V. Yeo,J. G. Yin, Z. D. Ying, Z. Yokoyama, A. Yoneda, S. Yonehara, E. Yoshino, J. Yuan, R. Z.

409 337 439 755 533 719 221 41 283 185 269 179 749 575 221 495 295 457

Zhai, P. C. Zhang, B. Zhang, L. M. Zhang, Q. J. Zhang, X. D. Zhang, Y Zhu, J.

203, 53,

47, 87, 307, 87, 53, 47, 307, 403, 87,

457 179 445 457 35 53, 179 53, 179, 203, 269

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