Chemical Processing of Ceramics (Materials Engineering, 8)

Chemical Processing of Ceramics MATERIALS ENGINEERING 1. Modern Ceramic Engineering: Properties, Processing, and Us...

Author: Burtrand Insung Lee | Edward J. A. Pope

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Chemical Processing of

Ceramics

MATERIALS ENGINEERING

1. Modern Ceramic Engineering: Properties, Processing, and Use in Design. Second Edition, Revised and Expanded, David W. Richerson 2. lntroductianto Engineering Materials: Behavior, Properties, and Selection, G. T. Murray 3. Rapidly SolidifiedAlloys: Processes StructuresApplications, edited by H0 ward H. Liebermann 4. Fiber and Whisker Reinforced Ceramics forStructuralApplicationS , David Belitskus 5. Thermal Analysis of Materials, Robert F. Speyer 6 . Friction and Wear of Ceramics, edited by Said Jahanmir 7. Mechanical Properties of Metallic Composites, edited by Shojiro Ochiai 8. Chemical Processing of Ceramics, edited by Burtrand 1. Lee and Edward J. A. Pope 9. Handbook of Advanced Materials Testing, edited by Nicholas P. Cheremisinoff andPaul N. Cheremisinoff 6

Additional Volumes in Preparation

Chemical Processing of

Ceramics edited 14y

Burtrand I. Lee Clemson University Clemson. South Carolina

Edward J. A. Pope MATECH Westlake Village, Calihrnia and University of Utah Salt LakeCity#Utah

Marcel Dekker, Inc.

New York*Basel*Hong

Kong

Library of Congress Cataloging-in-Publication Data

Chemical processingof ceramics I edited by Burtrand I. Lee, Edward J . A. Pope. cm.p. Includes bibliographical references and index. ISBN 0-8247-9244-0 (alk. paper) 1. Ceramicmaterials. I. Lee,BurtrandInsung. II. Pope,Edward J . A. TP810.5.C48 1994 666-dc20 94-25396 CIP The publisher offers discounts on this book when ordered in bulk quantities. For more information, write to Special SalesProfessional Marketing at the address below. This book is printed on acid-free paper.

Copyright 0 1994 by Marcel Dekker, Inc. All Rights Reserved. Neither this book nor any part may be reproducedor transmitted in any formor by any means, electronic or mechanical, including photocopying, microfilming, and recording, or by any information storage and retrieval system, without permission in writing from the publisher. Marcel Dekker, Inc. 270 Madison Avenue, New York, New York 10016 Current printing (last digit): 10987654321 PRINTED IN T I E UNITED STATES OF AMERICA

Foreword

The value of a book is not only in its contents but in the creative idea it further encourages. These are a few ideas about the past and the future that occur to me as stimulated by chapters in this book. The volume is a fine resource of references to the use of chemistry in materials and should, indeed, encourage thinking about such things. The early use of chemistry for particular compounds was usually not with materials chemistry in mind but rather was chemistry for chemistry’s sake. The very early nature of much chemistry, now rightly readdressed for materials, appears to surprise many who believe (or wish it to be believed) that their field is new. Many chemical methods are naturally revived with new materials goals in prospect. Mechanochemical studies and aerogels are in vogue at the moment, the former more than 100 years old and the latter 60 years old. Only very ocoasionally will someone, with enough time and energy to review the early literature, produce a historical perspective as Hlavacek did for thermite processing also known as fast propagating reactions. So we have come to realize that although the “spin” on chemistry for materials is new, the processes often have a long and steady historical progression. Also not new are the phenomena seen in nanostructures-all sol-gel (e.g., Chapter 13 “all aerogels are nanocrystalline”) and life chemistry is nanostructural, and chemists long have been able to prepare molecular clusters and nanocrystalline powers by vaporization or, better, by decomposition of precursors (typically hydroxides, carbonates, nitrates, acetates, citrates, and so on). Even the recent production of biphasic or polyphasic nanostructures from polymers had been preceded. by the decomposition of mixed crystals [e.g., CaMg(Co3)2, dolomite, to CaO and MgO]. Carbon and S i c fibers (NicalonB) are nanostructural. In this book, nanocrystalline cobalt from 1966 is mentioned (Chapter 18). Naturally, the nanostructural works of chemists could not be directly examined before the advent of electron microscopes, but more indirect iii

iv

FOREWORD

methods (sometimes forgotten today) had given us a good knowledge of what was going on. Although chemical synthesis has a long history, many new, simple syntheses are being discovered and many more are mentioned in this book, e.g., the use of alkali iodides to make the molybdenum bronzes (Chapter 3) and creative uses of electrochemistry (Chapter 15). Many new structure types, some related to known types, are yet to be found by suitable new methods as discussed by Rao in Chapter 3. For example, many regular intergrowth structures are little lower in free energy of formation than the separate-layer types standing alone and so can be quite difficult to attain. Nevertheless they may have unique and useful properties. For every rule there are exceptions, and this is peculiarly true in chemistry; it seems to be taken by many as gospel that precursorsto complex ceramics, chemically mixed at the molecular level, must always be preferable to other kinds of mixtures but “it ain’t necessarily so.” However, of all the possible methods of chemical mixing, the default recourse to the Pechini method seems curious as it is hardly the simplest or most desirable method in most cases. To elaborate: if unwanted intermediates are produced at lower temperatures on the way to the desired compounds, then the purpose of molecular mixtures may be defeated. This is true in some ferroelectrics and special routes involving deliberately unstable or transient intermediates can be preferable, as mentioned in Chapter 17. It was noted early that for the P-alumindmagnetoplumbitefamily, chemical mixing, which, indeed, produced products at lower temperatures (usually desirable), also led to syntactic intergrowth problems; this has recently been seen again in the high-temperature superconductors (HTSC). (There are no new ceramic phenomena in the HTSC.) In this case also, the production of special intermediates (e.g., Pb-doped-2122 en route to Pb-doped-2223) can be beneficial. Early alkoxide work was certainly done without modem materials in mind; rather the early sol-gel chemistry was with absorbents and catalyst supports in view. Much synthetic work has been done on alkoxides, but a peculiar theoretical schizophrenia exists about them and their relation to oxides. Studies show that bond lengths and angles are very similar in alkoxide clusters to those found in oxides and polyoxyanions. This is mentioned several times in the book (e.g., Chapters 1, 2, and 21). If it looks like a duck, etc., it probably is a duck, which is to say that the bonding in alkoxides is very similar to that in oxides. Alkoxides are not covalent and oxides ionic; they are similar. Modem X-ray methods to determine actual charges on “ions” in oxides show them to be very small. While modeling Madelung potential methods, etc., often work, they are as ballet or opera is to love, merely an artistic (and often useful) imaginative representation. This will become more apparent as workers are attracted to this field and as old shibboleths fade. Even with the extensive work on alkoxides described in this book, there are

FOREWORD

V

surprising new (old!) developments worth mentioning especially the use of polyhydric alcohols to attain unusual coordination states in silicon (viz. 5-fold coordinate with glycols and 6-coordinate with catechols). New synthetic possibilities will further enhance the power of alkoxide chemistry. The area of heterometallic alkoxides (as discussed in Chapter 2) is very important and there are many cheap ways of synthesizing them, including the reactions of chlorides and similar chemicals with already formed alkoxides. People who claim that alkoxides are too expensive to be used should realize that in any real commercial process the alcohols would be recycled. Difficulties with alkoxides-sol-gel, other than the silicon cases mentioned several times in this book, suggest to me that in the future syneresis, an area that surprisingly is little studied (mentioned in Chapter 17), should pay dividends. Those who may believesthat large-scale chemical methods are unlikely to be used for ceramics should consider the following: the earliest synthesis of Si3N4 was by the chemical reaction of Sic14 with ammonia and not the, perhaps, more familiar reaction of Si plus N2 (elemental Si was not readily available in the early days). Today the largest volume of Si3N4 powder (>300 tons per year) is made by that earliest reaction route and much of it in the future seems destined for automobiles. Carbothermal nitridation of Si02 to produce Si3N4 was only later done with fertilizers in mind. As discussed in Chapter 5, an unexpected related manifestation of this is the production of S i c by pyrolysis of rice hulls. It was for the longest time believed that p-Si3N4 was the high-temperature form, or could only be produced at higher temperatures than the a-form; chemical methods have revealed otherwise, as when pSi3N4 forms from prepolymers at 1150°C (Chapter 15). This exemplifies how chemical synthesis, often designed simply with the goal of making something, serendipitously reveals results in conflict with traditional wisdom that has been acquired via a different viewpoint. In chemical processing particular methods are rarely an answer to many problems but they may address niche uses of materials (the centrifugal pipe lining use of “thermite processing” is such a use). Although much early chemistry exists in the literature, and is not difficult to find, there is also the fact that much experimentation occurred that was never published. I know of all kinds of preliminary tests that were done and not followed up in the light of their perceived limited value at the time. It is inevitable that this ground will be gone over again and while the original workers cannot take credit for unpublished work, the modem “discoverers” should recognize that the oldsters knew much more than they let on, but inevitably they could not recognize the pertinence of their knowledge to the materials problems of today. Peter E. D.Morgan

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Preface

Despite many recent advances in material science and engineering, the performance of ceramic components in severe conditions is still far below the ideal limits predicted by theory. Modem ceramics have been primarily the products of applied physics and parallel the developments of physical metallurgy. The emphasis on the relation between behavior and microstructure has been fruitful for ceramic scientists for several decades, It has been recently realized, however, that major advances in ceramics during the next several decades will require an emphasis on molecular-level control. Organic chemistry, once abhorred by ceramic engineers trained to define ceramics as “inorganic-nonmetallicmaterials,” has become a valuable source of new ceramics. It has recently become known that as the structural scale in ceramics is reduced from macro to micro and to nano crystalline regimes, the basic properties are drastically altered. A brittle ceramic material has been shown to be partially ductile, for example. The impetus and the ultimate goal in chemical processing of ceramic materials are to control physical and chemical variability by the assemblage of uniquely homogeneous structures, nanosized second phases, controlled surface compositional gradients, and unique combinations of dissimilar materials to achieve desired properties. Significant improvements in environmental stability and performance should result from such nanoscale molecular design of materials. The unique properties of ceramics with comparable mechanical performance may be realized by the molecular level or nanoscale fabrication of the chemical building blocks of the materials. A number of books are available that deal with the chemical processing aspect of ceramic materials but most of them are conference proceedings. This book is written by many authors who are actively involved in the field of chemical processing of ceramic materials. The authors are from the intemational materials community-Japan, Germany, Korea, France, India, and the United States, where they practice chemical principles in the fabrication of superior ceramic materials. vii

viii

PREFACE

This book presents current development and concepts in chemical techniques for production of state-of-the-art ceramic materials in a truly interdisciplinary fashion. The twenty-three chapters are divided into six parts reflecting topical groups. It first discusses the starting materials-how to prepare and modify them in the molecular precursor stage. Powders are the most heavily used form of starting ceramic materials. Synthesis, characterization, and behavior of ceramic powders are presented in Parts I1 and 111. In the fourth part, forming and shaping of ceramic components via sol-gel technique are discussed. Fabrication of nonoxide ceramics is covered in Part V. In the last part of this book, several specific examples of classes of ceramic materials fabricated by chemical processing including thin films, membranes, and superconductors are presented. These classes of examples are chosen on the basis of the current demand and active research. The topics of basic principles of sub-gel technique, sintering, and postsintering processing are not reviewed in this volume since there are other excellent books dealing solely with these topics. Although this book is edited, it is organized to reflect the sequence of ceramic processing and the coherent theme of chemical processing for advance ceramic materials. Hence this book is suitable as a supplementary textbook for advanced undergraduate and graduate courses in ceramic science and materials chemistry as well as an excellent reference book for practicing chemists and materials scientists and engineers. A s shown by some of the data presented in this book, the results from chemical processing are not yet up to the real applications of ceramic materials. It is evident that, through further research, the full potential of chemical processing for ceramic materials that perform up to the theoretical limit can be realized. It is the authors’ and the editors’ desire to bring the ceramist and chemist closer together to produce superior ceramic materials. Burtrand I. Lee Edward J. A. Pope

Contents

Foreword Peter E. D.Morgan Preface Contributors

iii vii xi

1. Precursor Chemistry 1. Molecular Design of Transition Metal Alkoxide Precursors

Jacques Livage 2. Metal Alkoxides for Electrooptical Ceramics Liliane G. Hubert-Pfalzgraf

3 23

II. Powder Synthesis and Characterization

3. Chemical Synthesis of Metal Oxide Powders C. N. R. Rao 4.

Multicomponent Ceramic Powders T. Mah, E. E. Hemes, and K. S. Mazdiyasni 5. Chemical Synthesis of Nonoxides Christian Russel and Michael Seibold 6. Techniques for Characterization of Advanced Ceramic Powders S. G. Malghan, P. S. Wang, and V. A. Hackley

61 75 105 129

111. Powder Processing 7. Colloid Interface Science for Ceramic Powder Processing Hyun M. Jang 8. Ceramic Nonaqueous Particles Media in

Burtrand I. Lee 9. Synthesis and Dispersion of Barium Titanate and Related Ceramic Powders Ki Hyun Yoon and Kyung Hwa Jo

157 197 215

ix

CONTENTS

X

10. Rheology and Mixing of Ceramic Mixtures Used

in Plastic Molding Beebhas C. Mutsuddy

239

W . Sol-Gel Processing 11. Processing of Monolithic Ceramics via Sol-Gel

J. Phalippou 12. Bulk Optical Materials from Sol-Gel Edward J. A. Pope 13. Aerogel Manufacture, Structure, Properties, and Applications Jochen Fricke and Joachim Gross 14. Fractal Growth Model of Gelation Edward J. A. Pope

265 287 311 337

V. CeramicsviaPolymerChemistry 15. Nonoxide Ceramics via Polymer Chemistry Kenneth E. Gonsalves and Tongsan D. Xiao 16. Polymer Pyrolysis Masaki Narisawa and Kiyohito Okamura

359 375

VI. Processing of Specialty Ceramics 17. Processing of Lead-Based Dielectric Materials Hung C. Ling and Man F. Yan 18. Synthesis of Magnetic Particles Masataka Ozaki 19. Chemistry and Processing of High-Temperature Superconductors Shin-Pei Matsuda 20. Preparation and Properties of Tantalum Oxide Thin Films by Sol-Gel T. Ohishi 21. Crystalline and Amorphous Thin Films of Ferroelectric Oxides Ren Xu 22. Ceramic Membrane Processing C. Guizard, A. Julbe, A. Larbot, and L. Cot

Index

397 421

445

465

481 501

533

Contributors

L. Cot Centre National de la Recherche Scientifique, Montpellier, France Jochen Fricke Physikalisches Institut der Universiut, Am Hubland, Wurzburg, Germany Kenneth E. Gonsalves University of Connecticut, Storrs, Connecticut Joachim Gross Physikalisches Institut der Universitat, Am Hubland, Wurzburg, Germany

C. Guizard Centre National de la Recherche Scientifique, Montpellier, France V. A. Hackley National Institute of Standards and Technology, Gaithersburg, Maryland

E. E. Hermes Wright Paterson Air Force Base, Ohio Liliane G. Hubert-PfalzgrafUniversitk de Nice-Sophia Antipolis, Nice, France Hyun M. Jang Pohang University of Science and Technology, Pohang, Republic of Korea Kyung Hwa Jo* Yonsei University, Seoul, Korea

A. Julbe Centre National de la Recherche Scientifique, Montpellier, France A. Larbot Centre National de la Recherche Scientifique, Montpellier, France Burtrand 1. Lee Clemson University, Clemson, South Carolina Hung C. Ling AT&T Bell Laboratories, Princeton, New Jersey Jacques Livage Universite Pierre et Marie Curie, Paris, France *Currenr affiliation: Daewoo Corporation, Ltd.,

Seoul, Korea

xi

xii

CONTRIBUTORS

l: Mah UES, Inc., Dayton, Ohio S. G. Malghan National Institute of Standards and Technology, Gaithersburg, Maryland Shin-Pei Matsuda Hitachi, Ltd., Ibaraki, Japan

K. S. Mazdiyasni General Atomics, San Diego, California P. E. D. Morgan Rockwell Science Center, Thousand Oaks, California Beebhas C. Mutsuddy Michigan Technological University, Houghton, Michigan Masaki Narisawa University of Osaka Prefecture, Osaka, Japan

T. Ohishi Hitachi, Ltd., Ibaraki, Japan Kiyohito OkamuraUniversity of Osaka Prefecture, Osaka, Japan Masataka Ozaki Yokohama City University, Yokohama, Japan

J. Phalippou Universitb de Montpellier 11, Montpellier, France Edward J. A. Pope MATECH, Westlake Village, California, and University of Utah, Salt Lake City, Utah C.

N. R. Rao 1ndian.Instituteof Science, Bangalore, India

Christian Russel Universitat Jena, Jena, Germany Michael Seibold OSRAM GmbH, Augsburg, Germany

P. S. Wang National Institute of Standards andTechnology, Gaithersburg, Maryland Tongsan D. Xiao University of Connecticut, Storrs, Connecticut Ren Xu University of Utah, Salt Lake City, Utah Man F. Yan AT&T Bell Laboratories, Murray Hill, New Jersey Ki Hyun YoonYonsei University, Seoul, Korea

I PRECURSOR CHEMISTRY

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1 Molecular Design of Transition Metal Alkoxide Precursors Jacques Livage Universitk Pierre et Mane Curie Paris, France

I. INTRODUCTION Transition metal alkoxides are much more reactive toward hydrolysis and condensation than silicon alkoxides. This arises mainly from the larger size and lower ,electronegativity of transition metal elements. Coordination expansion becomes a key parameter that controls the molecular structure and chemical reactivity of these alkoxides. Hydrolysis and condensation rates of silicon alkoxides must be increased by acid or base catalysis, whereas they must be carefully controlled for the other metal alkoxides. The chemical modification of transition metal alkoxides leads to the development of a new molecular engineering. The chemical design of these new precursors allows the sol-gel synthesis of shaped materials in the form of fine powders, fibers, or films.

II. HYDROLYSIS AND CONDENSATION OF METAL ALKOXIDES Sol-gel chemistry is based on the hydrolysis and condensation of metal alkoxides M(OR),. These reactions can be described as follows [l]: M-OR M-OH

+ H20 ”-+ M OH + ROH hydrolysis + R 0 -M +M- 0 -M + ROH condensation

Condensed species are progressively formed, giving rise to oligomers, oxopolymers, colloids, gels, or precipitates. 3

4

LNAGE

Actually these reactions correspond to the nucleophilic substitution of alkoxy ligands by hydroxylated species XOH, as follows: M(OR), + AXOH _ j [M(OR),.AOX),I + xROH where X is hydrogen (hydrolysis), a metal atom (condensation), or even an organic or inorganic ligand (complexation). They can be described by a S$ mechanism [2]:

H” \ O-M-O’-R+

H

\

06+M”-OR/ X

/

X

Step 1

H” / XO-M--OXO-M+ROH \ R Step 3

Step 2

1. The first step corresponds to the nucleophilic addition ofnegatively charged H06 groups onto positively charged metal atoms M&. It leads to an increase in the coordination number of the metal atom in the transition state. 2. The second step is a transfer, within the transition state, of the positively charged proton toward a negatively charged OR group. 3. The positively charged protonated alkoxide ligand ROH is then removed.

The chemical reactivity of metal alkoxides toward hydrolysis and condensation mainly depends on the positive charge of the metal atom 6nn and its ability to increase its coordination number N [2]. Asa general rule, the electronegativity of metal atoms decreases and their size increases when going down the periodic table, from the top right to the bottom left (Table 1). The chemical reactivity of the corresponding alkoxides toward hydrolysis and condensation then increases.

Table 1 Electronegativity x, PartialCharge 6, Ionic Radius r, and Maximum Coordination NumberN of Some Metal Alkoxides

Alkoxide

X

6

Si(OiPr)4 Ti(OiPr)4 Zr(0iPr)rl Ce(OiPr)4 PO(OEt)3 VO(OEt)3

1.74 1.32 1.29 1.17 2.11 1S 6

0.32 0.60 0.64 0.75 0.13 0.46

?-(A) 4 7 8 4 6

0.40 0.64 0.87 1.02 0.34 0.59

N 6

MOLECULAR DESIGN

OF TRANSITION METAL ALKOXIDE

5

Silicon alkoxides are not very reactive. Gelation occurs within several days after water has been added [3]. Hydrolysis and condensation rates must be increased viaacid or base catalysis. The hydrolysis rate of Ti(OEt)4 ( h = lO-3/M/s) is about five orders of magnitude greater than that of Si(OEt)4 ( h = 5 x lO-g/M/s). Cerium alkoxides are very sensitive to moisture. They must be handled with care in a dry atmosphere; otherwise, precipitation occurs as soon as water is present [4]. Alkoxides of highly electronegative elements, such as PO(OEt)3, cannot be hydrolyzed under ambient conditions [5], whereasthe corresponding vanadium derivatives VO(OEt)3 readily give gels upon hydrolysis [6].

111. MOLECULARSTRUCTURE OF TRANSITION METAL ALKOXIDES Silicon alkoxides have beenwidelyused for thesol-gelsynthesis of silicabased glasses andceramics. Silicon remains fourfold coordinated (N = 4) in the precursor as well as in the oxide. All silicon alkoxides Si(OR)4 are therefore monomericandtetrahedral. Their reactivity decreases whenthe size of the alkoxy group increases; thisis mainly caused by steric hindrance factors, which play a major role during the formation of hypervalent silicon intermediates [l]. The situation is completely different with other metal alkoxides M(OR),. The coordination number N of the metal atom in the oxide M o a is usually larger than the oxidation state z. As a consequence, coordination expansion appears to be a general tendency for most metal alkoxides. Positively charged transition metal atomsMz+ tend to increase their coordination number by using theirvacant d orbitals to accept electrons from nucleophilic ligands (nucleophilic addition). This currently occurs via oligomerization (OR bridges), solvation or the formation of oxoalkoxides (p-oxo bridges) [2].

A.

Oligomerization

Sharing alkoxy groups is the easiest way for metal alkoxides to increase the coordination of the metal atom without changing their stoichiometry. In pure alkoxides, coordination expansion currently occurs viathe formation ofOR bridges. Therefore oligomeric as well as monomeric molecular precursors can be found. Oligomerization depends on physical parameters (concentration and temperature) and chemical factors(solventand oxidation state of themetal atom or steric hindrance of alkoxide groups) [7]. Bulky secondary or tertiary alkoxy groups tend to prevent oligomerization, and the steric hindrance of alkoxy ligands appears to be a major parameter. Oligomeric species [Ti(0Et)4ln(n = 2 or 3) have been evidenced for titanium ethoxide whereas Ti(OiPr)4 and Ti(OtAm)4 remain monomeric (Fig. 1). X-ray absorption experiments at the Ti-K edge have been performed on the

6

LNAGE

'OEt

Figure 1 Molecular structure of titanium alkoxides. (a) = TiOiPr4; (b) [Ti(oEt)4]3.

neat alkoxides in the liquid state [8]. Prepeaks can be seen in X-ray Absorption Near Edge Structure (=S) spectra just before the absorptionedge (Fig. 2a). They correspond to 1s + 3d electronic transitions and are very sensitive to the local symmetry around titanium. A sharp prepeak is observed in Ti(OtAm)4, suggestingthatTi is fourfold coordinated, whereastheweakerprepeakin Ti(OEt)4 should correspondto a fivefold coordination. Moreover, two Ti-0 distances (1.82 and 2.05 A) and Ti to Ti correlations (3.12 A) can be observed in Ti(OEt)4, evidencing the oligomeric structure of this alkoxide. Such correlations are not seen in the Extended X-ray Absorption Fine Structure (EXAFS) spectrum of Ti(OtAm)4, which exhibits a single Ti-0 distance (1.81 A; Fig. 2b). The chemical reactivity of metal alkoxides strongly depends on their molecular structure. Oligomeric titanium alkoxides, in which Ti has a higher coordination number, are less reactive, allowing better control of hydrolysis and condensation reactions. Monodispersed Ti02 powers are usually obtained from Ti(OEt)4 rather than Ti(OiPr)4 [9,10]. Oligomerization becomes more and more important as the N-Z difference increases, that is, as the oxidation state Mz+ decreases. Divalent metals give insoluble polymeric alkoxides [M(OR)2],, (#+ = Fe, CO, Ni, Cu, . . .). This was a real drawback for the sol-gel synthesis of high-Tc superconducting ceramics, such as YBa2Cu307-~Bulky ligands, such as 2-(2-ethoxy-ethoxy)ethoxide had then to be used to prevent oligomerization and obtain soluble copper oxide molecular precursors [l l].

B. Solvation Metal alkoxides are not miscible with water, so that sol-gel reactions must be performed in the presence of a common solvent, such as an alcohol. Coordination expansion can then also occur via solvation. Solvate formation is often

MOLECULAR DESIGN OF TRQNSITION METAL ALKOXIDE

7

Ti-0

.2 '

.¶

.6

.3

I

.

1wnm

.. :.

.

J

4

:i :g J

E

R

r 45.02 .18

4.96

5

Ti-0

.9 -

1.2

.6

-

.3

-

!

.!/ i'i/ I J

4.96

Ti(OEtI4

E I

4.98

I

5

I

5.02

Figure 2 X-ray absorption of titaniumalkoxides at the Ti-Kedge.(a) = X m S ; (b) = EXAFS (Fourier transform).

observed when alkoxides are dissolved in their parent alcohol. The stability of such solvates increases with the size and the electropositive character of the metal, that is, when going down the periodic table [7]. Molecular complexity increases with the size of the metal atom. Ti(IV), for instance, gives monomeric (Ti(OiPr)4 species, whereas Zr(IV) and Ce(1V) give dimeric species, [Zr(OniPr)4,iPrOnH]2and[Ce(OiPr)4(iPrOH]2. Such dimers can be crystallized from the solution. Their structure, studied byX-ray diffraction, shows that alcohol molecules are directly bonded to the metal atom to increase its coordination (Fig. 3) [12,13]. Because coordination expansion could occur either via alkoxide bridging or solvation, the molecular complexity of metal alkoxides can be tailored by an appropriate choice ofsolvent.[Zr(OnPr)41n oligomers are formed (n I4) in nonpolar solvents, such as cyclohexane, allowing slow hydrolysis rates and the formation of clear gels. Less condensed solvates are formed in propanol (n = 2); hydrolysis becomes much faster and leads to precipitation [14].

8

LNAGE Pro H

PrO

OPr

H OPr

Figure 3 Molecularstructure of [Ce(OiPr)4, PirOHIz.

C.

Formation of Oxoalkoxides

Alkoxide groups can also bridge two different metal atoms, leading to the formation of heteroalkoxides, which are often used a precursors for the sol-gel synthesis of multicomponent ceramics [15]. Condensation can even go one step further, leading to p-oxo bridges via ether elimination. M-OR+RO-M"-0-"M+ROR This typically occurs when alkoxide solutions are heated under reflux to provide better mixing of the components. Condensation is favored by the smaller size of oxo ligands and their ability to exhibit highercoordination numbers, M"oxo ligands, for instance,havebeenbeenobservedin Na2FeaO(OCH3)18; 6CH30H [16]. Oxo bridges favor the coordination expansion of metal atoms, and oxoalkoxides are more stable thanthe corresponding alkoxides and of course less reactive toward hydrolysis and condensation [17]. Large and electropositive metals are known to give oxo-alkoxides, such as Pb40 (OEt)6. This Pb(I1) precursor undergoes complete dissolution in ethanol when Nb(OEt)5 is added, giving rise tothecrystallization of [Pb604(OEt)4][Nb(OEt)5]4.Such heterometallic alkoxides have the correct Pb/Nb stoichiometry for the sol-gel synthesis of PbMgln NbmO3 (PMN) ceramics [18]. Heterometallic alkoxides are often more soluble than their parent alkoxides, a property that can be advantageously used for the sol-gel chemistry of nonsoluble alkoxides. They also provide molecular precursors withthecorrect M I M stoichiometry inwhichsome M-0-M' bonds are already formed [15]. Alkoxide bridges are usually hydrolyzed during the sol-gel synthesis, but oxo bridges are strong enough not to bebroken.CrystallineBaTiOscanbe obtained at rather low temperatures when a mixture of Ti(OiPr)4 and Ba(0iPr)z is heated under reflux before hydrolysis [19]. It was shown recently that heterometallic tetrameric species [BaTiO(OiPr)4,iPrOH]4arethen obtained in which Ti-O-Ba bonds are already formed in the solution [20].

MOLECULAR DESIGN OF TRANSITION METAL ALKOXIDE

9

W. CHEMICALMODIFICATIONOF TRANSITION METAL ALKOXIDES Most metal alkoxides are very reactive toward hydrolysis and condensation. They must be stabilized to avoid precipitation. The chemical control of these reactions is currently performed by adding complexing reagents that react with metal alkoxides at a molecular level, giving rise to new molecular precursors of different structure, reactivity, and functionality. Chemical modification is usually performed with hydroxylated nucleophilic ligands, such as carboxylic acids or P-diketones. In most cases complexation by XOH species can be described as a nucleophilic substitution, as follows: M(OR),

A.

+ xXOH +[M(OR),-AOX),I + xROH

CarboxylicAcids

Acetic acid (AcOH=CH3-COOH) is often used as an acid catalyst for the solgel synthesis of Si02. Adding a small amount of AcOH significantly decreases the gel time of silicon alkoxide systems. In similar conditions, this effect appears to be morepronouncedthan with inorganic acids, such as HCl [21]. However, the reverse effect is observed for Ti(OR)4. Adding acetic acid prevents precipitation and increases the gel time [22]. An exothermic reaction takes place when acetic acid is added to titanium alkoxide in a 1:l ratio. A clear solution is obtained. X-ray absorption experiments show that the coordination of titanium increases up to six(Fig. 4a), whereas two Ti-0 distances (1.80A and 2.06A) and Ti-Ti correlations (3.1 1 A) are observed on the ,Fourier transform of the EXAFS spectrum (Fig. 4b). Infrared spectroscopy can be used to study how acetate groups are bonded to tita-

(a)

Cb)

Figure 4 X-ray absorption at the Ti-K edge of pi(OiPr)3(OAc)]. (a) = XANES; (b) = EXAFS (Fourier transform).

WAGE

10

nium. The frequency shift AV between the two bands corresponding to the symmetric v, = 1450 cm" and antisymmetric v, = 1580 cm" stretching vibrations of the acetate group COO- is typical of bidentate bridging acetate groups (AV = 130 cm"). Acetate groups actually behave as complexing nucleophilic ligands and react with titanium alkoxides as follows: Ti(OiPr)4+ AcOH +[Ti(OiPr),(OAc)] + iPrOH During this stoichiometric reaction, the coordination number of Ti increases from four to six, and oligomeric species [Ti(OiPr)3(OAc)In ( n = 2 or 3 are formed (Fig. 5) Esterification occurs when more than 1 mol AcOH is added. Acetic acid in excess reacts with alcohol molecules released during complexation, providing the in situ generation ofwaterandgiving rise to more condensedspecies. Small oligomeric species are obtained in the presence of small amounts of water. Single crystals of hexameric Ti604(OiPr)~?(OAc)4,for instance, have been obtained upon aging an equimolar mixture of AcOH and Ti(OiPr)4 in a closed vessel (Fig. 6). Water, provided via esterification reactions arising from acetic acid in excess, leads to the slow hydrolysis of alkoxide groups, which are replaced by oxo bridges [23]. Most organic groups can be removed and clear transparent polymeric titanium dioxide gels are obtained in the presence of an excess of water. Upon hydroysis, alkoxy groups are removed first, whereas bidentate acetate ligands remain bonded to titanium. They prevent further condensation and increase the gelation time. The functionality of the new molecular precursor de-

Figure 5 Chemical modification of titanium alkoxides. (a) = ri(OiPr)3 (0&)]2; = [Ti(OiPr)3(acac)].

(b)

MOLECULAR DESIGN OF TRANSITION METAL ALKOXIDE

I1

0" OW

Figure 6 Molecularstructure bridges (hatched circles).

of [Ti604(0iPT)12(OAc)4]:Ti(solidcircles);

oxo

creases, leading to the growth of anisotropic particles. TizO(0Ac)rj oxoacetates, made of chainlike polymeric species, are obtained when an excess of acetic acid is added [24].

B.

P-Diketones

Strongly complexing P-diketones are currently employed to stabilize highly reactive metal alkoxides, such as W(OEt)6 [25]. Aluminum sec-butoxide, modified by ethylacetoacetate (etac), appears to be quite attractive as a precursor for the sol-gel synthesis of multicomponent ceramics, such as cordierite. Al(OsBu)z(etac) is more soluble and less reactive thanthe corresponding alkoxide [26]. Complexation is also observed with acetylacetone (acacH=CH3-CO-CHzCO-CH3). Its enolic form contains hydroxyl groups and reacts with metal alkoxides as a chelating ligond. Oligomers are not readily formed, and for a stoichiometric acacni = 1 ratio, the nucleophilic substitution leads to monomers in which Ti is only fivefold coordinated (Fig. 5b). Ti(OiPr)4+ acacH +Ti(OiP&(acac)

+ iPrOH

12

LWAGE

Oligomeric compounds are formed only upon hydrolysis [27]. Titanium coordination becomes six as soon as water is added, whereas Ti-Ti correlations can be seen on the Fourier transform of the EXAFS spectrum. Strongly chelating acac ligands are not removed upon hydrolysis (except at low pH). Condensation is then prevented, and only small oligomers are formed [28].

V. MOLECULARENGINEERING OF METALALKOXIDE PRECURSORS Metal alkoxides react with nucleophilic complexing ligands. Two chemical parameters can then be used to design new molecular precursors. Metal alkoxides react with water molecules (hydrolysis) and nucleophilic species (complexation). The first reaction is followed by condensation and leads to the formation of larger oligomeric or polymeric species, whereas complexing ligands prevent condensation and favor the formation of smaller species. It therefore becomes possible to design sol-gel-derived materials by controlling the following two chemical parameters: the nature and amount of complexing additive XOH (x = X / M ) and the hydrolysis ratio (h = H20/M). A large variety of oligomeric species can then be obtained upon hydrolysis and condensation. Molecular clusters, chain polymers, or colloidal particles can be synthesized, depending on the relative amount of hydrolysis and complexation. A s a general rule, more condensed species are obtained as x decreases and h increases.

A. HYDROLYSIS RATIO The hydrolysis and condensation of metal alkoxides M(OR)z leads to oligomers in which metal atoms tend to acquire their maximum coordination number. The formation of these species can be controlled by the hydrolysis ratio h = H20/M. In the presence of a large excess of water (h >> z). all alkoxide groups are removedand colloidal species are formed.Theylead to hydrous oxides MOzI2.xH20 similar to those synthesized from aqueous solutions. The adsorption and dissociation of water molecules at the oxide/water interface leads to the formation of charged particles. All alkoxide groups are not removed when h I z and chain oxopolymers are formed. They can be advantageously employed for drawing fibers or making coatings. The dielectric constant of the organic solvent is rather low, the surface charge of polymeric particles is very small, and vander Waals interactions prevail. For very low hydrolyis ratios (h -c 1) condensation is mainly governed by the formation of p o x 0 and alkoxo bridges. Solute molecular oxoalkoxides are then formed and can often be isolated as single crystals from the solution.

I3

MOLECULAR DESIGN OF T M S I T I O N METAL ALKOXIDE

Their structure appears tobe closely related to that of the corresponding polyanions formed in aqueous solutions. The first hydrolysis product ( h = 0.6) of titanium ethoxide, for instance Ti704(OEt)20, contains [Ti70241 units isostructural with MO7@& [29]. Larger species, such as Ti1008(0Et)24 ( h = 0.8) and Ti16016(OEt)32 (h = l), are formed when more water is added [30,31]. They have been obtained upon the controlled hydrolysis of Ti(OEt)4 and isolated as single crystals from the solution (Fig. 7). The molecular structure of these species in the solution has been studied by x-ray diffraction and 1 7 0 nuclear magnetic resonance using enriched water as a reagent [32]. Niobium ethoxide leads to Nb8010(OEt)20, which exhibits the same structure as the paratungstate [33]. In the presence of an organic base, NMwOH, a decamer, is obtained [Nblo028(NMe4)6*6H20], whichis isostructural with the decananadate polyanion (V10028)6- [34].

B. Complexation The chemical reactivity of metal alkoxides toward hydrolysis and condensation can be modified by complexation. Less electronegative alkoxide ligands are hydrolyzedpreferentially,whereasstronglybonded complexing groups are more difficult to removed. They prevent condensation, and gelation or precipitation rates decrease upon complexation. Complexing ligands behave as termination reagents. They cause the effective functionality toward condensation to be reduced, resulting in less condensed and more anisotropic polymeric species. Precipitation can then be avoided, and gelation is promoted.

b Figure 7

Structure of somemolecularclustersobtained via thecontrolled hydrolysis of Ti(OEt)4. (a) = [Ti704(0Et)20]; (b) = [Ti1008(OEt)24].

14

LNAGE

The reaction of titanium alkoxides with acetylacetone leads to several molecular compounds.[TiO(acac)z]z is formed in the presence ofan excess of acetylacetone or upon hydrolysis of Ti(acac)2(OR)2. Single crystals have been isolated;andx-raydiffraction experiments show dimers with sixfold coordinated Ti atoms linked through oxygen atoms. Strongly complexing acac ligands cannot be hydrolyzed easily, and condensation does not go any further (Fig. 8) [35]. Lager molecular species can be obtained with smaller ainounts of acetylacetone. Single crystals of Tii8022(0Bu)26(acac)2were obtained recently [36]. They are made of 18 [Ti061 octahedra sharing edges or comers and correspond to x = 0.1. h = 1.2. Complexing acac ligands remain outside the Ti18022 core of the molecule (Fig. 9). These examples show that condensation can be tailored via complexation and hydrolysis. Gels are obtained in the presence of an excess of water and for small values of x (x I0.3, h 2 10). Clear sols are obtained when Ti(OnBu)4 is hydrolyzed in the presence of acetylacetone. The mean hydrodynamic diameter of colloidal particles was measured by quasi-elastic light scattering. It appears to increase from 2 to 40 nm as the hydrolysis ratio increases from h = 1 to h = 4 for x = 0.3). It decreases from 40 to 4 nm when the amount of acetylacetone increases from x = 0.3 to x = l (for h = 4) [28]. Similar results have been found with zirconium alkoxides. Solvated dimeric species are formed when zirconium propoxide is dissolved in its parent alco-

Figure 8 Molecular structure of [TiO(acac)2]2 according to Ref. 35: Ti (solid); p OXO bridges (hatched); other atoms (C and 0) (open).

-

MOLECULAR DESIGN OF TRANSITION METAL

Figure 9

ALKOXIDE

15

Molecular structure of [Ti18022(0Bu)26(acac)2](h = 1.2 and x = 0.1).

hol. Tetameric Zr40(0nPr)1o(acac)4 species are obtained with acetylacetone (x = acacnr = 1) in the presence of a very small amount of water ( h = 0.2). They have been isolated as single crystals and their structure determined by x-ray diffraction (Fig. 10) [37]. Larger molecular clusters, such as [Zr1006(OH)4(0nPr)l8L6], with kallylacetoacetate, can be formed by decreasing x (x = 0.6) and increasing h (h = l) (Fig. l l ) [38]. Monodispersed zirconia powders have been prepared by Rinn and Schmidt [39] via the acid hydrolysis of Zr(OiPr)4 in ethanol in the presence of hydroxypropylcellulose as a stabilizer to avoid aggregation and acetylacetone as a complexing agent. Particles from about 200 to 10 nmin diameter were obtained depending on the relative amount of each reagent. Dimericsolvated solute species [Cez(OiPr)s(iPrOH)z] are formed when cerium isopropoxide is dissolved in its parent alcohol. These solutions are very sensitive to moisture and can be stabilized with acetylacetone. Molecular clusters are formed when a small amount of water is added. Red crystals have been

LNAGE

16

0

Zr

e

b 0 (oxo bridge)

Figure 10 Molecularstructure of [Zr4PO(OPr)lo(acac)4] (X = 1 and h = 0.2).

grown from a solution in which x = acac/Ce = 2. They correspond to hexameric species [Ce604(0H)4(acac)l2] in which the oxohydroxo core (Fig. 12) is isostructuralwiththe inorganic analog [Ce604(OH)4(S04)6], whichwas observed in aqueous solutions of cerium sulfate [ 131. The hydrolysis of P-diketonate-modified cerium(IV) isopropoxide leads to the formation of colloidalsolutions or gels. The complexationratio (x = acac/Ce) appears to be the key parameter to tailor the size of cerium oxide particles. Precipitation is observed when x e 0.1, whereas sols are obtained when 0.1 < x e 1. The mean hydrodynamicdiameter of these particles decreases from 450 to 15 A when x goes from 0.1 to 1. Only solute molecular clusters are formed when the complexing ratio becomes larger than 1 (1 Ix I2).

VI.ORGANICALLYMODIFIEDTRANSITIONMETAL ALKOXIDES Organically modified silicates can be conveniently synthesized from precursors, such as R4Si(OR), which contain nonhydrolyzable Si-C bonds so that organic moieties are not removed during the hydrolysis-condensation process. Organic groups R can behave as either network modifiers or network formers

MOLECULAR DESIGN

0

Zr

Figure 11

e

OF TRANSITION METAL

ALKOXIDE

17

0 (oxo & hydroxo bridges)

Molecular structure of [ZrloOs(OPr)ls(acac)6] (x = 1 and h = 0.8).

when polymerizable organic ligands are used. They allow the formation of an organic network in addition to the inorganic. The organic groups impart new properties to the organic network, such as flexibility, hydrophobicity, or refractive index [40]. The chemistry of hybrid organic-inorganic gels is mainly developed around silicon-containing materials. Titanium and zirconium alkoxides are sometimes used as cross-linking reagents for the condensation of difunctional silicon precursors, such as [(CH3)2Si(OC2H5)2]. However, these alkoxides do not really behave as cross-linking agents. They lead to the formation of Poly Dimethyl Siloxane (PDMS)-Ti02 nanocomposites in which Ti02 nanoparticles are embedded in a PDMS matrix [41]. This certainty arises from the higher reactivity of these alkoxides and their well-known catalytic activity toward the condensation of siloxanes. Therefore, PDMS chains are formed on one side and Ti02 or ZrO2 nanoparticles on the other side. Si-O-Ti bonds have not yet been clearly evidenced. The catalytic role of titanium and zirconium alkoxides to-

18

LNAGE

Figure 12 Molecularstructure of [Ce604(OH)4(acac)l2].

ward the condensation of siloxanes is quite interesting for film formation. The presence of long chains in the film makes the material more flexible, allowing the deposition of crack-free transparent coatings several micrometers thick. Hybrid organic-inorganic compounds containing transition metal species were recently synthesized using organically modified polyoxometallates linked, together through W-O-Si-C bonds [42]. The organically modified precursor is obtained via thereaction of GSiW11039 witha trichlorosilane RSiCl3, R=vinyl (CH=CH2), allyl (CH2CH=CH2), or 3-methacryloxypropyl [-CH2(CH2)2OOCC(CH3)=CH2]. Each [SiW1104o(SiR)2]4- carries two binfuctional R groups, available for organic polymerization reactions performed in the presence of radical initiators. Organically modified transition metal oxide gels are difficult to synthesize. Transition metals have a lower electronegativity than silicon. The M-C bond becomes more polar and would be broken upon hydrolysis. Complexing hydroxylated organic ligands must then be used to form nonhydrolyzable M-0-C

MOLECULAR DESIGN OF TRANSITION METAL ALKOXIDE

19

bonds between the metal atom and the organic species. Complexation appears to be the only way to synthesize hybrid gels involving transition metal oxides. However, very few results have been reported so far in literature. According to the results reported previously, carboxylic acids R-COOH could be conveniently employed as complexing species. Prehydrolyzed titanium butoxide (H2013 = 1) oligomers have been complexed with unsaturated organic acids, such as cinnamic acid(C6H5-CH=CH-COOH).Copolymerizationwasthenperformedwith styrene in the presence of benzoyl peroxide. Transparent brown polymers were thus obtained. They are very stable against water and cannot be dissolved in most organic solvents [43]. Acrylic or methacrylic acids can also be used as polymerizable complexing ligands [a]. They were recently reported for the sol-gel synthesis of zirconium oxides via the copolymerization of zirconium oxide sols and organic monomers [45]. Carboxylates are rather weak complexing ligands, however, and most of them are removed upon hydrolyis when an excess of water is added. Chelating P-diketones are stronger complexing ligands. Organically modified Ti02 gels, which exhibit photochromic properties, have been synthesized from an allyl acetylacetone-modified Ti(OnBu)4 alkoxide. A double polymerization process was initiated via the partial hydrolysis of alkoxy groups and the radical polymerization of allyl functions. The polymerization of the allyl function is slow, however, and organic polymerization is not very effective [46]. New approaches have been explored using functionalized chelates with both a strong chelating end and a highly reactive methacrylate group, such as aceto acetoxy ethyl methacrylate and methacryl amido salicylic acid (Fig. 13).

lor

HC

\

HC

I

0

Figure 13 Chelating ligands for the synthesis of Z r O 2 hybrid gels. (a) = Aceto acetoxy ethyl methacrylate (AAEM); (b) = methacyl amido salicylic acid (MASA).

20

LNAGE

Zirconium-oxopoly-polyAAEMcopolymers have been synthesized via the chemical modification of Zr(OnPr)4 precursors by M M . Both polymerization reactions were run simultaneously, leading to hybrid organic-inorganic polymers. The zirconium oxo core is made of oxoalkoxo AAEiM-modified species in which Zr is likely in a sevenfold coordination, as in monoclinic zirconia. Zirconium oxo species are chemically bonded to polymeric methacrylate chains via the P-diketo complexing groups [47]. The complexation ratio ( A A E W r ) seems to control the structure and texture of these hybrid gels, giving rise to more or less open structures. Organic and inorganic polymerization processes are presumably related, and such a hybrid material could be described as an interpenetrating polymer network (Fig. 14). Both organic and inorganic polymerizations of zirconium propoxide modified by methacryl amido salicylate were performed using similar procedures. This MASA ligand exhibits a stronger complexing power and a higher steric hindrance. Therefore, for equivalent complexation ratios (MASA substitutes two alkoxo groups), simultaneous polymerizations of methacryl amido salicylate-modified zirconium precursors lead to smaller zirconium oxopoly methacryl amido salicylate copolymers [47].

Figure 14 A hybrid organic-transitionmetaloxidenetwork.

MOLECULAR DESIGN OF TRANSITION METAL ALKOXIDE

21

Other routes could be used, such as the following: Copolymerizationbetweenhybridinorganic-organicsystems,zirconiumoxopoly-MM, andathirdorganicpolymerizablecomponent,suchas styrene. The complexation of zirconiumalkoxideswithligandscarryinganazobis group from which radical polymerization could be initiated Moreover,hybridorganic-inorganiccopolymerscouldbesynthesizedfrom many other metal oxide gels (rare earth, aluminum and others).

REFERENCES 1. Brinker, C. J., and Scherer, G.W., Sol-Gel Science, Academic Press, New York, 1989. 2. Livage, J., Henry, M., and Sanchez, C. Prog. Solid Sate Chem., 18, 259 (1988). 3. Klein, L., Annu. Rev. Muter. Sci., 15, 227 (1985). 4. Toledano, P., Ribot, F., and Sanchez, C., C.R. Acud. Sci. Fr., 311, 13 15 (1990). 5. Livage, J., Barboux, P., Vandenbore, M. T., Schmutz, C., and Taulelle,F., J. NonCryst. Solids 147-148, 18 (1992). 6. Nabavi,M.,Sanchez,C.,andLivage, J., Eur. J. SolidStateInorg. Chem, 28, 1173 (1991). Metal Alkoxides, Academic 7. Bradley,D. C., Mehrotra, R. C.,andGaur,D.P. Press,London,1978. 8. Babonneau, F., Doeuff, S., Leaustic, A., Sanchez, C., Cartier, C., and Verdaguer, M., Inorg. Chem., 27, 3166 (1988). 9. Barringer, E. A., and Bowen, H. K., J. Am. Cerum. Soc., 65, C-l99 (1982). 10. Barringer, E. A., Bowen, H. K., Langmuir, 1, 414 (1985). 11. Scozzafava, M. R., Rhine, W. E., and Cina, M. J., Better ceramics through chemistry, IV, Muter. Res. Soc. Symp. Proc., 180, 697 (1990). 12. Vaartstra, B. A., Huffman, J. C., Gradeff, P.S., Hubert-Pfalzgraf, L., Daran,J. C., Parraud, S., Yunlu, K., and Caulton, K. G.,Inorg. Chem., 19, 3126 (1990). 13. Toledano, P., Ribot, F., and Sanchez, C., Acta Crystullogs, C46, 1419 (1990). 14. Kundu D., and Ganguli, D., J. Muter. Sci. Lett., 5, 293 (1986). 15. Caulton, K. G., and Hubert-Pfalzgraff, L., Chem. Rev., 90, 969 (1990). 16. Hegetschweiler, K., Schmalle, H. W., Streit, H. M., Gramlich, V., Hund, H. U., and Emi, I., Inorg. Chem., 31, 1299 (1992). 17. Livage, J. and Sanchez, C., J. Non-Cryst. Solids 145, 11 (1992). 18. Hubert-Pfalzgraf,L.,Papiemik, R., Massiani,M.C.,andSepte,B.,Betterceramics through chemistry, IV, Muter. Res. Soc. Symp. Proc., 180, 393 (1990). 19. Mazdiyasni,K. S., Dolloff, R. T., and Smith, J. S., J. Am Cerum. Soc., 52, 523 (1969). 20. Yanovsky, A. I., Yanoskaya, M. I., Limar, V. K., Kessler, V. G., Turova, N. Y., and Struchkov, Y. T., J. Chem Soc. Chem. Cornmuc., 1605 (191). 21. Pope, E. J., and Mackenzie, J. D., J. Non-Cryst. Solids, 87, 185 (1986). 22. Doeuff, S., Henry, M., Sanchez, C., and Livage, J., J. Non-Cryst. Solids, 89, 206 (1987).

22

WAGE

23. Doeuff, S., Dromzee,Y.,Taulelle, F., andSanchez,C., Znorg. Chem., 28, 4439 (1989). Muter. Res. Bull., 25, 1519(1990). 24. Doeuff, S., Henry,M.,andSanchez,C., 25. Unuma, H., Tokoka, T., Suzuki, Y., Furusaki, T., Kodiara, K., and Hatsushida, T., J. Mater. Sci. Lett., 5, 1248 (1986). 26. Babonneau,F.,Coury,L.,andLivage,J. Non-Cryst. Solids, 121, 153(1990). Chem. Mater., I , 240(1989). 27. Uaustic, A.,Babonneau,F.,andLivage,J., 28. Uaustic, A., Babonneau, F., and Livage, J., Chem. Mater., I , 248 (1989). 29. Watenpaugh, K.,andCaughlan,C.N., Chem. Commun., 2,76 (1967). F. S., 30. Day, V. W., Eberslcher,Klemperer,W.G.,Park,C.W.,andRosenberg, Chemical Processing ofAdvunced Materials (L. L. Hench, and J. K. West, eds.), Wiley, New York, 1992, p. 257. 31. MossetA.,andGaly,J., C.R. Acad. Sci. Fr., 307, 1747(1988). 32. Day, V. W., Eberspacher, T. A., Klemperer, W. G., Park, C. W., and Rosenberg, F. S., J. Am. Chem. Soc., 113, 8190 (1991). 33. Bradley, D. C., Hurthouse, M. B., and Rodesiler, P. F., J. Chem. Soc. Chem. Commun., 1112 (1968). 34. Graeber, E. J. and Momson, B., Acta Crystallogr., B33, 2136 (1977). 35. Smith,G.D.,Caughlan,C. N., andCampbell, J. A., Znorg. Chem., 11, 2989 (1972). 36. Toledano,P., In, M.,andSanchez,C., C.R. Acad. Sci. Fr., 313, 1247 (1991). C.R. Acad. Sci. Fr., 31 1, 1161 (1990). 37. Toledano, P., In, M., and Sanchez, C., 38. Livage,J.Sanchez,C.,andToledano,P., Metaloxideclustersand colloids, Mater. Res. Soc. Symp. Proc., San Francisco 272, 3 (1992). 39. Rinn,G.,andSchmidt,H., CeramicTransactions (G.L.Messing,E.R.Fuller and H. Hausner, eds.), Am. Ceram. Soc., I, 23 (1988). 40. Schmidt,H., J. Non-Cryst.Solids, 73, 681(1985). 41. Dir6, S., Babonneau, F., Sanchez, C., andLivage,J., J. Mater.Chem., 2, 239 (1992). 42. Judeinstein,P., Chem.Mater., 4, 4 (1992). 43. Suvorov,A.L.,andSpaski, S. S., Proc. Acad. Sci. USSR, 127, 615(1959).

44. Schubert, U., Arpac, E., Glaubitt, W., Helmerich, A., and Chau, C.,Chem. Mater.,

4, 291 (1992). 45. Nass, R., and Schmidt, H., Sol-Gel Optics, SPIE, 1328, 258(1990). J. Non-Cryst. Solids, 100, 46. Sanchez, C., Livage, J., Henry, M., and Babonneau, F., 650 (1988). 47. Sanchez,C.,andIn,M., J. Non-Cryst. Solids 147-148, 1(1992).

2 Metal Alkoxides for Electrooptical Ceramics M a n e G. Hubert-Pfalzgtaf Universid de Nice-Sophia Antipolis Nice, France

1.

INTRODUCTION

Ceramics are one of the most recent classes of electrooptical materials. Table 1 summarizes the various compositions associated with transparent electrooptical ceramics [l]. These materials are mostly ferroelectric lead-containing titanates or niobates.Amongthe various types of precursors ofmetal oxides (e.g., carboxylates and nitrates), alkoxides have especially attractive features. These include high purity, the lability of the coordination sphere, which allows their tailoring according to specific applications, their easy transformation into oxides with formation of volatile by-products [2,3], and their ability to form homogeneous solutions under a large variety of conditions, and, for multicomponent systems, via heterometallic alkoxides [4]. As a result of their versatility, they can meet the requirements for sol-gel as well as metal organic chemical vapor-phase deposition (MOCVD) applications [5]. This chapter deals with the synthesis and general properties of homo- and heterometallic &oxides that are or might be used in chemical routes to electrooptical ceramic materials. Some considerations concerning structure, reactivity, and tailoring of their properties are also given. Emphasis is given to the most recent results.

II. HOMOMETALLICALKOXIDES For all elements involved in electrooptical ceramics, their homometallic alkoxides, of the general formula M(OR)n, or oxoalkoxides, MO(OR),, are either 23

on

24

HUBERT-PFALZGRAF

Table 1

Composition of ElectroopticalCeramics

Composition La)(Zr, (Pb, Ti103 PSN Ti103Nb)@ Pb(Sc, (Pb, La)(Hf, (Pb, Sr)(Zr, Ti)O3 (Pb, Ba, Sr)(Zr, Ti)@ (Pb, La, Li)(Zr, Ti)@ Ba, (Pb, Sn)(In, Zr, Ti)@ (Pb, M)(Zr, Ti)@; M = Bi, Sr (Pb, Bi)(Zr, Ti)03

PLZT M(Ta, Nb)03; M = Li, K PLHT PSZT (Pb, La)(Mg, Nb, Zr, Ti)@ PLMNZT PBSZT (Ba, BLTN Nb)@ La)(Ti, (Pb, PLLZT La)Nbz06 PBLN PSIZT (Pb, La)(Zn, Nb, Zr, Ti)@ PBLNZT PMZT (Sr. Ba)Nb206 SBN PBZT

commercially available or can be readily prepared in high yield (strictly anhydrous conditions are required for theirhandling. Table 2 lists their general properties and main synthetic routes.

A.

'

Synthesis

A variety of synthetic pathways, depending on the starting material, are generally available for most elements. Their value can be quite different in terms of selectivity of the reaction and yield and purity of the resulting metal alkoxide. The most electropositive metals (alkaline, alkaline-earth metals, yttrium, and lanthanides) can be oxidized directly by alcohols. The rate of the reaction is largely dependent on the surface properties of the metal and on the nature of the alcohol, especially its acidity and the bulkiness of the OR group. Thus the reaction proceeds easily with alkaline metals (Li, Na, and K), the primary alcohols being more reactive than the secondary or the tertiary. The same variation is observed for the alkaline-earth metals, the reaction being the least favored for magnesium [6,7]. In some cases, it might be necessary to catalyze the reaction by mercury or by gaseous ammonia. Ammonia was observed to be efficient toward barium when classic catalysts (Hg and 12) failed; it also avoids contamination [49]. Activation of yttrium and lanthanides by trace amounts of mercury(I1) salts (chloride or acetate) is generally required, although sonochemistry can be a way to activate the metals in the absence of mercury derivatives [50]. Nonotable reaction is observed between yttrium and lanthanides and methanol, ethanol, or t-butanol, but the reaction proceeds well with isopropanol and functional alcohols, such as 2-methoxyethanol [51,52]. Synthesis of metal alkoxides directly from the metal can lead to degradation reactions either of the solvent (tetrahydrofuran, THF, for instance [ 181) or of the reactant 1131; they appear especially easy with barium. Alcoholysis reactions applied to alkoxides (alcohol interchange reactions) or to metal silylamides generally make it possible to overcome these limitations.

METAL ALKOXIDES FOR ELECTROOPTICAL CERAMICS

25

These reactions are generally achieved in the presence of an excess of alcohol, and thus the alkoxides might be obtained-in solution as solvates

[m.

(113:

M+ROHexcess+

M(OR)n(ROH)x+;H2

(1)

The lability of the alcohol molecules and thus the stability of the solvates is quite variable. Such formulas as Ba(OR)2 or M(OR)3 (M = In, Y, Ln) have been traditionally used to describe alkoxides of these electropositive elements. Many of these alkoxides have now been more precisely defined by x-ray diffraction studies, which have revealed quite different formulas. In fact, the removal of the alcohol molecules (under vacuum) often leads to oxoalkoxides. Pentanuclear oxo or hydroxo units seem to be a common feature for tri- and divalent metals and are illustrated by MsO(OiPr)13 (M= In, Sc; Ln = Y, Yb, . , .) and Bas(OH)(OR)s adducts [R = Ph, 3,5-tBuzCtjH3, CH(CF3)2], and even P-diketonates [53]. It is thus important to emphasize that the formulation, the properties, and the reactivity of metal alkoxides based on large elements may be different according to their “history”-used in situ or isolated [50]. Substitution reactions are another general approach to metal alkoxides, according to Eq. (2):

MG,

+ nROH +M(OR),(ROH), + nGH

(2)

G = NR2, R, H, Cl, . . . . These substitution reactions are especially attractive for obtaining a pure product when the by-product is volatile and thus the anionic ligand G, an amide group NR2, an alkyl group R, or a hydride H. Alcoholysis reactions of trimethylsilylamino derivatives M[N(SiMe3)2In are quite a versatile route to a variety of metal alkoxides, although they require the prior synthesis of the metal silylamide derivatives (commercially available only for Li and Na). This procedure is particularly attractive for metals whose alkoxides are poorly soluble, such as lead, zinc, or iron isopropoxides, since contamination with halide or alkali residues can be overcome. It also allows a large choice of R groups, and lanthanide tertiobutoxide derivatives, for instance, become accessible [39]. The use of appropriate (bulky or functional) alkoxide or aryloxide groups has also opened a way to soluble and/or volatile lead(II) and zinc derivatives [23,29]. Reactions according to Eq. (3) are generally extremely fast and may favor the formation of oxo derivatives, as observed for bismuth [U. N(SiMe3)2],

+ nROH+

M(OR),

+ nHN(SiMe3)2

(3)

Alcoholysis of metal alkyl derivatives is based on the elimination of a volatile by-product as well:

MR, + nR‘OH +M(OR’),

+ nRH

(4)

E

P Q

a g

m

B B

z"z"

B

z"

E B

P

z"

Bm

%

e ....

m

8 2 .-aM m 0

c)

2m 0 c)

x

8 z.. 3 2

.-9ax

c a 8

27

28

Q

d

i

d

e

s

0 .P - 0 -

M

"'-

B

v)

s

29

30

HUBERT-PFALZGRAF

It is of practical use for metals whose alkyl derivatives are commercially easily available, namely, lithium and magnesium. Metal hydrides can be used as starting material mainly for alkali and alkaline-earth metals. AlthoughZn(0R)z alkoxides derived from classic OR groups are insoluble soluble tetranuclear zinc alkyl or hydridoalkoxides [ZnX(OR’)]4 (X= R or H) can be obtained starting from zinc alkyls or hydride [30,31]. Finally, other general routes involve halides, mostly chlorides, as starting materials, according to Eq. (5):

M = alkaline metals. Substitution by alcohols in the presence of a base is the most general route to early transition metal alkoxides, especially for groups IV (Ti, Zr, and Hf) and V (Nb and Ta), and they are commercially available, at least for some R groups. Less labile halides (lead, bismuth, or lanthanide chlorides) require the use of alkali metal alkoxides to achieve substitution. Even so, the reaction may proceed with a poor yield and/or offer final products containing halide or alkali metals, the latter being present either as an impurity or as part of a definite product (heterometallic species). When halides are used as a starting material for metal alkoxides, contamination by chlorides can also result from a limited purification: variable amounts of Cl or Na are often found in commercial titanium isopropoxide Ti(OiPr)4. Substitution reactions applied to acetates generally result in a poor selectivity and thus a low purity of the metal alkoxides; their interest is thus limited despite the low cost of the‘acetates. Reactions between hydrated acetates and refluxing 2-methoxyethanol have been used as a way to obtain acetatoalkoxides in sol-gel processing, but most products are poorly characterized. Electrosynthesis by anodic dissolution ofthe metal in absolute alcohol appears for many elements to be a promising, inexpensive way to obtain large amounts of metal alkoxides [54]. The process goes smoothly and has good current yields for metals, such as scandium and lanthanides, and early transition metals (Ti, Zr, and Nb) using tetrabutylammonium bromide as an electrolyte. Oxo or hydroxo compounds are obtained for more electropositive metals (Mg and alkaline-earth metals), however, probably as a result of alcohol dehydration reactions. Finally, alcohol exchange reactions are a means of obtaining more appropriate precursor, starting from commercial or easily accessible metal alkoxides (see Sec. KC).

METAL ALKOXIDES FOR ELECTROOPTICAL CERAMICS

B.

31

PhysicalProperties and StructuralAspects

Most of the metal alkoxides of interest for electrooptical ceramics are solids (less often liquids) that can be purified by recrystallization, sublimation, or distillation. They are all moisture sensitive, and handling under an inert atmosphere and with anhydrous solvents is thus required. Their unequivocal characterization and formulation are best achieved by x-ray diffraction studies (on monocrystals). Studies on solutions (molecular weight data, nuclear magnetic resonance, NMR, with ‘H, 13C, or metal nuclei) are a means either to establish whether the solid-state structure is retained or, in the absence of x-ray data, to establish the molecular structure and eventually stereolability [48]. Mass spectrometry provides information on the stability of the oligomers or the heterometallic species in the vapor phase. The information gainedby infrared spectroscopy is limited; the technique is mostly useful for the identification of solvates M(OR)n(ROH)x (vOH absorption 3400-3100 cm-1 or of chemically modified (heteroleptic) alkoxides (probe for the vC0 stretching of P-diketonate or carboxylate ligands, for instance). The solubility and volatility of homoleptic alkoxides are mainly determined by the degree of oligomerization m,which depends on the nature of the metal (size and thus coordination number) and on the steric demand of theOR group. Methoxides are thus the least favorable. With the exception of magnesium, lead, and zinc and dn transition metals, isopropoxides are reasonably soluble for all metals of interest for electrooptical ceramics. The insolubility of the first can be overcome by addition of another metal alkoxide (see Sec. 111) or by using a more bulky ligand, such as aryloxide (2.6-di-tert-butylphenoxide)or triphenylsiloxide PhsSiO, or functional alcohols (for instance, OHC2H& with X = OR’ for alcoxyalcohols or X = NR’2 for alkanolamines), which provide intramolecular donor sites [52].Although the aryloxide or triphenylsiloxide group; generally lead to soluble species (often monomeric with a low coordination number for the metal), they might be of limited interest as a result of a low ceramic yield and high organic content of the resulting oxides. Fluorinated OR groups lead to an increase in volatility, reduce the tendency to hydrolysis, and favor the formation of adducts [19]. The bulky tert-heptyloxide groups (R = CMeEtiPr) generally represent a good means of achieving volatility [5]. The metals involved in the composition of electrooptical ceramics are large, oxophilic, and electropositive. High coordination numbers are thus generally required. Recent years have seen, as a result of x-ray studies, the reformulation of many alkoxides based on di-or trivalent large elements as oxo- or hydroxoalkoxides. It is wellknownthat M(OR)n alkoxides are often oligomersdimers, trimers, or tetramers-as a result of the OR group acting as a doublebridging (p2) or triple-bridging ligand (p3), which allows the metal to attain its usual coordination number. Oxo 02- or hydroxo OH- ligands are even more

32

HUBERT-PFALZGRAF

versatile and can behave as p2 or p3 ligands but can also accommodate four, five, or even six metals. As a result, oxo or hydroxo ligands appear to be a means of achieving the high coordination numbers required by the large di- or trivalent element in the absence of neutral ligands. Barium seems to be especially prone to formation of oxo or hydroxo derivatives, and this observation appears in the earlier literature as a tendency to give carbon-deficient or unstable metal alkoxides. The size effect of the mental is illustrated by comparison of Sr and Ba phenoxides: the strontium derivative is a tetranuclear nonoxo unit in which the metal is pseudooctahedral [16]; the barium derivative is a pentanuclear oxo or hydroxo aggregate, which allows Ba, which is larger than Sr, to be sevencoordinate for the basal metal atoms [ 181. Polydentate alcohols, such as triethanolamine, can encapsulate the metal and drastically improve the stability toward hydrolysis [14]. Figure 1 summarizes the basic [M(OR)& oligomers, as well as the oxo or hydroxo aggregates that are observed. Dimeric units (Fig. 1A) are illustrated by [M(OR)5]2 (M= Nb, Ta) or [M(OiPr)4(iPrOHh (M = Zr, Hf); [Pb(OtBu)2]3 and La3(OtBu)gL2 are typical examples of openand closed trinuclearunits (Fig. 1B and C), respectively; tetramers display a [Ti(OMe)4]4 type of structure (Fig. 1D); (also found for divalent or trivalent metals but with additional neutral ligands as for [Nd(OiPr)3(iPrOH)]4or Sr4(0Ph)8(PhOH)z(THF)6) or a cubanelike, (Fig. 1E) as for [K(OtBu)]4 or [ZnR(OR')]4 [S]. The structurally characterized penta- and hexanuclear aggregates are all oxo or hydroxo clusters.M4(p&o)o6 cores (Fig. 1F) havingan adamantanelike structure are found for lead and barium alkoxides. The M5(pS-o)Oy (Fig. 1G) core seems to display a special stability for large di- or trivalent metals. Hexanuclear aggregates derive from the pentanuclear by the addition of a sixth

Figure 1

Basicaggregates for M(OR),lm alkoxides or oxoalkoxides.

33

METAL ALKOXIDES FOR ELECTROOPTICAL CERAMICS

X = OR,R

E

PR G

M5°14

OR

OR

I

I

H Figure 1

6R

Continued

metal (Fig. 1 H ) or by condensation of adamantane units. Oxo or hydroxo aggregates are mostly highly soluble despite their nuclearity-the OR groups ensure a lipophilic layer-and they are also sometimes volatile. NMR and mass spectrometry data have shown that these aggregates are generally retained in solution as well as in the gas phase. They result from C-0 bond cleavage reactions (with formation of alkenes) andor oxidation reactions (especially for

34

HUBERT-PFMZGRAF OR'

bridging-chelating

bridging P2 - q'

P2 - ? l 2

triply- bridging P3

M

triply-bridging P3 -y'

terminal

-q2

Figure 2

Coordination modes of 2-methoxyethanol.

alkaline-earth metals) and are present before hydrolysis; in some cases pb(II)], further condensation is observed by recrystallization, especially when dialkyl ethers or siloxanes are easily formed [24,25]. 2-Methoxyethanol displays numerous coordination modes (Fig. 2) but is basically a bidentate and assembling ligand; [M(OC2H40Me)n]malkoxides are thus large oligomers (m= 9, M = Ca; m = 10, M = Y), sometimes infinite polymers [M = Bi, Pb(II)]. Despite the large values of the nuclearity, however, they are soluble (with the exception of lead) because of cyclic (M = Y) or compact structures (M = Ca) or dissociation of the polymer in solution, as observed for Bi [32]. Aryloxides are, with the exception of barium, either monomers or dimers [30,38]. Alkoxide ligands OR and to a lesser extent trialkylsiloxo R3SiO and aryloxo OAr are n donors [3]. l 7 donation is especially favored for do elements (Ti, Zr, I%,. Ta, and lanthanides). Asthe OR ligand undergoes oxygen p to metal d n'bonding, the M-0 distances can become very short for terminal OR groups (4 covalent radii); the M-0-C angles open up to nearly linear M-0-C units. The n bonding as well as the thermal stability of the M-0-C bond decreases with the metal oxidation state and by moving to the late transition metals and is illustrated by more acute M-0-C angles.

C. Chemical Modifications In addition to using bulky OR groups, reduction of the molecular complexity can also be achieved by a polar solvent. Since formation of complexes is quite

35

METAL ALKOXIDES FOR ELECTROOPTICAL CERAMICS

limited for metal alkoxides because of their preferred autoassociation, giving oligomers [48], one of the few appropriate ligand or polar solvents is an alcohol [Q. (6)]. This is illustrated by the fact that many metal alkoxides are obtained as solvates (the M-OHR interaction is often stabilized byH bonding with an adjacent OR group): the group IV isopropoxides (M(OiPr)4(ihOH)]z (M = Zr, Hf) are typical examples [42]. It should also be noted that since the structure may be different according to the solvent [niobium alkoxides (R = Me, Et) are dimers in toluene but are monomeric Nb(OR)s(ROH) adducts in the parent alcohol] [48], the reactivity can be modified, especially toward other metal derivatives (Fig. 3). 1 -[M(0R),lm m

+ xROH

1 2 [M(OR),(ROH),,],

m’ m

(6)

The electronegative alkoxo or aryloxo groups make the metal atoms highly prone to nucleophilic attack. The alkoxides easily react with the protons of a large variety of organic hydroxy compounds, such as alcohols, silanols, glycols, carboxylic acids, P-diketones (Fig. 3). These exchange reactions convert homoleptic M(OR), into heteroleptic alkoxides M(OR),-XZ, and thus reduce the functionality of the precursor (2= OR). They also illustrate the use of alkoxides as versatile starting materials and as precursors whose coordination sphere can be easily modulated for specific applications. Such modifications of the coordination sphere of the initial alkoxide are often achieved by the use of the so-called additives, which are in fact chemical modifiers [3,55a]. 1. Alcoholysis Alcoholysis reactions are generally performed to “tailor” the volatility or the rheology, to decrease the hydrolysis rates, or, finally, to increase the solubility of the metal alkoxides. As a synthetic route, alcoholysis reactions with classic alcohols are usually achieved n refluxing cyclohexane, the reactions being promoted by azeotropic distillation of the alcohol [2]. In contrast with silicon alkoxides, however, alcohol interchange reactions proceed to some extent (depending onthe OR group and on the stoichiometry) at room temperature for most elements [S]. This has been shown, for instance, by metal N M R on niobium and titanium alkoxides. “Uncontrolled” alcohol interchange reactions may thus occur when a metal alkoxide is used in an alcohol different from the parent alcohol. Such reactions lead to a modification (decrease) of the apparent functionality of the precursors, which in turn can affect the properties (morphology, porosity, and so on) of the resulting oxide. Functional alcohols, such as alcoxyalcohols-mainly 2-methoxyethanol-or alkanolamines (diethanolamine and triethanolamine), are actuallythemost commonly used to slow hydrolysis rates and thus to “stabilize” metal alkoxides. Although such exchange reactions occur at room temperature, Some heat-

36

U

0

v

x e p:

0

::

W

C

p:

h

0 E v

0

X

V

0

1 % X

::

ea

p:

n

0 v

n

x m ea X U

/-\

E

U

HUBERT-PFALZGRAF

METAL ALKOXIDES FOR ELECTROOPTICAL CERAMICS

37

ing may be necessary for completion of the reaction, especially if stoichiometric amounts of reactants are used. Thus the reaction between zirconium isopropoxide [Zr(OiPr)4(iPrOH)]2 and four equivalents of 2-methoxyethanol leads to Zr(OCZH40Me)3(0iPr) a room temperature. The presence offunctional alkoxo groups around the metals generally results in an increase in its coordination number, and thus hydrolysis becomes more difficult [57]. Alcoholysis reactions using 2-methoxyethanol generally also offer more soluble products [52]. A s already mentioned,this observation is no longer valid for lead(I1) since the methoxyethoxide derivative; are polymeric and insoluble [23]. The reactivity of a variety of glycols (CH&(OH)2 toward titanium alkoxides, zirconium isopropoxide, and niobium and tantalum ethoxides has been studied by Bradley et al.[2]. Complete substitution generally resultsinpoorly soluble products; such derivatives asM(OR),~X[O(CH~)~O]~ generally remain soluble and sometimes volatile. Further reactivity, for instance with acetylacetone, has established that the Ti-OR bond is more labile than the Ti-glycolate bond. Trialkylsilanols R3SiOH (R = Me, Et) are also a way of improving solubility and controlling hydrolysis [58]; however, they have rarely beenused, probably as a result of their limited stability and thus commercial availability. 2. Reactions with Carboxylic Acid and P-Diketones Substitutionreactionswith carboxylic acids, such as acetic acid, or with Pdiketones, mainly acetylacetone,are other means of controlling hydrolysis rates by decreasing the functionality of the precursor. Carboxylic acids RCO2H are often added to the sol-gel system as an acid catalyst [59,60]. In fact, they are not merely a source of H+, but a chemical modifier, and carboxylate groups substitute for the OR groups. Although P-diketonate and acetate ligands both generally lead to an increase in the metal coordination number, their behavior is quite different. Carboxylates act as assembling and oxo donor ligands, and thus have a tendency to increase the nuclearity of the aggregates; diketones are chelating ligands and thus decrease the oligomerization. This behavior is well illustrated by the modification of titanium alkoxides.Reactionbetween [Ti(OEt)& and acetylacetone gives monomeric and dimeric species (according tox-ray absorption near-edge structurestudies[61], and condensation into hexanuclear oxo clusters is observed when acetic acid is used (x-ray structures of Ti6(~2-o)2(C13-o)2(oR)8(oAC)8 (R = Et. iPr, nBu) [46]. 3. Hydrolysis Hydrolysis reactions of alkoxides proceed by nucleophilic attack of the metal M s . (7)], and polymerization occurs via the reactive hydroxoalkoxides [Eqs. (7b) and (7c)l. Hydrolysis-polycondensationreactions are governed by numerous factors (Table 3). The metals involved in the composition of electrooptical ceramics are electropositive and oxophilic, and thus the hydrolysis of their ho-

HUBERT-PFALZGRAF

38

Table 3 HydrolysisParameters of Homo-andHeteroleptic Alkoxides Electronegativity of the metal and polarity of the M-0-Cbond Nature of alkoxo groupR Modifies the molecular complexity Rate increases with chain lengthening Sensitive to hydrolysis TertiaryR > secondaryRprimaryR OR > OSiR3 pH (acid or basic catalysis) Solvent and dilution Temperature Degree of hydrolysis h (h = [HzO]/M(OR)n) h < n: fibers, chains, coatings h c 1: molecular clusters h > n: gels, tridimensional polymers Modified precursors M(OR)n-xZx(Z= OH, OAc, B-dik,...) Rate decreases with Functionality of the precursor (number of the OR groups) Increase in metal coordination number Hydrolytic susceptibility OR > OAc > B-dik

moleptic alkoxides is often extremely facile and much easier than for silicon. Introduction of other ligands 2 slows hydrolysis rates and promotes anisotropic growth. M(OR),

+ H20

-

H\

0 - hf(OR),.1 H- --OR

2M(OH)(OR),-1+ (OR),-$4M(OH)(OR),-I + M(OR),*

/

> M(OH)(OR),-l(7a)

0 -M(OR),-1 + H20 (OR)n-lM- 0 “M(OR),-1 + ROH

(7b) (7C)

The early stages of the hydrolysis-condensationof [M(OR),Im can be characterized byx-ray diffraction for oxophilic metals (Table 4). These include centrosymmetric polynuclear oxoalkoxides of titanium, zirconium, andniobium. Their structure is often related to that of polyoxoanions: Ti704(OEt)20 is isostructural with Mo70246, and Nbg010(OEt)20 has a cagelike Structure that compares well with that of the paratungstate H~W120421G.These oxoalkoxides are soluble as a result of their closo structure, and 1 7 0 NMR has been developed as a technique to follow the hydrolysis of titanium(1V) alkoxides with

39

40

HUBERT-PFALZGRAF

H217O. Since enrichment in 1 7 0 occurs selectively at their oxide (as opposed to alkoxide) oxygen sites, the different types of oxo ligands (h,n = 2-4) can be determined, the most encapsulated oxo ligands being the most shielded [63]. Similar types of molecular oxo clusters are observed for alkoxides modified by acetylacetone [65a,66]. Although the early stages of the hydrolysis-polycondensation are extremely rapid, removal ofall alkoxide ligands at room temperature generally requires quite a large excess of water [57]. Hydrolysis studies have also been performed on several lead(1I) alkoxides: Pb(OR)2 (R= CMefit, CHMeCH2NMe2) in THF [23]. The hydrate 3PbOeH20 was obtained at low temperature as the kinetic product; its conversion to massicot and then to litharge is catalyzed by an increase in the pH (which may be caused by alkali metal alkoxide residues). Since 3PbOeH20 is less stable than massicot and litharge, it could act as a more reactive component in ceramic oxide preparations.

D. Thermal Decomposition Thermolysis reactions have mostly been studied for group IV metal alkoxides, although metal oxide depositshavebeen obtained for many metals [5]. Oxoalkoxides are often formed asintermediates,but mechanistic studiesremain scarce [68]. The decomposition can be enhanced by hydrolysis as a result of either residual water on the surface or dehydration reactions of tertiary alcohols.

111.

HETEROMETALLIC ALKOXIDES

Another aspect of the lability of the metal alkoxide bond is the easy formation of heterometallic or mixed-metal species [4]. Such compounds have also been called “double alkoxides” in comparison with double salts; however, their description as heterometallic is more appropriate in view of their covalent character and the existence of species having three different metals. The formation of heterometallics is a means of overcoming the poor solubility of polymeric metal alkoxides (Zn and Pb), of achieving homogeneity at a molecular level for multicomponent systems, and of providing a better homogeneity of the final material. If they display a convenient volatility and stability, they can also be used as “single-source” precursors (in which the different elements are incorporated in a single molecule) for MOCVD applications and thus improve the transport of poorly volatile elements and reduce possible interdiffusion phenomena [5]. Table 5 collects the heterometallic alkoxides, associating elements present in the formulation of electrooptical ceramics, although the stoichiometry may be quite different from those required for materials.

METAL ALKOXIDES FOR ELECTROOPTICAL CERAMICS

A.

41

Formation of HeterometallicAlkoxides

Heterometallic alkoxides are generally obtained either by reactionbetween Lewis acids and bases or by condensation with elimination of volatile or insoluble by-products. 1. Lewis Acid-Base Reactions Formation of heterometallic alkoxides by simple mixing, often at room temperature and in nonpolar solvents, is one of the most general routes and corresponds to the "complexation" step in the sol-gel process. This general reaction [Eq. (8)] appliestonearlyallalkoxides,withthe exception of thesilicon Si(OR)4:

yM'(0R')nI

+ M(0R)n +MM'y(OR)n(OR)nty

(8)

Determination of the stoichiometry of the reaction is best achieved if it is conducted in a solvent in which one of the reactants is poorly soluble. Since the heterometallic species is more soluble than the parent alkoxide, the progressive dissolution of the most insoluble species allows control of the stoichiometry between the two different metals. Metal N M R (7Li and 93Nb) may also be used as tool to follow the reaction, as shown for LiNb(OR)6 [91]. The following reactions, which all proceed at room temperature, are examples of this approach: L[Mg(OiPr)2], m M(OR)~ + "(OR)

+ 2Nb(OiPr)5 +MgNb2(OiPr)l2 250c > "'(OR)6

(9) (10)

M = Nb, Ta; M' = Li, Na, K, . . . Alkali metal alkoxides are often involved in the preparation of metal alkoxides [Eq.51; heterometallic alkoxides can thus be a side product if the stoichiometry of the reaction is poorly controlled ( o n ) . For electropositive metals, such as barium, the alkoxide might be generated in situ: 2Ba + [Zr(OiPr)4(iPrOH)]2 dBa2Zr2(pOR)2(OR)18+ H2

(11)

Systematic studies of the influence of the R group and the solvents on the stoichiometry of such reactions as Eq. (8) remain limited, although the Turova group has studied ternary-phase diagrams [M(OR)n, M'(OR)n, and solvent] and proposed compositions for many systems. The Ba-Ti system is one of the best characterized and has shown dependence on the R group (oxo compound with R = zW,nonoxo compound for R = Et) as well as on the stoichiometry of the reactants:

5 E .e

4 .. h E

2

P

0

C

B

e,

.^

z m C

Y

.e

8

%

S X -

52%

F? E

43

44

HUBERT-PFALZGRAF

[BaTi2(OEt)g(EtOH)4][H(OEt)2] +

+

Addition reactions similar to h.(8) can also be observed between anhydrous acetates and alkoxides. Such reactions can even occur at room temperature, but they may require heating (depending on the lability of the metal acetate bond) and/or be less selective and give redistribution products Fomometallic acetatoalkoxides, M(OR)n-dOAc)x] as well. Mg(0Ac)z + 2Nb(OiPr)s 250c > MgNb2(OAc)2(OiPr)lo

(14)

Although metal b-diketonates can be solubilized in the presence of a metal alkoxide and thus form heterometallic species, at least-as intermediates, most reactions are more in favor of the formation of redistribution products [M(OR)n-x(P-dik)x]as the final result [92]. 2. Elimination of Volatiles The formation of heterometallic alkoxides is sometimes driven by that of a volatile by-product, mainly a dialkyl ether. Typical examples are obtaining the oxoalkoxides Ba4Ti404(0iPr)16(iProH)~ and Pb6Nb404(0Et)24:

When acetates are involved, esters can be the volatile by-product and thus favor the formation of the mixed-metal species; this approach can overcome the noncommercial availability of some alkoxides [89]: 3Pb(OAc).y3H20 + Zn(OAc)z.2HzO

> P~Zn~(OAc)4(0Cz&OMe)4 + (17)

As a general observation, it is important to note that reaction between alkoxides and other metal derivatives, such as acetates, can be strongly dependent on the solvent and/or the addition order of the reactants [84]. Thus, no reaction is observed between anhydrous Pb(OAc)2 andNb(0Et)s in THF at room temperature, but the acetate dissolves readilyin toluene at room temperature in the presence of the niobium alkoxide with formation of a Pb-Nb species [go]. Another general feature is that in heterometallic acetatoalkoxides obtained by this route, the metalintroduced as an alkoxideappearspredominant. More condensed species are obtained if the reaction is achieved at higher temperatures [93]: Pb(0Ac)z + [Zr(OiPr)4(iPrOH)]z 250c> PbZr3(~O)(~OAc)z(~OipT)5(OiPr)~ + '

(18 )

METAL ALKOXIDES FOR ELECTROOPTICAL CERAMICS

45

In some cases, the stoichiometry between the two metals M and M' in the isolated species is different from that of the reaction medium, suggesting that other metal species are present [Eq. (12)].

3. Elimination of an Insoluble By-product Such reactions are generally based on the elimination of an alkali salt according to Q. (19): MCl,

+ M"M',(OR),

+ nM'C1

"+MIM'y(OR)y]n

(19)

M"= Li, Na, K, TI(I)3 Although they allow control

of the stoichiometry between the M and M', metals, they have not been much developed so far for species-associating metals involved in electrooptical ceramics. They are mostly limited totheuse of the complex alkoxide anionZrz(OiPr)s-,andLa-Zr species were recently built up by this route [87]. In contrast with the preceding synthetic routes based on mixing simple compounds, Q. (19) implies the prior synthesis of a heterometallic alkoxide based on an alkali metal and further removal of the salt.

B. PropertiesandStructuralAspects Mixed-metal alkoxides are generally covalent [4]. As a result, they are soluble in organic nonpolar solvents (often less in alcohols), generally more than the parent alkoxides. Some can besublimed without disproportionation, the volatility following the same variation as for the homometallic alkoxides (OtBu > OzR > OEt > OMe).Mass spectrometry, for instance, established thatheterometallic LiNb fragments were retained in the gas phase for LiNb(OEt)6. Their structure is mainly determined by the overall nuclearity of the aggregate (MxMy and derives from the basic structural moiety in which the different metals display the usual coordination numbers expected for the bulk of the OR [4]. Numerous mixed-metal homoleptic alkoxides for which the stoichiometry betweenthetwo metals is 1: 1 have a structurethat derives from that of [Ti(OMe)4]4 (Fig. 1D). &iNb(OEt)6], is an infinite polymer; this explains its poorsolubility,but it is also favorable for obtaining fibers [70]. Trinuclear mixed-metal species generally display a triangular Mkf2(p3-OR)2(p-OR)3core (type Fig. IC), the coordination sphere eventually being completed by neutral ligand L as for SrTi2(0iPr)s(iPrOH)3. Association of such triangular units via bridging OR groups or via incommon apex gives such compounds as [BaZr2(p-OR)4(OR)4]2(Fig. 4J) or a 1:4 stoichiometry, such as BaZr4(OR)i8' (type Fig. 4K), respectively [81]. Thestructure of the Pb-Nb oxoalkoxide, Pb6Nb404(OEt)24, derives from that of the lead, with the oxo ligands acting as donors toward Nb(OEt)5 moities (Fig. 4L), which in turn increase the lead coordination number by three double-bridged OR groups. Similarly,

46

HUBERT-PFALZGRAF

Ba4Ti404(0iPr)ia(iP1OH)~ (Fig. 4M) can be considered a Ba4(~-0)4tetrahedron with oxo groups bearing Ti(OR), moeities. MgNb2(p-OAc)2(p-OiPr)4(0iPr)6 corresponds to an open polyhedron of the type shown in Fig. IB, the acetate ligands connecting the central Mg atom to the neighboring Nb ones. The presence of chelating or assembling ligands may modify the basic arrangement between the different metals, however.

C. Reactivity Data on the reactivity of mixed metal alkoxides are limited. Their reactivity remains dominated by the lability of the metal alkoxo bond, but a problem that needs to be addressed is that of the maintaining or not of the heterometallic unit, of the stoichiometry between the different metals, and finally of the ho-

METAL ALKOXIDES FOR ELECTROOPTICAL CERAMICS

47

mogeneity, at a molecular level when they are reacting. Dissociation into the parent alkoxides upon dissolution of the heterometallic species can be a spontaneous process, even in a nonpolar solvent,andoccurs, for instance, for Pb604(OEt)e[Nb(oEt)s]4, asshown by207Pb NMR [79].Suchdissociation phenomena can be promoted by a polar solvent (alcohol or THF), depending on the nature of the metals, and thus the choice of the solvent becomes even more important than for homometallic alkoxides [4]. Common modifiers usedinthesol-gel process are functionalalcohols, acetic acid, and acetylacetone. Alcoholysis reactions are the most documented; KiNb(OEt)6Imfor instance, was converted to LiNb(OC2H40Me)rj in the presence of 2-methoxyethanol [91]. Assembling ligands, such as 2-methoxyethanol or acetic acid, seem more favorable than P-diketones for the stabilization of mixed-metal alkoxides. However, the poorly controlled addition of an additive may sometimes modify the mixed-metal species by “extracting,” at least partially, one metal as an insoluble product or as a complex. Information on the hydrolysis of mixed-metal alkoxides is even scarcer than for the homometallic, Partial hydrolysis of [LiNb(OEt)6lo0leads to the dimeric heterometallic alkoxide [LiNbO(OEt)4(EtOH)]2, in whichboth the stoichiometry between the two metals and their coordination numbers are retained [93]. On the other hand, the partial hydrolysis of solutions of methoxyethoxides of Ba and Ti (1:l molar ratio) offers BwTi13018(OC2H40Me)24 (30% yield). Its structure corresponds to a tetrahedron ofBaO3unitssurimposedon a TiOs(Ti03)n core; this Ti13042 core is related to that of the aluminum salt [Al1304(0H)24(H20)]7+ [83].The important modification of the stoichiometry between the two metals illustrates the complexity of the sol-gel chemistry of multimetallic systems and shows that important structural rearrangements can occur.

IV. APPLICATION TO CERAMICS A.

Sol-Gel Process

PLZT and PNM are, with L i m o s and LiTaOs, the systems related to electrooptical ceramics that have been the most studied.Chemical low-temperature routes are particularly attractive for lead-containing materials, since lead oxide is quite volatile in comparison with other metal oxides; thus control of the stoichiometry of lead-based materials is tedious. To overcome the difficulty inhandling alkoxides and their availability, commonly accessible salts, such as carboxylates (acetates or the more soluble 2ethylhexanoates), carbonates, nitrates, or hydroxides, have often been used in conjunction with metal alkoxides (mostly n-propoxides or n-butoxides for Ti and Zr and ethoxides for Nb or Ta), the solvent being an alcohol. The poor re-

48

HUBERT-PFUGRAF

activity of barium acetate is overcome by adding acetic acid [46a] or by using carbonate or, more often, barium hydroxide Ba(OH)~8Hz0[94]. Commercial ionic precursors are often available as hydrates, and this must betaken into account when speculation about the intermediates that may form is proposed. Carboxylate groups (acetates or 2-ethylhexanoates) are generally removed only at relatively high temperatures (400°C); they appear to favor porous materials and grain boundaries [95]. Additives for hydrolysis are either acids (0.1-0.2 M m03 or acetic acid) or ammonia solutions. Although leadoxide has also been used [96], acetate hydrates are the precursors most often used for the introduction of lead or lanthanum in the multimetallic systems required for PLZT materials,and they are often associated with the useof 2-methoxyethanol [97-991. This solvent appears as a way to achieve deshydratation by refluxing (boiling point 124"C), as well as condensation (elimination of organic esters as has been shown by infrared (vCOZ = 1730 cm-1) [loo]. Studies on the influence of the precursor remain limited. X-ray absorption fine structure analysis of the sol and gel precursors of PZT (lead acetate in methanol, n-propoxides of Zr and Ti, and acetic acid) have shown that an uniform composition exists in both states on a microscopic level with Ti-O-Z linkages [100a]. The choice of the titanium precursor Ti(OR)4, R = nF'r or iPr, was reported to have an effect on the microstructure development, particularly for PbTiOs films [97]. Clusters and grains were observed in the n-propoxide and were attributed to a less continuously cross-linked networks compared with films derived from the isopropoxide. Using the modified alkoxide Ti(OiPr)z(acac)z instead of Ti(OiPr)4 allowed an increase in the thickness of crack-free PbTiOs films up to 1 m under base-catalyzed conditions [loll. Using the same titanium precursor associated with lead acetate and pentane-lJ-diol instead 2of methoxyethanol permitted the production of films up to approximately 5pm in thickness by repeated coatings before fuing [ 1021. All alkoxide routes appear to be limited toobtaining PNM [88a] and lithium or potassium niobate or tantalate [102a]. For the latter, the formulation of the mixed-metal alkoxide MM'(OR)6, M = Nb, Ta, is in agreement with that of the material. Epitaxial thin films of LiNbO3 on sapphire could be obtained [103]. The useof a mixed-metal MgNbz(0R)lz species (R = Et [105], CzH40Me [104]) has been shown to promote the formation of the perovskite phase for PNM.

B. MOCVD Obtaining electrooptical thin films by MOCVD techniques has been less investigated, although volatile oxide precursors-alkoxides 'or p-diketonatesexist for nearly all elements. For PbTiOs films titanium isopropoxide has been

METAL ALKOXIDES FOR ELECTROOPTICAL CERAMICS

49

associated with an alkyl derivative, PbEU using argon as the carrier gas. No additional oxygenwasnecessary. The film deposited at a quartz substrate heated at 500°C was shown to be conducting and dense, with good surface morphology and strong [l001 texture direction [106]. Pb(0tBu)z was used by Trundle and Brierley as the volatile lead source with Ti(OiPr)4, giving 5 pm thick PbTiOs deposited at 450"C, annealed .in air at 800-900°C [107]. Although these authors could not deposit PbSc0.5Tm.503 (PST) directly from the metal precursors, the PST perovskite phase was obtained in atwo-stage process: deposition of cubic ScTaO4 by MOCVD using Ta(0Et)s and a scandium-fluorinated P-diketonate Sc(fod)3 (fodH = heptafluorodimethyloctanedione), followed by diffusion of PbO from a surface layer [5]. The heterometallic alkoxide LiNb(0R)a has the correct stoichiometry for depositing LiNbO3. Although LiNb(0R)a is volatile, control of the deposition parameters is sometimes hampered by disproportionation reactions. The problem has been overcome by using a lithium P-diketonate Lithd and Nb(0Me)s [5]. LiNbO3 has been deposited on a variety of substrates (-45O"C), but epitaxial layers require annealing in oxygen at higher temperatures.

V.

OUTLOOK

Metal alkoxides are readily available for all elements involved in electrooptical ceramics. The selection of appropriate OR groups-bulky or functional and thus polydentate-allows adjustment of their physical properties: solubility or volatility. An almost unlimited number of mixed-metal compositions is accessible under mild conditions, often at room temperature, by mixing metal alkoxides and carboxylates or P-diketonates. The molecular composition of these solutions can be quite complex and comprise mixed-metal species with different stoichiometries. To date, 2-methoxyethanol has been used predominantly for the sol-gel processing of thin layers and thus for the chemical modification of the metalalkoxides. However, methoxyethanol is teratogenic and can cause neurological and hematological damage, even at the ppm level. Alternative solgel systems, based for instance on other difunctional and polyfunctional alcohols and on hydroxyacids, for example, should be developed to promote the sol-gel process as a practical method. More systematic studies regarding the relationship between precursors, influence of additives, and properties of the final material are also required.

ACKNOWLEDGMENTS The author is grateful to the CNRS for financial support, and to the contributions and enthusiasm of coworkers, whose names are listed in the references.

50

HUBERT-PFALZGRAF

ABBREVIATIONS acacH acetylacetone (2,5-pentanedione) acetate OAc ArR R-substituted phenyl group:2,6-RzCaH3 ArR$ R-substituted phenyl group:2,6-Rz-4-MeCsH3 ArF phenylgroupsubstitutedbyCF3 groups: 2,4,6-(CF3)3C,jH2 Bz benzyl CH2C6Hs P-dikHP-diketone = (O)CRCH2C(O)R thdH 2,2,6,6-tetramethylheptane-3,5-dione

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Sakashita, Y . , Ono, T., Segawa, H., Tominaga, K., and Okada, M., Preparation and electrical propertiesof MOCVD-deposited PZT thin films,J. Appl. Phys., 69, 8352 (1991).

88a. Munozaguado, M. F., Gregorkiewitz, M., and Larbot,A., Sol-gel synthesisof the binary oxide (Zr, Ti)Oz from the alkoxides and acetic acid in alcoholic medium, Muter. Res. Bull., 27,87(1992), Laaziz, I., Larbot, A., Julbe, A., Guizard, C., and Cot, L., Hydrolysis of mixed titanium and zirconium alkoxides by an esterification reaction, J. Solid Sture Chem., to be published; Laaziz, I., Larbot, A., Foulon, 1. D., and Cot, L., zr2Ti404(oiPr)8(oiPr)2( OAC),~, submitted for publication. 89. Francis, L. F., Payne, D. A., and Wilson, S. R., Crystal structure of a new lead zinc acetate alkoxide, PbzZnz(OCzH40Me)4(0Ac)4, Chem Murer., 2,645-647 90. 91.

(1990).

Veith,M.,Hans, J., Stahl, L., May,P.,Huch,V.,andSebald, A., Bimetallic alkoxygermanate(II)-stannate (II) andplumbate(II), Z. Nuturjiorsch. B, 46b, 403-424 (1991).

Eichorst, Howard, K. E., and Payne, D. A.,NMR investigations of lithium alkoxide solutions, J. Non-Cryst. Sol., 121, 773 (1988). 92. Sirio, C., Poncelet, O., Hubert-F'falzgraf, L. G., Daran, l. C., and Vaissermann, J., Reactions between copper P-diketonates and metal alkoxides as a route sol-to uble and volatile copper(I1) oxide precursors: Synthesis and molecular structure Polyof Cu4(p-3,nl-OC2H40iPr)4(acac)4 and (acac)Cu(p-OSiMe~)zAl(OSiMe~)~, hedron, 11, 177 (1992). 93. Turova N. Y., Turevskaya, E. P., Kozlova, N. I., Rogova, T. V., Kessler, V. G., and Kucheiko, S. I., Transition metal alkoxides-precursors of oxides materials, J. Non-cryst. Sol., in the press (1992). 94. Thule, P. P., Raghavan,S., and Risbud, S. H., Comparison of Ba(OH)2, BaO and

of bariumtitanatebythealkoxide Baasstartingmaterialsforthesynthesis method, J. Am. Cerum Soc., 70, C108 (1987). 95. Hsueh, C. C., and McCartney, M. L., Microstructural development and electrical properties of sol-gel prepared lead zirconate-titanate thin films,J. Muter, Res., 6, 96.

2208 (1991).

Yanovskaya, M.I., Turevskaya, E. P., Turova, N. Y., Dambekalne, M. Y . , Kolganova, N. V., Ivanov, S. A., Segalla, A. G., Belov, V. V., Novoselova, A. V., andVenevtsev, Y. N., Transparentceramic(Pb, La) (Zr, Ti)03 preparedby alkoxy technology, Inorg. Muter., 23, 584 (1983).

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97.Budd,K.D.,Dey, S. K.,andPayne,D.A.,Sol-gelprocessing of PbTiO3, PbZrO3, PZT and PLZT thin films, Br. Ceram. Proc., 36, 107 (1985). 98.Ramamurthi, S. D.,andPayne,D.A.,Structuralinvestigations of prehydrolized precursors used in the sol-gel processingof lead titanate,J. Am. Ceram. Soc., 73, 2547 (1990). of finePbTiO3powdersby 99.Cheng,M. J., ZhaoZ.,andQuangD.,Preparation hydrolysis of alkoxide, Chem. Mater., 3, 1006 (1991). 100. Seth, V. K., Schulze, W. A., and Condrate, R. A., Sr. Vibrational spectral characterization of a lead lanthanum zirconate titanate during various stages of solgel processing, Specrro. Lett., 24, 1299 (1991). 100a.AhlfAanger, R., Bertagnolli,H.,Ertel, T., Kolb, U., PeterD.,Nass,R.,and Schmidt, H., First evidence of the preformation of an inorganic network in solgel processing of lead zirconate titanate, obtained by EXAFS spectroscopy,Ber. Bun. Gesel. Phys. Chem., 95, 1286 (1991). 101. Milne, S. J., and Pyke, S. H., Modified sol-gel process for the production of lead titanate films, J. Am. Ceram. Soc., 74, 1407 (1991). 102. Philipps, J. F., and Milne, S. J., Diol-based sol-gel system for the production of thin films of PbTiO3, J. Muter. Chem., 1, 893 (1991). 102a.Amini,M.M.,and Sooks, M.D.,Preparation of single-phaseKNbo3using bimetallic alkoxides, Better Ceramics Through Chemistry W , 180, 675. (1990). 103. Nashimoto, K., and Cima, M. J., Epitaxial LiNbo3 thin films prepared by a solgel process, Mater. Lett., IO, 348 (1991). 104. Francis, L. F., Oh, Y. J., and Payne, D. A., Sol-gel processing and properties of leadmagnesiumniobatepowdersandthinlayers, J. Muter. Scien., 25, 5007 (1990). 105. Chaput, F., Boilot, J. P., Lejeune, M., Papiernik, R., and Hubert-Pfalzgraf, L. G., Low-temperature route to lead magnesium niobate, J. Am. Ceram. Soc., 73, 1355 (1989). 106. Kwak, B. S., Boyd, E. P., and Erbil, E., Metalorganic chemical vapor deposition of PbTiOs thin films, Appl. Phys. Lett., 53, 1702 (1988). 107. Trundle,C.,andBrierley,C. J., Precursors for thinfilmsoxidesbyPhotoMOCVD, Appl. Surf: Sci. 36, 102 (1989).

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POWDER SYNTHESIS AND CHARACTERIZATION

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3 Chemical Synthesis of Metal Oxide Powders C. N. R. Rao Indian Instituteof Science Bangalore, India

1.

INTRODUCTION

Synthesis of oxide materials provides an excellent case study of the contribution of solid-state chemists to materials synthesis [1,2]. Although tailoring oxides of the desired structure and properties remains the main goal of solid-state chemistry and materials science, it is not always possible to do so. One can evolve a rational approach to the synthesis of solids [3], but there is always the element of surprise encountered not so uncommonly. A well-known example of an oxide discovered serendipitiously is NaMo,06 containing condensed MO, octahedral metal clusters [4]. This was discovered in an effort to prepare the lithium analog of NaZn2M0308. Another such serendipitous discovery is that of the phosphorus-tungsten bronze Rb$8W320112, formed by the reaction of phosphorus present in the silica of the ampule during the preparation of the RbW bronze [5]. Novel solids of the type cu$o6s8, called Chevrel phases, were also discovered accidentally. There are many examples of rational synthesis. A good example is Sialon, in which Al and oxygen were partly substituted for Si and nitrogen in silicon nitride, Si3N,. The fast Na+ ion conductor Nasicon was synthesized based on understanding the coordination preferences of cations and the nature of oxide networks formed bythem. The zero-expansion ceramic C%.5Ti2P3012,possessing the Nasicon framework, was later synthesized based on the idea that the property of zero expansion would be exhibited by two- or three-coordination polyhedra linked in to leave substantial empty space in the network [3].

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Traditionally, oxides have been prepared by the so-called ceramic method, which involves mixing and grinding various powders and heating them at high temperatures, with intermediate grinding when necessary. The trend nowadays is to avoid such a brute force method to obtain better control of structure, stoichiometry, and phasic purity. Some of the chemical methods of synthesis of oxides that are especially noteworthy are (1) the precursor method, (2) coprecipitation and soft chemistry routes, (3) the sol-gel method, (4) intercalation and ion exchange, and (5) topochemical methods. We examine examples of synthesis by some of these methods and briefly describe the synthesis of certain novel oxide materials, including superconducting cuprates, it should be noted that a variety ofconditions, often bordering on the extreme, such as very high temperatures or pressures, very low oxygen fugacities, and rapid quenching, have been employed in oxide synthesis. The low-temperature chemical routes, however, are of greater interest.

II. PRECURSOR METHOD In the ceramic method, diffusion distances between the reacting cations are rather large. The diffusion distances are markedly reduced by incorporating the cations in the same solid precursor. Synthesis of complex oxides by the decomposition of compound precursors has been known for some time. For example, thermal decomposition of L~CO(CN)~-SH,O and LaFe(CN),.H,O in air readily yields LaCo0, and LaFeO,, respectively. BaTiO, can be prepared by the thermal decomposition of Ba[TiO(C204),], and LiCrO, can be prepared from L I [ C ~ ( C ~ O ~ ) ~ ( H , O Unfortunately, )~]. compound precursors are not always easy to find. In such instances, precursor solid solutions can be used effectively. Carbonates of such metals as Ca, Mg, Mn, Fe, CO, Zn, and Cd are all isostructural, possessing the calcite structure. We can therefore prepare a large number of carbonate solid solutions containing two or more cations in different proportions [6,7]. The rhombohedral unit cell parameter uR of the carbonate solid solutions varies systematically with the weighted mean cation radius. Carbonate solid solutions are ideal precursors for the synthesis of monoxide solid solutions of rock salt structure. The carbonates can be decomposed in vacuum or in flowing nitrogen to yield single-phase solid solutions of monoxides of the type Mnl-pxO (M = Mg, Ca, CO, or Cd) of rock salt structure. Oxide solid solutions of M = Mg, Ca, and CO would require 770-970 K for their formation; those containing cadmium are formed at even lower temperatures. The facile formation of the oxides of rock salt structure by the decomposition of calcite carbonates is a result of the close (possibly topochemical) relationship between the structures of calcite and rock salt. The monoxide solid

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solutions can be used as precursors for preparing spinels and other complex oxides. Besides monoxide solid solutions, a number of ternary and quaternary oxides possessing novel structures can be prepared by decomposing carbonate precursors containing the different cations in the required proportions, in air or oxygen. Thus, one can prepare Ca,Fe,O, and CaFe204 by heating the corresponding carbonate solid solutions in air at 1070 and 1270 K,respectively, for about 1 h.Ca,F%O, is a defect perovskite with ordered oxide ion vacancies and has the well-known brownmillerite structure (Fig. 1) with the Fe3+ ions in alternate octahedral (0)and tetrahedral (T) sites. Two new oxides of similar compositions, Ca$0,05 and CaCo204, have been prepared by decomposing the appropriate carbonate precursors in oxygen atmosphere around 940 K. Unlike Ca,Fe,O,, in C%Mn,O,, anion vacancy ordering in the perovskite structure gives a square-pyramidal (SP) coordination around the transition metal ion (Fig.1). One can also synthesize complex oxides of the type Ca2FeCo05, Ca,Fe,,Mno~405, Ca3Fe2Mn08, and so on, belonging to the A#,O,,, family by the carbonate precursor route. In the Ca-Fe-0 system, there are several other oxides, such as CaFe407, CaFel2Olg, and CaFe204(FeO), (n = 1, 2, 3), that can, in principle, be synthesized starting from the appropriate carbonate solid solutions and decomposing them in a proper atmosphere. A good example of a multistep solid-state synthesis achieved starting from carbonate solid solution precursors is provided by the Ca2Fez-$4nXO, series of oxides (Fig. 2). The structure of both the end members, Ca2F%0, and Ca2Mn,05, are derived from that of the perovskite (Fig. 1). Solid solutions be-

-

-C - L a -

b

(0)

A

-O ( b)

Figure 1 Structure of (a) C%F905 (brownmillerite) and (b) C%Mn,O,. Oxygen vacancies are shown by open circles.

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I

JI

BM

Figure 2 Ca3F%MnOs (I) is reduced topochemically to Ca2Fe4BMm05 transforms to the brownmillerite (BM) structure on annealing in vacuum.

(m; (II)

tween the two oxides are expected to show oxygen vacancy ordered superstructures with Fe3+ in octahedral and tetrahedral coordinations and Mn3+ in square-pyramidal coordination, but they cannot be prepared by the ceramic method. These solid solutions have indeed been prepared starting from the carbonate solid solutions, Ca2F~-Mnx(C03),. The carbonates decompose in air around 1200-1350 K to give perovskitelike phases Ca2Fe--$fnxOcy (y < 1). The compositions of the perovskitelike oxides prepared from the carbonate precursors with x = 2/3 and 1 are Ca3Fe+4n8 and Ca3Fel,Mnl,08.,. X-ray and electron difraction patterns show that they are members of the A,,Bn03,,-1 homologous series of anion vacancy ordered superstructures with n = 3 (A3B308+J. Careful reduction of Ca3Fe2Mn08 andCa3Fe1.5Mn1.508.1in dilute hydrogen at 600 K yields Ca3Fe2Mn07.5 and Ca2FeMn05 (CagFe1.5Mn1.507.5),respectively (Fig. 2). During this step, only Mn4+ in the parent oxides is topochemically reduced to Mn3+, and Fe3+ remains unreduced. The mostprobable superstructure of Ca3F%Mn07.5 involves SP, 0, andT polyhedra along the b direction. On heating in vacuum at 1140 K,however, it transforms to the more stable brownmillerite structure with only 0 and T coordinations (Fig. 2). A variety of complex metal oxides of perovskite and related structures can be prepared by employing hydroxide, nitrate, and cyanide solid solutions precursors as well [ 8 ] . For example, hydroxide solid solutions of the general formula Ln1-Jblx(OH)3, where Ln = La or Nd and M = Al, Cr, Fe, CO or Ni) and Lal-x-,"~"y(OH)3 (where M' = Ni and M" = CO or Cu), crystallizing in the

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65

rare earth trihydroxide structure, can be decomposed at relatively low temperatures (-870 K) to yield LaNiO,, NdNiO,, LaNil-xCox03, LaNil-xCuxO,, and others. Anhydrousalkaline-earth metal nitrates A(N03)2 (A = Ca, Sr, Ba)and Pb(NO,), are isostructural. One can therefore readily prepare nitrate solid solutions of the formula Al+.Pbx(N03)2,which are ideal precursors for the preparation of such oxides as BaPbO,, B%PbO,, and Sr2Pb0,. Oxides of the type LaFe0.5C00.503and 5Nd&003, which cannot be madereadilyby the ceramic method, have been prepared by the decomposition of cyanide solid solutions: LaFeo,Coo.5(CN)6.5.H20 and L% 5N+.5C~(CN)6.5-H20, respectively.

Ill. LA2C0205 AND La2Ni205 BYTOPOCHEMICAL REDUCTION Both LaNiO, and LaCoO, crystallize in the rhombohedral perovskite structure, and the anion-deficient nonstoichiometry of these oxides is interesting. Occurrence of the homologous series La,Ni,O,,, on the basis of a thermogravimetric study of the decomposition of LaNiO, was proposed some time ago. It was not known, however, whether a similar series exists for cobalt. Controlled reduction of LaNiO,andLaCoO, in hydrogen shows the formation of La2Ni2O5 and L%Co205, representing the n = 2 members of the homologous series LaB,03,1 (B = CO or Ni). La2Ni205.can be prepared by the reduction of LaNiO, at 600 K in pure or dilute hydrogen [9] and La2C0205 by the reduction of LaCo0, in dilute hydrogen at 670 K. Both oxides can be oxidized back to the parent perovskites at lowtemperatures. Neither La2Ni2O5 nor La2C02O5 can be prepared by the solid-state reaction of La203 and the transitionmetaloxide.X-rayand electron diffraction data reveal that La2C0205 adopts the brownmillerite structure. The x-ray diffraction pattern of La2Ni205 is different from that of La2C0205 but could be indexed on a tetragonal cell related to cubic perovskite. The formation of La2C0205 and L%Ni205 by the reduction of LaCoO, and LaNiO,, respectively, is caused by the topochemical nature of the reduction process. The reduction of the high-temperature superconductor YB%Cu30, to YBa2Cu306 is also a topochemical reaction. Similarly, many of the reactions involving insertion of atomic species into host oxides are topochemical.

IV. MO,-xW,O,BYTOPOCHEMICALDEHYDRATION Many of the modem developments in solid-state chemistry owe much to the investigations carried out in MOO, and WO, the crystallographic shear planes being a major discovery. WO, crystallizes in a Reo3-like structure, but MOO,

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possesses a layered structure. MOO, can be stabilized in the WO, structure by partly substituting tungsten for molybdenum. These solid solutions can be prepared by the ceramic method (by heating MOO, and WO, in sealed tubes around 870 K) or by the thermal decomposition of mixed ammonium metallates. These methods do not always yieldmonophasicproducts,however, owing to the difference in volatilities of MOO, and WO,. We therefore sought to prepare Mo1,Wx03 by the topochemical dehydration of Mol,Wx03.H20. Topochemical dehydration, it is to be noted, is a gentle process. Since Mo03-H20 and WO,.H,O are isostructural, the solid solutions between them are prepared readily by adding a solution of MOO, and WO, in ammonia to hot 6 M HNO,. The hydrates, Mol,Wx03.H20, crystallize in the same structure as Mo03.H20 and W03.H20, with a monoclinic unit cell. The hydrate solid solutions undergo dehydration undermild conditions (around 500 K), yielding M o ~ - ~ W ~which O ~ , crystallize in the Reo3-related structure of WO,. Thermal dehydration of W03.H20 and M o , - ~ W ~ O ~ Hgives ~ O rise to oxides of Reo, structure. MOO, obtained by the dehydration of Mo03.H20, however, has a layered structure; this reaction is reported to be topotactic. The nature of the dehydration reaction was studied by an in situ electron diffraction study in which the decomposition occurs as a result of beam heating [IO]. Electron diffraction patterns clearly show how W 0 3 - H 2 0transforms to WO, topotactically with the required orientational relationships. The mixed hydrates, M O ~ - ~ W ~ O , . H also ~ O , undergo topotactic dehydration with similar orientational relations. What is interesting is that when the dehydration of Mo03.H20 is carried out under mild conditions (e.g., electron beam heating), MOO, in the Reo, structure is formed instead of the expected layered structure. The Reo, structure of MOO, is metastable and is produced only by topotactic dehydration under mild conditions. We believe that the preparation of Reo3 like MOO, by mild chemical processing is significant.

V. BRONZES OF MO,-flxO, REACTION

BYASOLID-STATE

Oxides of the type AxW03 (A = alkali metal) are known as tungsten oxide bronzes. These are readily prepared by the insertion of the alkali metal into WO,. The corresponding molybdenum bronzes are more difficult to prepare. High pressures and electrochemical methods are generally employed to synthesize some of them. A simple solid-state reaction between an alkali iodide and MOO, or Mol,WxO3 (under dry conditions) has been found to yield such molybdenum oxide bronzes [l l]. The following reaction represents a simple means of making these bronzes:

CHEMICAL SYNTHESIS OF METAL OXIDE POWDERS

VI.

67

INTERCALATION

Intercalation chemistry has become an important aspect of solid-state chemistry, and a large variety of guest molecules are accommodated in the cavities, cages, channels, or interlayer spaces of host solids. Many oxides act as hosts, and interesting properties emerge from this property. Many of the intercalation reactions are topochemical in nature. One of the interesting intercalation reactions involving oxides is that involving chemical or electrochemical intercalation or disintercalation of alkali metal ions in oxides of the type LiMO,, where M = V or Co. When Li, for example, is removed from LiVO,, the resulting VO, remains in the metastable structure of the parent LiVO,. Other such metastable structures of oxides have been prepared by the preferential removal of intercalated species [l].

VII.

ION EXCHANGE

Ion exchange is employed effectively to synthesize new oxides [l]. Thus Na in p-alumina is readily exchanged by other cations. Such exchange reactions are well known in silicates, especially zeolites. Examples of simple ion-exchange reactions are as follows: U 1 0 2 + AgN03 p "AgA102 + KNO3 a - LiFeOz + CuCl +CuFeOz + LiCl NaCrO;! + LiN03 +LiCrOz + NaCl It has been possible to exchange Li in LiNbO, by protons to obtain HNbO,. Such alkali metal-proton exchange reactions are common in layered oxides (e.g., H2Ti307 and HLaNbO,).

VIII.

ALKALI FLUX METHOD

One of the early examples of the use of strong alkaline media for the synthesis of oxides is that of Pb2Ru2-.Pbx07y, which has the pyrochlore structure. The method stabilizes higher oxidation states of metals. Alkali carbonate fluxes have been traditionally used to prepare many oxides (e.g., LaNiO,). Recently, superconducting La,CuO, was prepared by the reaction of L%O, and CuO in a molten mixture of NaOH and KOH around 520 K [12]. Superconducting Ba,_.K$iO, has been prepared by the use of molten KOH [13].

IX.

INTERGROWTH STRUCTURES

Several systems form chemically well-defined recurrent ordered intergrowth structure with large periodicities, rather than random solid solutions with vari-

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able composition. However, the ordered intergrowth structures themselves frequently show the presence of wrong sequences. The presence of wrong sequences or lamellae is best revealed by a technique more suited to the study of local structure.High-resolution electron microscopy (HREM) enables direct examination of the extent to which a particular ordered arrangement repeats itself: the presence of different sequences of intergrowths, often of unit cell dimensions. Selected area electron diffraction, which forms an essential part of HREM, provides useful information (not generally provided by x-ray diffraction) regarding the presence of supercells caused by intergrowth or defect ordering. Many systems forming ordered intergrowth structures have come to be known in recent years [14]. These systems generally exhibit homology. If the M O 3perovskite structure is cut parallel to the (1 10) planes, slabs of the composition An-lBn03, are obtained; if these slabs are stacked, an extra sheet of A is introduced, giving rise to the family of oxides of the general formula A#n03n+2. Typical members of this family are Ca2Nb2O7 (n = 4) NaCa4Nb,017 ( n = 5), and Na,Ca,Nb,O, (n = 6). High-resolution electron microscopy and x-ray studies show that an ordered intergrowth structure with n = 4.5 with the composition NaCa8Nb903*corresponds to alternate stacking of n = 4 and n = 5 lamellae. In Fig. 3 we show the lattice image of the ordered intergrowth in NaCa8Nb903,. Between n = 4 and 4.5, a large number of or-

Figure 3 High-resolution electron microscopic image of the n = 4.5 memberinthe series.

A,,B,,O,,,

CHEMICAL SYNTHESIS OF METAL OXIDE POWDERS

69

dered solids are found, with the b parameter of the unit cell ranging anywhere from 58.6 A in the n = 4.5 compound to a few thousand angstroms in longer period structures. These solids seem to belong to the class of infinitely adaptive structures. There is a family of oxides of the general formula Bi@A,1B,03,3, discovered by Aurivillius, in which the perovskite slabs, (An-1Bn03n+1)2-,n octahedra thick, are interleaved by (Bi202)2+layers (Fig. 4). Typical members of this family are Bi2W06 (n = l), Bi3Til,Wo.,0, (n = 2), Bi4Ti3CrO12(n = 3), and Bi,Ti3Cro15 (n = 4), andtheyhavebeeninvestigatedindetailby HREM. These oxides form intergrowth structure of the general formula Bi4Arn+&3,,, +n03(rn+fl)+6; involving alternate stacking of two Aurivillius oxides with different n values; the method of preparation' simply involves heating the compo-

0 0

eo 0

eo 0 0

Bi7M5021

Bi9M7027

"

Figure 4 The f i i t three members of a homologous series structures of Bi4Am+,&lm+f103~m+fl~4 formed by the Aurivillius family of oxides, where A is also takento be Bi. Bi cations are shown as filled circles and oxide ions as open circles. BO, groups are shown in polyhedral form.

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nent oxides or carbonates of metals. Ordered intergrowth structures with (m,n) values of (1, 2), (2, 3), and (3, 4) have been fully characterized by x-ray diffraction and HREM. What is most amazing is that such intergrowth structures with long-range order are indeed formed, although either member (mor n) can exist as a stable entity. These materials seem to be truly representative of recurrent ordered intergrowth. The periodicity found in the recurrent intergrowth solids formed by the Aurivillius family of oxides is indeed impressive. The relative ease withwhich WO, forms tetragonal, hexagonal, or perovskite-type bronzes by interaction with alkali and other metals is well known. The new family ofintergrowthtungsten bronzes (ITB) involving the intergrowth of nWO, slabs and one to three strips of the hexagonal tungsten bronze (HTB) is of relevance to our discussion here. In these intergrowth tungsten bronzes of the general formula MxW03, x is generally 0.1 or less, and depending on whether the HTB strip is one or two tunnels wide, ITB are classfied as belonging to (0, n) or (1, n) series (Fig. 5). HTB strips of two-tunnel width seem to be most stable in ITB, and many ordered sequences of the (0, n) and the (1, n) series have been identified. Recently, ITB phases of Bi have been characterized, and in this system the HTB strips are always one tunnel wide, Displacement of adjacent tunnel rows caused by the tilting of WO, octahedra often results in doubling of the long-period axis of the ITB. Evidence for the ordering of the intercalating Bi atoms in the tunnels has been found in terms of satellites around the superlattice spots in the electron diffraction patterns. Among the other systems exhibiting ordered intergrowth, special mention must be made of hexagonal barium femtes MpYq (M = BaFeI2Ol9 and Y = Ba2Me,2022, where Me is Zn, Ni, Mg, and so on). A large number of intergrowth structures of this family have been identified.

X.

SUPERCONDUCTING CUPRATES

The discovery of a superconducting cuprate with a T, above 77 K created a sensation in early 1987. Wu et al., who announced this discovery first, made measurements on a mixture of oxides containing Y, Ba, and Cu. In this laboratory, we independently worked on the Y-Ba-Cu-0 system on the basis of solid-state chemistry [15]. We knew that Y,Cu04 could not be made and that substituting Y with Ba in this cuprate was not the way to proceed (unlike in La2-€!axCu04). We therefore tried to make Y3Ba3Cu5OI4by analogy with the known La3Ba3Cu6OI4and varied the Y/Ba ratio as in Y3,rBa3+xC~6014. When x = 1, we obtained YBa2Cu3 (T, 90 K). We knew the structure had to be that of a defect perovskite from the beginning, because of the route we followed for the synthesis. We briefly examine some preparative aspects of the various types of cuprate superconductors [16]. The cuprates are ordinarily made by the traditional ce2:

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71

I I ,6)

Figure 5 The (1, n) intergrowth tungsten bronzes. Hexagonal tunnels of HTB strips separate the WO, slabs shown in polyhedral form. (After Kihlborg, 1979.)

ramic method (mix, grind, and heat), which involves thorough mixing of the various oxides and/or carbonates (or any other salt) in the desired proportion and heating the mixture (preferably in pellet form) at a high temperature. The mixture is ground again after some time and reheated until the desired product is formed, as indicated by x-ray diffraction. This method may not always yield the product with the desired structure, purity, or oxygen stoichiometry. Variants of this method are often employed. For example, decomposing a mixture of nitrates has been found to yield a better product in the 123 compounds of the type YBa2Cu307 by some workers; some others prefer to use BaO, in place of BaC03 for the synthesis. The sol-gel method has been conveniently employed for the synthesis of 123 compounds, such as YBa2Cu30, and other cuprates. The sol-gel method

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provides a homogeneous dispersion of the various component metals when a solution containing the metal ions is transformed into a gel by adding an organic solvent, such as glycol or an alcohol, often in the presence of other chemicals, such as organic amines. The gel is then decomposed at relatively low temperatures to obtain the desired oxide, generally in fine particulate form. Materials prepared by such low-temperature methods may need to be annealed or heated under suitable conditions to obtain the desired oxygen stoichiometry as well as the characteristic high T,. A total of 124 cuprates of the type YBa2Cu,08, lead cuprates, and bismuth cuprates have all been made by this method; the first two are particularly difficult to make by the ceramic method. Coprecipitation of all the cations in the form of a sparingly soluble salt, such as carbonate, ina proper medium, followed by thermal decomposition, has been employed by many workers to prepare cuprates. One of the problems with the bismuth cuprates is the difficulty in obtaining phasic purity (minimizing the intergrowth of the different layered phases). The glass or the melt route has been employed to obtain better samples. The method involves preparing a glass by quenching the melt; the glass is then crystallized by heating it above the crystallization temperature. Thallium cuprates are best prepared in sealed tubes (gold or silver). Heating T1203 with a matrix of the other oxides (already heated to 1100-1200 K) in a sealed tube is preferred by some workers. It is important that thalliumcuprates are not prepared in open furnaces since T1203, which readily sublimes, is highly toxic. To obtain superconducting compositions corresponding to a particular copper content (number of Cu02 sheets) by the ceramic method, one often must start with various arbitrary compositions, especially with the T1 cuprates. The real composition of a bismuth or a thallium cuprate superconductor is not likely to be anywhere near the starting composition. The actual composition can be determined by analytical electron microscopy and other methods. Heating oxidic materials under high oxygen pressures or in flowing oxygen often becomesnecessary to attain the desired oxygen stoichiometry. Thus h 2 C u 0 4 and La&a,-,Sr,Cu2O6 heated under high oxygen pressure become superconducting,with T, of 40 and 60 K, respectively.With the 123 compounds, one of the problems is that it loses oxygen easily. It therefore becomes necessary to heat the material in an oxygen atmosphere at an appropriate temperature below the orthorhombic-tetragonal transition temperature. Oxygen s t 6 ichiometry is not a problem with the bismuth cuprates, however. The 124 superconductorswere first preparedunder high oxygenpressures. It was later found that heating the oxide or nitrate mixture in the presence of N%02 in flowing oxygen is sufficient to obtain 124 compounds.Superconducting Pb cuprates, on the other hand, can only be prepared in presence of very little oxygen(N2witha small percentage of 02).For the electron superconductor Nd2-xCexCu04, it is necessary to heat the material in an oxygen-deficient at-

CHEMICAL SYNTHESIS OF METAL OXIDE POWDERS

73

mosphere; otherwise, the electron given by Ce merely gives an oxygen-excess material. It may be best to prepare Nd2_,Ce,Cu04 by a suitable method (say, decomposition of mixed oxalates or nitrates) and then reduce it with hydrogen. Many thallium cuprates, as prepared by the sealed tube method, have excess oxygen; they become superconducting only on heating in vacuum or hydrogen. Several other novel strategies have been employed for the synthesis of superconducting cuprates. For example, a Eu-Ba-Cu alloy precursor has been oxidized to obtain EuBa$u,O, [17], and hyponitrite precursor has been employed to prepare YBa,Cu,O, [18]. More interesting in the synthetic strategies are those in which structural and bonding considerations are involved in the synthesis. One such example is the synthesis of modulation-free superconducting bismuth cuprates [19]. Superconducting bismuth cuprates, such as Bi2CaSr2Cu208, exhibit superlattice modulation. Since such modulationhad something to do with the oxygen content in the Bi-0 layers and lattice mismatch, Bi3+ was substituted partly by Pb2+ to eliminate the modulation without losing the superconductivity.

XI.

CONCLUDINGREMARKS

The area of oxide synthesis has become trulyextensive, with newertypes of materials being prepared everyday by employing a variety of novel methods [20]. We have touched on only some of the important methods, and there are many more. For example, we have not discussed the commonly employed coprecipitation method (involving the simultaneous precipitation of the cations as carbonates and oxalates, for example) followed by decomposition or the sol-gel method. This method does not always give precursor solid solutions, but coprecipitated materials on decomposition readily give the desired oxides. Another novel method of preparing oxide powders is by the combustion method, involving the spontaneous ignition of a mixture of metal nitrates in presence of a fuel, such as urea or glycine [20]. This method can be employed to prepare powders of most oxides, including dielectrics and superconductors. The use of templatemolecules to synthesizeporous solids involving silicates andphosphates is noteworthy. The ingenuity with which properties of oxides are modified drastically by appropriate substitutions or by the modification of the structure forms an important part of synthetic strategies. Other methods of interest are those involving fine particles, vapor transport, and electrochemical methods

1201.

ACKNOWLEDGMENT The author thanks the Department of Science and Technology and the Indo-European Economic Community collaborative program for support.

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REFERENCES 1. Rao, C. N. R., and Gopalakrishnan, J., New Directions in Solid State Chemistry, Cambridge University Press, London, 1989. 2. Rao, C. N. R., and Gopalakrishnan, J., Acc. Chem. Res., 20, 228 (1987). 3. Roy, R., Solid State lonics, 32133, 3, (1989). 4. Torardi, C. C., and McCarley, R. E.,J. Am. Chem. Soc., 101, 3963, (1979). 5. Giroult, J. P., Goreand, M., Labbe, P. H., and Raveau B., Acta Crystallogr., 1336, 2570, (1980). . 6. Vidyasagar, K., Gopalakrishnan, J., and Rao, C.N.R., Inorg. Chem., 23, 1206, (1984). 7. Rao, C. N. R., Gopalakrishnan, J., Vidyasagar, K., Ganguli, A. K., and Ramanan, A.,J. Mater. Res., 1, 280 (1986). 8. Vidyasagar, K., Gopalakrishnan, J., and Rao, C. N. R., J . Solid State Chem., 58, 29 (1985). 9. Vidyasagar, K., Reller, A., Gopalakrishnan, J., and Rao, C. N. R., J. Chem. Soc. Chem. Commun., 7 (1985). 10. Ganapathi, L., Ramanan, A., Gopalakrishnan, J., and Rao, C. N. R., J. Chem. Soc. Chem. Commun., 62 (1986). 11. Ganguli, A. K., Gopalakrishnan, J., and Rao, C. N. R., J. Solid State Chem., 74, 228 (1988). 12. Ham, W. K., Holland, G. F., and Stachy, A. M., J . Am. Chem. Soc., 110, 5214 (1988). 13. Schneemeyer, L. F., Thomas, J. K., and Siegrist, T., Nature, 335, 421 (1988). 14. Rao, C.N.R., and Thomas, J. M., Acc. Chem. Res., 18, 113 (1985). 15. Rao, C. N. R., Ganguly, P., Raychaudhuri, A. K., Mohan Ram, R. A., and Sreedhar, K.,Nature, 326, 856 (1987). 16. Rao, C. N.R., Nagarajan, R., and Vijayaraghavan, R., Supercond. Sci. Technol., 6, 1(1992). 17. Matsuzaki, K., Inone, A., Kimura, H., Aoki, K., and Masumoto, K.,Jpn. J . App. Phys., 26, L1310 (1987). 18. Horowitz, H. S., Mclain, S. J., Sleight, A. W., Druliner, J. D., Gai, P. L.,Van Kavelaar, M. J., Wagner, J. L., Biggs, B. D., and Poon, S. J., Science, 243, 66 (1989). 19. Manivannan, V., Gopalakrishnan, J., and Rao, C.N.R., Phys. Rev., B.43, 8686 (1991). 20. Rao, C.N.R., Mater. Sci. Eng., in print (1992).

4 Multicomponent Ceramic Powders T. Mah UES,Inc., Dayton, Ohio

E. E. Hermes Wright Paterson Air Force Base, Ohio

K. S. Mazdiyasni General Atomics, San Diego, California

1.

INTRODUCTION

The chemical processing of ceramics, especially ceramic powder syntheses, has drawn a considerable amount of attention over the past two decades. The reason for this is the demand for reliable and advanced ceramic components for high-performance applications. This increased interest in the chemical synthesis of ceramic powders is illustrated by the vast number of publications and ongoing investigations [ 1-91. This chapter focuses only on the syntheses of oxide ceramic powders. The precursors of these powders are metal-organic compounds, mainly metal alkoxides. The different varieties of ceramic powders synthesized by mixed metal alkoxide precursors are the focal points of this chapter. Numerous references on the fundamental principles of sol-gel processing are available in the literature and within other chapters of this book [10-12]. This chapter is divided into three major sections: (1) metal alkoxide precursor synthesis methods used by the present authors, (2) silicate powder syntheses, and (3) nonsilicate powder syntheses.

II. PRECURSOR SYNTHESIS In this section, only the following synthesis methods are described ammonia, metal/alcohol, metal halide/alcohol, ester exchange, and alcohol exchange. De75

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tailed synthesis techniques for metal alkoxides can be found in Metal Alkoxides by Bradley et al. [ 101.

A.

Ammonia Method

The simplest andmost economic method for the large-scale productionof alkoxides involves the addition of anhydrous metal halides to a mixture of anhydrous alcohol in a diluent (benzene, n-hexane, or toluene) in the presence of anhydrous ammonia [13,14]:

whereM is many of the transition metals, such as titanium, zirconium, hafnium, tantalum, and niobium, and R is the organic group. The transition metal tetrakis-isopropoxide is readily purified by fractional distillation or recrystallization. The removal of NH4Cl by filtration is always cumbersome and time consuming: the ammonia method may be carried out in the presence of amides or nitriles. For this particular method, the metal alkoxide separates out as the upper layer, and the ammonium chloride remains in the solution of the amide or nitile at the lower layer; thus, the filtration step is eliminated [13,15].

B.

Metal/Alcohol Reaction Method

The alkoxides of some metals of interest, such as aluminum, yttrium, and the rare earths, can be made by reaction with alcohol using a reactioncatalyst (e.g., HgC12 and HgI2) [16-181: M + ROH

HgC12 > M(OR), exothermic

+ EH2 2

where n is the valence of metal M, and R is the organic group. Most alkali or alkaline-earth metal alkoxides can be prepared by this reaction without using reaction catalysts.

C.

Metal Halide/AlcoholReaction Method

A typical example of this method is the synthesis of various alkoxides of sili-

con [19]: S i c 4 + 4ROH

exothermic

'Si(OR)4 + 4HC1

where R is the ethyl, n-propyl, isopropyl, and so on, group.

MULTICOMPONENT CERAMIC POWDERS

D.

77

EsterExchangeReaction

A valuable method for converting one alkoxide to another is an ester exchange reaction [20]. This method is particularly suited for the preparation of the tertiary butoxide from the isopropoxideand r-butyl acetate. The reaction is as follows:

M(OR)4 + 4R’OOCCH3 _ _ j M(OR’)4 + 4ROOCCH3 where M is zirconium, titanium, hafnium, and thorium, for example, R is the isopropyl group, and R’ is the tert-butyl group. Since there is a large difference between the boiling points of the esters, the fractionation is simple and quite rapid. Another advantage appears to be the lower rate of oxidation of the esters compared with that of alcohol.

E.AlcoholExchangeReaction Substitution of other branched R groups with the lower straight-chain alcohols has been carried out in an alcohol interchange reaction [lo]:

M(OR)4 + 4R‘OH-

M(OR‘)4 + 4ROH

The distillation of azeotrope drives the reaction to completion. Compounds prepared by this method [21] include the secondary pentoxides, secondary hexoxides, secondary heptoxides, tertiary butoxides, tertiary heptoxides of transition metals, such as yttrium, and also rare earth elements. Increasing the molecular weight and branching of the attached organic group result in the rise in the melting point and stability toward hydrolysis and thermal decomposition. These characteristics are very important in controlling the hydrolysis of multicomponent mixed alkoxides.

111.

ALKOXY-DERIVEDCERAMICPOWDERS

The oxide ceramic powders produced through mixed-metal alkoxides have many advantages over powders prepared conventionally; some of the advantages are lower temperature processes, higher purity, more homogeneous distribution of constituents, finer particle size, and easier compositional alteration. Most metal alkoxides exposed to moisture and/or heat cause decomposition of the alkoxide and thus provide forming methods for fine ceramic powders (notable examples for decomposition are thermal decomposition and hydrolytic decomposition). Mazdiyasni[22,23] reported the direct pyrolysis of the metal alkoxides, which form very fine ceramic powders. The overall thermal decomposition of a liquid zirconium tertiary butoxide to ZrO2 is as follows: Zr(OR)4

*

S

ZrO2 + 2ROH + olefin

MAH ET AL.

78

The thermal decomposition method is very rapid, and the products formed are volatile olefins, alcohols, and fine ceramic powders. The hydrolytic decomposition of metal alkoxides, with subsequent dehydration, has been used to form many types of fine ceramic powders. The decomposition of the metal-organic is initiated via the hydrolytic reaction of water, with subsequent thermal dehydration of the resulting precipitates. The general decomposition is described as a two-step process:

M(OR), M(OH),

-

+ nH2O +M(OH), + nR(0H) MO,/;! + "nH 2 0 2

hydrolysis

dehydration

The hydrolysis reaction usually occurs at room temperature and dehydration occurs below 600°C. resulting in the formation of very fine ceramic particles, that is, 2-5 nm [22]. This method has been successfully used to make high-purity submicrometer-sized oxides from several metal alkoxides [23,24]. Focus on multicomponent oxide powder synthesis through the two-step hydrolysis and dehydration of metal alkoxides constitutes the remainder of this chapter. When mixtures of alkoxides are hydrolyzed, the different hydrolysis and polycondensation rates of each individual alkoxide can cause complications. These complications may lead to local inhomogeneities. If one species is hydrolyzed faster than the other, which is almost always the case, then differential precipitation occurs during hydrolysis, leaving the other species substantially unreacted. Examples of these are found in the following sections of this chapter. The classes of ceramic powder are divided into two categories (silicates and nonsilicates) and are discussed separately in the following sections.

A.

Silicate Powders

Within the category of silicate powders there are two different varieties: (1) glass and/or glass-ceramics and (2) crystalline. The first variety includes various compositions of glasses and glass-ceramics (e.g., lithium aluminosilicate and magnesium aluminosilicate), and mullite and zircon are typical examples of silicates that belong to the latter variety. 1. Multicomponent Silicate Glass and Glass-Ceramics Hydrolyses of multicomponent metal alkoxides were reported by Dislich [25,26] for silicate glass and by the present authors [27,28] for glass-ceramics. Dislich [25] synthesized the eight-component glass through a complexation of mixed alkoxides, followed by hydrolysis andcondensation. The processing flowchart for the synthesis of magnesium aluminosilicate (MAS) (with Li, Zr, and Nb) glass-ceramic oxide powder, hydroxides, and their mixtures for ce-

MULTICOMPONENT CERAMIC POWDERS

79

ramic matrix composite matrix material is shown in Fig. 1 [27]. Synthesis of the individual alkoxides was accomplished using one of the appropriate techniques previously described. The individual alkoxides combined into proportions to yield the desired oxide composition. Solid alkoxides, such as Mg(OC2H5)2 and AI(OC3H7)3,are mutually soluble in liquid alkoxide, such as Si(OCzH5)4. As the metal alkoxides are combined together and refluxed, the possibility exists that some of the alkoxides may form complex mutimetal-organic ligands distributed evenly throughout the solution. Within these mixedmetal alkoxides, the formation of double alkoxide between Mg(OCzH5)z and AI(OC3H7)3 is a strong possibility. In fact, the existence of the double alkoxide, MgA12(0R)g, is well known [10,29], indicating that the distribution of the multiple-metal hydroxides, in the solution after hydrolysis, may not be completely uniform at the molecular scale as a result of the preferential hydrolysis of some metal alkoxides and/or metal-alkoxide complexes. Although it has been documented [30] that the MgA12(OR)g double alkoxide does not break down to its constituents during hydrolysis, the resulting multicomponent metal hydroxides should have, at minimum, a uniform compositional distribution.

2. Multicomponent Crystalline Silicates There are a vast number of publications about mullite (3A1203*2Si02), ranging from synthesis and processing to its applications andproperties [31,32]. The interest and importance of mullite as a candidate material for various applications has been manifest by the number of symposiums on the topic. Hy-

Figure 1

Processingflowchart for thesynthesis of magnesiumaluminosilicate.

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MAH ET AL..

drolytic decomposition of the mixed-metal alkoxide route to synthesize stoichiometric mullite powder was carried out by one of the present authors approximatelytwo decades ago [33].Aluminum tris-isopropoxide and silicon tetrakis-isopropoxide were synthesized. The mixed alkoxides were refluxed in excess isopropyl alcohol for 16 h before hydrolysis to ensure thorough mixing. The hydroxyaluminosilicate was prepared by slowly adding the alkoxide solution to ammoniated triply distilled deionized water according to the reaction 6Al(OC3H7)3 + 2Si(OC3H7)4 + xH20 H20

+

NH3

> 2A13Si(OH)l3 .xHz0 + 26C3H7OH

The resulting hydroxyaluminosilicate was repeatedly washedwith propyl alcohol and dried in vacuum at 60°C for 16 h: vacuum 2A13Si(OH)13 F 3&03 60°C

*

dry iso-

2Si02 + 13H20

At this stage of preparation, the mixed oxide was amorphous to x-ray diffraction. Transmission electron microscopy ("EM) photomicrographs of the as-prepared powders, calcined at 600°C statically for 1 and 24 h and dynamically for 24 h, are shown in Fig. 2. In Fig. 2a, needlelike crystallites of the very fine asprepared particulates are evident. The crystallite growth (agglomeration of small particles followed by rapid growth) occurs during the calcination process, leading to very large but well-defined acicular or prismatic particulates [(Fig. 2b-d)l. The as-prepared powders are extremely active, with a surface area of -550 m2/g; however, the surface area is reduced to 280 m2/g when the powders are calcined at 600°C for 1 h. High-temperature x-ray diffraction (m) studies showed that the crystallization of stoichiometric mullite took place between 1185 and 1200°C. However, the previously mentioned hydrolytic reaction is very sensitive to experimental conditions, such as the presence of acid or base catalysts, reaction temperature, and molar ratio of alkoxides to H2O. Three critical experimental parameters (pH, temperature, and reaction time) that affected the rate of hydrolysis were systematically studied by Paulick et al. (34). The volume ratio of 100 ml mixed alkoxide [Al(OC3H7)3 and tetraethylorthosilicate (TEOS)] solution to 500 m1of an H20/CH30H (3:l ratio) solution was used for all the experiments. It was found that in an acid environment (pH 2) the hydrolytic decomposition was completed within 7 h at room temperature. Stoichiometric mullite was the only product obtained, as evidenced by both elemental analysis and x-ray diffraction. Under alkaline conditions (pH lo), the degree of hydrolytic decomposition was both time and temperature dependent. The TEOS was only partially hydrolyzed even after heating at 70°C for 20 h. The oxide powder obtained was composed of 65% A1203 and 35% mullite. The degree of hydrolysis is only moderately affected by temperature, and it is not affected by the reaction time.

MULTICOMPONENT CERAMIC POWDERS

81

The difference between acid-versus base-catalyzed hydrolysis of mixed TEOS and Al(OC3H7)3 can be accounted for by the different reaction mechanisms of TEOS hydrolysis [l l]. Nucleophilic hydrolysis and condensation in alkaline solutions tend to produce highly cross-linked species, which may not be completely hydrolyzed. In contrast, in acid solution, the electrophilic reaction mechanism favors the production of weakly cross-linked species that tend to be completely hydrolyzed [35]. The specific surface area of the as-calcined mullite powder obtained under acid hydrolysis conditions is 197 m2/g. Clusters of particles develop during concentration of the hydroxide slurry. These particles polymerize on further drying, causing bound -OHgroups in the clusters, thereby resulting in hard agglomerates instead of fine powder. One can make a compromise through the partial hydrolysis of TEOS [36-391 and then add AI(OC3H7)3 to accomplish complete hydrolysis and yet achieve fine powder sizes.

B.

Nonsilicate Powders

A few varieties of nonsilicate oxide ceramic powderssynthesizedthrough

alkoxide processing are presented here. These oxide powders were developed for the use of electronic, optical, and high-temperature structural applications. For each material, we start with a brief description of the synthesis, which is followed by powder characteristics (e.g., particle sizes, morphologies, and size distributions) and densification behavior andsome properties of dense material. 1. Spinel The precursor alkoxide syntheses for spinel (MgA1204) were carried out through the metal/alcohol reaction methodusing HgC12 as acatalyst. This methodwas described in the previous section. The two metal alkoxides, Al(OC3H7)3 and Mg(OC2H5)2, were mixed together and gently heated in isopropyl alcohol until solid Mg(OC2H5)2 was completely dissolved and reacted, producing a clear solution, and then allowed to cool to room temperature [34]. During the refluxing process, magnesium and aluminum double alkoxide was formed, The formation and structural characterization of the MgA12(0R)8 double alkoxide can be found in references [10,29,30]. The double alkoxide was readily hydrolyzed by water, without acid or base catalyst. No differential hydrolytic decomposition of mixed alkoxide was found in this system to hinder the formation of ultrafine homogenous spinel powder. The XRD analysis (Fig. 3) of the ultrafine powder, obtained after calcination at 550"C, shows amorphous characteristics. The powder was calcined at 1000°C for 1 h, and XRD analysis showed broad peaks of spinel, which is a characteristic diffraction pattern for fine particles. Distinct, sharp XRD peaks of spinel can be observed after the powder is hot pressed at 1550°C. Elemental analysis of the powder calcined at 550°C gave a 1:l MgO/A1203 ratio. Emission spectrographic analysis of the as-calcined powder indicated that no measurable impurities were pre-

Figure 2 Transmissionelectronphotomicrographs of mullitepowders (a) as-prepared, (b) calcined at 6OOOC for 1 h, (c) calcined for 24 h with tumbling, and (d) calcined for 24 h without tumbling.

MULTICOMPONENT CERAMIC POWDERS

83

84

MAH ET AL.

MULTICOMPONENT CERAMIC POWDERS

85

sent. The surface area of the as-calcined powder was 260 m2/g, which corresponds to an average particle size of 6.5 nm. The transmission electron photomicrograph of the as-calcined powder (Fig. 4) showed the particle size to be of the order of 10 nm, which is in agreement with the calculated value. The average particle size of the powder, heat treated at 1000°C for 1 h, increased to about 30 nm, and also the particles were faceted. 2. Electronic Ceramics A few multicomponent electronic ceramics, fabricated through metal alkoxide processing, are described in this section. PZT and PLZT. The development of polycrystalline lead lanthanum zirconate-titanate (PLZT) electronic ceramic monoliths, which fully transmit incident light, requires methods for controlling stoichiometry, impurity content, porosity, grain size, and so on. Alkoxy-derived PLZT powders were prepared by hydrolytic decomposition of mixed-metal alkoxides [40]. The Zrand Ti alkoxides were synthesized by the ammonia method, and the lanthanum trisisopropoxide was synthesized by themetal/alcoholreaction method, which were described in the previous section. The Pb alkoxide was prepared by the reaction of anhydrous lead acetate, Pb(C2H302)2, with sodium isoamyloxide, NaOCsHl I [40]: Pb(C2H302)2 + 2NaOR

-

Pb(OR)2 + 2NaC2H302

where R is the isoamyl group. The solvated lead alkoxide solution was filtered to a clear, colorless solution. PLZT powder is typically prepared by dissolving the mixed-metal alkoxides in a mutual solvent, such as isoamyl alcohol, followed by hydrolytic decomposition to yield a nominal zirconate-titanate molar ratio of 65:35 containing 10atom%La. The preferred general formula of Pb1-~Lax(Zr,Ti~)1-~/403 is referred to in the literature[41,42]. The resulting hydroxide was washed repeatedly with high-purity water and then with isopropanol. The washed hydroxide was dried undervacuum at 60°C toyield white amorphous powder. The as-preparedPLZTpowderswere calcined at 500°C for 30 minutes to 1 h and then ground in a B4C mortar to break the large agglomerates. TEM photomicrographs of the powder calcined at 500°C for 15 min~tesand 1 h is shown in Fig. 5. The cubic symmetry of the material is quite evident; the electron diffraction patterns indicate a cubic structure, which is in agreement with the X R D analysis. Because of the rapid outward diffusion of PbO in this system, calcination at a higher temperature and for a longer time at 500°C invariably resulted in massive agglomeration or partial sintering of the particle. Thermograms (TG) for the as-prepared PLZT powders are shown in Fig. 6. The TG of the powders, unwashed and washed withisopropyl alcohol, was run inan ambient atmosphere from room temperature to 1000°C. Both types of

86

MAHETAL..

Figure 4 Transmission electron micrographsof dispersed spinel powder calcined at (a) 550°C and (b) 1000°C inair for 1 h.

Figure 5 Transmissionelectronmicrographs 500°C for (a) 15 minutes and (b) 1h.

of PLZT particulatescalcinedat

88

MAHETAL.

powder showed an initial 5% weight loss from 70 to 11O"C, caused mostly by the loss of alcohol. The additional weight loss observed within this temperature range is attributed to the loss of surface-absorbed water. The isopropanol washing was effective in removing this water, as evidenced by a difference of -15% in the weight loss in this region. Continued weight loss occurs as the temperature is increased and the remaining carbonaceous material is removed. The differential thermal analysis (DTA) curve for the as-prepared PLZT powder is also included in Fig. 6. Dehydration of the powder results in the endotherm observed at 110°C. The exothermic peak at -340°C is attributed to the nucleation and growth of the very fine as-prepared particles to large crystallites. No further peaks were observed while heating to 1200°C. When the sample was cooled and rerun, no additional peaks were noted, indicating retention of the crystalline phase as formed. The microstructure, typical of a PLZT body prepared by cold pressing and sinteringunderthestated experimental conditions, is showninFig.7a;the compact was thermally etched at 800°C for 30 minutes. The fine-grained microstructure is quite uniform, with internal and grain boundary porosity virtually nonexistent. Bodies fabricated similarly but sintered at higher temperatures

30

r

I

252 5.

S

,i

-

200 . 2 v) v)

'15 'c1 5 c I

-

_"

DTA

!2

G

I

/

K L IO O -

0/ 0

"_"". ~/clltormc-w*w€o

L

TGA

O

I

W

1

l

1

I

; 0TEMPERATURE 2 b o 2('c) ;0&&4bosso TG

Figure 6

Bzz

R

S

and

DTA of Alkoxy-derived PLZT

TG and DTA of akoxy-derived PLZT powder.

I

I

MULTICOMPONENT CERAMIC POWDERS

89

(W Figure 7 Scanning electron microscopy (SEM):typical microstructure of PLZT prepared by cold pressing and sintering at (a) 1120OC for 8 h and (b) 1220°C for 4 h.

90

MAHETAL.

(<1200"C), exhibited a microstructure with a much larger grain size (Fig. 7b) and no significant improvement in optical quality. Titanates. Various types of titanate powders, both doped and undoped, were synthesized through the simultaneous hydrolytic decomposition of appropriate metal alkoxides. One of the important concerns for the development of ferroelectric materials, such as BaTiO3, are the impurities. Although the presence of such impurities may not be a priori adverse, it is important to understand their influence on the final properties or to eliminate them completely, if necessary. The metal alkoxide processing is appropriate tosatisfy these purposes by controlled doping of certain elements to ultrapure material. Stoichiometric highpurity submicrometer BaTi03 powders were prepared through a mixed-metal alkoxide process[43]. Titanium tetrakis-tertiary amyloxide wasprepared through the alcohol exchange method:

-

T i c 4 + 4C3H70H + 4NH3 Ti(OC3H7)4 + 4CsH110H

C6H6 3 Ti(OC3H7)4 + 4 N h C 1 5°C

reflux 24 h

Ti(OCsH11)4+ 4C3H7OH

The barium bis-isopropoxide waspreparedby the metal/alcohol reaction method. Appropriate amounts of these alkoxides were dissolved in a mutual solvent, such as isopropyl alcohol or benzene, for a barium and titanium molar ratio of 1:1. The solution was refluxed for 2 h with vigorous stirring before the hydrolysis reaction. Drops of deionized triply distilled water were slowly added to the solution, which was continuously stirred. The reaction was carried out in a COzfree atmosphere. The hydrated oxide was dried in vacuum or in a dry helium atmosphere at 50°C for 12 h. At this stage the oxide was a finely divided, stoichiometric titanate with 50-150 A (maximum agglomerate size
I

Figure 8 Transmission electron micrographs of BaTi03 powder (a) as prepared and (b) calcined at 700°C.

92

MAH

W

W

P A R T I C L E diameter, ('A)

Figure 9

Particle size distribution of high-purityBaTiO3.

Figure 10 SEM microstructure of deeplyetchedBaTi03.

ETAL.

MULTICOMPONENT CERAMIC POWDERS

93

type of material. The synthesis of BaTiO3 powders doped with various rare earth oxides (e.g., La203, Nd203, and Sc2O3) has been reported [44,45]. The advantage of rare earth element doping to BaTiOs is that the grain size can be reduced considerably. A TEiM replica photomicrograph (Fig. 11) of h 2 0 3 doped BaTiOs shows the absence of pores withinthe grains andthe grain boundaries. The microstructure shows a uniform fine grain size (1-1.5 pm). Furthermore, no segregation of a second phase or impurities are visible along the grain boundaries. However, the Sc203-doped BaTiOs specimens were of relativelylowdensity (5.4 g/cm3)and contained a second-phase dendritic structure with very small bubbles characteristic of a glassy phase. 3. Yttria-Stabilized Zirconia and Hafnia

XRD analysis over the temperature range 250-120O0C, upon heating the submicrometer ZrO2 powders synthesized through hydrolytic decomposition of Zr

Figure 11 Electronmicrograph of LazO3-doped BaTi03, cold pressedandsintered at 1300°C for 4 h in air.

94

MAHETAL.

alkoxide, has been reported [46,47]. The following sequence of phases was observed as the temperature increased to 1200°C: amorphous, to metastable cubic, to metastable tetragonal, to tetragonal + monoclinic, to monoclinic + tetragonal, and finally to tetragonal. Calcination studies were undertaken to observe the growth of primary crystallites into large particles and the influence upon growth of tumbling in different atmospheres, such as air, oxygen, hydrogen, and vacuum [46-48]. Figures 12 and 13 are typical electron micrographs' showing, the effect of calcination at 500°C in air for 24 h without tumbling and calcination at 750°C in air for 24 h with tumbling, respectively. Tumbling is seen to be an effective method of controlling crystallite growth; the atmospheric conditions have a negligible effect during the presintering stage of processing these powders. Similar experiments were carried out for Hf02[49], and DTA analysis showed the nucleation and growth of small crystallites at 420"C, leading to conversion to monoclinic HfO;! at 480°C. Y203-stabilized (6-7 mol%) ZIQ and HfO2 powder production and sinteringstudieshavebeencarriedout[50,51]. Zirconium, hafnium,and yttrium alkoxides were synthesized using one of the appropriate methods described previously. Hydrolytic decomposition followed by similar procedures have already been described. The typical particle size distribution of of Y2Os-stabilized (7 mol%) Hf02 after 24 h calcination at 750°C is shown in Fig. 14. The crystallite size was measured directly from the electron photomicrographs. The powder had a mean particle size of -30 nm; 90% of the particles were less than 35 nm in diameter.

Figure 12 Transmissionelectronmicrograph of ZrO2 powder calcined at 500°C in air for 24 h without tumbling.

MULTICOMPONENT CERAMIC POWDERS

'0.1 A

m

95

,

Figure 13 Transmission electron micrograph of ZrO2powder calcined at 750°C in air for 24 h with tumbling.

Rhodes and Haag [52] studied the particle size distribution of Y2Og-stabilized (6 mol%) ZrO2,and typical cumulative particle size distribution curves for two different batches of powders are shown in Fig. 15, which is a common

96

MAHETAL.. 9991 99.8

-

99 98

80

-

60 50 40 3070

v)

0 W

t-

4 lv)

z

4

K

W

-I -1

4

I I I I I

LOT 4 5 - 6 5 -

F

LOT44-6.5-550

20-

I

l-

-I

-

90-

N

I

-

95-

W

I

IO

-

5-

I U)

;r

2I-

0.20.10.050.01a005 0.5

Figure 15 Particle size distribution of 6.5 mol% Y203-ZrOz "zyttrite."

characteristic of most ceramic powders prepared by this technique. The effect of calcination temperature on the specific surface area (measuredby the Brunauer-Emmett-Teller (BET) method) of the Y203-stabilized Hf02 is shown in Fig. 16. The as-prepared powders are extremely active, with surface areas of 210 m2/g. The surface area decreased rapidly up to 40O0C, mainly because

MULTICOMPONENT CERAMIC POWDERS

97

of the nucleation and growth of thefine particles to large crystallites. From 400 to 750"C, the surface area decreased only slightly to 120 m2/g. Above 800"C, the surface area decreased rapidly with increasing temperature to 10 m2/g at 1200°C. This curve suggests that a calcination temperature of 750°C is desirable to retain a relatively highly surface-active powder for the best sintering results. Figures 17 and 18 show the microstructures of a ZrO2 with 6 mol% Y2O3 (sintered at 1500°C for 3 h) and Hf@ with 7 mol% Y203 (sintered at 1650°C for 4 h). The electron photomicrographs show negligible internal porosity and insignificant porosity at the grain boundaries. These materials sintered to the excess of 95% of theoretical densities in a matter of few minutes at several hundred degrees lower thannecessary for the conventional ceramic powder routes. Both materials are fine grained, and the grain sizes ranged from 2 to 5 pm. Such grain sizes and microstructures ordinarily cannot be obtained for oxides as refractory as ZrO2 or HfO2, which are typically sintered at temperatures ranging from 1800 to 2200°C and firing times of 24-48 h or longer. Such a conventional sintering process makes achivement of low porosity concomitant while maintaining a fine, uniform grain size extremely difficult.

98

MAHETAL.

i

l

Figure 17

Electron micrograph (replica ' E M ) of cold-pressed and sintered 6 mol%

Y203-zr02.

4. Y3A1501-2 (Yttrium Aluminum Garnet, YAG)

Numerous reports [53-62] regarding the synthesis of YAG have appeared in the literature, and in most cases the final product was either a mixture of various yttrium aluminate phases or stoichiometric YAG of undesirably large particle size. The phase inhomogeneity in the final product, through alkoxide processing, was caused by the formation of a Y-A1 double alkoxide, which is a very stable compound and, hence, extremely resistant to hydrolysis [63]. For the Y and A1 isopropoxides, the two react to form a stable Y-A1 double isopropoxide according to the reaction Y(OiPr)3 + 3Al(OiPr)3

-

Y[Al(OiPr)&

The formation of this double alkoxide further complicates the precursor chemistryinvolved in the synthesis. Since Y and A1 isopropoxides are initially mixed in a ratio corresponding to the stoichiometry of YAG ( 3 5 = Y/Al), the reaction between the two alkoxides results in a solution containing a mixture of Y-Al double isopropoxide (1:3) and excess Y-isopropoxide. During hydrolysis, Y-isopropoxide hydrolyzes much faster than the double alkoxide, causing precipitation of Y-(OH)x to occur, which eventually leads to inhomogeneity in the final product.

MULTICOMPONENT CERAMIC POWDERS

99

I

Figure 18 Electronmicrograph(replica "EM) of cold-pressed andsintered 7 mol% Y203-Hf02.

27Al nuclear magnetic resonance ( N M R ) studies revealed that the coordination of aluminum in the mixture of yttrium and aluminum isopropoxides was predominantly tetrahedral, suggesting the formation of a Y-AI double alkoxide [63]. To avoid the double alkoxide formation, one of the isopropyl groups was replaced by ethyl acetoacetate (EAA) to reduce the reactivity of yttrium isopropoxide with aluminum isopropoxide. Equimolar quantities of yttrium isopropoxide and ethyl acetoacetate were mixed and refluxed under isopropanol for 6 h. This solution was then added to a stoichiometrically measured amount of aluminum isopropoxide. *TA1 NMR spectra of Al-isopropoxide mixed with EAA-modifiedY-isopropoxideshowed no evidence of the formation ofthe double alkoxide. Infrared spectroscopy provided convincing evidence for the substitution of the isopropyl group. Hydrolysis of the solution mixture was carried out by the dropwise addition of the alkoxide into a water (excess)-isopropanol bath at room temperature. The transparent sol was then aged for 4 h to allow condensation to occur. Fourier transforminfraredspectroscopy showed that during hydrolysis, isopropyl groups were the first to be removed.

MAHETAL..

l00

The gel obtained was then treated to evaporate the solvent, dried at 125°C for 8 h, and calcined at 950°C for 1 h. X-ray diffraction analysis of the powder obtained showed peaks corresponding to single-phase crystalline YAG. Complementing the x-ray results, DTA analysis showed only one exothermic peak occurring at 920"C, corresponding to the formation of YAG. The resulting powder is very fine (submicrometer) and, more importantly, is a single phase.

IV. SUMMARYAND CONCLUSIONS Powder synthesis using the alkoxy precursor technique exhibits processing flexibility not available in traditional high-temperature solid-state reaction of the mechanicallymixed,andmilled oxides and/or the inorganic salt constituents. With proper process control, impurities can be reduced to very low levels, yielding the desired electrical, optical, thermal, and mechanical properties. In addition, selected elements may be purposelyadded in controlled amounts to influence these properties. The alkoxide decomposition process for the preparation of single-phase and homogeneous mixed oxides of high purity and fine grain size has been demonstrated and extended to include a variety of mixed-phase systems. Extremely high purity, in excess of 99.95%, can be obtained for a wide variety of oxides. Particle size ranges below 50 nm are easily achieved while maintaining this exceptionally high level of purity for noncrystalline glasses and crystalline oxide ceramics. Because of the high reactivity of the alkoxy-derived powders, low processing temperatures can be used to densify oxide monolith or as matrices for fibrous reinforced ceramic composites with very fine grain size. In addition to processing advantages, dopants can be added more easily in the liquid phase, resulting in a composition tailored to the user's need. The control over purity and homogeneity afforded by alkoxy-derived powderprocessing allows control of the characteristics of the grain boundaries, which are likely to be central to the issues of electrical, optical, and structural properties of ceramics at service temperatures. Finally, contrary to popular belief, chemically derived powder is an economically feasible and viable method for the production of large-volume raw material for manufacturing reliable high-performance, advanced high-technology ceramic components.

REFERENCES 1. Vincenzini, P. (ed.), Ceramic Powders. Proceedings of the 5th International Meet-

ing on Modem Ceramics Technologies, Elsevier, Amsterdam, 1983. 2. Messing, G.,Mazdiyasni, K. S., McCauley, J. W., andHaber,

R. A. (eds.), Ceramic Powder Science, Advances in Ceramics, Vol. 21, American Ceramics Society, Westerville, OH, 1987.

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IO1

3. Brinker, C. J., Clark, D. E., and Ulrich D.R. 4.

5. 6. 7. 8.

9. 10. 11. 12.

(eds.), Better Ceramics Through Chemistry I , M R S Symposium Proceedings, Vol. 32, Pittsburgh, PA, 1984. Brinker, C. J., Clark, D. E., and Ulrich, D. R. (eds.), Better Ceramics Through Chemistry II, M R S Symposium Proceedings, Vol. 73, Pittsburgh, PA, 1986. Brinker, C. J., Clark, D. E., and Ulrich, D. R. (eds.), Better Ceramics Through Chemistry III, M R S Symposium Proceedings, Vol. 121, Pittsburgh, PA, 1988. Zelinski, B. J. J., Brinker, C. J., Clark, D. E., and Ulrich, D. R. (eds.), Better CeramicsThroughChemistry N,MRSSymposiumProceedings,Vol. 180, Pittsburgh, PA, 1990. Hench, L. L., and Ulrich, D. R., (eds.), Ultrastructure Processings of Ceramics, Glasses, and Composites, Wiley & Sons, New York, 1984. Hench, L. L., and Ulrich, D. R. (eds.), Science of Ceramic Chemical Processing, Wiley & Sons, New York, 1986. Mackenzie, J. D., and Ulrich, D. R. (eds.),Ultrastructure Processingof Advanced Ceramics, Wiley-Interscience, New York, 1988. Bradley,D.C.,Mehrotra,R.C.,andGaur,D.P., MetalAlkoxides, Academic Press, (London), 1978. Ller, R. K.,The Chemistry of Silica, Wiley & Sons, New York, 1979. Brinker, C. J., and Scherrer, G. W., Sol-Gel Science, Academic Press, San Diego,

1990. 13. Bradley, D. C., Mehrotra, R. C., and Wardlaw, J., J . Chem. Soc., (London), 1963 (1953). J. S.. J. Am. Ceram. Soc., 50, 532 14. Mazdiyasni, K. S., Lynch, C. T., and Smith, (1967). 15. Herman, D. F., U.S. Patents 2,654,770 and 2,655,532 (1953). 16. Mazdiyasni, K. S., Lynch, C. T., and Smith, J. S., Inorg. Chem., 5, 342 (1966). 17. Brown, L. M., and Mazdiyasni, K. S., Inorg. Chem., 9, 2783 (1970). 18. Adkins, H.,and Cox, J., J. Am. Chem. Soc., 60, 1151(1938). 19. Bradley, D. C., Mehrotra, R. C., and Wardlaw, J., J. Chem. Soc. (London), 5020 (1952). 20. Mehrotra, R. C., J. Am. Chem. Soc., 76, 2266 (1954). T., Mazdiyasni, K. S., andSmith, J. S., in Decomposition of 21. Lynch, C.

Organometallic Compounds to Refractory Ceramics, Metals and Metal Alloys(K. S. Mazdiyasni, ed.), University of Dayton Press, Dayton, OH, 1967. 22. Mazdiyasni, K. S., Ceram. Int., 8, 42 (1982). 23. Mazdiyasni, K.S., Lynch, C. T., and Smith, J. S., J. Am. Ceram. Soc.,48(7), 372 (1965). 24. Yoldas, B. E., J. Mat. Cer. Sci., 12, 1203 (1977). 25. Dislich, H., Angew. Chem. Int., Ed. Engl., 10, 363 (1971). 26. Dislich, H., in Transformation of Organometallics into Common and Exotic Materials (R. Laine, ed.),NATOAS1 Series E, No. 141, Nijhof, Dordrecht, 1988, pp. 236-249. 27. Hemes, E. E.,and Mazdiyasni, K. S., in NASA Conference Publication 2406 (J. D. Buckley, d . ) , NASA, Langley, VA, 1985, pp. 217-228.

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28. Mazdiyasni, K. S., Paulick, L. A., and Hermes, E. E., in NASA Conference Publication 2406 (J. D. Buckley, ed.), NASA, Langley, VA, 1985, pp. 191-195. 29. Mehrotra,R. C., J . Non-Crystal. Solids, 100, 1(1988). 30. Jones, K., Davies,T.J.,Emblem, H.G., andParkes, P.,in MaterialsResearch Soc. Symp. Proceedings, Vol. 73 (C. J. Brinker, D. E. Clark, and D. R. Ulrich, eds.), Materials Research Society, Pittsburgh, PA, 1986, pp. 11 1-1 16. 31. Somiya, S., Davis,R. F., andPask,J.A.(eds.), CeramicTransactions, Vol. 6, Mullite and Mullite Matrix Composites, American Ceramics Society, Westerville,

OH, 1990. 32. Symposium F MulliteProcessing,Structure,andProperties, 43rd PacificCoast Regional Meeting, Am. Cerum. Soc. Bull. 69(9), 1526 (1990). 33. Mazdiyasni, K. S., and Brown, L.M., J . Am. Cerum. Soc., 55(1l), 548(1972). 34. Paulick, L.A., Yu, Y. F., andMah, T., inAdvancesinCeramics,Vol. 21 (G. 35. 36. 37. 38. 39. 40. 41. 42. 43.

Messing, K. S. Mezdiyasni, J. W. McCauley, and R. A. Haber, eds.), American Ceramics Society, Westerville, OH, 1987, pp. 121-129. Keefer, K.D., inMaterialsResearch Soc. Symp.Proc.,Vol. 32, MaterialsResearch Society, Pittsburgh, PA, 1984, p. 15. Thomas, I. M., U.S. Patent 3,791,808(1974). Yoldas, B. E., J . Muter. Sci., 12, 1203 (1977). Yoldas, B. E., J . Muter. Sci., 14, 1843 (1979). Brinker, C. J., and Mukerjee, S. P., J . Muter. Sci., 16, 1980 (1981). Brown, L. M., and Mazdiyasni, K. S., J . Am. Cerum. Soc., 55(1l), 541(1972). Haertling, G. H., and Land, C. E., J. Am. Cerum. Soc., %(l), 1 (1971). Haertling, G.H., J . Am. Cerum. Soc., 54(6), 303 (1971). Mazdiyasni, K. S., Dolloff, R. F., and Smith, J. S., J . Am. Cerum. Soc., 52(10),

523 (1969). 44. Mazdiyasni, K. S, and Brown, L. M., J. Am. Cerum. Soc., 54(1l), 539 (1971). 45. Mazdiyasni, K. S., and Brown, L.M., Am. Cerum. Soc., 55(12), 633 (1972). 46. Mazdiyasni, K. S., Lynch, C. T., and Smith, J. S., J. Am. Cerum. Soc., 49(5), 286 (1966). 47. Phillippi, C. M., and Mazdiyasni, K. S., J . Am. Cerum. Soc., 54(5), 254 (1971). 48. Mazdiyasni, K. S., in M R S Symposium Proceedings, Vol. 32 (C. J. Brinker, D. E. Clark,and D. R. Ulrich,eds.),MaterialsResearchSociety,Pittsburgh,PA, 1984, pp. 175-186. 49. Mazdiyasni, K. S., and Brown, L. M., J. Am. Cerum. Soc., 53(1), 43 (1970). 50. Mazdiyasni,K. S., Lynch,C.T.,andSmith,J. S., J . Am. Cerum. Soc., 50(10),

532 (1967). 51. 52. 53. 54. 55. 56.

Brown, L. M., and Mazdiyasni, K. S., J. Am. Cerum. Soc., 53(1l), 590 (1970). Rhodes, W. H., and Haag, R.M., Tech. Report AFML-TR-70-209(1970). Messier, D.R.,and Gazza, G. E., Am. Cerum. Soc. Bull., 51(9), 692 (1972). De With, G., and van Dijk, H.J.A., Muter. Res. Bull., 19, 1669 (1984). Takamori, T., and David, L.D. Am. Cerum. Soc. Bull., 65(9), 1282 (1986). Hardy, A.B., Gowda,G.,McMahon, T. J.,Riman,R.E., Rhine, W. E., and

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Bowen, H.K., in UlwaswuctureProcessing of AdvancedCeramics (J.D. MacKenzieandD. R. Ulrich,eds.),WileyandSons, NewYork, 1984, pp.

407428. 57. De With, G., Philips J. Res., 42, 119 (1987). 58. Haneda, H., Watanabe, A., Matsuda, S., Sakai, T., Shirasaki, S., and Yamamura, H.,in Proc. 4th International Symposium on Science and Technology of Sinter-

ing,Tokyo,Japan (S. Somiya, M. Shimada, M. Yoshimura,and R. Watanabe. eds.), 1987, pp. 383-386. 59. Inoue, M., Otsu, H., Kominami, H., and Inui, T.,J. Am. Ceram. Soc., 74(6), 1452

(1991). A., J. Mater. Sci. Lett., 10, 101 60. Yamaguchi, O., Takeoka,K.,andHayashida, (1990). 61. Yamaguchi, O., Takeoka, K., Hirota, Takano, H., and Hayashida, A., J . Mater. Sci., 27(5), 1261 (1992). 62. Apte, P., Burke, H., and Pickup, H., J . Mater. Res., 7(3), 706 (1992). 63. Sambasivan, S., Keller, K., Sim, S., and Mah, T., J. Mater. Sci., in press (1993).

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5 Chemical Synthesis of Nonoxides Christian Russel Universitit Jena Jena, Germany

Michael Seibold OSRAM GmbH Augsburg, Germany

1.

INTRODUCTION

In the last two decades a great deal of interest has been focused on nonoxide ceramics (see, e.g., Refs. 1-5) that is mainly because of their potential use as high-temperature, high-strength materials. The main efforts have been conducted in the field of silicon carbide and nitride ceramics. In addition, aluminum nitride has attracted attention because of its high thermal conductivity, whichwillbecomeincreasinglyvitalin future substrate applications. Other nonoxides, such as boron carbide and various transition metal carbides, nitrides, and borides, are prospective materials for cutting tools and grinding media. All these nonoxides exhibit strong covalent bonding, high strength, and high hardness. They are extremely difficult to sinter because of low self-diffusion coefficients. Therefore, most of the nonoxide powders must be densified either by hot pressing or by using sintering additives. The fabrication of dense ceramics by hot pressing or hot isostatic pressing methods need extensive equipment. The high-temperature properties, however, are preserved. Sintering additives lead inmany cases to the presence ofglassy phases within the microstructure and hence harmed mechanical properties, such as reduced hightemperature strength. Until recently, the industrial production of nonoxide powders was predominantly based on two different high-temperature processes: carbothermal re105

106

RUSSEL AND SEIBOLD

duction [6,7] and direct synthesis from the elements. In the past few years, routes via the gas phase, which are also high-temperature or even plasma reactions, have increasingly been investigated. Powders synthesized at high temperatures show low sinterability caused by large crystallite sizes. Generally they cannot be pressurelessly densified without sintering additives. Moreover, routes via tractable polymer precursors have been studied. They can be used for the production of ceramic fibers, coatings, and monoliths and also for the formation of powders with tailored morphology and hence high sinteractivity. In contrast to oxides, no common reaction path for the synthesis of nonoxide ceramics via tractable precursors is yet known. Therefore, in the past few years many different synthesis reaction schemes for both nitrides and carbides have been discussed in the literature. Although silicon carbide and nitride have been under intense investigation (e.g., Ref. l), many other reactions leading to other nonoxides, such as aluminum nitride, boron nitride, boron carbide, and some transition metal carbides and nitrides, have been reported (e.g., Refs. 1-5). Polymeric compounds even allow the pressureless densification of nonoxide powders without the use of sintering additives in some cases.

II. HIGH-TEMPERATUREROUTES TO NONOXIDES Three main processing schemes for the formation of nonoxide ceramics are known, carbothermal reaction, direct synthesis, and gas-phase synthesis, which usually require reaction temperatures well above 1200°C. Carbothermal reaction as well as the direct synthesis from the elements can be varied over a wide range byusing different carbon sources, such as graphite or other carbon sources.

A.

Carbothermal Reduction

Conventional routes are high-temperature processes that usually require temperatures ranging from 1200 to 1600°C. Oxides and carbon are used as starting compounds, and the reactions are carried out in a nitrogen or inert gas atmosphere for the formation of nitrides (e.g., Refs.6 and 7) and carbides [8-IO], respectively.

It has been mentioned that this reaction can be carried out not only using elemental carbon but also withany carbon precursor, such as phenolic resins. They provide a more intensified mixing of the starting sources. Another varia-

CHEMICAL SYNTHESIS OF NONOXIDES

I07

tion of this method is the combined use of sol-gel technique together with complexation agents, such as furfuryl alcohol [IO], which act as the carbon source during calcination. This method has been used for the formation of TiN and favors much lower formation temperatures than usually required.

B.

Direct Synthesis

Another common route for the production of both nitride and carbide powders is direct synthesis from the elements. For the formation of nitrides, the elements are heated in a nitrogen atmosphere. Me + $N2

MeNx

All nitrides suitable for potential application in the field of materials science show a highly negative Gibbs free enthalpy of reaction according to Eq. (3). Kinetic hindrances, must be considered, however. This gains importance if the metal of focus exhibits a low melting point; in this case the nitridation temperatures are usually restricted below this value. Formation of aluminum nitride, for example, using this method is difficult to achieve because at temperatures belowthe melting point the nitridation rate is far too low for a commercial process. In Ref. 12, however, an interesting pathway is shown: a thin aluminum nitride layer was formed by nitridating at temperatures below the melting point at the surfaces of very finely grained aluminum powder; the temperature was subsequently raised above the melting point, and the nitridation process was completed. Notethat nitridation is a diffusion-controlled process, and hence the temperature cannot be raised until the thickness of the nitride layer around every aluminum particle prevents joining of the aluminum particles. Nitridation of other elements, such as boron, silicon, or transition metals, can easily be achieved at temperatures below the melting point. From a kinetic point of view, the particle size should be as small as possible. The surface of each particle is covered by a thin oxide layer, however, leading to oxygen contamination. In contrast to carbothermal reactions, direct nitridation does not entail an oxygen removal mechanism. Using coarse particles, however, requires remarkably higher temperatures or soaking times. For silicon the reaction rates strongly depend on the amount of impurities [13]. To obtain carbides, the elements are heated together with elemental carbon or another carbon source in an inert gas atmosphere. Me + XC ”+ MeC, Phenolic resins or many other C-containing compounds can be used for this purpose [9,10],

108

RUSSEL AND SEIBOLD

C.Gas-PhaseReactions Gas-phase reactions for the formation of nonoxide powders have been known for a long time. In the last few years many of these routes were reinvestigated and other precursors have been developed. This is mainly because of their potential to yieldvery fine grained nonoxides, withmean grain sizes in the nanometer range. Routes via the gas phase have also been intensively studied for the direct formation of coatings that is, CVD (chemical vapor deposition). All reaction paths basically considered as a CVD process can be tailored for the formation of powders; the processing parameters must be optimized. Metal halides with covalent bonding, which thus can be easily evaporated, and ammonia [14,15]are used. MeX,

4n + -NI43 3

MeN,D

+ nNH4X

(5)

The productionofsiliconnitride [16,17]andaluminumnitride [18-211 has been extensively studied; plasma reactions [21-231 were also reported. Instead of ammonia nitrogen may be used [24,25].Hydrocarbons and metal halides for the formation of carbides are worth mentioning [26,27].

111.

POLYMERIC ROUTES TO MAIN GROUP ELEMENT CARBIDES

A.

Silicon Carbide

A s mentioned, the sinterability of pure silicon carbide strongly depends on the properties of the powder used. Powders made by high-temperature routes are predominantly composed of different hexagonal Sic polytypes (e.g., 2H, 4H, and 6H). Incontrast, S i c produced by polymeric routes almost exclusively consists of p-Sic. In principle, Sic can be formed by the pyrolysis of either polysilanes or polycarbosilanes. The preparation of polysilanes is described in detail in Chap. 2 and was first reported by Kipling in 1924 [28]. The concept most commonly used now was developed by Yajima et al. [29],who used dimethylchlorosilane as a starting substance.

R

I I

Cl-Si-Cl

+Me

R Alkali metals and Me (liquid Na/K alloys are frequently cited [29-321) act as reducing agents of the chlorosilanes. It has been found that the polysilane

CHEMICAL SYNTHESIS OF NONOXIDES

109

transforms to a polycarbosilane [see Eq. (7) and Refs. 2,4, and 29-32] during thermal treatment.

In the past few years, similar reactions using numerous dichlorosilane derivatives were investigated to yield preceramic polymers. Some routes use several different types of chlorosilanes as starting materials (e.g., Refs. 33-35) Analogous to the preparation of polysilanes, polycarbosilanes can also be synthesized using a similar route. Alkyl chloroalkyl chlorosilanes are used as starting materials for this purpose [2]. R

I I

Cl"si"CH2Cl+

Na, K

(8)

R The reduction of dichlorosilanes or chloroalkyl chlorosilanes stimulates the formation of linear chains, whereas the same reaction using trichlorosilanes or similar compounds may support cross-linked products. Cross-linking can also be obtained by the use of vinyl silanes [33-351 or Si-H-modified compounds. A silylation reaction occurs at higher temperatures, also leading to branched products. Routes via polysilanes or polycarbosilanes favor products with comparably low molecular weight in many cases (Fig. 1). During calcination, high quantities of silicon compounds are evaporated and subsequently low ceramic yields are obtained. A notable weight increase in the ceramic residue is sometimes found if the pyrolysis is carried out under either reflux or pressure (see, e.g., Ref. 2). High ceramic yields are derived only with highly cross-linked polymeric products, whereas high molecular weights are not a sufficient prerequisite. A high amount of cross-linking can be achieved by using trichlorosilanes, chloroalkyl dichlorosilanes, and Si-H- and vinyl-modified compounds. For example, the reaction of ClCH2Si(CH3)2Cl with metallic potassium results in only traces of S i c [2]; the reaction of bothClCH2Si(CH3)2Cl and CH3SiCl3 yields ceramic residues of 30.8 wt% [34]. Coreduction of (CH3)3SiCl, (CH3)2SiCH=CH2, and (CH3)SiHClz with potassium allows a ceramic yield of 57 wt% [34]. In some cases, the use of dichlorosilanes containing aromatic substituents has led to an increase in ceramic product: for example, for the coreduction of(CH3)2SiC12 and (C&)2SiC12 with sodium, a ceramic yield of 54% was observed [2].

RUSSEL AND SEIBOLD

- " a

I

6

4

2

0

-2

chemical shift. ppm

Figure 1 29Si nuclear magnetic resonance spectra of different polycarbosilane precursors (molecular weight distributions Mnranging from 320 to 1750) with respect to treatment temperatures.

The pyrolysisof polysilanes and polycarbosilanes is usually carried out using inert gas (e.g., argon) as pyrolysis atmosphere. A general problem asse ciated with the pyrolytic formation of carbides is the desired stoichiometry of the calcined products: in contrast to nitrides, excess carbon cannot be evape rated during calcining; it may therefore contaminate the powders obtained as an elemental impurity and thus influences the physical, especially mechanical and electrical, properties of the sintered ceramic bodies. The volatiles evaporated during pyrolytic treatment of carbosilanes to form a network structure are H2 and CH4, and they depend onthe structure of the polycarbosilane used (Fig.2).

CHEMICAL SYNTHESIS OF NONOXIDES

112

RUSSEL AND SEIBOLD

Silicon-containing volatile compounds are predominantly SiH4 and H3SiCH4. High cross-linking, however, suppresses the evaporation of Si compounds.

B.

Boron Carbide

Some routes for the preparation of boron carbide ceramics were investigated recently. The highest ceramic yield (80 wt%) [5] was obtained using the thermal decomposition of boronylpyridine. Other routes, such as the pyrolysis of triphenyl boron or 1,2-C2BloH12,resulted in ceramic yields of only 2 and 4 wt%, respectively [4]. It should be said that many boron nitride precursors also yield boron carbide if calcined in an inert gas atmosphere.

W. POLYMERICROUTES TO MAIN GROUP ELEMENT NITRIDES A.Silicon

Nitride

Silicon nitride has found many applications in which structural features play an important role. Conventionally derived silicon nitride is fairly difficult to sinter, and pressureless sintering requires large amounts of sintering aids (e.g., 8% A1203 and 2% Y203). The fiist “chemical route”, which has been applied for the formation of silicon nitride at comparably low temperature, is diimid precipitation, already reported in 1830 [36]. During the reaction of silicon tetrachloride with liquid ammonia, the intermediate, Si(NH2)4 was presumed. A precipitation of different polymeric products [e.g., the diimid, Si(NH)2] subsequently occurs, which finally yields silicon nitride. This reaction involves a multiple-step mechanism. Simplified chemical reactions are shown in Eqs. (9) to (11) (for a more detailed description, see, e.g., Refs. 37-40). Sic14 Si(NH2)4 + 8NH3

+4 m C l

(9)

Equation (9) shows that high quantities of ammonium chloride are formed during the reaction that must be evaporated during calcining. In addition, a highly exothermic reaction is much easier to control if a solvent, such as n-hexane [38], is used. The required formation temperature of only 500°C is fairly low. Until recently, diimid precipitation has been the focused of numerous papers,

CHEMICAL SYNTHESIS OF NONOXIDES

I13

and in principle the formation of Si3N4, with much better sintering properties than conventionally produced powders, is possible. One should take into consideration that the precipitate is completely insoluble, however, and therefore, the formation of powders with controlled morphology is difficult and must be determined by the first reaction step. The route does not enable the production of fibers or coatings. Routes via polysilazanes, in contrast, lead to soluble or fusible intermediates that are prerequisites for the formation of coatings and fibers. Although a route enabling the formation of polysilazanes was already reported by Stock in 1921 [41], numerous routes, all leading to polysilazane intermediates, have been discussed in the literature only recently [4249]. Dichlorosilanes and ammonia or alkyl amines are used as starting compounds. R

I I

Cl-Si-Cl

+ 3NH3

(12)

_j

R Inanalogy to the formation of polysilanes, products with low molecular weights are obtained, usuallya mixture of cyclomers (with n = 3-6) and oligomers (with n < 10). Pyrolysis of these products thus leads to the evaporation of silicon-containing compounds and, in many cases, to almost negligible ceramic yields. Cross-linking andhigher molecular weights in principle can be obtained by using trichlorosilanes and vinyl- or Si-H-modified starting materials. The chemical reactions that occur are summarized in Eqs. (13) and (14). These reactions, in principle, might be used to tailor the polysilazane structures, similar to the chemistry of polysiloxanes

I

I

Cl-Si-Cl

+ 3nC1-Si-C1I

c1

R

I

-m

I -Si-CH=CH2+H-SiI R

F

R

R

M12R,

-NH3RCl

I I

].

N-Si N-Si-N Si-N

-NH

l I

"Si-CH2-CH2-Si-

R

(13)

I I

114

RUSSEL AND SEIBOLD

Further cross-linking reactions, leading to products of high molecular weight, have been reported by Seyferth et al. [1,50]. Thus the formation of two-dimensional linked polymers occurs if low-molecular-weight oligomers containing Si-H and N-H bonds have reacted with potassium hydride (residual potassium hydride can be removed by adding CHJ). High ceramic yields of about 85 wt% furthermore prove the necessity for cross-linked structures. Equation (15) shows the reaction mechanism proposed by Seyferth et al. [1,50]. The initial oligomer was synthesized byadding ammonia to chlorosilane and converted without further cross-linking into a negligible quantity of ceramic products during calcining.

Another method for the formation of cross-linked products is the catalytic dehydrogenation of Si-H and N-H bonds using Ru3(C0)12 as a catalyst [43].

R

I

nx H-Si-H

I

H

+ ax[RSiHNH],

_j

[RSiHNH],

(16)

Catalyst

I

I1

with R = phenyl, ethyl, or hexyl. The reaction mechanism was assumed as follows: first linear intermediates (I) and subsequently bridged, branched, or ring polymers (11) are formed (for a detailed discussion, see Ref. 43). Because of the high degree of cross-linking, high ceramic yields are obtained. Similar to polysilanes in terms of synthesis and calcination behavior, a large varietyof polysilazanes have already beenreported. The condensation of CH3NH2 with (CH3)2SiC12 or H2SiC12 resulted in ceramic yields of about 40% [47,51]. When HSiC13 and [Si(CH3)3]2NHwere used, the ceramic yield was in the range of 70-90% [44,49]. Using hydrazine N2H4 as a nitrogen donor and triethylamine instead of [Si(CH3)3]2NH ammonia, the ceramic yields were 72% with HCH3SiC12 as chlorosilane compound (Fig. 3) [47]. With the method of

115

CHEMICAL SYNTHESIS OF NONOXIDES

too W &-

cc

2 9)

W

80

60

40 20 0

200

400

600

800

l000

temperature, 'C

Figure 3 Thermogravirnetric analysis of a polysilylhydrazine precusor, pyrolyzed in nitrogen atmosphere.

Seyferth using potassium hydride to stimulate higher degrees of cross-linking, ceramic yields of 85% could be achieved [1,51]. The pyrolysis of polysilazanes results in fairly different chemical compositions of the residue; the calcination gas is varied. Although this effect strongly depends on the type of polysilazane, some more or less general conclusions can be drawn. Calcining in an inert gas atmosphere, such as argon, results in the formation of a highly carbon-contaminatedsilicon nitride or even in SiC/Si3N4 composites or solid solutions, depending on the calcination temperature. In contrast, if the calcination is carried out in ammonia or nitrogen, the product is pure silicon nitride. The formation of oxynitrides can be achieved by calcining in nitrogen atmosphere with substantial oxygen content or by oxidizing the polysilazane before or after calcining in nitrogen (Fig. 4). The main weight loss during calcination usually occurs in the temperature range between 200 and 6OO0C, strongly depending on the type of polysilazane used. For example, polyphenylsilazane exhibits its mainweight loss in a temperature range of 350-600"C; poly-(N-methy1)silazane loses weight in a two-stepprocess at temperatures of about 250 and 380°C (Fig. 5).

B. Aluminum Nitride Aluminum nitride powder synthesized by high-temperature routes can be sintered to a density of more than 97% of the theoretical density by adding calcium or yttrium compounds as sintering aids [52-541. Therefore, new processing routes to aluminum nitride predominantlyaim at powderswith lower quantities of cationic impurities. In addition, the formation of aluminum nitride coatings [S-581 or fibers [59] has become subject of extensive research.

116

RUSSEL AND SEIBOLD 8OO0C " " . = . W

/

.-5

1550°c

v)

aJ

U

.C

1700OC

40

35

30

25

20

l5

BR4GG angle, degrees

Figure 4 X-raydiffraction ( X R D ) patterns of a polysilazane-polysiloxane"hybrid" polymer with respect to different pyrolysis temperatures.

The first low-temperature route for the formation of aluminum nitride was discovered in the 1950s by Wiberg with Amberger and May [60,61]. Three-dimensionallypolycondensed iminoalanes [e.g., Al(NR)l.s] are formed using aluminum hydride and ammonia or primary organic amines as starting materials. It should be added that aluminum hydride is extremely difficult to handle

I17

CHEMICAL SYNTHESIS OF NONOXIDES 0

Q 0

00

5: 0

cr

a 0

0

200

400

600

1000

800

temperature, 'C

Figure 5 TGA profile of a phenyl-substituted polysilazane (x) compared with hexyl groups containing polymers.

because of its sensitivity to air and moisture. Nevertheless, these routes have been reinvestigated during the last few years r62-651. Fairly similar reactions can be carried out using aluminum alkyls as initial compounds [66-69]. AIR'3

+ 3NH2R

-3RH

Al(NHR)3

-1SNH2R

'A(m)1.5

(17)

with R' and R = H or alkyl. Tractable AlN precursors cannot be produced by adding aluminum chloride to ammonia or primary amines because, in contrast to silicon chloride, a substitution does not occur and different adducts (AlCl3 xNH3 with 1 5 x < 4) are formed, which can be evaporated without decomposition. If, alkali amides are used instead of ammonia or amines, however, the substitution takes place and highly polycondensed polyiminoalanes are formed [68].

-

3KNH2

+ AlCl3

-3KCl

AI(NH2)3

-1SNH3

'Al(m)1.5

(18)

Polyiminoalanes can also be synthesized using a high-pressure route, first reported by Cucinella et al. [70]. Another, fairly different route was recently reported by Schleich [71]. condensation of the starting materials, hexamethylsilazane and aluminum chloride, resulted in the formation of the dimeric and soluble compound ClAlNHSi(CHs)3. By heating to temperatures above 20O0C,ClSi(CH3)3 is evaporated and the polymeric compound (CIAINNH)n [72] is formed, which can be transformed to aluminum nitride by calcining in

RUSSEL AND SEIBOLD

118

an inert gas atmosphere. This route involves the disadvantage that silicon contamination cannot be completely avoided. Calcining polyiminoalanes usually results in the formation of aluminum nitride at comparably low temperatures. If the calcination is carried out in ammonia, carbon contamination is easier to avoid, calcining in argon usually results in the formation of black carbonaceous residues in the range of 10-20 wt% carbon. In general, crystallization is promoted by using ammonia and is prevented if argon is used.

C.

Boron Nitride

Most routes to boron nitride referred to in literature make use of borazole derivatives [4,73-801. A large variety of different ligands, R and R’, have already been mentioned.

R

I

N

R-B’

‘B-R

The easiest way to obtain boron nitride from these precursors involves only heat treatment: for example, borazoles undergo condensation reactions without previous evaporation. The ligand R should either be -NH2 or NHR, and a variety of different ligands R’ are possible. The ceramic yields are found in the range of 26% (R’ = H, R = NHCsHs [4]) and 55% (R’ = H, R = NH2 [4]). Although the starting materials themselves are not polymeric, comparably high ceramic yields are obtained. With R’ = SiR3 and R = NH2, spinnable compounds have been obtained, which might open up possibilities for the formation of boron nitride fibers. In some cases, the formation of white boron nitride with relatively a low oxygen and carbon content could be achieved [75]. Recently, the preparation ofwhiteboron nitride from polyborazylene wasreported [79,80]. Infrared spectra (Fig. 6) of pyrolyzed residues were found to be close to those reported for CVD samples of hexagonal boron nitride. For the B-chloroborazoles (R = Cl), however, the use of condensating agents is required. Either [(CH3)3Si]2NR [73,77,78] or(CH3)Si-C = CSi(CH3)3 [73] have already found use. The borazole rings are connected via BNR-B and B - M - B bonds, respectively, and trichlorosilane [73,76] is formed. In some cases, a solid of gellike consistency is obtained. Calcining of these polymers leads to relatively high degree of carbon contamination, even in an ammonia atmosphere. Calcining under inert gas often results in the formation of BN/B4C composites with desirable properties.

CHEMICAL SYNTHESIS OF NONOXIDES

I19

wavenumbers, cm”

Figure 6 Infraredspectrum of a themolyzed borazeneoligomerheated in vacuum to 275OC.

In analogy to the route of Wiberg for the formation of polyiminoalanes and subsequently of aluminum nitride [61], NaBH4, BCb, and ammonia may be used as starting materials. The intermediate formation of diborane, B2H6, was assumed [81] in this case.

The polymeric compound (BNH2),, is formed as an intermediate and may be thermally decomposed to boron nitride. The synthesis is difficult to carry out because all compounds and intermediates are extremely sensitive to air and moisture. Polymeric adducts of decaboranes, B10H12, are used in many other studies. Diphosphines or diamines (L”) [81-SS] could act as condensating agents that favor the formation of B-L”L-B groups. Most of these polymers are stable toward oxidation and hydrolysis and soluble in polar organic solvents, such as dimethylformamide ( D m or hexamethylphosphorus triamide (HMPT). During pyrolysis in either ammonia or argon, white boron nitride in 70% ceramic yieldand boron carbonitride with ceramic yields ofup to 90% are formed, respectively.

120

V.

RUSSEL AND SEIBOLD

POLYMERICROUTES TO TRANSITIONELEMENT CARBIDES AND NITRIDES

Transition element carbides and nitrides are applied as cutting tools because of their extreme hardness and wear resistance. In some cases nitrides and carbides (e.g.,of titanium) form solid solutions over the entire compositional range; other transition metal nitrides and carbides exhibit fairly different structures and are not completely soluble. Carbon contents within the range of few percentage points usually do not influence the mechanical properties of transition metal nitrides, and vice versa. Hence, completely carbon-free nitrides or nitrogen-free carbides are not required, especially for the titanium compounds. Until now, polymeric routes to transition metal nitrides or carbides have not been as numerously reported in the literature as those of the main group elements. They have been developed in most cases to produce thin coatings on various substrates or fibers. Seyferth and h4ignari [86] reported a route to titanium nitride and titanium carbide via tetrakis-dialkylamine titanium compounds. A polymeric or oligomeric product can be obtained by substituting the secondary amine by a primary amine [87,88]. The initially formed tetrakis-alkylamine polycondenses and leads a double-bridged polymer or oligomer. Calcining in argon or ammonia results in the formation of titanium carbide or titanium nitride, respectively.

I

R

I

R

I

R

I

R

Analogously to silicon nitride, titanium nitride precursors can also be prepared by the reaction of titanium tetrachloride with fluid ammonia [89,90]. This leads to the precipitation of highly polymeric products, which can be transferred to metal nitrides or carbonitrides by calcining in ammonia or an inert gas atmosphere. Zirconium and niobiumpolymers were prepared by fairly similar routes, the ammonolysis of the tetrakis-dialkylamines [88] or the chlorides [89,90] in liquid ammonia. After pyrolysis, nitrides or carbonitrides were obtained in high ceramic yields. The formation of titanium nitrides or carbonitrides via the calcination of titanium acetilides, such as Ti4(C=C)(NH)7 or NaTi4(W)(NH)7NH2 at temperatures I 8OOOC was reported by Maya [91].

CHEMICAL SYNTHESIS OF NONOXIDES

121

It should be mentioned that sol-gel routes have also been described in the literature, which finally yield titanium nitride [l l]. Additives, such as furfuryl alcohols, are used in this case, which decompose during the calcination in ammonia and form amorphous carbon. Subsequently carbothermic reactions occur, and finally titanium nitride is formed.

VI.

ELECTROCHEMICAL ROUTES TO CARBIDES AND NITRIDES

A large variety of precursors have been prepared lately with an electrochemical route [92-991. The metal is anodically dissolved in an organic electrolyte, which can be directly converted into a ceramic coating [55-581 or powder without further purification. The method was first reported by Seibold and Russel [92] for the formation of aluminum nitride, who used an electrolyte consisting of propylamine, acetonitrile, and tetrabutylammoniumbromide as a supporting electrolyte to achieve sufficient electrical conductivity. The electrodes were formed by sheets of various metals. The assumed chemical reactions at the electrodes are as follows: Me xNH2R + xexNH2R + Me

-

MeX++ exNHR- + H2

+Me(=),

x2

+ ;H2

The metal is oxidized at the anode, the propylamine is reduced at the cathode, and the corresponding anion and gaseous hydrogen are formed. As an intermediate, monomeric Me(NHR), is assumed. These intermediates condense to oligomeric or polymeric compounds.

The electrolytes were dried and subsequently calcined. The main weight loss occurs at temperatures in the range of 150-350°C (e.g.,Ref.92).A further slight weight loss occurred up to 550"C, and then the weight remained nearly constant. It should be noted that the thermogravimetric analysis TGA profiles were fairly similar and only slightly influenced by the type of metal. The ceramic yields, however, strongly depended on the type of metal used and were in the range of 20-50 wt% [92,96,98,99].

I22

RUSSEL AND SEIBOLD

Figure 7 XRD patterns of aluminum nitride precursors calcined in ammonia at differenttemperatures: (a) driedprecursor; (b) 600°C; (c) 800OC; (d) 1100°C; and (e) 1400°C.

Allprecursors are amorphous up to calcination temperatures ofaround 600°C. At higher temperatures, in most cases powders with extremely small crystallite sizes of around 2 0 4 0 nm are formed (Fig. 7). A further increase in calcination temperature promotes crystal growth.Withaluminum nitride, a white powder with a low oxygen and carbon content is obtained [97]. Other main group element precursors exhibit fairly different behaviors: Mg and Ca precursors yield metal cyanamide [99]. Calcination of the transition element precursors (Fig. 8) results in the formation of nitrides, carbonitrides, or carbides. For the titanium-containing precursors, TiN/TiC solid solutions can be obtained [96]; the quantity of carbon strongly depends on the calcination atmosphere applied (argon, 31 wt%; ammonia, 5.1 wt%). In contrast, chromium-containing precursors are pyrolyzed in ammonia to CrN; calcining in argon leads to the formation of Cr3C2 [98]. Other products formed are ZrN, TaC, NbC, MoC, and WC. Y-containing precursors lead to the formation of YNTy0.45Co.55 mixtures (Table 1) [99]. Because of the small

CHEMICAL SYNTHESIS OF NONOXIDES

70 *

60

50

40

30

123

20

28 (CUL)P

Figure 8 XRD patterns of various ceramicproductsobtainedaftercalcining:(a) CaCN2; (b) CrN (c) Cr3C2; (d) YN (open circles); Yo.45Co.55 (closed circles); and (e) TaC.

crystallite sizes, the powders obtained exhibit veryhighsinterability. A s already shown for aluminum nitride powders, pressureless sintering without the addition of sintering aids at 1750°C results in densities > 98% of theoretical density (t.d.) [92]. The oxygen contents of the pyrolyzed powders were in the range of 1.5-2.0 wt% for aluminum nitride, calcium cyanamide, and the transition metal compounds investigated. The electrochemical synthesis of polymeric precursors is characterized by a broad versatility. Most of the metals can be handled easily and electrolytically dissolved and transferred into polymericprecursors; thus this processing scheme seems to be a general approach and an alternative to already known methods. In addition, all precursor solutions can be mixed with one another, and therefore a wide range of new composites or solid solutions is possible in the near future.

124

RUSSEL AND SEIBOLD

Table 1 NonoxidesDerivedbyElectrochemicallySynthesizedPrecursors Metal dissolved Ceramic yield

(%)

AI Ti

42 20

zr

41 40

Cr

Ta Ca Mg Y

31 34 21 39

Calcination gas Product NH3 NH3 N2 NH3 NH3 N2

AIN TiN Ti(C, N) ZrN CrN

Ar

C4C2 TaC CaCN2 MgCN2

NH3 NH3 NH3 NH3

cm2

m,y0.45c0.55

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SYNTHESIS CHEMICAL.

OF NONOXIDES

127

71. Schleich, D. M., U.S. Patent 4,767,607 (1988). J . Mater.Sci.Lett., 9, 222-224 72. Riedel, R., Petzow,G.,andKlingebielU., (1990). 73. Maya, L., and Angelini, P., J. Am. Ceram. Soc., 73, 297 (1990). R., Ceram. Eng. Sci.Proc., 1171 74. Bender,B. A., Rice,R. W., andSpann,J. (1985). 75. Tanigushi, I., Harada, K., and Maeda, T., Japan Kokai 76 53OOO (1986). 76. Narula, C. K., Paine, R. T., and Schaeffer, R., in (C. I. Brinker, D. E. Clark, and D.R.Ulrich,eds.), M R S Pr. Vol. 73, BetterCeramicsThroughChemistry I1 Pittsburgh, PA, 1986, pp. 383-388. Chem. Mater., 3, 77. Paciorek,K. J.L., Nakahara,J. M., andHoferkamp,L.A., 83-87(1991). 78. Rye, R. R., Tallant, D. R., Borek, T. T., Lindquist, D. A., and Paine, R. T., Chem. Mater., 3, 286-293 (1991). 79. Fazen, P. J., Remsen, E. E., and Sneddon, L. G., Polymer Preprints, 32, 544-545 (1991). 80. Paine, R. T., Janik, J.F., Borek, T. T., Lindquist, D. A., Duesler, E. N., Smith, D. M., Kodas, T. T., and Datye, A. K., Polymer Preprints, 32, 546-547 (1991). W., in M R S Proc.Vol. 121, BetterCeramics 81. Seyferth, D., andSmithRees, 82. 83. 84.

85. 86. 87. 88. 89. 90. 91. 92. 93. 94. 95.

Through Chemistry Ill (C. J. Brinker, D. E. Clark, and D. R. Ulrich, eds.),Pittsburgh, PA, 1986 pp. 449-454. Seyferth, D., and Rees, W. S., Chem. Mater., 3, 1106-1 116 (1991). Schroeder, H., Reiner, J. R., and Knowles, T. A., Inorg. Chem., 2, 393 (1963). Niedenzu, K., and Buschbeck, K.-C., Gmelin Handbook of Inorganic Chemistry, 8th ed., Vol. 54, Boron Compounds: B-H Compounds, Springer, Berlin, 1979, pp. 151ff. Maya, L., in Better Ceramics Through Chemistry III (C. J. Brinker, D. E. Clark, and D. R. Ulrich, eds.) M R S hoc. Vol. 121, Pittsburgh, PA, 1988, pp. 455-460. Seyferth, D., and Mignani, G., J. Mater. Sci. Lett. 7, 487-488 (1988). Lappert, M. F., Power, P. P., Sanger, A. R., and Srivastava, R. C.,Metal and Metalloid Amides, Ellis Horwood, Chinchester, 1980. Brown, G. M., and Maya, L., J . Am. Ceram. Soc., 71, 78-82 (1988). Maya, L., Inorg. Chem., 25,4213-4217 (1986). Maya, L., Inorg. Chem., 26, 1459-1462 (1987). Maya, L.,in M R S Symp. Proc. Vol. 73, Better Ceramics Through Chemistry I1 (C. J. Brinker, D. F. Clark, and D. R. Ulrich, eds. Pittsburgh, PA, 1986, p. 401. Seibold, M., andRussel, C., inMRSSymp. Proc. Vol. 121, BetterCeramics Through Chemistry I l l (C. J. Brinker, D. F. Clark, and D. R. Ulrich, eds.), Piasburgh, PA, 1988, pp. 477-482. Seibold, M., Viemeusel, U., and Russel, C., in Ceramic Powder Processing Seience (H.Hausner,G.C.Messing,and S. Hirano,eds.),DeutscheKeramische Gesellschaft, Koln, 1989, pp. 173-179. Seibold, M., and Russel, C., J . Am. Ceram. Soc., 72, 1503-1505, (1989). Russel,C.,Hofmann,T.,Kulig,M.,andSeibold., Silikattechnik, 40, 425-429, (1989).

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6 Techniques for Characterization of Advanced Ceramic Powders S. G. Malghan, P. S. Wang, and V. A. Hackley National Institute of Standards and Technology Gaithersburg, Maryland

1. IMPORTANCE OF POWDERCHARACTERIZATION The practical performance of a ceramic component, the microstructure of the ceramic, and powder behavior during processing are strongly dependent on the physical and chemical characteristics of starting powders [l]. In addition, in the manufacture of advanced ceramic components, detailed information on powder characteristics is required to achieve reproducibility and cost competitiveness. The choice of starting powder characteristics depends on the intended microstructure and application of the final ceramics. For instance, although total oxygen may be one of the most important measurements in silicon nitride powders because oxygen participates in the formation of a grain boundary silicate phase, the total Y2O3 content may be the most desired parameter in zirconia powder processing because it controls phase composition and transformation. The requirements of powder characterization are complicated by the fact that a large number of properties must be defined to understand its characteristics completely. A review of Table 1 shows that this task is not simple because the powders are complex materials consisting of numerous features. Therefore, ceramists are faced with the question of what characteristics of powders to measure for a given process. Commonly, an answer to this question lies in the development of parametric relationships between the powder properties and their influence on the performance of the ceramic. The next question is how to measure the powder characteristics. A program has been in progress 129

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Table 1 Properties of AdvancedCeramic Powders That Constitute Their Characterization Physical properties Specific surface area Primary particle size and size distribution Agglomerate size and size distribution Porosity, total quantity and pore size distribution Density Phase composition Crystalline phases, quantity and identification Amorphous material, quantity Chemical composition Major element concentration Minor impurities (10 ppm to =l%) Trace impurities (10 ppm) Inorganic elements Organic elements Composition of impurities Surface composition Major elements Minor elements Trace elements Inorganic species Organic species Crystallinity

since 1985 under the auspices of the International Energy Agency to develop procedures for the analysis of ceramic powders [2].An overall goal in this task is to define procedures that provide repeatable and reproducible data. The development of standard reference materials (SRM) stands at the forefront of this task since the SRM are necessary for the improvement of measurement quality [3]. In this chapter, we address the available methods for the characterization of ceramic powders, and selected methods have been briefly described. Primarily, for these methods, a description of operating principles, type of data to be obtained, and measurement limitations are presented. In addition, nuclear magnetic resonance (NMR) and x-ray photoelectron spectroscopy ( X P S ) techniques have been described in more detail because their treatment, as applied to powders, is not available elsewhere. Thermal techniques, such as differential thermal analysis, are not covered because of their specialized use on powders. For interested readers, a number of references have been cited for additional information.

CHARACTERIZATION OF ADVANCED CERAMIC

POWDERS

I31

II. PHYSICAL PROPERTIES The major physical properties of ceramic powders constitute size distribution of primary particles and agglomerates, specific surface area, density, porosity, and morphology (e.g., shape, texture, and angularity).

A.SizeDistribution In the measurement of particle or agglomerate size distribution, the major distinction is that agglomerates are made of primary particles. Hence, often there is a need to determine the true size of the underlying primary particles. In such cases, the size distribution of particles may be determined by a number of techniques, as shown in Table 2, which lists the commonly used methods [3]. With

Table 2 Commonly Used Methods of Particle Size Analysis, Nominal Size Ranges, and Measurement Parameters Nominal particle Measurement Method Coarse particles > 10 pm, sieving Dry Wet Fine particles < 10 pm Field scanning Optical microscopy Electron microscopy Gravity sedimentation Pipette Photoextinction X-ray absorption Radiation scattering, Laser diffraction, scattering Stream scanning Resisitivity Optical Ultrasonic attenuation Column hydrodynamic chromatography Sedimentation fieldflow fractionation Laser Doppler velocimetry Centrifugal sedimentation Photoextinction Mass accumulation X-ray absorption

>l0

Geometric

>2 0.5-1000

Image

0.01-10 1-100 0.5-100 0.1-130

Stokes

0.03-900

Geometric

0.05-500

Dynamic/Stokes

1-500 1 0 0 1 m 0.1-1 .o 0.01-1.0 0.01-3.0 0.05-100 0.05-25 0.1-5

Dynamic/Stokes

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such a large number of choices, the selection of a system for a given application depends on a number of criteria, including the size range of the application, throughput of the instrument, accuracy, precision, reproducibility, resolution, versatility, cost, and manufacturer support [4]. The three methods described in the following text are microscopy, gravity sedimentation, and light diffraction. Their selection is based purely on their technical diversity.

1. Microscopy Three commonly used direct viewing techniques are optical, scanning electron, and transmission electron microscopy. Optical microscopy is used in the size range 1-150 pm for the determination of morphology, agglomeration, and size distribution using automated counting devices. Despite significant improvements in optics, a major disadvantage is its depth of focus, which is about 0.5 pm at x1000 [5]. Scanning electron microscopy (SEM) is a versatile technique in which a beam of electrons at 5-50 keV scans the specimen surface. The resultingx-rays, backscattered electrons, andsecondary electrons are detected and analyzed by a number of techniques. Magnifications of up to ~100,000can be achieved at resolutions finer than 20 nm. Depth of focus and magnification are inversely proportional. Automated image analysis systems are available to analyze particle size distribution.However, the major issues remain sample preparationand dispersion of powder, that is, particles tend to agglomerate when deposited on SEM stubs and thereby do not provide complete and true data for primary particles. Intransmission electron microscopy (TEM), 5-0.001 pm particles deposited on thin membranes are examined at 10-100 times better resolution than that in SEM. TEM is more often used as an analytical tool for surface characterization and texture rather than for routine particle size distribution. 2. Gravity Sedimentation The underlying principle behind the gravity sedimentation method is Stokes’ law, which describes the relationship between the settling velocity of particles in a fluid medium of known density and viscosity [6]. A number of instruments are available in which the settling velocity of particles is measured by x-ray absorption, light absorption, and density changes. One such instrument is the Sedigraph* by the Micromeritics Corporation in which the particle size distribution is determined by x-ray absorption [6]. This instrument offers an excel-

*Certain trade names and company products are mentioned in the text or identified in illustrations to specify adequately the experimental procedure and equipment used. In no case does such identification imply recommendation or endorsement by the National Institute of Standards and Technology, nor does it imply that the products are necessarily the best available for the purpose.

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lent measurement capability in the 0.2-50 pm range of a relatively concentrated suspension containing up to 6% by weight powder in an appropriate dispersion fluid. The data are presented as a continuous plot of percentage weight Of particles smaller than the stated size or other convenient formats. heparation of an appropriate dispersion is a key step in obtaining repeatable data [7]. The dispersion preparation for a given powder requires two Sets of data:(1) the amount and intensity of energy application to break agglomerates into primary particles; and (2) the pH at which the isoelectric point, the point at which the particles carry a net zero charge, occurs in a given aqueous solvent containing a surfactant for stabilization of particles. For each powder, it is necessary to develop an appropriate data for these parameters. 3. Light Scattering Instruments based on light scattering employ a laser beam to radiate particles in a stream, and the resulting scattered light is analyzed using the Fraunhofer and Mie theories to obtain size distribution data [g]. When the size of particles is very small with respect to the wavelength of light or when the refractive index of the particles is very close to that of the dispersion medium, such as that of a liquid to form a suspension, the relatively simple equations of Raleigh and Gans can be employed. For colloidal suspensions, however, it is necessary to resort to the Mie theory, which includes the restrictive conditions that the particles be spherical and intrinsically isotropic. With the advent of desktop computers, the application of the Mie theory has become a reality for polydisperse systems. Mie formulated equations to delineate light scattering by particles by considering electrical fields within and outside each particle. The equations for scattering intensity for each scattering angle involved refractive index difference between particle and suspending medium, the wavelength of incident light, and the spherical diameter of the particle. For particle diameter in the range of %o-10 times the wavelength of incident light, scattering from different portions of the particle is out of phase, which results in interference and reduced intensity. The net effect is an angular distribution of scattered light in the forward direction. This information is utilized in Mie theory computations. A number of instruments are available by which 0.03-900 pm particles can be measured in a matter of several minutes from dispersed suspensions. One major advantage of these systems is their ability to handle particulate systems containing different densities and refractive indices. A requirement of these instruments is the refractive index data, which may not be readily available for many powders. If the light-scattering particulate systems are not monodisperse, the particle size distribution obtained from these systems represents an average value, related but not identical to the weight-average particle size distribution. These averages are expected to differ depending on particle shape, particle size distribution, and degree of anisotropy.

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B.

Specific SurfaceArea

Particles consist of both internal and external surface area. The external surface area represents that caused by exterior topography, whereas the internal surface area measures that caused by microcracks, capillaries, and closed voids inside .the particles. Since the chosen surface area technique should relate to the ultimate use of the data, not all techniques are useful for fine powders. The commonly used approaches are permeametry and gas adsorption according to the Brunauer, Emmet, and Teller (BET) equation [g]. Because of simplicity of operation and speed of operation, permeametry methods have received much attention. The permeametry apparatus consists of .a chamber for placing the material to be measured and a device to force fluid to flow through the powder bed. The pressure drop and rate of flow across the powder bed are measured and related to an average particle size and surface area. Especially for porous powders, permeametry data include some internal surface area, thus decreasing their value. In the BET method, the volume or weight of adsorbed gas as a function of partial pressure is measured. Most commonly, NZgas is adsorbed by the powder at liquid N2 temperature from a gas stream of NZand He. The NZadsorbed and later desorbed is measured by thermal conductivity in one type of equipment. The BET equation describes the equilibrium between vapor and adsorbate in which multilayer adsorption occurs. The BET equation is typically written in the form PIP0

._

1 '

+

"

V(1- PIP,)

V,C

(C - 1)P& vmc

where PIP0 is relative pressure (the ratio of the equilibrium pressure to the saturation pressure at a specific temperature), V is the amount of gas adsorbed, and Vm represents V for monolayer coverage. The constant C is related to the energies of adsorption and gas liquefaction. Sample preparation is the key to obtaining reproducible data. Sufficient outgassing, minimal surface contamination, and the absence of microporosity are important aspects of the proper analysis of surface area. BET measurements are carried out in either singlepoint or multipoint mode, in which multipoint data are normally higher and more representative than those of the single point [2].

C.

Density

Theoretical, true, and tap density are three types of densities associated with ceramic powders. Theoretical density is determined by atomic composition and lattice parameters; tap density is determined by a prescribed procedure, which

CHARACTERIZATION OF ADVANCED CERAMIC POWDERS

I35

consists of measuring the volume of a known weight of powder after a certain number of mechanical taps. Therefore, tap density provides a relative measure of fill density or degree of compactibility of the powder. In the determination of true density, the powder volume is determined by the volume of gas displaced by a known weight of powder. Helium pycnometry is the commonly used method for this application. Helium is the most frequently used gas because of its inertness and small size, which enables it to penetrate even the smallest pores. Therefore, contribution tothe volume by pores can beaccounted for the measurement of true density.

D. Porosity Porosity in a powder can come from both closed and open pores. Two primary methods for the determination ofporosityare gas adsorption andmercury porosimetry. It is assumed that gas adsorption is favored in small capillaries because ofthe overlapping surface potentials, which result in capillary condensation.Pores from 1.5 to 100 nm in diameter are determined using the Kelvin equation, which relates capillary radius r to the ratio of the vapor pressure P and the equilibrium vapor pressure of the same liquid over a plane surface Po, as follows:

where y is the surface tension of the liquid, V is the molar volume of the liquid, 8 is the contact angle between the liquid and the wall, R is the gas constant, and T is the absolute temperature [lo]. Adsorption in the microporosity range ( r c 1.5 nm) often exhibits a Langmuir typeof isotherm, normally a characteristic of monolayer adsorption. The second principal method for porosity measurement is mercury porosimetry, in which a nonwetting liquid, such as mercury, is forced into capillaries of radius r. The force Foutcaused by interfacial tension y is calculated as F , ~ = 2 m y COS e

(3)

where 8 is the contact angle for mercury. The applied pressure P is related to the force driving the mercury Fin into a capillary and is given by -Fin

= P m2

Equating the two forces results in the Washburn equation [l l]:

(4)

MALGHAN ET AL.

136

Mercury porosity measurements are carried out in the pressure range 14 Pa (ambient) to 415 Pa, which corresponds to pore radii from 2 nm to 200 pm.

E.

Morphology

The morphological analysis of particles constitutes the measurement of size, shape, and texture, which describes the surface profile of a particle image. The primary methods of morphology analysis are microscopy techniques, depending on the size of particles. In recent years, a large number of software packages for processing image data fromautomated microscopes have enabled faster accumulation and interpretation of data. However, the major drawbacks of microscopic techniques are representativeness of the sample examined and two-dimensional examination of the particle surface. The data from the shortest dimension may not be adequately represented.

111.

BULKCHEMICAL COMPOSITION

A large number of techniques are available, depending on the specific need to

analyze major chemical components, minor components, nonmetallic impurities, and metallic impurities (see Table 3). Further, the applicability of a technique depends on the concentration of impurities in the powder, a problem of considerable magnitude for ultrapure powders. A generalproblemwith the techniques'that require dissolution of the powders is their resistance to chemical attack and lack of complete dissolution. The two major methods of placing the powders in solution are acid dissolution and flux decomposition followed by dissolution [2,3]. The nonavailability of standard reference powders and procedures is a serious drawback to all the analysis methods. It is not possible to cover all the available techniques; therefore, only selected techniques with applicability to bulk and minor impurities are described here.

A.

InductivelyandDirectlyCoupledPlasma

These techniques, used for qualitative and quantitative measurements, utilize excitation of atoms in a plasma to obtain their characteristic emissions that are analyzed by photodetection. The powder is dissolved, and the resulting solution is injected into the plasma. The fast speed of analysis and applicability to a wide range of elements are strong points, but these methods are subject to errors resulting from powder dissolution. Currently, a number of industrial laboratories are addressing this issue.

B.

Atomic AbsorptionSpectroscopy (AAS)

A A S is a versatile method in which an analyte is atomized in a flame, thus emitting spectral lines. The characteristic spectral lines that correspond to the

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137

Table 3 Methods for BulkChemicalandImpurityAnalysis of Metallic and Nonmetallic Impurities in Ceramic Powders Bulk chemical analysis X-ray fluorescence spectroscopy Atomic absorption spectroscopy Inductively coupled plasma emission spectroscopy Direct-current plasma emission spectroscopy Arc emission spectroscopy Gravimetry Combustion Kjeldahl Impurities Neutron activation analysis Mass spectrometry Electrochemical Coulometry Selective-ion potentiometry Potentiometric titration Argentometric Ion chromatography Nuclear magnetic resonance Electron paramagnetic resonance energy required for an electronic transition from the ground state to an excited state are emitted. The absorption of radiation from the light source depends on the population of the ground state, which is directly proportional to solution concentration. The absorption is measured by the difference in transmitted signal in the presence and absence of the test element [ 121. However, this technique is also subject to errors resulting from powder dissolution.

C.

X-rayFluorescenceSpectroscopy

The powder in the form of fine particulates or dissolved in solution is excited in an x-ray source, and the characteristic fluorescence intensity is analyzed for qualitative and quantitative data. Standards are required for the conversion of x-rayintensities to absolute concentrations. The preparation of appropriate standards that include background similar to that in the solution tobe analyzed is critical to obtaining accurate data.

D.

Nuclear Magnetic Resonance and Electron Paramagnetic Resonance (EPR)

In principle, NMR utilizes the atomic nuclear transition, induced by radio frequency irradiation, between two quantized nuclear energy levels in a magnetic

138

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AL.

field. This transition caused by resonance in a magnetic field occurs in all atomic nuclei, except in atoms with even numbers of both mass and atomic number. The nuclear spin numbers are zero for atoms withboth mass and atomic numbers that are even, and consequently, nuclear energy splitting does not occur in a magnetic field and these atoms are not active in NMR. All other isotopes in the periodic table are NMR active. ‘H, W , 15N,19F,27Al, W i , 31P, 47Ti, 63Cu, 69Ga, 89Y, and 91Zr are some examples. For an atom with a nuclear spin number I , there are 2 I + 1 quantized nuclear energy levels: -I, -(I - l), . .. (I - l), I. These levels are practically equal in energy outside a magnetic field. If a magnetic field is applied to this atom, however, the nuclear energy split into 2 I + 1 levels separated by (yh/27c)H, where H is the magnetic field strength, h is the Planck constant, and y is the nuclear magnetogyric ratio. The nuclear energy E is expressed by

E=--

rh m,H 2x

where MI = I, (I - l), .. . , -(I - l), -I. If a sample is irradiated by a radiofrequency energy equal to the energy separation, resonance occurs, the nuclei at the ground state are excited to a higher energy state, and an absorption signal is observed. The NMR is therefore an absorption spectroscopy, not an emission spectroscopy (e.g., X P S ) . The magnetogyric ratio is a constant for a specific isotope of interest. However, the magnetic field H is an effective field to the nuclei including the applied external field, the field induced by the electrons around the nuclei, and even the field produced by other parts of the molecule. For Si atoms, for example, 95.3 atom% are 28Si, which have a nuclear spin I = 0 and are not NMR active ( 2 I + 1 = 1, only one energy level; therefore, no resonance can occur). There are 4.7 atom% 29Si in natural abundance, however, and 29Si has a nuclear spin I = H. In an external magnetic field of 9.4 T, for example, the 29Si nuclear energy is split into two levels separated by approximately 79.5 MHz, with mr = H and -H and with H lower in energy. In addition, the actual magnetic field experienced by the 29Si nuclei in Sic, Si3N4, and Si02 is different and not exactly 9.4 T because the electron density and distribution around the Si nuclei in these compounds are different. This difference makes it possible for us to observe the absorption signals at different field strengths and is the “chemical shift.” Chemical shift is a comparative value and is most often referenced to tetramethylsilane [TMS, Si(CH3)4]. This compound can be used as a reference to three frequencies: ‘H, W , and 29Si. The frequency difference between the sample and that of TMS is often divided by the frequency of the spectrometer to give a value in parts per million (ppm), which is spectrometric frequency independent. Other factors, such as spin-spin interaction with neighboring atoms, affect Si nuclear energy level splitting. For nuclei with I 2 1, the nuclear quadrupole

CHARACTERIZATION OF ADVANCED CERAMIC POWDERS

139

effect also alters the energy level splitting. A more in-depth coverage of N M R theory is found elsewhere [13,14]. N M R spectroscopists are also interested in nuclear spin relaxation times. The relaxation time measures the time required for an excited nucleus to return to the ground state. Two types of relaxation times are involved: spin-lattice relaxation time T i , the time constant for thermal equilibrium between the nuclei and crystal lattice, and spin-spin relaxation time T2, the time constant for thermal equilibrium between nuclei themselves. Information on molecular dynamics can be obtained from these relaxation times. Generally, Ti = T2 for lowviscosity liquids and Ti >> T2 for solids. A combination of information in molecular dynamics (from relaxation times), molecular structure (from spinspin interaction), molecular identification (from resonance frequency and chemical shift), and spin density (from signal intensity) make the NMR an extremely versatile tool. The application of NMR spectroscopy to ceramic powders is a relatively new area, but it has potential. N M R application in solids is difficult because of large dipole-dipole interactions that result in line broadening and thus poor resolution. However, by the application of modem instrument technology and discoveiy of pulsed Fourier transform N M R , scientists have developed various line-narrowing techniques to improve resolution [15]. One of the most often utilized techniques is magic angle spinning (MAS). Molecules in solids are not mobile and tumbling as they are in the liquid or gaseous state. This induces large dipolar interactions between two spin centers. If we physically spin the sample, at several kilohertz, the anisotropic part of the nuclear dipole is eliminated. The dipole-dipole interaction is a function of (3 cos2 8 - l), where 8 is the angle between the static magnetic field and internuclear vector. Setting this term to zero, the “magic” angle is 54”44’ [16]. N M R has been applied most successfully for high-temperature superconductors, YBa2Cu307-6 [17-211. The studies involve mainly 89Y,63Cu, and 65Cu NMR and nuclear quadrupole resonance of this compound below, above, and at the critical temperature. Nuclear relaxation, Knight shift, and crystal structure are often examined at these temperatures. 29Si and *7Al N M R have been applied to silicates, aluminosilicates, and surface-adsorbed molecules on ceramic materials [22]. N M R can also be a powerful technique to determine the surface area [23], porosity, pore size, and molecular diffusion into the crystallattice [24]. The technique is basedon the different relaxation times (or linewidths) of the mobile molecules and adsorbed molecules, the latter losing certain degrees of freedom. The crystal structure of B a T i e has been characterized by 137Ba, 47Ti, and 49Ti NMR at various temperatures up to the Curie point [25]. Structural changes in the alkaline-earth silicate glasses induced by phosphorus were also studied by 31P and 29Si NMR [26]. 29Si MAS NMR is a very useful technique to study the crystal-phase composition and transition in silicon nitride and carbide [27,28]. The a phase of

MALGHAN ET AL..

140

these compounds has more than one Si crystal site; while p phase has only one. This yields different line patterns for different phases. MAS NMR has been used to study high-temperature reactions, sintering, the formation of surface oxide and suboxide, and reaction products with the sintering aid [29-341. The theory of EPR (also known as electron spin resonance) is very similar to that of the NMR except the frequency falls into gigahertz region because the mass of electrons is smaller. Electrons are negatively charged, so that the splitting of energy is inverted in its order compared with the nuclei [35]. Since only molecules with unpaired electrons are EPR active, the application of EPR to ceramics is limited.However,EPR is a unique technique for characterizing powders with dangling bonds (as in carbon atoms), impurities, and defects.

E. Combustion This technique is limited to such elements as carbon, nitrogen, and oxygen that can be liberated in gaseous form during decomposition. The concentration of gas is determined by a property of the gas. Commonly a flux is used to aid the combustion [3,36].

IV.

PHASE COMPOSITION

A.

X-rayPowderDiffraction(XRPD)

This is the most widely used method for the determination of the phase composition of powders. The x-ray diffractometer contains a source of monochromatic x-rays that irradiate the sample and are diffracted from atomic planes and detected. The angle of diffraction of x-rays by the crystalline planes is characteristic of the crystal structure, and the intensity of scattered radiation is characteristic of the atomic composition. In recent years, automated data processing has enabled higher accuracy andspeed. A number of problems are encountered in the quantitative determination of phases in fine powders. Some of these are overlap of phase peaks (e.g., in silicon nitride), orientation of grains, and presence of coarse particles. The last produces distortion of the diffraction data. A number of standard reference materials for XRPD have been developed for use in improving the quality of data [37].

V. A.

SURFACECHEMICAL COMPOSITION BY SPECTROSCOPY X-ray Photoelectron Spectroscopy and X-ray-Induced Auger Electron Spectroscopy (XAES)

X P S (also known as electron spectroscopy for chemical analysis) and XAES provide not only the surface elemental composition but also reveal the oxida-

CHARACTERIZATION OF ADVANCED CERAMIC POWDERS

141

tion states of the elements. X P S , especially, has emerged as one of the most important techniques for studying the chemistry of the ceramic-powdersurface as a result of oxidation at high temperatures. In X P S , when soft x-rays, such as MgKa (1253.6 eV), impinge on a powder surface, photoelectrons with binding energy BE are ejected from the surface with kinetic energy (KE). The relationship among the irradiation energy hv (1253.6 eV in the case of MgKa), KE, and BE is as follows: KE=hv-BE-eS

(7)

where is the work function of the spectrometer. Since is a constant for a specific spectrometer, hv is known and KE can be measured; therefore, BE can be calculated. Often, residual carbon in the form of hydrocarbon (-CH) is found on the powder surface as a contaminant, and the C 1s photoelectron peak at 284.5 eV binding energy is used as a reference to calculate the binding energies of other signals. From the binding energy of the emitted electrons, the surface elemental composition and chemical state can be derived. Detailed information on surface analysis techniques, comparison of XPS and Auger electron spectroscopy (AES), and photoelectron binding energies can be found in the respective references [3840]. For quantitative surface composition, the photoelectron cross sections (or atomic sensitivity factors) are required to convert the signal intensities to atomic ratios. In other words, one photoelectron from C 1s gives a different spectral intensity from that of a Si Is, for example. Scofield cross-sectional values are often used for this purpose [41]. When two spectral signals from different oxidation states of an element (for example, Si 2p from Si3N4 and its surface oxide Si02) can be resolved because of a large difference in binding energy, the surface film thickness can be calculated if the photoelectron mean free paths data are available [4247]. When photoelectrons are emitted from the core level as a result of a photoelectronic process, the surface becomes ionized and unstable. The ion relaxes its energy by drawing an outer electron to fill the inner orbital vacancy left by the photoelectron, and a second electron is emitted by the excess energy. This second emitted electron is an Auger electron (more precisely, x-ray-induced Auger electron, X A E S , or bremsstrahlung-excited Auger electron). A comparison of the photoelectron and Auger processes is shown in Fig. 1. It can be concluded that the kinetic energy of a photoelectron is directly proportional to the energy of irradiation, but the kinetic energy of an Auger electron is independent of the radiation source and possesses kinetic energy equal to the difference between the energy of the initial ion and the doubly charged final ion. The energy of neither electron can exceed the energy of the ionizing photons. Sometimes the difference in kinetic energies of two Auger electrons from two different oxidation states is larger than that of the photoelectrons, and thus the

142

-

MALGHAN ET AL.

L1 OR 2s

PHOTON 0 0

0

f

PHOTOELECTRON

0/

K OR 1s

/

f

p AUGER ELECTRON

12.3

OR

“-ece---- L, OR b I

i

I

I K OR 1s

Figure 1 Comparison of photoelectron and Augerprocesses. Augerspectrum provides better resolution for surface chemical studies, as shown for S i c and Si3N4 powders [42,43,48]. S i c and Si3N4 powders have been studied extensively by X P S or x-ray-induced AES [42,43,48-571. Powders, whiskers, and platelets were treated at elevated temperatures, and surface analysis was carried out by measuring Si 2p X P S or Si KLL AES intensities for Si02 and S i c (or Si3N4). Subsequently, surface oxide film growth rates were measured and surface oxidation activation energies were calculated [42,43,48-511. In some cases, the effects s f the presence of ymia, a sintering aid, and boron, an impurity, were studied [43,48,49]. The effects on surface oxidation chemistry caused by different powder processing routes were also studied [50]. Surface composition, including oxygen and oxynitide concentration, is important information required for nonoxide powder processing. In a recent study, several commercial silicon nitide powders were analyzed for total oxygen at high temperatures, and the “surface” oxygenwasmeasured by AES. The “bulk” oxygen was then calculated based on these data [52]. Commercial S i c

king

CHARACTERIZATION OF ADVANCED CERAMIC POWDERS

143

powders were examined for surface composition and contamination by X P S [53,54]. Ultrafine powders grown bya radiofrequency plasma process were found to have a thin oxide layer; those grown in a vapor-liquid-solid process had a thicker silica layer [53]. Surface-sensitive techniques, such as XPS and AES, also contribute to the understanding of moisture effects, surface coating, and interfacial chemistry of ceramic powders [S-591.

B.

Infrared and RamanSpectroscopy

For a polyatomic molecule containing n atoms, the total degree of freedom is 3n and there are 3n - 6 modes of vibration if the molecule is nonlinear. This shows that molecular vibration in a polyatomic molecule can be very complex. Infrared (R) and Raman are two complementary techniques to study molecular vibration and identify unknown species. There are two fundamental types of molecular vibrations: stretching and bending. Stretching, in which twoatoms increase or decrease their distance but remain in the same bond axis, can be either symmetrical or asymmetric. Bending is a molecular deformation and can be scissoring, rocking, wagging, or twisting. Figure 2 shows these vibrations for a group of atoms. Molecular bending generally requires less energy than stretching and is thus observed at lower frequency. The frequency of stretching depends on the bond strength and mass of the atom attached and can be calculated from

where 2) is stretching frequency in cm-1, c is the velocity of light, Mx and M y are the mass in grams of the two atoms attached to the bond, and k is the force constant in dyneshentimeter of the bond.

Asymmetric

Symmetric

Scissoring

'Figure 2 Vibrations for a group of atoms.

Twisting

144

MALGHAN ET AL..

Each of these vibrations has a characteristic frequency and can occur at quantized frequencies only. When IR light of the same frequency is incident on the molecule, the energy is absorbed by the molecule and the amplitude of the particular mode increases. However, this absorption occurs only if this vibrational mode can cause a change in the molecular dipole. Consequently, not all vibrational modes are IR active and the molecular symmetry plays a key role in the reduction of IR spectrum patterns. In addition to these fundamental vibrations, overtone peaks may also be observed with much reduced intensity at two, three times, and so on, the wave numbers, the sum of twoor three times the wave numbers, or the difference between two wave numbers. Detailed IR spectroscopic theory and group theory can be found elsewhere [60-62]. Surface moisture is a problem of concern in ceramic powders, and IR has been used to characterize the surface groups of -OHand -H [58,63,64]. IR was also applied to characterize chemically bound hydrogen in chemical vapordeposited silicon nitride at various ammonia-silane ratios [65]. Surface silicon dioxide on S i c powders was determined by photoacoustic IR and diffuse reflectance IR spectroscopy [66,67]. IR spectroscopy was also used to study the surface oxidation of S i c and Si3N4 [68,69]. Raman spectroscopy is complementary to R.The vibration that is inactive to IR, because ofa high degree of molecular symmetry or lack of dipole change, may be detected by Raman. Raman depends on a change in the polarizability during the vibration instead of the electric dipole moment. When a monochromatic light with frequency v impinges on a molecule, the scattered light has the same frequency if the scattering is elastic. This is Rayleigh scattering. However, some of the light is scattered with v f v’ because of inelastic interactions. Lines with lower frequency are known as Stokes lines And are formed by the loss of energy to the molecules during scattering. Lines with higher frequency are known as anti-stokes lines and are formed by energy gain from the molecules during scattering. The Raman frequency v’, is completely independent of the incident light v. Rather, it is characteristic of the molecular vibration energy. By measuring the characteristic Raman frequency, one can identify the unknown molecules [60]. Application of Raman scattering to ceramic powder characterization is a new research area, and the data are limited.

VI.INTERFACEANALYSISIN

SUSPENSIONS

A suspension of finely dispersed particles is thermodynamically unstable; this system is inclined to lower its free energy through flocculation. The stability and rheology of a powder in a suspension depend on the nature of the solid/ solution interface, particularly on the electrical properties of this region [70]. This interface may be described as consisting essentially of two layers (Fig. 3):

CHARACTERIZATION O F ADVANCED CERAMIC POWDERS

145

Shear Plane

0

0 0 Solution

Figure 3 The solid/aqueous solution interface and electrical double layer for a hydrolyzed oxide surface: G, charge density; v, electrostatic potential; +, cation; -, anion, 0 , surface; c, compact layer; d, diffuse layer.

a compact layer near the surface consisting of potential determining adsorbed species and complexed ions, balanced by a diffuse layer of counterions in solution. Suspension stability is a consequence of mutual repulsion between similarly charged double layers. It is therefore desirable to measure the electrical potential as a function of solution or powder conditions to optimize slurry properties during wet processing. A complete surface chemical characterization should also include information about the type and density of surface sites and the interaction of solution species with the surface. In this section, we cover only in situ techniques for the analysis of aqueous powder dispersions, because these systems are most relevant to ceramic powder processing. The focus is on electrokinetic methods for the measurement of particle electrostatic potential. In addition, a brief overview of available surface chemical characterization techniques is given. Additional details on this topic are given by other authors in this book.

MALGHAN ET AL.

I46

A.

ElectrokineticCharacterization

At some distance from the particle surface (usually identified as the beginning of the diffuse layer in Fig. 3), a hydrodynamic shear plane exists that is characterized by the potential. The magnitude of ( is directly related to dispersion stability [71]. For oxides, hydroxides, and related materials, is strongly influenced by solution pH and electrolyte concentration and may be modified by surface-active species, such as oxyanions and polyelectrolytes. The key parameter characterizing a powder surface is the isoelectric point pHg. Under pristine conditions (i.e., no surface contamination), PHiep defines the solution pH at which = 0 and the particles exhibit a net surface charge of zero. This point can be identified from an acid-base titration curve in which is plotted against pH. Shifts in PHiep may result from changes in surface chemistry that occur during aging or chemical treatment of powders [ S ] , for instance, or caused by the chemical adsorption of solution species [72]. The potential may be obtained from measurements of particle mobility using electrokinetic techniques, such as electrophoresis or sedimentation potential [70]. Electrophoresis, the standard technique for submicrometer particles, is based on the movement of charged particles in response to an applied electrical field. Optical scattering methods are used to measure the distribution of particle velocities for a given field strength, and may then be calculated using the Henry equation,

<

<

<

<

<

c

where q and E are the viscosity and dielectric constant of the medium, P E is the electrophoretic mobility, and the functionfdepends on the thickness of the double layer ~ - 1relative to the particle radius a. The Henry equation is applicable to many powder slurry systems in which the double layer is thin relative to the particle radius (KU >> l), although more exact solutions exist for the calculation of over a broad range of KU values [70]. Because electrophoresis uses optical detection, this technique is limited to the analysis of dilute systems; however, the recent development of electroacoustic methods has extended analysis to concentrated slumes containing up to 50% vol/vol solids [73]. The electroacoustic effect is the response of charged particles to an applied alternating electrical or acoustical field [74], in contrast to the static field employed in electrophoresis. The acoustical response results from relative vibratory motion between particle and medium if the two phases differ in density. If an alternating electrical field is applied, charged particles vibrate in a back-and-forth motion in phase with the applied field, producing a sound wave whose pressure amplitude is proportional to the particle mobility and <.This technique is termed electrokinetic sonic amplitude (ESA). Alternatively, if an ultrasonic wave is applied, the particles vibrate at the sound

<

<

.

CHARACTERIZATION OF ADVANCED CERAMIC POWDERS

I47

frequency, producing a measurable electrical field that is proportional to the particle mobility. This is termed the ulfrasonic vibration potential (WP). ESA is directly proportional to the dynamic mobility, whereas UVP requires knowledge of the high-frequency conductivity of the slurry. For this reason, ESA is frequently used for routine analysis of powder slumes under processing conditions [75]. An example of the application of ESA to powder characterization is given in Fig. 4, where the effect of surface cleaning on a Si3N4 powder is evidenced by a shift in pHiep of the aqueous slurry. This shift is caused by a reduction of the surface oxide layer thickness.

B. Overview of Interface Analysis Techniques The most important group of surface properties that characterize an aqueous powder suspension are its acid-base properties. The reaction of surface species with H+ and OH- in solution controls the surface charge and potential as a function of pH. The surface charge and point of zero net charge may be determined from potentiometric [76] and conductometric [77] titrations. In a potentiometric titration, pH is measured as a function of the amount of base or acidadded. Titration curves for the suspension and solution blank are subtracted to obtain a proton uptake curve at several ionic strengths. The common intersection point of this set of curves denotes the point of zero charge, which, once known, maybeused to calculate charge densities at any pH. Solution conductivity may also be used as an indicator in acid-base titrations. Here changes in the slope correspond to reaction end points. Isotopic exchange of radioactively labeled compounds is useful for characterizing the labile nature of complexes formed between surface and solution species and the availability of surface sites [78]. Surface complexes and bonding mechanisms may be identified using in situ spectroscopic methods, such as cylindrical internal reflection Fourier transform infrared [79]. This technique is sensitive to IR-active vibrations of the interfacial region and is less inhibited by the strong OH absorption interference normally encountered when IR is performed in aqueous environments. The interference problem is avoided by using blank subtraction methods and an internal reflection element that reduces the signal from bulk water by sampling only a microscopic distance into the adjacent suspension. Another technique is extended x-ray absorption fine structure [go]. This method allows measurement of the local structural environment around species complexed at the solid/solution interface and yields interatomic distances and the coordination number of the nearest neighbors. Calorimetry [81] measures the enthalpic contributions to interfacial chemical reactions and, in combination with potentiometric titration, can be an effective tool for characterizing powder surface chemistry. The heat of adsorption provides information regarding the affinity and interaction mechanism between solution species and the powder surface.

148

MALGHAN ET AL. 0.9

cleaned surface 0.6 n

3 \

0.3

E &

pc

E

0.0

W

2 W

-0.3

as-received

-0.6

-0.9

4

5

6

7

0

Q

10

PH Figure 4 Electrokinetic sonic amplitude @SA) analysis of Si3N4 slumes before and after powder surface cleaningby Soxhlet extraction. The isoelectric pointsare indicated in parentheses.

VII.

SUMMARY

This review of physical, bulk chemical, surface chemical, and spectroscopic techniques for the characterization of powders shows that no one or two techniques can provide all the necessary details regarding powder properties. In fact, each method reveals a distinctly different characteristic of the powder. Often, a choice must be made in terms of the selection and measurement of relevant properties of a powder. Whenone is faced with this question of which specific properties to measure, the relationships between powder properties and their effect on the final microstructure and properties of the ceramic should be explored. Currently, the quality and reproducibility of data are significantly affected by the unavailability of standard methods and standard reference materials. Efforts are underway in different organizations around the world to alleviate this problem.

CHARACTERIZATION OF ADVANCED CERAMIC POWDERS

149

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37. Cline, J. P., XRD standard reference materials-their characterization and uses, in Accuracy in Powder Diffraction 11, NIST Special Publication, 1992. 38. Czandema, A. W. (ed.), Methods of Surface Analysis, Elsevier, New York, 1975. 39. Carlson, T. A., Photoelectron and Auger Electron Spectroscopy, Plenum Press, New York, 1975. 40. Muilenberg, G. E. (ed.), Handbook of X-ray Photoelectron Spectroscopy, PerkinElmer Corporation, Eden Prairie, Minnesota, 1979. 41. Scofield, J. H., Hartree-Slater subshell photoionization cross sections at 1254 and 1487 eV, J . Elec. Spec. Relat. Phenom., 8, 129 (1976). 42. Wang, P. S., Hsu, S. M., and Wittberg, T. N. Oxidation kinetics of silicon car-

bide whiskers studied

by x-rayphotoelectronspectroscopy, J . Mater. Sci., 26,

1655 (1991). 43. Wang, P. S., Hsu, S. M., Malghan, S. G., and Wittberg, T. N., Surface oxidation

44. 45. 46. 47.

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kinetics of Si3N4-4%Y203powdersstudied by bremsstrahlung-excitedAuger spectroscopy, J. Mater. Sci., 26, 3249 (1991). Powell, C. J., Energy and material dependence of the inelastic mean free path of low-energy electrons in solids, J. Vac. Sci. Technol., A3(3), 1338 (1985). Powell, C. J., and Seah, M. P., Precision, accuracy, and uncertainty in quantitativesurfaceanalyses by Auger-electronspectroscopyandx-rayphotoelectron spectroscopy, J. Vac. Sci. Technol., A8(2), 735 (1990). Tanuma, S., Powell, C. J., and Penn, D. R., Calculation of electron inelastic mean free paths. II. Data for 27 elements over the 50 - 2000 eV range, Surf. Interfac. Anal., 17, 911 (1991). Tanuma, S., Powell, C. J., and Penn, D. R., Calculation of electron inelastic mean free paths. m. Data for 15 inorganic compounds over the50-2000 eV range, Surf. Interfac. Anal.,17, 927 (1991). Wang, P. S., Malghan, S. G., Hsu, S. M., and Wittberg, T. N., Surface oxidation of silicon carbide platelets as studied by x-ray photoelectron spectroscopy and bremsstrahlung-excited Auger electron spectroscopy, Surf. Inte~ac.Anal., 18, 159 (1992).

Wang, P. S., Malghan, S. G., Hsu, S. M., and Wittberg, T. N., Oxidation of surface-treated S i c platelets studied by X P S and bremsstrahlung-excited AES, Surf. Interfac. Anal., 20, 105 (1993). Wang, P. S., Malghan, S. G.,Hsu, S. M., andWittberg,T. N., Effects of asilicon nitride powder processing on surface oxidation kinetics, submitted to J. Mater. Res., (1993). Pugh, R. J., Surface chemical analysis of oxidized ultrafhe a-Sic powdersby electron spectroscopy, J . Coll. Interfac. Sci., 138, 16 (1990). Jenett, H., Bubert, H., and Grallath, E., Comparative surface and bulk analysis of oxygen in Si3N4 powders, Fresenius Z. Anal. Chem., 333, 502 (1989). Taylor, T. N., The surface composition of silicon carbide powders and whiskers: An XPS study, J . Mater. Res., 4(1), 189 (1989). Rahaman, M.N., Boiteux, Y., and De Jonghe, L. C., Surface characterization of silicon nitride and silicon carbide powders, Am. Ceram. Soc. Bull., 15(8), 1171 (1986).

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111 POWDER PROCESSING

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7 Colloid Interface Science for Ceramic Powder Processing Hyun M. Jang Pohang University of Science and Technology Pohang, Republic of Korea

1.

INTRODUCTION

Interfacial phenomena at metal oxide/water interfaces are fundamental to various phenomena in ceramic suspensions, such as dispersion, coagulation, coating, and viscous flow. The behavior of suspensions depends in large part on the electrical forces acting between particles, which in turn are affected directly by surface electrochemical reactions. Therefore, this chapter first reviews fundamental concepts and knowledge pertaining to electrochemical processes at metal oxide powder (ceramic powder)/aqueous solution interfaces. Colloidal stability and powder dispersion and packing are then discussed in terms of surface electrochemical properties and the particle-particle interaction in a ceramic suspension. Finally, several recent examples of colloid interfacial methods applied to the fabrication of advanced ceramic composites are introduced.

II.

INTERFACIALELECTROCHEMICALCHARACTERISTICS

A.

Origins of Surface Charge at the Ceramic Powder/ Water Interface

Most substances in contact with an aqueous medium acquire a surface electrical charge caused by a redistribution of charge in the interfacial region. This process is described as the formation of an electrified interface, namely, an electrical double layer. 157

JANG

158

Several types of redistribution of charges are possible [1,2]: 1. The ionization of ionogenic groups; proteins acquire their charge mainly by ionization of -COOH and -NH2 groups to give -COO- and -NH3+. 2. Unequal dissolution of oppositely charged ions of which the solid phase may be composed (e.g., silver halide particles suspended in water). 3. Unequal adsorption of ions of opposite charge by the solid phase. 4. Adsorption and orientation of dipolar molecules. With regard to ceramic powder (inorganic metal oxides), most experimental studies have revealed the special role of hydrogen and hydroxide ions in determining both the surface charge and potential of all inorganic oxides. It was observed that small changes in concentration of acid or base are able drastically to change the sign and magnitude of the surface charge or potential [3]. On the other hand, the addition of other simple inorganic electrolytes has a smaller effect and usually does not change the pH of the suspension at which the charge and the potential of the oxide are zero [3]. Therefore, the electrical double layer at the ceramic powder/aqueous solution interface is considered established by the unequal adsorption of H+ and OH-. The adsorption mechanism of H+ and OH- has long been attributed to the amphoteric reaction of the surface hydroxide groups -M-OH [4]. This is schematically represented as

-M-OH+H+ -M-OH

2

K F ! !

-M-OH:

-M-O- + H +

(1)

where -M-OH2+ and 44-0-represent positive and negative surface sites. An alternative to the H+ and OH- adsorption mechanism, suggested by Parks and de Bruyn [S], involves the adsorption of metal-hydroxo complexes derived from the hydrolysis products of materials dissolved from the solid oxide. In this case the charge is also caused by amphoteric dissociation of hydroxide groups, but the hydroxide groups are considered a part of the adsorbed hydroxo complexes rather than the real surface. The distinction between these two mechanisms is a fine one, the end result being the same or very similar, depending on how the hydroxo complexes are adsorbed. The main difference is that, in the adsorbed hydroxo complex mechanism, the surface must first dissolve sufficiently to provide enough hydrolysis products and establish the surface charge upon their adsorption. The experimental fact that charged oxide/water interfaces are readily identified through such techniques as streaming potential, in which the concentration of hydrolysis products is very unlikely to be significant, indicates that these species are not indispensable to establish the electrified interface. Therefore, the surface

159

SCIENCE COLLOID INTERFACE

dissociation reactions given in Eq. (1) must represent the most important charging mechanisms, but some contribution to the surface charge from the adsorption of soluble hydroxo species cannot be completely ruled out. The surface charge density 00 on oxides is then given by the difference in adsorption density of hydrogen and hydroxide ions r w and rOH- respectively, and it is defined as

where F is the Faraday constant, 00 is commonly expressed in pC/cmz, andrH+ and ToH- are in mol/cm*. Since the adsorption of hydrogen and hydroxide ions determines both the surface electrical charge and the surface coulombic potential yo, they are called the potential-determining ions (PDI) for metal oxides [4,5]. When these surface parameters (i.e., the surface charge density and the surface coulombic potential) are zero, the solid is said to be at its point of zero charge (PZC). Since the PZC is conveniently represented by - log (activity of the potential-determining ions), it designates the equilibrium suspension pH when the surface charge density 00 at the ceramic particles/aqueous solution interface becomes zero. Thus, this value gives the solution conditions under which the surface is uncharged and the suspension is destabilized. Therefore, a suspension is stable only if the pH is kept sufficiently far from the PZC. Approximate values of the PZC of various ceramic materials are listed in Table 1. Table 1 Point of Zero Charge (PZC)

of Various Aqueous Ceramic Suspensions

pH

Material Silica, quartz Tin oxide (cassiterite) Titania (rutile, anatase) Zirconia Chromium oxide Goethite (FeOOH) Boehmite (AIOOH) Hematite (synthetic) Alumina (aand y) Calcite (CaC03) Magnesia Mullite (3A1203 . 2Si02) Cordierite

of PZC

2-3 4.5 6 -6.5 6-7 6.8-7.6 7.8-8.8 8-9 8.5-9 9.5 12 7.5 2-3

160

B.

JANG BasicElectricalDouble-LayerTheory

In the preceding section the origin of the surface charge at the ceramic powder/water interfacewasdiscussed. To maintainelectroneutrality,the surface charge is then balanced by an equal and opposite charge in the mobile aqueous phase (counterions). This compensating solution charge is created by the adsorption (or condensation) of electrolyte ions at the solution side of the interface. The distribution of the interfacial charges and the consequent variation in the electrostatic potential is the concern of the theory of the electrical double layer (EDL). The well-known Gouy-Chapman-Stem-Grahame(GCSG) model, which has been successfully applied to mercury/solution interfaces by Grahame [6]and to AgI/solution interfaces by other investigators [2], is briefly outlined here. The characteristics of the EDL and the problems associated with a directapplication of theGCSGmodel to themetal oxide (ceramic powder)/aqueous solution interface are then discussed in the next two sections. In the GCSG model, inaddition to the surface charge itself, the solution side of the electrical double layer (or simply the diffuse layer) is regarded as consisting of two regions: an outer diffuse region, in which the ions undergo thermal motion and are distributed according to the equilibrium Boltzmann distribution, and an inner region, where any adsorbed ions are relatively immobile. The Gouy-Chapman theory of the diffuse layer at a charged interfacial region is based on the well-known Poisson-Boltzmann equation:

where v is the mean coulombic potential at any point in the diffuse layer, ni 0 is the number of ions of type i per unit volume in the bulk solution, zi is their valency, including sign, e is the electron charge, &b is the dielectricconstant of themedium, EO isthepermittivity of vacuum (8.854 x 10-I2F/m), k is the Boltzmann constant, and T is the absolute temperature. For a single symmetrical z/z electrolyte and a flat interface, Eq. (3) is readily solved with the boundary conditions ~ ( x =) y ~ dat x = 0 and = @/& = 0 at x = (bulk solution), to give

-

where

COLLOID INTERFACE SCIENCE and

(E;;)

I61

112

K=

= 3.29 x 109z@

m-’ at 298 K

where C is the bulk electrolyte concentration in mol/dm3. The quantity IC’ is the center of gravity of the diffuse layer charge and is generally referred to as the “thickness” of the diffuse layer or the Debye-Huckel length. The diffuse layer charge o d is the charge contained in a cylinder of unit cross-sectional area extending from x = 0 to x = and is given by

For curved interfaces, Q. (3) cannot be solved analytically. For spherical interfaces, the Debye-Huckel approximation can be applied, provided that zew c 25 mV. Alternatively, solutions may be obtained by numerical methods with the aid of a computer, as have been fully compiled for spherical double layers by Loeb et al. [7]. A schematic representation of the inner region of the double layer model is shown in Fig. 1. Figure l b describes the distribution of counterions and the potential profile w ( ~from ) a positively charged surface. The potential decay is caused by the presence of counterions in the solution side (mobile phase) of the double layer. The inner Helmholtz plane (IHP) or the inner Stem plane (ISP) is the plane through the centers of ions that are chemically adsorbed (if any) on the solid surface. The outer Helmholtz plane (OW)or the outer Stem plane (OSP) is the plane of closest approach of hydrated ions (which do not adsorb chemically) in the diffuse layer. Therefore, the plane that corresponds to x = 0 in Eq. (4) coincides with the OHP in the GCSG model. The doublelayer charge and potential are defined in such a way that 00 and WO, op and wp, and o d and w d are the charge densities and mean potentials of the surface plane, the Stem layer (IHP), and the diffuse layer, respectively (Fig. 1). On the analogy of the Langmuir adsorption isotherm, one can derive the following equation (the Stem equation) for op caused by the adsorption of ions at the IHP (or ISP):

where Ns is the number of available adsorption sites per unit area, No is Avogadro’s number, M is the molecular weight of the solvent, z* is the valence of cations or anions, including sign, and nf is the bulk concentration of the ionic species in ions cm-3. The terms I$+and Q- are the specific adsorption potentials

I62

JANG

(ISP) (OSP)

f

Potential

0

P P+Y Distance "I>

shear plane 0

Potential

IHP

zeta potential reversal 0

0

P P+Y Distance

+

Figure 1 Theelectricaldoublelayer:(a)theinnerregion of theelectricaldouble layer, showing the location of the M P , O W ,potentials y, and charge densities o;(b) distribution of counterions and the potential profile y ( x ) from the positively charged solid surface; (c) distribution and potential drop across the double layerin the presence of specifically adsorbed ions.

SCIENCE COLLOID INTERFACE

I63

that take into account any noncoulombic contribution to the adsorption energy, changes in the state of solvation of the adsorbed ions, and covalent, van der Waals, or hydrogenbondingbetween the adsorbate and the surface.When these noncoulombic interactions significantly contribute to the total free energy of adsorption (zeyp + $), adsorption is said to be specific [ 1,6], and the extent of adsorption then depends on the nature as well as the charge of ions. An important consequence is that a specifically adsorbed ion can adsorb against electrostatic repulsion (i.e., I$l>lzeypl), and the surface charge may be exceeded by the specifically adsorbed charge (lop1 >loot, as shown in Fig. IC), therefore reversing the diffuse layer charge. The charge densities at the surface plane, the IHP (or ISP) and the OHP (or OSP), are related by the necessity for overall electrical neutrality:

The potential drop between each plane is assumed linear, and the integral capacity per unit area Ki of the zone between the surface and the IHP is then given by

where &I and P are the dielectric constant and the distance shown in Fig. 1. The integral capacity KO per unit area between the IHP and the OHP is given bY

where ~2 and y are the dielectricconstant and the distance between theIHF' and the OHP, respectively. Therefore, &S. (5) through (9) relate the six variables op, 00,od,yp, yo, and y d of the GCSG model of the electrical double layer. Hence, if one of the variables is known, then the others can be calculated, provided that values for the parameters N,, $*, Ki,and KO are also known. In the simple GCSG model already described, the charge densities are assumed uniformly smeared over each plane rather than in the form of discrete ions. Levine and co-workers [8] modified the GCSG model by introducing the discreteness of charge effect to account for various phenomena, mostly relating to adsorption at the mercury/water and silver iodide/water interfaces. Their theory shows that the electrostatic work of adsorption at the IHP is not zeyp but

I64

JANG

where \VA is the micropotential in the hole created by the rearrangement of the other adsorbed ions and $‘p is the self-atmosphere potential. The calculation of $‘p is rather complicated and depends on the physical properties of the inner region. An interesting consequence of the discreteness of charge theory is that it predicts that a maximum may be observed in the OHP potential lyd. This is a possible explanation for the maxima sometimes observed in electrokinetic potentials as the surface potential is increased. The work of Wiese et al. [g] on silica is a representative example of the applicability of this theory. Among the six interfacial variables discussed inthis section, the surface charge density 00, the surface potential yo, and the potential at the OHP yd (usually called the diffuse layer potential), are most important in characterizing interfacial properties. The three remaining variables (i.e., op, yp, and od) can be estimated using Eqs. (5), (7), and (8) if 00, yo, and yd are known exactly. 00 can be determined experimentally by the potentiometric titration method, and detailed explanation of the potentiometric titration is given, for example, by Yates [lo]. The estimate of yo for the ceramic powder/aqueous solution interface is discussed in the next section. y d is perhaps the most important interfacial electrochemical parameter since it is closely correlated with the kinetic stability of a given colloidal suspension and it can be conveniently determined (approximately) experimentally. One method of estimating y d involves applying an electrical field across a dispersion of particles and measuring the resulting particle velocity. This technique, known as electrophoresis, allows measurement of the so-called potential or slipping plane potential. The slipping plane (or shear plane) arises as solid particles move through the fluid, causing the counterions in the diffuse layer to slip away and leaving some of the interfacial charges uncompensated (electrokinetic effect). Although the location of the shear plane is not known exactly, it is likely to be located at a small distance farther from the surface than the O W , so that Iyd. Various careful studies showed that the identity of with yd is reasonable for low potentials (for yd less than -1OOmV at 1C2 M electrolyte concentration). At higher potentials, particularly at high electrolyte concentration, the shear plane probably shifts farther from the surface, and 161 is expected to be less than yd. In systems in which polymer adsorption is known to occur, the slipping plane is located rather far from the surface and the electrokinetic effects are reduced. The solution condition at which the 6 potential (or yd) reverses sign has been termed the isoelectric point (IEP) or, simply, the point of 6 potential reversal (PZR). Although both the PZC and the PZR usually give rise to a zero 6 potential, the PZR is a result of surface charge compensation by a strong adsorption at the IHP (00 + op = od = 0 in Fig. l), but at the PZC there is no surface charge (00= 0 in Fig. 1). More recently, the electrophoretic laser light-scattering method was used to

<

<

<

SCIENCE COLLOID INTERFACE

I65

c

estimate diffusion coefficients and potentials of small colloid particles by measuring the Doppler shift of the frequency of the scattered light caused by the velocity of the scattering centers [ 1 l]. In addition, the electrokinetic sonic amplitude (ESA) method is actively applied to the measurement of electrophoretic mobility (thus, potential), especially for a concentrated suspension. In this method, an alternating electrical field is applied to the suspension. This createsrelativemovementbetween the colloidparticlesandthe surrounding medium, generating a sound wave at the same frequency as the electrical field. The acoustic energy generated is proportional to the electrophoretic mobility. A correlation between these ESA measurements and the potential values was established through this relationship [12].

c

c

C.

Surface Potential of a Charged Interface

Because they are electrical insulators, the net potential difference across a ceother ramic powder/water interface yo cannot be directly measured. On the hand, the other two important interface variables (GO and yd) can be measured or estimated, as discussed in the previous section. Therefore, one should be able to calculate or reasonably estimate yo from other experimentally observable quantities for a complete characterization of interfacial properties. In this section we first introduce the so-called Nemst approximation for the surface potential yo. The limitation of the Nemst equation is then discussed, and the modified Nemst equation for yo of a ceramic powder/aqueous solution interface is subsequently introduced. The general derivation of an equation for the potential difference across an interfacehasbeenthoroughlytreated by Parsons [l] andalso by Overbeek [13]. The starting point is the condition for thermodynamic equilibrium between the two bulk phases a and p, given by

where jila is the total work required to move charged species i from a point in a vacuum and infinitely distant from the phase a to a point in the bulk phase a.This quantity ji is called the electrochemical potential and is given by

F:

=pp*+ RT In ay + Z , . F $ ~

where jiP is the standard chemical potential, aia is the chemical activity of species i in phase a,Zi is the charge of that species, F is the Faraday constant, and I$P is the Galvani or inner potential of the phase a. Btrub6 and de Bruyn [l41 pointed out that electrochemical equilibrium between both bulkphases and the surface phase must be considered when the surface potential y o is introduced. Hence, the equilibrium condition to be satisfied is

I66

JANG

where the superscript S refers to the surface phase. Using Eqs. (12) and (13), the inner potential difference between the surface phase and the bulk phase p (mobile liquid phase) can be written as

where

On defining the surface potential yo as the difference in the inner potential between the surface phase and the bulk phase p, Eq. (14) becomes

where aPBand a$ are the activities of species i in the bulk phase p and the surface phase S, respectively, when WO = 0. A s discussed in Sect. ILA, it has been experimentally observed that H+ and OH- are the potential-determining ions for metal oxides, even when these ions are not present in the solid phase. By analogy with Eq. (15), the general form of the Nemst equation describing the surface potential of metal oxides can be written as

where the superscript for the bulk aqueous phase has been omitted and m is the activity of H+ in the bulk aqueous phase. Only if the surface activity of hydrogen ions is constant and independent of y o does Eq. (16) become the often assumed Nernst equation for metal oxides:

where pH0 = - log a$, the pH at which yo = 0 (i.e., PZC). A number of authors [1+16], however, have questioned the validity of this equation for metal oxides. It is now generally recognized that the Nemst approximation for metal oxides is doomed to failure. The problem is in the assumption that the activity of the potential-determining ions on the surface is independent of the surface potential (zeroth order approximation). For example, as the bulk pH decreases

COLLOID INTERFACE SCIENCE

167

and deviates from the PZC, the surface charge density, and thus the surface potential, increases [Eq. (l)]. The probability of finding positively charged sites adjacent to a given positively charged site (-M-OH2+) then increases. This increases repulsive interaction energy on the positively charged sites. The increased repulsive interaction energy enhances the chemical potential and, thus, the activity of positive sites significantly. Therefore, the surface activity of hydrogen ions (a*) or positive sites is no longer constant, and one cannot derive the Nemst equation. An alternative derivation of Eq. (16), based on the amphoteric reactions of the surface hydroxide groups, given as Eq. (l), has been proposed by Levine and Smith [16].Using the random mixing approximation for the distribution of thethreetypes of surface sites [i.e.,neutral,positive,and negative sites, as shown in Eq. (l)], these authors were able to derive an expression for the dependence ofon the oxide surface. The derivation has also been outlined by Smith [2],and the modified Nernst equation can be written as RT F

\vo = 2.303 -(pH,

RT 2F

r-MO-

- pH) + -In r-MOH;

The first term of the equation clearly represents the Nernst term. Therefore, Eq. (18) can be regarded as a modified version of the Nernst equation relating to the solution pH and taking into account the variation in the surface activity of H' with the surface charge. To calculate y o from the solution pH, it is necessary to know the surface density of the positively and negatively charged sites, which in turn depends on the surface charge density 00 and the two surface reaction constants (Kaand Kb) that appeared in Q. (1). Since Eq. (18) is based on the random distribution of the surface sites (first-order approximation), it still does not correctly evaluate the variation of the surface activity as the surface charge density varies. For a better approximation of \vo, the term that takes account of the variation in the surface activity in terms of the probabilities of nearest neighbor pairs, should be included. Within the framework of the random mixing approximation, Eq. (18)is also valid for a more complicated system in which the specific adsorption of multivalent electrolyte ions (i.e., adsorption at the IHP, shown in Fig. 1) does occur [17].In this case, however, r-~oin Eq. (18), for example, should be read as the density of (purely) negatively charged sites not occupied by the specifically adsorbed cations.

D.

Interfacial Electrochemical Characteristics of a Ceramic Powder Dispersion

As mentioned previously, potentiometric titration is the most widely used experimental technique to study the electrical double layer on insoluble ceramic

JANG

I68

powder (metal oxides). From the titration data, the net adsorption density of the potential-determiningions ( h + - ToH-) and, hence, the surface charge density can be determined. Further information can be obtained from electrokinetic potential potential) and colloidal stability:these properties are closely related to the charge distribution in the electrical double layer. The potentiometric titrations were first applied by Bolt [181 to the study of double layers on oxides with silica sols and later by Parks and de Bruyn [5] with &Fe203 (hematite) suspensions. Their investigations confirmed the potential-determining role of H+and OH- and indicated that the oxide double layers are significantly different from the well-characterized classic double layers at the AgI/ and HgJsolution interfaces. To illustrate the characteristic surface charge data for metal oxides, a selection of results from several laboratories and for several oxides in 0.1 M 1:1 electrolyte is depicted in Fig. 2. The results are shown in terms of CO versus ApH curves, where ApH = pH - pHpzc. A curve of 00 versus pAg for the AgI/O.l M KF interface is included for comparison. The principal feature of

(c

0.1 M

--

& , / ,,.-

-3

-2

-I,

I

c

.ys'/J

..S

A ~ H or

Aphg

"

.......... ol- Fe24/KCI PZC

88 8

/

4

Ti02/NaNO3* PZC 5.85

/?/// /.*

."

3

2

l

IO

---

1"-

15

8-5

Y - A ~ O ~ / N O C PZC I * 8-5 4-Fe00H/KCIt PZC 7.6

S i , Cab"Sil/KCI,

PZC 3

Figure 2 Experimental surface charge density data for various metal oxides and AgI

as a function of ApH(ApAg for AgI) for 0.1 M 1:l electrolyte.

I69

SCIENCE COLLOID INTERFACE

the plots in Fig. 2 (and all other surface charge results) is that the magnitude of 00 for oxides other than Si02 is considerably greater than that for the corresponding AgI/solution interface. The various ceramic powder dispersions all show essentially the same type of 00 versus ApH behavior [lo]. The much higher charges on the metal oxides and the steep rise with increase in ApH require some explanation, especially because electrokinetic measurements on metal oxide systems suggest that the diffuse double-layer potentials u/d are quite low comparable to and even less than those of AgI systems. For example, fora silicaand M KC1 suspension, Abendroth [l91 found that at pH 9 the surface charge density was -10 pC/cm2, which is considerably higher than the diffuse layer charge of 1.9 pC/cm2 calculated with E@ ( 5 ) from the corresponding 6 potential of 65 mV. Hence, at pH 9, lod/oolis approximately 20%. This example is typical of most metal oxide data [10,20]. This difference in charge (i.e., lap1 = -0.81001) must be balanced by adsorption of electrolyte ions within the region inside the hydrodynamic slipping plane. To account for the apparent strong adsorption of electrolyte ions at the ceramic powder/water interface, Yates et al. [10,21] proposed the formation of "ionpairs" at charged surface sites (the site-bindingmodel for the metal oxide/aqueous solution interface). This model and its modifications have been successfully applied to many oxide/aqueous solution interfaces in the presence of simple monovalent inorganic ions [21-231. For a ceramic powder surface in a simple electrolyte solution (e.g., W03 and NaCl), the formation of ion pairs can be represented as

- "OH

- "OH

+Y +

C

+ H+ + X-

"0-

e

*..Y++ H+

- M-OH,'

*.*X-

(19)

where YX denotes simple electrolyte present in the bulk solution phase. The formation of ion pairs or "surface complexes" readjusts the acid-base equilibria and could thereby affect the surface charge. According to the sitebinding model, the surface charge 00 determined from the proton balance represents the net number of protons released or consumed by all surface reactions, not merely theformation of the ionized surface sites -MO-and -MOH2+. Therefore, the surface charge density as determined by the potentiometric titration should be read

Increasing electrolyte concentration causes additional binding of counterions until equilibrium is reestablished, subject to theeffects of the electrostatic field. In the site-binding model, therefore, the formation of surface complexes is the principal mechanism by which protons are released or consumed by a ceramic powder surface in an aqueous electrolyte solution. The concentration of

I 70

JANG

ionized surface sites K-MoH~+ and T-Mo- in Eq. (1) or (2)] is generally small in comparison with the surface complexes (r"MoH2+ x-and T-hfo p). This is partially a result of the perturbation of surface equilibria by the electrostatic field. A s the surface charge and potential become more negative, the release of a proton by the simple dissociative reaction given as Eq. (1) is opposed by an increasing electrostatic repulsion. However, the formation of such surface complexes as - M O +'l is not retarded in an equivalent manner because of the attractive electrostatic term for thecounterion. A similar argument can be made for a positively charged surface: Thus, one of the significant points of the sitebinding model is the inclusion of surface complexation in the development of surface charge. In this manner, the surface complexation (site-binding) model can provide a mechanism for the development of surface charge in addition to the role of protons and hydroxide ions in oxidelaqueous solution systems. Figure 3 is a schematic representation of the surface complexation model for an oxide/aqueous solution interface using the NaCl electrolyte asan example of a counterion [22]. Despite successful applications to many ceramic powdedwater interfaces, the site-binding model has several ambiguities in its description of the nature S..

a-.

a..

Figure 3 The surface complexation model for a metal oxide/aqueous solution interface using NaCl electrolyte as an example of counterions.

171

SCIENCE COLLOID INTERFACE

of counterions at interfaces. That more than 80% of the counterion charges reside inside the shear plane does not necessarily indicate a strong site binding of counterions with oppositely charged surface groups at the ceramic powder/ aqueous solution interface. There is no unambiguous evidence or picture of the actual distribution or location of the counterions inside the shear plane because of the lack of knowledge of the mean distance of separation between the surface plane and the shear plane (d = p + y in Fig. 1) for the oxide/water interface. In addition, little is known about the nature of counterion interaction with the oxide surface (i.e., real chemical binding versus atmospheric condensation caused by electrostatic attraction). In view of these uncertainties in the description of the oxide/aqueous solution interface, Jang and Fuerstenau [24] sought an in situ spectroscopic method to obtain molecular-level information on the nature of the ceramic powder/ aqueous solution interface. The nuclear magnetic resonance (NMR) spin-lattice relaxation rate of quadrupole nuclei was used to probe the nature of counterion interaction with the oxide surface [24]. The model system studied was a silica/NaCl electrolyte suspension system since NaCl is a typical electrolyte present in aqueous solutions and both Na+ and C1- have nuclear quadrupole moments (quadrupole nuclei). Measuring the spin-lattice relaxation rates ofNa nuclei at two different NMR resonance frequencies and applying the theoretical equation for the chemical exchange of quadrupole nuclei between two different sites (interface and bulk regions) [25], they could estimate the quadrupole coupling constant and the correlation time zc of the interfacial Na ions [24]. In this way, the mode of counterion-silica interaction (i.e., 430- Na+) could be delineated since the electrical field gradient between two oppositely charged surface atoms (0- and Na') is proportional to the quadrupole coupling constant of the interfacial Na ions. This study clearly showed that the main driving force for the adsorption of Na ions stems from purely electrostatic interaction (atmospheric condensation), without a significant short-range chemical interaction. Therefore, the ion-pair formation represented by (19) is not a correct picture of the SiOdaqueous solution interface. Nevertheless, as mentionedpreviously, more than 80% of the counterion charges reside inside the slipping plane. One possible explanation of this discrepancy is that the location of the shear plane is rather far from the surface of charged particle. In this way, most counterion charges can reside within the slipping plane. In this case, however, they do not necessarily undergo a strong site-binding interaction with the surface groups. Theoretical analysis of the interfacial structure of a TiOdaqueous solution interface based on data for surface charge densities and electrokinetic potentials showed that the location of the shear plane is 10-15 8, away from the surface [20], which agrees with the conclusion drawn from the analysis of NMR data.

x

-e-

m.

I72

JANG

111.

COLLOIDALSTABILITY AND DISPERSION OR CONSOLIDATION

A.

Colloidal Stability

The importance of a high green density with a uniform microstructure to improve the various properties of sintered technical ceramics has been documented in numerous studies. Thus, in recent years, extensive efforts have been made to increase the homogeneity of green compact microstructures [26]. Among these, a colloidal consolidation route using a kinetically stable slip with a narrow particle size distribution has received increasing attention [27-291. It has been well known that the stability of a given lyophobic colloid sol (likea ceramic powder dispersion), according to the DLVO(DeryaguinLandau-Verwey-Overbeek) theory [13], is determined by the balance between the repulsive and attractive forces the particles experience as they approach each other. If the stability is caused by the particle charge (electrostatic stabilization), the repulsion force depends on the degree of electrical double-layer overlap. The attractive force is caused by the London-van der Waals interaction. The total potential energy of interaction VT can readily be adjusted by altering the magnitude of the repulsion, either by increasing the ionic strength of the solution or by changing the surface potential or surface charge 00 on the particles. Using the DLVO theory, the total interaction energy of two spherical particles of radius a can be written as [11,13,30]

and Aa 12So

r=2a+So

(23)

where So is the surface-to-surface distance of separation between two approaching particles and k l is the Debye-Huckel length [defined in Q. (4)]. v0 in Eq. (22) is the surface potential and can be replaced, with a good approximation, by the Stem potential v p or the potential, in practice, under the condition of the thermal Brownian collision of two particles [30]. Equation (22) is valid for a constant, low surface potential and large ka. A in Q. (23) is the effective Hamaker constant of a given colloidal dispersion and depends on the physical properties of the particles and the dispersion medium. Using the geometrical mean approximation [31], this can be expressed as A = (Asin- ApIR), where As and Ap are the Hamaker constant for solvent-solvent interaction and that for particle-particle interaction, respectively. The Hamaker constants can

c

173

SCIENCE COLLOID INTERFACE

be calculated theoretically using optical dispersion data [32,33]. These values can also be determined experimentally, and these experimental methods include [33] (1) coagulation, (2) equilibrium film thickness, (3) surface tension, and (4) attractive London dispersion force. Figure 4 is a schematic representation of the total potential energy of interaction versus particle separation. The dotted curve represents the contribution derived from the repulsive electrical double-layer interaction VR, and the dashed line denotes the contribution of the attractive London-van der Waals interaction VA. The primary maximum and the shallow secondary minimum in the potential energy curve are denoted Vmax and Vmin, respectively. The overall potential energy barrier to coagulation in the primary minimum (permanent coagulation) is given by the expression Vm, - Vmin. One should also realize that the primary minimum is bounded by the contribution of the steep Born repulsion caused by the overlapping interaction of the electronic cloud in the atoms and the structural force caused by the bounded layers of liquid at the

I

l 0

I

I 1

I 8

0 I

I l 0 I

! I

I 0

I

INTERPARTICLE DISTANCE (S, 1 Secondary MinimumcV,,,i)

/'

/ I van der Waals (VA) I I C/Primary Minimum

Figure 4 Total potential energy of interaction versus particle separation.

I74

'

JANG

particle surface. Also note that the coagulated state, in which two particles are in the deep primary minimum, is thermodynamically more stable than the dispersed state (So = -). Thus, colloidal stability achieved by the increase in the V,, does not correspond to thermodynamic stability. In most cases, ceramic powder dispersions exhibit this type of potential energy curve. Therefore, a well-dispersed ceramic suspension (or a dispersion of finely divided solid particles, in general) corresponds to a kinetically stable state (lyophobic colloid) without possessing a true thermodynamic stability. Contrary to this general notion that dispersions of finely divided solid particles cannot be thermodynamically stabilized, Stol and de Bruyn [34] discussed the conditions necessary for obtaining a thermodynamically stable dispersion (TSD) of solid particles in an aqueous solution medium. They stressed the role of the adsortion of potential-determining ions in lowering the interfacial free energy y to promote spontaneous dispersion and subsequently reasoned that for simple inorganic solids a decrease in y by about 200 mN/m relative to its value at the PZC may be sufficient to yield a TSD. However, the existence of thermodynamically stabilized dispersions of inorganic solids has not yet been demonstrated experimentally. Figure 5 shows the calculated potential energy of interaction VT of A1203 particles (a = 0.25 pm, A = 4.5 x lO-*OJ, and 0.01 M ionic strength) as a function of the surface-to-surface distance of separation for various conditions of potential in an aqueous suspension. Note that the height of the potential energy barrier increases quite sharply as the potential becomes larger than a certain critical value (-30 mV in Fig. 5). Therefore, the potential is a very good index of the magnitude of the repulsive interaction between colloid particles. Because of this, measurements of potential are most commonly used to assess the stability of a given colloidal sol. A concept of critical colloidal stability can be deduced from the DLVO equations introduced in this section. Using &S. (21), (22), and (23), the total interaction energy between two approaching spheres can be rewritten as

c

c

c

In(1 + e-Go ) - where

and

m=

AK

24m,,ebc2

c

I75

COLLOID INTERFACE SCIENCE

loo

80

r

0.01M KNOa

AhOa

60 a

c Y,

40

5’20

0

k

-20

-

40 IO

20

30

40

60

Se(nm)

Figure 5 Calculated potential energy of interaction VT of A1203 particles as a function of thesurface-to-surfacedistance of separationfor various conditions of the potential.

c

Deryaguin [l l] showed that Eq. (24) exhibits a minimum value and also a maximum (at smaller So) provided m is less than about 0.5. It is quite likely that rapid coagulation occurs near the point at which the total interaction energy rapidly decreases. Deryaguin further showed that a rapid decrease in stability can be expected when m < 1/6, which occurs when

For a given system, the quantities A and &b are fixed, so that varying the electrolyte (in type and concentration) should produce instability at a particular value of &K (Eilers and Korff rule). To the extent that 6 measures the potential characterizing the diffuse part of the electrical double layer, it ishardly surprising that it should provide a good description of the dispersion-coagulation process. A much more strict test of the dispersion-coagulation process requires an understanding of the region of slow coagulation. The overall kinetic stability

176

JANG

and the process of slow coagulation of a given colloidal dispersion can be obtained by determining the stability ratio W. When the potential is high enough to produce a significant potential energy barrier opposing coagulation, the rate of coagulation caused by the thermal Brownian encounter is slowed by a factor W. The stability ratio (ortheretardation factor) is formally defined as

<

[WO1

where kr and ks are the second-order rate constant of the rapid coagulation (Smoluchowski limit) and the rate constant of the slow coagulation caused by the existence of the potential energy barrier between two approaching particles, respectively. V d r ) is the total potential energy of interaction between two approaching particles at distance r[= 2u + So in Eq. (22) or (23)]. The stability ratio can be qualitatively viewedas the ratio of the total numberof particle collisions for a fixed time interval to the number of collisions forming permanent flocculates (or coagulates). To generate a stable sol requires a W value larger than lo3 for moderately concentrated sols. The behavior of marginally stable sols (W = 1-20) has been studied by Wiese and Healy [35]. The 6 potential and the corresponding stability ratio of A1203 suspension in this stability region are shown in Fig. 6. Note that, at a given concentration of sol particles and electrolyte, the stability ratio rapidly increases at a certain particular suspension pH and, thus, potential, supporting the concept of the critical stability (Eilers and Korff rule) mentioned previously. Further discussion of the relation between the critical colloidalstability and thepacking density of green body derived fromthe corresponding suspension is provided in the next section. These results also suggest that there is a close parallelism between the stability ratio and the potential, especially near the critical stability region. Indeed, a semiquantitative correlation between the stability ratio and the potential can bemade using the DLVO theory. Reerink and Overbeek [36] provided the following theoretical justification based on the normal distribution approximation for the shape of potential energy curve:

<

<

logW = kl logC+k, where kl = -

2.15x 1O9y;a Z

and

2

<

20

I

I

I

o o IS q1-I. A

10

0 30 q l.', lo"

0030qr'.

-

I

3

I O - ~ UK N O ~ LI KNO3

I O - ~ MK N O ~

-

5-

-

2-

-

1-

-

S B

0

2.

ii 5

(W

f8 55

I

80

1

1

90

I 95

pH

Figure 6 (a)[potentialas a function of pH for A1203 (0.15 a)in an indifferent ratio for thesame electrolyte(notspecificallyadsorbingion)solution.(b)Stability ,41203 sol as a function of pH.

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where C is the concentration of the bulk ions in moles per liter, a is the radius of particle in meters, and z is the valence of the counterions (electrolyte in the bulk phase). \~roin this equation can again be replaced by the l,potential. If the 6 potential is less than 51.5 mV at 25"C, e\1ro/2kT (= eCj2kT), which appeared in Eq. (27), can be expanded using the Maclaurin expansion (1:l electrolyte). Then Eq. (27), under this condition, can be rewritten as log W =

-2.15 X 109ae2c2 log c + k2 16(kT)2

e

(for 6 < 50 mV) should give a straight line, withthe slope being a constant positive value (log C 0), if the ionic strength of the medium remains constant. Figure 7 shows the log W versus IQ' diagram for the a-cordierite (2Mg0.2A1203-5Si02; Baikowski International Co., mean radius ~ 0 . 4pm) suspension in the.presence of 0.001 M KNO3. W C in Fig. 7 denotes the critical stability point at which there occurs a rapid increase in the kinetic stability of the dispersion. As shown in Fig. 7, there is a close linear correlation between log W and IQ' for IC1 < 45 mV. In calculating the repulsive interaction energy between colloidal particles [Q. (22)] it is usually assumed that a complete equilibrium exists between the particle surface and the dispersion medium at any separation between the particles. This is not necessarily correct, however, because the adjustment of equilibrium takes a finite time and this time may be longer than the time involved in a collision or longer even than the coagulation time. The time needed for adjusting the structure of the diffuse double layer (the relaxation time of the double layer) is equal to the average time needed for the displacement of ions across the double layer. By comparing the different time scales involved in the collision of colloid particles (i.e., double-layerinteraction,surface-charge adjustment, Brownian collision, and coagulation), Overbeek [30] concluded that the double-layer r e laxation is fast even compared withthe shorter time of Brownian encounter, but the surface-charge adjustment is too slow to follow the single encounter. In general, it may even be too slow to adjust the surface charge within the coagulation time. Consequently, in calculating the interaction energy of two particles, the condition of constant charge is probably a better approximation than thatof constant potential, and intermediate cases may be appropriate in practice. If the constant charge approximation is more appropriate than the constant potential approximation during the collision of particles, VR in Q. (22) should be replaced by the expression derived by Wiese and Healy [37],

Thus, a plot of log W with respect to

vR(00)= v R ( \ I ~ ) - ~ T E O E , ~ \ I ~ ;

In (1 - e(-21GTo))

(29)

where VR (00) denotes the potential energy of repulsion caused by the overlap of the diffuse double layers under the condition of constant surface charge.

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COLLOID INTERFACE SCIENCE 3

2

log W

1

C 2000

I

1000

I

I c (mv) l*

1

3000

Figure 7 Correlation of log W of cordierite (2Mg0.2A1203.5Si02)suspension with the square of the 6 potential.

Equation (29) also predicts that VR (00) > VR (\v) for a given suspension condition. Using Equation (22) for the expression of V R ( ~ Equation ), (29) can be rearranged as

vR(o0) = -2m0~,a\vtIn (1- e(-a~))

(30)

This equation is again valid for a low surface potential and large m.

B.

Powder Dispersion and Consolidation

Figure 8 shows a correlation of the stability ratio of an A1203 suspension (Sumitomo AKP Hp, 0.01 M ionic strength) with the packing density of A1203 green body prepared by gravimetric sedimentation. The packing densities of A1203 as a function of ILI’ and the surface charge density GO are also plotted

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in Fig. 8. The result clearly shows that the stability ratio or the 6 potential should be larger than a certain critical value for a green body with high packing density. A s shown in Fig. 8, the critical stability ratio W Cand the critical 6 potential LCof A1203 suspension are 15.8 and 36 mV, respectively. The equilibrium pH corresponding to the critical behavior in the packing density is 7.3 in the presence of 0.01 M KNO3; that is, WC,LC, and Icb(crit)l simultaneously occur at pH 7.3. These kinds of critical behaviors were also observed for centrifugally casted ZrO2 green bodies [38] and seem to be essentially independent of casting methods. The incipient point for a rapid increase in the potential energy barrier (at the 6 potential of -35 mV, or between 30 and 40 mV, shown in Fig. 5) is consistent with the critical 6 potential for a rapid increase in the packing density (36 mV, Fig. 8).

0

0

0.5

1.O

1.5

2.0

logw Figure 8 Correlation of packing density of an A1203 green body prepared by gravimetric sedimentation with log W , lclz, and lool.

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The critical behavior in the packing density when it is plotted as a function of the square of 6 potential was first demonstrated by Aksay and co-workers [29,39]. They showed that there occurs a significant decrease in the packing density as the interparticle binding energy at the cutoff point E, which is approximately inversely proportional to the square of the 6 potential, is increased. They interpreted the main cause of this rapid decrease in the packing density as the retention of third-generation void space in the colloidal solids formed by gravitational settling [29]. Assuming the sediment to be composed of clusters of a fixed number of particles and the sedimentation density to be inversely proportional to the square of the maximum radius of a cluster (aggregate of particles), they later showed [40] that the packing density exhibits a critical behavior when it is plotted as a function of the binding energy or the square of the 6 potential. The incipient point of stability (or instability) in the 6 potential can be estimated using Eq. (25), and this is 40 mV for A1203 under the condition of 0.01 M ionic strength (IC' = 3.04 nm). A remarkable similarity between the 6 potential for the calculated incipient colloidal stability (40 mV) and the critical 6 potential for a rapid increase in the packing density (36 mV; Fig. 8) indicates that a rapid increase in the packing density occurs only when the colloidal suspension nearly reaches the point of its incipient stability. Scanning electron micrographs of iwo ,41203 specimens sintered at 140OOC for 4 h are shown in Fig. 9. The specimen fabricated using a kinetically stable slip (6 = 56 mV and log W = 1.6 at pH 3) shows a homogeneous microstructure with an essentially pore-free state (Fig. 9a). On the other hand, the sample prepared under the condition of low kinetic stability (6 = 4 mV and log W = 0 at pH 9) exhibits a nonuniform microstructure, with the large secondgeneration pores [41] induced by a bridging between agglomerated units in the early stage of sintering (Fig. 9b). The colloidal approach also can be successfully used for a low-temperature sintering of technical ceramics. For example, Yeh and Sacks [42] could prepare an agglomerate-free, fine-sized a-AhO3 suspension. Homogeneous green bodies with -69% relative density and -10 nm median pore radius were prepared by slip casting the well-dispersed suspensions at pH 4 (6 = 75 mV). Samples could be sintered at 1150°C to a relative density > 99.5% and an average grain size of 0.25 pm. It was also found that the shrinkage during drying and sintering is considerably reduced compared with the well-known boehmite solgel process. The 6 potential (or interfacial electrochemical parameters, in general) has also proved to be a valuable guide to the understanding of behavior of a complicated multicomponent metal oxide dispersion. Cordierite is a multicomponent metal oxide system (2Mg0-2A1203-5Si02) and has a wide range of potential applications stemming from its important properties oflowthermal

I82

Figure 9 Microstructures of a-Al203 specimens prepared at (a) pH The compacts were sintered at 140OOC for 4 h.

JANG

3 and (b) pH 9.

expansion and dielectric constant coupled with high chemical and thermal stability. However, preparing dense cordien'te ceramics has long been a problem because of the narrow sintering range near its incongruent melting point [43]. Thus, systematic investigation of the interfacial characteristics and colloidal stability of cordierite dispersion was recently made to fabricate cordierite ceramics having a dense and uniform microstructure by a uniform colloid process [M].The electrophoretic mobility, which is proportional to the l, potential of suspension for a fixed electrolyte concentration and particle size, of aqueous cordierite suspension is shown inFig. 10. As designated inFig. 10, the cordierite (Baikowski Int. Co., average particle radius -0.4 pm, specific surface area -3 m*/g)/aqueous solution interface is characterized by three distinct points of the l, potential reversal. The result is quite striking because this type of l, potential reversal in a natural suspension of metal oxide has not been reported until now. As shown in Fig. 10, the PZR 1 of a cordierite suspension is consistent with the IEP ofa fumed silica (Aerosil 200; Degussa Co., Belgium) suspension. This indicates that the interfacial properties of an aqueous cordierite suspension in an acid region are governed by the surface silanol sites (-SiOH) and therefore suggests that the PZR 1 corresponds to the inherent IEP of an aqueous cordierite suspension. The anomalous increase in the l, potential above pH 3.5 leads to a second l, potential reversal (PZR2) at a higher pH value. This was attributed to the readsorption of dissolved aluminum hydroxy species via the surface-induced hydrolysis [U],for example,

COLLOID INTERFACE SCIENCE

183

- MOH+ Al(H20),3+ e -MO-.*.AlOH(H20)~+ 2H+

- MOH +Al(H20)T e 440----Al(OH)2(H20),++3Ht c

As the pH of a suspension increases further, the potential decreases again, revealing another point of potential reversal at PZR 3. A careful analysis of the data indicated that the surface-induced adsorption of the molecular Al(OH)3 and the subsequent formation of the hydroxylated aluminum surface sites are responsible for the PZR 3 [U].Figure 1l a is a scanning transmission electron micrograph of the cordierite core coated with the aluminum hydroxide layer of approximately 15 nm thickness. The uniform surface-induced coating of ultrafine scale aluminum hydroxide was achieved by an excess addition of aluminum salt [e.g., Al(N03)3] to the suspension at a pH below the PZR 2 and

c

3

7

5

9

11

PH Figure 10 Electrophoretic mobilities of cordierite and fumed silica as a function of suspension pH without addition of external electrolyte.

184

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Figure 11 (a) Scanning transmission electron micrograph of the cordierite core uniformly coated with the aluminum hydroxide layer; (b) EDAX spectrum of the core region shown in a; (c) EDAX spectrum of the coated layer shown in a. slowly adjusting pH level above the PZR 3 to induce the surface-induced precipitation. The energy-dispersive analyticalx-ray (EDAX) spectrum clearly indicates that the coated layer (Fig. l l c ) is exclusively composed of aluminum hydroxide.

SCIENCE COLLOID INTERFACE

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When particles are suspended in a nonaqueous medium, the electrical double layer is difficult to develop (because of the instability of ions in a nonaqueous medium), and therefore inmost cases the electrostatic stabilization represented by Eq. (22) is not expected. In this case, stabilization of a colloidal dispersion can still be achieved, under suitable thermodynamic conditions, by the adsorbed polymer chains. Indeed, polymers have been widely used since antiquity to stabilize colloidal particles against coagulation. When two surfaces, from which flexible long chains are protruding into the solution, come close together (Fig. 12), there occurs a sharp repulsive interaction (steric repulsion) above the so-called 8 point [45]. Two effects seem to contribute to the steric repulsion. In the narrow gap between the surfaces the long polymer chains lose some of their conformations (volume restriction effect). This results in a partial loss of chain entropy, in an increase in the free energy, and thus in an entropic repulsion. Furthermore, the concentration of polymer segments in the gap increases, and this osmotic effect results in another contribution to the repulsion. Figure 12 illustrates this schematically. When two adsorbed polymer layers interpenetrate each other as schematically shown in Fig. 13, the repulsive energy of interaction under the condition of constant segment density can be written as [45]

where 61S0126, o is the weight of stabilizing moieties (polymer chains) per unit surface area, v2 is the partial specific volume (volume per unit weight) of the moieties, v1 is the volume of solvent molecule, No is the Avogadro number, and 6 is the thickness of the adsorbed polymer layer (Fig. 13). X is the

Figure 12 Thevolumerestrictioneffect(a)andtheosmoticeffect tion by adsorbed or chemically bound long chains.

(b) in stabiliza-

186

JANG

P

overlap (interpenetration) region

Figure 13 Stericstabilizationcaused by theoverlap of adsorbedpolymerlayers upon close approach of dispersed solid particles.

Nory-Huggins interaction parameter and is less than X for the steric repulsion or above the 0-point. Therefore, in steric stabilization in a nonaqueous suspension, Eq. (22) for VR should be replaced by Eq. (31). Since the steric repulsion increases sharply (hard-sphere type of repulsion) when two adsorbed polymer layers interpenetrate .each other under the condition of x IX, this provides a very efficient means of colloidal stabilization in nonaqueous solvent systems. If the adsorbedpolymer is a polyelectrolyte, which carries i charges per monomer unit, and the particles are suspended in an aqueous solvent containing CSmol of 1:l valence salt per unit volume (liter), the repulsive energy of interaction expressed in Eq. (31) should be modified as [46]

where Amix is defined in Eq. (31) and Vu is the molar volume of a monomer unit. Thus, for electrosteric stabilization in an aqueous suspension, should be smaller than % + i2/(4VuCS). The term T/(4VuCS)represents the charge-charge repulsion arising from the intermixing of the charged polymeric chains when two particles approach each other.

x

SCIENCE COLLOID INTERFACE

I87

W. FABRICATIONOFCERAMIC COMPOSITES Various methods of colloid interface science are now extensively applied to advanced ceramic processing, including powder synthesis, powder dispersion and consolidation, and microstructural control of green and sintered bodies (especially for ceramic composites). In the author's judgment, the three most important applications of colloid interface science to modem ceramic powder processing are (1) the synthesis, dispersion, and packing of monodisperse (or uniformly sized) powder via the solution-precipitation reaction of relevant metal alkoxides, initiated by Barringer et al. [26,27]; (2) the controlled transformation and sintering of sol-gel derived bodies by the submicrometer-sized colloidal seeds, pioneered by Messing and coworkers [47,48]; (3) the nanometer-scale multiphasic sol-gel processing of multicomponent ceramics, proposed by Roy and Roy [49,50]. Since there are many references for these, I do not review and discuss these in this chapter. Instead, discussion is provided of several recent examples of the applications of colloid interface science to the fabrication of advanced ceramic composites. In the fabrication of an A1203/SiC whisker composite, Sacks and coworkers demonstrated that colloidal processing can be effectively used to prepare homogeneous green and sintered bodies by pressureless sintering [51]. By using the combination ofpH adjustment to -4 and polyelectrolyte addition, A1203 particles and S i c whiskers could be codispersed at a high overall solid concentration (-50 ~01%)while maintaining the relatively low viscosities (-100 mPas at the shear rate of 10 S") desired for casting operations. Green bodies with a high relative density (6670%) were obtained with S i c whisker contents in the range of 5-30 ~01%.Although densification was inhibited by the S i c whiskers, significantly higher sintered densities were obtained by colloidal processing than by conventional processing. Samples with 15 vol% whiskers could be pressurelessly sintered to -97% relative density with zero open porosity and -1.6 pm matrix average grain size [52]. Therefore, these results suggest that the colloidal processing route can be used successfully in the fabrication of S i c whisker-reinforced matrix composites by pressureless sintering. As discussed previously, the DLVO theory predicts that the van der Waals attractive potential combines with the long-range repulsive electrostatic potential to produce the net interparticle potential, which can be either repulsive or cohesively attractive, depending on the magnitude of the repulsive potential. If there exists a very short range ( 4nm) repulsive potential instead of the longrange repulsive potential caused by the diffuse electrical double layer, the approaching particles are weakly attracted to one another but highly repulsive when pushed together [53]. A schematic diagram of the interparticle potential corresponding to this situation is shown in Fig. 14. Under certain conditions, the basal surface of mica is known to develop a short-range repulsive potential called the hydration (or solvation, in general)

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Coagulated

o"0 0 ,V

non-touching

network

Waals

Potential

Figure 14 Theformationof a weakly attractive non-touching network in the presence of a short-range hydration (solvation) potential.

potential. Using the surface force apparatus, Pashley [54] showed that when mica is immersed in an aqueous solution containing a high concentration of potassium salt, the K+ion in its structural surface site can be strongly hydrated by water molecules. When no long-range repulsive potential is present, surface force measurements showed that the net surface potential has a functional form similar to that shown in Fig. 14, that is, weakly attractive at separation distances > -2 nm and highly repulsive at smaller separations. More recently, Lange and coworkers [55] showed that a short-range repulsive potential can be also developed on the surfaces of A1203 particles. When sufficient salt was added to the slurry, it was noted that the dispersed suspension (as a result of long-range electrostatic repulsion at a low pH) was altered to what superficially appeared to be a coagulated network. Unlike a flocculated network produced by changing the pH to the PZC,the.particles in this attractive network (coagulated network) were easily arranged, as deduced from the rheological data. This suggests that the particles slip into the potential energy well of shallow depth to produce an attractive but nontouching particle network called a coagulated network, as schematically shown in Fig. 14. They [53] further demonstrated that the mixed A1203/ZrO2 green bodies formed from welldispersed suspensions exhibited significant phase segregation, whereas those formed with coagulated slumes under the presence of the short-range hydration potential exhibited uniform microstructures throughout. Since the control of rheological properties (e.g., shear thinning behavior with a suitable viscosity value) and the minimization of phase segregation (differential settling) are very important in the fabrication of ceramic composites from slumes, the processing scheme based on the weakly coagulated slurry seems to be a promising potential approach. Similarly, in the fabrication of a ceramic composite, Bleier and coworkers [56,57] found that the uniformity of the relative phase distribution can be greatly improved when the potentials of the unlike particles are of the same sign but differ greatly inmagnitude. For example, in the fabrication of an

c

COLLOID INTERFACE SClENCE

189

~203-Zxfkcomposite using a colloidal processing route, the unlike particles (i.e., A1203 and ZrO2) experience a weak, mutual association in the pH range 5-6. However, when they are more similar electrostatically (e.g., in magnitude of potential) at a lower pH, they do not mutually associate and the spatial uniformity of the ZrO2 phase in the resulting green body and sintered composite suffers [56,57]. They concluded that the observed phenomena are associated with a weak attractive interaction in the secondary minimum in the potential energy curve (Fig. 4) and suggested that the secondary minimum for the interaction between the dissimilar oxide particles is a useful, quantitative guide for predicting conditions under which a homogeneous distribution of constitutive phases (also without extensive agglomeration) can be obtained. Since they also emphasized a weak attractive interaction between the nontouching particles in a shallow potential energy well, their approach very much resembles the previously discussed suspension processing that utilizes the concept of the short-range repulsive hydration potential. To eliminate fundamentally problems associated with phase segregation and agglomeration in the fabrication of ceramic composite, Jang and Moon [58] proposed another colloidal processing scheme in which a thermodynamic theory of interfacial electrochemical phenomena is applied. In this approach, a kinetically stable colloidal suspension for the matrix phase (after sintering) is first prepared. Then, precursor precipitates (e.g., zirconium hydroxide ina Al203-ZrO2 composite) for the dispersed phase are individually coated on the matrix-phase colloid particles. The selective, interface-induced nanometer (or submicrometer) scale coating of the precursor on the matrix-phase particle was termed surface-induced coating. In the fabrication of zirconia-toughened alumina (ZTA), for example, one should carefully examine the following three points to set up optimal conditions for the surface-induced coating: (1) the pH range for the formation of kinetically stable A1203 dispersion, (2) the pH range for the enhanced interfacial concentration of ionic species (e.g., zirconium cation and hydroxide anion) needed for a selective formation of the precursor phase for ZrO2 only at the interface; (3) the pH range in which the other two conditions overlap each other. The enhanced interfacial concentration of the ionic species i for the surfaceinduced coating on the matrix phase particle was derived based on the interface electrochemical equilibrium and can be expressed as [58].

c

where i can be either zirconium cation or hydroxide anion (for the fabrication of ZTA), Ci(r) is the concentration of the ionic species i at a distance r from the charged A1203 surface, Ci(m) is the corresponding bulk concentration, and y ( r ) is the mean coulombic potential at r. @so~vrefers to the change in the sol-

190

JANG

vation energy of ionic species upon adsorption, and denotes any possible short-range chemical interaction free energy (or any free energy change other than the electrostatic and the solvation free energy). A semiquantitive calculation based on the preceding equation clearly indicated that the charged A1203 interface, under suitable conditions can induce a significant interfacial enrichment of both hydroxide and zirconiumions and thus provides a thermodynamic basis for selective surface-induced nucleation at the charged colloid interface. Based on the preceding proposition, homogeneous Al203-ZrO2 composite powder was fabricated by the surface-induced coating of the ultrafine precursor on the kinetically stable colloid particles of A1203 [58]. The composite prepared by this processing scheme was characterized by a uniform spatial distribution of the dispersed ZrO2 phase and by the absence of large ZrO2 grains formed fromhard ZrO2 agglomerates. The composite also showedahighly uniform grain size distribution ofboth the dispersed ZrO2andthe matrix A1203 phases. The uniform grain size distribution of the matrix phase indicates that the homogeneous coating of the ultrafine precursor particles effectively pins the A1203 grain boundaries. More recently, this processing scheme was also applied to the fabrication of Al203-ZrO2-SiCwhisker composites [59] with uniform microstructures and large values of Klc (-12.5IMPa-m'R). The concept on which the surface-induced coating route is based is illustrated in Fig. 15 using the A1203-ZrO2-SiC whisker system as an example. Sacks and coworkers [60,61] proposed another novel approach for enhancing densification and controlling microstructure development of multicomponent ceramics and ceramic-ceramic composites based on the processing of composite particles via a sol-coating route. In this method, they prepared cores of crystalline materials (A1203, Si3N4, Sic, and ZrO2) with outer coatings of amorphous silica. Silica coatings were made using the hydrolysis-condensation route of silicon alkoxides (e.g., tetraethylorthosilicate).Before deposition of the coating, the core powders were fractionated using colloidal size separation. Powder compacts prepared with composite particles showed enhanced densification caused by viscous flow of the amorphous coating and the increased threshold concentration for percolation (connected network) of the crystalline phase. An example is illustrated in Fig. 16, in which the core is alumina and the coating is silica. As shown schematically in Fig. 16, if the size of the composite particles is selected properly, this composite structure should viscously sinter at a relatively low temperature to form an alumina/siliceous glass composite. At a higher temperature, the core and coating would react to form mullite. This process was referred to as transient viscous sintering. The amorphous sol-coating approach was also successfullyused to fabricate mullite-matrix composites [60,61] by mixing the silica-coated alumina particles with other silica-coated particles (e.g., zirconia particles for mullite/zirconia composite and S i c particles for mullite/SiC composite).

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COLLOID INTERFACE SCIENCE Kinetically Stable A G and S i c 0 Colloidal Suspensions

Unifonn. Mixed Colloidal Suspension of Al&'SiCoN)

with SubmicronAParticles attached to Sic whisker via Heterointeraction

Nanometer Scale Surface-induced Coating of the Precursor for the Dispersed Z Q Phase

Sintered Composite Characterized by Uniform Spatial Distribution of the Dispersed Phases and few Hard Agglomerates

Figure 15 Homogeneousfabrication of an A1203-Zr02-SiCwhiskercomposite by surface-induced coating. The preparation of a kinetically stable mixed suspensionand the subsequent compression of the electrical double layer, while maintaining the surface potentials \v0 near constant values, proved to be a useful strategy for the fabrica-

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3 4 4+

2s14

-

3az4 +

m

___)

3Al~O,*2SQ(M~llit0)

Figure 16 Transient viscous sintering of mullite using amorphous silica-coated alumina particles.

tion of ceramic composites. For example, based on the electrokinetic properties of aqueous silica, boehmite, and zro;? dispersions, cordierite-Za composites were fabricated by a mixed colloidal processing route [62]. In this method, stable boehmite, silica, and Zro.2 suspensions were f i s t prepared separately and then mixed while continuously stirring at pH 3-4, at which the 6 potential of the mixed suspension exhibits its maximum. The mixed suspension was quite stable 45 mV). The addition of Mg(N@)2 solution to the mixed suspension then increased the ionic strength and caused it to gel quickly by the compression of the electrical double layer (decrease in the Debye-Hiickel length without significantly decreasing vo).This essentially eliminated the problems associated with the phase segregation (or differential settling) and produced a composite powder with a uniform spatial distribution of constitutive phases. The fabricated composite was characterized by a dense and homogeneous microstructure andbyauniform spatial distribution ofsubmicrometer-sized tetragonal Zro;! particles throughout the matrix [62]. This processing scheme was also successfully applied to the fabrication of S i c whisker-reinforced lithium aluminosilicate matrix composites [63].

(c -

REFERENCES 1. Parsons, R.,Modern Aspects of Electrochemistry, Vol. 1 (J. 0”. Bockris and B. E. Conway, eds.), Butterworths,London, 1954, p. 103. 2. Smith, A. L., Electrical phenomena associated with the solid-liquid interface, Dispersion of Powders inLiquids (G.D. ParFtt, ed.),AppliedScience,London, 1985, p. 99. 3.Modi, H. S., andFuerstenau, D. W., Streamingpotentialstudies on corundum (a-Al203) in aqueous solutions of inorganic electrolytes,J. Phys. Chem., 61, 640 (1957). 4. Parks, G. A., Theisoelectricpoints of solid oxides, solid hydroxides, andaqueous hydroxo complex system, Chem. Rev., 65, 117 (1965).

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5. Parks, G. A., and de Bruyn, P. L., The zero point of charge of oxides, J. Phys. Chem., 66, 967 (1962). 6. Grahame, D. C., The electrical double layer and the theory of electrocapillarity, Chem. Rev., 41, 441 (1947). G. The Electrical Double 7. Loeb,A. L., Wiersema,P.H.,andOverbeek,J.T. Layer Around a Spherical Colloid Particle, M.I.T. Press, Cambridge, MA, 1961. 8. Levine, S., Mingins, J., and Bell, G. M., The discrete-ion effect in ionic doublelayer theory, J . Electroanal. Chem., 13, 280 (1967). and Healy, T. W., Discreteness of charge and solva9. Wiese, G. R., James, R. 0.. tion effects in cation adsorptionat the oxide/water interface,Disc. Faraday Soc., 52, 302 (1971). 10. Yates, D. E., The structure of the oxide/aqueous electrolyte interface, Ph. D. Thesis, University of Melbourne, Australia, 1975. 11. Hunter, R. J., Zeta Potential in Colloid ScienceLPrinciples and Applications, Academic Press, New York, 1981, Chaps. 4 and 6. 12. Babchin,A. J., Chow, R. S., and Sawatzky, R. P., Electrokinetic measurements by electroacoustical methods, Adv. Colloid Interfac. Sci., 30, 11 1 (1989). 13. Overbeek, J. Th.G., The interaction between colloid particles, in Colloid Science (H. R. Kruyt, ed.), Elsevier, Amsterdam, 1952, p. 245. 14. BCruM, Y. G., and de Bruyn, P. L., Adsorption at the rutile-solution interface. I. thermodynamic and experimental study, J. Colloid Interfac. Sci., 27, 305 (1968). 15. Hunter, R. J., and Wright, H. J. L., Dependence of electrokinetic potential on concentration of electrolyte, J . Colloid Interjac. Sci., 37, 564 (1971). 16. Levine, S., and Smith,A. L., Theory of the differential capacity of the oxide/aqueous electrolyte interface, Disc. Faraday Soc., 52, 290 (1971). 17. Jang, H. M.,andFuerstenau,D.W.,Thespecificadsorptionofalkaline-earth cations at the rutile/water interface, Colloids Surfaces, 21, 235 (1986). 18. Bolt, G. H., Determination of the charge density of silica sols, J . Phys. Chem., 61, 1166 (1957). J. Colloid 19. Abendroth, R. P., Behavior of a pyrogenic silica in simple electrolytes, Interfac. Sci., 34, 591 (1970).

20. Jang, H. M., The natureof counter-ion adsorption at the oxide/water interface, Ph. D. Thesis, University of California, Berkeley, 1986. 21. Yates, D. E., Levine, S., and Healy, T. W., Site-binding model of the electrical double layer at the oxide-water interface, J . Chem. Soc. Faraday Trans. I , 70, 1807 (1974). 22. Davis, J. A., James,

R. O., and Leckie, J. O., Surface ionization and complexation at the oxide/water interface, J. Colloid Interfac. Sci., 63, 480 (1978). 23. Houchin, M. R., and Warren, L. J., Surface titrations and electrokinetic measurements on stannic oxide suspensions, J. Colloid Interfac. Sci., 100, 278 (1984). 24. Jang, H. M., and Fuerstenau, D. W., The nature of simple monovalent cation-silica interaction as reflected in the spin-lattice relaxation time of =Na, Langmuir, 3, 11 14 (1987). 25. Bull, T. E., Nuclear magnetic relaxation of spin-312 nuclei involved in chemical exchange, J. Magn. Reson., 8, 344 (1972).

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JANG

26. Barringer,E. 27. 28. 29. 30. 31. 32.

A., andBowen, H. K., Formation,packing,andsinteringof monodisperse Ti02 powders, J. Am. Ceram. Soc., 65, (12) C199 (1982). Bamnger, E. A., Jubb, N., Fegley, B., Pober, R. L., and Bowen, H. K., Processing monosized powders, in Ultrastructure Processing of Ceramics, Glasses, and Composites (L. L. Hench and D. R. Ulrich, eds.), Wiley,New York, 1984, p. 315. ner, R. K.,Inorganic colloids for forming ultrastructures, in Science of Ceramic Chemical Processing (L. L. Hench and D.R. Ulrich, eds.), Wiley, New York, 1986, p. 3. Aksay, I. A., and Kikuchi, R., Structuresof colloidal solids, inScience of Ceramic Chemical Processing (L. L. Hench and D. R. Ulrich, eds.), Wiley, New York, 1986, p. 513. Overbeek, J. T. G., Recent developmentsin the understanding of colloid stability, J. Colloid Interfac. Sci., 58, 431 (1977). Hiemenz, P. C., Principles of Colloid andSurface Chemistry, 2nd ed., Marcel Dekker, New York, 1986, Chap. 11. Gregory, J. The calculation of Hamaker constants, Adv. Colloid Interfac. Sci., 2,

396 (1969). 33. Visser, J., On Hamaker constants: A comparison between Hamaker constants and Lifshitz-van der Waals constants, Adv. Colloid Interfac. Sci., 3, 331 (1972). 34. Stol, R. J., and de Bruyn, P. L., Thermodynamic stabilization of colloids, J. Colloid Interfac. Sci., 75(1), 185 (1980). 35. Wiese, G. R., and Healy, T. W.,'Coagulation and electrokinetic behavior of Ti02 and A1203 colloidal dispersions, J . Colloid Interfac. Sci., 51, 427 (1975). 36. Reerink, H., and Overbeek, J. T. G., The rate of coagulation as a measure of the stability of silver iodide sols, Disc. Faraday Soc., 18, 74 (1954). T. W., Effect of particle size oncolloidalstability, 37. Wiese, G.R.,andHealy, J. Chem. Soc. Trans. Faraday Soc., 66,490 (1970). 38. Jang, H. M., and Lee, K. G., Effects of kinetic stability of colloidal dispersion on

thegreenmicrostructureanddensificationbehavior of zirconia, in Ceramic Transactions Vol. 12, Ceram. Powder Science 111 (G. L. Messing, S. Hirano, and H. Hausner, eds.), Am. Ceram. Soc., Westerville, OH, 1990, p. 383. mi39. Aksay,I.A.,andSchilling,C.H.,Colloidalfiltrationroutetouniform crostructures, in Ultrastructure Processing of Ceramics, Glasses,and Composites (L. L. Hench and D. R. Ulrich, eds.), Wiley, New York, 1984, p. 439. 40. Shih, W. Y., Aksay, I. A., and Kikuchi, R., Reversible-growth model: Clusterclusteraggregationwith finite bindingenergies, Phys. Rev. A, 36(10),5015

(1987). 41. Zhao, J., and Harmer, M. P., Effect of pore distribution on microstructure development II. First- and second-generation pores, J. Am. Ceram. Soc., 71(7), 530 (1988). 42. Yeh, T.-S.,and Sacks, M.D., Low-temperaturesintering of aluminumoxide, J. Am. Ceram. Soc., 71(10), 841 (1988). 43 Suzuki, H., Ota, K., and Saito, H., Preparation of cordierite ceramics from metal alkoxides (Part II). Sintering, Yogyo-kyokai-shi,95(2), 170 (1987). 44. Jang, H.M.,and Lee, S. H., Interfacial characteristics and colloidal stability of cordierite dispersion, Langmuir, 8(7), 1698 (1992). I

COLLOID INTERFACE SCIENCE 45. 46. 47. 48.

49. 50. 51. 52.

195

Napper, D. H., Polymeric Stabilization of Colloidal Dispersions, Academic Press, New York, 1983, Chapts. 10 and 12. Takahashi, A., and Nagasawa, M., Excluded volume of polyelectrolyte, J . Am. Chem. Soc., 86, 543 (1964). Kumagai, M., and Messing, G. L., Controlled transformation and sintering of a boehmite sol-gel by a-alumina seeding, J . Am. Ceram. Soc., 68(9), 500 (1985). Messing, G. L., Kumagai, M., Shelleman, R.A.,andMcArdle,J.L.,Seeded transformations for microstructuralcontrol in ceramics, Science of Ceramic Chemical Processing (L. L. Hench and D.R. Ulrich, eds.), Wiley, New York, 1986, p. 259. Roy, R. A., and Roy, R., New metal-ceramic hybrid xerogels, in Abstracts Ann. Meeting Mater. Res. Soc., Boston, MA, 1982, p. 377. Roy, R., Ceramics by the solution-sol-gel route, Science, 238, 1664 (1987). Sacks, M. D., Lee H.-W., and Rojas, 0. E., Suspension processing of A1203/SiC whisker composites, J. Am. Ceram. Soc., 71(5), 370 (1988). Lee,H.-W.,and Sacks, M.D., Pressureless sintering of Sic-whisker-reinforced A1203 composites. I. Effect of matrix powder surface area, J . Am. Ceram. Soc.,

73(7), 1884 (1990). 53. Lange, F. F. New interparticle potential paradigm for advanced powder process-

54. 55. 56.

57. 58. 59.

ing, in Ceramics Transactions, Vol. 22, Ceramics Powder Science N (S.-I. Hirano, G. L. Messing, and H. Hausner, eds.), Am. Ceram. Soc., Westerville, OH, 1991, p. 185. Pashley, R.M., DLVO and hydration forces between mica surfaces in LP, Na', K+,and CS+ electrolyte solutions: Correlation of double-layer and hydrationforces with surface cation exchange properties, J . Colloid Znterfac. Sci., 83, 531 (1981). Velamakanni, B. V., Chang, J. C., Lange, F. F., and Pearson, D. S., New method for efficient particle packing via modulation of repulsive lubricating hydration forces, Lungmuir, 6(7), 1323 (1990). Baik, S., Bleier, A., and Becher, P. F., Preparation of A1203-zfl2 composites by adjustment of surface chemical behavior, Mater. Res. Soc. Symp. Proc., Vol. 73, Better Ceramics Through Chemistry (C.J. Brinker, D. E. Clark, and D. R. Ulrich, eds.), Mater. Res. Soc., Pittsburgh, PA, 1986, p. 791. Bleier, A., and Westmoreland, C. G., Effects of pH and particle size on the processing of and the development of microstructurein alumina-zirconia composites, J . Am. Ceram. Soc., 74( 12), 3100 (1991). Jang, H.M., andMoon, J. H.,Homogeneousfabricationanddensification of zirconia-toughenedalumina(ZTA)composite by thesurface-inducedcoating, J . Mater. Res., 5(3), 615 (1990). Jang, H. M., Moon, J. H., and Jang, C. W., Homogeneous fabrication of Al2O3ZrOz-Sic whisker composite by the surface-induced coating, J . Am. Ceram. Soc.,

75( 12), 3369 (1992). 60. Sacks, M. D., Bozkurt, N., and Scheiffele, G. W., Fabrication of mullite and mul-

lite-matrix compositesby transient viscous sintering of composite powders, J . Am. Ceram. Soc., 74(10), 2428 (1991). 61. Sacks, M. D., Scheiffele, G. W., Bozkurt, N., and Raghunathan, R., Fabrication of ceramics and composites by viscous sintering of composite particles, Ceram-

196

JANG ics Transactions, Vol. 22, Ceramics Powder ScienceN (S.-I. Hirano, G. L. Mess-

ing, and H. Hausner, eds.), Am. Ceram. Soc., Westerville, OH, 1991, p. 437. H.M.,Homogeneousfabricationanddensificationof cordierite-zirconiacomposites by amixedcolloidalprocessingroute, J . Am.

62. Lim, B. C., andJang,

Ceram. Soc., 76(6), 1482 (1993). 63. Jang, H. M., Kim, K. S., and Jung, C. J., Development of Sic-whisker-reinforced

lithium aluminosilicate matrix composites by a mixed colloidal processing route, J . Am. Ceram. Soc., 75(10), 2883 (1992).

Ceramic Particles,in Nonaqueous Media Burtrand 1. Lee Clemson University Clemson, South Carolina

1.

INTRODUCTION

Properties of final ceramic products made from fine powders are often dependent on the state of dispersion of the paiticles in the liquid media. The particles interact with each other and with what is in the medium. Some of the liquid molecules can adsorb onto the solid surface. Some solid surfaces can be ionized to exhibit surface charges. Ceramic particles in organic liquid media behave differently from those in aqueous media. Many known ceramic-processing techniques involve fine particles dispersed in nonaqueous media. Some of the better known examples are tape casting of barium titanate and aluminum nitride powders and tape or slip casting, extrusion, and injection molding of ceramic high Tc superconductor powders. The theories of colloidal stability used for aqueous slurry systems may also apply to nonaqueous systems. According to the Derjaguin equation [l], the repulsive energy can be written as V, = 2 m y d In(l+ e-KH)

(1)

where: VR= electrostatic repulsive energy between two like particles E = dielectric constant of medium a = particle radius 197

198

LEE

\vd =

electrical potential at the diffuse layer

H = interparticle distance K

= Debye-Huckel parameter

The attractive energy, on the other hand, is given by

vA -"AQ 12H

where: VA = attractive energy between two identical particles A = Hamaker constant

The repulsive energy between two particles of size Q in a medium is a function of the dielectric constant of the liquid medium. This means a smaller VR for media of organic solvents with smaller E than that of water. The dielectric constants of some common liquids are listed in Table 1. The Debye-Hiickel parameter K in Q. (1) also depends on E by

where:

e = electron change NA = Avogadro's number C = electrolyte concentration 2 = charge of the ionic species

.

The double-layer thickness 1 / is ~ therefore larger for larger E. However, the ionic strength, C in nonaqueous media is much smaller than that in aqueous media, hence larger VK. The attractive interaction depends largely on the Hamaker constant A, as shown in Eq. (2). The larger the value of A, the greater is the attractive energy between the particles. The net interaction energy is then the sum of VR and VA. Equations (1) and (2) show that a better dispersion of fine particles must come from systems having large \vd and small A. Fowkes and Lee and Rives [2,3] showed that ion formation in a nonaqueous medium is appreciably higher than predicted by Eq. (1). Steric stabilization of colloidal particles is achieved by adsorption of long-chain molecules on solid surfaces [4]. For an effective steric stability by the adsorbed molecules, there are several basic requirements: (1) long chain of the adsorbate molecule, (2) strong adsorption of the adsorbate on the solid particle surface, and (3) good solubility of the chain in themedium [4]. Regarding the chain length of a steric stabilizer, there is the question of how

CERAMIC PARTlCLES IN NONAQUEOUS MEDIA

199

Table 1 RelevantPhysicalProperties of Solvents Dielectric Solvent Chloroform Ethanol Methylene chloride Water Tetrahydrofuran Acetone Ethyl acetate Toluene

Viscositya (ea-s) 0.58 l .2 0.45 1 .o 0.55 0.32 0.46 0.59

Ammd

(J x 10-20)

Acidlbasec HPCPb constant 4.8 25 9.1 78.3 7.4 21.5 6.4 2.4

aAt 20°C. bHydrogen bonding cohesion parameter. cBased on Drago’s C and E values [K!], (strong) 1 dcalculated Hamaker constants [5]. Source: From Refs. 15 and 16.

6.3 20 9.6 40.4 7.4 6.7 11

8.9 1.6

A1

A2 A3 A4

B1 B2 B3 B4

7.6 6.4 7.4

204

6.6 6.7 8.4

4 (weak).

long the chain must be for effective steric stability. It has been shown by a number of experimentalists that the chain length need not be so long as predicted by the theory [2-51. Relatively small molecules, such as fatty acids and some surfactants, are widely used as dispersants. Calvert et al. [6] showed that the maximum packing density of silica suspension in hexane was achieved by a silane surfactant with chain length of carbon number 12. A chain length longer than 12 did not improve the packing density. Surfactants that produce electrostatic charges on ceramic particles exhibit large changes in the viscosity and the suspension stability. When steric stabilization is combined with electrostatic stability, the results can be remarkable [3,5]. The more additives used in ceramic slurries for processing, the more complex interactions among the components in the system are expected. Better understanding of these interactions is necessary to control particle dispersion and processing in nonaqueous media. The interactions are often indirectly assessed by the rheology of the suspension systems, the 6 potentials, and the green body characteristics. This chapter presents some of the results and an explanation of the rheological and related behavior of alumina, silica, and portland cement particles in various organic solvent media.

II.

RHEOLOGY AND ELECTROSTATICCHARGE

A.

Alumina

The surface of alumina in aqueous environment is well characterized [7-lo]. It has amphoteric sites, that is, acidic sites of aluminum and basic sites of hy-

200

LEE

droxyl. However, the isoelectric point and the point of zero charge of pH = 9 indicate that the surface is weakly basic in water. Although basic molecules can adsorb on an alumina surface via the acidic aluminum sites, more effective adsorption of dispersant molecules is expected for the hydroxyl groups via hydrogen bonding. Figure 1 shows the viscosity of 47 vol% alumina suspensions having the mean particle size of 1.5 pm (Alcoa A-152) in three different organic liquids in the presence of 5 wt% linolenic acid (LNA) per alumina. At low shear rates, they all exhibit shear thinning. Ethyl acetate (EtAc) exhibits the highest viscosity, followed by toluene (Tol). LNA, which is an acid, can adsorb on alumina via acid-base interaction or hydrogen bonding. The structural formula of LNA and those of other organic dispersants are given in Fig. 2. The interaction between the solvent and LNA cannot be responsible for the viscosity curves because the most basic solvent, tetrahydrofuran (THF) [1l], would interact with LNA the most. The acid-base interaction is based on Drago's work on Lewis acid-base interaction energies [12]. Fowkes [2,13,14] expanded the concept to ceramic processing. The relative acidity and basicity of solvents used are listed in Table 1. The ceramic powders used are listed in Table 2. The rheogram in Fig. 1 suggests two possibilities: one is the largest amount of LNA adsorption onto an alumina surface in THF and the least in EtAc for steric stability. The other possibility is that the alumina surface is the most highly charged in THF and the least in EtAc for electrostatic stability. Our electrophoretic deposition experiments yielded alumina particles deposition on the positive electrode, which signifies that the alumina surface in THF was

t. U

E

>

IO

'

0 3.4 6.8 10.2 13.6 17 20.4 23B SHEAR RATE (I/S)

Figure 1 Rheogram of 47 vol% alumina in threedifferentsolvents LNA per A-152 alumina.

with 5 W%

201

CERAMIC PARTICLES IN NONAQUEOUS MEDIA

Linolenic Acid (LNA), C-18 0 €I$ = CH

II

-C

OH I

- 0 - C%CH - q

H I -N

- (CHJ,Si(OEt),

N-(3-Acryloxy-2-hydroxy-propyl)-3-aminopropyl~ethoxysilane(AHAS), 15 memberChain

CH3 \

“C-

l

CH3 x=14.3

1,

0

Figure 2 Chemicalstructure of dispersantmolecules.

negatively charged. In EtAc a smaller amount of alumina was deposited on the positive electrode. This means that EtAc is weakly basic relative to the surface of alumina, and THF is-more strongly basic. Alumina in toluene yielded an even smaller deposit of alumina on the positive electrode than in EtAc. This agrees with the basicity scale of Drago shown in Table 1.

Table 2

Characteristics of CeramicPowders Mean particle Surface area (p) alumina) (mVg

Source Powder Alumina A-l52 AKP-15 Silica Portland cement, type I

Alcoa Sumitomo Stiiber Holnam Santee

1.S

0.65 0.5 10

1.2 3.8 7.9 0.16

0

202

LEE

Although toluene and EtAc are both listed as weak bases, this is the relative basicity to the solid surface. Hence, one can conclude that alumina is a weaker acid in toluene and EtAc, but a stronger acid in THF. The electrophoretic deposition of alumina in the presence of 5 wt% LNA per alumina yielded a deposit on the negative electrode in toluene and on the positive electrode in THF and EtAc. Although LNA is an acid (Fig. 2) relative to the alumina surface, the basic media THF and EtAc overrode the acidity of 5 wt% LNA, causing the negative charge on the alumina surface. In toluene, on the other hand, LNA reversed the charge on the alumina surface. This could be possible because the surface charge of alumina in toluene was nearly zero, shown by the very small amount of deposit of alumina particles on the positive electrode in the absence of LNA. Figure 3 sdows viscosities of 60 vol% AKP-l5 alumina suspensions as a function of shear rate in THF and chloroform (Chl), all with 5 wt% LNA. Both solvents in the presence of LNA exhibited that the suspensions are pseudoplastic with Chl as a better solvent than THF for alumina. A strong positive charge of the alumina surface in chlorofomLNA and a strong negative charge in THFLNA were indicated by the electrophoretic deposition. The surface charges and 6 potentials of alumina are determined primarily by the solvent media. As given in Table 1, choloroform has been classified as a strong acid and THF as a strong base, which is supported by the electrophoretic deposition results.

10000 h

'p

P

a

E

v

1

0

200

400

600

800

1000

Shear Rate (I/sec) Figure 3 Rheogram of 60 vol% alumina in THF and chloroform with 5 wt% LNA per AKP-15 alumina.

CERAMIC NONAQUEOUS PARTICLES INMEDIA

203

Figure 4 is a rheogram of 60 vol% A-152 alumina in THF and toluene at low shear rates in the presence of 2.5 wt% per alumina of aminopolyisobutylene (APIB);(Fig. 2). AnAPIB molecule consists of an amine terminating group and cyclic dicarbonyl joined to a long alkyl chain. The amine functional groups are expected to behave as a base. The rheogram shows that toluene is a better solvent than THF for dispersing alumina at very high solids loading. In toluene, the alumina surface was shown to be negatively charged. In THF, it was charged weakly positively, indicated by a small deposit of alumina on the negative electrode. This disagrees with explanation that THF is a strong base in the presence of LNA; the surface charge of alumina is primarily affected by the liquid media, as seen in Fig. 3. In the presence of APIB, however, the charge is reversed. This indicates significantly strong interactions between APIB and the alumina surface. Figure 5 is a rheogram of 60 vol% AKP-15 alumina in chloroform and in THF at higher shear rates in the presence of 5 wt% APIB. Aslightly lower viscosity of the slurry with THF than with chloroform is shown. A-152 alumina in methylene chloride was deposited in large amount on the negative electrode under an applied electrical field. This agrees with the acid nature of methylene chloride regardless of the nature of APIB. A smaller deposit of alumina in THF on the negative electrode indicates that the positive potential on the alumina surface in THF is smaller than the oxitive potentialin methylene chloride. Since chloroform is a stronger acid than methylene chloride [ll], the low viscosity of the alumina slurry in THF must be caused by a mechanism other than electrostatic, that is, steric stabilization. The sedimentation heights of alumina sediments, shown in Fig. 6, indicate that a strong base "IF and a strong acid chloroform are equally effective in stabilizing alumina particles with LNA and with an aminosilane, AHAS (see

a

IO'

I 0

3.4 6.8 10.2 13.6 17 20.4 23.8

SHEAR RATE (Vs)

Figure 4 Rheogram of 60 vol% aluminain THF and toluenewith 2.5 wt% APIB per A-152 alumina.

204

LEE

f v)

S E

IO

>

F

I

,O

400 600 800 SHEAR RATE ( m )

200

1000

Figure 5 Rheogram of 60 vol% alumina in THF and chloroform with 5 W% APIB per AKP-15 alumina. 60 -I Y

c

50

c

.P 40 0) I

.-c 3al .-xE 0

30

LNA 0 AHAS

c

20 10

v)

0

EtOH Acetone THF

CHCl3

Hexane

Fig 2). This is shown by the more compact sediments. Alumina with an amphoteric surface was stabilized by dispersants and strongly acid or strongly basic solvents by acombined steric and electrostatic mechanism. Hexane, which is a nonpolar solvent with no hydrogen bonding capability, is shown to be a poor solvent for AHAS, which is a highly polar dispersant. Hexane would introduce a low 5 potential as well; hence the stability had to rely solely on the steric stabilization contribution. When the adsorbate AHAS is not highly soluble in the medium, the dispersant is ineffective as a steric stabilizer. Despite the high dielectric constant of EtOH, it is shown to be an ineffective medium with LNA. EtOH as a weakly acid solvent introduced nearly a zero potential

CERAMIC PARTICLES

205

IN NONAQUEOUS MEDIA

on alumina, although the adsorption isotherm showed a large amount of LNA, =20 mg/g Al203, on the alumina surface. Thus, for the most effective stabi-

lization, a strong electrostatic contribution in addition to the steric contribution becomes essential.

B. Silica The surface of silica is characteristically acidic because of the silanol groups. Figure 7 is a rheogram of 30 vol% silica suspensions in various media. EtOH exhibited the lowest viscosity and EtAc the highest for silica prepared by base hydrolyzing silicon tetraethoxide followed by calcination at 600°C. The potentials measured by electrokinetic sonic amplitude showed very a low negative value of -15 mV for silica in EtOH, as shown in Fig. 8. According to Drago (Table l), EtOH is a Lewis acid and is not expected to ionize the acid silica surface appreciably. The lowest viscosity must then be from a steric contribution by the adsorbed EtOH via hydrogen bonding and/or esterification (silicon ethoxide formation). A weakly basic solvent, EtAc did not adsorb and did not cause the silica surface to be highly charged. Acetone, as a more strongly basic solvent than EtAc, exhibited a lower viscosity than EtOH for silica. As shown in Fig. 8, the negative potential of silica in acetone is relatively high. In the presence of the organic dispersants LNA or AHAS, the viscosity of the 30 vol% silica suspension in acetone (Fig. 9) decreased substantially. This shows that AHAS is the more effective dispersant than LNA for silica in acetone, despite the small potential, less than around -10 mV in Fig. 8. Based on the chemical structure of AHAS given in Fig. 2, AHAS was believed to be a basic dispersant via the amino group, ester group, and ethoxy groups, representing a multifunctional adsorbate. The -OH group may act as an

<

<

<

I

8

IL,

3.4

’

IO 0.34

L

SHEARRATE

(Vs)

Figure 7 Rheogram of 30 vol% silica in four differentsolventswith no dispersant present.

206

LEE

-100

Zeta P 0 Viscosity EtOH Acetone Ace+LNA Ace+AHAS

Figure 8 6 potentialand viscosity of 35 vol%silicasuspensionsinvariousmedia. The viscosity values are at the shear rate of 110 Vs.

1

'lo00 I

YJ

B

L

-

100

"P- W/D Polymer

>

c

v)

0 0

Linolenic

+ AHAS IO

U,

> I

0.34

0

. ,: 3.4

&L

34

SHEAR RATE (Vs)

Figure 9 Rheogram of 30 vol% silica in acetone with 5 wt% LNA and AHAS per silica compared with a rheogram of no dispersant. acid site, however, making the molecule amphoteric. Although strong hydrogen bonding andlor chemisorption of AHAS via Si-0-Et is possible, the molecule must have reversed the charge of the silica surface in acetone. One possible explanation for this is the base-catalyzed enolization of acetone, making the enol form of acetone more acidic. LNA, on the other hand, expected to behave as an acid adsorbate, did not appreciably affect the t; potential. Thus, the lowest viscosity of the silica suspension in the presence of LNA must be from the combined effects of a steric effect by LNA and an electrostatic effect by acetone. For AHAS, this is a more steric contribution than electrostatic. In EtOH, LNA is shown to be a superior dispersant to AHAS (Fig. lo), and in THF (Fig. 1l), AHAS is superior to LNA (Fig. 1l), similarly to the finding in acetone (Fig. 9). LNA in chloroform, shown in Fig. 12, exhibit a poor rhe-

207

CERAMIC PARTICLES IN NONAQUEOUS MEDIA

-

1000

E

W / O Polymer Linolenic AHAS

Y

>.

8

10; I

I

Q34

L

3.4 SHEARRATE

(Vs)

34

Figure 10 Rheogram of 30 vol% silica in ethanol with

5 wt% LNAand AHAS.

1000 1 ?

2

d

100

!z >. v)

0 V

r \

" .C c -

IO

I 0.34

e

: ,: 3.4

Linolenic AHAS

34

SHEAR RATE (Vs)

Figure 11 Rheogram of 30 vol% silica in THF with 5 wt%LNAandAHAS. 1000

'" If

E

100

LNA-CHCIS

"-c

LNA-THF

-c AHAS-CHCIS

Y

---e-- AHAS-THF --e-- APIB-CHCIS

>.

k

8 V

'-c

IO

--b

II

APIB-THF

>

I

0

200

400

600

SHEARRATE

800 (I/sec)

1000

Figure 12 Rheogram of 35 vol% silica in Chloroform and THF with 5 wt% LNA, AHAS, and APIB.

208

LEE

ological behavior. The adsorbent silica is acidic and the solvent and dispersant are basic. This combination is shown to be ineffective for dispersing silica. AHAS, on the other hand, is shown to be effective in both choloroform and THF. This may be because of the amphoteric nature of AHAS, discussed previously. Aminopolyisobutylene is also shown to be effective in both chloroform and THF, although not as low in viscosity as AHAS. APIB was previously mentioned as a basic dispersant for the amino groups and carbonyls. It may be that the basic APIB adsorbs on silica very well and is very soluble in both chloroform and THF. To consider the three-way interactions between dispersant-solvent,solvent-solid surface, and dispersant-solid surface, one must consider the relative acidity or basicity of all the participants. Is chloroform a stronger acid than silica? Is THF a stronger base than APIB? Such questions are relevant and need to be answered. It may be speculated that silica is a stronger acidthan chloroform; hence the basicAPIB preferentially interacts with silica over chloroform. If APIB is a stronger base than THF for silica, then it preferentially interacts with silica over THF. However, the exact relative acidity or basicity of the three interacting participants, that is, APIB, solvent, and silica, is unknown. Hence, it is difficult to assess the degree of the interaction quantitatively.

C.

Portland Cement

It is well known that portland cement is strongly basic because of the large proportion of calcium hydroxide. Figure 13 shows the viscosity of cement slurries in various liquids as a function of Brookfield shear rates. At 24 vol% cement in organic solvents, it showed lower viscosities than that in water. A strongly basic liquid medium, THF, exhibited the lowest viscosity with Newtonian be-

1200

Acetone

UI

a

S

900

t

g 0

Y

600

300

12

30

60

RPM

Figure 13 Rheogram of 24 vol% portland cementin various media.

209

CERAMIC NONAQUEOUS PARTICLES INMEDIA

havior. This is also difficult to assess based simply on the Drago and Fowke's acid-base interaction theory, because the cement surface is basic and the THF medium is also basic. Using the relative acid-base scale in Table 1,THF is acid compared with cement and introduces a larger 1 / with ~ a smaller ionic strength and higher electrostatic contribution than for the cement surface in water to stabilization of the suspension [Eqs. (1) and (3)]. In the presence of LNA, EtOH, which most resembles water chemically, shows the 'higher viscosity in Fig. 14. The cement in all other liquid media exhibitedlowviscosities. In particular, the strong acid chloroform and strong base THF show Newtonian low viscosities. The reasoning given for silica in Fig. 12 must apply here as well, that is, the relative acidity-basicity in the suspension among the three participants. When water is added to the cement slumes in an organic medium, chloroform in the presence of LNA (Fig. 15), the viscosity changed little up to 10 vol%of H20. A water volume above this destroys the stability. Figure 16 shows the relative diametral tensile strength for specimens from cement dispersed in various liquid media. The diametral tensile strengths were determined by applying compressive loadings diametrically on cylindrical specimens 3 x 3 cm in diameter and height formed from 41.4 vol% cement and 17 vol% organic solvent in water. The relative strengths were calculated by dividing the measured values by the strength of the specimen formed from 100% water. The strength of the cement body formed from chloroform-water showed more than a threefold higher strength and a slightly higher strength than EtOH and water. This shows the effectiveness of nonaqueous liquid in forming superior cement components through improved dispersion.

% l_"" \_" S 2000

" " " " " " " " " "

%

3

6

12

30

60

RPM

Figure 14 Rheogram of 24 vol% portlandcementin LNA per cement.

various mediawith 5 wt%

LEE

210

12 RPM

60 RPM

0'

I

5

I

1

15

IO VOLUME 9'0 of H,O

Figure 15 Viscosity of 35 vol%portlandcementwi Ith 5 wt% LNA in Chloroform as a function of water content.

l

CHCI,

1

EtOHw/LNA E

t 0

O

H

m , I

, 2

,

, 3

,

, 4

RELATIVE STRENGTH

Figure 16 Relative diametral tensile strengths of 27 day cured portland cement cast from various liquid media.

111.

HAMAKERCONSTANT

In Deryagin, Landau, and Verwey (DLVO) theory the magnitude of attractive energy is represented by the Hamaker constant A, as shown in Eq. (2). Two like particles in a liquid medium exhibit attraction arising from van der Waals forces. The net Hamaker constant representing the attraction between the two particles in the medium is given by the geometric mean of the two materials in Eq. (4):

CERAMIC PARTICLES

IN NONAQUEOUS MEDIA

21 I

where the subscripts p p = particle-particle and m m = medium-medium. For a small App/m, that is, better dispersion, one wants the value of App to be close to that of A-. The A- values calculated from the surface tension of the liquid [2,5] are listed in Table 1. If the dispersoid particles are covered by another material, the value of App approaches that of the adsorbate material. In this case, for small App/m, one can select a dispersant having a value ofA close to thatof the liquid medium. Unfortunately, accurate determination of the Hamaker constant for a given material system is not yet well established. Since the Hamaker constant is a function of the polarizability of a molecule, LNA, having three double bonds,would have a relatively high value (>S) [5,17]. Plots of the viscosity of alumina suspension as a function of the Hamaker constant of the liquid medium, made by Rives [IS], show that the viscosity minimum at a Hamaker constant = 9 for LNA in acidic solvents and = 8 for LNA in basic solvents. The values of Am in Table 1 show that such solvents as toluene, methylene chloride, chloroform, and THF would be better media for smaller APplm and water the worst. This practice should apply to mixed solvent systems to tailor the Hamaker constant, solubility, electrostatic charge, and so on. This was presented in our recent work elsewhere [19].

IV.

SUMMARY AND CONCLUSIONS

Ceramic particles can be well dispersed in nonaqueous solvent via electrostatic repulsive forces. In the presence of organic dispersants, the electrostatic repulsive forces can be augmented by steric hindrance from the adsorbed dispersant molecules. The combined hindrances to attraction and coagulation can be very effective in dispersing fine particles and increasing solids loading. For alumina powders, THF is found to be a good medium with an acid dispersant LNA. Chloroform, which is an acidic solvent, is found to be an even better solvent for alumina dispersion with LNA. For a basic dispersant, APIB, toluene is found to be good choice of solvent as well as TI-IF.Chloroform is also found to be a good choice. For silica powders, ethanol is a good medium, but in the presence of an aminosilane, AHAS, the more basic solvents acetone and THF are better dispersionmedia.A strong acid Chloroform was also found to be good with AHAS. For a basic dispersant, APTS, both THF and Chloroform are good dispersion media. THF and Chloroform were found to be excellent in obtaining a low viscosity of cement slurries. Cement specimens formed from 16.7% chloroform in water produced a 3.4-fold stronger cement cast. It may be possible to reduce

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the amount of water currently in use in the field to improve the strength of cement concretes. For better dispersion, the relative acid-base interactions of the three participants-liquid medium-adsorbate, adsorbate-solid surface,andsolidsurfaceliquid medium-should be considered simultaneously rather than one pair at a time. To reduce the attractive forces between the particles, tailoring the effective Hamaker constant of the particle surface can be made by choice of dispersant and solvent. Mixed dispersants and/or mixed solvents may be considered for the purpose. Depending on the choice of dispersant, the role of dispersants may be combined with therole of binders to decrease the content of the organic additives.

ACKNOWLEDGMENTS The author acknowledges the contributions of data from his graduate and undergraduate students at Clemson University, C. Calhoun, J. Rives, and U. Paik. The financial support from the National Science Foundation and the State of South Carolina is also gratefully acknowledged.

REFERENCES 1. Shaw, D. J., Introduction to Colloid and Surface Chemistry, 3rd ed., Butterworths, London,1985,p.186. 2. Fowkes, F. M.,Dispersionofceramicpowdersinorganicmedia,in Ceramic Powder Science Advances in Ceramics, Vol. 21 (G.L. Messing, K. S. Mazdiyasni, J. W. McCauley, and R. A. Haber, eds.), Am. Ceram. Soc., Westerville, OH, 1997, p. 41 1. 3. Lee, B. I., and Rives, J. P., Dispersion of alumina in nonaqueous media,Colloids Surfaces, 56, 25 (1991). 4. Napper, D. H.,Polymeric Stabilization of Colloidal Dispersions, Academic Press, NewYork,1993. 5. Johnson,R. E., andMomson,W. H., Jr., Ceramicpowderdispersioninnonaqueous systems, in Ceramic Powder Science, AdvancesinCeramics,Vol.21 (G. L. Messing, K. S. Mazdiyasni, J. W. McCauley, and R. A. Haber, eds.), Am. Ceram. Soc., Westerville, OH, 1987, p. 323. 6. Calvert, P. D., Lalanandham, R. R., Parish, M. V., Fox, J., Lee, H., Pober, R. L., Tormey, E. S., and Bowen, H. K., Dispersion of ceramic particles in organic liquids, in Proc. Mae. Res. Soc. Symp., Vol. 73, (C. J. Brinker, D. E. Clark, and D. R. Ulrich, eds.), Materials Research Society, Pittsburgh, PA, 1986, pp. 579-584. 7. Griffiths, D. A., and Fuerstenau, D. W., Effect of pH and temperature on the heat of immersion of alumina, J. Colloid Interfac. Sci., 80, 271 (1981). 8. Robinson, M., Pask, J. A., and Fuerstenau, D. W., Surface charge of alumina and magnesia in aqueous media, J. Am. Cerarn. Soc., 47, 516 (1964).

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of 9. Lee, B. I., and Hench, L.L., Electrophoretic behavior and surface reactionssolgel derived alumina, Colloids Sufaces, 17, 21 (1987). 10. Kiselev, A. V., and Lygin, V. I., Infiared Spectra of Sugace Compounds, Wiley and Sons, New York, 1975, p. 254. 11. Okuyama, M., Garvey,G., Ring, T. A., and Haggerty, J. S., Dispersion of silicon carbide powders in nonaqueous solvents, J. Am. Ceram. Soc., 72, 1918'(1989). 12. Drago, R. S., Vogel, G. C., and Needham, T. E., A four parameter equation for predicting enthalpies of adduct formation,J. Am. Chem. Soc., 93, 6014 (1971). Rubber 13. Fowkes,F.M.,Acid-basecontributionstopolymer-fillerinteractions, Chem. Technol., 57, 328 (1984). 14. Fowkes, F. M., and Mostafa, M. A., Acid-base interactions in polymer adsorption, Ind. Eng. Chem. Product Res. Dev., 17, 3 (1978). 15. Barton, A. F. M., CRC Handbook of Solubility Parameters and Other Cohesion Parameters, CRC Press, New York, 1983. 16. Weast, R. C. (ed.), CRC Handbook of Chemistry and Physics, 52nd ed., Chemical Rubber Co., Cleveland, OH, 1971. 17. Osmond, D. W. J., andWaite,F.A.,ThetheoreticalbasisforthestericstabiDispersion Polylization of polymer dispersions prepared in organic media, in merization in Organic Media (K.E. J. Barrett, ed.), Wiley and Sons, New York, 1975, pp. 9 4 . 18. Rives, J. P.,Dispersion of alumina powder in nonaqueous media via steric and electrostaticrepulsiveforces, MS. Thesis,ClemsonUniversity,Clemson,SC,

1990. 19. Lee, B. I., and Paik, U., Dispersion of alumina and silica powders in nonaqueous media: Mixed solvent effects, Ceram. h r . , 19, 241 (1993).

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9 Synthesis and Dispersion of Barium Titanate and Related Ceramic Powders Ki Hyun Yoon and Kyung Hwa Jo* Yonsei University Seoul, Korea

1.

INTRODUCTION

To improve the electrical and mechanical properties of advanced electronic ceramics, it is known that the methods of powder preparation and control of the starting materials, that is, grain size and size distribution, have a direct effect on the material properties. There are various kinds of chemical methods for ceramic powder preparation via a liquid phase, which make it easy to control the properties of the powder product and to prepare fine particles. The sol-gel method has some merits, such as highpurity, homogeneity, stoichiometric composition, and infinite flexibility [1,2] using a variety of organic raw materials [3,4]. However, it is necessary to reduce the agglomeration and to control the interaction between particles in suspension. The high reactivity results in shortening the sintering cycle and developing stability against inhomogeneous grain growth. Therefore, research on suspension behavior is necessary, but this can proceed only at a finite pace. The extent to which individual particles exist as aggregates obviously has a major influence on the behavior of the suspension during processing and on the properties and performance of the final product. Therefore, the primary emphases are on the preparation and dispersion of fine powders. The extension of most of the general concepts to high-density *Current affiliation: Daeww Corporation, Ltd., Seoul, Korea

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and homogeneous grain growth need to be assured [5]. The sintered samples, which are homogeneously packed with fine spherical particles, have high density, a reduced sintering cycle, and enhanced stability against abnormal grain growth. It is believed that by preparation of fine powders and manipulation of the dispersion characteristics, a uniformly packed, high green density compact with stable dispersion will result. From controlling the dispersion characteristics, enhancement of sintered density and stability from abnormal grain growth may be expected.

II. PREPARATION OF FINE POWDERS: PREPARATION BY THE SOL-GEL METHOD Whena multicomponent alkoxide is used, the multicomponent or complex alkoxide can be prepared by reacting a combination of single alkoxides or by adding soluble inorganic salts to single or complex alkoxides. The species from an inorganic salt can be incorporated into both multicomponent alkoxides and the alkoxide sol structure itself. If the atom or ion from the salt is to be incorporated into the sol, it must be properly dispersed and then reacted to form the multicomponent oxide. Utilizing an inorganic salt in this manner requires obtaining a stable solution of the alkoxide and the salt. Metal salts are less expansive than alkoxide, so the use of metal salts compensates for a disadvantage of the sol-gel method. Barium titanate is a well-known dielectric material of technological importance for high-technology ceramics and is normally synthesized byconventional solid-state reaction. For the sol-gel-derived powder, the high surface area of the dried gel results in very high reactivity, which permits low-temperature processing, Therefore, the weight loss that is a problem inthe BaTiOg-PbTiO3 system may be reduced during calcination since vaporization of the PbO is expected to be less. By starting with well-mixed solutions or sols, chemical homogeneity on the molecular scale can be obtained. High purity can be maintained because the successive steps common to many conventional ceramic processes can be avoided [6]. The nature of the solvent can likewise play a critical role in the synthesis. The solvents are not inert with respect to the reaction sequence. It is important to choose solvents for the complete solubility of alkoxides or salts. Solvents play a direct role through the process of alcohol exchange. The solvent and raw materials were dissolvedcompletely with vigorous stirring before the hydrolysis reaction at a certain temperature. The solution was refluxed for many hours. Generallythis temperature is near the boiling point of the solvent. The solvent must be dried before use to prevent partial hydrolysis. Because the alkoxide and salts are extremely moisture and carbon dioxide sensitive [7], they may be handled under an N2 atmosphere. The early work of

SYNTHESIS AND DISPERSION OF BaO-Ti02

217

Mazdiyasni [S] utilized barium isopropoxide derived from high-purity Ba metal; much earlier, Flaschen [g] used Ba(OH)2 as a starting material. Ritter et al. [lo] studied the alkoxide-based synthesis of phases in aBaO-Ti02 system. Inorganic water-soluble salts, such as Bach and Ba(N03)2, are considered but not investigated because of the potential problems of crystallization of these reagents during drying .of the gels and contamination of the final electroceramic by the residual anionic species [1 l]. For PbTi02, there are two possible structures of product formed by the reaction of a lead salt and a titanium alkoxide [12]. The hydrolysis is performed for metal alkoxide precusors. It is performed by introducing the alkoxide into excess water under vigorous stirring. The purity of water seems to be very important as well, so hydrolysis is performed generally with deionized water. During hydrolysis, the input method of water is steam or slow dripping. The total amount of water for hydrolysis has an effect on particle size [13]. The rate of hydrolysis depends on the nature of both the metal ions and the alkoxide groups. The hydrolysis rate generally decreases with increasing length of the n-alkyl chain and with increasing branching of the alkyl group, presumably for steric reasons [14]. It is also expected that the hydrolysis rate should increase as the metal ion is more electropositive. In the sol-gel method, the amount of acid usedfor peptization influences the shape and size of the particle [3]. Acid and base additions are generally specified in terms of pH. In this case, however, it is observed that the type of acid plays a much more important role than the pH of the system. The anions of the acid must be noncomplexing or very weakly complexing with metal ions at these low catalytic concentrations. The amount of acid in relation to the metal must not be large enough to prevent the formation of a continuous metal bonding through oxygen. For example, acrylic acid used for peptization in the (Ba, Pb)TiOs system has an ionization constant [4] of 5.6 x 10-5 at 25°C (generally others have ionization constants below 10-4). Thus, the acrylic acid satisfied the general requirement for an acid for peptization: it must not be too large to prevent the formation of a continuous bonding of metalions of the alkoxide throughhydroxides. After hydrolysis and peptization, gel is formed. The last step, drying, plays a very important role in maintaining a fine powder and has direct effects on the calcination step. The relation between the drying method and the surface area of powders is important. There are many drying methods, such as air, vacuum, and freeze drying. This point is illustrated with the examples of (Ba0.2Pbo.s)Ti03. In powder preparation, the surface reactivity, that is, the specific surface area, plays a very important role. Specific surface areas are often determined by the BET (Brunauer, Eminett, and Teller) method. Typical BET plots of the sol-gel, CM0 (calcined mixed oxide method), and MSS (molten salt synthesis method) [l51 derived powders are shown in Fig. 1.

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Figure 1 BET plotfor(Bao.2Pbo.8)TiO3: (triangles) sol-gelderived (freeze dried); (squares) sol-gel derived (vacuum dried); (filled circles) CM0 derived (air dried); (open circles) MSS derived (air dried).

Table 1 shows that the specific surface area is dependent on the synthesis method for the powders. In general, sol-gel-derived (Bao.2Pbo.8)Ti03 powders have specific surface areas greater than those of CMO-derived powders. Generally, the smaller the particle size, the greater is the specific surface area. Thus, it can be expected that the particle size of sol-gel-derived powders is smaller than that of CMO- and MSS-derived powders. When sol-gel-derived powders are dried,the surface areas of air-dried,vacuum-dried,and freezedried powders are 36.4,51.9, and 54.6 m2/g respectively. Freeze-dried powders have the highest value. These results are similar to the value for vacuum-dried powders prepared only from alkoxides by Mazdiyasni et al. [l61 Freeze-dried powders, after heat treating at 600"C, exhibited a surface area value of 19.9 m2/g, which is higher than those obtained by, Mazdiyasni [161 and Rehspringer and Bernier [17], 15 and 16.8 m2/g, respectively, both preparedusing only alkoxide. Metallurgical micrographs for sol-gel-derived (Bao.2Pbo.8)Ti03 prepared by air drying (Fig. 2) reveal that the air-dried powder has the strongest

219

SYNTHESIS AND DISPERSION OF BaO-Ti02 Table 1 BET Values(SurfaceArea) for (Bao.flb0.8jTiO3 Sol-gel CM0

MSS

powder Calcined Dried gel

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Synthesis method

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dry

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(17)

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50.0 6.2-16.8 15.019.9

agglomerate and the freeze-dried powder has the weakest agglomerate. The results agree with the BET values shown in Table 1. The freeze-drying method may be divided into freezing, sublimation, and desorption steps. The sublimation step is affected by the equilibrium condition of composition versus temperature and pressure. The advantages of this method are that it is easy to control particle size, composition, and input of minor components in the systems homogeneously, low-temperature processing is possible, and evaporation of volatile elements is prevented . The disadvantages are the need for a cold trap at very low temperatures and a vacuum system, so this method is difficult to apply to mass production. The resulting monosized and submicrometer particles are easier to process into uniform green microstructures, which results in easier control of the microstructure during densification. This means uniformityof the particle size

Figure 2 Metallurgicalmicroscopephotographs for sol-gel-derived (Bao.flbf~.8) TiO3: (a) air drying, (b) freeze drying, (c) vacuum drying.

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and distribution of the voids, which can be measured by a particle size or pore size analyzer. Figure 3 shows the BaTiO3-PbTiO3, particle size distributions as a function of synthesis method. In Fig. 3, the sol-gelderived powders have the narrowest particle size distribution and smallest mean diameter. The particle size and distribution can also be observed with scanning and transmission electron microscopy (SEM and TEM). X-ray diffraction of (Ba0.2Pbo.8)Ti03 in Fig. 4 shows that the powder is initially amorphous. Samples were calcined at various temperatures, and a progressive reduction in peak width with increasing calcination temperature can beseen. The x-ray diffraction patterns of solgelderived andCMO-derived(Ba0.2Pbo.8)Ti03 are shown inFig. 5. For CMO-derived powders, the characteristic peaks of unreacted BaC03 are found at 600"C,andthey decrease with increasing the treatment temperature to 900°C. At 9OO"C, the powder is completely a single phase. The degree of crystallinity appears relatively good even at 600°C in the solgel method, indicating that this method has a synthesis temperature 200°C lower than that of the CM0 method. Tine infrared absorption spectra have been obtained for powder samples dispersed in pressed KBr disks to identify the molecular structure of their components, specifically organic materials. Figure

Particle

diameter [pm)

Figure 3 Particle size distributions of ( B ~ I - J . ~ P ~ o .prepared ~ ) T ~ Oby ~ differentmethods. (a) Sol-gel, (b) CMO,(c) MSS.

SYNTHESIS AND DISPERSION OF BaO-Ti02

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2 8

Figure 4 X-ray diffraction patterns of sol-gel-derived (Ba0.2Pb0.8)Ti03:(a) 60O0C, (b) 500°C, (c) 400"C, (d) 200°C.(e) 100"C, (0 dried gel.

6 shows the infrared (IR) absorption spectra of (Bao.2Pbo.8)Ti03 heated to various temperatures. The absorption bands around 3450 cm-1 (peak 1) are caused by the stretching vibration of -0-H bonds. With increasing heat treatment, a single C bond diminishes at lo00 cm-1 in size and a double C bond at 1650 cm-l.Peaks2and 3 shift to lower wave numbers, and the intensity of the bands decreases with increasing temperature, suggesting gradual evaporation of the residual organic compounds. The absorption band (peak 4) corresponds to the inorganic metal characteristic peak in the sol-gel-derived BaTi03-PbTi03 solid solution. The broad absorption band at 530 cm-1 [l81 corresponds to the characteristic peaks of BaTiOs. Lead titanate possesses two characteristic absorption bands at 580 cm and 400 cm-1 [19]. Peak 4 is broad and appears to include a PbTiOs peak in the region; there is a large component of Pb in the composition of (Ba0.2Pbo.8)Ti03. The breadth of peak 4 narrows with increasing heat treatment temperature.The sharper peak indicates strengthening of the bonds[20,21]. In the same frequencyrange, the sharp peaksuggestsan increase in crystallization. Figure 7 shows infrared absorption spectra compared withCMO-derived,MSS-derived,and sol-gel-derived (Ba0.2Pbo.8)Ti03. For

YOON AND JO

222

70

60

50

40

30

20

Figure 5 X-ray diffraction patterns of (Ba0.2Pl~o.g)Ti03:(a) 600°C, 1 h (sol-gel derived), @) 900°C. 2 h (CM0 derived),(c)600°C, l h (CM0derived); * (circles) BaC03.

CMO-derived powders heated to 900°C, for MSS-derived powders heated to 8OO0C, and for sol-gel-derived powderheated to 600°C, the characteristic peaks of the solid solutions indicate complete transition to the final oxide. The IR results show a tendency similar to that of the x-ray diffraction patterns mentioned before. The formation of a fully crystalline product appears to be essentially complete at a relatively lowtemperature. Higher temperatures promote grain growth, resulting in an increase in average crystallite size but with no further crystallization. One can also observe a degree of grain growth by SEM and

SYNTHESIS AND DISPERSION OF BaO-Ti02

223

Wavenumber (*102cm"

Figure 7 Infraredtransmissionspectra of (Ba1).2Pb&8)Ti03afterheating(a) 90OoC, 2 h (CM0derived), (b) 800"C, 6 h (MSS derived),(c) 600"C, 1 h (sol-gel derived). (From Ref. 7.)

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(a )

Figure 8 SEM photographs of (Bm.flbo.8)Ti03 prepared by variousmethodsand sintered at 1200°C (bar = 1 p): (a) CM0 derived, (b) sol-gel derived (nondispersed), (c) sol-gel derived (dispersed with ;2% PMMA).

"EM. Homogeneous and fine particles lead to a good sintered body. In a good sintered body, grain grows homogeneously. Figure 8 shows SEM photographs of BaTi03-PbTi03 solid solution from various synthesis methods. In the solgel-derived specimens, the grain size is smaller compared with CMO-derived specimens. However, in the well-dispersed system, which is explained in the next section, the grain grows homogeneously, resulting in high density.

111.

DISPERSIONCHARACTERISTICS OF FINEPOWDERS

A.

Dispersion Characteristics at Various pH Values

The electrophoretic mobility is proportional to the 6 potential and is thus a qualitative measure of colloid stability. Therefore, to find how to prevent aggregation and accelerate dispersion, the variation in suspension behavior with pH, which determines the surface charge, and the effect of dispersion on the sintered specimens are the subjects of this study. l, potential measurements can provide important guidance on the preparation of colloidal dispersions. A plot of 6 potential against pH for the sol-gel-derived (Bao.2, Pbo.8)Ti03 gel is given in Fig. 9 [22]. The point of zero charge (PZC) is found to occur at (Ba2+, Pb2+, and T i 4 ion concentrations of 2.0 x 10-7, 5.01 x 10-7, and 1.0 x l0-8M, respectively, or &, = - log [Ba2+] = 6.7, ppb = - log [Pb2+] = 6.3, and h i = -log [Ti&] = 6.0 [23]. A point of zero charge (6 = 0) with positive and negative charge branches is shown in Fig. 9. The 6 potential simply decreases to zero with an increase in pH to 7 and becomes increasingly negative with further increases in pH. Thermogravimetric analysis of (Ba0.2Pbo.8)Ti03 powders is shown in Fig. 10.The drastic weight loss is attributed to dehydration and the

SYNTHESIS AND DISPERSION OF BaO-Ti02

Figure 9

225

c potential as a function of pH for sol-gel-derived (Bao.zPbo.s)Ti03.

loss of volatile organic solvent residue from the solution at high temperature for pH 11and 13. At the isoelectric point, pH 7, the weight loss is small. Mazdiyasni [24] studied the preparation of fine particles and their applications to perovskite materials. In this report, a well-dispersed colloidal system with good reactivityaccelerated the removalof organic residue from the solution and maintained the deagglomerative conditions. Extensive reviews of the charge potential behavior of colloidal systems have been given by Hunter [25]. The most important dissociable groups are of the strong acid (sulfate or sulfonate), weak acid (sulfite or carboxyl), weak base (amine), and strong base (quaternary ammonium) types, and they may occur alone at the surface or in various combinations. These combinations form zwitterionic surfaces, and then the driving force removed from the system is replaced by repulsive and interactive forces. The specific surface areas [22] of sol-gel-derived (Bm.2Pb0.8)Ti03powders are 5 m2/g for pH 7 and 98-100 m2/g for pH 11and 13. That pH 7 is found to be highly agglomerated, and those at pH 11and pH 13 indicate good dis-

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Figure 10 Thermogravimetry function aas of pH for sol-gel-derived (Ba0.2Pbo.g)Ti03:(solidline) pH 1; (dash-dotline) pH 3; (dash-doubledot) pH 5; (dashed) pH 7;(circles) pH 9; (squares) pH 11; (triangles) pH 13. (From Ref. 22.)

persion. Metallurgical microscope photographs of the sol-gel-derived (Bao.2Pbo.8)Ti03at various pH are shown in Fig. 11. The results correlate with the weight loss and BET values. For all operations that involve suspensions of powders in liquids, the processing behavior is dominated by the rheological properties of the suspension. Suspension rheology is determined largely by the agglomeration of particles within the liquid. Most suspensions of ceramic interest are shear thinning, and their properties depend upon plastic viscosity. Plastic viscosity depends on the state of agglomeration. Figure 12 shows a plot of viscosity as a function of shear rate and pH for sol-gel-derived (Bm.2Pb0.8)Ti03 powders. In the samples at pH 3-9, the shear-thinning phenomenon is apparent, demonstrating that

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Figure 11 Metallurgical microscope photographs of sol-gelderived (Bao.2Pb0.8)Ti03:(a)pH 1, (b) pH 3, (c) pH 5, (dl pH 7, (e) pH 9, (0 pH 11, (g) pH 13. (From Ref. 22.)

r

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0.0

Y

=

0.02 0.01

0

10

20

30 40 Shear Rate (sec" )

50

60

Figure 12 Suspension viscosity plotted against shear rate for sol-gel-derived 1; (filledcircles) pH 3; (open triangles)pH 5; (Bao.$b0.8)Ti03:(opencircles)pH (filled triangles) pH 7 ; (open squares) pH 9; (filled squares) pH 11; (exes) pH 13. (From Ref. 22.)

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there is a weak agglomeration in the suspension. At higher pH or greater negative potential value, the viscosity decreases with increasing shear rate. One of the most obvious effects is the creation of large voids caused by poor packing around aggregates and huge voids caused by the bridging of aggregates. Such voids, much larger than the surrounding grains, cannot be removed during sintering: there is no driving force for elimination of such oversized pores. A more subtle effect occurs when the packing density is not uniform, having denser regions and more porous regions. This leads to locally inhomogeneous sintering and microstructures with completely dense patches and more porous areas. In the dense regions, there is no porosity to impede grain boundary migration, and grain growth is rapid. Sintering kinetics is affected by the state of agglomeration as well. X-ray diffraction patterns of the calcined powders and the sintered specimens of the sol-gel-derived (Ba0.2Pbo.8)Ti03 are shown in Fig. 13. In the pH 1 pattern, not only the characteristic peaks of the solid solution but also second-phase peaks are found from the HCl used to adjust the dispersion characteristics. For the pH 13 specimen, however, only the characteristic peaks of the solid solution appear. At low pH, the degree of crystallinity is reduced and the powder is not synthesized completely. Figure 14 shows infrared absorption spectra of sol-gel4erived (Bao.2Pbo.8)Ti03 calcined powder for various dispersives in pH. The intensity of the hydrogen bond and the characteristic bond of the organic compound in the dried gel decrease as the pH is increased. Peaks 2 and 3 become sharper with an increase in pH and shift to higher frequencies as the crystalline phase becomes more stable. Figure 15 shows the apparent and relative density for

c

70

20

70

2 8 (deg)

20

Figure 13 X-ray diffractionpatterns of sol-gelderived (Bao.2Pbo.s)TiO3dispersed in (a) pH 1, (b) pH 7, (c) pH 13.

4000 3OOO 1500 2000

1000

Wavenumm (cm”)

500 400

Figure 14 Infraredtransmission of sol-gelderived and [email protected])Ti03: (a) pH 1, (b) pH 7 , (c) pH 13.

Figure 15 Apparentand relativedensityplottedagainstpHforsol-gel-derived (Bao.2Pbo.8)Ti03 sintered at 1200°C for 0.5 h.

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230

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I

Figure 16 SEM photographs of(Ba0.2Pbo.g)Ti03sinteredat pH 1, (b) pH 7, (c) pH 13. (Bar = 1 pm.)

1200°C for 0.5 h: (a)

(Ba0.2Pbo.8)Ti03 sintered at 1200°C for 0.5 h as a function of pH. At pH 7, where the dispersion is poor, the density is the lower than at all other pH levels. The effect on the sintering mechanism of agglomerate size and particle size inyttria-stabilized zirconia hasbeenstudied by Rhodes [26]. According to Rhodes’ report, an agglomerate with a very open arrangement of crystallites may develop large pores or pore clusters and leave large lenticular voids that are difficult to close. These results decrease the density. Reeve [27] predicted a similar phenomenon in BeO. At pH 11 and 13, with good dispersion characteristics, the sintered density is high. SEM photographs of the sol-gel-derived (Ba0.2Pb0.8)Ti03as a function of pH are shown in Fig. 16. At pH 13, the grain size is homogeneous; aggregation appears and the sinterability decreases at pH 7.

B. Dispersion Characteristics as a Function of Polymethyl Methacrylate (PMMA) Solvents should have a solubility parameter close to that of the organic polymer chain to maximize extension of the attached chain into the liquid. The extent of polymer adsorption from a solution is determined by the balance of three interactions: polymer-solvent, polymer-adsorbent, and solvent-adsorbent. In good solvents (i.e., solvents in which polymer-solvent contacts are energetically favored over polymer-polymer and solvent-solvent contacts, called 0 solvents), the polymer chains on one particle repel the polymer on another. In contrast, a poor solvent is such that segment-segment contacts are energetically favored, leading to association of the polymer molecules. Since osmotic pressure is a colligative property, intermolecular association leads to fewer osmot-

231

SYNTHESIS AND DISPERSION OF BaO-Ti02

ically active species in solution, and this leads to a negative deviation from ideality. The polymer is tied to the surface at a number of points, but for some of its length it is able to extend into the solution. Segments attached to the surface form trains, which are separated by loops: theends of the polymer are usually able to extend into the solutions as tails. Complete dissolution of the polymer allows the outer segments to extend into the solution and the inner segments to become attached to the particle surface. Polymers can be good dispersants only in suitable solvents. In BaTiOsand related material systems, PMMA is a popular dispersant [28]. PMMA is dissolved completely when nhexane and benzene are used as the dispersing medium for this system [28]. The results, illustrated in Fig. 17 in terms of sedimentation density, show an optimum dispersion in the 70% benzene and 30% hexane mixture. Above 40% n-hexane, the 8 point is reached and the sedimentation density decreases drastically. The 8 point is as follows: with gas molecules, polymer chains can show either positive or negative deviations from ideal behavior. It follows that a temperature can be found at which the solvency leads to ideal osmotic pressure behavior, at least up to a polymer concentration of a few percent. Under 8 solvency conditions, polymer molecules ofamillion molecular weight behave as if they are ideal small molecules. At this temperature, the effects of attractive and repulsive interactions balance one another. This implies

F :

8 c

50

e

401

100 0

I

90

10

I

I

Benzene ( % 1 Hexane (%)

I

60 40

I

50 50

I

Figure 17 Sedimentation of (Bao.zPbo.8)TiO3powderdispersedinsolutions of various PMMA contents in benzene-hexane mixtures: (a)0%.(b) OS%, (c) 1.O%, (d) 1 S%, (e) 2.0%, (f) 3.0%.

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YOON AND JO

that the polymer chains can telescope one another without change in the Gibbs free energy. Thus, the 8 point represents the transitional point with respect to segmentsolvent interactions. At this point, the polymer segments change from exhibiting a net mutual repulsion to a net mutual attraction. A fuller description of this change would involve consideration of segment-segment, segment-solvent, and solvent-solvent interactions. The adsorption of polymer molecules at aninterface, which determines the segment density distribution function, is critically dependent upon many factors, which include [29] (1) the chemical constitution of the polymer, (2) the chemical nature and geometrical shape of the interface, (3) the chemical composition of the solvent, (4) the mode of attachment of the polymer chains to the surface, and (5) the surface density of the polymer molecules at the interface. It is the subtle interaction of these diverse factors that determines the conformation of a polymer at an interface and thus the steric interactionina particular system. The effect ofPMMA concentration on (Ba0.2Pbo.8)Ti03 sedimentation density in 70% benzene and 30% hexane is shown in Fig. 18. The results show the best dispersion at about 2% PMMA concentration, with a slight decrease in packing at higher concentrations, which can also affect weight loss during firing. Dehydration and removal of the organic residue increased with increasing the concentration of polymer to a critical concentration, above which these decreased. The weight loss is not proportional to the concentration of polymer as a dispersant. Poly(viny1 alcohol) in aqueous solution takes up much space, which results in brisk dehydration. This can be confirmed by Fig. 19. Optical micrographs for sol-gel-derived (Ba0.2Pbo.8)Ti03powders dispersed in 70% benzene and 30% hexane mixture as a function of PMMA concentration are shown here. Optimum dispersion ap-

r

0

l 2 PMMA concentration( %

3

Figure 18 Effect of PMMA content on (Ba0.2Pbo.8)Ti03 sedimentationin 70% benzene and 30% hexane mixture.

SYNTHESIS AND DISPERSION OF BaO-Ti02

233

Figure 19 Metallurgicalmicrographsfor sol-gelderived (Bac~Pbr~)Ti03 powder dispersed in 70% benzene and 30% hexane as a function of PMMA content: (a) 0%, (b) OS%, (c) 1.0%, (d) 1.5%, (e) 2.0%, (0 3.0%.

pears at about 2% PMMA concentration. This result supports the preceding theory. There are numerous ways of determining the concentration of the polymer in the adsorptionmedium. These include (1) gravimetric analysis by direct weighing of the polymer after vaporization of the solvent, (2) spectrophotometric analysis using the visible, infrared, and ultraviolet regions of the spectra, (3) the refractometric method, (4) viscometric methods, (5) densometric methods, (6) titrimetric analysis, (7) use of radioisotopes, and (8) measurements of the change in other physical constants, such as turbidity, and dipole moment. Viscosity measurement is used to determine the quality of dispersion in suspensions. A well-dispersed system usually shows a Newtonian viscosity with no pseudoplasticity. The viscosity of such a system is lower than that of poorly dispersed systems. A plot of viscosity against dispersant concentration typically shows a rapid decrease in viscosity and then a leveling at a concentration required for monolayer coverage of dispersant on particle surfaces. At higher dispersant concentrations the viscosity either rises again or remains low. The degree of agglomeration in dilute particle suspensions can also be found

YOON AND JO

234

by particle size measurements using a photon correlator; agglomeration leads to an increase in the average apparent particle size. Figure 20 represents plots of suspension viscosity against shear rate for solgel-derived (Ba0.2Pb0.8)Ti03 as a function of PMMA concentration. The viscosity decreased with increasing shear rate in the 0% PMMA solution. In this chapter the suspension with no polymer addition shows extensive low shear rate aggregation, as indicated by the high shear-thinning behavior. The viscosity becomes independent of shear rate, that is, Newtonian behavior observed with an increase in the polymer. For the interaction between polymers and a mica sheet, explained in detail by Klein [30], the viscosity of the suspension with polymer washigher than that of the suspension without polymer. The reason for this is the high molecular weight of the polymer itself. Molecular packing and sedimentation density play a crucial role in determining the characteristic of the green and sintered specimens. The high packing density satisfies the conditions for a pore-free and dense, sintered body. Figure 21 shows the apparent and relative densities of sol-gel-derived (Bao.2Pbo.8)Ti03 sintered at 1200°C for 0.5 h after dispersion in PMMA solution. The sintered density increases with increasing the PMMA concentration to 2% PMMA, above which it decreases. The highest sedimentation density at 23% concentration is char-

Y

Shear Rate (sec")

Figure 20 Suspensionviscosityplottedagainst shear rate for sol-gelderived (Bao.zPbo.8)Ti03 as a functionof PMMA content: (a) 0%,(b) OS%, (c) 1.0%, (d) 1.5%, (e) 2.0%, (f) 3.0%. (From Ref. 28.)

235

SYNTHESIS AND DISPERSION OF BaO-Ti02

I

0

1

0

0

l 2 3 PMMA concentration (% )

Figure 21 Apparentandrelativedensityplottedagainst gel-derived (Bao.2Pbo.8)Ti03 sintered at 1200°C for 0.5 h.

PMMA content for sol-

acteristic of a well-dispersed suspension. Addition of a polymer enhances the strength, flexibility, and workability of ceramics in its green state, before sintering. It might be preferable for the polymer to function not only as a dispersant but also as a binder. The effect of a dispersant on green and sintered specimens has been studied by Calvert et al. [31]. The agglomeration causes inhomogeneous grain growth and void formation. The residual pores and agglomerates in the sintered body act as microflawsthatadversely affect mechanical and dielectrical properties. InFig. 22 the dielectric constant of (Ba0.2Pbo.8)Ti03sintered at 1200°C for 0.5 h is given as a function of PMMA concentration. The packing density increases with increasing PMMA concentration [28]. A good packing density results in high sintered density and high dielectric constant. However, at higher PMMA concentrations (>2%), poor dispersion behavior causes decreasing density and results in a decrease in the dielectric constant.

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r

Temperature (OC)

Figure 22 Dielectric constantplottedagainsttemperatureforsol-gelaerived (Bao.2Pbo.g)Ti03sinteredat 1200°C for 0.5 h withaPMMAcontentof OS%, (c) 1.0%, (d) 1.5%, (e) 2.0%, (f) 3.0%.

(a) 0%, (b)

REFERENCES 1. Sakka, S., Sol-gel synthesis of glasses: Present and future, Am. Cerum. Bull. 6 4 , 1463 (1985). 2. Segal, D. L.,Sol-gel processing: Routes to oxide ceramics using colloidal dispersionsofhydrogenoxidesandalkoxideintermediates, J. Non-Cryst. Solids, 63, 183 (1984). 3. Dislich, H., and Him, P., History and principles of the sol-gel process, and some new multicomponent oxide coatings, J. Non-Cryst. Solids, 48, 11(1983). 4. Jo, K. H., Yoon, K. H., Preparation of sol-gel derived (Bao.zPbo.g)Ti03 powders, Muter. Res. Bull., 24, 1 (1989). 5. Ueyama, T., Wada, H., and Kanako, N., Pulverization and dispersion technique J. Am Cerum Soc. foragglomeratedparticlesofaluminapowderinaslurry, Commun., 71, C-74 (1988). 6. Afsten, N. J., Sol-gel derived transparent IR-reflecting IT0 semiconductor coatings and future applications, J. Non-Cryst. Solids, 63, 243 (1984).

SYNTHESIS AND DISPERSION OF BaO-Ti02

237

7. Jo, K. H., and Yoon, K. H., Characteristics of the (Bal-xPh)Ti03 powders prepared by various synthesis methods, J. Kor. Cerum. Soc., 27, 127 (1990). 8. Mazdiyasni, K. S., Fine praticle perovskite processing,Am. Cerum. Soc. Bull., 63, 591 (1984). 9. Flaschen, S. S., Preparation of BaTiOs by chemical methods, J. Am. Cerum. Soc., 77, 6194 (1955). 10. Ritter, J. J.. Roth, R. S., and Blendell, J. E., Alkoxide precursor synthesis and characterizationofphaseinthebarium-titaniumoxidesystem, J. Am.Cerum. Soc., 69, 155 (1986). 11. Phule,D.P.,andRisbud, S. H.,Sol-gelsynthesisofbariumtitanatepowders Adv.Cerum. Mat., 3, 183 usingbariumacetateandtitanium(1V)isopropoxide, (1988). 12. Grukovich, S. R.,andBlum, J. B., UltrastructureProcessing of Ceramics, Glasses and Composites, John Wiley and Sons, New York, 1984, p. 152. K., Formation,packing,andsinteringof 13. Baninger,E.A.,andBowen,H. monodisperse Ti02 powders, J. Am. Cerum. Soc., 65, C-l99 (1982). Metal Alkoxide. Academic 14. Bradley,D. C., Mehrotra,R.C.,andGaur,D.P., Press, NewYork, 1970, p. 55. 15. Yoon, K. H., Oh, K. Y., and Yoon, S. O., Influence of synthesis methods of the FTCR effect in semiconducting BaTi03, Mat. Res. Bull., 21, 1429 (1986). 16. Mazdiyasni, K. S., Dolloff, R. T., and Smith, J. S., III, Preparation of high purity submicron barium titanate powders, J. Am. Cerum. Soc., 52, 523 (1969). 17. Rehspringer,J. L., and Bemier, J. C., Mat. Res. Soc. Symp. Proc., Pittsburgh, PA., 1986, p. 67. 18. Last, J. T., Infrared-absorption studies on barium titanate and related materials, Phys. Rev., 105, 1740 (1957). 19. Congshen, Z., Lisong, H., Fuxi, G., and Zhonghong, J., Low temperature synthe-

sis of ZrO2-TiO2-SiO2 glasses from Zr(N03)45H20, Si(OC2H5)4 and Ti(OC4H9)4 by the sol-gel method, J. Non-Cryst. Solids, 63, 106 (1984). 20. Gottardi,V.,Gulielini, M., Bertoluzza,A.,Fagnano,C.,andMorelli,M.A., J. Non-Cryst. Ramanandinfraredspectraonsilicagelevolvingtowardglass, Solids, 48, 117 (1965). Mat. Sci. Eng., 44, l(1980). 21. Bowen, H. K., Iron diffusion in iron-aluminate spinel, 22. Jo, J. H.,Kim,E. S., and Yoon, K. H., Dispersion characteristics of sol-gel derived (Bao.flb0.8)Ti03 at various pH, J. Mat. Sci., 25, 880 (1990). 23. Lauf, R. J., and Bond, W. D., Fabricationof high field zinc oxide varistor by solgel processing. Am. Cerum. Soc. Bull., 63, 278 (1984). on measure24. Mazdiyasni, K. S., Effects of dispersion of barium titanate powder ments in an electrical-sensing-zone particle size counter, Am. Cerum. Soc. Bull., 57, 448 (1978). 25. Hunter, R. J., Zetapotential in Colloid Science, Academic Press, New York, 1981, p. 134. on sintering yittria-stabi26. Rhodes, W. H.,Agglomerateandparticlesizeeffects lized zirconia, J. Am. Cerum. Soc., 6 4 , 19 (1981). 27. Reeve, R. D., Non-uniform shrinkage in sintering,Am. Cerum. Soc. Bull., 42,452 (1963).

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28. Jo, K. H., and Yoon, K. H., Dispersion characteristics of sol-gel derived (Bao.2pbo.8)Ti03 as a function of PMMA, J. Mat. Sci., 26, 809 (1991). 29. Napper,D. H.,Polymeric Stabilimtion of Colloid Dispersions, Academic Press, NewYork,1983, p. 197. 30. Klein, J., Forces between mica surfaces bearing layer of absorbed polystyrene in cyclohexane, Nature, 288, 248 (1980). 31. Calvert, P., Lanardham, R., Parish, M., and Tormey, E.,Dispersants in Ceramics, JohnWileyand Sons, New York, 1984, p. 249.

10 Rheology and Mixing of Ceramic Mixtures Used in Plastic Molding Beebhas C. Mutsuddy Michigan Technological University Houghton, Michigan

1.

INTRODUCTION

The plastic forming of ceramic shapes is based on the application of external forces to a mixture of ceramic powder and binder. The plastic mixture deforms and flows under applied stresses. The external forces cause the plastic mix to be adjusted to any die or mold, which dictates the eventual shape. Therefore, the flow behavior of such mixes during plastic forming has a major effect on the quality of the ceramic parts. Again, the flow stability of the mix during forming depends on the homogeneity of the ceramic powder and binder mixture. The primary focus of this chapter is to describe the factors that influence flow behavior and mix homogeneity of a plastic mixture containing a high volume fraction of ceramic powder as a filler and a relatively small volume fraction of organic polymeric binder. The emphasis is placed on mixes in which flowability is achieved only above the ambient temperature.

II. RHEOLOGY The flow of a plastic mass through a die or into a mold is dictated by its rheological behavior. Rheology describes the deformation and flow of a material under the influence of applied stresses. The rheological response of a fluid is generally expressed as viscosity. The viscosity of a fluid is a measure of the

239

240

MATSUDDY

internal resistance offered to the relative motion of different parts of the fluid. Viscosity is defined as Newtonian when the shearing force per unit area z between two parallel planes of fluid in relativemotion is proportional to the velocity gradient dvldx between the planes; in other words, dv

z=q-

dx

where:

q = coefficient of viscosity z = shear stress = forcelarea = Nlm2 = Pa dvldx = shear rate y = m/s/m = S-1 q = dynamic viscosity = zly, (Pa=s) For many fluids, especially if highly concentrated and/or if the particles are asymmetric, deviations from Newtonian flow are observed.The main causes of non-Newtonian flow are the formation of a structure throughout the fluid and the orientation of asymmetric particles caused by the velocity gradient. Non-Newtonian fluids may be classified (shown in Fig. 1) as follows: Pseudoplastic, in which the shear stress depends on the shear rate alone. Power law, in which the shear stress is not a linear but an exponential function of shear rate. The rheological expression for a power-law fluid is T = Ky"

(2)

where n is the power-law index and K is known as the coefficient of viscosity. Here, we have three different situations: (1) if n = 1, or unity, the fluid is characterized as Newtonian; (2) if n < 1, less than unity, the fluid is considered pseudoplastic; and (3) the fluid is characterized as dilatant if n > 1. This lastcharacteristic is sometimes encountered with filled systems when the filler concentration is very high, so that the corresponding scarcity of a continuous phase makes it difficult for the dispersed particles to slide over one another. The lack of adequate lubrication by the continuous phase produces frictional voids and sets up structures among the particles under constraint. These structures become increasingly resistant to deformation as the applied shear rate is increased: the viscosity increases with increasing shear rate. Fluids in which no deformation occurs until a certain threshold shear stress is applied, in which upon the shear stress z becomes a linear function of shear rate y. The characteristics of the function are the slope (viscosity) and the shear stress intercept (yield value) zy The rheological expression for this type of material, known as a Bingham solid, is 2"Zy

=qy

(3)

PLASTIC MOLDING RHEOLOGY AND MIXING

241

ORatenl with Yield Pdnt

t

Q) Q)

2 ti

Pseudoplastic with

Yield Point

Shear Rate

-

Figure 1 Classification of non-Newtonianfluids.(After J. S. Reed, Introduction to the Principles of Ceramic Processing, Wiley-Interscience, New York.) Only when the applied stress is greater than the yield stress do viscous flow and shear deformation become manifest. Fluids that show a decrease in shear stress (hence in apparent viscosity) with time are termed thixotropic and are often observed with ceramic slurries. Rheologists have long believed that all fluids are viscoelastic in behavior. A s a result, the deformation of any fluid from the imposition of a stress is the

sum of an elastic deformation, which is recoverable, and viscous flow, which is not recoverable. For fluids of low viscosity at moderate rates of shear, the elastic recovery is extremely rapid and the relaxation time is extremely short. A s a result, the elastic portion of the deformation is too small to measure, and the fluid is considered simply viscous. When viscoelastic fluids are stressed, some of the energy involved is stored elastically, various parts of the system being deformed into new nonequilibrium positions relative to one another. The remainder is dissipated as heat, various parts of the system flowing into new equilibrium positions relative to one another. The elastic nature ofa fluid is characterized by dynamic mechanical or stress relaxation techniques. Dynamic mechanical (oscillatory) testing is a procedure inwhicha sample is sinusoidally strained and the resultant stress is measured. The she& stress z varies with the same frequency as the shear rate

MATSUDDY

242

y but not in phase with it, the phase angle between z and y being equal to n/2 for anon-Newtonianfluid. Because the normal stress differences are even functions of the shear rate, they vary at twice the applied frequency; they also suffer different phase shifts. Measuring the amplitudes of the stresses and their phase shifts with respect to y is a standard way of studying viscoelastic effects in non-Newtonian fluids. These measurements do not change with time unless the fluid exhibits thixotropy. In measuring viscoelastic effects with an oscillating plate method, the percentage of strain must be selected carefully, because the strain is proportional to the total angle at which the plate is turning. The larger the angle through which the plate oscillates, the greater is the strain. If the angle is too large, the fluid being measured may undergo internal cracking or segregation.

A.

Viscosity of HighlyFilledFluids

Figure 2 shows the flow curves for concentrated aqueous suspensions of 30 vol% 3Y-zirconia and 70 vol% alumina [l]. Figure 2 illustrates near-Newtonian flow characteristic of suspensions with filler concentrations as high as 56 ~01%.Newtonian flow is generally encountered with highly filled suspensions when the polymeric binder concentration in the suspension is fairly low ( 4 ~01%)and the suspension is highly dispersed (discussed later). Figure 3 shows a gradual decrease in viscosity at different ceramic powder concentrations with

p

E: v,

120

60

0 0

15

30

45

60

75

Shear Rate,

Bo

105

120

135

150

sec"

Figure 2 Flow curves for concentratedaqueoussuspensions of 30 vol%3Y-zirconia and 70 vol% alumina.

243

PLASTIC MOLDING RHEOLOGY AND MIXING 10000

A

v)

*

4 I

l00 100

..

x

55v/O

0

50 V/O

Sic + 40 V/O EEA

SIC + 4 5 V/O E E A Sic + 50 V/O EEA

40 V/O

Sic + 60 V/O EEA

r?

I

10000

l000

Shear Rate (per

60 V/O

S)

Figure 3 Decrease in viscosity at different concentrations of silicon carbide in EEA with increasing shear rate at 150°C.

increasing shear rate at 150°C temperature. The flow behavior is typically pseudoplastic. The power-law index n of these fluids should be less than 1. The suspension consists ofa silicon carbide powder (Starck A-10) and ethylene ethyl acrylate (EEA) polymer (40-60 ~01%).Figure 4 shows (shear stress)ll2 plotted against (shear rate)ln for 55 vol% silicon carbide in polyethylene at 150°C. The extrapolated intercept at zero shear rate gives the Bingham yield value .z, The general phenomenon of dilatancy was observed by Reynolds in 1885. Reynolds [2] observed that a mixture of beach sand and water confined in a balloon would dilate if deformed; that is, the total volume of the mixture would increase under pressure, causing the liquid level in the balloon to be lowered. In addition to the dilation, such a mixture would exhibit increasing resistance to stirring as the rate of stirring was increased, until rupture of the pseudosolid material that resulted occurred. A more general meaning of the term “dilatant fluid” can be given by extending it to all materials showing increasingly apparent viscosity with an increasing rate of shear. Sometimes this is referred to as inverted plasticity to differentiate between volumetric dilation versus no volumetric dilation. Typical dilatant flow curves for two concentrations of A-16SG alumina and paraffin wax are shown in Fig. 5. Reynolds’ approach to dilatancy has some limitations. For example, Reynolds’ approach specifies hexagonal close pack-

244

N

2m

a

" MATSUDDY

0

Shear R a t e 1/2 , Sec 112

Figure 4 (Shear stress)ln against(shear rate)ln for 55 vol% siliconcarbidein 45 vol% polyethylene at 150°C.

ing. Such close packing is represented byavoid volume of approximately 28%. Therefore, it is expected that the volume concentration of dispersed, spherical particles must approach 72% for dilatancy to occur. In reality, dilatancy has been observed in fluids with solid concentrations well below 72% (Fig. 5). Besides, the approach does not account for other factors, such as the wetting characteristics of the suspension fluid for the dispersed solid, the suspension fluid viscosity, and the effect of particle size of the dispersed solid.

B.

Viscoelastic Behavior of Highly Filled Fluids

Figure 6 shows plots of the storage modulus versus oscillation frequency for samples with S i c concentrations of 40, 50, 57, and 63 ~01%.A s the solids

245

PLASTIC MOLDING RHEOLOGY AND MIXING

I 10

1

1

I 1 1 1 1 1 .

i 02

I

-1 Shear Rate, Sec

I

I

I

I l l l j 10

Figure 5 Dilatant flow curves for two concentrations of A-16SG alumina and paraffin wax, 37 and 35 ~01%.at 150°C.

loading increases, at lower frequencies the storage modulus increases as a function of shear rate. The S i c filler increases the rigidity of the polymer. The increases in rigidity implies less mobility of the macromolecular chains under the influence of the applied shear stress. In a highly concentrated suspension (say above 60 vol%), there is particle-particle interaction analogous to the chain entanglements in polymer solutions, showing a tendency for a rapid decrease in storage modulus that is indicative of nonuniform dispersion and agglomeration. This suggests that the storage modulus decreases because the agglomerates do not absorb much energy during interparticle movement.

C.

Influence of Temperature on Fluid Flow

Temperature has a significant effect on the flow behavior of liquids and slurries, including polymer melts, solutions, and filled suspensions. Viscosity de-

246

MATSUDDY r

I

109

-

IO8

- ii

I

I I I Ill1

I

I

1 I 111'1

I

-

V

0

0

b 106

-

V

V

m m mm-

V

m

m

V

m

P

X

54

0

X

0

t

0-

m-

v-

lo5

I

I I 1 1 l l1 l

P@

V

V

X-

1

Ig,

X

X

104L1

I IIIII-

-

h

m a 107 v

I

I

1

40 V/O

Sic + 60 V/O EEA

5 0 V/O SIC

+

50"I/O

EEA

57 V/O SIC + 4 3 V/O EEA 6 3 V/O sic + 37 V / O EEA

I I I I l l10 1

I

I

I I I Ill

Frequency, d l =

Figure 6 Storage modulus versus oscillating frequency for 40, 50, 57, and 63 vel% silicon carbide with EEA at 200°C.

creases with increasing temperature for both Newtonian and non-Newtonian fluids, primarily because of a decrease in intermolecular forces. The increased flexibility of a polymer backbone chain also contributes to reduce viscosity. The relation between viscosity and the temperature of a filled suspension can be expressed by the modified Eyring equation, which allows for variations in the free volume of the fluid structure:

In this equation h is Planck's constant, No is Avogadro's number, Vm is the molar volume of the fluid phase, E is the activation energy needed to overcome the potential barrier between equilibrium positions, y is a constant in the range 0.5-1.0, V0 is the van der Waals volume of the molecule, and V is the average volume per molecule in the fluid. Because of the exponential form of the viscosity equation, mixtures of ceramic powder in thermoplastic binder are only significantly sensitive to temperature change above 70°C [3]. In an injection molding operation, it is usual to plot In q versus T for a constant shear rate, not a In q versus 1/T plot at a

PLASTIC MOLDING RHEOLOGY AND MIXING

247

353 363 373 Temperature(K) Figure 7 Effect of temperature on viscosity at a constant rate of shear of 100 S-’. includes 55.0 vol% A1203 and 45 vol% ethylene-vinyl acetate. (After Ref. 3).

constant force. An example of one such plot for 55.0 vol% A1203 loading in 45 vol% ethylene-vinyl acetate copolymer is shown in Fig. 7. The linearity of sucha plot indicates the process sensitivity to temperature change, and the slope provides a measure of the activation energy for viscous flow. For a filled suspension, the effect of temperature on particle-particle interaction must be considered, in addition to the effect of temperature on inter- and intramolecular interaction in polymers. This temperature effect depends on the level of loading. If the temperature is sufficiently high to allow migration of binder from particle-particle interfaces, then the suspension may exhibit dilatancy in place of pseudoplasticity. The effect of temperature is similar to that of shear rate on polymer rheology, but its effect on particle-particle interaction is unpredictable. Both high temperature and high shear rate weaken polymerparticle interaction. However, it is important to remember that the extent of the effect (viscosity and interaction reduction) for a particular change in shear rate or temperature is not predictable. Both factors, exerting their effects at different rates, contribute to the complexity of the system.

D. Effect of FillerConcentration Concentration effects and interparticle forces are of rheological significance with injection molding mixes. Increasing the filler concentration (volume fraction) of a suspension increases the suspension relative viscosity q~ = q/qo, where q is the suspension viscosity and q o is the continuous-phase viscosity.

248

MATSUDDY

In the dilute region (CD O.l), q~is a linear function of CD and for spheres is found to obey the Einstein relation [4]

(5)

q R= l + 2.5$ As CD increases, particle-particle interactions become significant and Eq. (5) takes the form of a general power-series equation in volume fraction [5]:

-=1+2.5t$+kt$'+ rl TO

+

The second-order coefficient k need not be a constant. In general, k is a complex function of the variables that determine particle interactions. DeBruijn [6] assumed that Q. (6) was a quadratic and that the relative viscosity T R would become infinite when the volume fraction reached the value for cubic closepacked spheres. The value of k then becomes 4.7, but it increases slowly as the particles are more disordered. According to Vand's [7] assumption, however, if pairs of particles separate along rectilinear paths after collision, then k has a value of 7.4. A setback with this approach is that at higher concentrations a large number of terms in Eq. (6) are required to calculate T R , with the greatest influence moving to the terms with high powers of CD. Allowing for interaction between particles and for interactions between particles and continuous phase, Mooney [8] proposed the expression

and introduced k as l/CDmax.In the limit CD + 0, Eq. (7) reduces to 1 + B@. Thus, B can be interpreted as an Einstein coefficient and should be between 2.5 and 5, depending on the degree of agglomeration and the extent of elec~ m; therefore, k-* trochemical forces between the particles. As CD + l/k, r ) + refers to the maximum possible value of CD, that is, the value for which the suspension loses mobility, k is known as a crowding factor. Equation (7) provides a reasonable fit; as the data for an A1203-EVA (ethylene-vinyl acetate) copolymer (with M n = 3500 and melt flow index = 200, the melt flow index defined in terms of the amount of polymer that flows through a die in a given period of time: the higher the flow rate and thus the lower the viscosity, the higher the melt flow index) show the 4 2 0 3 had an irregular shape and agglomerated with a mean particle size 0.62 pm. Relative viscosities at four powder loadings are given in Table 1. Taking logs of Eq. (7),

B$ lnqR = l-kt$

249

PLASTIC MOLDING RHEOLOGY AND MIXING Table 1

RelativeViscosity Determinedat a Shear Rate of 100s-1 of A1203-EVA Molding Mixes vol.% A1203

Relative viscosity

1.83 41.0 48.0 3.91 57.0 63.0

6.25 12.55 49.90

In 7

) ~

2.53 5.30

Source: After Ref. 3.

A plot of this equation, shown in Fig. 8, can be used to find the slope at x, y, and

@m:

B 1 - 0.41/$, B 11.43 = 1 - 0.48/$, 7.5 =

at x at Y

Hence B = 2.46, which is a close fit to the Einstein constant 2S@m = 61.396, which is less than perfect close packing of the powder, an effect contributing to shape irregularity and agglomeration.

5 CI

F4~E

v

3 2 -

1 -

0.1 0 . 2 0.3

0 . 4 0.5 0.6

0.7

Volume Fraction Figure 8 Plot of <<m.

m.(8) using data from Table 1 can provide the slope at

X,

y, and

250

111.

MATSUDDY

MIXING

The objective in the mixing process is to disperse uniformly a relatively small amount of binder in a much larger quantity of fine powder and to achieve a homogeneous mix-one free of agglomerates that contains the optimum ceramic and binder content while maintaining sufficient fluidity for molding. It is a critical process step in producing high-quality, reliable injection-molded ceramic parts. Nevertheless, it has always been regarded as a simple operation, partly because relatively crude methods of mixing can be made effective if the operation is continued over an extended period of time. Mixing involves the transport of the material in the mixer to produce the desired spatial arrangement of the individual component.

A.

Mixing Mechanism

According to Lacey [9], the transport mechanisms can be grouped into the following steps 1. Diffusion mixing: If two materials are placed in a container, after a suffi-

cient length of time has elapsed the molecules of both materials commingle as a consequence of both the concentration gradient and random molecular motion. The result is the formation ofa uniform mixture ona submicrometer scale by molecular diffusion. The type of diffusion process usually encountered in practical mixing situations results from velocity gradients within the fluid. This diffusion process by itself, however, is too slow for the preparation of ceramic molding mixes. 2. Laminar mixing: The term “laminar” is applied to any operation that increases the randomness ofthe spatial distribution of the binder constituents or particles in a system, without reducing their size. In laminar mixing, the binder constituents pass from a less to a more random arrangement. 3. Dispersive mixing: An important application of dispersive mixing is the incorporation of the ceramic powder into the molten organic binder. During dispersion, the agglomerates of the constituent elements of the secondary component of the mixture are ruptured. Whenever there is relative motion between a fluid and a particle, drag forces act on the particle’s surface. If the fluid velocity is greater than the particle velocity, these drag forces tend to accelerate the particle. If the particle velocity is greater than the fluid velocity, they tend to retard its movement. If the particles are spherical and if the Reynolds number (Reynolds number represents the ratio of inertial forces to the viscous force-dominated regime; that is, a low Reynoldsnumber, to one that is inertial force dominated, that is, ahigh Reynolds number) for the flow of a particle is low, then according to Stokes’ law, the drag force exerted on this particle is equal to 6xReqV, where V is the

PLASTIC MOLDING RHEOLOGY AND MIXING

251

relative velocity between the fluid and the particle, Re is the Reynolds number, 71 is the viscosity of the suspension, and 6n is a geometric factor for a very large diameter tube. High shear stresses promote dispersion in dispersive mixing. In any given system, there is some stress, defined as the critical stress, below which dispersion cannot occur.However,when the shear stress is just slightly larger than the critical stress, only agglomerates with a favorable initial orientation become dispersed. If the flow in the mixer is completely unidirectional, only agglomerates with a favorable initial orientation are ruptured. The others are aligned with the flow, and no further dispersion takes place evenwith continued mixing. However, if the mixer causes the flow pattern to change direction continually, further dispersion can occur as mixing continues. Ultimately, the shear stresses created position all the remaining agglomerates in an orientation favorable for rupture to occur.As mixing proceeds, all three mechanisms function to some extent.

B. Effect of Mixing Time In mixing operations it is essential to know the time required to achieve a homogeneous mix, often referred to as degree of mixedness. Information on the rate of mixing would allow us to understand the basic mechanisms of mixing and the stability of the binder components and to select an appropriate mixing device. In the initial stages of mixing, the binder components change progressively from solid or liquid to semifluid and finally to a fluid state. Mixing mechanisms at this stage are likely to be involved with interdiffusion, mutual solubility of binder components, and binder (molten) interaction with ceramic particle surfaces. Agglomerates are formed (1) as a result of the change in interfacial tensions between the powder particles andmelted binder components and (2) as the entire mix slowly turns into a viscous mass. Subsequent mixing of the viscous mass is essentially dispersive or shear. Since mixing tends toward an equilibrium state, it is easy to see that its rate of attainment is likely to be of the form

where c is the rate constant and t the time, so that X approaches its limiting value asymptotically, as in Fig. 9. From Eq. (9), a plot of -log, (1 - x ) against time should give a straight line of slope c, and the effectiveness of dispersive mixing is indicated by the value of this constant. Although the experiment suggests a relationship of this form, nonetheless it cannot necessarily be taken as complete verification of any particular mechanism that invokes an exponential

252

MATSUDDY

t Figure 9 Rate of completemixing of anattainment asymptotic relation with time.

of equilibriumstateshows

function of time, because temperature and speed (rpm) [lo], as shown in Figs. 10 and 11, can also significantly influence the mechanism of mixing. It is evident from Fig. 10 that the shear mixing becomes progressively less effective as the temperature is increased and as the mix becomes less viscous. On the other hand, the low torque generated at 10 rpm shows (Fig. 11) that the shear forces are insufficient to provide uniform dispersion of the binder in the mix. Further, oxidation of binder components in air with mixing time have been recorded [11,121 from gel-permeation chromatography and infrared spectroscopy (Fig. 12). Figure 12a compares the spectra obtained with polyethylene before and after kneading. Changes in infrared spectra associated with kneaded polyethylene at 160°C for 1 h can be attributed to oxidative degradation. The degradation of polyethylene binder during kneading is further substantiated from the observed shift in the chromatographic peaks (Fig. 12b). Using process tolerance capability (defined as three times the standard deviation CT divided by the final product dimension L) as an indicator of mixing efficiency, Billiet [l31 showed the time dependence of the mixing operation (Fig. 13). Insufficient mixing time negatively affects the process tolerance capability. There is an optimum mixing time for which this capability is greatest; that is, the tolerance that can be held confidently reaches the lowest value. Extending mixing time beyond this time period produces a negative effect on tolerance capability. It has been suggested that the time region I represents a net randomizing area inwhich increased mixing time has the effect ofhomogeneizing the mixture and the region 11 is a symmetrical region in which period doubling occurs and the mixture shows areas of inhomogeneity.

C.

Role of ProcessingAids and Dispersion

Shear force alone cannot break down agglomerates. A s a first step, it is essential to remove the moisture from the powder before preparing a mix. Further-

253

PLASTIC MOLDING RHEOLOGY AND MIXING

.i - 1

50 vol% A1203 Mixing Conditions: 125%, 200 rpm

600

500

v

i-

4001 300

4007

50 ~01%AI2O3

Mixing Condilions: 150”C, 200 rpm

I

200

-

l00

-

0 200

50 V O W AI2 0, Mixing Conditions: 175°C. 200 rpm I

1

I

I

1

I

-

1

I

I

I

I

50 voi% A1203

I

0

I

220°C. Conditions: Mixing

100 -

0

I

I

5

1

10

15

200 rpm I

20

I

I

25

I

I

30

I

TIME (min)

Figure 10 Plots of torque versus mixing time for 50 vol% A1203 and 50 vol% polyethylene mixed at different temperatures. (After Ref. 9.)

254

MATSUDDY 500 50 Val% Ai2 03 Mixing Conditions: 150°C,10 rpm

400

300 W

3 0

U

p

200

o lo

Q

15 5

0

10

30

25

20

TIME (min) 500

50 ~ 0 1 % AI2 0, Mixing Conditions: 150eC, 200 rpm

4OQ-

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W 3

0

300-

-

U

p

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100-

0 - L 0

1

I

15 5

I

1

10

*I

I

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1 20

1

I

25

l

I

30

TIME (min)

Figure 11 Plots of torque versus mixing time for 50 vol% A1203 and 50 vol% polyethylene as a function of speed. (After Ref. 10.)

I

PLASTIC MOLDING RHEOLOGY AND MIXING

10

255

-

> t-

M

-l

c.

m Q 5m 0

a (L

---- A2-PE.60x

0-

lo3

I I ~ I I I I I I I ~ I I I IIIIII J loc IO* lo6 MOLECULAR WEIGHT

I I II

(b)

Figure 12 (a) Infraredspectra ofpolyethyleneillustratingtheeffectofmixing, (After Ref. 11.) (b) Molecularweightdistribution of polyethylene andalumina-polyethylene mixes. (After Ref. 11.)

256

MATSUDDY

Region I

Mixing Time Figure 13 Regionof optimum mixing time.(After Ref. 13.)

more, it is often necessary to premill the powder with a processing aid that modifies the surface characteristics of the powder and breaks down the agglomerates. A brief account is presented here to amplify the role of processing aid(s) in achieving an agglomerate-free homogeneous mix. In general, dispersion forces give rise to an attraction between particles in a fluid whose polarizability differs from.that of the particles. The interaction potential is inversely dependent on particle separation, and the magnitude of the interaction is described by the Hamaker [l41 constant. The well-known DLVO [15,16] theory describes the stabilization of aqueous suspensions of particles on the basis of coulombic repulsions between charged particle surfaces. It is not clear to what extent charged double layers are effective in nonaqueous suspensions. This is the subject of another chapter in this book. An alternative approach to overcoming these forces in organic fluids is to attach or graft polymer chains to the particle surfaces such that steric hindrance, the so-called entropic repulsion, prevents the particles from approaching one another more closely than the extended length of these chains in the surrounding media. The dispersion process can be illustrated in three distinct stages. First, wetting, in which the powdedair interface is replaced by the powderlfluid interface; second, mechanical disruption, as discussed earlier, in which some of the agglomerates are broken down into smaller particles; and third, stabilization of the resulting dispersion against flocculation.

PLASTIC MOLDING RHEOLOGY AND MIXING

257

Wetting is the displacement from a surface of one fluid by another. It therefore involves three phases, at least two of which must be fluids. Three types of wetting can be distinguished (1) spreading, (2) adhesional, and (3) immersional. For wetting, the function adhesion tension, y cos CP, is involved, where y is the surface tension of the liquid and CP is the contact angle of the liquid with the solid surface. The author used a simple method of measurement of adhesion tension of a number of processing aids in contact with powders. This was based on measurement of the rate of unrestricted penetration of the processing aid into a column of A16-SG A1203 powder in a small glass tube. The adhesion tension y cos CP is then estimated from '" -1=

t

r ycosq k2 271

where l is the distance of penetration in time t to the surface tension y, contact angle CP, and viscosity 71 of the advancing liquid, r is the radius of the capillaries between particles, and k is a tortuosity constant. In this equation, r/k2 is a constant depending upon the packing of the column and was eliminated by the use of a standard packing. Table 2 presents the calculated adhesion tension y cos CP for processing aids. It has been noted that the higher the measured adhesion tension, the more rapid is the rate of dispersion during the mixing operation. The technique is not suitable for use with polymers that are solid at room temperature, however, because a linear relation between the square of the distance of penetration and time is not attained. For high-melting polymers and waxes, the contact angles of polymers to ceramic substrates have been measured as a function of temperature and time using a goniometer.Figures 14 and 15 illustrate the effect of temperature on contact angle for different polymers [10,17]. Table 3 presents some data for thermoplastic resins and waxes with alumina and silicon nitride substrates. The atactic polypropylene has given by far the best wettability of polymers and waxes. The paraffin and most of the microcrystalline waxes give low contact angles for both alumina and silicon niTable 2

CalculatedAdhesionTension for a Selected Processing Aidand A16-SG Alumina y cos 0

Processing Adhesion aid tension Corn oil Camphorated cottonseed oil Peanut oil Olive oil Pine oil Olein Menhaden oil

22.0 17.3 12.4 8.5 7.6 20.7 24.8

TEMPERATURE (C* 1

Figure 14 Contact angle of resin to alumina substrate (A) AF'P, 10,000,20,000mol(D) ecularweight; (B) M P , 6000 molecularweight;(C)low-densitypolyethylene; EVA (30% vinylacetate);(E), (F) medium-densitypolyethylene; (G) polystyrene. (After Ref. 17.)

tride, whereas the contact angles for polyethylene and polystyrene withalumina and silicon nitride are relatively high. Besides temperature, it can be assumedthatthesubstrate surface characteristics,suchas adsorbed species, porosity, cleanliness, and morphology, have a profound effect on the wetting of polymers and waxes. A similar influence might be expected with fine powders. There is clearly scope for further work on the surface nature and wettability of fine powders, so that progress in the development of improved dispersability in injection molding mixes can be placed on a more rational basis. The mechanical dispersion as applied to mixing has already been discussed. The final stage of the dispersion process is that of stabilization of the dispersed particles against flocculation. It is not intended to give a detailed account of the theoretical basis of dispersion stabilization, because this is extensively covered inthe literature, at leastat lower particle concentrations in

259

PLASTIC MOLDING RHEOLOGY AND MIXING 60

Polyethylene/Alp03 50

40

30

20

10

0 0 20

10

30

40

50 70

60

80

90

TIME (mln)

Figure 15 Contact angle as a function of time for polyethylene on alumina substrate at different temperatures. (After Ref. IO.)

low-viscosity liquids (e.g., water or solvents). Very little information is available on dispersion stability of suspensions with particle concentrations above 50 ~01%.An unstabilized dispersion flocculates by collision of the primary particles and thereby reduces the number of particles P in the system: dP = 81tDRp’ -dt

where R is the collision radius of the particles, D is the diffusion coefficient, and t is the time. D can be expressed as

for particles of radius a in a fluid of viscosity 1 1 ,where k is the Boltzmann constant and T is the absolute temperature. It has been shown [l81 that the time required to reduce the number of particles to half the initial value t112 is given by the equation 412

311 =-

4KTP0

260

MATSUDDY

Table 3 Wetting Angles of Thermoplastic Resins and Waxes with Ceramic Substrates Contact angle 0, (“1 Thermoplastics and waxes

Si3N4

Temperature (“C)

Polyethylene NA-250 Polystyrene D-210 Atactic polypropylene Waxes Paraffin Shell 120 Microcrystalline 18654 Microcrystalline Shell 400 Microcrystalline C1035 Microcrystalline Mobil

200 200 160

-33 -40 -8

-36

160 160 160 160 160

-15 -14 -30 -20 -20

<S

-

<5 <5 <5 <5

At higher particle concentrations (-55 vol%), tln decreases rapidly. Thus, the fluids become unstable at higher particle concentrations, and some force of energy barrier is needed to avoid flocculation by collision of primary particles.

of Mixedness

D.Degree

Besides deagglomeration, it is also essential to assess the extent of homogeneity, that is, the degree of mixedness. Because of the random nature of the mixing process, statistical analysis is most frequently used. This concerns primarily measurement of the standard deviation or the variance of the spot samples taken from a mixture. The criteria are then expressed in different forms of the standard deviation of the variance. So far, the statistical analysis of mixedness in injection molding mixes has not been fully developed. However, Billiet [l31 made a significant observation while he was trying to determine the maximum mixing error Em, that is, the error in injection molding mixes that would result in the loss of the targeted process tolerance capability T, and showed the need for extreme care in weighing and handling at the mix preparation stage. He used the following expression to relate maximum mixing E@ and process tolerance capability T:

IEd= G 3T

where is the volumetric powder loading. Measurement of mixedness is often subject to appreciable experimental and sampling error. Therefore, experimental verification of the validity of the theoretical criteria for the degree of mixedness of a mix is not always satisfactory.

PLASTIC MOLDING RHEOLOGY

261

AND MlXING

Simulation of a mixing process by using a computer provides a better alternative. Monte Carlo simulation techniques are often utilized. Monte Carlo techniques are numerical methods that involve sampling from statistical distributions, either theoretical or empirical, to approximate the real physical phenomena without reference to the actual physical systems. For the general discussion of Monte Carlo methods, readers are referred to Hammersley and Handscomb [l91 and Tocher [20].

E. Other Factors The rate of mixing and the degree of mixedness are functions of many variables relating to characteristics of the powder, mixing equipment, and operating conditions. Some aspects of powder characteristics were discussed earlier in the light of rheology. The author has also published [21] some information on mixing equipment. Readers are well advised to refer to these publications. Again, there is no systematic analysis of the operating conditions for preparing injection molding mixes. On a broader basis, careful attention must be given during the mixing operation to a number of factors, such as weighing of each constituent added, the ratio of volume of the mix to that of the mixer, method, sequence, and rate of adding constituents, and mixer speed.

REFERENCES 1. Novich, B. E., Lee, R. R., Franks, G. V., and Ouellette, D., Quickset injection molding of high temperature gas turbine engine components, inRoc. 27th Auto-

motive Technology Development Contractors’ Coordination Meeting, Dearborn, MI, 1980, pp. 311-318. 2. Reynolds, 0..Philos., Mag., 8, 20 (1885). 3. Mutsuddy, B. C., Influence of powder characteristics on the rheology of ceramic injection molding mixtures, Proc. Bri. Cerum. Soc. Fabrication Sci., 3, 117-137

(1983). 4. Einstein, A., in Investigation on the Theory of Brownian Movement Furth, ed.), Dover, New York, 1956, p. 55. 5. Goodwin, J. W., The rheology of dispersion, J. Coll. Sci., 2, 246-293 (1975). 6. DeBruijn, H.,The viscosity of suspensions of spherical particles, J. Rec. Trav. Chim., 61,863-874 (1942). 7. Vand, V., Viscosity of solutionsandsuspension,Iand II, Phys. Chem., 52, 277-314 (1948). 8. Mooney,M,,Theviscosity of a concentrated suspension of spherical particles, J. Coll. Sci., 6, 162-170(1951). 9. Lacey, P. M. C., Developments in the theory of particle mixing, J. Appl. Chem., 4, 257-268 (1954). 10. Dow, J. H., Sacks, M. D., and Shenoy, A. V., Dispersion of ceramic particles in polymer melts, Cerum. Trans., Z(A), 380-388 (1988).

(R.

262

MATSUDDY

11. Takahashi, M., Suzuki, S., and Nitanda, H., private communication. 12. Hunt, K. N., Evans, J. R. G., and Woodthorpe, J., The influence of mixing route

on the properties of ceramic injection molding blends, Br. Ceram. Trans. J., 87,

17-21(1988). 13. Billiet, R. L., Injectionmolding of advanced PM materialsinSEAsia, Metal Powder Rep., 45(5), pp 326-332 (1990). 14. Hamaker, H. C., Physika, 4, 1058 (1937). 15. Deryagin, B. V., and Landau, L.,Acta Phys. Chim URSS, 14, 633 (1941). 16. Verwey, E. J. W., and Overbeek, J. Th. G., Theory ofthe Stability ofLyophobic Colloids, Elsevier, Amsterdam, 1948. 17. Saito, K., Tanaka, T., and Hibino, T., Method of producing a ceramic article by injection molding, U.S. Patent 4,000,110 (1976). 18. Overbeek, J. T. G., in Kruyt, Colloid Sci., l (1952). 19. Hammersley, J. M., and Handscomb, D. C., Monte Carlo Methods, Methuen & Son, London, 1964. 20. Tocher, K. D., The Art of Simulation, English Univ. Press, London, 1963. 21. Mutsuddy, B. C., Equipment selection for injection molding, Bull. Am. Ceram Soc., 68(10), 1796-1803(1989).

IV SOL-GEL PROCESSING

This Page Intentionally Left Blank

l1 Processing of Monolithic Ceramics via Sol-Gel J. Phalippou Universitk de Montpellier II Montpellier, France

1.

INTRODUCTION

The sol-gel process has attracted increasing interest over the last decades because it permits the development ofnew materials andnew shaping routes (fibers,thinfilms,andnearnet shape objects)withgoodhomogeneityand good purity. Moreover, the sol-gel process is expected to be very effective for ceramic materials, which must be sintered into a dense body. The very fine structure of the gel enables the sintering to be achieved at low temperatures compared with green bodies prepared by classic ceramic technology. This is particularlytrue for amorphous monolithic ceramics, which represent the largest family of compounds synthesized from a sol-gel. In this case, sintering, which occurs by viscous flow, is easy to perform and does not require long thermal treatment. It is important to emphasize that the sintering temperature increases with the size of the elementary particles (or pores) constituting the connected solid part (Fig. 1). Most experiments reported to date have been devoted to monolithic silica materials. Fundamental scientific work is most easily done for a pure component. At the same time, silica is inherently interesting because hydrolysis and polycondensation reactions lead to siloxane bonds that build up a strong network, with irreversible chemical bonds. The main disadvantage of gels is their high porosity. This property may be used to improve the thermal insulation of aerogels, but in most cases it is con265

266

PHALJPPOU

- 0

O\!

v

t

L

20

i 30

0

800

400

1200

T(OC)

Figure 1 Linearshrinkageforsilicaxerogelspreparedbytwo-stepacid-catalyzed hydrolysis (A) and acid-base
II. CAUSES OF CRACKING IN GELS Cracking of a gel can occur at each step of the process, from synthesis to the final densification of the material.

SOL-GEL PROCESSING OF MONOLITHIC CERAMICS

267

10'

Figure 2 Young modulus and modulus of rupture of silica aerogels as a function of the aerogel bulk density. (From Woignier and Phalippou [2].)

A.

Cracks Appearing in Wet Gel

Cracks may appear during the gelation step. In this state the gel begins to form and cracking is mainly related to several effects. I t is well known that gels prepared under basic conditions show a high tendency to crack in the first instants after gelation occurs. Under basic conditions the rate of polycondensation is high, producing a strong exothermic effect. As a result of this temperature rise, the liquid inside the pores of the gel expands more than the network and tries to escape the solid part of the gel. Because of the low permeability of the gel, the liquid cannot escape freely and the induced pressure results in an increase in the geometrical dimensions of the solid gel part [3]. This effect is not expected to occur in gels with large pores and highly permeability. Also, strong gels do not expand much and can resist the stress induced by liquid expansion. In this case, only a small expansion of the solid network is expected. Cracks may then suddenly appear even in the absence of measurable strain.

268

PWPPOU

It was shown [3] how expansion measurements enable the permeability of the gels to be evaluated (Fig. 3). The increase in the volume inside the pores of a gel is due to the high thermal expansion coefficient of the liquid and the flux of liquid (time dependent), which is proportional to the pressure gradient, the liquid viscosity, and the gel permeability. The change in the volume of a liquid is related to the volumetric strain of the network. A more common cause of gel cracking in the wet state is related to the aging effect. The aging effect begins after gelation and is characterized by several features (polymerization, syneresis, and so on) that induce structural and textural changes. During aging, in particular for acid and neutral silica gels, a high degree of shrinkage occurs. If the solid part of the gel sticks to the walls of the container, cracking occurs; this phenomenon is readily avoided by using fluoropolymers or polypropylene containers, which permit no adhesion with gels. The shrinkage of the gel is related to hydrolysis and condensation, which proceed well after the gelation point. Syneresis liquid is compressed within the pores of the gel. Liquid is then in compression and consequently induces the solid part to be under tension, which may lead to fracture of the gel. When the network of the gel is full of liquid, cracking is always related to the permeability and to the mechanical strength. Note that, as demonstrated be-

4

i /T:

> 34

(L

'l

-

l l l \

so 2 : 4-; a

d

0

F

0

'

'

I

l

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;

...."I

.

'\\"+\I\

32 n

0

c I

4

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... ................(..........p ..... ...\......_. 28 t

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-2

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1

40

I

I

80

I 26

Figure 3 Dimensional variation in the solid part of a gel (dashed line, left ordinate) subjected to a temperature cycle (solid curve, right ordinate). Time is expressed on the abscissa. The gel expands first during the heating run and recovers its starting dimensions when submitted to an isothermal treatment. Cooling from 33°C induces a gel contraction. (From Scherer [3].)

SOL-GEL PROCESSING OF MONOLJTHIC CERAMICS

9

269

fore, a highly porous and weak gel may not crack if its permeability is high enough to allow the liquid to escape without friction. These effects were also found during the temperature rise of a gel treated in an autoclave. Scherer [4] demonstrated thatheating rate, aging,and gel dimension play very important roles in the Occurrence of cracking. Cracks can also appear during the pressure release in the autoclave. In the supercritical drying process, the gel is subjected to high temperature and high pressure. When the critical point is reached, the pressure of the autoclave is decreased while the temperature is kept constant. At this instant, the pressure applied to the supercritical fluid is equal to that within the pores. The supercritical fluid has a very low density and viscosity compared with thatof the liquid at room temperature; however, the low permeability of the gel resists the flow of the supercritical fluid out of the gel. In other words, if the supercritical fluid release is performed too fast a pressure gradient appears. In this case the supercritical fluid within the gel, which is in compression, suddenly expands and the solid part suffers tensile stress. Experiments show that cracking depends on the pressure release rate, on the nature of the gel (basic or neutral), and on its geometrical dimensions. Cracks can also occur when a gel prepared in an alcoholic medium is placed into another liquid. This effect is particularly seen when alcohol is removed by water or acetone [5]. Dimension variations in the sample are expected to be produced by this solvent exchange. The change in surface energy between the solid and the liquid, the osmotic pressure, or differences between diffusion rates of liquids are believed to induce this effect. To avoid this kind of cracking a slow solvent exchange can be performed. An exchange using Soxhlet equipment may be useful in overcoming this difficulty [6].

B.

CracksAppearingDuringDrying

The drying of gels has been extensively investigated [7], and only a brief survey is given here. First, it is important to note that cracking is always related to the establishment of capillary forces, which appear when the solid part of the gel comes in contact with the vapor. The pressure change versus radius of the meniscus is given by Laplace’s law, which stipulates that the pressure variation is related to the mean curvature x: x=-

1

RI

1 +R,

of the meniscus at a given point. In the relationship

270

PHALIPPOU

' ~ L Vis

the liquidvapor interfacial energy. hp is positive if one passes through the curved surface from the convex side. It is noteworthy that the mean curvature x may be evaluated in any way provided that radii R1 and R2 are orthogonal. This effect, which applies to microscale texture, induces a macroscopic evolution of the network. At the onset of drying, the surface of the liquid is flat. The curvature increases as the liquid of the gel evaporates. The liquid is then in tension, and as a consequence the solid part of the gel is under compression. This effect causes the gel network to shrink. That shrinkage continues as long as the solid network (depending on the nature of the gel) is not stiff enough to resist the compressive stress. During this shrinkage, the permeability of the gel, which is low, decreases further. While shrinkage continues, the quantity of liquid removed by evaporation is balanced by the liquid flow to the surface. Darcy's law gives the liquid flux J due to the pressure gradient VP: J=DVP IlL

This flux obviously depends on the permeability D and on the liquid viscosity q ~Since . the evaporation rate Ev must equal J at the surface,

Because it is the pressure gradient that causes differential strain and cracking, it is evident that fast evaporation and low permeability are detrimental. Cracks usually appear at the end of the constant rate period when the shrinkage stops and the meniscus recedes into the pore. At this point, the radius of the meniscus is minimal and depends on the pore radius r, and contact angle 0: x=- 2 cos

e

'P

Capillary stresses then reach a maximal value. It appears that cracks may be avoided if the liquid evaporation rate is very slow. This was the first way to achieve monolithicity in gels, but it is evident that this drying process is very long (months). Moreover, the dry gel (xerogel) exhibits a very low permeability, which may present manydisadvantages when the xerogel must be heated further. A convenient way to avoid capillary forces is to perform supercritical drying. If the process is carried out carefully, a monolithic aerogel is obtained in avery short time (1 day). Moreover, the aerogel has apermeabilitymuch higher than that of its corresponding monolithic xerogel.

271

SOLGEL PROCESSING OF MONOLITHIC CERAMICS

C. Cracks Appearing

During Thermal Treatment

In most cases, xerogels are to be densified to obtain material for specific applications. They are very brittle, thus densification improves their mechanical properties. The thermal treatment of a xerogel must be carried out with care. At low temperatures the xerogel, which contains unreacted organic groups, must be gently oxidized. This treatment is often performedusing air, butlong-chain residues require long treatments under oxygen fully to achieve oxidation. The released gas contains C02, H20, and also CO [8] and some other low-weight organic molecules [g]. Darkening of the xerogel indicates a poor oxidation treatment. Difficulties may be encountered in xerogels exhibiting a high specific surface area. The escape of gas by-products is hindered by the low xerogel permeability. Experiments performed on silica xerogels show that an expansion of the gel dimension is observed. This expansion depends on the previous heat treatment of gels (Fig. 4). An explanation may be offered to take account of this behavior. The escape of gas by-products is hindered by the low xerogel permeability. The produced gas cannot escape easily. Its pressure increases, and thus the xerogel expands. A fast heating rate can lead to a cracked xerogel. To avoid this phenomenon, the thermal schedule must be optimized. However, more experiments are necessary to confirm this proposed explanation. 6

I

I

I

~

o-o-

I

I

300 O C

-4 -6

'

0

I

100

I

200

I

T ( " C ) 300

I

400

Figure 4 Linearthermalexpansion of thesamexerogel previously heattreatedat different temperatures for 2 hs. (From Kawaguchi et al. [lo].)

272

PHALJPPOU

Regarding permeability, aerogels .are more permeable than xerogels. However, cracks can also occur in aerogels during the oxidation stage. Aerogels prepared by hypercritical drying of alcohol are hydrophobic because they are totally esterified during the autoclave process. They contain a high quantity of organic groups, and oxidation, even though easy to perform because of the open structure, produces a large amount of gases. Gas production resulting from the oxidation reaction is not the only reason for the observed phenomenon. Often a small amount of liquid remains condensed inside the smallest pores as a result of capillary condensation. Above a temperature of about 12O-15O0C, the liquid, for instance water, evaporates and creates a similar effect. Previously oxidized aerogel samples expand when they are heat treated as a consequence of adsorbed water evaporation. They shrink again in a small temperature range because of the adsorption of air moisture when cooled in No shrinkage is recorded if cooling is performed under vacuum. Obviously, the expansion is heating rate dependent [ 1l]. It is noteworthy that to avoid cracking during drying, gels with large pores are desirable. However, this is not the case for sintering. Of course, the most desirable gels for sintering are materials showing the best compromise between permeability, strength, and high green density. Cracking at higher temperatures is rarely observed for xerogels or aerogels for whichsintering occurs by viscous flow. Duringthesinteringtreatment, however, if a gel that is not totally dense is taken from the furnace and cooled to room temperature, it again adsorbs air moisture. It can crack during a subsequent heating run.

air.

111.

ROUTES TO AVOID CRACKING

In the previous section we have seen the most important parameters that influence the monolithicity. A slow evaporation rate ora supercritical drying process in any case can lead to monolithic samples. However, it is possible to modify some parameters to increase the ability of a gel to retain its monolithicity. First, as expressed by Laplace’s law, the pressure is lowered when y ~ v decreases. The addition of surfactant compounds [l21 in small quantities can be used. Alcohol or acetone has a lower surface tension than water ( y ~ v= 73 erg/cm2).However,their pressure vapor is highand these liquids evaporate rapidly.More advantageous liquids are formamide or dimethylformamide @Mm. These liquids evaporate well after water, leaving the gel surface covered by a liquid with a lower y ~ than v water. When such a gel is treated above the temperature of decomposition of formamide or DMF and then cooled to room temperature, water is resorbed and the gel cracks very rapidly. In fact, lowering y ~ decreases v the maximum tension in the liquid at the gel surface. The total stress on the gel results from the pressure gradient, however, which

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mainly depends on Ev and T L . With respect to this last parameter, formamide is a bad solvent if drying is not performed at high temperatures to reduce qL. Lowering y ~ may v be effectuated by a solvent exchange, but because of the low permeability of the gel this treatment requires a long time to be achieved effectively. Laplace's law indicates that any treatment that induces an increase in the pore size is beneficial to monolithicity. Several methods are available to enlarge the pore size. Dissolution and redeposition phenomena are currentlyused to reach this goal. The dissolution of convex asperities is balanced by deposition on concave regions. The net result is an increase in the neck size between the elementary particles and the smoothing of asperities. The smallest pores disappear first. Several parameters may be used to increase the dissolution phenomenon. At room temperature amorphous silica is insoluble (5 ppm) in alcohol, but its solubility is around 70 ppm in water [13]. Under basic conditions the solubility of silica increases. The temperature is a parameter that can be used to increase silica solubility,and these experiments are advantageously performedunder pressure in 'an autoclave.However,the temperature must belowenough (<180°C) to avoid silica-quartz transformation. From the point of view of chemistry, it is probably possible to create silicic acid in the pore liquid. In this case, the amount of silica in the liquid remaining constant, a deposition into crevices happens and dissolution is prevented. To our knowledge this kind of experiment has not yet been reported. Drying control chemical additives (DCCA) (such as formamide, oxalic acid, and glycerol)[141 have often been suggested to increase the monolithicity because the pore size of the xerogel is larger than that of usual xerogel. Furthermore, the pore size distributionis narrower. In fact, to support this assumption, the pore size measurement would have to be done on the wet gel (by mean of thermoporometry) [5]. These experiments have not yet been reported. The reported uniformity of the pore size distribution may originate Erom DCCA Chemistry, which modifies the aggregation process.On the other hand, during drying, because alcoholis removed first, these compounds can stick to the silica surface. As a consequence they increase the network strength, which hinders extended shrinkage of the gels during drying. Consequently, an increase in the mean pore size is observed. It was speculated that a high permeability favors monolithicity. Increasing the pore size is usually performed by washing in an acid(HF) or basic medium. However, other techniques present the advantage of not removing the material. For example, permeability is increased if large pores are created in the gel texture. The easiest method is to add fillers to the gel. Fumed silica, Stober silica particles [15], and Ludox (a colloidal silica solution) are the fillers most often used for silica gels. Fillers decrease shrinkage during drying and prevent the collapse of pores.

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It must be noted that the permeability of an assembly of fine particles decreases when a small amount of coarse particles is added [16]. Permeability increases again for a high filler quantity. For gels in which elementary particles are far frombeing closely packed, the filler quantity necessary to increase permeability should be lower than 40 ~01%. It isimportant to note that thesize, shape, and size distribution of pores play a very important rolein permeability [17]. Onthe other hand, permeability seems mainly governed by the connectivity of the largest mesopores, which is difficult to evaluate. Permeability also depends on the tortuosity, whichis a parameter related to the effective length of a pore. The enlargement of the pore size (if realized without washing) is often accompanied by an increase in the neck diameter between elementary particles. Aging also increases the gel connectivity [l81 and leads to the disappearance of microporosity [ 191. All these phenomena increase the gel strength and favor their monolithicity during drying. Changing the wettability of the liquid can also be used to decrease the capillary forces. Creating a water-repellent surface is a means to reach this goal. Two kinds of bonds, Si-H and Si-F, make the silica surface hydrophobic or partially hydrophobic. Precursors bearing these special functional groups can be found in the literature [20,21]. Finally, it is possible to diminish the pressure gradient caused by low gel permeability by acting on the rheological properties of the liquid filling the pores. The lower the viscosity, the lower is the pressure gradient. C02 is a good solvent obeying this condition. Furthermore, C02 exhibits a very low ~ L V value. The question of the original location of the cracks must be adressed. Cracks during drying occur from microflaws expected to be present on the surface of the gel. Regarded as a whole, the network is under compression stress. However, Scherer [22] demonstrated that at the extremity of the flaw tips the solid part remains in tension. We must keep in mind that the pore volume of a gel generally consists of three families of pores: micropores, mesopores, and macropores. If a straightforward relation between capillary forcesand the pore size is applied, the highest stresses should be associated with the micropores. However, it is not clear that failure appears at this particular location. Macroscopic stressesmay induce failure near macropores, where the mechanical strength of the network is low. During drying the largest pore size empties first. When identical dried gels were placed into liquids having different surface energies, they spontaneously crack. According to Laplace’s law, micropores, which induce the highest stresses, should crack first. Thermoporometry measurements lead to the conclusion that cracking does not induce a change in the microporous or mesoporous volumes. This result does not depend on the na-

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ture of liquid used to perform the thermoporometric measurement (water or decane) [23]. A possible explanation is offered. The gel consists of an assembly of clusters connected in such a way that macroporesseparate them. Micro- and mesopores are most probably located in the clusters, and thus stresses induced by capillary forces are higher inside clusters. However, the mechanical strength of the cluster is high enough to resist cracking. Cracks may occur at the connection between clusters. Consequently, it is speculated that micropores and mesopores in the vicinity of macropores must be avoided to increase the probability of obtaining a monolith. It is noteworthy that capillary rise rate is ‘highest for macropores. During soaking and for a given time thereafter, differential stresses, as explained before, arise between filled macropores and incompletely filled micropores. During soaking the stresses instantly reach high values, and gels crack. With such an explanation, cracks wouldbeof small extent because these differential stresses are applied to a microscopic volume. On the other hand, during soaking the region of the gel invaded by the liquid shrinks, and consequently the differential strain between the wet and dry parts leads to cracking. This effect, which is on a macroscopic scale, accounts for the observed macroscopic crack.

W. MONOLITHICCERAMICSFROMASOL-GEL Uncontrolled drying of gels leads to a powder consisting of grains. Each individual grain sinters easily, but the grains do not sinter together. These powders present by themselves interesting properties when they are modified and prepared for a special application, such as abrasives [24]. It is obviously possible to transform these powders into monolithic material by classic forming techniques. Melting [25] and hot-pressing [26] techniques are the most commonly used. Room temperature uniaxial pressing or isostatic pressing are common ceramic-forming techniques, but they are not widely used to make monolithic ceramics from gel powders, although they are successfully applied to “fumed silica” powders, which present some similar properties with gel powder [27,28]. It is more advantageous to start with dried monolithic gel to avoid contamination problems linked to the container or pressing equipment. Initially, dried monolithic aerogels are easily obtained from supercritical solvent evacuation [29,30]. This technique has been applied to several materials, such as silica, alumina [31], mullite [32], and cordierite [33]. Dried monolithic xerogels have been obtained by numerous methods, which are discussed here.

A.

Monolithic Ceramics from Aqueous Solutions

This section includes the different routes to obtain monolithic ceramics from mixtures of mineral compounds and mainly aqueous solutions. The peculiar characteristic of these processes is the presence of fine solid particles in sus-

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pension in an aqueous medium. The fine solid particles in each case are very reactive. Their origins are different. Colloidal silica (Ludox HS-40) is a stable suspension of fine silica particles having a mean size of 150-200 A. The pH of the solution is around 9.5. These particles are obtained from a water-soluble glass and then purified to remove the major part of alkali ions. Shoup developed a process in which a solution of potassium or sodium silicate (80-90 wt%) is added to the colloidal silica (10 wt%). The potassium silicate solution contains mixtures of polysilicic anions, which deposit on colloidal particles if the pH of the solution is lowered [34,35]. The increase in the pH is performed by a strongly basic organic solution, such as formamide. Gel formation is easier with sodium silicate; potassium silicate is preferred to facilitate further removal of alkali ions. Washing is performed using several baths, allowing a slow decrease in the pH of the solution. The ability of these gels to be dried without cracking is due to the narrow pore size distribution and a mean pore size above 600 A. The mean pore size may be varied by changing the weight percentage of silica sol in alkali silicate solution. The colloidal silica particles play the role of nuclei for network gel formation. Densification of the dried gel is carried out at 1350°C under a helium atmosphere. Rabinovich et al. [36,37] proposed a method that requires a double process. The starting material is a fumed silica (Cab-0-Sil) consisting of fractal silica clusters. A dispersion of fumed silica in water is prepared under acid conditions to increase the flowability of the slurry in such a way that the suspension presents a high silica content. A first gelation is then realized by increasing the temperature to 60°C. The obtained gel is the result of hydrogen bonding between species located at the surface of silica particles. This kind of gel has obviously poor mechanical properties and can be transformed into a dispersion under the action of a high shear stress. The gel is dried up to 150°C. It cracks, and consequently lumps of gels are obtained. The gel fragments are dispersed using a high-speed blender and gelled again. Drying may then be carried out without cracking. The second step enhances the permeability by introducing “man-made” macropores (5000 A) between the lumps. It is noteworthy that the specific surface area of the gel is not much affected by the first drying step. Consequently, small pores remain in the lumps. They facilitate the subsequent sintering step, but macropores interfere. It was recently reported that two parameters may be optimized to improve glass quality. The gel lumps are gently densified, increasing their mechanical strength. In addition, reducing the size of the lumps leads to small interaggregate pores. A thermal treatment at 950°C followed by a “ball milling” process for 16 h is preferred. With this procedure, the mean aggregate size is about 3 pm and the large pore size determined by mercury penetration is close to 0.1

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pm. The two-step process including calcination of the fragmented solid was also used to provide a slip for casting bodies of different shapes [38]. Fumed silica aerosil (0x50)having a relatively low specific surface area (50 m2/g) can be dispersed into a pure water solution using vigorous stimng [39]. This preferred grade of fumed silica corresponds to individual particles having a wide size distribution. For a given percentage of solid, such a suspension exhibits the lowest viscosity compared with other solutions consisting of a narrow size distribution of particles [40]. It is then possible to obtain dispersion up to 50% in weight percentage of silica particles. In such a case, only a small shrinkage is expected during the drying step. Gelation, which is very rapid, is performed by adding small quantities of a soluble fluoride compound. NH4Fseemsto give thebestresults[41]. However, theroleplayedinthis process by the F- ion is not well understood. Fluoride ions are assumed to displace hydroxyl ions at the surface of silica particles and to establish hydrogen bonds with the slightly acid hydrogen of the hydroxyl groups. However, if we remember the hydrophobic property of Si-F or Si-H, another explanation can be offered. The Si-F groups behave like Si-H, for which the chemistry is betterunderstood. The Si-H group makes the other groups sharing thesilicon atom more labile [20]. If a fluoride atom replaces a hydroxyl group, other vicinal hydroxyl groups are more reactive. Polymerization can then occur from these highly reactive hydroxyl groups; fluoride atoms are not involved in this reaction. Moreover, the presence of Si-F groups allows the preparation of gels with low adsorbed water content. This feature has been demonstrated for Si-H groups. Infrared spectra within the range 2-5 pm are not reported for gels obtained with a fluoride reagent. It should be also very important to evaluate the surface tension between the water and the fluorinated gel to appreciate the role of this parameter on the retention of monolithicity. Indeed even if the mean pore radius of OX50 gel is large compared with other silica grades, it appears very small (300 W) compared with other methods, allowing monolithicity. Permeability is probably not the limiting parameter. The Si-F bond is very stable, and it is not destroyed during sintering. Sintering treatment is advantageously carried out by zone sintering at a rate of 10 d m i n u t e . The maximum furnace temperature is between 1100 and 1500°C [42]. Peptization of mineral precipitates in a suitable pH range is a method to obtain colloidal particles in aqueous solution [43]. Gelation is sometimes possible by a slow solvent evaporation [M].Monolithic materials are obtained if the drying step is carefully controlled. Usually drying requires a very long time. Moreover, the monolithic ceramics are frequently hydroxides, and monolithicity is rarely conserved during thefurtherthermal treatments because of the chemical or crystalline transformations that take place [45].

278

B.

PHALIPPOU

Monolithic Ceramics from an Organic Solution

There are two families of monolithic gels obtained from organic solutions. The first is basically identical to those obtained in aqueous medium. Similarities come from the mineral compound, which undergoes gelation; the solvent and additives are organic compounds. The second refers to gels prepared from alkoxides in organic solvent. This kind of gel is the most popular and has been extensively reported in the literature. 1. Gels from Colloids The interest in the method developed by Scherer and Luong [46] is that it applies well to commercial "fumed" silica; moreover, it can be applied to other oxide particles [47] obtained from flame oxidation. The oxide particles are typically 10-300 nm in diameter. TiO2, Al2O3, and Si02 can be produced. A dispersion of such particles is carried out in a nonaqueous solvent. For low organic solvents y ~ is v 20-30 ergkm2 and induces low capillary forces during drying compared with those generated by water. However, the increase in the mechanical strength of the network requires a large amount of solid particles in the organic dispersion. Gelation is prevented first by coating the particles with dispersants, such as l-decanol or 1-propanol. These alcohols are adsorbed at the particle surface and create a steric barrier that hinders collision with neighboring particles. Such an approach requires careful control of submicrometer particle size, the nature of the solvent, and the nature of the dispersing agent. Moreover, the relative amounts of the different compoundsmustbe optimized [47]. For example, fumed silica is welldispersed in chloroform, but l-decanol is the most efficient dispersant. For TiO2, oleic acid is preferred as dispersant; for A1203 the best results are obtained with methanol. Gelation is performed by deprotonation of surface hydroxyl groups. NH3 or amines in very low concentration effectively induce the gelation of the nonaqueous suspension. A quaternary ammonium ion is formed. It establishes a link between neighboring negatively charged surface sites. The pore sizes of the gel are close to that of the particles and lie between 600 and 1000 A. An acceptable permeability and a low capillary force caused by the nature of the solvent allow preparation of dried gels in a short time. Some attempts to evaluate gel strength as a function of the solvent were done using a spherical indenter. Unfortunately, nothing was reported on crack appearance. However, it was reported that gels become stronger as the amount of fumed silica increases [48]. 2. Gels from Alkoxide Solutions Tetramethoxysilane, tetraethoxysilane, aluminum-sec-butoxide, and isopropyl titanate are among the best known alkoxides. They have been used for a long

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time because of their ability to form a monolithic gel produced by hydrolysis andpolycondensation reactions [49-511. The advantage of alkoxides is that chemistry can be used to modify the starting compounds to include elements that will aid in attaining the desired property. One can say that material research began to develop with the idea that molecular engineering may be induced directly or oriented from starting compounds. Ormosil and Ormocer [52] and carbon- [53,54] or nitrogen- [20,55] containing glasses are some applications of this concept. However, monolithicity remains a problem because the pore diameter of all polymeric gels is very small. Capillary stresses are high, and only a long drying treatment leads to crack-free xerogels. Using this technique, monolithic gels have been obtained. They sinter at very low temperature and exhibit physical properties identical to those obtained with high-temperature processes. The high potentialities of polymeric gels stimulated many laboratories to investigate new ways to control drying. It was disclosed that some organic compounds, including dimethylformamide [56], formamide, glycerol, and oxalic acid [14], allow shorter drying schedules. This technique was used successfully to prepare binary and ternary systems [57]. Drying, however, which depends on gel volume, requires several days. Moreover, organic compounds, especially glycerol [58], are sometimes very difficult to remove. Another approach to overcome cracking is to mixapartially hydrolyzed alkoxide solution with a suspension of fumed silica [59]. The colloidal suspension acts as a filler and reduces shrinkage during drying. Consequently, during solvent evaporation the remaining liquid enters larger pores because the shrinkage is smaller compared with that of the same gel without fillers. The mean pore diameters and pore volume of dried gels increase with the filler weight percentage (Fig. 5). Unfortunately, nothing is known of the evolution of the linear drying shrinkage versus filler concentration. Byimproving the different parameters thatplaya role in monolithicity, preparation refinements can be found. A base-catalyzed alkoxide solution leads to fine particles of silica with a specific surface area of 250-350m2/g. This powder is mixed with fluorinated silicon ethoxide solution. Both filler and fluorine allow one to obtain monolithic silica gels [60]. Gels containing fillers are now receiving more attention, and research on doped gels has extensively developed since it was demonstrated that this method allows low-loss optical fibers to be prepared 1611. Another way to preserve the monolithicity of polymeric gel is to replace alcohol with solutions having higher and higher water contents. This method achieves hydroxylation of silica surface particles; in addition, an increase in network connectivity is expected, as in aging [62]. This reaction is temperature dependent. Treatment withacidwater,whichremovestheresidual Si-OR

280

PHALIPPOU 0.20

a :90 wt?, 0.16 n

b: 65 wt ? , c:4owt%

m 0.12

d :20 wt?,

E

e : OwtV,

1 '. d)

-E g g

B

0.08 0.04

o 0

0 0 0 10

a+ b+ c +

" 100

Pore dlameter

( i)

1000

Figure 5 Textural properties (pore volume and pore diameter) of dried xerogels containing different amounts of fumed silica (aerosil 0x50).(From Toki et al. [59].)

group more effectively, allows monolithic gel to be obtained with a drying rate of 1.4 g/h. With this procedure the monolithicity arises from the strengthening of the network, not from the ' y ~ vchange. However, the time necessary to obtain large monolithic gels may be long, because both solvent exchange and isothermal treatment at 60°C require several hours to complete.

V.

SINTERINGTREATMENT

When completely dense materials are to be obtained, porous monolithic gels must be densified by increasing the temperature. Residual organic removal and crystallization of amorphous gel are well-known phenomena and may be accomplished or avoided by using specific thermal schedules. However, when an amorphous gel is transformed into a glass, a new phenomenon often occurs. During sintering and sometimes during thermal treatment carried out at high temperature, bubbles appear in the glass. The analysis of the gases in bubbles reveals principally the presence of NZ, H20, and C02 [63]. All these gases are relatively insoluble in the glass structure and do not diffuse easily. When the viscosity of the glass is low enough, the gas pres-

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281

sure inside microbubbles tends to increase and their radii to grow. Thermal treatments performed under a helium atmosphere [50,64] partially decrease the intensity of this phenomenon. It was speculated that the foaming effect is directly related to the water content of the gel [65]. Infrared (IR) spectroscopy shows that during sintering the water content is lower at the surface (hot zone) and higher at the core (Fig. 6) [66]. Because of to the closure of external pores, water is trapped and bubbles appear first at the center of the body. Moreover, the glass viscosity decreases as its water content increases. Both phenomena favor the bloating effect. To reduce this effect, it was demonstrated that a chlorination treatment is very efficient [65]. Chlorination may be carried out using different compounds (SOCl2, Cl2, and cC14). This treatment, which is easily done because of the open texture of the gel, is advantageous for two main reasons. First, the diffusion path inside the elementary particles constituting the gel network is very short. Consequently, such impurities as transition metal cations diffuse very quickly toward the surface of the particles. They are removed at the surface, where they react with chlorine [67], since most metal chlorides are gaseous at the temperature of treatment. They are eliminated in the gas (He or Cl2) flow. Hence, chlorination treatment provides a purification process. Second, as previously mentioned, the chlorine or chloride compounds decrease the water content. Silanol groups are transformed into Si-Cl groups. It is noteworthy that this newbond is notvery stable. The chlorinated gel readily reacts with atmospheric water as the gel is removed from the furnace [65]. Obviously, the sintering treatment must immediately follow the chlorination. An excessive chlorine content in the glass structure gives rise to another bloating effect when the fully dense gel is heated at temperatures near 18002000°C. In this temperature range the Si-Cl bond is no longer stable. Reboiling occurs, because of the low solubility of chlorine in silica melt. To avoid this phenomenon, flushing with a He-02 mixture is often proposed. However, using fluorine as a dehydroxylating agent is preferred. Si-F is hydrophobic, stronger than Si-Cl, and even at high temperatures fluorine is not released from the glass structure (see Refs. 1-1 1 in Ref. 68). This property is currently being investigated in the field of optical fibers because the introduction of fluoride atoms into the glass structure is a means to decrease the refractive index [69,70].

VI.

CONCLUSION

Most of the parameters related to cracking of gels are now well known. Optimization of these parameters seems to be unavoidable to control the drying stage, which appears to be the most crucial stage as far as monolithicity retention is concerned. Monolithicity will always be the result of the best compro-

PHALIPPOU

Figure 6 IR spectra of a silica glass issued from a gel. The sample is a cylinder of 22 mm in diameter and 0.15 mm in thickness. Spectrum a corresponds to analysis of glass near the surface, spectrum b to the core. (From Woignier and Phalippou [66].)

mise between drying rate, geometrical dimension, permeability, strength, and pore size distribution for a gel immersed into a given liquid. Most of the experiments performed until now have been devoted to silica or silicate systems, and it is evident that other chemical systems must be explored. However, the physical principles are the same for every system, so the previous results obtained on pure silica should allow monolithicity to be achieved.

ACKNOWLEDGMENTS I thank Elsevier Publishing and Rivista della Stazione Sperimentale del Vetro for allowing publication of the figures. Iam indebted to my collegues Gerard Orcel, George Scherer, and Thieny Woignier for helpful discussions of this work.

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33. Vesteghem, H., Di-Giampaolo, A. R., and Dauger, A., J. Mater. Sci. Lett., 6, 1187 (1987). 34. Shoup, R.D., Controlledporesilicabodiesgelledfromsilicasol-alkalisilicate mixtures, Colloid lnterfac. Sci., 3, 63 (1976). 35. Shoup, R.D., and Wein, W.J.,U. S. Patent 4,059,658(1977). 36. Rabinovich, E. M., Johnson, D.W., Jr., MacChesney, J. B., and Vogel, E. M., J. Am. Ceram., Soc., 66(10), 683, 688, 693 (1983). 37. Rabinovich, E. M., Johnson, D. W., Jr., MacChesney, J. B., and Vogel, E. M., J. Non-Cryst. Solids, 47, 435 (1982). 38. Bihumiak, P. P., Brandes, L. H., and Guile, D. L., U. S. Patent 4,042,361 (1977). 39. Clasen, R., J. Non-Cryst. Solids, 89, 335(1987). 40. Farris, R.J., Trans. Soc. Rheo., 12, 281(1968). 41. Clasen, R., Glastech. Ber., 61, 119(1988). 42. Clasen, R.,Glastech. Ber., 62, 234(1989). 43. Ramsay, J. D. F., Chem. Soc. Rev., 15, 335(1986). 44. Yoldas,B. E., J. Mater. Sci., IO, 1856(1975). 45. Assih,T.,Ayral, A., Abenoza,M.,andPhalippou,J., J. Mater. Sci.,23, 3326 (1988). 46. Scherer, G. W., and Luong, J. C., J. Non-Cryst. Solids, 63, 163(1984). 47. Scherer, G. W., in Better Ceramics Through Chemistry, MRSProceeding (C. J. Brinker, D. E. Clark, and D.R. Ulrich, eds.), Vol. 32, 1984, p. 205. 48. Bonner, F. J., Kordas, G., and Kinser, D. L.,J. Non-Cryst. Solids, 71, 361 (1985). 49. Yamane, M.,Aso, S., and Sakaino, T., J. Mater. Sci., 13, 865(1978). 50. Kawaguchi, T., Hishikura, H., Iura, J., and Kokubu, Y., J. Non-Cryst. Solids, 63, 61 (1984). 51. Matsuyama, I., Susa, K., Satoh, S., andSuganuma,T., Ceram. Bull., 63, 1408 (1984). 52. Schmidt,H., in Better Ceramics Through Chemistry, M R S Proceeding,(C. J. Brinker, D. E. Clark, and D. R. Ulrich, eds.) Elsevier, Amsterdam, Vol. 32, 1984, p. 327. 53. Baney, R.H., in Ultrastructure Processing of Ceramics, Glasses and Composites (L. L. Hench and D. R. Ulrich, eds.) Wiley, New York, 1984, p. 245. 54. Chi, F. K., Ceram. Eng. Sci. Proc., 4, 704(1983). 55. Brinker, C. J., and Haaland, D. M., J. Am. Ceram. Soc., 66, 758 (1983). 56. Adachi, T.,and Sakka, S., J. Mater. Sci., 22, 4407 (1987). 57. Orcel,G.,Ph.D.Dissertation,University of Florida, 1987. 58. Hench, L. L., in Better Ceramics Through Chemistry, M R S Proc. (C. J. Brinker and D. R. Ulrich, Eds.), Elsevier, New York, Vol. 32,1984, p. 101. 59. Toki,M.,Miyashita, S., Takeuchi, T., Kanbe, S., andKochi, A., J. Non-Cryst. Solids, 100, 479 (1988). 60. Okazaki, H., Kitagawa, T., Shibata, S., and Kimura, T., J. Non-Cryst. Solids, 116, 87 (1990). 61. Kitagawa, T., Shibata, S., and Horiguchi, M., Electron. Lett., 23, 1295(1987). S., J. Non-Cryst. Solids, 100, 236(1988). 62. Mizuno, T., Nagata, H., and Manabe, 63. Yanagisawa, O., Tanaka, C., Kokubu, Y., andSuzuki, Y., J. Non-Cryst. Solids, 38-39, 599 (1980).

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Satoh, S., Susa, K.,Matsuyama, I., and Suganuma, T., J. Non-Cryst. Solids, 55,

455 (1983). 65. Phalippou, J., Woignier, T., and Zanycki,J., in Ultrastructure Processing of Ceramics, Glasses and Composites (L.L. Henchand D. R.Ulrich,eds.),Wiley, New York, Chap. 7, 1984, p. 70. 66. Woignier, T., and Phalippou, J., Riv. Staz. Speri. Vetro, 5, 47 (1984). 67. Clasen, R., Glastech, Ber., 63, 291 (1990). 68. Rabinovitch, E. M., Krol, D. M.,Kopylov, N. A., and Gallagher, P. K.,J. Am. Ceram. Soc., 72(7), 1229 (1989). J. Non-Cryst. Solids, 100, 435 69. Tsukada,T.,Shinmei,M.,andYokokawa,T., (1988). 70. Tsukuma, K.,Yamada, N., Kondo, S., Honda, K., and Segawa, H., J. Non-Cryst. Solids, 127,191(1991).

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12 Bulk Optical Materials from Sol-Gel Edward J. A. Pope MATECH and University of Utah

Salt Lake City, Utah

1.

INTRODUCTION

In the past decade, substantial progress has been achieved in demonstrating the potential of sol-gel processing for the fabrication of large-scale, bulk optical components. Despite a wide array of technical challenges, “proof of concept” of the use of sol-gel to fabricate near net shape lenses, laser glasses, solid-state dye lasers, nonlinear optic elements, photochromic materials, andgradientindex (GRIN)lenses have allbeen demonstrated. Key technical challenges have included attaining molecular level homogeneity of dopants, gel pore size control, drying of gels without cracking, hydroxyl removal (in some cases), and sintering. Many of these issues are described in detail in other chapters. In all instances, the attainment of good optical quality materials, free from defects and manifesting high optical transparency, has been paramount to the successful demonstration of the technology. In this chapter, recent progress in several key areas is reviewed. These areas are catagorized by material classification rather than by end-use application:(1) bulk silica optics; (2) optically active doped silica glasses; (3) gel-polymer composites; (4) organically modified silicates (ormosils);and (5) gradientindex glasses. These represent the five most significant developments in the area of bulk optical materials by the sol-gel process to date.

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II. BULK SILICA OPTICS The fabrication of large silica glass elements by sol-gel has been a major ambition of many research groups in both industry and academia for well over a decade. By conventional melt casting, silica glass requires temperatures well in excess of 2000°C for processing. Casting glass at such high temperatures is very difficult, even for simple shapes. Achieving the final desired shape of the glass, such as a lens or mirror blank, requires expensive grinding and polishing steps that are labor intensive and time consuming. The net result is that silica optical components by traditional processing methods are quite expensive. The sol-gel route, on the other hand, permits near net shape casting of the silica gel at room temperature. This can typically be cast in plastic containers or machined wax molds. Assuming the gel survives the aging and drying steps, sintering the xerogel into fully dense glass can be achieved at temperatures as low as 1200"C, far lower than the temperatures required for conventional melt casting. After sintering, only light polishing is typically required to render the glass into its final desired shape. This is because the initial gel can be cast into complex shapes of exacting geometry. Thus, solgel silica should, in principle, be far more amenable to mass production than conventional melt-cast silica. Another significant advantage of sol-gel over melt casting of silica is purity. The chemical reagents employed in sol-gel processing can be purified by distillation and other means to as low as sub-ppm impurity levels, particularly with regard to metal ion contaminants. Furthermore, the open pore structure of the dried gel permits organic and hydroxyl impurities to be removed by heat treatment in oxygen and chlorine atmospheres at intermediate temperatures before sintering.Through the use ofhigh-purity starting materialsand careful heat treatment regimes, sol-gelaerived silica glass can contain impurities on the level of a few ppb. In contrast, most melt-cast silica glass typically contains impurities of the order of a few hundred ppm. Unfortunately, as with most emerging technologies, there are several serious technical challenges to be solved. The most significant of these has been the tendency of gels to crack during the drying process as a result of capillary pressure. Several different approaches to solving this problem have emerged in recent years. All strategies to reduce the tendency of gel cracking and increase monolith size have beenbasedupon one or more of the following: (1) increasing the mechanical strength of the gel; (2) controlling the pore size and size distribution of the gel network to reduce capillary stress; and (3) altering the drying conditions. One of the first successful approaches to fabricating large silica glass bodies is based upon the use of colloidal silica particles to form the initial gel [l-51. Colloidal silica particles can be purchased commercially or fabricated via the peptization of silicon alkoxides under acid conditions. Gelation can be in-

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Figure 1 High purity fused silica window (27 X 3 cm thick) by ceramic casting of gel particles. (Courtesy of Robert Shoup, Coming.)

duced by either solvent evaporation or increasing the pH. Colloidally derived silica gels tend to have large pore diameters, ranging from a few hundred angstroms to several thousand angstroms. These large pore dimensions help reduce capillary stresses during drying. Another important advantage of colloidal gels is their relatively good mechanical properties as a result of a typicallyhigh volume fraction of silica in the final wet gel [3]. Very large silica glass optical elements have been successfully prepared by this route (see Fig. 1). A potential disadvantage to colloidally derived silica glass is the difficulty in achieving true molecular level homogeneity of dopant ions, such as lanthanide laser ions. If the desired product is pure silica, this is not a problem. Sintering of colloidally derived silica gels is typically conducted between 1400 and 1700°C. Another approach to the fabrication of silica glass is to utilize silicon alkoxides as the initial precursors. Alkoxide-derived silica gels typically contain pore diameters of the order of several hundred angstroms or less. Alkoxide-derived gels also contain a lower oxide content than colloidal gels in the initial wet gel state.Alkoxide-derived silica gels therefore experience amuch greater tendency toward cracking during drying.

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GELATION 4OoC, 48 hrs. METHOD A (Standard Process)

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Figure 2 The two methods used to vary the pore characteristics of the gel-silica matrices. (Reproduced from Ref. 9.)

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One strategy to improve the mechanical properties of silica gel, increase the initial oxide content of the wet gel, and reduce gel shrinkage during drying has been to introduce “filler” into the initial gel solution [6,7]. These fillers are typically “fumed” silica powder, such as Cab-0-Sil, or alkoxide-derived particulate matter.Seiko-Epson has successfully prepared dense silica glass plates 20 x 20 x 1 cm by this process [6,7]. One key disadvantage to this approach is that the addition of filler material introduces microscopic defects in the final dense glass, such as micrometer-sized bubbles and particulates. This has limited its usefulness for such applications as photoresist masks, which require high optical quality. Substantial progress in fabricating high-purity silica glass monoliths from alkoxide-derived silica gels has been achieved through aging the gels before drying [%lo]. As illustrated in Fig. 2, aging can be conducted in the initial pore liquid or withbase-catalyzed aqueous solutions. Typically, this aging process can take up to 60 days to complete at ambient pressure under 100°C. The largest dense gel-glass samples prepared from this process are under 10 cm in diameter. The effect of aging the gel is to improve its mechanical properties by driving the polycondensation reaction of network formation toward completion. Another effect is to increase the average pore diameter. Another more aggressive method of enhancing the mechanical properties of wet gels and increasing pore diameters is hydrothermal aging (HTA) [ 11,121. In this procedure, the pore liquid in the gel is replaced with distilled water and the gel is placed inside an autoclave. The autoclave treatment greatly accelerates network polycondensation and the increase in pore diameter (see Table 1). By autoclaving, temperatures in excess of the ambient pressure boiling point of water can be achieved without drying the gel [ 1l]. By controlling the time and Table 1 DryGelPoreDiameter as a Function of Hydrothermal Aging Conditions temperature DryHTA (h)

duration HTA Sample (“C) Example 1 Example 2 Example 3 Example 4 Example 5 Example 6 Example 7 3.0Example 8

70 100 125 150 200 1581150 150

Source: From Ref. 11.

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1.o 1.o 1.o 1.o 1.o 2.0

gel pore diameter (8) 132 163 263 511 993 3138

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temperature of hydrothermal aging conditions, pore diameters can be controlled between 130 and 3000 A without changing the initial gel solution conditions. Combining HTA treatment with careful drying procedures has resulted in extremely large silica monoliths from alkoxide-gels (see Fig. 3) [13]. Alkoxide gels are typically sintered between 1150 and 1400°C. Much progress has been made during the past decade in fabricating highquality monolithic silica glass for optical applications. Taking advantage of the

Figure 3 A dry monolithicxerogel.(Courtesy CA.)

of YTC America, Inc., Camarillo,

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unique capability of sol-gel of near net shape casting, one inch silica glass lenses are now commercially available [14]. In the near future, more silica glass products based upon this technology should reach the marketplace.

111.

OPTICALLY ACTIVELANTHANIDE-DOPEDSILICA GLASSES

Lanthanide-doped silica glasses, particularly those doped with neodymium, erbium, and europium, have a wide range of potential applications as solid-state lasers and fiber-optic amplifiers. Lanthanide ions can be readily incorporated by melt casting into low softening temperature silicates, borates, phosphate, and halide glasses, but their usefulness is limited by the poor chemical durability and thermal shock resistance of these glass systems. The solubility of lanthanide ions into silica by melt casting is extremely low. The sol-gel process offers the possibility of incorporating higher dopant concentrations into silica than can be achieved by conventional technology.

A.

Doped SilicaGlass

Neodymia-doped glasses have commercial applications as high-power laser systems and fiber-optic amplifiers [15,16]. Currently available laser glasses are low softening temperature silicates, phosphates, borates, and fluoroberyllates [15]. These glasses have typically high thermal expansion coefficients, between 70 and 125 x l e 7 / "C, and poor thermal shock resistance. Silica glass, with its high glass transition temperature (1 lOOOC), low thermal expansion coefficient (5 x 10-7 / "C), and low nonlinear index of refraction (9.5 x -14 esu), is a highly attractive candidate as a neodymium host glass [17-191. Several attempts, by melt casting [15], chemical vapor deposition [16],andsol-gel [17-271, have been made to fabricate neodymium-doped silica glass. Until recently, none of these attempts have been entirely successful in achieving good fluorescense behavior for dopant levels greater than 0.5 wt%. Neodymium fluorescesat about 1.06pm in most glass and crystalline systems [16]. In silica, however, its fluorescence peaks at 1.088 pm [15-19,21,28,29]. This anomolous fluorescence behavior for neodymium in silica is caused by its unusually low coordination number of six in a pure silica host [18,19,27]. In most glasses and crystals, neodymium coordinationis seven or nine [15]. It is relatively easy to fabricate neodymium-doped silica glass with as much as 5 wt% neodymium oxide by the sol-gel process, but it is extremely difficult to obtain good lasing properties [18]. Three key criteria for fabricating a highquality laser glass. are (1) molecular level homogeneity in the distribution of the lanthanide ion; (2) low hydroxyl content; and (3) high glass transparency. Achieving all three of these criteria simultaneously has been difficult. Most sol-gel-derived silica and doped-silica glass contain several thousand ppm hy-

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droxyl impurities. This is because sol-gel is an aqueous-basedprocess and residual silanols in the dried gel have a relatively high binding energy, making them difficult to remove entirely during sintering. The conventional process for removing hydroxyl from porous silica relies upon exposing the porous gel to chlorine gas at elevated temperatures (500-9OO0C), followed by oxidation in oxygen. This method, although it reduces hydroxyl levels to below 1 ppm, results in the formation of neodymium oxychloride, NdOC1, and subsequent crystobalite formation during sintering [17-191. Another effect of chloration is a significant volatilization loss of neodymium [29]. Such samples exhibited very short fluorescence lifetimes (e10 p). Another method of hydroxyl removal is called in situ dehydroxylization, in which a dehydroxylating agent, fluorine, is incorporated into the initial gel solution [18,19,28]. Upon sintering in either helium or vacuum, the fluorine incorporated into the gel performs the same function as chlorination at elevated temperatures, removing hydroxyl impurities. Hydroxyllevels can be reduced to well below the 1 ppm level without causing devitrification or neodymium loss [17-19,281. A fluorescence lifetime for the Nd203-Si02 system in excess of 200 vs has been observed [18,19,28]. This is the highest fluorescence lifetime observed without the introduction of codopants. In Fig. 4, infrared absorption spectra are presented for dense sol-gel-derived neodymia-silica glass containing high levels of organic and hydroxyl impurities (Fig. 4a and b) and containing less than 1 ppm hydroxyl (Fig. k ) . In addition to hydroxyl impurities, clustering is also a serious problem associated with neodymium doping. One innovative solution to clustering phenomena has been codoping with aluminum [21-26,291. This has the effect of suppressing cluster formation and increasing the oxygen coordination around the neodymium [27]. The fluorescence maximum is correspondingly shifted to about 1.063 pm [29]. Most attempts employing this method have failed to reduce hydroxyl levels sufficiently to achieve fluorescence lifetimes in excess of 125 p [21-261. Recently, however, fluorescence lifetimes in the 300-400 ps range have been reported for the Ndz03-AI203-Si02 system, and lasing action has been demonstrated [28]. Muchprogresswasrecently achieved in molecular level homogeneityof neodymium doping inthesol-gel silica system. In addition, innovative new methodsof hydroxyl removal have proved quite effective. In the words of Thomas and coworkers, “substantial progress is still required to obtain material that is suitable for a practical laser system” [28].

B.

Er-Doped Silica Glass

Er3+-doped silica glass has tremendous potential as fiber amplifiers in long-distance, high bit rate optical communication systems [30]. The 1.55 pm emission of Er3+ can be pumped at 0.98 and 1.45 pm by laser diodes. Intense interest

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Wavenumbercl /cm) Figure 4 Infrared absorption spectra for (a) sol-gel-derived glass containing both organic and hydroxyl impurities; @) dense neodymia-silica glass with approximately2500 ppm hydroxyl content; and (c) dense 1 wt % neodymia-silica glass containing less than 1 ppm hydroxyl. (Reproduced from Ref. 28.)

exists in using Er-doped fibers to regenerate digital signals used for transcontinental communications and for fast data transmission over fiber-optic networks [31]. CVD-derived Er-doped silica fiber amplifiers are currently planned for deployment in transoceanic cables in the mid-1990s [31]. One limitation of current Er-doped fibers by CVD is low doping levels (about 0.01 mol%). This requires fairly long fibers for signal amplification, of the order of 500 m [32]. One method of reducing the lengthof fiber-optic amplifiers is to increase dopant concentration. The amplification factor can be increased by lo00 if up to 5% erbium is added. Recently, Sumitomo Electric demonstrated the successful fabrication of erbium-doped silica glass by the sol-gel method, with Er concentrations more than 30 times higher than those with CVD-derived fiber

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[32]. The challenges in suppressing concentration quenching, because of clustering,andhydroxyl quenching, because of silanol are similartothose encountered in neodymium-doped gel-glasses.

C.

Eus+-Dopedand W+-Doped SilicaGlass

Trivalent europium has been demonstrated to produce a bright red fluorescence, which peaks at about 615 nm in glasses. More recently, europium was successfully incorporated into gel-derived glasses [33,34]. Possible applications include solid-state lasers in the visible and fiber-optic amplifiers. The fluorescence spectra of Eus+-doped silica is presented in Fig. 5. The 615 nm peak associated with the DO + 7F2 transition state is clearly evident, along with three weaker side bands at 592,654, and 705 nm. The excitation spectra for the 615 nm fluorescence peak is typical of lanthanide elements, exhibiting relatively narrow excitation bands (see Fig. 6) [35-371. Hexavalent uranium fluoresces brilliant green, peaking between 500 and 540 nm in most glass hosts. The brilliance of uranium fluorescence far exceeds that of the most efficient organic dyes. The fluorescence emission spectrum of hexavalent uranium-doped silica gel-glass is presented inFig.7. The peak maximum is at 533 nm [351. In the excitation spectrum of europium (Fig. 6), there is a moderately strong peak at 532 nm, which is also where the fluorescence maximum of uranium exists (Fig. 7). This overlap provides a convenient opportunity for energy transfer between uranium and europium. Codoped silica gel was prepared with a Eu/U ratio of 20:l [35]. The excitation spectrum for the 615 nm europium emission is showninFig. 8 for the codoped samples. For samples withthe same europium content, the fluorescence emission is nearly doubled when sensitized by uranium at ideal pumping wavelengths [35]. At nonideal pumping

700*0

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Figure 5 Emission spectrum of europium-doped silica gel-glass.

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Figure 6 Excitationspectrum of europium-doped silica gel-glass. 120.0

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Figure 7 Emissionspectrum of uranium-doped silica gel-glass.

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Figure 8 Excitation spectrum of 615 nm emission of uranium-sensitized, europiumdoped silica gel-glass.

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frequencies, europium emission was increased by two orders of magnitude [35]. This represents the first time that energy transfer between uranium and europium has been demonstrated in any material. These same results were observed in gel-derived silica-organic polymer composites more recently [46-48].

IV. TRANSPARENTGEL-POLYMEROPTICAL COMPOSITES Employing the sol-gel route to prepare dense glasses and ceramics requires the sintering of dried porous xerogels at moderately high temperatures. Another route to nonporous monolithic materials is to fabricate gel-polymer composites. In this process, the dried porous gel is impregnated with an organic monomer that can be polymerized in situ by chemical catalysts, ultraviolet irradiation, or heating at temperatures below 100°C [38-43]. By this process, composite materials comprised of a gel-derived inorganic phase and an organic phase can be fabricated at relatively mild temperatures [38-43]. The low-temperature processing conditions required for the fabrication of these new composite materials permits the incorporation of organic molecules, such as laser dyes, as well as inorganic dopants, such as lanthanide ions [44-49]. In Fig. 9, the processing procedure for fabricating organic dye-doped gel-polymer composites is outlined. Because the composite is comprised of two phases, both of which are derived from solution, the organic dye can be incorporated in either the gel phase or the polymer phase W]. Unlike conventional composite materials, which are typically opaque, gelpolymer composites can be extremely transparent [41,42]. This is primarily the result of the extremely small phase dimensions of the gel-structure, of the order of 100 A or less. Rayleigh scattering can be less than 0.1% using 5000 A light through an optical path length of 1cm [41]. In addition to high optical transparency, many properties of the composite, such as density, refractive index, Young’smodulus, modulus of rupture strength hardness, and abrasion resistance, can be varied over a wide range by controlling the ratio of gel phase to polymer phase [41,42]. In Fig. 10, the variation in density as a function of volume fraction polymer phase is presented for silica gel-polymethyl methacrylate (PMMA) composites [41]. Density can range from 1.2 glml (pure PMMA) to 2.2 g/ml (pure silica). Refractive index can also be modulated by varying the ratio of gel to polymer content (see Fig. 11). The effect of the relative volume fraction of gel to polymer is even more dramatic for elastic modulus, which can be controlled over two orders of magnitude (see Fig.12). Thus, both optical and physical properties can be controlled over a wide range in this composite system while maintaining good optical transparency. The high-transparency, low-temperature processing requirements and good mechanical properties of transparent gel-polymer composites make them an at-

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MIX SOLUTION OF PRECURSOR, WATER, SOLVENT,& CATALYST

-ADD

DYEFOR MASS-DYED)

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IMPREGNATE WITH MONOMER I

POLYMERIZE (CATALYST, UV, OR THERMAL) TRANSPARENT POLYMER-DYE-OXIDE GEL BASED COMPOSITE

i

Figure 9 Reparation procedure for transparent gel-polymer composites (Reproduced from Ref. 44.)

tractive host material for a wide range of optically active organic molecules and inorganic ions. Thus far, laser dye molecules [44,46-49], nonlinear optic dye molecules [44,45], electrically conductive polymer molecules [46-48], and optically active lanthanide and actinide ions [46-48] have all been successfully doped into gel-polymer composites. Lasing of these materials has been demonstrated for perylene-doped silica-PMMA composites [49]. The electrooptic d.c. Kerr effect has also been measured for NLO-doped composites [45]. The most recent advance in new optical composites has been the demonstration of strong luminescence of lanthanide and acticide ions in transparent composites r46-481. In Fig. 13, the strong luminescence of Eu3+-doped silica

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Volume Fraction PMMA Figure 10 Density of transparent silica-PMMA composites content. (Reproduced from Ref. 41.)

as a function of PMMA

gel-PMMA is presented [46-48]. Uranium sensitization of europium has also been demonstrated in optical composites [46-48]. Recent advances in composite manufacturing techniques permit the fabrication of large optical elements. In Fig. 14, a wide range of organic dye-doped composites are presented. The largest of these measures 30 cm in diameter by 3.5 cm in thickness. Organic dyes represented include coumarin-314T, fluorescein, rhodamine 6G, rhodamine B, and poly@-phenylene vinylene).

V.

ORGANICALLY MODIFIED SILICATES

Another new class of materials developed in recent years is the organically modified silicates (ormosils) [50-531. As their name implies, ormosils are silicate materials that have been modified at the molecular level with an organic component. These materials are often refered to as molecular composites or nanocomposites [53]. They combine the hardness and strength of an inorganic glass with the fracture toughness of a polymer material [50]. Unlike the composites of the previous section, however, ormosils are prepared from one solution. In this regard, ormosils are more like copolymers than composites, except

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Figure 11 Refractiveindexasafunction of Ph4MA contentfortransparentsilica-. PMMA composites. (Reproduced from Ref. 41.)

0.0

0.5 Volume Fraction PMMA

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Figure 12 Elastic modulus as a function of PMh4Acontent for silica-PMMA composites. (Reproduced from Ref. 41.)

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Figure 13 Fluorescenceemissionspectrum of europium-dopedsilica-PMMAcomposite.

Figure 14 Largemonolithicdye-dopedsilica-PMMAcomposites.(Courtesy MATECH, Westlake Village, CA.)

of

that instead of copolymerizing two organic monomers, one of the monomers is an inorganic network former. The debate over whether ormosils are composites or copolymers will probably continue for sometime to come. What cannot be debated, however, are the tremendous number of potential applications of ormosils in optics [50-561. Ormosils are prepared in a manner very similar to conventional sol-gel solutions, except that in addition to the conventional alkoxide precursor, another organically modified precursor is added [50]. For example, ormosils have been

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successfully preparedby combining tetraethoxysilane (TEOS) withpolydimethylsiloxane (PDMS) plus solvent, catalyst, andwater[59]. One of the unique features of this type of process is that the properties of the ormosil can be varied, brittle to rubbery, simply by adjusting the ratio of TEOS to PDMS 1531. Ormosils can possess excellent transparency when processing conditions are optimized. Moreover, they are typically nonporous. Their properties, as well as their low-temperature processing, make them ideal candidates as optical host media for a wide range of optically active species. These include laser dyes [51,53,55], photochromic molecules [52], second-harmonic-generating dielectric oxides [54], semiconductor quantum dots [53], and metallic quantum dots [54]. Already, lasing, second-harmonic generation, andphotochromism have been successfully demonstrated in ormosils. This new class of material has a bright future in bulk optical applications.

VI.

GRADIENT-INDEXGLASSES

Gradient-index lenses are a relatively recent phenomenon. They are currently used in commercial photocopiers, Fax machines, and fiber-optic couplers. The commercially available GRIN lenses are made by an ion-exchange process in which a glass rod is immersed in a molten salt for a set period of time. The metal ion in the salt interdiffuses with a metal ion in the glass, producing a gradient in refractive index. GRIN lens made by this method tend to be less than 3 mm in diameter. The sol-gel method offers the possibility of fabricating gradient-index lenses of much larger diameter [57-68]. Already, a 13 mm diameter GRIN glass lens has been fabricated [61]. The sol-gel process for fabricating GRIN lens is a predominantly low-temperature process (see Fig. 15). First, a multicomponent gel-forming solution is mixed and cast. After gelation, the gel is leached in an acid solution to generate a composition gradient. The leaching process is halted by immersing in a “fixing” solution. The gel is then carefully dried and sintered into GRIN glass [57-68]. Several different compositions of GRIN glass have beenattempted. The mostextensively researched is the lead-potassium-borosilicate system developed by Yamane et al. [57-61]. GRIN lens with an index gradient as high as 0.05 and 13 mm in diameter have been fabricated from this system. One liability of the lead-containing GRIN glass is concern over the production of leadcontaining toxins or hazardous waste during manufacturing. Another GRIN system is the binary titania-silicate system of Konishi and coworkers [62,63]. Small 2 mm diameter lens of low ~ - have n been produced with good index profiles [62,63]. The low solubility of titanium dioxide in silica limits the maximum refractive index gradient in this binary system.

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ITnEssw

EFFECTS

1) MIX CHEMICALS

GEL STRUCTURE 8 STRENGTH POROSITY

;EACnlNGf=d BATH

2) POUR INTO MOLDS AND ALLOW TO GEL AND AGE

NETWORK

3) COMPOSITION

GEL

LEACH

FORMATION

TO

GRADIENT

SIZE OF DELTA-N ”-

PROFILE SHAPE 4 ) ARREST LEACHING

BY SOAKING GEL IN FIXING MEDIA.

FIXING

BATH

U Be

5) DRY GEL CAREFULLY

GRIN LENS PHYSICAL DIMENSIONS UNIFORM STRAIGHTNESS AND ROUNDNESS

-

FURNA

6) SINTER DRIED GEL INTO GRIN GLASS

,

&

Figure 15 Processingsteps from Ref. 67.)

in fabricatinggel-derived

GRIN glass. (Reproduced

BULK OPTICAL MATERIALS FROM SOLGEL

305

Radial Position (mm) Figure 16 Refractiveindexprofileof GRIN glass rod. (From Ref. 68.)

a 5 mm diameiertitaniumaluminosilicate

To increase the index gradient for the titanium-containing system, modifiers, such as aluminum, are added to increase titanium solubility [64-68]. Refractive index gradients as high as 0.07 have been achieved in titanium aluminosilicate GRIN glasses [66,67]. The superb refractive index profiles that can be achieved in this ternary system are exemplified by the profile presented in Fig. 16 [68]. This system, once sintered, exhibits superb stability. GRIN rods of titania aluminosilicate glass have been drawn, by fiber-drawing techniques, into smaller diameter rods without degradation of the refractive index profiles [68].

VII.

SUMMARY

A wide variety of bulk optical materials can be prepared by sol-gel processing. These include large silica optics, laser glasses, solid-state dye laser materials, nonlinear optic materials, photochromic ormosils, and GRIN glasses. In addition, novel materials containing quantum dots of semiconductors and metals can also be prepared by sol-gel. Many of the applications of these materials have yet to be explored. In the years to come, many more new materials will be invented using sol-gel, resulting in new optical components and applications.

REFERENCES 1.Shoup, R. D., Controlledpore.silicabodiesgelledfromsilicasol-alkalisilicate mixtures, Colloid Znterjk. Si., 3, 63. (1976).

306 2. Shoup, R. D., and Wein, W. J., U. S. Patent 4,059,658 (Coming Glass) (1977). 3. Scherer, G.W., and Luong, J. C., Glasses from colloids, J. Non-Cryst. Solids, 63, 163-172 (1984). 4. Rabinovich, E. M.,Johnson,D.W.,Jr.,MacChesney,J.B.,andVogel,E.M., Preparation of transparent high silica glass articles from colloidal gels, PartsI, II, and III, J. Am. Ceram. Soc., 66(10), 638-699 (1983). Ceram. Bull., 70(9), 5. Shoup,R. D., Gel-derivedfusedsilicaforlargeoptics, 1505-1510 (1991). 6. Seiko-Epson, K. K.,Japanese Patent 61186226 (1986). 7. Seiko-Epson, K. K.,Japanese Patent 61168542 (1986). 8. Wang, S. H., and Hench, L. L., Drying control additives for rapid production of large sol-gel monoliths containing transition and rare earth elements, in Science of Ceramic Chemical Process (L. L. Hench and D. R. Ulrich, eds.), (John Wiley and Sons, New York, 1986), pp. 201-207. 9. Lui, S., and Hench, L. L., Control of the texture of gel-silica monoliths by aging treatments, in Sol-Gel Optics II (J. D. Mackenzie, ed.), Proc. SPIE 1758, (1992) pp.14-25. 10. Hench, L. L., Wilson, M. J. R., Balaban, C., and Nogues, J. L., Sol-gel processingoflarge silica optics, in Ultrastructure Processing of Advanced Materials (D.R.UhlmannandD.R.Ulrich,eds.),(WileyandSons,NewYork,1992), pp.159-177. 11. Pope, E. J. A., Sano, Y., Wang, S., and Sarkar, A., Sol-gel process for glass and ceramic articles, U.S. Patent 5,023,208 (June 11, 1991). 12. Sano, Y.,Wang, S. H., Chaudhuri, R., and Sarkar, A., Silica glass monoliths from Sol-Gel Optics, U , Vol.1758 alkoxidegels;anoldgamewithnewresults,in D. R. Ulrich, eds.),SPIE,Billingham,WA,1992, (J.D.Mackenzieand pp.113-124. 13. Wang, S., Kirkbir, F., Chaudhuri, S. R., and Sarkar, A., Accelerated subcritical dryingoflargealkoxidesilicagels,in Sol-Gel Optics, II, Vol.1758(J.D. Mackenzie, ed.), SPIE, Billingham, WA, 1992, pp. 113-124. 14. Geltech sales brochure, GelSil porous glass and lenses. J. Non-Cryst. Solids, 123, 15. Weber, M. J., Science and technology of laser glass, 208 (1990). 16. Mears, R. J., et al., Neodymium-doped silica single-mode fibre lasers, Elect. Lett., 21(17), 738 (1985). 17. Pope, E. J. A., and Mackenzie, J. D., Nd-doped silica glass. I. Structural evolution in the sol-gel state, J. Non-Cryst. Solids, 106, 236 (1988). 18. Pope, E. J. A., and Mackenzie, J. D., Nd-silica laser glass, presented at the 4th Int. Conf. on Ultrastructure Processing, Tucson, M , February 22, 1989. of neodymia-silica glass, Ph.D. Dissertation, 19. Pope,E.J.A.,Sol-gelprocessing University of California, Los Angeles, 1989. Science of Chemical Ceramic Processing 20. Wang, S. H.,andHench,L.L.,in (L.L.HenchandD. R. Ulrich,eds.),WileyandSons,NewYork,1986, pp.201-207. 21. Berry, A. J.,andKing, T. A., Characterization of dopedsol-gelderivedsilica hosts for use in tunable glass lasers, J. Phys. D, Appl. Phys., 22, 1419 (1989).

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22. Moreshead, W. V., Nogues, J. L. R., and Krabill, R. H., Preparation, processing, and fluorescence characteristicsof neodymium-doped silica glass prepared by the sol-gel process, J. Non-Cryst. Solids, 121, 267 (1990). 23. Fujiyama, T., Hori, M., and Sasaki, M., Preparationof Nd-doped silica glasses by the sol-gel method, J. Non-Cryst. Solids, 121, 273 (1990). 24. Fujiyama, T.,etal.,SilicaglassdopedwithNdand A1preparedbysol-gel method, J. Non-Cryst. Solids, 135, 198 (1991). 25. Abramoff, B., et al., Synthesis of neodymium-aluminum doped silica xerogels, in 26. 27.

28. 29. 30. 31.

Roc. 5th Int. Conf. on Ultrastructure Processing, Wiley and Sons, New York, in press. Abramoff, B., et al., Preparation and characterization of Nd-AI doped glass from gels, to be submitted. Mathur, A., and Pye, L. D., EXAFS investigation of Nd3+ environment in sol-gel derived NdzOs-SiOz glasses, presented at the Symposium on Solid-state Optical Materials, 93rd AnnualMeeting of theAmericanCeramicSociety,Cincinnati, OH,April 29,1991. Pope, E. J. A., and Mackenzie, J. D., Sol-gel processingof neodymia-silica glass, J. Am. Ceramic Soc., in press. Thomas,I.M.,Payne, S. A.,andWilke, G. D.,Preparation,opticalandlaser Sol-Gel Optics, II, Roc. SPIE 1758 properties of Nd-dopedsol-gelsilica,in (J. D. Mackenzie, ed.), 1992, pp. 622629. Mita,Y.,Yoshida, T., Yagami,T.,andShionoya, S., Luminescenceandrelaxation processes in Er3+-doped glass fiber, J. Appl. Phys., 71(2), 938-941 (1992). Desurvire, E., Lightwave communications: The fifth generation, Sci. Am. January,

114-121(1992). 32. Laser Focus World, Companies making news, June, 120 (1992). 33. Devlin, K., Kelly, B. O., Tang, Z. R., McDonagh, C., and McGilp, J. F., A struc-

spectural study of the sol-gel process by optical fluorescence and decay time troscopy, J. Non-Cryst. Solids, 135,8-14(1991). O., Luminescence 34. Campostrini, R., Carturan, G., Ferrari, M., Montagna, and Pilla, of Ed+ ionsduringthermaldensification of Si02 gel, J. Mater. Res., 7(3),

745-753 (1992). 35. Pope, E. J. A., Fluorescence behavior of organic dyes, europium, and uranium in sol-gelmicrospheres,in Sol-Gel Optics, II, Proc. SPIE 1758 (J. M.Mackenzie, ed.), 1992, pp. 360-371. 36. Pope, E.J. A., Microwave processing of silica and doped silica glasses by sol-gel in Sol-Gel Optics, II, Proc. SPIE 1758 (J. M. Mackenzie, ed.), 1992, pp. 26-39. J. Am. Ceram Soc., in 37. Pope,E. J. A.,Multifunctionalsilica-gelmicrospheres,

press.

38. Pope, E. J. A., and Mackenzie, J. D., Porous and dense composites by sol-gel, in

Tailoring Multiphaseand Composite Ceramics (R.E.Tressleretal.,(ed.), (Plenum Press, New York, 1986), pp. 187-194. 39. Pope, E. J. A., and Mackenzie, J. D., Oxide-nonoxide composites by sol-gel, in Better Ceramics Through Chemistry, (C. J. Brinker et al., eds.),M R S Symposium Vol. 73, Materials Research Society, Pittsburgh, PA, 1986, pp. 809-814. 40. Pope, E. J. A., and Mackenzie, J. D., Novel composite materials for space struc-

308

41. 42.

43.

44. 45.

POPE tures and systems, in Proceedings of the 32ndInternational SAMPE Symposium, Vol. 32 (R. Carson et al., eds.), SAMPE, Anaheim, CA, 1987, pp. 760-771. Pope, E. J. A., Asami, M., and Mackenzie, J. D., Transparent silica gel-PMMA composites, J. Mat. Res. 4(4), 1018-1026 (1989). Pope, E. J. A., Asame, M., and Mackenzie, J. D., Properties of transparent silica gel-PMMA composites, in Multicomponent ultrajine microstructures (L. E. McCandishetal.eds.), M R S Symposium,Vol. 132, MaterialsResearchSociety, Pittsburgh, PA, 1989, pp. 105-1 10. Abramoff, B., and Klein, L. C., Mechanical behavior of(po1y)methylmethacrylate impregnatedsilicagels,in Ultrastructure Processing of Advanced Materials (D.R.UhlmannandD.R.Ulrich,eds.),WileyandSons,NewYork 1992, pp. 401-407. Pope, E. J. A., and Mackenzie, J. D., Incorporation of organic dyes in polymeroxide composites, Mat. Res. Soc. Bull., 12(3), 29-31 (1987). Che, T.M., Carney, R. V., Khanarian, G., Keosiuan, R. A., and Bono, M., Electro-optical dc Kerr effects and solid-state deuterium NMR studies of stable gelderivedglassorganicpolymercomposites, J. Non-Cryst. Soli&, 102,280-287

(1988). 46. Pope,E.J.A.,Sol-gelopticalnanocomposites,presentedatthePacificCoast Regular Meeting Am. Ceram Soc., San Francisco, CA, November 2, 1992. 47. Pope, E. J.A.,Transparentsol-gelnanocomposites,presentedat95thAnnual Meeting Am. Ceram. Soc., Cincinnati, OH,April 18-22,1993. 48. Pope, E. J.A.,Sol-gelopticalnanocomposites,presentedat7thInternational Workshop on Glasses and Ceramics from Gels, Paris, France, July 19-23,1993. 49. Reisfelt, R., Brusilousky, D., Egal, M., Miron, E., Burshtein, Z., and Ivri, J., Pery-

50.

51. 52. 53. 54.

55. 56.

lene dye in a composite sol-gelglass-a new solid-state tunable laser in the visible range, in French-Isreali Workshop on Solid State Lasers, Proc. SPIE 1182, (1988), pp. 230-239. Schmidt, H.,New type of non-crystalline solids between inorganic and organic materials, J. Non-Cryst. Solids, 73, 681-691 (1985). Dunn,B.,Mackenzie,J.D.,Zink,K.I.,andStafsudd, O., Solid-statetunable lasers based on dye-doped sol-gel materials, in Sol-Gel Optics, Proc. SPIE 1328 (J. D. Mackenzie and D. R. Ulrich, eds.), (1990), pp. 174-182. Yamanaka, S. A., Zink, J. I., and Dunn, B. S., Photochromism of sol-gel glasses Optics,II, Proc. SPIE 1758 containing encapsulated organic molecules, Sol-Gel in (J. D.Mackenzie, ed.), 1992, pp. 372-380. Li, C. Y.,Tseng, J. Y., Morita, K., Lechner, C. L., Hu, Y., and Mackenzie, J. D., in Sol-Gel Optics, II, Proc. SPIE 1758 (J. D. Mackenzie, ed.), 1992, pp. 410419. Cheng, C. H., Xu, Y., Mackenzie, J. D., Chee, J.K., and Liu, J., Second-harmonic generation in metal oxide/onnosil nanocomposites derived from sol-gel processing,in Sol-Gel Optics, 11, Proc.SPIE 1758 (J.D.Mackenzie,ed.), 1992, pp. 48549 1. Altman, J. C., Stone, R. E., Wishida, F., and Dunn, B. S., Dye-activated onnosils ibid in Sol-Gel Optics, 11, Proc.SPIE 1758 forlasersandopticalamplifiers, (J. D. Mackenzie, ed.), 1992, pp. 507-518. Tseng, J.'Y., Li, C. Y.,Takada, T., Lechner, C., and Mackenzie, J. D., Optical

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57. 58.

59.

60. 61. 62. 63. 64.

65.

66.

67.

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properties of metal-cluster-doped ormosil nanocomposites, in Sol-Gel Optics, II, Proc. SPIE 1758 (J. D. Mackenzie, ed.), 1992, pp..612-621. Yamane,M.,Caldwell,J.B.,andMoore, D. T., Preparation of gradient-index glass rods by the sol-gel process, J. Non-Cryst. Solids, 85, 244-246 (1986). Yamane, M., Kawazoe, H., Yasumori, A., and Takahashi,T., Gradient-index glass rods of PbO-K20- B203 -Si02 system prepared by the sol-gel process, J. NonCryst. Solids, 100, 506-510 (1988). Yamane, M., Yasumori, A., Iwasaki, M., and Hagashi, K. rod of large diameter and large delta-N, in (B. J. J. Zelinski, C. J. Brinker, D. E. Clarkl, and D. R. Ulrich, eds.), GRIN Better Ceramics Through Chemistry, W ,MRS Symp. hoc. 180, 1990, pp. 717-725. Yamane,M.,Gradient-indexmaterialsbythesol-gelprocess,in Ultrastructure Processing ofAdvanced Materials (D.R. Uhlmann and D. R. Ulrich, eds.), Wiley and Sons, New York, 1992, pp. 509-517. Yamane, M., Yasumori, A., Iwasahi, M., and Hayashe, K., Gradient index materials by the sol-gel process, in Sol-Gel Optics, Roc. SPIE 1328 (J. D. Mackenzie and D. R. Ulrich, eds.), (1990), pp. 133-144. Konishi, S., Sol-gel derived r-GRIN doped silica lenses, inSol-Gel Optics, Proc. SPIE 1328 (J. D. Mackenzie and D. R. Ulrich, eds.), 1990, pp. 160-166. Konishi, S., Shingyouchi, K., and Makashima,A.,r-GRIN glass rods prepared by the sol-gel method, J. Non-Cryst Solids, 100, 511-513 (1988). Che, T. M., Caldwell, J. B., and Mininni, R. M., Sol-Gel derived gradient index optics, in Sol-Gel Optics, Roc. SPIE1328(J. D. Mackenzie and D. R. Ulrich, eds.), 1990, pp. 145-159. Caldwell, J. B., Che, T. M., Cruse, R. W., Mininni, R. M., Nickles, R. E., Warden, V. N., and Banish, M. A., Studies on the reproducible production of GRIN glass rods by a sol-gel process, in Betfer Ceramics Through Chemistry W , (B. J. J. Zelinshi et al., eds.), M R S Proc. 180, 1990, pp. 727-732. Banish, M. A., Che, T. M., Caldwell, J. B., Mininni, R. M., Soshey, P. R., Warden, U. N., and Chin, H. H., The effect of gel aging on the physical and optical properties of a gel derived GRIN glass, Better Ceramics Through Chemistry, V M. J.Hampden-Smith,W.G.Klemperer,andC.J.Brinker,eds.),MRSSymp Roc. 271, €992, pp. 535-540. Che, T. M., Soshet, P. R., Banish, M. A., Caldwell, J. B., Javidi, M., McCallum, I., Mininni, R. M.,andWarden, V. N.,Optimization of a gel derived gradient indexmaterial,in Sol-Gel Optics, 11, hoc. SPIE 1758 (J. D. Mackenzie, ed.), 1992, pp. 193. Banish, M. A., Caldwell, J. B., Che, T. M., Mininni, R, M., Soshey, P. R., Warden, V. N., and Pope, E. J. A., Gradient index fiber optic preforms by a sol-gel method, in Submolecular Glass Chemistry and Physics (P. Bray and N. J. Keidl, eds.), Proc. SPIE 1590, 1991, pp. 8-13.

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13 Aerogel Manufacture, Structure, Properties, and Applications Jochen Fricke and Joachim Gross Physikalisches Institut der Universitat, Am Hubland Wiirzburg, Germany

1.

INTRODUCTION

Aerogels are highly porous nanostmctured materials [ 1-51. A special variety of aerogels is even transparent or translucent. As the name implies, aerogels consist mostly of air.They were made for the first time by Kistler at Stanford University in 1931 [6].At that time, Kistler undertook to dry wet gels without shrinkage. Such gels consist of a network of interconnected, entangled chains, which are embedded in a liquid. The network itself occupies a volume fraction of only a few percentage and thus conventional drying in air leads to a dramatic shrinkage of the gel. To avoid this effect, Kistler used a technique called supercritical drying. It is performed in an autoclave, a pressure vessel. The materials Kistler obtained had a spongelike open structure and a porosity of up to 99%. Kistler produced aerogels of silica, titania, alumina, and iron oxide. He also studied some of the most intriguing properties, including their low thermal conductivity [7]. However, at that time nobody seemed to be interested in the application of aerogels as new materials. This situation is completely different today. The demand for new materials is large for many reasons: insulating foams blown with CFC (chlorofluorocarbons) are known to be environmentally dangerous because the chlorine released from the CFC in the upper atmosphere destroys the ozone layer. Insulating fiber materials are considered a health hazard because they release small fiber fragments. which can be inhaled and deposited in the lung tissue, where Manuscript received August 1992

31 l

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FRICKE AND GROSS

they may cause cancer. For these and other reasons, aerogels today are a target for materials research. If we keep in mind that aerogels have extremely small pores in the 1-100 nm region (conventional porous materials have pore structures in the micrometer and millimeter range), we may anticipate that aerogels .have intriguing properties and allow surprising new applications. The nanostructure of aerogels causes much less scattering of light and sound waves as in materials with a larger pore structure. The nanostructure also causes the thermal conductivity in such materials to be extremely small. Last but not least, the high specific inner surface [2] of the nanostructure of aerogels promises interesting technical applications with respect to adsorption and catalysis. The synthesis of aerogels in a wet chemical sol-gel process allows adjustment of many parameters to optimize, aerogels for specific applications. It is possible, for example, to grow an extremely fine polymeric network or a more colloidally structured coarser skeleton [S]. The skeleton may consist of inorganic materials, of organic materials, or a mixture of both [2]. The inner surface may be hydrophilic, consisting basically of OH groups, or hydrophobic if methyl groups cover the skeleton [9]. Furthermore, a mixture of various oxides can be converted into an aerogel [lo] with extremely uniform composition. Mixed oxide aerogels can be reduced to obtain catalytically active metal surfaces on a silica surface [ 1 l]. Via sintering, the density of aerogels can be increased in a controllable way [12]. Pyrolysis of organic aerogels [l31 leads to pure carbon aerogels [14], which are electrically conductive and thus may become a new class of electronics materials. On the other hand, the production of aerogels is far from simple: it is difficult to understand and control the sol-gel process and the drying process. However, because the successful tailoring of the nanostructure has been demonstrated by several research groups, a broad variety of new materials is available and many technical applications have become feasible [2,15].

II. PRODUCTION OF AEROGELS A.

Silica Aerogels

Kistler made aerogels of water glass (sodium silicate), which he mixed with a catalyst to gelify this material. Teichner in Lyon used an organic precursor, tetramethoxysilane (TMOS), for the first time in the 1960s [16]. This last technique was employed by Poelz and Riethmuller [l71 and Henning and Svensson [ 1 S] for many years to manufacture aerogels for application in high-energy physics. It is also possible to use tetraethoxysilane in ethanol instead of TMOS and methanol [19]. This process has the advantage of using less toxic ingredients; however, reaction times are longer and the results generally are not as convincing as in the TMOS route.

313

AEROGELS

H

H

H

H

Figure 1 Model for base-catalyzedsilicaaerogel.

The gelation [20] requires tworeactions. The first is called hydrolysis, which is the detachment of methoxy (-W&) groups from TMOS and their substitution by OH groups. The next step is called condensation. Here, two molecules already containing OH groups react with each other, forming an Si0-Si bond and thereby releasing water. After more and more molecules have aggregated, a nanoscopic cluster is formed. These clusters then stick together to form larger entities, finally forming the coherent gel body(Fig.1). The structure and the density of the resulting gel can be influenced by the pH value of the starting solution, the temperature, and the concentration of the ingredients. This one-step process (all chemicals are mixed together at the same time) yields aerogels with densities between 50 and 250 kgJm3. The density range can be extended by use of a two-step process [21,22]. In the first step, a condensed silica (CS) precursor is produced by use of a substochiometric amount of water. It is stabilized by destilling off the alcoholic solvent. In the second step, the CS is diluted with acetone or acetonitrile and water with catalyst. This initiates gelation, and after drying, aerogels of 3-500 kg/m3 density result (note that the density of air is 1.3 kgJm3 at 1 bar and 2OOC) [23].

B.

OrganicAerogels

A few years ago, Pekala and the staff of the Lawrence Livermore National Laboratory (LLNL) [ 131 made organic aerogels. Pekala et al. employed organic reactions that proceed through a sol-gel transition. What is more, the interac-

FRICKE AND GROSS

314

tion of the resulting wet gel with the liquid solvent is weak enough that virtually no shrinkage occurs upon drying. The aqueous polycondensation of resorcinol with formaldehyde and of melamine with formaldehyde leads to the formation of organic gels. In principle, the reaction pathway, the nanostructure, and the properties of these new materials are close to those of their inorganic counterparts. A s in inorganic gels, the polymerization conditions, such as pH, reactant ratios, and temperature, influence the structure-property relationships typical of aerogels [24]. The polymerization of resorcinol is possible because this trifunctional monomer is capable of linking with formaldehyde in the two-, four-, and/or six-ring position. The substituted resorcinol rings condense with each other to form 3-20 nm large clusters in solution. The cluster size is controlled by the catalyst concentration in the resorcinol-formaldehyde mixture. A suitable catalyst is sodium carbonate. The clusters possess surface functional groups, such as CH2OH, which react further to form a dark red gel. Typical resorcinol to catalyst (R/C) values are 50-300. Colorless and transparent organic aerogels can be made from the aqueous polycondensation of melamine (2,4,6-triamino-s-triazine)with formaldehyde [13]. Melamine is a hexafunctional monomer capable of reacting at each of the amine hydrogens. Under basic conditions formaldehyde adds to the positions just mentioned to form hydroxymethyl (-CH20H) groups. In the second part of the polymerization the solution is acidified to promote condensation of these intermediate species, leading to gel formation. Both the resorcinol-formaldehyde (RF)and melamine-formaldehyde (MF) gels are transformed into aerogels by supercritical drying with respect to carbon dioxide. Because water is not miscible with liquid C02, the aquagels are first exchanged with acetone andthenprocessed inside atemperature-controlled pressure vessel. The critical point of C02 is low enough that no degradation takes place during drying [13]. If RF aerogels are pyrolyzed at temperatures of about 1050°C in a nonoxidizing atmosphere, they are turned into black carbon aerogels [14]. Pyrolysis leads to volume shrinkage of about 70% and a weight loss of about 50%.

C.

Drying

Kistler knew that air drying of a wet gel causes dramatic shrinkage, warping, and cracking of the monolith. The reason is that as soon as the liquid begins to evaporate from the gel, surface tensions around the gel body create concave menisci in the gel’s pores. The meniscus is the concavely shaped conture at the surface of a liquid contained in a pore or a tube. A s evaporation proceeds, the menisci retreat into the gel body, and compressive forces build up around its perimeter. Consequently the gel contracts [23]. To prevent shrinkage the gel must be dried at high temperatures and high pressures in an autoclave [25]. When the temperature and the pressure in the vessel are increased above a crit-

AEROGELS

315

ical threshold, the liquid becomes supercritical. In such a fluid every molecule can move about more or less freely, and surface tension no longer exists. If this supercritical fluid is slowly released from the autoclave, a highly porous piece of aerogel is left behind. After cooling to ambient temperature, the aerogel can be removed from the pressure vessel. The critical parameters of methanol commonly used in the production of aerogels are 240°C and about 80 bar [25]. Supercritical drying with respect to methanol is not without risk. Several years ago, the main gasket of a 3000 1 autoclave in the Swedish production facility at Lund failed, and more than 1000 l hot methanol leaked into the building. The following explosion destroyed the entire facility. A much safer method of drying gels is supercritical drying with respect to C02, which has a critical temperature of only 31°C. Besides, carbon dioxide is nonflammable. However, before drying the original solvent must be exchanged with liquid CO2 in the pressure vessel. This must be done with care not to degrade aerogel quality. Another method ofdrying gels without shrinkage is freeze drying [26]. Until now, however, only powdery gels with rather high densities have been obtained. Subcritical drying to make aerogels is also performed. Several groups worldwide are pursuing this method, which requires strengthening of the network by aging, a reduction in capillary forces by diminishing the gel-liquid interaction, and the use of a liquid with especially small surface tension. No definite results on subcritical drying have been published yet, but increasing efforts are being made to circumvent costly autoclave drying [27].

111.

CHARACTERIZATION AND PROPERTIES

A.

Structural Investigations

Before we discuss the aerogel structure itself, we give a brief survey of methods commonly used to reveal their structural features. Generally, three techniques can be distinguished. 1. Adsorption or Penetration Measurements (BET, Pycnometry,

and Porosimetry) BET adsorption and desorption is usually performed with liquid N2 and is a widespread technique for determining the specific surface area of porous materials [28]. In fact, only surfaces that are accessible to the NZ molecules are detected. The Kelvin equation correlates the curvature of the liquid surface with the applied partial pressure and pore size distributions can be derived [29]. However, this method is successful only for pore structures below about 20 m. Thus, in aerogels with typical pore sizes in the 1-100 nm range, only a fraction of the total available pore space is detected. For an aerogel with a den-

316

FRICKE AND GROSS

sity of about 100 kg/m-3 the detected pore space is of the order of 2 cm3/g instead of the expected value of about 10 cm3/g [30]. For pycnometry measurements helium is used. The sample is pressurized with helium, and a subsequent expansion of this gas into a precisely defined volume results in a pressure drop. The pressure measurements allow derivation of the skeletal volume of the sample, that is, the volume occupied by the solid phase, and thus the skeletal density can be determined [31]. Skeletal densities of silica aerogels were found to be in the range of 1700-2100 kg/m3 [32]. Mercury porosimetry is not suitable to determine the pore volume of aerogels. Since Hg is a nonwetting liquid, it penetrates the porous body only if pressure is applied. The higher the pressure, the smaller are the pores that can be filled. Because of the low compressive modulus of aerogels at low Hg pressure, only compression of the aerogel body as a whole occurs instead of penetration. Hence, pores that are larger than those present in reality are being “derived.” At elevated pressures the aerogel body is penetrated; however, its delicate structure is likely to be destroyed.

2. Direct Methods (ElectronMicroscopy) Scanning and transmission electron microscopy (SEM and TEM) techniques yield direct images of the structure. Thus, morphological features, such as particle shapes and particle arrangements, can be recognized. In principle, even estimates of particle sizes can be obtained, although the derivation of size distributions is rather tedious as usually only two-dimensional projections are available. A special problem with aerogels is the destructive interaction of the in-pinching electron beam with the tenuous aerogel skeleton. Even for small electron beam currents the deposited energy may cause local sintering. Thus TEM or SEM pictures must be interpreted with caution. A successful electron microscopy study was performed recently by Bourret [33]. 3. Indirect Methods (Scattering Techniques

and Nuclear Magnetic Resonance (NMR)) Most powerful and nondestructive are those methods that employ scattering. The specimen is irradiated either with monoenergetic photons (light and x-rays) or neutrons. The amount of scattered intensity depends on the concentration of the scattering units and the contrast between the solid phase and the voids. In general, the scattered intensity is measured as a function of scattering angle 8 or of scattering wave vector q = (4nh) sin e/2. Typical wavelengths in x-ray measurements are 0.1-0.2 nm and the structural sizes extend from 1 to about 100 nm, so that the scattered intensity must be detected at small scattering angles (0.01-loo) close to the primary beam. Thus this technique is called smallangle x-ray scattering (SAXS) [34]. A typical scattering curve for a disordered material consisting of a network of particles is shown in Fig. 2. The characteristics of the scattering pattern can be related to the different structural fea-

oms

31 7

AEROGELS continuum

particle

surface

network

cluster "R"

r

log scattering vector

Figure 2 Typical scatteringcurve foraerogels

(see text).

tures on the respective scale length. At the smallest scattering vectors corresponding to a scale length much larger than the inhomogeneities, the x-ray intensity constructively interferes, and thus a large constant scattering intensity is observed. The absolute value is determined by the number and size of the scattering units and the contrast. At increasing q values, the largest inhomogeneities are resolved. Since the phase shifts between scattered waves become larger, the intensitygradually drops. This scattering regime is called the Guinier region. The radius of gyratiou R can be estimated from the q values at which the drop occurs. For even larger q values, a more or less extended linear decay is observed in the log I versus log q plot, which can be explained in term of fractals. The scattered intensity in this range obeys a power law:

I

q-d

where d is the mass fractal dimension. The Montpellier group [35] has succeeded in making aerogels that are fractals over two orders of magnitude in scale length. Upon further increase in the scattering vector, the scale length approaches the particle diameter. A second crossover is observed from which the size of these building blocks can be estimated [36]. For even larger q values scattering becomes sensitive to the particle surface. Again a power-law decay is observed: 1

cc qDs-6

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FRICKE AND GROSS

Dsis the surface fractal dimension, which is 2 for a smooth surface; this leads

to a decay I a 4-4 and is known as Porod’s law [37]. If the surface is fractally rough, that is, 2 c D, c 3, a less steep decrease in the scattering intensity with q vector is observed. For even larger q values, the structure on the atomic and molecular levels is fiially probed. Because these subunits show some shortrange order, broadened Bragg peaks are observed. Further information from x-ray scattering can be extracted if a two-phase media model is applied. In this model, the porous body is assumed to consist of two phases with constant mass (or electron) density. The solid phase has a mass density p, that occupies a volume fraction Os,and the voids with pp = 0 occupy a volume fraction 0p= 1 - 0,. By the application of this model, the specific surface area and the mean chord length of each phase can be derived, provided an absolute intensity measurement is made [38]. NMR is used to probe the structure of sols and gels on the molecular level [39]. This is a very common analytical method for chemists and it is not further detailed here.

B. AerogelStructure Typically, the specific inner surface of aerogels is very large. Values of up to 1000 m2/g have been reported, depending mainly on the production parameters.SAXS-derived surface areas are in general about 10-50% larger than BET-derived areas, indicating a considerable amount of pores inaccessible to NZ molecules [12]. It seems that wet gels exhibit much larger surface areas than their dried counterparts as a result of partial destruction or rearrangement of the skeletal structure during supercritical drying. The skeletal densities derived for different aerogels with different methods still show a considerable data spread. Values of 1700-2100 kgm3 were reported for silica aerogels [32], whichmust be compared to p = 2200 kg/m3 for nonporous vitreous silica. Thus, a residual porosity of 5-25% for the skeletal material is implied. Base-catalyzed silica aerogels as well as RF and carbon aerogels with high WC ratios are composed of more colloidal entities (“string of pearls”), with diameters of about 4-7 nm [40] for one-step silica aerogels, 1-2 nm for two-step silica aerogels [41], and 10-20 nm for RF aerogels [42]. The pore diameters between those chains are typically 10-30 nm for the two-step silica aerogels [41] and up to 100 nm for the other species. Most of the smaller pores are probably located inside and on the surfaces of the chains. However, apart from low-density colloidal aerogels, they usually show no extended linear power-law decay of the scattering intensity. This means they are probably neither fractals nor hierarchical structures. On the other hand, neutral and acid-catalyzed silica aerogels and RF-aerogels with low WC ratio exhibit a more polymerlike microstructure with smaller chain diameters but less interconnectivity. A s a result,

319

AEROGELS

they shrink to a much higher extent during drying than their collodial counterparts. In double-log SAXS curves they show a linear decrease in intensity over a large range. Thus they may be considered fractals over up to two orders of magnitude in scale length [43].

C.

Thermal and InfraredOpticalProperties

Heat is transported in these highly porous materials in three ways: via gaseous conduction within the pores, via solid conduction along the tiny chains, and via infrared radiation [M]. Gaseous conduction in aerogels is partially suppressed because these materials have structures in the 1-100 nm range, and on the other hand, the mean free path of air molecules is about 70 nm at normal pressure and temperature [45]. Thus in aerogels, collisions of air molecules among themselves occur about as often as collisions with the pore walls. This causes a dramatic reduction in the gaseous thermal conductivity in aerogels. Typical are values of only 0.005-0.010 W/(m-K) [46] compared with 0.026 W/(m.K) for nonconvecting air in micrometer- and millimeter-sized pores of typical insulating foams (Fig. 3). The solid conduction is strongly reduced by the tenuous, highly branched skeleton. Thus in silica aerogel the solid conductivity is about 0.005 W/(mK), which is about 200 times smaller than in vitreous silica, its nonporous counterpart. The thermal infrared radiation is absorbed to some extent by the molecular building units. Although the reduction in radiative heat transfer by Si02 aerogels is sufficient below 20°C, radiation leakage increases

x

4

0.oq

-1

t0.1

1

10

100

1000

p, I mbar

Figure 3 Total thermalconductivity h of various RF aerogelsversusgaspressure ps: p = 82 kdm3 (circles); p = 157 kdm3 (triangles); p = 303 kdm3 (squares).The data were measured with the hot-wire method at room temperature.

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FRICKE AND GROSS

dramatically above room temperature. To keep radiation leakage small also for higher temperatures, we suggested doping the SiOz-skeleton with carbon black, a powerful infrared absorber [47]. Although the aerogel loses its optical transparency by this measure, its thermal insulation is strongly enhanced. Total conductivities for nonevacuated aerogels of the order of 0.014 W/(mK) are achievable. This is considerably lower thaninCFC-blown polyurethane (PU) foams,the best nonevacuated thermal insulators available today [48]. To reveal the thermal properties of aerogels, stationary hot-plate measurements are usually employed [45]. In such a measurement two equal aerogel specimens are sandwiched between a hot plate and two cold plates. If the electrical power fed into the hot plate and the temperature difference between the hot and the cold plates, as well as the thickness of the specimens, are known, the thermal conductivity can be derived. For the thermal characterization of opacified aerogels, the faster nonstationary hot-wire method can also be used. In this case a thin platinum wire is embedded into the aerogel specimen and a constant power is delivered into the wire, which also serves as a temperature sensor. From the temperature increase in the wire as a function of time, the thermal conductivity of the aerogel specimen can be determined [49]. Generally, all the measurements are accompanied by infrared optical measurements using a Fourier-transform infrared spectrometer. The spectrally transmitted infrared power is a measure of the absorption and the scattering of the aerogel and thus of the thermal resistance of such a material. Typical spectra for a nonopacified pure Si02 aerogel and a carbon black-doped Si02 aerogel are given in Fig. 4. Pure Si02 aerogel shows strong absorption above 8 pm, but thermal radiation can penetrate the aerogel skeleton easily below 8 pm. The integration of a dopant can increase absorption in this spectral region by two orders of magnitude [49]. It is worthwhile to note that Si02 aerogels are thermally stable to temperatures of about 600°C. Above this temperature sintering starts and the skeleton begins to shrink. Organic RF aerogels are stable to temperatures of about 80°C. Carbon aerogels can be used to 2000°C in nonoxidizing atmospheres. If oxygen is present carbon aerogels may be set on fire even at slightly elevated temperatures. This holds especially if a catalyst, for example platinum, is present [51].

D. MechanicalProperties The mechanical properties of aerogels are crucial for almost any application because they are the main factors influencing both the handling and machining of the material, as well as its durability in rough environments. In addition, there is fundamental interest in understanding the variation in elastic properties with density to compare experimental results with theoretical calculations on fractals and percolation clusters.

321

AEROGELS lOOOc

,

I

I

I

I

I

I

,

A

wavelength A I pm

Figure 4 Spectral specific extinction (absorption) of pure silica aerogel (dashed), silica aerogel doped with 5% carbon black (dotted), and RF aerogel (solid). Note that the “transmission window” in pure silica aerogel between 3 and 5 pm leads to a dramatic increase in thermal radiative transport and renders pure Si02 aerogels ineffective as thermal insulators above 1OOOC. The integration of an opacifier markedly improves the thermal resistance of Si02 aerogels.

1. Measurement Techniques To probe the mechanical properties of aerogels, dynamical and static methods are employed. The measurement of the sound velocityby the pulse-echo method provides relatively high precision. However, this method requires some expertise and sensitive instrumentation. This is caused by relatively high acoustic attenuation coefficients (compared with nonporous materials).Because most aerogels are destroyed by liquid contact agents, the acoustic coupling must be done with either a nonpenetrating adhesive or without a contact agent, using highly polished surfaces and sufficient mechanical pressure [52]. Typically, sound velocities are in the range from 50 to 300 m/s; this results in acoustic wavelengths of the order of 1 mm at frequencies around 1 MHz. As stated in Sec. II, the pore diameters are in the submicrometer range. This allows us to treat aerogels as homogeneous at these frequencies. Generally aerogels are also isotropic. In this case, two elastic constants (e.g., cl1 and CM) suffice to describe the elastic properties of the material. These constants are directly accessible via longitudinal and transverse sound velocities; that is, c11 =$p

c, =c;p

(1)

In the very high frequency range, Brillouin scattering may be used to measure the sound velocities of aerogels [53]. Static mechanical properties of aerogels have been measured mainly by the three-point flexure test [54] or by direct longitudinal compression [%l. The

322

FRICKE AND GROSS

main problems with these methods are the need for relatively large, crack-free samples and, in the Young's modulus experiment, defined contact between sample and pistons. Usually the Young's modulus Y or the modulus of compression K are determined. By stressing the samples until failure occurs, the modulus of rupture (MOR) may be determined [56]. Acoustic attenuation in an aerogel untilnowwas solely determined by acoustooptical techniques. The pulse-echo method proved not to be suitable for this purpose because of millimeter-sized inhomogeneities generally present in the material. At lower ultrasonic frequencies (0.5-10 MHz), Raman-Nath diffraction of laser light by plane acoustic waves was used to measure the decay of acoustic energy inside an aerogel specimen [57]. This method relies on the density variations produced by the sound wave that create a phase pattern inside the specimen. The intensity of the first order of the diffracted light beams is proportional to the sound intensity at the intercept of light and sound beam. In a log plot of sound intensity versus distance from the transducer, the slope is a measure of the attenuation. At higher frequencies, Bragg reflection of light by the sound wave is used for the same purpose [58]. Above about 0.5 GHz, Brillouin scattering provides an elegant technique of measuringphononmean free paths (which are inversely proportional to the sound attenuation coefficient) by evaluating the width of the Brillouin lines [53]. However, as a result of finite spectrometer resolution and contrast, this method is limited to aerogel densities above 180 kgjm3. 2. Sound VelocityandModuli In general, the most prominent parameter influencing the mechanical properties of aerogels is porosity l7 or density p [59]. The two quantities are correlated via the solid density p0 of the skeleton material:

For aerogels made by the same process but with different densities, elastic moduli as well as sound velocities scale with density according to modulus = pa

and

c, c,,

= pp

where a and p are scaling exponents. They are related to each other, using Eqs. (1) and (3), by

p=- a-l 2

Because the elastic properties of aerogels are defined by two independent elastic constants, there are also two independent scaling exponents a.However, it

323

AEROGELS

e / kgm-3 Figure 5 Scaling of the elastic modulus c11 of aerogels versus density p. Materials: Si02 (triangles);sintered Si02 (dots); carbon(diamonds); MF (circles); RF (squares). Open symbols denote evacuated aerogels; filled symbols denote aerogels in air.

was found that the Poisson ratio v of most aerogels is close to 0.2 and is independent of density [52]. Thus, only a single a is obtained. Typical values are in the range a = 2.5, .... 3.8 [55, 59, 601, depending on material and production process (Fig. 5). The scaling exponent of the MOR is usually considerably lower: values around 2.6 were reported [54]. In many cases, the scaling behavior of aerogels can be attributed to fractal properties of the materials [61]. A prerequisite for this explanation is that the S U S curves exhibit a constant slope over more than one order of magnitude in scale length. Even then, however, it is not trivial to indentify the parameters of theoretical fractal models with measured quantities. A concise analysis by Woignier et al. showed discrepancies and contradictions between different sets of quantities [60]. For most base-catalyzed aerogels, especially at higher densities (above 100 kgJm3). the fractal range in the scattering curves is smaller than one order of magnitude [62]. However, a scaling of elastic properties with density is still observed. A nonfractal model uses a density-dependent fraction of dangling bonds in the skeleton. For densities varying over more than one order of magnitude, this model reproduces the experimental data with realistic geometrical parameters [63]. A s already mentioned, the sound velocity is measured with much better precision than the elastic moduli. Therefore, the generally small influence of some other parameters on the acoustic properties of aerogels is experimentally accessible. It is striking that in aerogels of different materials, the sound velocity was found to be lowered by uniaxial compression [52]. In normal solid ma-

324

FRICKE AND GROSS

terials, the opposite effect is observed at very high stresses. Because the Young's modulus of aerogels is some four orders of magnitude smaller than of nonporous solids, this elastic nonlinearity is rather easilymeasured. Even though fused silica exhibits an effect with the same sign, it has now been confirmed by measurements on carbon and organic aerogels that the nonlinearity is caused by bending deformation of the nanostructural entities of the skeleton [W.

Higher density aerogels have elastic moduli well above 10 MPa, but for lower densities elastic moduli become small enough to allow the atmospheric air pressure to become noticeable. The gas inside the aerogel pores influences both the total (dynamic) modulus and, at very low aerogel densities, the total density of the system. According to a simple model (Kelvin-Voigt model), the sound velocity becomes

The index 0 indicates properties of the evacuated aerogel. K is an effective adiabatic exponent for the system gas-skeleton. For most aerogels it is very close to unity. Only at extremely low aerogel densities does it become both larger and pressure dependent [65]. In Fig. 6, the variation in Ci with gas pressure is depicted for air and SF6 for p = 5 kdm3 aerogel. 3. Ultrasonic Attenuation and Creeping

Onlyavery few measurements on the acoustic attenuation of aerogels are available as yet. They all deal with transparent silica aerogels for the experimental reasons already discussed. At very high frequencies in the gigahertz range, Brillouin scattering allows the detection of a crossover from propagating to localized modes (phonon-fracton crossover) [47]. The crossover frequency corresponds to sound wavelengths equal to typical pore sizes of the aerogel skeleton. At lower frequencies, at which the phonon mean free path is much larger than the wavelength, pure acoustic propagation of ordinary sound waves is expected (and observed). This may also be expressed by stating that the attenuation per wavelength is very small compared with unity. Bragg scattering of laser light by plane acoustic waves yielded a linear increase in attenuation versus frequency in the range 50-90 MHz [58]. In the lower megahertz range, by Raman-Nath diffraction a similar frequency dependence was found. In addition, the attenuation was observed to decrease up to a factor of 7 in vacuum [57]. Thus, the ultrasonic attenuation in silica aerogels is to a major extent caused by water adsorbed at the inner surface. The frequency dependence of ultrasonic attenuation as stated yields a constant loss tangent (tan 6), which is the inverse of the so-called quality factor Q.

325

AEROGELS

I

2o0

200

.

.

400 Pga8

600

800

1000

/ mbar

Figure 6 Sound velocity ci as a function of gas pressure pg= in a two-step aerogel (p=5.3 kdm3): evacuation of air (open triangles); venting with SF6 (filled triangles). The solid curves are fit to the data according to Eq. (5). The shape of the SF6-curve is

influenced by the high specific heat and density of this gas. These quantities are also accessible through creep measurements, yielding information on the attenuation, or dissipation of elastic energy, at very low frequencies 1661. Measurements of this type are in progress. Preliminary results indicate that the loss tangent of silica aerogels may be essentially independent of frequency over more than 10 orders of magnitude.

E.

Optical Properties

Pure Si02 aerogels have an airy appearance and contain only small optical imperfections. They gleam either bluish if viewed against a dark background or yellowish if held against a white light source. This effect is known as Rayleigh scattering, which also renders the sky above us blue and tints the sunset yellow. Rayleigh scattering is caused by the nanostructure, the extremely small inhomogeneities in the aerogel skeleton. Optical characterization of aerogels is commonly performed using a scattering apparatus in which a laser beam interacts with the specimen and a detector is rotated around the irradiated volume section. Such measurement allows us to derive structural information. Structural units comparable to the wavelength of the laser light cause a strongly forward peaked scattering pattern [67]; entities which are much smaller than the wavelength cause more or less isotropic scattering. In addition to these scattering experiments, directional hemispherical measurements are also performed

FRICKE AND GROSS

326

inwhich the total scattered intensity is detected via an integrating sphere arrangement. Such measurements are important if aerogels are to be employed in solar energy applications.Aswe have discovered in a systematic study, aerogels madewithpH 13 show the highest optical transparency [68]. The solar transmission for a variety of aerogels versus density is shown in Fig. 7. At 630 nm the specific extinction e of the best aerogels is of the order of 0.1 m2kg [69]. Thus an aerogel with a density p = 100 kg/m3 has a photon mean free path at that wavelength of = (ep)" = O.lm

(6)

Because of their low density, aerogels have a refractive index very close to 1. Experimentally one finds that

n - 1= 2.1 x 104p

kg/m3

(7)

in agreement witha simple theoretical estimation [17]. For p = 100 kg/m3, one obtains n = 1.02. The numerical factor in this equation depends on the number of OH groups on the inner surface. For pure silica one expects 0.9 x 10-4. As a result of the small refractive index, light enters and leaves a piece of aerogel almost without reflective losses and refractive effects. The dependence

L

60 -

E

I

1

40 -

~ o o " " t " 100 " " " t "

200

300 g I kgm-3

Figure 7 Directional-hemispherical solar transmission tdh as a function of density p for different aerogels: granular aerogel (diamond); low-density aerogel tiles from LLNL (L. W. Hrubesh, squares); aerogel tiles from DESY, Hamburg(G. Poelz, filled circles); specimens made by P. Wang (University of Wiirzburg) at pH 13 (open circles) and pH 11 (asterisks). Thickness of all specimens was 20 mm.

327

AEROGELS

of the index of refraction on the density is often used to probe the local density variation in aerogel monoliths. Most employed is the minimal deflection method at 90" comers to determine absolute values of n. Small variations are detected very sensitively by interferometric methods [70].

W . APPLICATIONS All applications of aerogels make use of their high porosity, which is responsible for the low index of refraction, the small Young's modulus, the low acoustic impedance, the low thermal conductivity, and the excellent accessibility of the inner surface. In addition, in some applications the high optical transparency is of importance. One can recognize in principle three types of applications for aerogels. In the first type (high-technology applications), only very small quantities of aerogels are necessary and thus the price of these materials is of minor importance. The second type includes the mass applications with quantities of thousands and tens of thousands cubic meters. In this case the price per liter or kilogram aerogels must be as low as possible. There is a third type of application in which the quantities needed are of the order of cubic meters; however, the price can be high because the potential user is capable of paying for rather expensive aerogel tiles and no cheaper alternatives exist. The last application is typical of the use of aerogels in Cherenkov detectors used in the high-energy physics community with a rather generous budget, at least in the past. On the other hand, the demand for relatively large high-quality aerogel tiles promoted the efforts to scale up the production of these materials [17,18]. The small index of refraction of aerogels allows determination of the momentum of relativistic particles within a momentum range that is covered neither by compressed gases nor by liquids. The in-pinching relativistic particle produces an electromagnetic shock wave, the cone angle of which depends on the ratio of particle velocity and light velocity within the medium. If sintered aerogels are included, a variation in the index of refraction n between 8 x 10-3 < n - 1 0.4 is possible. Several high-energy research groups have used aerogel Cherenkov detectors in storage rings and accelerators or for the investigation of cosmic rays [71].

A. Aerogels

as Thermal Insulators

1. AerogelWindows One of the most promising applications of Si@ aerogels is their use as transparent thermal insulation. Extensive studies of transparent aerogels have shown that the thermal resistance as well as the solar transmission are sufficiently high to make these materials most suitable for passive solar usage [72]. However,

328

FRICKE AND GROSS

even aerogel windows for conventional housing applications can prove to be economical if they can be produced cheaply enough and, on the other hand, energy prices are not too low to encourage energy savings at the cost of higher investments. Convincing examples of the use of granular silica aerogels as translucent insulations are in a two-family house in Ardon, Switzerland and a house recently built in Freiburg, Germany [73]. The energy consumption for heating purposes in such houses is exceptionally low, typically 300 1 oil per year in a MiddleEuropean climate. In such applications an aerogel layer sandwiched between two glass panes is installed in front of a massive house wall, painted black on the outside. The solar radiation penetrates the aerogel insulation and is absorbed by the surface of the wall. Because of the excellent thermal insulation of the aerogel layer, most of the heat generated at the wall surface is usable as heating energy in the house; only a small amount of this heat is lost to the environment (Fig. 8). Other interesting applications for aerogels are in "frosted" windows [74] or insulating covers for solar panels [75].

2. Opaque Aerogels as Substitutes for CFC-Blown Insulating Foams Most promising is the application of aerogels as opaque insulations in refrigerators or heat-storage systems. Until now CFC-blown polyurethane foams have been employed. The CFC must be abandoned for environmental reasons, however, because they are responsible for the destruction of the ozone layer in the upper atmosphere and contribute considerably to the greenhouse effect. shade

n

Figure 8 Passive solar heatingwith a translucentaerogelinsulation in front of a blackened house wall. Even on cold but sunny days heat gain can be achieved.

329

AEROGELS

-- 0.04

:l 2

0.03

x

:zc D 3 0

0.02 0.01 0

- 50

0

50 150

4 100 200 temperature I O C

Figure 9 Thermal conductivity h as a function of temperature for a CFC-blown PU foam (triangles), an opacified Si02 aerogelpowder(squares),anopacifiedmonolithic Si02 aerogel (circles), and a monolithic RF aerogel (exes). The solid lines axe guides to the eye.

Typical thermal conductivities for PU insulations freshly blownwith Freon (F1 1) are h = 0.02 W/(mK) [76] at 300 K. Such low conductivities are achievable because F11 has a thermal conductivity of only 0.008 W/mK compared with nonconvecting air at h = 0.026 W/(m-K). Because PU foams have a combined solid and radiative conductivity of about 0.008 W/(m.K), air-filled PU foams have totalthermal conductivities above 0.03 W/(mK).The need for nonhazardous, noninflammable thermal insulants with conductivities in the 0.02 W/(m-K) range or below is obvious. The substitution of PU foams by Si02 aerogels is possible, if the infrared radiative heat transfer through the highly porous material can be reduced sufficiently. A s is shown in Fig. 9, at 300 K nonevacuated powders and monolithic opaque aerogels have thermal conductivities of about 0.018 and 0.014 W/(m-K), respectively. Evacuation reduces these values to below 0.009 W/(mK). The best values for aerogel powders at 300 K are 0.002 W/(m-K). Thus these materials can be considered most suitable for application in high-performance insulations, for example in a NaS battery or in other high-technology insulations for which large dwell times are required.

B.

AcousticImpedance Matching

Very similar to the corresponding optical case, sound waves crossing a plane interface between two materials are partially reflected. However, in the optical

330

FRICKE AND GROSS

analogy the reflected power is determined by the refractive index of the two materials, which varies only between 1.0 and about 3.0 The analogous acoustic quantity, the acoustic impedance Z = pc (c = sound velocity), spans an interval of more than five orders of magnitude (e.g., Z = 400 kd(m2.s) for air versus lo7 kd(m2.s) for ceramics). This reduces usable acoustic energy and thereby signal-to-noise ratio in applications in which the sound waves are generated and detected in piezoelectric ceramics but propagate in gases (e.g., in acoustic ranging applications). Similar to the optical analogy, the problem is solved by acoustic quarterwave matching layers (Fig. 10). Silica aerogels with densities of around 300 kdm3 have the ideal acoustic impedance to match a piezoelectric transducer to air [77]. In addition, they exhibit rather low attenuation, as opposed to many porous, polymer materials usedfor this purpose until now. An increase in sound transmittance by more than 30 dB was achieved in a relatively simple arrange ment without optimization. More elaborate designs, eventually including multiple layers combining different materials, will probably result in transducer systems optimized with respect to output power, sensitivity, and bandwidth.

C.

CarbonAerogels asSupercapacitors

For this application carbon aerogels with as high a specific inner surface as possible are used. The carbon aerogel specimens are cut into about 1 mm thick

-

__c

__c

ultrasound waves piezo ceramic

aerogel

air

zP

za

ZO

Figure 10 Aerogelquarterwaveimpedancematchinglayer

for a piezoceramic ultrasound transducer. For maximized output power the acoustic impedance of the aerogel layer must be Z, = (Zpa)ln.Typical thickness d is around 0.1 mm forf = 1 MHz.

AEROGELS

331

wafers using a diamond blade dicing saw. The aerogel wafers are then immersed in an aqueous 4 M KOH solution for at least 1 day to allow the KOH solution to diffuse into the pores. Two aerogel wafers of identical composition, diameter, and thickness are pressed together with two nonporous carbon cylinders as electrical contacts inside a Teflon cylinder. The two aerogel wafers are separated by an electrolyte-wet microporous glass separator. The whole arrangement is pressed together to achieve low-resistance contacts. If a voltage is applied to this structure, an electrochemical double layer develops at the inner surface of the carbon aerogel, which has an extremely high capacity. Typical values are about 3 x 104 F k g carbon. The energy density is up to 25 W g and the power density up to 10 kW/kg, which is in excess of all other capacitors and batteries. Charge-decharge cycling of well over 4000 has been demonstrated in such devices. It seems realistic to retain well above 85% of their original capacity after lo5 cycles [78].

D. AerogelsinElectronicCircuits Because of their low density, Si02 aerogels also have low dielectric constants E of the order of 1.1 compared with 2.1 for massive nonporous silica or 9 for alumina. In high-frequency electronic circuits, the dielectric constant E governs the speed of signal propagation: the higher E, the lower the transfer rate. Thus low E are especially important in complex and fast computers. Instead of using alumina or polyimide substrates, aerogel layers could considerably enhance the performance of hybrid circuits. However, the problem of providing a heat sink for the active components has yet to be solved, because aerogels are known to be excellent thermal insulators.

E.

OtherApplications

The nanostructure of aerogels could be used for gas-filtering purposes in the range from 20 to 100 nm [79]. In contrast to other materials for this purpose, because of their high porosity aerogels allow a relatively high gas flow. The pore size distribution can be influenced via drying control chemical additives W]. The inner surface of aerogels is more accessible to gas molecules than that of xerogels, with much lower porosity. Thus, aerogels are much better suited as catalysts or as carriers for catalysts [81,82]. This important application was one of the main reasons for the comeback of aerogels at the end of the 1960s. Another application that benefits from the small pores and large inner surface of aerogels are photo-luminescent light sources [83]. Until now, in these devices a radioactive gas (e.g., tritium) was let into a glass tube coated with a phosphor at the inner surface. Integrating the phosphor in an aerogel of low density increases both the intensity and total output of the system.

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FRICKE AND GROSS

Micrometeorites traveling at high speeds in space are of interest to astrophysicists. With an “aerogel trap,” they can be decelerated and finally stopped without being destroyed [84]. The aerogel trap also offers low weight in relatively large volume. Because of its transparency, both the direction and speed of the particle can be determined and the particle may be easily located and examined in situ. Last but not least, aerogels may serve as target wick materials in laser fusion. L120/B203 [ M ] and organic aerogels 1861 were made for this purpose, which exhibit both low density and excellent homogeneity.

ACKNOWLEDGMENTS The aerogel research at the University Wurzburg was supported by the German Bundesministerium f i r Forschung und Technologie (BMFT) in Bonn. Specimens were provided by L. W. Hrubesh and R. W. Pekala (LLNL), S. Hennig (Airglass, Staffanstorp), and G. Poelz (DESY, Hamburg).

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2nd Int. Symposium Aerogels (Revue de Physique AppliquCe, Colloque C4), Edi-

tions de Physique, Les Ulis, France, (1989). 5. Fricke, J. (guest ed.), Proceedings of the 3rd Int. Symposium Aerogels, J. NonCryst. Solids, 145 (1992). 6. Kistler, S. S., Nature, 127, 741 (1931). 7. Kistler, S. S., and Caldwell, A. G., Ind. Eng. Chem., 26, 658 (1934). 8. Schaefer, D.W., in Proceedings of the 2nd Int. Symposium Aerogels (Revue de

J. Plalippou, J. Pelous,and Physique AppliquCe, ColloqueC4)(R.Vacher, T. Woignier, eds.), Editions de Physique, Les Ulispou France 1989 p. 121. 9. Schwertfeger, F., Glaubitt, W., and Schubert, U.,J . Non-Cryst. Solids, 145, 85

(1992). 10. Heinrich, T., Raether, F., Tappert, W., and Fricke, J., J . Non-Cryst. Solids, 145, 55 (1992). ’ 11. Gardes, G.,Pajonk,G.,andTeichner, S. J., Bull. Soc. Chim. Fr., 9-10, 1327 (1976). 12. Emmerling, A., Gross, J., Gerlach, R., Goswin, R., Reichenauer, G., Fricke, J., and Haubold, H.-G., J. Non-Cryst. Solids, 125, 230 (1990). 13. Pekala, R. W., Alviso, C. T., and LeMay, J. D., Organic Aerogels: A New Type

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14.

15. 16. 17. 18. 19. 20. 21.

22. 23.

24. 25. 26.

27.

28. 29. 30. 31. 32. 33. 34. 35.

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of Ultrastructured Polymer, 5th Ultrastructure Proceeding Conference, Orlando, Florida, February 1991. Pekala, R. W., and Alviso, C. T., Carbon Aerogels and Xerogels,in Novel Forms of Carbon, Renschler, C. L., Pouch, J. J., and Cox, D.M. (eds.), MRS Symp. Proc., 270, 3 (1992). Fricke, J., and Emmerling, A., J . Am. Ceram. Soc., 75,2027 (1992). Teichner, S. J., in Aerogels, Springer Proceedings in Physics, Vol. 6 (J. Fricke, ed.), Springer-Verlag, Heidelberg, 1986, p. 22. Poelz, G., and Riethmuller, R., Nucl. Instr. Meth., 195, 491 (1982). Henning, S., and Svensson, L., Phys. Scr., 23, 697 (1981). Russo, R., and Hunt, A. J., J . Non-Cryst. Solids, 86, 219 (1986). Brinker, C.J., andScherer, G. W., Sol-Gel Science, Academic Press, Boston, 1990. Schaefer, D.W., Keefer, K. D., Ashley, C. S., Pearson, R. K., and Thomas, I. M., in Physics and Chemistry of Porous Media I1 (J.R. Banavar, J. Koplik,and K. W. Winkler, eds.), AIP, New York, (1987), p. 63. Hrubesh, L. W., Tillotson, T. M., and Poco, J. F., 5th Conf. Ultrastructure Processing, Orlando, Florida, February 17-21, 1991. Scherer, G . W., in Sol-Gel Science and Technology (M.A. Aegerter, Jaffelicci, Jr.,M., Souza, D. F. Zanotto, E. D., eds.), World Scientific, Singapore, (1988), p. 181. LeMay, J. D., Mat. Res. Soc. Symp. Proc., 207, 21 (1991). Henning, S., in Aerogels, Springer Proceedings in Physics, Vol.6 (J. Fricke, ed.), Springer-Verlag, Heidelberg, 1986, p. 38. Egeberg, E. D., and Engell, J., in Proceedings of the 2nd Int. Symposium Aerogels (RevuedePhysiqueAppliquCe,Colloque C4) (R.Vacher,J.Phalippou, J. Pelous, and T. Woignier, eds.), Editions de Physique, Les Ulis, France, 1989, p. 23. Einarsrud, M.A., Farbrodt, L. E., and Haereid, S., A New Silica Xerogel Material with low Density, talk given at 3rd Int. Symp. Aerogels, Wunburg, Germany, September 30 to October 2, 1991. Brunauer, S., Emmett, P. H., and Teller, E., J. Am. Chem. Soc., 60, 309 (1938). Gregg, S. J., and Sing, S. W., Adsorption, Surface Area and Porosity, Academic Press, New York, 1982. Schuck, G., Dietrich, W., andFricke,J., in Aerogels, SpringerProceedingsin Physics, Vol. 6 (J. Fricke, ed.), Springer-Verlag, Heidelberg, 1986, p. 148. Woignier, T., and Phalippou J., J . Non-Cryst, Solids, 93, 17 (1987). Ayral, A., Phalippou, J., and Woignier, T., J. Mater, Sci., 27, 1166 (1992). Bourett, A., Europhys. Lett. 6, 731 (1988). Glatter, O., and Kratky, G . (eds.), Small Angle X-ray Scattering, Academic Press, London, 1982. Vacher, R., Woignier, T., Phalippou, J., Pelous, J., and Courtens, E., in Proceedings of the 2nd Int. Symposium Aerogels (Revue de Physique Appliquk, Colloque C4) (R.Vacher, J. Phalippou, J. Pelous,and T. Woignier,eds.),Editions de Physique, Les Ulis, France, 1989, p. 127.

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36. Schaefer, D. W., Mater. Res. Soc. Symp. Proc., 79, 47 (1987). 37. Porod, G., Kolloidz. 124, 83 (1951). 38. Porod, G.. in Small Angle X-ray Scattering (0. Glatter and G. Kratky, eds.), Academic Press, London, 1982,p. 17. 39. Devreux, F., Boliot, J. P., Chaput, F., and Lecomte, A., Mater. Res. Soc. Symp. Proc., 180, 211 (1990). 40. Emmerling,A.,Gerlach,R.,Goswin,R., Gross, J., Reichenauer, G., Fricke, J., and Haubold, H. G., J . Appl. Cryst., 24, 768 (1991). 41. Hrubesh, L. W., Thomas, T. M., and Poco, J. F., Mater. Res. Soc. Symp. Proc., 180, 315 (1990). 42. Pekala, R. W., Alviso, C. T., LeMay, J. D., in Chemical Processing of Advanced Materials, Hench, L. L., West, J. K. (eds.) John Wiley, New York, 1992. p. 671. 43. Schaefer. D. W., in Proceedings of the 2nd Int. Symposium Aerogels (Revue de Physique Appliqute, Colloque C4) (R. Vacker, J. Phalippou, J. Pelous, and T. Woignies, eds.), Editions de Physique, Les Ulis, France, 1989, p. 121. 44. Fricke, J., in Aerogels, Springer Proceedings in Physics, Vol. 6 (J. Fricke, ed.), Springer-Verlag, Heidelberg, 1986, p. 94. 45. Buttner, D., Caps, R., Heinemann, U., Hummer, E., Kadur, A., Scheuerpflug, P., and Fricke, J., in Aerogels, Springer Proceedings in Physics, Vol. 6 (J. Fricke, ed.), Springer-Verlag, Heidelberg, 1986, p. 104. 46. Fricke, J., Hummer, E., Morper, H.-J., and Scheuerpflug, P., inProceedings ofthe 2nd Int. Symposium Aerogels (Revue de Physique AppliquBe, Colloque C4)

(R.

Vacher, J. Phalippou, J. Pelous, and T. Woignies, eds.), Editions de Physique,Les Ulis, France, 1989. p. 87. 47. Kuhn, J., Lu, X., Arduini-Schuster, M. C., and Fricke, J., Carbon Black as Infrared Opacijier and Thermal Insulant Carbon '92, Int. Conf. on Carbon, Essen, Germany, June 22-26, 1992. 48. Fricke, J., Arduini-Schuster, M. C., Biittner, D., Ebert, H.-P., Heinemann, U., HetJ., andLu, X.. Thermal Conductivity, 21, 235 fleisch,J.,Hummer,E.,Kuhn,

( l 990). 49. Anderson, P., and BLkstram, G., Rev. Sci. Instr., 17, 205 (1976). 50. Lu, X.,Wang,P.,Arduini-Schuster, M. C.,Kuhn,J.,Buttner, D., Nilsson, O., Heinemann, U., and Fricke, J., J. Non-Cryst. Solids, 145, 207 (1992). 51. Eberle, W., private communication 52. Gross, J., Reichenauer, G., and Fricke, J., J . Phys. D,21, 1447 (1988). 53. Courtens, E., Pelous, J., Phalippou, J., Vacher, R., and Woignier, T., Phys. Rev. Lett., 58, 128 (1987). 54. Woignier, T., Phalippou, J., Hdach, H., and Scherer, G . W., Mat. Res. Soc. Symp. Proc., 180, 1087 (1990). 55. Pekala, R. W., Hrubesh; L. W., Tillotson, T.M., Alviso, C. T., Poco, J. F., and LeMay, J. D., Mat. Res. Soc. Symp. Proc., 207, 197 (1991). 56. Woignier, T., and Phalippou, J., J . Non-Cryst. Solids, 100, 404 (1988). 57. Altmann,H.,Schlief, T., Gross, J.,andFricke,J.,in Ultrasonics International '91, Conference Proc., Butterworth Heinemann, Oxford, (1991), p. 261.

AEROGELS 58. 59. 60. 61. 62.

63.

64. 65. 66. 67.

68. 69. 70. 71. 72. 73. 74. 75. 76. 77. 78.

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Nouailhas,B.,Michard, F., Gohier, R., andZarembowitch, A.. Proc. 11th Int. Cong. on Acoustics, Paris, (1983), p. 179. Gross, J., and Fricke, J., J. Non-Cryst. Solids, 145, 217 (1992). Woignier, T., Pelous, J., Phalippou, J., Vacher, R., and Courtens, E., J . Non-Cryst. Solids, 95/96,1197(1987). Vacher, R., Woignier, T., Phalippou, J., Pelous, J., and Courtens, E., in [4], 127 Boukenter, A., Champagnon, D., Dumas, J., Duval, E., Quinson, J. F., Rousset, J. L., Serughetti, J., Etienne, S., and Mai, C., in Proceedings of the 2nd Int. Symposium Aerogels (Revue de ,Physique Appliquk, Colloque C4) (R.Vacher,J. Phalippou, J. Pelous,and T. Woignier,eds.)Editions de Physique, Les Ulis, France, 1989, p. 133. Gross, J., Ph.D. Thesis, University of Wiirzburg, Germany, (1992). Gross, J., and Fricke, J., Phys. Rev. B, 45, 12776 (1992). Gross, J., and Fricke, J., J . Acoust. Soc. Am., 91, 2004 (1992). Kjartansson, E., J. Geophys. Res., 8 4 , 4737 (1979). Beck, A., Gelsen, O., Wang, P.,and Fricke, J., in Proceedings ofthe2nd Int. Symposium Aerogels (Revue de Physique AppliquCe, ColloqueC4) (R. Vacher, J. Phalippou, J. Pelous,and T. Woignier,eds.)EditionsdePhysique, Les Ulis, France, 1989, p. 203. Beck, A., Wang, P., Komer, W., and Fricke, J., Components of Transparent InsulationSystems,2ndWorldRenewableEnergyCongress,Reading,England, September 13-18,1992. Beck, A., Caps, R., and Fricke, J., J . Phys. D,22, 730 (1989). Hrubesh, L. W., and Alviso, C. T., Mat. Res. Soc. Symp. Proc., 121, 703 (1988). Rasmussen, J. L., in Proceedings of the 2nd Int. Symposium Aerogels (Revue de Physique AppliquCe, Colloque C4) (R. Vacher, J. Phalippou, J. Pelous, and T. Woignier, eds.) Editions de Physique, Les Ulis, France, 1989, p. 221. Goetzberger, A.,and Withver, V.,in 3, Aerogels, SpringerProceedings in Physics, Vol. 6 (J. Fricke, ed.), Springer-Verlag, Heidelberg, 1986, p. 84. Goetzberger, A., Transparent Isolation Technologyfor Solar Energy Conversion, 2nd Ed., Fraunhofer-Institut fur Solare Energiesysteme, Freiburg, Germany,1991. Wittwer, V.,in J . Non-Cryst. Solids 145, 233 (1992). Svendsen, S., in J . Non-Cryst. Solids 145, 240 (1992). Kuhn, J., Ebert, H.-P., Arduini-Schuster, M. C., Biitmer, D., and Fricke, J., Znt. J . Heat Mass Transfer, 35, 1795 (1992). Gerlach, R., Krauss, O., and Fricke, J., in Ultrasonics International '91, Conference Proc., Butterworth Heinemann, Oxford, (1991), p. 323. Mayer, S. T., Pekala, R. W., and Kaschmitter, J. L., J . Electrochem. Soc., 140(2),

446 (1993). 79. Cooper, D. W., Sci. Technol., 7, 371 (1989). 80. Hench, L. L., in Ultrastructure Processing of Glasses, Ceramics and Composites (L. L. Hench and D. R. Ulrich, eds.), John Wiley, New York, (1984), p. 52. 81. Pajonk, G.M., Applied Catal., 72, 217 (1991). 82. Hoang-Van, C., Pommier, B., Harivololona, R., and Pichat, P., in [5], 250

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83. Ashley, C. S., Reed, S. T., Brinker, C. J., and Walko, R. J., Aerogel Composites for Radioluminescent LightlPower Sources, 5th Int. Conf. Ultrastructure Processing, Orlando, Florida, February 1991. 84. Hrubesh, L. W., Report UCRL-2134, Lawrence Livermore National Laboratory, Livermore, California, 1989. 85. Brinker, C . J., Ward, K. J., Keefer, K. D., Holupka, E., Bray, P. J., in Pearson, R. K.,in Aerogels, SpringerProceedingsinPhysics,Vol.6 (J. Fricke,ed.), Springer-Verlag, Heidelberg, 1986, p. 57. 86. Hair, L. M.,Pekala, R. W., Stone, R. E., Chen, L., and Buckley, S. R., .I. Vac. Sci. Technol., A6, 2559 (1988).

14 Fractal Growth Model of Gelation Edward J. A. Pope MATECH Westlake Village, California and University of Utah Salt Lake City, Utah

1.

INTRODUCTION

A fractal is not a specific object or material but a paradigm for describing the morphologyofa seemingly random shape or growth process.Many things, from snowflakes to electrical discharge patterns, can be described by fractal concepts [l-31. In this work, the evolution of gels from solution employing fractal growth concepts is presented. In classifying the polymerization of monomeric units in solution to form fractal species, there are two extremes to the same spectrum; bounded by reaction-limited growth, the diffusion rate of the monomer is rapid relative to the polymerization rate. In aggregative processes, the “sticking coeffkient” is low. This process tends to produce particles of relatively uniform density throughout with fractally rough surfaces. These are called surface fractals [4,5]. Applicable growth models include the Eden growth model, poisoned Eden growth, and percolation theory [5]. A possible example in silica gels is the growth of “colloidal silica particles” under conditions of cluster nucleation followed by aggregation, described by Bogush and Zukoski [6]. In their system, monomer evolution and cluster nucleation was slow relative to the aggregation process, resulting in particles of fairly uniform cross-sectional density and size. Experimental evidence suggests that a simple diffusion-limited growth model is inapplicable. 337

POPE

338

In diffusion-limited growth, diffusion is slow relative to polymerization. The fractal species produced have a morphology in which the center of the fractal has the highest density, with a sharp decrease in density with increasing radius. These are termed mass fractals [S]. Applicable models include percolation theory under appropriate boundary conditions and the fractal growth model described here. Examples may include I-F-catalyzed silica gels, in which the reaction rate is rapid, such as those prepared by Pope, Mackenzie, and Rabinovitch and Wood [7]. The “clustering of clusters” with a high sticking coefficient is another example of a diffusion-limited gelation process [ 8 ] . A noteworthy exception to the preceding generalizations is reaction-limited growth in which hydrolysis of the monomer is incomplete [4]. In such a scenario, the unhydrolyzed sites are “poisoned,” or nonreactive, because of the presence of large alkoxy ligands, thereby leading to a more “ramified” structure. The conditions of reaction-limited growth and diffusion-limited growth are summarized in Fig. 1. Also presented are the characteristic morphologies of the evolving particles.

PARTICLE MORPHOLOGY

GROWTHMECHANISMCONDITIONS

Reaction Limited Growth

a) Reaction rate low relative to diffusion rate. Low sticking coefficient. Typicallylong gelation time. Surface Fractals or Colloids b) Reaction rate slow relative to ditfusion rate,but hydrolysis of monomer incomplete.

Mass Fractals

DiffusionLimited cally Coefficient. Growth

Reaction ratelast relative to sticking Diffusion High rate.

short Gelation time. Mass Fractals

Figure 1 Summary of conditions arising from the gelation processes.

FRACTAL GROWTH MODEL OF GELATION

339

II. FRACTAL DENSITY The growth in mass or molecular weight of a fractal object follows the relationship

in which the mass M is proportional to the fractal radius R to the power factor D, the fractal dimension [l-5,B-1 l]. For a material of uniform density, the fractal dimension is 3. In other words, the mass of the object is proportional to its radius cubed. The density of a fractal, however, decreases with increasing radius. This decrease in density can be calculated. By recognizing that a nonfractal object increases in mass by the relation

m = p. (t).r; and the fractal increases in mass by

m = p( +-D which when set equal to one another yield

where F is the relative fractal density at radius r. The m and p0 are the core radius and core density, respectively. The core acts mathematically as a reference point for calculating the decrease in density as the fractalincreases in size. The core can also represent a physical element in the fractal construction, such as one Si04 tetrahedron, which is the epicenter of the fractal’s radius of gyration for a silica system. In Fig. 2, the relative fractal density is plotted as a function of increasing radius for selected values of fractal dimension. A fractal dimension of 1.0 represents linear chain growth, whereas a fractaldimension of 3.0 represents three-dimensional nonfractal growth. True fractal growth is bounded by these two extremes in fractal dimension. A s can be seen in Fig. 2, the decrease in density can be severe. For example, a fractal growing with a dimension of 2.0 has a relative density of 0.25 its core density at only 4 times its core radius. Fractal dimensions of approximately 2.0 havebeen observed for silica gels evolved under both mildly acid and mildly basic conditions [5,12].

340

POPE D=3

i3D growth)

l-

G

z W

C3

W

CC J

U I-

o a a L

0

J

:

l

:. 2 3

4

5

6

(chain growth) I : 8 9 1 0

;.

I

7

RELATIVERADIUS(r/ro)

Figure 2 Fractal relative density versus relative radius as a function of variations in fractal dimension.

111.

CRITICALFRACTAL RADIUS

As a result of the rapid decrease in density as a function of increasing fractal radius, the fractal density may become so low as to be non-self-supporting. A s a consequence, there is be a limiting, or critical, fractal radius, r* beyond which particle growth is not favored or greatly retarded. Correspondingly, there is a critical fractal density E* that defines this radius, such that

where r* depends not only upon F* but also upon the fractal dimension. Clearly, increasing the fractal dimension, which indicates increased cross-linking, increases the critical fractal radius. In Fig. 3, the change in critical fractal radius is plotted against increasing fractal dimension for a critical density of 0.1. The relative density is a measure of the volume element bounded by rg and r* Although r* represents the critical fractal radius, the relative density is not zero at that point, inasmuch as it approached zero asymptotically. Therefore,

341

FRACTAL GROWTH MODEL OF GELATION

a J

a l-

o 4

a

L

-l

6

0 t a

o

’ 0 .o 1

1.5

l

I

2.0

2.5

1

FRACTAL DIMENSION (D)

Figure 3 Calculated critical fractal radius versus fractal dimension for fractal relative density.

a 0.1 critical

the critical fractal density should be interpreted as corresponding to the critical (minimum) density of the fractal volume element bounded by rg and r*.

W. MOLECULAR WEIGHT AND FRACTAL NUMBER DENSITY

To utilize fractal concepts to quantify meaningful parameters in the evolution

of gel structure, certain fundamental concepts must be established. For example, if the mass concentration of monomer Q is constant, as when a finite quantity of precursor is added to solution and hydrolyzed, then the number of particles is inversely proportional to their average molecular weight, such that

N=-NOQ M where N = number of fractal particles and M = average molecular weight.

342

POPE

It has also been established by Cannel and Aubert that for silica gels the molecular weight increases exponentially with time [13]. Their experimental results show that

M = ~,,e{qmt}

(7)

where MO = initial molecular weightof monomer, Qm = molecular weight growth factor, and t = time. By combining Eqs. (7) and (6), the change in N with respect to time is quantified by

suchthat the number of particles should decrease inverse exponentially. It should be noted,however,that these relations are applicable to systems in which the monomer has been fully hydrolyzed at the onset of reaction and the average molecular weight includes contributions of monomeric units as well as the larger particles that evolve. These relations are not accurate for systems in which the monomer is either hydrolyzing slowly or is being added slowly to a solution throughout the gelation process.

V.

VISCOSITY RELATIONSHIP

A simple relationship was derived by Einstein [l41 to calculate the viscosity of

a suspension of fine particles in a fluid, such that 71 = 710(1+ LC)

(9)

where qo= viscosity of fluid, C = volume concentration of particles, and L = a geometrical shape constant. The Einstein equation applies for dilute systems in which the volume concentration of the dispersed phase is less than 0.10, which is typically the case with gel solutions. For spherical particles, the geometrical constant L is 2.5. The geometrical constant differs for oval, whiskershaped, or other geometries [15-171. It is assumed that until gelation occurs, the fractals and even large fractal aggregates represent a discrete particulate phase in solution of an unknown shape factor. Unlike hard, noninteractive particles, however, shear rate affects viscosity as a result of bond breakage of the large fractal aggregates during viscosity measurement [18]. Hence, the viscosity values of this derivation should be considered the “true” viscosity of solution with negligible shear rate effects. Equation (9) can be rendered into aphysically more meaningful form through several substitutions. The volume concentration of the particle phase

FRACTAL GROWTH MODEL OF GELATION

343

is equivalent to the number of particles per unit volume multiplied by their average particle volume, such that (10)

C=NVp where

vp=-

M

%Pp

and p p = the average density of a particle. From Eq. (4), the density of a growing fractal is a function of its radius and fractal dimension, and hence Eq. (9) becomes

where qs= reduced viscosity. From Eq. (6) we already know that L, N,M, and Q are related. Thus Q can be solved:

Q=

[rlsPo('o/')3-D] r

which can be further rearranged to

The physical significance of this last relationship is that, as the fractals grow larger, they become less dense. Even though the'mass concentration of oxide precursor is constant, the volume concentration of fractal particles is increasing because of the decreasing density of the fractals themselves! Cannel and Aubert demonstrated that the particle radius increases exponentially by

where r is the radius exponential rate constant. Through the substitution of Eq. (5) into (4), a simple exponential relationship for reduced viscosity is obtained, such that qs=-eQL

Po

(3-D)qrr

POPE

344

Thus,

A graph of In qs versus time should yield a straight line for a fractal growth regime.

VI.

LIGHT SCAlTERlNG

The relative intensity of light traversing through a porous gel is given by the relation

where I refers to the pore radius, n is the index of refraction, 5 is the wavelength of light, and V, is the volume fraction of particles. By substituting Eqs. (l),(4), (7), and (1 1) into Equation (18)and applying the appropriate manipulations, the resulting relation for light scattering as a function of time is

which can be reduced to

in which the scattering cross section S is simply expressed as S = Be[qm+ ( 6 - ~ k r ] t

(21)

Thus, the relative intensity of light traversing a gelling system would follow an inverse double-exponential function of time. Unlike viscosity, in which only the size of the clusters evolving from a solution plays a significant role, both cluster size and molecular weight exhibit a profound effect on light scattering.

VII.

REFRACTIVE INDEX

The refractive index for a transparent two-phase mixture can be calculated based upon the net polarizabilities of the two constituent phases, as follows:

FRACTAL GROWTH MODEL OF GELATION

345

which reduces to

where nu is the index of refraction of the solution, nb is the index of refraction of the evolving cluster, and V, is the volume fraction of clusters in the solution. Substituting Eqs. (1) and (7) into Eq. (1 1)and then substituting into Eq. (23) yield.

which can be further simplified to

n2 -1 -

”

+

+(3-D)9,1,

n2+2

Thus, the index of refraction of a gelling system increases exponentially as a function of time and the fractal dimension.

VIII.

OVERVIEW

Three time-dependent property relationships have thus far been derived for a gelation system. These are for viscosity, light scattering, and refractive index. It should be noted that these relationships have been derived to the gel point. Past gelation, other processes related to the “aging,” or postgelation reactions, occur that are not treated by these formalisms. The three relationships are summarized in Table 1.

Table 1 Time-Dependent Property Relationships

rate property Growth Constant Equation Viscosity

q - 1=

Scattering cross section

A = - QL

Po

q0

S = Beqd

POPE

346

When comparing the relationships in Table 1, several similarities become apparent. First, all three properties behave exponentially with time to the gel point. In all three cases, both the fractal dimension of the gelation process and the radial growth rate factor, which relates to cluster size, play a key role. For the scattering cross section and refractive index, however, the molecular weight growth factor also appears to influence the time dependence strongly. This may be easily understood when considering that rheology is a surface-dominated property in which the effective cluster size is critical. In optical properties, the mass of material through which the light traverses is also important.

IX. STRUCTURALBOUNDARYCONDITIONS IN GEL-DRIVED SYSTEMS In the evolution of solids from solution, a wide spectrum of structures can be formed. In Fig. 4, a simple schematic representation of the structural boundary condition for gel formation is presented. At one extreme of the conditions, linear or "nearly linear" polymeric networks are formed. For these systems, the functionality of polymerization$ is nearly 2. This means there is little branching or cross-linking. The degree of cross-linking is nearly 0. In silica, gels of this type can be readily formed by catalysis with HCl or HNO3 under conditions of low water content (less than 4 mol water to 1 mol silicon alkoxide). The "ideal" fractal dimension for such a linear chain structure is 1. The pheMinimum

"Polymeric" "Colloidal" Cluster Linear Nearly Network Growth f-2 f-4 DC 0 D - l D-3

-

In Between

"Fractal" Growth & Aggregation 2
Figure 4 Structural boundaryconditions for gel formation.

Maximum

Uniform Particles Formation DC

- 1.0

FRACTAL GROWTH MODEL OF GELATION

347

nomenon of chain entanglements, however, usually results in an experimentally measured value greater than 1, by small-angle x-ray scattering and small-angle neutron scattering. At the other end of the spectrum, fine colloidal particles are formed. These particles typically have relatively uniform densities from center to outer surface. The fractal dimension D for this type of growth approaches 3. The formation of colloidal particles is essentially nonfractal. Dilute aqueous solutions under base catalysis are the most common means of forming such particles from alkoxides. The functionality ofpolymerization is nearly 4 under these conditions. The degree of cross-linking within the colloidal particles approaches unity. Most gels, however, form structures between these two extremes. The hydrolysis and polycondensation of metal alkoxide precursors yield an initial formation of "fractal" clusters, which upon reaching a critical size begin to form a continuous network or skeleton via cluster-cluster aggregation. Both the initial cluster formation and the subsequent aggregation of clusters are fractal in nature. There is a very distinct difference between these two processes, however, as is more evident in Sec. X.

X.

HF-CATALYZEDSILICA: AN EXAMPLE SYSTEM

The fractal growth model of gelation described in the preceding sections has been applied to silica gels made from solutions catalyzed by hydrofluoric acid. If, in fact, gel formation is fractal in nature and the preceding derivations are indeed correct, then gel structure and proper information should be obtainable through the measurement of viscosity and light scattering versus time. In this study, silica gels were prepared from a standard formula composed of 21 ml tetraethoxysilane, 22 ml ethanol, 6.8 ml distilled water, and X g hydrofluoric acid (50% concentrated). The values of X were 0.1,0.125,0.15, 0.175, 0.2, 0.3, 0.375, 0.5, and 0.6 g. Viscosity was measured using a Brookfield model DV-11 digital viscometer. Light scattering was measured with an Aerotech LSlOR HeNe laser (25 mW) light source and an Aerotech Model 71 laser radiometer using an LF1 100-3/6 detector. Both measurements were plotted simultaneously on a Brookfield Model 1202-000 dual-pen chart recorder. Fourier-transform infrared (FTIR)measurements were made usingaPerkinElmer 1720 FT-IR. For all samples in the HF-catalyzed silica gel series, viscosity and light scattering were measured simultaneously. In Fig. 5, the natural logarithm of reduced viscosity is plotted versus time for the silica gel catalyzed by 0.2 g HF. The shape of the curve in Fig. 5 is similar to that of the curves plotted for the other samples in the series. Only the slopes of each region and the location of the transition differ from sample to sample. Based upon previous work, region

POPE

348

Transition

5 41

Figure 5 Naturallogarithm alyzed silica gel.

of

I

reducedviscosity plotted versus time for W-cat-

I has been determined as the cluster growth regime and region I1 is the cluster-cluster aggregation regime. In Fig. 6, the light-scattering data for the silica gel sample catalyzed with 0.5 g HF is presented. Once again, the curves for all samples were quite similar in their basic shape, the only significant differences being the slopes and positioning of the transition intercepts. Unlike viscosity, however, the light-scattering data can be collected after gelation. In Fig. 6, the intercept point between region I1 and the postgelation regime coincides exactly with the measured gel point. Within a short period of time after gelation, the transmitted light intensity begins to asymptotically approach a steady-state value. In Fig. 5, two regimes, one for cluster growth and another for cluster aggregation, were identified. Associated with each regime is a slope, which corresponds to a growth rate factor (3 - D)Qr. In Fig. 7, the ratio of the cluster aggregation rate factor to the cluster growth rate is plotted versus HF catalyst concentration for the silica series. As catalyst concentration increases, the ratio of these two growth rates decreases to a catalyst concentration of 0.5 and then increases slightly at a catalyst level of 0.6. In Fig. 6, it was pointed out that the transmission intensity approaches a steady-state value. It was found by following the change in light scattering over several days time that the steady-state value is reached well before 24 h after gelation. In Fig. 8 , the steady-state value of relative transmission intensity is plotted for the entire silica series as a function of catalyst concentration.

FRACTAL GROWTH MODEL OF GELATION

1.48,

*

n h

v

1.46.-

‘ r U

1.44.

5

1.42..

v

Region I Cluster Growth

1.401.38.1.36 12

13

I I I I I

I

14

Region I1 Cluster Aggregation

2\ I

I

I

15 17

349

\

Post Gelation Growth

1

16

18

19

20

I 21

Time(minutes)

Figure 6 Doublelogarithm of light transmission plottedversustimefor alyzed silica gel.

HF-cat-

When the two curves presented in Figs. 7 and 8 are compared, it quickly becomes apparent that they are somehow inversely related. In Fig. 9, the optical intensity data from Fig. 8 are plotted versus the ratio of the cluster-cluster aggregation rate to the cluster growth rate (CCAWCGR) from Fig. 7. The relationship between the relative rates of cluster aggregation to cluster growth and the light scattering of the gels appears to be linear. The light-scattering values represent an experimentally measured value, but the ratio CCAWCGR is derived completely from the assumptions of the fractal growth model described previously and summarized in Table 1. In an attempt to explain the strong correlation observed between the light scattering of our gel samples and the calculated ratio between cluster aggregation rate and cluster growth rate, estimated particle diameters were calculated for each sample based upon the method outlined earlier, in which the particle size can be estimated from the cluster growth rate (region I) and the intercept between regions I and 11. The results of these calculations are presented in Fig. 10, in which estimated particle diameter is plotted versus catalyst concentration. The shape of this curve corresponds well to the light-scattering data, indicating that the particle sizes estimated from viscosity data can be used to predict the light scattering of the gel samples. To understand better the mechanisms of gel formation and, ultimately, to begin to “design” gel structures, viscosity and light scattering as a function of

POPE

350

Ol

0.1

0.2

0.3 0.6 0.4 HF Concentration(gm/SOcc)

i

0.5

Figure 7 Ratio of cluster-clusteraggregation rate to cluster growth rate (CCAR/ CGR) as a function of HF concentration.

0.0

0.7

h

g

2-

0.6

0.5

0.4

c

0.2

0.3 0.4 HF Concentration(gm/SOcc) Path Length=4.2cm

0.5

1

0.6

Figure 8 Optical transmission of gel samples as a function of HF concentration.

FRACTAL GROWTH MODEL OF GELATION

351

2.2-

2.0

.'

-

\

h

0

v

1.8-

1.6= 0.085 R

-

1.66

/ 1 . 2 1 0

/ , 10

20 30 R(CCAR/CGR)

40

Figure 9 Plot of inversetransmissivity versus ratio of cluster-clusteraggregation rate to cluster growth rate (CCAWCGR).

Figure 10 Estimatedparticlediameter as a function of HF concentration.

I

50

352

8

0

rr

POPE

FRACTAL GROWTH MODEL OF GELATION

353

354

POPE

time for a HF-catalyzed silica gel system have been systematically examined. A strong correlation has been observed between the structural information obtained from viscosity data and light scattering. From the viscosity versus time data collected for each system, the molecular weight and cluster size can be calculated as a function of time to the gel point using the equations derived in the first few sections of this chapter. From these calculations, it is possible to model the evolution of the gel structure dynamically. In Fig. 11, the evolution of an HF-catalyzed gel as a function of time is presented, starting with the unhydrolyzedmonomer molecule tetraethoxysilane at time zero and proceeding, in stages, to the final gel structure at 16 minutes. Also presented is the effect of normal drying, in which the fractal characteristics of the individual clusters are all but eliminated. These types of models and measurements provide an invaluable insight into the structural evolution of gel networks. From measurements of viscosity and light-scattering behavior, sufficient understanding is gained to begin to tailor gel structures for specific applications.

REFERENCES 1. Mandelbrot,B., Fractals, Form,Chance,and Dimension, Freeman,SanFrancisco, 1977. 2. Stanley, H. E., and Ostrowsky, N. (eds.), On Growth and Form, Nijhoff, Boston, 1986. 3. Ohrbach, R., Science, 231, 814 (1986). 4. Keefer, K. D.,in Better Ceramics Through Chemistry 11, MRS Symp. Proc., Vol. 73, M R S , Pittsburgh, PA, 1986, p. 295. 5. Scheafer, D. W., and Keefer, K. D.,in Better Ceramics Through Chemistry II, MRS Symp. Proc., Vol. 73, MRS, Pittsburgh, PA, 1986, p. 277. 6. Bogush, G.H., and Zukoski, C. F., IV, in Ultrastructure Processing of Advanced Ceramics, Wiley, New York, 1988, p. 477. L., in Better Ceramics Through Chemistry II, 7. Rabinovich, E. M., and Wood, D. M R S Symp. Proc., Vol. 73, MRS, Pittsburgh, PA, 1986, p. 251. 8. Kolb, M., Botet, R., Julien, R., and Henmann, H. J., in On Growth andForm (H.E. Stardey and N. Ostrowsky, eds.), Nijhoff, Boston, 1986, p. 222. 9. Langer, J. S., Rev. Mod. Phys., 52, p. 1(1980). 10. Whitten, T. A., and Sander, L. M., Phys. Rev. Lett., 47, 1400 (1981). 11. Meakin, P., Phys. Rev., A27, 604, p. 1495 (1983). 12. Schaefer, D.W., and Keefer, K. D., Phys. Rev. Lett., 53, 1383 (1984). 13. Cannel, D. S., and Aubert, C., in On Growth and Form (H.E. Stardey and N. Ostrowsky, eds.), Nijhoff, Boston, 1986, p. 187. 14. Einstein, A., Ann. Phys., 19, 289 (1905);34, 591 (1911). 15. Hermans, J. J., Flow Properties of Disperse Systems, North-Holland, Amsterdam, 1953.

FRACTAL GROWTH MODEL OF GELATION

355

16.Sobotha, Z., Rheology of Materials andEngineering Structures, Elsevier,Amsterdam,1984. 17. Kao, D.T.Y., in Handbook of Fluids in Motion, Ann Arbor Science, Ann Arbor, MI, 1983. 18. Rabinivich, E. M., and Kopylov, N. J., in Ultrastructure Processing of Advanced Ceramics (G.A. Bogush and C. F. Zukoski, N,eds.), Wiley, New York, 1988, p. 281. 19.Yasumori, A., Yamane,M.,andKawaguchi, T., in Ultrastructure Processing of Advanced Ceramics (G.H. Bogush and C. F. Zukoski, N,eds.), Wiley,New York, 1988, p. 355. 20. Pope, E. J. A., andMackenzie, J. D.,J. Non-Cryst. Solids, 101, 198; 106,242 (1988).

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v CERAMICS VIA POLYMER CHEMISTRY

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15 Nonoxide Ceramics via Polymer Chemistry Kenneth E. Gonsalves and Tongsan D. Xiao University of Connecticut Storrs, Connecticut

1.

INTRODUCTION

There is considerable interest in the synthesis of ceramic powders and in the fabrication of ceramic fibers for the development of high-temperature and highstrength nonoxide ceramic composites that exhibit resistance to oxidation, corrosion, and thermal shock [l]. The development of innovative synthesis and processing technologies required for thefabrication of components of such composites is therefore of major interest, because traditional methods, such as sintering, often cannot readily be applied to high-temperature structural materials to produce complex shapes, especially in fiber form. The major disadvantage for the synthesis and fabrication of high-strength nonoxide ceramic materials by conventional methods is that they often require high temperatures (>150O0C) and pressures [2].For example, Si3N4 was prepared by direct nitridation of elemental Si or by the reaction of Sic14 with NH3, and S i c powders were prepared by the reaction of sand with coke in a high-temperature reactor [2,3]. Similarly,inthe preparation of most ceramic coatings, conventional chemical vapor deposition (CVD) is employed. A CVD reaction requires a significantly high temperature to break the molecular bonds of the reactants and to combine them into desired products. Inthesynthesis ofSi3N4 coatings [4-lo], for example, hazardous or corrosive reactants, such as SiF4-NH3, SiH4NH3, and S i C l W 3 , are used to prepare amorphous or a-phase Si3N4.Although effective and reliable, the CVD process is extremely slow, wasteful of

359

360

GONSALVES AND H A 0

expensive and hazardous reactants, and limited to materials that can withstand the high temperatures involved. Although plasma-enhanced CVD systems can be operated at a temperature as low as 3Oo0C, processing parameters, such as radio frequency (RF)power, W, electrode spacing, total pressure,and substrate heating, are all interrelated and difficult to control. Also, the kinetics of the reaction rate is slow at such low temperatures and the pressure at which one can maintain a discharge is restricted. One potential solution to these problems involves the use of organometallic polymers as precursors to the desired nonoxide ceramic products. FYocessing ceramics by the pyrolysis of organometallic polymers is thus very versatile: the limitations associated with other conventional processing techniques, such as melting and sintering, can be eliminated [ll]. Technologically important ceramics with novel compositions and superior structuresare feasible by this method. Successful polymer pyrolysis requires identifying a suitable polymer, devising an efficient synthesis route with control of yield, molecular weight, and purity, and developing pyrolysis conditions that yield a practical quantity and quality of the desired ceramic end product. Ceramic yield is typically measured by the weight ofthe ceramic product as a percentage of the starting polymer weight. The quality of the ceramic product is expected to be a function of polymer composition and structure and processing conditions. Quality is defined by the composition of the product, the amount and character of the voids and cracks, the existence or absence of grain structure, and crystallinity. The shaping of polymers into forms that subsequently can be conserved during pyrolysis is a technology that has been extensively employed in the manufacturing of graphite fibers and carbon-carbon composites. The art of producing complexly shaped ceramics through the organometallic polymeric precursor route involves the synthesis of precursors that can be shaped byemploying production: scale operations followed by pyrolysis to produce the desired properties within the shape [12]. In the production scale, the chemical route required to convert the preceramic monomers or polymers to the final ceramic products should also be relatively uncomplicated and not overly expensive. Toxic and hazardous chemicals should be avoided, as should hazardous operations. More specific requirements of a preceramic polymer are as follows: It should be soluble andor fusible, so that it can be processed by conventional polymer processing procedures. It has the appropriate rheological properties for the application in question. Its pyrolysis gives a high ceramic residue yield (>70%). Its elemental composition is such that its pyrolysis gives a ceramic of an acceptable chemical composition. The ceramic obtained in its pyrolysis has the desired microstructures for the application in question.

CERAMICS NONOXIDE

VIA POLYMERS

361

Historically, this method of synthesis of potential preceramic polymers has passed three stagesof development with quite different goals.It began with the efforts of Stock and Somieski in 1912 for the preparation and classification of general properties of polysilazane [13-151. The potential for the commercial use of polysilazane in the 1950s and 1960s promoted studies on their synthesis.Interestin organometallic polymer precursors (preceramics) emergedin 1975 with the work of Yajima et al. [16-171, who found that poly(dimethy1silylene) could be used as a precursor to Sic. Research in this field, however, grew slowly at first after this reawakening, but very rapidly in the last decade. Withtheirefforts, Seyferth et al.[18,19] demonstrated thepossibilitythat under controlled pyrolysis conditions, polysilazanes can lead to ceramics of appropriate physical and chemical properties. Analogous work in the synthesis and processing of other nonoxide ceramics, aswellas silicon-based ceramic materials, has been extensively reported in the UnitedStates, Europe, and Japan. The physical and chemical properties of organometallic (preceramic) polymers [20] depend greatly upon the nature of the organic group bonded to the metal species. The molecular weight of the organometallic polymer depends upon the exact method of synthesis as well as the starting materials. In the pioneering work of Seyferth et al. [18,191, organosilazane precursors were prepared by the ammonolysis of dichlorosilane toyield a polysilazane oil, [H2SiNHIx. This oil, on thermal pyrolysis up to 1150"C, produced a crystalline solid ceramic product of a-Si3N4 and P-Si3N4-mixture.The precursor gives a very clean end product but undergoes rapid hydrolysis and subsequent decomposition to a glassy product after 3-5 days. The problem of the chemical instability of the precursor was eliminated by incorporating a methyl group into this structure. This chemical method has therefore opened new avenues for the synthesis of nonoxide ceramics with enhanced properties. The area of preceramic polymer chemistry is now about 20 years old. In this time a number of preceramic polymer systems have been developed. These systems include polymer precursors for Si3N4, Sic, AlN, BN, B4C, TiB2, and TiN. In this discussion, siliconbased nonoxide ceramics havebeen generally excluded; they are dealtwith more extensively in a related chapter by Professor Okamura.

II. THEORETICALBACKGROUND As mentioned, important factors that govern the properties of the end products include composition and molecular weight, as well as the characteristics of the metal elements and the organic group functionalities. These interrelated factors determine the fusibility, the percentage of the ceramic yield, and the optimal properties of ceramic products. Therefore, the starting precursors must be cho-

GONSALVES AND XIAO

362

sen carefully in the chemical conversion synthesis of nonoxide ceramics because the final ceramic product depends on the composition ofthe starting polymeric precursor as well as the pyrolysis conditions. The structure of the starting precursors is also important with regard to eliminating or diminishing unwanted by-products andbecause their nature greatly influences the chemistry during pyrolysis [ 11,20,21].

A.

Structure-PropertyRelationship

A high ceramic yield on polymer pyrolysis is important not only for economic viability but also for compensating significant density changes in going from polymer to ceramics. Linear polymers, however, give negligible ceramic yield because of reversion reactions, that is, the generation of large volatile molecules and cyclics [ l 1,22-241. The liberation of such large volatile molecular fragments is responsible for the rapid weight loss. This feature is particularly clear for polymers with backbone structures, such as the.polysiloxanes, which have a relatively low ceiling temperature (Tc). Linear polymers, on the other hand, can increase their ceramic yield if they are pretreated with ultraviolet radiation, for example, which causes cross-linking [25,26] and preventsthe backbiting chain scission mechanism from occumng. Polymer structures containing rings (or cages) give good ceramic yield. Ring structures slow the kinetics of reversion reactions because the back-reaction is slowed by the sterically hindered structures that require multiple bond rupture to produce volatile molecular fragments. The cross-linking necessary for the preceramic structure buildup in general involves localized reactions and the evolution of smaller molecules, such as H2 or CH4, during pyrolysis. Thus, conceptually, linear ring polymers constitute a potentially attractive class of preceramic polymers, combining processability associated with linear polymers with reasonable ceramic yields as a result of the presence of rings. Branched-ring polymer structures also give high ceramic yields, and typical ceramic yields can be as high as 85% [ 1l]. Although branched polymers produce high ceramic yields, “green” ceramic fibers formed from these polymers are generally mechanically weak and brittle because of the combination of low molecular weight and the branched-ring structure, which restricts polymer chain mobility [ 1l]. Pyrolysis yield affects the density of the final ceramic products. Typical pyrolysis yields of carbon-based polymers are 4040%. The change from starting polymer densities of about 1 g/cm3 to final densities of 2 gkm3 or more commonly results in about 30% porosity. Thus, unless there is substantial shrinkage as a result of sintering or plastic deformation, a substantial amount of porosity is generated. Any densification is counterbalanced to some extent by outgassing pressure, which is determined by .pyrolysis conditions, especially the rate of heating and the cross-sectional dimension, that control the rate of

363

NONOXIDE CERAMICS VIA POLYMERS

gas evolution. Thus, the least porosity and cracking are presentedif these ceramics are obtained in powder or fiber forms because of their small diameters [ll].

B. Ceramic Systems The following discussion covers the chemical synthesis of ceramics derived from organometallic polymers: BN, A N , TiN, Tic, and TiB2. It should be emphasized here that some of the syntheses involve starting materials, monomers, and intermediates, as well as polymers, that are oxidatively unstable and/or susceptible to hydrolysis. These syntheses therefore generally require inert atmospheres and the extensive use of vacuum (Schlenk t y p e )line or dry-box techniques. This makes it obvious that collaborations between synthetic chemists and materials and ceramic scientists and engineers is important. Here we outline a selected number of synthetic routes to preceramic polymers. '1. B(N, C) Ceramics Initially, BN was synthesized by the pyrolysis of boric acid and urea in the presence of ammonia [27]. In this reaction, boric acid presumably reacts with urea to form a urea-boric acid complex gel, which on pyrolysis in an ammonia atmosphere results in the formation of BN [28]. Precursors for BN recently received the greatest attention, with emphasis on borazine-derived polymers. It has been recognized that borazine (1) can form polymeric chains of B and N, which can form BN ceramics on thermal pyrolysis [29]. The polymerization gives a graphitelike hexagonal layered structure (2): I

H

I

I

H-B

/N\

B-H

I

H-N

I

\/" I

H

-1

-2

This graphite form ofBNtransforms into a cubic, diamondlike form under pressure at 1800°C,which is claimed to surpass diamond inmechanical strength [30]. Combinations of B-trichloroborazine and hexamethyl-disilazane lead to the formation of gels, which uponthermolysis give a hexagonal BN material. Similarly, on pyrolysis B-trianilinoborazine,B-triamino-N-triphenylborazine, and B-triaminoborazine also give BN ceramics with the presence of C impurities. The pyrolysis of the B-triamino-N-triphenylborazine(3) is [26,31]

GONSALVES AND XIAO

364 Ph I

W - B

/N\

Pm2

I

I Ph -N

250°C

\dWPh -PhNHZ I

(B2N2PhNH2),

>6oO0C

BN

Apart from the impurity content, the ceramics produced by these precursors are apparently rich in B and poor in N [24]. These chemical routes have been extensively reviewed by Paine and Narula [29] and Pouskouleli [26]. Sneddon et al. [31] developed poly(viny1 borazine) and poly(borazy1ene). along with a hybrid inorganic-organic copolymer, poly(viny1 borazinehtyrene), as precursors for BN. The synthesis of B(N, C) ceramics via polyureidoborazines has also been reported[32].Here, triisocyantotrimethylborazine was synthesized from trichlorotrimethylborazine and AgNCO. The latter was derived by the reaction of BC13 with CH3NH2HC1. The polyureidoborazine was then obtained by reacting aliphatic diamines with tri-isocyantotrimethylborazine at ambient temperature. B(N, C) materials were obtained by the pyrolysis of the polymers in N H 3 to 1200°C. Higher boron hydrides have also beenutilized as polymeric precursors for BN. The pyrolysis of poly(2-vinylpentaborane) in a stream of ammonia at 1000°C gave BN in high yield [33]. However, pyrolysis in argon resulted in the formation of boron carbide in high yield[34]. Polymers of the type [B10H12 diamine], prepared byreactionof decacarborane with various diamines, on pyrolysis in ammonia to 1000°C yield amorphous BN and on heating to 1500°C a crystalline material [35].

-

Professor Seyferth and his group at the Massachusetts Institute of Technology (MIT) recently reported the synthesis of polyborasilazanes [36]. These are at an early state of development and have shown good performance in coatings for the protection ofcarbon-carbon composites against high-temperature oxidation. They are preparedby the reactionof cyclic [CH3(H)SiNHIn oligomers with a BHg-Lewis base adduct, such as H3B SMe2, H3B Me2NH, and H3B THF, a process in which a network polymer of borazine rings linked by polysilazane units is formed. This also is a very flexible synthetic procedure in that the boron content of the polymer (and thus in the final ceramic) can be

-

-

NONOXIDE CERAMICS VIA POLYMERS

365

varied from small to substantial. A potentially useful feature of the boron nitride ptoduced is that it retards crystallization of the silicon nitride that is also formed in the pyrolysis of the poly(borasi1azane). These polymers are being investigated in industrial laboratories in the United States and Japan. Boron-containing organosilicon polymers have also been synthesized via a sodium coupling reaction of silicon and boron halides or alkylboron halides in hydrocarbon solvents[37].Onpyrolysis,mixed ceramics comprising Sic, SiBx, and Si-(B, C) can be obtained. 2. AlN Ceramics

A well-known donor-acceptor reaction takes place betweenAlR3andNR3 [38]: AlR,+NR,-R,Al-NR, Depending on R (which can be H), monomers, dimers, trimers, and even octamers, or more complicated rings are produced. On the pyrolysis of these intermediate preceramic precursors, AlN ceramics can be obtained according to R3Al-NR3

-

AlN + by-products

(5)

Similarly,reactions of Al sources with amines [39]canleadtopolymeric N-alkylamino alanes, such as

LiAlH,

+ RNH3Cl-

(AlHNR),

+ 3H2 + LiCl

(7)

where n ranges from 2 to 35. Usually structures with n = 4, 6, or 8 are obtained. To deposit AlN layers for surface wave acoustic electronics, the metal organic chemical vapor deposition method has been used [40-43]. In a typical reaction, Me3A12 was used to react with ammonia at about 1200°C to produce high-purity AlN layers. The reaction scheme is Me6A12+ 2NH3 Me3Al-NH,

2Me3Al-NH,

(8)

AlN + 3CH4

(9)

Inan analogous reaction,AlEt3was also reacted withammoniatoproduce a thermoplastic polymeric precursor,[(EtAlNH),(Et2AlNH2)b (EtgAl),.],, (Q + b/c = 50), which at 1600-1800°C gave a crystalline AlN ceramic [U]. Other Al-containing compounds, such as AlCl3, also react with ammonia to produce aluminum-ammonia complex gels, which on further pyrolysis in the presence of ammonia lead to A1N ceramic materials [28,45,46].

-

GONSALVES AND XIAO

366

A polymeric iminoalane gel has been reported [47-49], which on pyrolysis in ammonia to 1100°C yields crystalline A1N.In this unique electrochemical approach, metallic Al was dissolved anodically in a mixture of primary amine andacetonitrilethat contained a tetraalkylammonium salt asthe supporting electrolyte. Removal of solvent and heating the residue to 150°C yielded the precursor gel.

3.OtherNonoxide Ceramics Ti-containing “monomers,” (R2N)4Ti (R = C2Hs or CH3), can be prepared by the following reactions [50,51]: R,NH TiCl,

-

+ CH,Li R,NLi (10)+ CH, + R,NLi -(R,N),Ti + LiCl

(1 1)

Pyrolysis of this monomer (alone or in a polyacrylonitrile matrix) results in the formation of Ti-C-N ceramics. Recently, the preparation of TiB2 was reported by pyrolyzing a mixture of boron and a polymeric precursor. Thermal decomposition of furfuryl alcohol and hydrolysis of Ti(0Bu)lr produces C and TiO2. Therefore, heating a mixture of boron, titanium butoxide, and furfuryl alcohol results in the the generation of B, C, and TiO2. The carbon then causes the reduction of Ti02 to form TiB2 as the final product, according to [52] 2C + 2B + TiO, = TiB,

+ 2CO

(12)

SU and Sneddon [53] reported the formation of Ti and Zr borides by solidState reactions involving the respective metal oxides intimately dispersed in a decaborane-dinitrile polymer at temperatures greater than 1450”C, presumably according to

MO,

+ polymer dB,C + C + MO,

-

MB,CO

(13)

These approaches also have the potential of being able to synthesize ceramicceramic composites containing transition metal borides or silicides and carbides

[W.

111.

MATERIALSCHARACTERIZATION

Characterization serves as an indispensable adjunct to synthesis and processing activities. To this end, the preceramic polymers must be intensively characterized and the cnre kinetics of the pyrolytic processing studied extensively. The characterization portion is therefore generally composed of studies pertaining to (1) thermal processing and (2) materials properties. Table 1 is a summary of the pertinent instrumental techniques.

NONOXIDE CERAMICS VIA POLYMERS

367

Table 1 Summary of CharacterizationStrategies Measurement Material studies

.Molecular structure Microstructure Elemental composition Surface interface properties Thermal properties Rheology Pyrolytic processing studies

Chemical phenomena Physical phenomena

A.

IR.

NMR, MS, XRD SEM (EDAX), XRD, HRTEM Electron microprobe, SEM-EDAX Auger, X P S , IR microscopy, STM/AFM TMA, DMTA, TGA, DSc, DTA Rheometry GC/MS, DRIFTS, IES, electron microprobe XRD, TEM, SEM

PyrolyticProcessingStudies

1. Pyrolysis of Preceramic Polymers

The transformation of polymers to ceramic on pyrolysis have been extensively studied to develop ahigh degree ofunderstanding about the chemical and physical phenomena (cure behavior) at work during thermal processing by combining vapor-phase monitoring by gas chromatography/mass spectrometry (GCMS) with the solid-phase probe technique of x-ray diffraction (XRD). Diffuse reflectance Fourier transform infrared spectroscopy (DRIFTS) is also used to monitor chemical changes during the pyrolysis. 2. Vapor-Phase Monitoring by Pyrolysis Gas ChromatographyMass

Spectroscopy (PGCMS) A standard technique for the analysis of thermal decomposition and pyrolysis, PGCMS [55,56] can provide detailed information concerning the chemistry of these complex phenomena. When combining PGCMS with thermogravimetric analysis (TGA), it ispossible to obtain qualitative and quantitative chemical information during a thermometric analysis. Thus, the mechanistic details of pyrolysis reaction occurring in the bulk phase can be acquired by separating, identifying, and quantifying volatile species in an indirect analysis. Correlating chemical phenomena with specific phase transitions is an especially powerful advantage of PGCMS when combined with TGA analysis. For these studies, a high-temperature furnace thermogravimetric analyzer can be interfaced with a gas chromatograph, employing capillary chromatographic separations of volatiles with high-speed mass spectrometric detection (to 200 amu). 3. Bulk-Phase Studies by X-ray Diffraction

Although PGCMS is an extremely important analytical tool, it provides only an indirect probe of the chemistry of the pyrolysis. By x-ray powder diffrac-

368

GONSALVES AND XIAO

tion complemented by transmission electron microscopy ("EM), structural information, such as extent of phase transformation, grain size, and crystallinity, associated with pyrolytic processing can be obtained. CP-MAS nuclear magnetic resonance (NMR)spectroscopy is an additional tool for bulk-phase studies of pyrolysis,particularly for characterizing intermediate phasesthat are amorphous and cannot be characterized by XRD or TEM [57]. 4. Bulk-Phase Monitoring by Infrared Spectroscopy

Vibrational spectroscopy can potentially provide a large amount of useful information but is relatively underdeveloped in.the characterization of high-temperature materials. This is particularly true in the understanding and control of the processing steps, when real-time, noninvasive monitoring of the chemical state of the material is required. To obtain desired information, two sampling techniques currentlybeing exploited are (1) diffusionreflectioninfrared Fourier transform spectroscopy and (2) infrared emission spectroscopy (IES) [58]. The DRIFTS technique is the more developed of the two techniques in the sense that the theoretical background is relatively well understood and specialized DRIFTS accessories have been available commercially for almost 20 years. Two early influential papers by Fuller and Griffithsdemonstrated the analytical utility of DRIFTs [59,60]. Since then, DRIFTS has been demonstrated to be applicable to many difficult samples, such as ceramics, coal, and carbon fiber-epoxy prepregs. The signal-to-noise ratio attainable by DRIFTs is excellent (comparable to and sometimes exceeding thatbytransmission spectroscopy of thin semitransparent films), and the sources of distortion(e.g., specular reflection and refractiveindex dispersion) are understood and, to some degree, controllable. However, the DRIFTs technique does not meet the criterion of being noninvasive, and therefore the E S technique is of added interest.

B. Materials Studies 1. Molecular Structure Traditional instrumental techniques, such as nuclear magnetic N M R , mass spectrometry infrared (IR) spectroscopy, ultraviolet-visible spectrophotometry, and gas andliquid chromatography and size-exclusion chromatography, are used extensively for purity assessment and molecular structure and 'molecular weight measurements of monomers and polymers [61]. 2. Thermal Analysis [61] Thermogravimetric analysis can be utilized to determine the thermal stability of polymers indifferent environments and temperature ranges of interest. These results can then be compared with those of thermo-mechanical analysis (TMA)anddifferential thermal analysis(DTA) or differential scanning

NONOXIDE CERAMICS VIA POLYMERS

369

calorimetry (DSC) thermograms to establish thetemperature range in which the polymer has a stable melt:

1. A qualitative approach is to determine the range inwhichthe polymer melts or softens on the TMA. These data can then be compared with the TGA curve within the melting and softening range to observe any degradation or cure. 2. The DSC of the polymers can obtained from 50 to 1600°C to observe any first- or second-order transitions and the onset of chemical reactions. The DSC and TGA curves can be superimposed to differentiate between definite melting or glass transitions and the onset of degradation or cure. 3. By combination of TGA, D S c , and TMA data, polymers that have the POtential to melt can be obtained. These methods provide qualitative data on the thermal properties of the polymer, together with their melt flow character. These studies should be complemented by rheological measurements. It should be mentioned that the melt flow properties of these materials as a function of temperature or time have not been extensively studied. 3. Rheology [62,63] A complete rheological study of the polymer component serves to establish knowledge of the flow behavior of the material and thus contributes toward an optimal choice of melt processing methods. Mixing of the polymers and/or organometallics in the melt can be done with a Brabender plasticization mixer, and the degree of mixing can be controlled by the total energy input and monitored by the torque output. Correlations between the degree of mixing and the mixing conditions can be determined by analyzing the morphology of the mixture using microscopy techniques and relating this to torque measurements. The rheological behavior of the starting polymers and their mixtures can be determined, for example, with a Rheometrics mechanical spectrometer using parallel plates and an oscillatory shear deformation. The dispersion of the metal complex can be determined by measuring the effects on the zero-shear rate viscosity and comparing it to known relationships for filled polymers. Microscopic analysis of the same mixtures can provide a direct measurement of the particle size(s).

C.

Microstructureand Elemental Composition

1. Microstructures Information on the structure and homogeneity of the preceramic polymers or ceramics can be obtained by x-ray diffraction, scanning.electron microscopy (SEM), and transmission electronmicroscopy. A conventional x-raydiffractometer as well as one equipped with a position-sensitive detector can be uti-

370

GONSALVES AND XIAO

lized for XRD studies. SEM studies can provide surface morphology as well as elemental composition of the surface from energy-dispersive analysis of x-rays (EDAX). TEM studies can provide such information as microstructures, phases, and elemental composition in the interior of the materials, as well as material interfaces. For bulk materials and interface studies, samples can be sectioned and thinned for analysis using ion-milling techniques to obtain a desired thickness (4000 A). Another method for sample preparation for preceramic or ceramic coatings uses the replica extraction method. In this method, a tape, combined with an appropriate solvent or acid, is used to extract the surface of a thin-film section. The sample is then coated with carbon film and shadowed with C-Pt to a thickness of about 50-100 A. After the coating process, the sample is placed on a TEM grid and the tape is then dissolved using acetone. For powder materials, samples can be prepared using the solution method. In this technique, powders are usually suspended in deoxygenated alcohol, and a carboncoated grid is dipped into the solution for sampling. This technique, however, requires that the particle size be smaller than 1 pm to transmit the electron beam during TEM examination. 2. Elemental Composition

Electron microprobe analysis can be used for the elemental analysis of preceramic polymers and ceramics. However, since these materials are generally quite thermoxidatively stable and thus are not readily amenable to the traditional combustion approach to elemental determination, chemical analysis methods can be complemented by x-ray fluorescence techniques.

3. Surface Analytical Studies Surface properties of preceramic polymer or ceramic materials can be studied usingAuger electron spectroscopy, x-ray photoelectron spectroscopy ( X P S ) , scanning electron microscopy with EDAX, and infrared microscopy. Fractured specimens can be analyzed and depth-profiling techniques applied to determine any changes in the interface compared with the bulk and surface. An investigation at thenanolevel can be performedby scanning tunneling (STM) or atomic force microscopy (AFM). STM is similar to a roughness measuring device, whichusesa stylus to trace across a surface while mapping the up-down motion of the stylus touching the surface. The most important differences are that the STM is noncontact and has superior (factors of 100’) spatial sensitivity and resolution in all directions. STM can image not only individual atoms but also large features up to hundreds ofnanometers. The AFM is similar to the STM in being able to image samples using sharp tips to achieve high resolution. The advantage of the AFM is that no electrical conductivity is needed, so that insulating materials can also be measured.

NONOXIDE CERAMICS VIA POLYMERS

3 71

IV. APPLICATIONS Nonoxide ceramic fibers have been the major area of interest because of their high modulus values that can be utilized in fiber-reinforced composites. Ceramic fibers include graphite, BN, B4C, AlN, Si3N4, and Sic. The preparation of B4C fibers was first reported in 1969 [a]. These fibers were prepared by the high-temperature chemical conversion of graphite fibers in the presence of a BC13 and H2.These fibers have high modulus values of up to 30-70 million psi. Economy et al. also reported that BN fibers can be obtained by the nitridation of B203 fibers using NH3 gas at about 800°C. Fiber preparation via the polymer precursor route provides many desirable properties for use in continuous-fiber ceramic matrix composites intended for high-temperature uses in oxidative and nonoxidative environments [65]. These fibers, especially those having low electrical conductivity and good dielectric properties, are being investigated for use in radiation-transparent structures, such as radomes [66]. The advantages of the polymeric route to ceramic fibers include the ability to control morphology (amorphous or crystalline and control of crystalline size) and the ability to prepare continuous, fine-diameter fibers (<30 pm) suitable for weaving and knitting. One of the unique advantages of this method is the ability to prepare metastable compositions unobtainable by conventional methods. According to Lipowitz [66], the procedures for ceramic fiber preparation via an organometallic route using polymeric precursors can be generalized ina flowchart:

- -

Organometallic Spun Receramic RWtantS Fiber Polymer

Reactive Gases

I

Fiber

Ceramic Fiber

Pyrolysis Crosslinking Spinning Synthesis

Each process step should be carefully performed because it affects the final fiber properties. Desired fiber properties also depend on the type of spinning methodused. Melt spinning is the most common andpreferredmethod for fiber drawing. It involves the extrusion of a viscous molten polymer through a heated spinnerete. This method requires that the polymer being drawn must be thermally stable (have stable viscosity) at the melt spinning temperature to obtain uniform, high-quality, and high-strength fibers. Other possible methods include dry and wet spinning. After drawing, fibers must be well cured to prevent interfiber fusion during processing. During the curing step, polymer fibers are converted into a highly cross-linked, infusible gel of infinite molecular weight. Cross-linking can be

372

GONSALVES AND XIAO

accomplished by several methods, including thermal, chemical and radiation curing. After the curing step, the as-processed polymer fibers must be thermally treated to produce the desired ceramic fibers. During the pyrolysis step, the loss of volatile polymer components followed by thermal degradation involves simultaneous loss of a large volume of gas, along with a threefold or greater volume shrinkage, increase in density, and development of porosity. Another area in which preceramic polymers can be utilized effectively is as binders for ceramic powders in near net shaping fabrication processes, such as compression or injection molding with subsequent sintering. Alternatively, an active filler and a polymer [67,68], as reported by Greil and Seibold, can be used in such fabrication. Other potential applications of preceramic polymers is in the general area of coatings, especially for carbon-carbon composites [69], and in the synthesis of nanostructured ceramic particles and composites [70-731.

ACKNOWLEDGMENTS We are grateful to Professor Dietmar Seyferth (MIT) for his kind permission to use information provided by him to us and for the inspiration provided by his pioneering work in the area of preceramic polymers.

REFERENCES 1. Gonsalves, K. E.,in Inorganic and Metal-Containing Polymeric Materials (J. E. Sheats,C. E. Carraher,C. U. Pittman, Jr., M.Zeldin,andB.Currell, eds.), Plenum, New York, 1990, p. 173. 2. Seyferth, D., Wiseman, G. H., Schwark, J. M.,Yu,Y. F., andPoutasse,C. A., Am. Chem. Soc. Symp. Ser. No 360, 143 (1988). 3. Messier, D. R., and Croft,W. J., Properties of Solid State Materials, Vol. 7, 1982, p.131. 4. Morosanu, C. E.,Microelectron. Reliab., 20, 357 (1980). 5. Kingon, A. I., Lutz, L. J., Liaw, P., and Davis, R. F., J. Am. Ceram. Soc., 66, 551 (1983). 6. Airey, A. C., Clarke, S., andPopper,P., Trans. J. Br. Ceram. Soc., 22, 305 (1973). 7. Galasso, F., Kuntz, U., and Croft, W.J., J. Am. Ceram. Soc., 55(8), 431 (1973). 8. Kajima, K., Setaka, N., and Tanaka, H.,J. Cryst. Growth, 24-25, 183 (1974). 9. Niihira, K., and Hirai, T., J. Mater. Sci., 11(4), 593 (1976). 10. Oda, K., Yoshio, T., and Oka, K.,J. Am. Ceram Soc., 64(4), c8 (1983). 11. Wynne,K. J., and Rice, R. W., Annu. Rev. Mater. Sci., 14, 297 (1984). 12. Seyferth, D., MIT, private communication 13. Stock, A., and Somieski, K., Ber. Dtsch. Chem. Ges., 54, 740 (1921). 14. Brew, S. D., and Haber, C. P., J. Am. Ceram. Soc., 70, 361 (1948). 15. Osthoff, R. C., and Kantor, S. W., Znorg. Synth., 5, 61 (1957).

NONOXIDE CERAMICS VIA POLYMERS

3 73

16. Yajima, S., Hayashi, J., and Omori, M., Chem. Lett., 931 (1975). 17. Yajima, S., Okamura, K., and Hayashi, J., Chem. Lett., 931 (1975). J. Am. Ceram. Soc., 66, C13 18. Seyferth, D., Wiseman, G. H., and Pmdhomme, C., (1983). 19. Seyferth, D., and Wiseman, G. H., J. Am. Ceram Soc., 67, C132 (1984). 20. West, R, and Maxka, J., Am. Chem. Soc., 360, 7 (1988). K., PolymerCarbon-CarbonFibre Glass and 21. Jenkin,G.M.,andKawanura, Char, Cambridge University Press, London, 1976. 22. Wesson, J. P., and Williams, T. C., J. Polym. Sci Polym. Chem., 17, 2833 (1979). 23. Wesson, J. P., and Williams, T. C.,J. Polym. Sci Polym. Chem., 18, 959 (1980). 24. Wesson, J. P., and Williams, T. C., J. Polym Sci Polym Chem., 19, 65 (1981). I., and Yu, H., Am. Ceram. Soc. Bull., 62, 25. West, R., David, L. D., Djurovich, P. 889 (1983). 26. Pouskouleli, G., Ceram. Int., 15, 213-229 (1989). 27. O'Connor, T. E., J. Am. Ceram. Soc., 1733 (1963). J. Am. Ceram. Soc., 76(4), 987 28. Xiao, T. D., Gonsalves, K. E., and Strutt, P. R. (1993). 29. Paine, R. T.,and Narula, C. K., Chem. Rev., 90,73-91 (1990). 30. Wentorf, R. H., J. Chem. Phys., 26, 956 (1957). T., Remsen, E. E., and Beck, J. S., 31. Sneddon, L., Su, K., Fazen, P. J., Lynch, A. 32. 33. 34. 35. 36. 37.

in InorganicandOrganometallicOligomersandPolymers, (J. F. Harrodand R. M. Laine, eds.) Kluwer, Amsterdam, 1991, p. 191. Gonsalves, K. E., and Agarwal, R., Appl. Orgmetal. Chem. 2, 245 (1988). Mirabelli, M. G. C., and Sneddon, L. G., Inorg. Chem., 27, 327 (1988). Mirabelli, M. G. C., and Sneddon, L. G., J. Am. Chem. Soc., 110, 3305 (1988). Seyferth, D., and Rees, W. S., Jr., Chem. Mater., 3, 1106 (1991). Seyferth, D., and Plenio, H., J. A m Chem. Soc., 73, 2131 (1990). Jaffe, R. L., Riccitiello, S., Hsu,M. T., Chen, T.. andKomornicki,A., Polym

Preprints, Polym. Chem Div. Chem. Soc., 32(3), 489 (1991). 38. Davison, N., and Brown, H. C.,J. Am. Chem. Soc., 54, 316 (1942). 39. Cotton, F. A., and Wilkinson, G., Advanced Inorganic Chemistry, 4th ed., New York, 1980, p. 345. S., Tsubouchi,K.,andMikoshiba,N., Jpn. 40. Morita,M.,Uesugi,N.,Isogai, J. Appl. Phys., 20, 17(1981). 41. Eichhorn, G., and Rensch, W., Phys. Status Solidi, A69, K3 (1982). F. M.,andSimpson,W. I., J. Electrochem Soc., 42. Manasevit,H.M.,Erdmann, 118, 1864 (1971). 43. Maeda, T., and Harado, K. (Sumitomo Chemical Co., Ltd.), Manufacture of Aluminum Nitride, Japan Kokai, 78 69 700, January 1978. 44. Bolt, J. D., and Tebbe,F. N., Mater. Res. Soc. Symp. Proc., 108, 337 (1988); 121, (1988). 471 ' 45. Riedel, R., and Gaudl, K., J. Am. Ceram. Soc., 74, 1331(1991). J. Am. Ceram. 46. Twait, D. J., Lackey, W. J., Smith, A. W., and Hanigofsky, J. A., Soc., 72, 1510 (1990). 47. Seibold; M. and Russel, C., Mater. Res. Soc. Symp. Proc., 121, 477 (1988). 48. Distler, P., and Russel, C., J. Mater. Sci., 27, 133 (1992).

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49. 50. 51. 52. 53. 54. 55. 56. 57. 58. 59. 60. 61. 62. 63. 64. 65. 66. 67. 68. 69. 70. 71. 72. 73.

GONSALVES AND XIAO Tensel, I., and Russel, C., J. Mater Sci., 25, 3531 (1990). Gonsalves, K. E., and Agarwal, R., J. Appl. Poly. Sci., 36, 1659 (1988). Gonsalves, K. E., and Kembaiyan, K. T.,Solid State lonics, 32/33, 661 (1989). Jiang, Z., and Rhine, W. E., Chem. Mater., 4, 497-500 (1992). Su, K., and Sneddon, L. G., Chem Mater., 3, 10-12 (1991). Seyferth, D., Bryson, N., Workman, D. P., and Sobon, C. A., J. A m Ceram. Soc., 74 2687 (1991). Paciorek, K. J. L., Harris, D. H.,and Kratzer, R. H., J. Polym. Sci., Polym. Chem. Ed., 24, 173 (1986). Ballister, A., Garazzo, D., and Montando, G., Mucromol., 17, 1312 (1984). Babonneau, F., Livage, J., and Laine, R. M., Polymer Preprints, Polym.Chem. Div. Chem. Soc., 32(3), 579 (1991). Garton, A., Infrared Spectroscopy of Multicomponent Polymer Materials Hanser, New York, 1992. Fuller, M.P., and Griffith, P. R., Anal. Chem., 50, 1906 (1978). Fuller, M. P., and Griffith, P.R., Appl. Spectrom., 34, 533 (1980). Billmeyer, F. W., Textbook of Polymer Science, 2nd ed., Wiley, New York. Baney, R., Polymer Preprints, 25(l), 2 (1984). Tadmor, Z., and Gogos, C. G., Principles of Polymer Processing, Wiley, New York,1978. Beemsten, D. J., Smith, W. D., and Economy, J., Appl. Polym. Symp., No. 9,365 (1969). Mazdiyasni, K. (ed.), Fiber-Reinforced Ceramic Composites; Materials, Processing and Technology, Noyes Publications, Park Ridge, NJ, 1990. Lipowitz, J., Ceram. Bull., 70, 1888 (1991). Seibold, M.,and Greil, P.,Adv. Mater. Proc. (Germany), I , 641 (1990). Greil, P., and Seibold, M.,Chem. Trans. (Adv. Compos. Mater.), 19, 43 (1991). Gonsalves, K. E., and Yazici, R., J. Mater. Sci Left., IO, 834 (1991). Magee, A. P., Strutt, P. R., Gonsalves K. E., Chem. Mater., 2(3), 232 (1990). Gonsalves, K.E., Strutt, P. R., and Xiao, T.D., J. Adv. Mater., 3, 202 (1991). Gonsalves, K. E., Strutt, P. R.,Xiao,T. D., and Klemens, P. G., J. Mater. Sci., 27(12), 3231 (1992). Xiao, T. D., Gonsalves, K. E., Strutt, P. R.,andKlemens, P. G., J. Mater. Sci., 28, 1334 (1993).

16 Polymer Pyrolysis Masaki Narisawa and Kiyohito Okamura University of Osaka Prefecture

Osaka, Japan

1.

INTRODUCTION

The development of processing ceramics from polymer precursors has attracted great attention in recent years. In particular, inorganic polymers containing silicon are being actively studied as precursors for ceramics. S i c ceramics have the advantage of high-temperature stability in an oxidizing atmosphere. S i c is not readily sintered, however, and so is difficult to obtain in either fiber or film form by traditional inorganic processes. The polymer precursor method is available to produce continuous fibers, ceramic coatings, or sintered ceramic bodies. S i c fibers were produced using polycarbosilanes by Yajima et al. in 1975 [1,2]. Besides S i c fibers, Si-Ti-C-0 fibers prepared from a polytitanocarbosilane have been obtained by adding a titanium tetrabutoxide to polycarbosilane or polysilane [3]. S i c fibers (Nicalon) and Si-Ti-C-0 fibers (Tyranno) are manufactured on an industrial scale. Colorless silicon oxynitride fibers and silicon nitride fibers [4] have been obtained by the nitridation of polycarbosilanes in the author’s laboratory. Polymers used for ceramic precursor and the resulting ceramic fibers are listed in Table 1. These fibers have high tensile strength and heat resistance. In recent years, S i c fibers or their fabrics have been used to prepare fiber-reinforced ceramicmatrix composites. Such composites havehigh strength and high fracture toughness, even at elevated temperatures [5,6], and are a promising class of

375

NARISAWA AND OKAMURA

376 Table 1 CeramicsfromPolymerPrecursors Decade 1970

1980

1990

Polysilazanes Polycarbosilane Polyaluminoxiane Borazine Polyborosiloxane Polysilastyrene Polycarboranes siloxane Polysilane Polytitanocarbosilane Polysilazane Polysilazane Polysilazane Polycarbosilane Perhydropolysilazane Polymetallocarbosilane Polysilazane (boron) Aluminum nitride polymer Hybrid polymers

Six-N Si-C A1203 BN Si-B-Ca Si-C SiC-B4C Si4 Si-Ti-C4 Si-N-C Si3N4 Sic, SicN4, C Si-N-O, Si3N4 Si3N4 Si-"C4 Si-N-B AI-N SbNdAlN Si3NdN

structural materials for application in a high-temperature or radioactive environment. Besides the fibers, application of metallorganic polymers to heat-resistant coatings, ceramic moldings, or sintering aids is being developed. Other types of noble polymers, not polycarbosilane alone, are attracting attention for these uses. This review concerns the metallorganic polymers, the process of pyrolysis, and the property of the ceramics that are obtained.

II.

POLYMERPRECURSORSFORCERAMICS

The chemical structure of polymer precursors is not easily identified and ordinarily consists of ring and chain or branched-ring groups. To use as ceramic precursors, the polymer must meet the following conditions; 1. Ease of manufacture in a melt 2. Solubilityin organic solvent 3. Yield of pyrolysis beyond 50% 4. Ceramics with amorphous or microcrystalline phase

POLYMER PYROLYSIS

A.

377

PolymerPrecursors for SIC Ceramics

1. Polysilanes Polysilane is usually synthesized from dichlorosilane (RlR2SiC12) by dechlorination condensation reaction withsodium. Various types of alkyl or aryl groups can be added to (the changing monomers) main -Si-Si- chain. Polysilastyrene is a representative copolymer synthesized by this method of mixing the monomers [7]dimethyldichlorosilane and phenylmethyldichlorosilane[g]. Baney and Gaul synthesized methylchloropolysilanefrom disilanes (=Si-Si=) by redistribution reactions of silicon-chlorine/silicon-siliconbonds [9,10].Distilled mixed disilanes were catalytically rearranged with tetrabutylphosphonium chloride (1 %) to produce preceramic polymer. The polymer obtained is considered to have apolycyclic-Si-Si-chain. The structuralunits are =Si (CH3), =Si(CH3)2, and =Si(CH3)C1. Several new polymers were derived from methylchloropolysilane by substituting the remaining silicon-chlorine bonds by other chemical groups [ 1l]. 2. Polycarbosilanes Polycarbosilane is a general term for organosilicon polymers with -Si-C- bonds in the main chain. Manykinds of polycarbosilanes have been synthesized using various methods. Figure 1 illustrates some polycarbosilanes that have been prepared in the author's laboratory. Polycarbosilanes are usually synthesized by the thermal decomposition of monosilanes or disilanes and by ring-opening polymerization of disilacyclobutanes. Fritz et al. sythesized polycarbosilanes (PC-TMS) by heat condensation of tetramethylsilane at 700"C, circulating unreacted silane repeatedly in a continuous pyrolysis furnace [12]. This method is expensive and time consuming, however. As an alternative, a polycarbosilane (PC-A) was obtained from polydimethylsilane (PDS) [2].In the first step, PDSwas subjected to dechlorination condensation in xylene heated under N2 gas using sodium to produce white PDS powder. The PDS was then heated at 450470°C in an Ar atmosphere in an autoclave to give PC-A in a yield of over 60%. PDS can be also pyrolyzed and polymerized under N2 gas flow at normal pressure using a reflux condenser to avoid the use of the autoclave. A polycarbosilane (PC-N) was obtained in a yield of about 50% [13].A representative structural model of polycarbosilane is shown in Fig. 2. When pyrolyzing PDS under normal pressure, a few percent polyborodiphenylsiloxane (PBDPSO) [l41 is added so that the reaction proceeds catalytically, enhancing the rate. The resinous product (PC-B) was obtained in a yield of over 60% [15].PC-N and PC-B are now produced on an industrial scale.

378

NARISAWA AND OKAMURA

(PC-B)

(PC-A)

Figure 1 Preparation of various polycarbosilanes.

CH3 H CH2 CH2 CH2 CH2 SI \H\/ \/ \/ \ / \ Si Si Si Si Si

CH3 CH3

"

I

I

I

I

I

I

CH2 CH2CH H CH CH

I/

H'

'\/\/\H\/\/\ Si Si Si Si / \ /-\ /-\ /\

CH3 CH3 H H

'CH3 Si

/\

CH~CH~CHJCH~CH~CH~

Figure 2 Chemicalstructure of polycarbosilane.

379

POLYMER

Dimethyldichlorosilane was converted into not onlyPDS but also aring molecule, dodecamethylcyclohexasilane (DMCHS). From DMCHS, a polycarbosilane was obtained [1,2]. Besides dimethyldichlorosilane, other types of dichlorosilane were used to produce polycarbosilane. Dechlorination condensation reaction of vinylmethyldichlorosilane and monosilane mixture yields a polycarbosilane [16,17]. The molecular structure of polycarbosilane is difficult to represent precisely. However, from measurements of the molecular weight, the intrinsic viscosity, infared and ultraviolet spectroscopy, proton, carbon, and silicon nuclear magnetic resonance (1H,13C, and 29SiNMR), and chemical analysis, the structure of polycarbosilane is found to be represented by three simple units. These are silicon bonded with four carbon atoms (SiC4), silicon bonded with three carbon atoms (SiQH), and silicon bonded with x carbon atoms and 4 - x silicon atoms (SiCSi4 "x,x = 1,2, or 3). These are shown by the following structural units [181: CH3

I

I

-CH2-Si-CH2-

-CH2-Si-CH:!-

I

I

(SiC3H) H

(SiC4) CH3 CH3

I

I

-CH2-Si-Si-Si-Si-Si-

I

t

CH3 CH3 CH3 CH3 CH3

I

l

l

I

l

l (SiC$3iCbx)

For PC-TMS, PC-A, and PC-N, the fractions of Sic4 and SiC3H are large; the fraction of SiCSi4-x is very small. For PC-B, on the other hand, the fraction is in the order Sic4 = SiCSi4x > SiC3H. The fractions of the three units in some typical polycarbosilanes are listed in Table 2. Polytitanocarbosilane(PTC) is synthesized by heating a mixture of PDS and PBDPSO with titanium alkoxide. Cross-linking by the titanium compound occurs simultaneously with the conversion of PDS to polycarbosilane [3,19]. 3. Polysiloxanes By the pyrolysis of polymethylsiloxane generally used as silicon resin, a small amount of S i c ceramic is obtained. However, polydiphenylborosiloxane [20] or polydiphenyl-glycerosiloxane is available as a precursor inayield of 40-50%, which was synthesized by dechlorination condensation of diphenyl-

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NARISA WA AND OKAMURA

Table 2

Fraction of Structural Unitsin Polycarbosilane

Polycarbosilane

Sic4

SiC3H

PC-A470 PC-B-3.2

0.53

0.47

0

PC-B-5.5

0.40 0.44

PC-TMS

0.82

0.15 0.15 0.18

0.41 0

0.44

Source: From Ref. 18.

dichlorosilane, boric acid,andglycerin. To achieve a greater ceramic yield with polysiloxane, thecarbon content in the polymermustbeincreased. A phenyl group is more appropriate than a methyl group for the side chain. In particular, PBDPSO is used for the binder in IR radiating coating materials, not only as a polycarbosilane synthesis accelerator. The structure of PBDPSO is C6H5

I

-si

I

/

0-

- o - B'

\

C6H5

0-

Recently, polysiloxane whose composition is (MeSiO1.5)0.75-x(PhSi00.5)~ (Me2ViSi00.5)0.25was synthesized by theDow Coming Corporation. This polymer is reported to be good as a sintering aid [21].

B.

PolymerPrecursors for Si3N4 Ceramics

1. Polysilazanes Polysilazanes for precursor polymers were synthesized from alkylchlorosilanes and alkylamines by Verbeek [22], and the process was characterized by Penn et al. [23]. The polymer was obtained by the polymerization of tris(N-methylamino)methylsilane [CH3Si(NHCH3)3], synthesized by the reaction of methylamine with methylmchlorosilane:

+

CH3SiC13 6CH3NH2 + CH,Si(NHCH,),

+3[CH3N+H3]C1-

A general formula for the polysilazane obtained is [CH3(CH3NH)Si(CH3N)]x [CH3Si(CH3N)1.sly, and its chemical structure is considered branched, partly incorporating a ring structure. Seyferth synthesized soluble polysilazanes in high yield by ammonolysis of SiH2Cb in polar solvents [24], such as dimethyl ether or dichloromethane. The ammonolysis of CHsSiHC12 also yields silazane oligomers, which can be poly-

POLYMER

381

merized by KH catalyst. The structure unit of polymer is -CH3SiH-NH-, and the -CHSiH-NH- rings are combined by SizNz units. A soluble polysilazane was obtained by the ammonolysis ofan adduct between SiHzCIz and pyridine [25]. A formula for the polymer is (SiHzNH)o.46(SiH2)0.36(SiI-I3)0.i~,and a structural model is shown in Fig. 3. This polymer is produced as the precursor for the Si3N4 fiber and Si-N coating by Tonen Corporation.

2. Polydisilazanes Polydisilylazanes have been synthesized by the reaction of mixeddisilanes with hexamethyldisilazane through silicon-chlorine/silicon-nitrogenbond redistribution reactions [10,26]. The formula [(cH3)2.6(si2)i.ONHi.5]ii was deduced during (reaction) at 250°C. 3. Nitridation of Polycarbosilanes Heat treating polycarbosilane in an NH3 atmosphere yields amorphous Si3N4 ceramics [27].If polycarbosilane is cured by oxygen before heat treating, amorphous Si-N-0 is obtained [4]. The oxygen content is correlated with the structure of the amorphous phase.

C.

Polymer Precursors for Other Types of Ceramics

The synthesis of boron nitride (BN) fiber has been attempted. The thermoplastic precursor for BN fiber was obtained by the reaction between B-tris(methy1amino)borazine and lauric amine [28,29]. Aluminum nitride polymer has been actively studied as the precursor for AlN ceramics [30,31].

Figure 3 Chemical structure of perhydroplysilazane. (From Ref. 25.)

382

111.

NARISAWA AND OKAMURA

SYNTHESISOFCERAMICS FROM POLYMER PRECURSORS

Various types of ceramic fibers, coatings, and moldings from precursor polymers are produced and commercialized on an industrial scale. The properties of ceramics are correlated with the microstructure formed in the thermal decomposition process. Detailed information about polymer pyrolysis is now required to control the qualities of the ceramics obtained. Figure 4 shows weight loss curves for polycarbosilane, polysilazane, and polysiloxane, which are representative precursors for ceramics. The schematic mechanisms of each polymer pyrolysis are [32] >IOOo"C

(-R,R,Si-CH,"),

>1OOo"C

(-R,R,Si-W-),

> SicI+, > Si,C,N,

>16000c> >16oooc

SiC+C

polycarbosilane

> S i c + C + N,

polysilazane

The ceramic yield of polycarbosilane is excellent above 1600°C compared with the other polymers, but the yield is low below 1500°C. The thermal decomposition of polycarbosilane in particular has been well studied because of its widespread use for fibers and coatings. Various kinds of ceramic fibers obtained from polycarbosilane are shown in Fig. 5.

A.Synthesis

of SICFibers

Figure 6 shows the gas evolution and structural change in polycarbosilane during heat treatment [33]. The dotted line indicates that the polycarbosilane was cured in air. r

l

Temperature/'C

Figure 4 Weight loss curves for plycarbosilane (A), polysilazane (B), and polycarbosilane (C). (From Ref. 32.)

POLYMER PYROLYSIS

383

Electron ifradiation curing

Oxidation curing

/ \

/ \

Heating in NHS

Heating in Ar or in vacuum

Heating in NHJ

Heating in Ar or in vacuum

I

J

I

I

Silicon oxynltride fiber

Silicon carbide fiber (Si-C-0)

Figure 5 Preparation of various ceramics from polycarbosilane.

In the first step, the polycarbosilane is concentrated by distillation to remove low-molecular-weight components and to adjust its spinnability. A typical distillation temperature is 280"C, and the number-averaged molecular weight of the resulting polymer is 1500-2000. The polymer is spun into fibers by a melt spinning method. At this stage, the spun polycarbosilane fibers are very fragile. The spun fibers are cured by oxidation in air at a temperature ofup to 100-200°C or by 'y-ray or electron beam irradiation in a vacuum or an inert atmosphere. The curing process is necessary to permit the conversion of polycarbosilane fiber to silicon carbide without softening during heat treatment. Oxidation curing is achieved by cross-linking the polycarbosilane with oxygen bonds, such as Si-0-Si or Si"c, whereas radiation curing forms direct Si-Si or Si-C bonds [34]. Heating polycarbosilane, whether cured or uncured, results in the evolution of H:! and CH4 gas as the heating temperature is increased. The thermal decomposition product from 500 to 800°C has a structure intermediate between that of organic and inorganic compounds. The product contains many radicals, and the pyrolyzed fiber keeps its flexibility. Although the thermal decomposition of cured polycarbosilane fiber. is completed in the neighborhood of 8OO"C,

384

-c(

a

d

0

NARISA WA AND OKAMURA

POLYMER

385

its strength continues to increase with temperature. Above 100O"C, the structure of the thermal decomposition product changes from the amorphous to the micro-crystalline state, and the tensile strength of pyrolyzed fiber achieves a maximum value at around 1200°C. The schematic representation of atomic arrangement in pyrolyzed Sic fiber through oxidation curing is showninFig.7. The formula for thefiber is Sic& (x = 1.2-1.4, y = 0.3-0.4). Above 1300"C, evolution of CO gas and hence weight loss occur with the p-Sic crystallization, resulting in a drop in strength. CO gas evolution around 1300°C occurs only with oxidation-cured polycarbosilane. Figure 8 shows the effect of electron curing on the mechanical properties of the fiber. Both Young's modulus and the tensile strength at high temperature are improved remarkably [35]. Raman spectroscopy is useful for studying @-Sic crystallite andCa-ring structure in Sic fiber above 1200°C [36], whereas infraredspectroscopy is useful for studying polymer structure below 1000°C. Raman spectra of oxidationcured Sic fibers obtained by heat treatment at various temperatures were measured, and these are shown in Fig. 9. The integrated Raman intensities of each components can be estimated from the width and amplitude of the individual curves in Fig. 9. The Sic intensity includes curves for the amorphous and crystalline states. The Sic component is almost constant to 1200"C, beyond which it decreases before increasing abruptly at 1500°C. Carbon components with a Ca-ring structure increase withheat treatment to 1400°Candthensuddenly decrease at 1500°C. These results suggest that excess carbon in Sic fibers precipitates to form carbon particles up to 1400"C, which then disappear above 1500°C. The CO gas evolution between 1400 and 1500°C may compete with the precipitation of the carbon. In the Sic fiber with the highest tensile strength (i.e., that treated at 1200"C), a small number of carbon and S i c microcrystallites (1-2 nm) are considered distributed uniformly in carbon-rich amorphous Sic [36]. The carbon component intheRamanspectravirtually vanishes for heat treatment temperatures > 1800"C, whereas the Si-0 line remains up to 1900°C. Thus the amount of excess carbon is expected to be less than or comparable to that of oxygen contained in Sic. Thus, from these combined experimental results, it is considered that the strength of S i c fibers decreases with the crystallization of p-Sic and the elimination of excess carbon and oxygen contained in the fibers. Recently, ESR spectra of Sic fibers and polycarbosilane without curing gave interesting information about the thermal decomposition process [37]. The quantity of radicalsin polycarbosilane shows two peaks at 800 and 1200°C during thermal decomposition, whereas ordinary organic polymers show only one peak [38]. The ESR spectra of Sic fiber cured by irradiation show behav-

386

NARISAWA AND OKAMURA

Si: Figure 7 Atomic arrangement in Sic fiber pyrolyzed by oxidative curing. iorsimilar to that of polycarbosilane,exceptabsolute quantity of radicals. However, ESR of fiber cured by oxygen shows only one large peak at the second position, 1200°C. For polycarbosilane without curing, the peak width reachesamaximum ataround 700°C and aminimum at 1100°C. Beyond

387

POLYMER

r

300

3.01

rI

d)

loot 1200

1300

1400

1500

Temperaturd 'C

I T

I

&

I

I

I

T

1200 1300 1400 1500

Temperature/'C

Figure 8 Tensile strength and Young's modulus of Sic fibers. (a) Oxidation curing (02content 10.8%). (b) Electron curing( 0 2 content 4.0%). (c) Electron curing( 0 2 content 1.9%). (d) Electron curing (@ content < 0.5%). (From Ref. 43.)

11OO"C, the peak width gradually increases. The first peak at 800"C, indicating the existence of many dangling bonds, seems to correspond to the finishing point of structural change from organic to inorganic amorphous, and the second peak at 1200°C corresponds to the finishing point of H2 gas evolution and Cg-ring precipitation from carbon-rich amorphous Sic.

B. Synthesis of SiC-TC Fibers SiC-Tic fibers have been obtained by the pyrolysis of a polytitanocarbosilane (PTC-0) [39]. PTC was synthesized by adding titanium alkoxide to polycarbosilane in xylene. The PTC was melt spun and cured in air at 180°C. The cured fibers were heated in N2 gas at temperatures in the range 800-1500°C. Ube Industries Ltd. has prepared Si-Tic-0 fiber (Tyranno) on an industrial scale [40]. Figure 10 shows the tensile strength, Young's modulus, and specific resistance of Tyranno fibers. Tensile strength is maximum for the 1300°C treatment compared with the maximum at 1200°C for S i c fibers. This superior heat resistance of Tyranno fiber seems to be caused by the bonding of excess carbon with titanium or the increase in the overall bonding strength of the constituent elements in the fiber by the addition of titanium.

388

NARISAWA AND OKAMURA

1

0

I

200

800 1400 Raman shift (cm”)

2000

Figure 9 Ramanspectra of Sic fibersobtainedbytheheattreatment of polycarbosilanefibers at (a) 1000°C, (b) 1200”C,(c) 1400°C. (d) 1500”C,and (e) 1700°C (From Ref. 36.)

POLYMER PYROLYSIS

389

-2

P

200 -

f !Vi i P

U

-a Y)

3

U

2 100-

.Y)

C

r-"

:P ,

(a)

800

*

(c)

,

I

,

I

.

I

L

(b)

I

1000

.

1200

I

I

1400

I

Temperature ("C)

Figure 10 (a)Tensilestrength, PTC fibers. (From Ref. 40.)

I

0 700 900 110013001500 Temperature ("C)

Temperature ("C)

1l"I

P

(b) tensile modulus, and (c)specificresistance

of

390

--

NARISAWA AND OKAMURA

C. Synthesis of Silicon Nitride and Silicon Oxynitride Fibers from Polycarbosilanes

The nitridation of polycarbosilane (with ammonia) begins at about 500°C and terminates almost completely at about 800°C. There is scarcely any carbon in the nitride obtained at 1400°C. Chemical analysis indicates its composition to be Si3N3.72, almost that of pure silicon nitride, Si3N4 [39]. Up to 1300"C, the X-ray diffraction pattern of nitride fiber is broad and characteristic of the amorphous state, whereas those at 1400°C indicate crystalline a-SisN4 (Fig. 11). Both silicon nitride and silicon oxynitride fibers have been obtained by the heat

120O0C

13OO0C

20 (degrees)

Figure 11 X-ray diffraction patterns of silicon nitride obtained by the nitridation polycarbosilane. (From Ref. 34.)

of

POLYMER PYROLYSIS

391

treatment of electron-cured and oxygen-cured polycarbosilane fibers, respectively, in an NH3 gas flow [34]. Both nitride fibers are colorless, transparent to visible light, and amorphous.

D. Synthesis of Silicon Nitride Fibers from Polysilazanes Continuous stoichiometric silicon nitride fiber was produced by the pyrolysis of perhydropolysilazane.A transparent colorless silicon nitridefiber, which has high strength, high modulus, and high thermal stability and is suitable to reinforce plastics, metals, glasses, and even ceramics, was obtained. The fiber was characterized by Fourier transform infrared (FTIR) spectroscopy, X-raydiffraction and 29Si NMR. The surface characterization of the fiber was also conducted using scanning electron microscopy, X P S , and FTIR [41].

E. Synthesis of Sic Whisker from a Biological Source Sic whisker has been obtained by the pyrolysis of rice shell. The precursor, rice shell, contains 15-20wt% Si02. The rice shell is heat treated at 700°C yields an intermediate. The Sic whisker, whose diameter is 0.1-.5 p,is obtained by pyrolysis of the intermediate at 1500°C. The whisker contains 75% a-Sic and 25% p-Sic. The ceramic yield is only 10%; however, a larger yield was achieved by using Fe as catalyst [42].

F.

Synthesis ofSi3N4-SiCCeramics

Si3N4 ceramics using polycarbosilane have been prepared by Kurosaki Yogyo (Kuroceram-N). In the first step,polycarbosilane is mixed with Si powder with a diameter below 10 p and cured by oxygen. By pyrolysis in a N2 atmosphere, the sample is converted into Si3N4 ceramics. Nitridation of Si yields Si3N4, and the pyrolysis of polycarbosilane yields Sic and excess carbon. The excess carbon is trapped by unreacted Si and changed Sic. The ceramics obtained consist mainly of Si3N4 with a few percent Sic. By changing the quantity of polycarbosilane, the microstructure of the ceramic can be controlled through the form of the residualSic. A microgap between Si3N4 and Sic crystallites gives flexibility to the material, and high thermal shock resistance is achieved.

G.

SynthesisofHeat-ResistantCoatings

Tyranno Coat was obtained from Tyranno polymer and widely used [43]. The virtures of Tyranno Coat are as follows: 1. The polymer can be solubilized in organic solvent and used in the form of

a solution or slurry.

N4RISAWA AND OKAMURA

7.,

L

0

250

Tyranno polymer

500

750

1000

Temperature/'C Figure 12 Weight loss curve of Tyranno Coat (in air). (From Ref. 43.)

2. The polymer is thermoplastic and is easily manufactured because the softening point is about 300°C. 3. The polymer forms persistent plastic coatings below 400°C and changes into S i c ceramics above 500°C.

The weight loss of Tyranno Coat is shown in Fig. 12 compared with that of phosphate and polysiloxane. By dispersing ZrO2 particles in Tyranno Coat, highly radiative, highly heat resistant, and high thermal shock-resistant coating is obtained, which can avoid the radiation loss from surface reflection. Besides radiative coatings, Tyranno Coat is generally used to prevent the corrosion and oxidation of metals. Polyborodiphenylsiloxane is also used as a binder for infrared coatings. The coating is high in quality in particular for far-infrared rays. Recently, a coating of perhydropolysilazane on stainless steel has attracted attention. The perhydropolysilazane is converted into silicon nitride in an N H 3 atmosphere at 600°C or to a transparent siliconoxynitride in humid air beyond 150°C. This material is making progress in overcoming some weakness in the expansion mismatch to the coated metal and in contraction during pyrolysis.

W. CONCLUSIONS Several kinds of ceramics prepared from metallorganic polymers have been described. By controlling the precursor and conversion process, polymer precursor ceramics have the possibility to form novel types of structure, such as an interpenetrated microstructure, that is different from that of earlier ceramics.

POLYMER PYROLYSIS

393

Ceramic fibers are the most advancedmaterials in the polymer precursor ceramic field. These fibers have high tensile strength and heat resistance and so are promising as reinforcement fibers for composites, such as fiber-reinforced plastics, metals, and ceramics. Detailed investigation of the structure and nature of precursor polymers and the conversion process are in progress. Many interesting phenomena occur during the conversion of polymers, and the study of these will expand not only with regard to ceramic fibers but also to other ceramic materials with refractory, magnetic, and electronic properties. Therefore it can be predicted that new ceramics with a variety of properties will be prepared in the near future.

REFERENCES 1. Yajima, S., Hayashi, J., and Omori, M., Chem Lett., 931 (1975). 2. Yajima, S., Okamura, K., Hayashi, J., and Omori, M., J. Am. Ceram Soc., 59, 324 (1976). 3. Yajima, S., Iwai, T., Yamamura, T., Okamura, K., and Hasegawa, Y., J. Mater. Sci., 16, 1349(1981). 4. Okamura, K., Sato, M., Hasegawa, Y., and Amano, T., Chem Lett.,2059 (1984). 5. Prewo, K. M., Brennan, J. J., and Layden, G. K., Am. Ceram. Soc. Bull., 65, 305, 322 (1986). 6. Cornie, J. A., Chinag, Y., Uhlmann, D. R., Mortensen, A., and Collins, J. M., Am. Ceram. Soc. Bull., 65, 293 (1986). 7. West, R., J. Organometal. Chem, 300, 327 (1986). 8. West, R., David, L. D., Djurovich, P. I., and Yu, H., Am. Ceram. Soc. Bull., 62, 899 (1983). 9. Baney, R. H., and Gaul, J. H., Jr., U.S. Patent 4,310, 651 (1982). 10. Baney, R. H., in Ultrastructure Processing of Ceramics, Glasses and Composites (L. L. Henchand D. R. Ulrich.eds.),JohnWileyandSons,NewYork, 1984, p. 235. 11. Baney, R. H., Gaul, J. H., Jr., and Hilty, T. K., Organometallics, 2, 859 (1983). D., Adv. Inorg. Chem Radiochem., 77, 349 12. Fritz, G.,Grobe,J.,andKummer, (1965). 13. Yajima, S., Omori, M., Hayashi, J., Okamura, K., Matsuzawa, T., and Liaw, C., Chem. Lett.,551 (1976). 14. Yajima, S., Hayashi, J., and Okamura, K.,Chem. Lett.,521(1977). T., Nature, 273, 15. Yajima, S., Hasegawa, Y., andOkamura,K.,andMatsuzawa, 525 (1978). 16. Schilling, C. L., Jr., and Williams, T. C., Polym. Preprints, 25, 1(1984). 17. Schilling, C. L., Jr., Br. Polym. J., 18, 355 (1986). 18. Hasegawa, Y., and Okamura, K., J. Mater. Sci., 21, 321 (1986). 19. Yamamura, T.,Polymer Preprints, 25(No 3), 8 (1984). 20. Yajima, S., Hayashi, J., and Okamura, K., Nature, 226, 521(1977). 21. Baney, R., Gaul, J., Jr., and Hilty, T., Mater. Sci. Res., 17, 253 (1984).

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22. Verbeek, W., U.S. Patent 3, 853, 567 (1974). 23. Penn,B.G.,Ledbetter,F. E., m, Clemons,J.M.,andDaniels,J. G.,J. Appl. Polym Sci., 27, 3751 (1982). G. H.,in UltrastructureProcessing of Ceramics, 24. Seyferth,D.,.andWiseman, Glasses and Composites (L. L. Hench and D. R. Ulrich, eds.) , John Wiley and Sons, New York, 1984, p. 265. 25. Arai, M., Sakurada, S., Isoda, T., and Tomizawa, T., Polym. Preprints, 28, 407 (1987). 26. Gaul, Jr., J. H., U.S. Patent 4, 340, 619 (1982). 27. Okamura, K., Sato, M., and Hasegawa, Y., Ceram. Int., 13, 55 (1987). F’roc. 1st Japan Intema28. Kimura,Y.,Hayashi,N.,Yamane,H.,andKitano,K., tional SAMPE Symposium, November 28 to December l, 1989. 29. Paine, R. T., JAPAN-US Joint Seminar on Inorganic and Organometallic Polymers, Nagoya, Japan, March 25-27, 1991, p. 78. 30. Hashimoto, N., Sawada, Y., Bando, T., Yoden, H., and Deki, S., J. Am. Ceram. Soc., 74, 1282 (1991). 31. Amato, C., Hudson, J., and Interrante, L. V., Mater. Res. Soc. Symp. Proc., 168, 119 (1990). 32. Atwel, W. H., Bums, G. T., and Zank, G . A., private communication PolymerPreprints, 33. Okamura,K.,Sato,M.,Matsuzawa,T.,andHasegawa,Y., 25(No l), 6 (1984). 34. Okamura, K., Sato, M., and Hasegawa, Y., 6th World Congress on High Tech Ceramics (CIMTEC), Milan, Italy, June 23-28, 1986. 35. Okamura, K.,Sato, M., Seguchi, T., and Kawanishi, S., Proceedings of the Third 21-24,1990, InternationalConferenceonCompositeInterface(ICCI-III),May p. 209. 36. Sasaki,Y.,Nishida,Y.,Sato,M.,andOkamura,K., J. Mater. Sci., 22, 443 (1987). 37. Shimoda, M.,Sugimoto,M.,Katase,Y.,Okamura,K.,andSeguchi,T., Muki Kobunsi Kenkyu Touronkui, 10, 76 (1991). 38. Singer, L. S., and Lewis, I. C., Appl. Spectrosc., 36(No l), 52 (1982). 39. Okamura,K.,Sato,M.,andHasegawa,Y., Proc. Fifth Int. Conf. on Composite Mater., San Diego, CA, July 29 to August 1, 1985, p. 535. 40. Yamamura, T., Hurushima, T., Kimoto, T., Shibuya, M., and Iwai, Y., 6th World Congress on High Tech Ceramics (CIMTEC), Milan, Italy, June 23-28, 1986. 41. Isoda, T., in Controlled Interphases in Composite Materials (H. Ishida, ed.), Elsevier, Amsterdam, 1990, p. 255. 42. Milewshi, J. V., Sandstrom, J. L., andBrown,W. S., in Silicon Carbide (R.C. Marshall, J. W. Faust, and J. R. C. E. Ryan, eds.), University of South Carolina Press, 1973, p. 634. 43. Nishihara, Y.,Kagaku to Kogyo, 40, 309 (1987).

Part VI PROCESSING OF SPECIALTY CERAMICS

This Page Intentionally Left Blank

17 Processing of Lead-Based Dielectric Materials Hung C. Ling AT&T Bell Laboratones, Princeton, New Jersey

Man F. Yan AT&T Bell Laboratories, Murray Hill, New Jersey

1.

INTRODUCTION

Historically, BaTiOs-basedmaterials have been the dielectricsof choice for use in multilayer ceramic capacitors. BaTiOs has a perovskite crystal structure and exhibits ferroelectric behavior. At its Curie temperature, the dielectric constant E may exceed 20,000. The effect of dopants on Curie point, microstructural control, and other processing parameters, such as powder conditioning (milling and solvents), green sheet processing (binder or casting), and firing temperatures and ambients, are reasonably understood. Multiple-cation substitution in the BaTiOs-based system has also been studied extensively withregardto preservation and cation ordering in the perovskite phase. In the 1950s, Smolenski and coworkers [ l 4 1 investigated many cation substitutions into PbTi03 in a search for new ferroelectric materials. In this kind of substitution, the general guidelines are that the ionic sizes should be comparable to those of W + ion and the combination must yield the same average charge as Ti& to maintain charge neutrality. Many such compositions take on the complex perovskite structure, and their properties have been compiled in Ref. [5]. With the promise of higher dielectric constants and lower firing temperatures, and hence lower electrode cost, lead-based perovskite compositions of the general form Pb(B1, &)Os have been studied extensively for application as capacitor dielectrics for the past 15 years. Of particular interest are compositions in which B1 is a divalent or trivalent cation such as Mg2+, Zn2+,Ni2+, 397

LING AND YAN

398

or Fe3+, and B2 is a high-valence ion, such as Ti&, M S + , Ta5+, or W+.The most commonly studied is Pb(Mg1/3Nb2/3)01/3 (PMN). These compounds exhibit a broad maximum in the dielectric constant, and the temperature of the dielectric maximum also increases with the testing frequency (Fig. 1). These are also known as the relaxor compounds. The origin of this broad maximum ‘is postulated as caused by a distribution of Curie points resulting from microcompositional fluctuation in theB-sitecations.Onlytwo compositions have been found to show ionic ordering by thermal annealing, Pb(SclnTaln)O3 and Pb(SclnNbt/2)03 [6,7].In these compounds, thedielectric response varies from the normal behavior in the ordered state (as in BaTi03) to relaxor behavior in the totally disordered state. Another characteristic of relaxors is the frequency dispersion inthedielectricloss (DF), which occurs at a slightly lower temperature than thedielectric maxima. The dielectric loss is also slightly higher than that of the normal ferroelectrics. Beginning with the work by Ohno and Yonezawa on PFN-PFW* systems in the late 1970s [ 8 ] , many multicomponent dielectric systems have been evaluated and put into manufacture. Some of the patented compositions developed for multilayer capacitor (MLC)applicationwererecentlysummarized by Shrout and Dougherty [9].Other compositions were developed for piezoelectric sensors and electrostrictive actuator applications [lo].Most of the compositions used for capacitor dielectrics are based on PFN [S],PMN [ 11-14], or PZN [15]. An important issue in the preparation of the complex perovskite compositions is the appearance of a cubic pyrochlore phase, whichmay occur as a dominant or minor phase co-existent with theperovskite phase and may appear or disappear depending on the processing conditions. The pyrochlore is more likely to occur, usually in higher proportion when using the conventional processing method of ball milling the starting oxide powders. Since ferroelectric behavior is exhibited in the perovskite phase, the pyrochlore phase is generally perceivedas detrimental toattainingthedesirableproperties of a pure perovskite phase.Atlow pyrochlore concentrations, an approximate lineardecrease in E with pyrochlore content has been observed [13,16].Figure 2 shows that in the series of compositions we studied [13],the decrease in the dielectric constant E follows the relation E ,,

=E0

-@

(1)

where E m a = 17,200,k = 1.6 x 105, and p = volume fraction of pyrochlore. Compositional modification and/or alternative processing methods, usually spe-

LEAD-BASED DIELECTRIC MATERIALS

399 +7949: R3670-4 Pbl-.-yMQ;NbZyO1+4y x = 0.2651; y =0.1566

W 4000Q

0 -40

I

I

-20

I

0

i

I

20

I

I

40

I

I

TEMPERATURE ( C 1

60

l

l

80

(a)

Figure 1 Dielectricconstant anddissipationfactorversustemperatureatdifferent frequencies in a PbO-MgO-Nb205 composition. (From Ref. 13.)

1 0

400

LING AND YAN

Figure 2 Maximum dielectric constants at 1 kHz versus pyrochlore concentration in the PbO-MgO-Nb205 ternary composition. (From Ref. 13.)

cific to a particular composition, may be developed to eliminate the pyrochlore phase from the end product. On the other hand, nonferroelectric pyrochlores may serve as technologicallyusefuldielectricsin other applications,suchas temperature-stable dielectrics or microwave dielectrics. It is important to understand the structural relationship between the perovskite and pyrochlore phases in the Pb-based systems and to elucidate the thermodynamic and kinetic factors that mayinfluence the phase equilibrium. Understanding these willallow a more systematic method of tailoring dielectric compositions to applications. Another important area in dielectric processing is the formation of dense, sintered bodies and theformation of thin layers (<25 pm) in the green (unfired) state for the fabrication of multilayer structures. Densification and microstructural development during ceramic processing are influenced by the characteris-

LEAD-BASED DIELECTRIC MATERIALS

401

tics of the starting powder and the microstructure of the green compacts. In a review paper published 10 years ago, Yan [l71 discussed the factors impacting on the properties of sintered electronic ceramics. In general, small particle size and narrow particle size distribution are required for densification to full density. Furthermore, deagglomeration of very fine powders before forming green compacts is also extremely important. It was shown [l81 that BaTiOs with a very small particle size ended up with a low sintered density because of an inhomogeneous agglomeration in the fine powders. Defect-free green compacts with uniform microstructure before sintering are necessary for sintered ceramics with high mechanical strength and good electrical properties. Flaws in the green compacts will most likely be magnified during the sintering process. The origin of these flaws may include foreign particles in the milling process, binder accumulation and incomplete burnout, and lamination defects in multilayer structures [19]. In electronic ceramics these flaws can later become the initiation sites for failure, such as electrode migration under temperature-humidity-bias conditions [20,21],or physical cracking under surface mount assembly processes. Other than the desirable property of high dielectric constant, the technology to form very thin green sheets of dielectrics and to cosinter with metals at low temperatures (
II. PHASE STABILITY In a study [22]on the crystal structures of perovskite and pyrochlore in compositions nominally Pb(B2+, B5+)03 and Pb(B3+, B5+)03, where B5+ was either Nb or Ta, we concluded that the pyrochlore phase has the B2+ or B3+ ion inherently incorporated as part of the structure. We also concluded that both the perovskite and pyrochlore phases are essentially made up of the same structural unit of (B’, B”)O6 octahedra. In the pyrochlore, the rigid octahedral framework

LING AND YAN

402

is stable, with a substantial deficiency in oxygen. If the average charge on the B site differs from 4, one can anticipate a corresponding change in the number of oxygen ions occupying the seventh site. The total oxygen ions can vary between 3.0 and 3.5 without affecting the stability of the pyrochlore structure, but changes in the lattice distortion and/or lattice parameter might be anticipated. Thus the general formula of the pyrochlore phase can be expressed as Pb(B’Jl”1 - &3OP, where 0 I p I 0.5. The incorporation of B’ cation in the Pb niobate pyrochlore structure was reported in the PbO-MgO-Nb205 system [16,23-251. Using the Pb0-ZnO-Nb205 system, we investigated the structural variation in the pyrochlore phase using x-ray diffraction (XRD).Furthermore, we developed aseries of low E dielectrics utilizing the pyrochlore structure [26]. The relative stability of the pyrochlore and perovskite phases depends on the specific combination of cations and processing temperatures. In one study [22] inwhich mixing andmillingof oxide powderswas the standard processing method, Nb tended to stabilize the perovskite phase more effectively than Ta as the B5+ cation. Within the Nb or Ta family, Zn, Ni, and A1 cations stabilize the pyrochlore phase, COcation favors the formation of perovskite, andMg and Mn cations result in about equal amounts of pyrochlore and perovskite regardless ofthereactiontemperature.WithFe3+ cation, the perovskite phaseincreases with calcination temperature, reaching 100% at 1100°C. Thus, the divalent and trivalent cations play an important role in determining the stability of competing pyrochlore and perovskite phases. Since the nominal compositions in this study were chosen to correspond to the perovskite structure, one suspects that kinetic barriers exist suchthata metastable pyrochlore phase forms instead of the more stable perovskite. If this is true, raising the calcination temperature should promote perovskite formation. This apparently is the case in Pb(FelnNbln)03 and Pb(Fe1/2Taln)03). However, other cation compositions do not show such a change. In fact, the PbO-ZnO-NbO5 composition usually forms 100% pyrochlore, which persists up to the melting temperature. Detail studies on the reaction sequences in the PbO-Fe203-Nb205-WO3 and PMN systems [27-301 suggested that the perovskite phase was formed through a sequence of reactions involving several lead-niobate pyrochlores. From extensive differential thermal analysis and XRD analysis, Inada [30] proposed the following reactions in PMN: 3PbO + 2Nb205 Nb,Pb,

0,

+ PbO

-

Pb3Nb4013

Pb, 2

Nb, Ol

Pb2Nb201+ )5 MgO + perovskite PMN + Pb3Nb4OI3

(2) (3)

(4)

He suggested repeating the process of calcination and ball milling to control PbO evaporation and obtain single-phase perovskite PMN. Slightly different

MATERIALS DIELECTRIC LEAD-BASED

403

reaction sequences were reported by Lejeune and Boilot [28] and Swartz and Shrout [27], which involved different forms of cubic or rhombohedralpyrochlore. In our own study, we have encountered only cubic pyrochlores. Similar reaction sequences canprobably be established for compositions with other cation combinations as well. However, it is not clear whether a generalized reaction sequence can be established for complex Pb compositions. There have been other attempts to generalize the tendency to form the pyrochlore phase in complex Pb compositions through correlations in ionic size, ionic bonding, and so on [31,32]. Two parameters were used: the Goldschmidt tolerance factor [33], defined as I=

‘Pb rO l

f i r B + ro

where rPb, rB, and ro are the respective ionic radii, and the electronegativity difference, defined as 2 where Xpb-0 is the electronegativity difference between Pb2+ and oxygen and XB - o the electronegativity difference between B-site cations and oxygen. Both rg - o and XB - o are calculated using the weighted average of B-site cations. Figure 3 is a plot of electronegativity difference versus tolerance factor for some AB03 compositions. Compositions withahigh electronegativity difference (-2.3) readily form a stable perovskite structure [BaTi03(BT), SrTi03(ST), and CaTi03(CT)] as the tolerance factor varies between 0.97 and 1.06. For Pb relaxors, the electronegativity difference is smaller (1.7-2.0) and the tolerance factor is between 0.94 and 1.00. To a first approximation, Fig. 3 suggests a sequence in which the perovskite formation is as follows: PCN < (PZN, PMN, PNN) c (PFN, PFW) < P T

(7)

Experimentally, Shrout and Halliyal [31] found that the ease of fabricating perovskites follows the sequence

PZN
(8)

Our results are represented by a slightly different sequence:

PZNcPNNcPMNcPCN
(9)

Figure 3 cannot satisfactorily explain either of these experimental sequences. Phase stability in lead-based systems is further complicated by forming solid solutions among relaxors (such as Pm-PFW) or with other ferroelectrics, such as BaTiOs. Since BaTi03 readily forms the perovskite phase, a number of di-

404

LING AND YAN

1.5

094

0.96

090

100

l02

I04

1.06

TOLERANCE FACTOR (11

Figure 3 Plot ofelectronegativitydifferenceversustolerancefactor ovskite compounds. (From Ref. 31.)

for some per-

electric systems have been developed empirically by adjusting the concentrations so that a pure perovskite phase is formed. Typical examples are adding BaTiOs to PZN or PNN. A summary of these compositions can be found in Shrout and Dougherty [ g ] . It is likely that thermodynamic guidelines cannot be relied on exclusively to select the appropriate cation combinations that will result in pure perovskite phase with high E. The development of Pb-based compositions is still based on an empirical approach of trial and error, but it may be useful to investigate the kinetic factors during ceramic processing to provide a scientific base for the formulation of lead perovskite dielectrics. Quantitative measurements of the diffusivity and solubility limits in these perovskite systems will be crucial for these scientific understandings. The formation of pyrochlore is also enhanced in PbO-deficient compositions, which is usually a result of the volatization of PbO during calcining or sintering. This can be seen in samples of stoichiometric compositions sintered above the optimum sintering temperatures. In other cases, the exterior surfaces may be richer in pyrochlore than the interior of samples, when the rate of densification of ceramics occurs faster than the evaporation of PbO. Thus it is important to control the sintering atmosphere or to adjust the compositions to account for PbO volativity.

LEAD-BASED DIELECTRIC MATERMS

111.

405

PROCESSINGALTERNATIVES

The propensity to form the pyrochlore phase in complex Pb-based compositions has led to the development of three alternative processing routes to the conventional mixed oxide method. The first method aims at achieving finer particle size and more homogeneous mixing of the starting materials to eliminate preferential reactions between PbO and Nb205 to form the lead niobate pyrochlores. The second route uses a two-step process in which Nb2O5 is first reacted with the other cation oxides to form a columbite (such as MgNb206) or wolframite (such as B1B204) phase, which is then reacted with PbO to form the perovskite phase. The third route uses excess PbO in the starting compositions, both to compensate for the volatization and to serve as a flux for lowering the sintering temperatures.

A.

Homogeneous Mixing

Inada [30] first proposed that to obtain single-phase PMN, it is necessary to repeat theprocess of calcining and milling a number oftimes. Lejeune and Boilot [28] reported the effect of various processing parameters on the compositionandnumber of phases formed in the traditional process of milling oxide powders. Replacing MgO by MgC03 as the starting material, repeating calination and milling up to three cycles, and replacing water by acetone as milling medium all increase the fraction of perovskite [volume fraction of perovskite defined asl(llO)pero/(l(llO)pero + I(222)pyrol from about 50 to 80%. These changes improveddispersability and reactivityamongthepowders. Other studies have used high-energy attrition or vibratory milling to produce similar improvements. The use of chemical processing (sol-gel) toincrease the homogeneity of powders have also been demonstrated [34,35]. Dielectric powders in the PMN-PNN-PT ternary system were synthesized by a hydrolysis and solution process. Raw materials of Pb, Nb, Ti, Ni, and Mg cations in the form of oxides, acetates, and alkoxides were mixed in an organic solvent. The resulting mixture was hydrolyzed to yield mixed-metal hydroxides. These hydroxides were washed with water, milled, calcined, and milled again. The aspreparedpowdersareshowninFig. 4a, versus conventionally prepared powders,showninFig.4b. The powders from the hydrolysis route have a smaller particle size, narrower size distribution, and less agglomeration than conventional mixed oxide powders. This permits the composition to be sintered at
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h

"

Figure 4 Scanningelectronmicrographsofdielectricpowderprepared alkoxide process and (b) a conventional process. (From Ref. 34.)

by (a) an

The sintered compacts from the solution-prepared powders are also chemically more homogeneous, withvery little variation in Mg or Ni distribution (Fig. 5). Previous reports showed particles of MgO distributed among the perovskite matrix. However, it is somewhat surprising that these authors were able to obtain only 93% perovskite phase via either the solution or the columbite route.

B.

Columbite Intermediate Phase

The columbite route prereacts MgO with m205 to form the columbite MgNb206 [27,31], which is then mixed and reacted with PbO:

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Table 1 Comparison of Solution-Prepared and Mixed OxidePrepared PMN-PNN-PT Powders

MixedProcess Solution Property 875°C 93%

Reaction complete Perovskite 93%' content Sintering temperature Green sheet density Sintered size Grain homogeneityb Chemical Peak E 12,000

1200°C <1000°C 10-15 pm 2-3

pm 4.5% 15,000

950°C 25 P 97.5%

5 P 2-10%

'Data on peroskite content was given from powders from the solution route and the columbite route (not simple oxide mixing and milling). Themical homogeneity measured bystandard deviation in elemental XRD intensity from 48 different spots for Mg, Ni, Nb, and Ti; largest variations are seen in Mg and Ni.

This is a two-step process, with a secondary milling step to mix MgNb206 with PbO for a second calcination before sintering. A systematic investigation of the role of different cations to form an intermediate phasewith m205 [36,37] shows that either the columbite (BNb206) or wolframite (BNb04) forms in the temperature range of 700-1000°C. The reaction temperature Tr to form 100%

o Alcoxlde Process A

Y

Mg

N ITI .

Oxide Process

Nb

ELEMENT

Figure 5 Elementalanalysis of sinteredspecimenmadeusinganalkoxideprocess and a conventional oxide process. Each symbol represents an average 16 of spots in one area on a polished specimen. The alkoxide powder shows greatly reduced variation in elemental intensity, especially in Mg and Nb. (From Ref. 34.)

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Table 2 Formation Temperature of Columbites and % Perovskite Formed at 900°C % Perovskite at 900°C

Columbitdwolframite MgNb206

NiNbzOa MIlNb206 ZnNb206 CrNb04

AIM04

99

Tr

Oxide Columbite Mixed

700

43 0

55 73 43 0

1000 1000

0 0

Incomplete

900 900 900

91

-

columbite or wolframite is shownin Table 2. On subsequent reactionwith PbO, this processing method resulted in nearly 100% perovskite phase in PMN and PNN at 9OO"C, versus a perovskite content of less than 50% by the mixed oxide technique. WithPMnN,the columbite route showed no improvement. Compounds containing ZnO, C1-203, and A1203 form either columbite (ZnNb206) or wolframite (CrNbO4 andAINb04) in reaction with Nb205; however, only the pyrochlore phase is formed when the precursor is reacted with PbO. The MgNb206 columbite phase reacts with PbO to form almost 95% perovskite at 700"C, with some pyrochlore formed between 500 and 600°C. The phase concentration in PMN as a function of calcination temperature is shown in Fig. 6. The reaction between NiJYb206 and PbO frst forms nearly 100% pyrochlore phase between 600 and 800°C. which then transforms into the perovskite phasebetween 800 and900°C(Fig.7). Since theprocessoftransforming pyrochlore into perovskite is quite sluggish, this may account for a lower perovskite concentration in PNN (91%) compared with PMN (99%), in which the perovskite is formed directly from the reaction between PbO andthe columbite phase. We also found that ZnNb206 and MgNb206 form a continuous solid solution of ZnxMgl-xNb206. For x < 0.5, subsequent reaction with PbO results in 99% perovskite phase at 1000°C [37]. Thus in two relaxor systems,PMNandPNN, the columbite processing route hasconsistentlybeen shown to increase substantially the amount of perovskite phase compared with the conventional mixed oxide route. This processcan also be usedto form nearly 100% perovskite in the Pb0-ZnxMg1-x0-Nb205 system, where x < 0.5. However, it is also clear that the columbite method cannot be applied to increase the perokvskite phase of all relaxor systems. Crystallographically, the cations in the columbite precursor occupy the octahedral sites of a hexagonal close-packed assembly of oxygen ions. The precursor also has a cation to anion ratio of 1:2, identical to that of TiO2. Since

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TEMPERATURE ('C 1

Figure 6 Phase concentration as a function of calcination temperature when PbO is reacted with MgNb206 columbite phase. (From Ref. 36.)

these cations occupy the octahedral sites in the perovskite structure, the reaction between the columbite precusor and PbO readily results in the perovskite structure. Kinetically, the prereaction between m205 with the other cations also reduces the reaction rate between PbO and m205 to form pyrochlore, which has been shown to be sluggish in converting to perovskite later. This is seen withPMN. For PNN, the explanation is not so clear. The columbite NiNb2O6 forms readily, but it reacts with PbO to form the pyrochlore phase first, which then transforms into perovskite. The Zn2+ cation favors fourfold coordination (tetragonal site), as in wurtzite ZnO, and it is impossible to pre-

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41 0

TEMPERATURE (OC)

Figure 7 Phase concentration as a function of calcination temperature when PbO is reacted with NiNbz06 columbite phase. (From Ref. 36.)

pare perovskite PZN via the ceramic route. This discussion indicates an interplay between the thermodynamic tendency ofcertain cations to form perovskite and the kinetic factors, whichmay be controlled through clever processing methods to enhance perovskite formation in some systems.

C. Use Of Excess PbO/MgO Although the columbite method can improve the perovskite concentration in PMN, the sintering temperature is still > 1200°C. It is desirable to reduce the sintering temperature so that less expensive electrode materials can be used. Excess PbO is commonly used as a sintering aid and, for relaxor compositions, to compensate for the volatization of PbO during the sintering step. Lejeune and Boilot [28,38] investigated the addition of excess PbO after the calcination

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41 I

step to PMN, PMN-PT, and PMN-PZN systems. With 6-8wt% excess PbO, PMN ceramics are sintered to 96% of theoretical density at 900"C, instead of 87% at 1000°C with no excess PbO. The pyrochlore is also reduced from 8% at 1000°C to 2% at 900°C with the excess PbO. As a result, the dielectric constant is increased from 9000 to 15500. Similar beneficial effects are shown in the other systems. We prepared and investigated the properties of ceramics in the PbO-MgOm205 system near the PMN composition [13,14]. The compositions are of the general form tPbTi03 + [(l - t)/(l - x - y ) ] (Pb1+,Mgx Nbzr01+4y), where 0.20exe0.28,0.14eye0.20,andOete0.10.Forx=y=0.2,thecomposition becomes tPT+ (1 - t)PMN. We prepared these compositions by a simple method of mixing and reacting the constituentoxide powders, and the extra processing step to form the columbite MgNb206 precursor is not necessary to attain favorable propertiesin these compositions. A rangeof compositions were identified that can be sintered at temperatures between 950 and 1000°C. Furthermore, these compositions have a much reduced pyrochlore concentration and thus a higher maximum dielectric constant (>18,000). Our approach is different from that of Lejeune and Boilot [28] in that excess PbO and MgO are incorporated in the starting compositions. A standard processing sequence is as follows. Each batch of oxide and carbonate powders, with an approximate weight of 850 g, was mixed with about 3.2 kg ZrO2 milling media and deionized water. Darvan C dispersant, 1% by weight of the powders, was added to decrease the viscosity of the slurries. The powder slurries were milled for 4 h andthenfiltered and dried inanoven overnight at 110°C. The driedpowderswere granulated through 20 mesh screen (850 pm opening) and then placed in covered silica boats and calcined at 750°C for 4 h. The calcined powders were granulated through a 20 mesh screen againbut were notmilled before compaction andsintering. Pressed disks were sintered at a temperature ranging from 900 to 1050°C for 2-8 h in a flowing oxygen atmosphere. Figure 8 shows the pyrochlore concentration in the ternary diagram. Compositions with pyrochlore concentration e 1% were prepared at 0.24 e x e 0.27 and 0.15 e y e 0.18. Figure 9 shows the dielectric constant and dissipation factor at 1 kHz versus temperature in five selected compositions with no PT ( t = 0). The dielectric constant have a broad maximum at around -5°C. The magnitude of the maximum depends on the pyrochlore concentration in these compositions. At a temperature slightly (15-20°C) below the Curie temperature, the DF shows a maximum value. At higher temperatures above the Curie temperature, DF is much smaller (~0.005). The addition of various components to the lead magnesium niobate compositions was found to be advantageous for the dielectric properties. Addition of PbTiOs is commonly used to raise the Curie temperature ofPMN composi-

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PYROCHLORE CONCENTRATION

Figure 8 Pyrochlore concentration in the PbO-MgO-Nbz05 ternary diagram at 980°C. (From Ref. 13.)

tions. Figure 10 shows the shifting of Curie temperature and the increase in E from 18,000 at t = 0 to 27,100 at t = 0.05 and to 29,000 at t = 0.10. Addition of PCN improves the dielectric loss. In the undoped composition, DF increase sharply on cooling below Tc, increasing from a value of 0.01 to 0.1 in a narrow temperature range of 20°C. Addition of PCN reduces the rate at which DF increases below Tc and also reduces the maximum DF to 0.06 over the entire temperature range. In these series of compositions, excess PbO and MgO are incorporated in the starting materials. It is important to understand how the nonstoichiometric starting compositions affect the final product in composition and in properties. We analyzed [39] a series of selected compositions by varying x and y in Pbl-Mg&b2yOl+4y through measurement of weight changes during sintering. Our data show that the PT and PCN components do not contribute to weight loss at 1000°C. We attributed the total PbO loss to thenonstoichiometric PMN components. From the weight loss measurements,we calculated the final molar ratio of PbO/MgO/Nb205 of the PMN component in the sintered composition.

413

LEAD-BASED DIELECTRIC MATERIALS 20000

+7949

-

0 -40

l

-20

I

1

0

20

I

40

I

60

1

80

-

I

XI

TEMPERATURE ( C )

(a) 0.12I

0

L

-

0 - , -40

(b)

1 , -20

1 0

,

I I fI I I 80 20 40 60

100

TEMPERATURE ("C)

Figure 9 (a) Dielectric constant and (b) dissipation factor versus temperature in five selected compositions in the PbO-MgO-NbzOs ternary diagram. The dielectric properties were measured at 1 kHz and 1.0 V. The magnitude of the maxima depend on the pyrochlore concentration. (From Ref. 13.)

Figure 11 shows that all the final compositions converge on the line that represents excess MgO added to the stoichiometric PMN (point A). For another series of compositions B-H, we maintained the relative concentration between Mg and Nb and varied the concentration of PbO. After sintering, compositions B-H all converged to the same final composition. These results indicated that

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414

TEMPERATURE (Cl

(b)

TEMPERATURE ( C 1

(b) Figure 10 (a)Dielectricconstantand (b) dissipationfactorinPMNcompositions 980°C in 0 2 for 4 h. with different PbTiOs concentrations. Samples were sintered at (From Ref. 13.)

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Figure 11 TernaryPbO-MgO-NbzO5phasediagramshowingthatthestartingcompositions (indicated by numbers and letters) all converge on the line linkingA and composition 10. This line represents stoichiometric PMN + excess MgO. (From Ref. 39.)

although the excess PbO in the initial compositions significantly influenced the densification, it was all evaporated after sintering. Measurements of dielectric constant and dielectric loss have not reveal any detrimental effect caused by residual PbO. The cubic pyrochlore can be eliminated and densification can be enhanced by the addition of excess PbO or MgO, but it is important to understand the processes involved so that reproducible results can be obtained for each formulation. Otherwise, variations in physical and dielectric properties may result even for nominally the same composition. There has been some concern that in using excess PbO, residue may be left in the grain boundaries [31], which may reduce E, cause mechanical weakness, and increase aging. It is important to ensure that proper processing eliminate the excess PbO through volatization, after the useful functions of enhanced sintering and prevention of pyrochlore formation have been performed. There has been a report in the literature of a PbO-rich grain boundary phase [24], but the specific sample was sintered at 900°C with 12% pyrochlore. Thus, the result is unlikely to be representative of the grain boundary structure of properly sintered relaxors. We reported on a

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transmission electron microscopy study of low-firing lead-magnesium-niobate ceramics with nonstoichiometric starting compositions [40]. Most importantly, x-ray spectrometry did not reveal a PbO phase in the grain boundary regions. The ceramic dielectrics exhibited a matrix with large (m7500 A) crystalline inclusions. The interplanar spacings of the matrix matched closely the x-ray diffraction values for the composition Pb(MgmNbm)Os The inclusionswere found to be rich in Mg, with small and varying amounts of cobalt, and the interplanar spacings matched closely the composition (Mg, Co)O. The compositionchanged abruptly to the matrix composition in crossing the inclusion/ matrix boundary with no interfacial variations. These observations are consistent with our results obtained from weight loss measurements, and we concluded that stoichiometric PMN is formed from the nonstoichiometric starting compositions, the excess MgO appearing asinclusionsandthe excess PbO vaporizing during sintering. The use of excess PbO to achieve densification is quite common in the sintering of PZT. Accurate control of the atmosphere in binder burnout and sintering is important to assure reproducible electrical, dielectric, and piezoelectric properties, The same understanding of processing and control is needed in the manufacture of relaxor dielectric materials and relaxor MLC using the excess PbO route. Successful implementation can result in MLC reliability equal to that of BaTiOs MLC [41].

IV. CONCLUSION The subject of lead-based dielectrics has occupied many researchers during the past 15 years. A large body of knowledge has been published in the literature on compositions, processing, and properties in thecomplex lead perovskite systems. A fundamental understanding of the stability relationship between the perovskite and pyrochlore crystal structures is critical to the successful preparation and utilization of these dielectrics in electronic components. Although attempts have been made to differentiate thermodynamic criteria from kinetic factors in determining the formation of the perovskite phase, no definitive relation can be relied upon for predictability. However, these studies have provided guidelines from which the empirical method of trial and error is used to develop compositions for practical applications. It is important to realize that although the perovskite structure is responsible for high E dielectrics, other low-firing and temperature-stable dielectrics can be successfully developed from the pyrochlore phase. The effectof forming a range of compositions from two or more end components (such as PFN-PFW) and the addition of a small amount of dopants (such as Mn and Ag) on processing and properties represents an area in which a separate review is desirable. In this chapter, we only briefly discussed the effect of adding PbTiO3 and Pb(C01/3Nby3)03.

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41 7

Focusing on alternative processing routes to the conventional mixed oxide method, we discussed homogeneous mixing and particle size reduction through mechanical or chemical (solution) means, forming the intermediate columbite phase before reacting with PbO, or adding excess PbO and/or MgO to thestarting compositions. Homogeneousmixingandparticle size reduction are germane to being able to form thin green layers in a multilayer configuration and should be viewed as applicable to all relaxor compositions. To a large extent, the method of columbite intermediate phase or excess PbO was developed in response to the challenge of forming a pure perovskite phase in the relaxor compositions. The columbite route is effective for PMN and PNN, with some applicability to thePZN-PMN system, inattaining 100% perovskite phase; however, the sintering temperature remains high, at around 1200°C. It may be possible to lower the sintering temperatures by combining mechanical attrition with the columbite method. Adding excess PbO to the starting compositions can achieve both lower sintering temperatures and a pure perovskite phase. It is important to control the processing steps such that all excess PbO is vaporized after sintering. This reduces variations in properties and results in electronic components meeting the performance and reliability requirements.

ACKNOWLEDGMENTS We thank W. W. Rhodes of AT&T Bell Laboratories at Murray Hill, New Jersey for supporting our work in the relaxor materials. Subsequent work on the processingandreliabilityof multilayer ceramic capacitors was achieved through manyuseful collaboration and discussions with our colleagues at AT&T Bell Laboratories Engineering Research Center at Princeton, New Jersey.

REFERENCES 1. Smolenski, G. A., A. I. Agranovskaya, A. I., Popov, S. N., andIsupov, V. A., Sov. Phys.-Tech. Phys., 3, 1981 (1958). 2. Smolenski, G. A., andAgranovchkaya, A. I., Sov. Phys.-Solid State, I , 1429 (1959). Sov. Phys.-Tech. Phys., 3, 1380 3. Smolenski, G. A.,andAgranovskaya,A.I., (1958). 4. Ismailzade, I. G.,Kristullograjiya, 5, 316 (1960). 5. Hellwege, K. H., and Hellwege, A. M., Vol. m/16a, Ferroelectrics and Related Substances, Lmdolt-Borstein (eds.), Springer-Verlag, 1981. 6. Stenger, C. G.,Scholten, F. L., and Burggraf, A. J., Ordering and diffuse transitions in Pb(ScTa)O3, Ceram. Solid State Commun., 32, 898 (1979). 7. Setter, N., and Cross, L. E., Thecontribution of structuraldisorder todiffuse phase transitions in ferroelectrics. J. Mater. Sci., 15, 2428 (1980).

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8. Yonezawa, M., Low-firing multilayer capacitor materials, Am. Cerum. Soc. Bull., 62, 1375 (1983). 9. Shrout, T. R., and Dougherty, J. P., Lead based Pb(BlB2)03 relaxors vs BaTiOs dielectrics for multilayer capacitors, in Cerumic Dielectrics: Composition, ProVol. 8 (H. C. Ling and M. F. Yan, cessing and Properties, Ceramic Transactions eds.), Am. Ceram. Soc., Westerville, OH 1990, p. 3. 10.Uchino,K.,Electrostrictiveactuators:Materialsandapplications,Cerum.Bull., 65, 4 (1986). 11. Swartz, S. L., Shrout, T. R., Schultz, W. A., and Cross, L. E., Dielectric properties of lead-magnesium-niobate ceramics,J. Am. Cerum. Soc., 67, 311(1984). 12. Lejeune, M., and Boilot, J. P., Optimization of dielectric properties of lead-magnesium-niobate ceramics,Am. Cerurn Soc. Bull., 64, 679 (1985). 13. Yan, M. F., Ling, H. C., and Rhodes, W. W., Preparation and properties of PbOMgO-NbzO5 ceramics near the PMN composition,J. Muter. Res., 4, 930 (1989). 14.Yan,M. F., Ling, H. C., and Rhodes, W. W., Effects of dopants on PbO-MgONb205 ceramics near the PMN composition, J. Muter. Res., 4, 945 (1989). 15.Yamashita,Y.,Furukawa, O., Kanai, H., Imai, M., and Harata, M., A low firing dielectricforMLCbasedonRelaxorBTceramiccomposite,inCeramicDielectrics: Compositions, Processing and Properties, Ceram. Transactions, Vol. 8 (H.C.LingandM. F. Yan, eds.), Am. Ceram. Society, Westerville, OH 1990, p. 33. 16. Chen, J., and Harmer, M. P., Microstructure and dielectric properties of PMN-pyrochlore diphasic mixtures, J. Am.Ceram. Soc., 73, 68 (1990). 17.Yan,M. F., Microstructuralcontrolintheprocessing of electronicceramics, Muter. Sci. Eng., 48, 53 (1981). 18. Carbone, T. J., and Reed, J. S., Microstructure development in BaTiO3: Effect of physical and chemical inhomogeneities, Am. Cerum. Soc. Bull., 58, 512 (1979). 19. Rice, R. W., Processing induced sources of mechanical failure in ceramics, Muter. Sci. Res., 11, 303 (1978). 20. Ling, H. C., and Jackson, A. M., Correlation of silver migration with temperaturehumidity-bias (THB) failures in multilayer ceramic capacitors, IEEE-CHMT, 12, 130 (1989). of electrode melting in multi21. Ling, H. C., and Chang, D. D., In situ observation layer ceramic capacitors, J. Muter. Sci., 24, 4128 (1989). 22. Ling, H. C., Yan, M. F., and Rhodes, W. W., Phase stability in Pb(B'2+B" S+)@ andPb(B'3+B" 5+)03 compositions, Ferroelectrics, 89, 69 (1989). 23. Shrout, T. R., and Swartz, S. L., Dielectric properties of pyrochlore lead magnesium niobate, Muter. Res. Bull., 18, 663 (1983). 24. Gm, E., Yamamoto, T., and Okazaki, K., Microstructure of lead-magnesium-niobate ceramics, J. Am. Cerum. Soc., 69, C188 (1986). 25. Chen, J., Gorton,A.,Chan,H.N.,andHarmer,M.P., J. Am.Cerum. Soc., 69, C303 (1986). 26. Ling, H. C., Yan, M. F., and Rhodes, W. W., Lead zinc niobate pyrochlore: Structure and dielectric properties, J. Muter. Sci., 24, 541 (1989). 27. Swartz, S. L., and Shrout, T. R., Fabrication of perovskite lead magnesium niobate, Mater. Res. Bull., 17, 1245 (1982).

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28. Lejeune, M., and Boilot, J. P., Formation mechanism and ceramic process of the ferroelectric perovskites: PMN and PFN, Ceram Int., 9, 119 (1983). 29. Kassarjian, M. P., Newnham, R. E., and Biggers, J. V., Sequence of reactions dur-

ing calcining of a lead-iron niobate dielectric ceramic,Am. Ceram. Soc. Bull., 6 4 ,

1108 (1985). 30. Inada, M., Analysis of the formation process of the piezoelectric PCM ceramics, Jpn. Natl. Tech. Rep., 27, 951 (1977). 31. Shrout, T. R., and Halliyal, A., Preparation of lead-based ferroelectric relaxors for capacitors, Am. Ceram. Bull., 66, 704 (1987). K., Di32. Furukawa, O., Yamashita,Y.,Harata,M.,Takahashi,T.,andInagaki, electric properties of modified PZN ceramic, Jpn. J. Appl. Phys., 24, 96 (1985). 33. Goldschmidt, V. M., Shrifier Norske Videnskaps-Akad. Oslo 1: Matemot. Naturuid. Klasse, No. 2 (1926). 34. Ochi, A., Utsumi, K., Mori, T., Yonezawa, M., Morishita, J., and Yoshimoto, T.,

A New Dielectric Material for Capacitors with High Specific Capacitance, CeF. Yan,eds.), Am. Ceram.Society, ramicTrans.,Vol. 8 (H.C.LingandM.

Westerville, OH 1990, p. 45. 35. Ochi, A., Mori, T., Nakanishi, M., Utsumi, 36. 37. 38. 39. 40.

K., Abe, S., and Yoshimoto, T., Y5V multilayer ceramic capacitors with high specific capacitance and low-equivalent series resistance, IEEE-CHMT, 8, 572 (1991). Yan, M. F., Ling, H. C., and Rhodes, W. W., unpublished research, 1987. Ling, H. C., Rhodes, W. W., and Yan, M.F. Method of preparing a ceramic composition, U.S. Patent 4,637,989 (1987). Lejeune, M., and Boilot, J. P., Low firing dielectrics based on lead magnesium niobate, Mater. Res. Bull., 20, 493 (1985). Ling, H.C., Yan, M. F., Jackson, A. M., and Rhodes,W. W., Effect of PbOevap oration on PbO-MgO-Nb205 based dielectrics, J. Mater. Res., 5, 629 (1990). Yegnasubramanian, S., and Ling, H. C., TEM studies of microstructure and compositionoflow-firinglead-magnesium-niobateceramics, J. Mater. Res., 7 , 197

(1992). 41. Chang, D. D., Electrical characteristics and reliability study of high-K lead-based

multilayer ceramic capacitors, in CeramicDielectrics:Composition,Processing and Properties, Ceramic Transactions, Vol.8, (H. C. Ling and M. F. Yan, eds.), Am. Ceram Society, Westerville, OH 1990, p. 88.

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Synthesis of Magnetic Particles Masataka Ozaki Yokohama City University Yokohama, Japan

1.

INTRODUCTION

Magnetic particles are of importance not only in industrial technology but also in our environment and in the functions of some biosystems as well as in scientific interest. Recently, magnetic particles have had wide applications in the biological and medical diagnostic fields. Ever since iron powders were employed for the magnetic recording system shortly before World War 11, extensive efforts have been made to produce better magnetic particles [1-41. Originally, their useof magnetic particles waslimited to audiotapes, but at present they are employed in a variety of forms for video and data recording in the form of tapes, cards, and flexible and rigid disks. The magnetic interactions between such particles are very strong, however, and stable dispersion is difficult to obtain. Small magnetic particles having a superparamagnetic nature are used for magnetic fluids [5,6]. Techniques are being developed to produce new magnetic particles and to introduce new functions to particulate materials by incorporating magnetic material. For example, Figlartz et al. developed a poly01 process [7]. Coating of magnetic particles by nonmagnetic materials, coating of nonmagnetic particles by magnetic material and a combination of enzymes on the surface of magnetic particles, for example, have beenaccomplished [8-lo]. Attempts have also been made to use magnetic particles in labeling and separation of cells [1l].

421

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OZAKI

In using fine magnetic particles, particle size is the most important parameter, as are other qualities, such as crystallinity and composition, since the magneticproperties of theparticlesare strongly influenced by these properties. Therefore, it is advantageous that the particles areformed uniformly in size and shape-monodisperse or in a narrow size distributioninthe desired size ranges. The preparation of colloidal particles with a narrow size distribution has been investigated by colloid chemists for a long time. However, the preparation techniques for magnetic particles were developed empirically as a result of strong industrial demands. The work by LaMer and Dinegar gave us some basic principles for the formation of monodispersed particles [121. According to their theory, to produce colloidal particles with a narrow size distribution, the nucleation and growth process must be carefully controlled. The conditions necessary for the formation of magnetic particles are essentially the same as for nonmagnetic particles, but some special precautions are necessary because of strong magnetic interactions among the particles. In producing monodispersed particles, the essential parameters are (1) sep(2) protection of aration of the nucleation process from the growing process, particles from aggregation, (3) a controlled supply of precursor material, and (4) temperature and pH in the solution. These parameters are intimately related, and sometimes it is difficult to separate them. Thus, the concentration of starting material, reaction temperature, and pH in the solution must be optimized. Dispersion reagents, such as surfactants and polymers, must be carefully chosen. A continuous supply of precursor material is sometimes made by the controlled decomposition of some other material. Readers may refer to the literature for articles on the formation of monodispersed particles [12-151. It is. known that magnetic particles can be found in the bodies of some biosystems. In 1975, Blakemore found that some bacteria have magnetic particles in their cells and navigate along geomagnetic fields using these magnetic particles [16]. It is also believed that certain animals have the ability to detect a magnetic field [17].

II. MAGNETIZATION CURVE AND SIZE OF THE MAGNETIC PARTICLE Any material can be magnetized more or less when it is placed in a magnetic field; therefore every particle shows more or less magnetic property. We confine magnetic particles that have proper permanent or induced magnetic moments that are strong enough to interact with other particles or strong enough to interact with an applied magnetic field with a magnetic interaction energy of the order of or larger than the thermal energy. Therefore, these magnetic particles show particular magnetic phenomena, such as aggregation, caused by magnetic interactions or orientation along applied magnetic fields.

SYNTHESIS OF MAGNETIC PARTICLES

423

When a magnetic material like iron is placed in a magnetic field, it is magnetized, as shown in Fig. 1, witha change in the applied magnetic field strength. Bulk material or powder is not magnetized originally unless it is exposed to a magnetic field, since the original directions of the magnetic dipoles in the magnetic domains of the bulk material or powder are random. Thus, the magnetization of a magnetic solid is brought about by the orientation of the dipoles. A s the magnetic field increases, the degree of orientation of the dipole increases and saturation magnetization is attained at perfect orientation of the dipoles. Even though the magnetic field decreases to zero, the material will have a certain magnetization, the so-called residual or remanent magnetization designated by M r in Fig. 1. A magnetic field of opposite direction is required to bring the magnetization to zero. This magnetic field is called the coercive force, frequently designated by Hc. Coercive force is essential for magnetic recording media, as is remanent magnetization. If we remove the magnetic field at the Hc, the magnetization of the material remains to some extent. A negative magnetic field slightly stronger than the Hc is required to bring the magnetization to zero. This magnetic field is called remanent coercive force. A s the size of the particle decreases, the number of magnetic domains in the particle decreases, changing the domain structure from multidomains to single domain. Generally, the Hc ofmultidomain particles is smaller thanthatof single-domain particles since the rotation of the magnetic moment in the former particles occurs easily from the domain walls. The coercive force of a single-domain particle is determined by so-called magnetocrystalline anisotropy together with shape anisotropy. The Hc caused by shape anisotropy increases with the increase in the aspect ratio if the particle size remains the same. Therefore, elongated single-domain particles are preferentially employed for magnetic recording media. A s we see later in this chapter, the particles available for magnetic media must be single magnetic domain particles having high saturation magnetization and proper coercive force.

Figure 1 Magnetizationcurve. Mr,remanencemagnetization; Hc, coerciveforce.

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OzAKl

If the particle becomes sufficiently small, the magnetic moment shows no preferential orientation as a result of the thermal agitation exhibiting some superparamagnetic property. Such particles have a very small coercive force and are not useful for magnetic recording material. Thus, small magnetic particles of a superparamagnetic nature are constituents of magnetic fluids. In a stable dispersion of fine magnetic particles, particles are free to rotate. The magnetic field H applied to a magnetic particle witha permanent magnetic moment m causes orientation along the direction of the field. The orientation may be disturbed by thermal agitation. The degree of the orientation can be related by the Langevin equation,

(

m,, = m cotwhere mVis the average magnetic moment and a = mH/kT, k being the Boltzmann constant and T the absolute temperature. Therefore, the magnetization of a stable dispersion of magnetic particles can be given by

M = nm,

(2)

where n is the number of the particles per volume.Thus, the dispersion of magnetic particles shows a magnetization-curve similar to that of paramagnetic material.

111.

MAGNETICINTERACTIONSBETWEENMAGNETIC PARTICLES

where m1 and m2 are the magnetic dipole moments of the particles, r is the vector joining thecentersof the twodipoles,and is themagneticpermeability of the vacuum. The magnetic dipole moment, assuming a single-domain particle, is expressed as m = Iov, Io being the magnetization per unit volume of the particle and v its volume. When two dipoles m1 and m2 are oriented in the same direction on the same line, Eq. (3) can be 'written as

where U0 is the minimal potential energy of the system. The magnetic interaction energy increases with the second power of its volume, .that is, propor-

SYNTHESIS OF PARTICLES MAGNETIC

425

tional to the sixth power of the radius. Therefore, the magnetic interaction energy increases rapidly as the size of the particle increases. The interaction between particles of strong magnetic material, such as magnetite or maghemite, is very strong, and stable colloidal dispersion is difficult to obtain unless the particle size is small, such as superparamagnetic particles. The magnetic interaction energy between superparamagnetic particles is reduced because the dipole moment in such particles has no preference for aparticular orientation as a result of free rotation of the magnetic moment by thermal agitation. Accordingly, the average interaction energy is expressed as U=-="

kT

x2 +-+1.54x104x6 7x2

6

1800

by Scholten and Tjaden [l91 for small values of x, where x = Uo/kT. For large negative values of x, 4x" 5 6 ~ - ~ u=x+2--+9 3 The magnetic interaction energy increases if the interaction works between multiparticles. The multiinteraction becomes important in a concentrated dispersion [20].

W. SYNTHESIS OF MAGNETIC PARTICLES FOR MAGNETIC RECORDING MEDIA Recently magnetic particles have been widely employed for recording media not only for audio and video recording systems but also for a variety of computer memories in the form of tapes, flexible and rigid disks, and cards. With the progress of these systems, magnetic particles having better recording performance have been pursued. Particulate magnetic recording materials are required to have (1) high coercivity, (2) largesaturation magnetization, (3)gooddispersibilityand orientability, (4) narrow distributions in size and shape, and (5) high chemical and mechanical strength [1-4,21,22]. The magnetic particles must have large saturation magnetization and proper coercive forcefor the magnetic recording media to ensure enough output signal strength and reliability of the storage of the signal. However, not many materials can satisfy all these demands. After the finding that elongated maghemite (y-femc oxide, y Fez03) particles showed superior magnetic property for recording performance, maghemite particlesattractedtheinterest of recording tape manufacturers. Since then, maghemiteparticleswerethemost popular material until cobalt-modified maghemite appeared. Recently, iron particles have been employed for 8 mm videocassette tapes, still camera disks, and high-quality audio tapes [23]. Much

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attention has been focused on fine barium ferrite particles for perpendicular recording materials as well as for use in videotapes to improve the recording properties [24,25].

A.

Aciculary-FerricOxideParticles

Since y-ferric oxide has a cubic structure, it cannot be directly formed in an elongated shape. Therefore, the particles are prepared from another kind of compound by solid-phase transformation. Although y-ferric oxide can be prepared from either P-FeOOH or y-FeOOH, most commercially available particles are produced from a-FeOOH (geothite) [21,22]. The original shape and size of a-FeOOH particles are kept almost the same throughout the dehydration and reduction process, followed by oxidation. For high-density recording, particles must be small, and particles are made to have a large aspect ratio to achieve a large Hc. Therefore, the size and shape of the starting a-FeOOH particles must be carefully controlled. However, theconditions and mechanism for the formation of oxihydroxides was not well known until Kiyama and Takada revealed them after extensive studies [26]. Today, goethite particles with desired sizes and shapes are available. They are described here. First, seed particles of goethite are prepared by the reaction Fe2++ OH-

+ KO, + a-FeOOH

The seed particles are then grown to desired sizes in the presence of iron and ferrous ions. Second, the goethite particles are dehydrated by heating, then reduced by hydrogeninto magnetite, andfinally oxidized into maghemite through the reactions heat a-FeOOH

H2

0 2

+a-Fe2O3 +Fe304_ jy-Fe20,

- 300°C 300400°C

- 25OOC

The temperatures and the length of the reduction and reoxidation time are optimized and carefully controlled. To achieve good dispersibility and good orientability, the surface of the particles is covered with nonmagnetic materials by chemical procedures before or after the reduction. The y-ferric oxide particles on the market are 0.2-0.5 pin length with aspect ratios of about 10. Typical y-ferric oxide particles are shown in Fig. 2a. When a-FeOOH particles are used for the starting material, it is very difficult to avoid pore formation completely in the dehydration process, which necessitates heating at high temperatures to decrease pores. It was expected that nonporous particles could be formed if elongated hematite particles were used for thestarting material. The first successful method for the production of elongated hematite particles was developed by Matsumoto et al. using hydrother-

SYNTHESS OF PARTICLES MAGNETIC

427

a

b

Figure 2 Transmission electron micrograph (TEM) of acicular y-Fe203 particles for audiotapes produced from a-FeOOH. Hc = 340 Oe. (b) E M of y-Fez03 particles modified by cobalt for VHS videotapes. Hc = 680 Oe. (Courtesy of Horiishi, Toda Kogyo Corp., Hiroshima, Japan.)

428

OZ4KI

mal reactions [U]. The maghemite particles prepared from the hematite particles were poreless and were more ellipsoidal thanother magnetic particles. According to Corradi et al., the good magnetic properties of the solids could be accounted for by the fact that the magnetic field lines in the particles were more parallel than in other particles [28]. They called the particles nonpolar (Np) particles. Maghemite particles with good magnetic properties could also be produced from spindle-type hematite particles prepared by the forced hydrolysis of ferric chloride solutions in the presence of small amounts of phosphate ions [29]. Arndt obtained maghemite particles of improved quality from acicular hematite particles prepared by doping tin in the hydrothermal transformation of precipitated Fe(OH)3 [30].

B. Cobalt-Modified y-Ferric Oxide Particles According to the theory of magnetism, an Hc of y-ferric oxide particles larger than 500 Oe is obtainable, but the highest Hc of the solids reported is 450470 Oe [31]. It was very difficult to increase the Hc higher than 500 Oe. However magnetic particleswith higher Hc were required for high-density recording media. It was well known that y-femc oxide particles containing cobalt had larger Hc. However, the temperature dependence of the Hc of the solids was large because the greater part of the Hc of such y-ferric oxide particles was a result of magnetocrystalline anisotropy rather than shape anisotropy. Great improvements in magnetic properties inHc were achieved by forming a thin layer of cobalt oxide compoundsonaciculary-ferric oxide particles[32]. The yFe203 particles were modified by immersing the particles in a solution containing Co2+ and Fe2+ followed by the addition of alkaline solution. Nowadays, most home videota’pes andaudiotapes of high bias position use this kind of particle. Figure 2b shows y-ferric oxide particles modified by cobalt.

C.

Iron Particles

Since iron has a large saturation magnetization of 216 emdg, it must be one of the best materials for magnetic recording media.Althoughironparticles were the fxst particulate materials employed for tapes, they were not usedafter the introduction of maghemite particles. This was partly because of the difficulty in overcoming corrosion with small iron particles. It did not take long until iron powders again attracted the interests of tape manufacturers as a material for a higher density recording system. Although iron particles can be prepared by variety of methods, most of the particles on the market areproduced from geothite through dehydration and reduction processes [23, 331. merefore, most processes are the same as the production processes for y-ferric oxide particles. That is, dehydrated geothite particles are reduced into ironratherthan magnetite, andthe surface of the

particles is oxidized carefully to prevent further oxidation by forming a thin oxide layer [34]. The saturation magnetization of the powders are 7040% pure iron, but an Hc higher than lo00 Oe can easily be attained. Special procedures, such as the addition of silicone oil or Ag or CO compounds during the reduction process, are effective in preventing the mutual sintering of particles, leading to the powder’s improved magnetic properties [35]. Figure 3 shows iron particles prepared from a-FeOOH. Elongated iron particles are superior particulate materials for magnetic recording media, but the problem associated with corrosion has not been completely overcome for long storage times.

D. Chromium DioxideParticles Chromium dioxide particles were developed as a high-quality magnetic recording material [36] and were used until cobalt-modified y-ferric oxide particles appeared on the market. The commercially available CrO2 is produced by the oxidation of Cr203 at high temperature and high pressure. The particles produced are single crystals and acicular in shape, with large aspect ratios and good magnetic properties. However,chromium dioxide particles are not widely

c

I & A T ” ’ ; n

-.

x

”

7

l

.

U

p

111

Figure 3 TEM of iron particles for 8 mm videotapes. Hc = 1540 Oe.(Courtesy of Horiishi, Toda Kogyo Corp., Hiroshima, Japan.)

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OZAKl

used today because of their high production cost in addition to their toxic properties. Although these particles are not much used at present, the technical term “CrO2 position” is still used for the bias position together with “high or metal position” for audio recorders and tapes.

E.

Iron Nitride Particles

By reducing dehydrated goethite particles in the presence of ammonia, iron nitride particles can be obtained [37]. They consist of variety of compounds having different FeM ratios in addition to FwN. These particle have a high saturation magnetization of about 190 emdg and a large coercive force, but there are some’pibblems with stability. Thus, they are not used at present for magnetic media: In ‘1972, Kim and Takahashi reported that Fe16Nz film was obtained by evaporating iron under a nitrogen atmosphere [38]. This compound showed a very large saturation magnetization of 2.8 clg per iron atom, which was larger than that of iron, where p~ is the Bohr magneton. However, there was some ambiguity about the existence of the compound until recently, when asingle-crystalFel6N2was reported [39]. Fe16N2 may be a desirable compound for recording materials if the compound is produced in the form of elongated fine particles. However, it seems very difficult to obtain the compound at present using chemical procedures.

F.

BariumFerriteParticles

In an ordinary recording system, the magnetic medium is magnetized along the direction of the motion of the magnetic medium, called the longitudinal recording method. In contrast, to this method, Iwasaki proposed a different type of recording method, called a perpendicular recording method, for a high-density recording system, in which the magnetic medium is magnetized perpendicular to the direction of motion of the magnetic medium [40]. Therefore, the easy axis of the magnetic particles for this medium must be oriented perpendicular to the supporting film. Platelet barium ferrite particles seemed a suitable material for this purpose since it has an easy axis perpendicular to the flat surface. Barium ferrite has been widely used for a precursor material for permanent magnets.However, the Hc of the particles was too large for the magnetic medium. Kubo and Hidehira developed new platelike barium ferrite particles having a proper coercive force by a glass crystallization method in solid-phase reactions [41]. Barium ferrite particles can also be produced by precipitation in aqueous solutions [42]. A transmission electron micrograph of platelike barium ferrite particles employed for perpendicular magnetic recording media is shown in Fig. 4. The perpendicular recording system is superior for high-density magnetic recording, and it is expected to be used in the future for computer memory and for high-definition video recorders.

SYNTHESIS OF MAGNETIC PARTICLES

431

4

R 6

m.. Figure 4 TEM of platelikebariumferriteparticlesforperpendicularrecording media. Hc = 750 Oe. (Courtesy of Horiishi, Toda Kogyo Corp., Hiroshima, Japan.)

Elongated magnetic particles are employed for the ordinary longitudinal recording system, as stated earlier. It was suggested by Lemke that the vertical components of the recorded signals play important roles as the recording density increases or as the recording wavelength becomes shorter [43]. Based on the same principle, videotapes with improved recordingproperties were made by mixing small amounts of barium ferrite particles and cobalt-modified particles [24].

V.COERCIVEFORCE AND DlSPERSlBlLlTY OF MAGNETIC PARTICLES FOR RECORDING MEDIA Coercive force is an extrinsic property that is influenced by many factors, such as size, shape, and packing density. Maghemite particles prepared from directly precipitated spindle-type hematite particles show a tendency to increase in coercive force with a decrease in particle size, and to increase in aspect ratios [29]. However, a quantitative relationship between the coercive force and size and/or aspect ratio is still not available. It is most desirable that measurements of coercive force be carried out for individual particles since particulate material contains particles of different

432

OZAKI

sizes, although they are highly monodisperse in size and shape.A novel method was developed by Knowels for the measurement of the rotation of magnetization of individual particles: he applied a pulse magnetic field to a particle dispersed in a viscous liquid and observed under a microscope the rotation of the solid in a critical magnetic field [M]. From the remanent coercive force o b tained, he explained the rotation of the magnetization in the particle using a fanning reversal mechanism. On the other hand, Aharoni showed that the rotation mechanism could be explained by a curling model using reported experimental data for individual particles [45]. Recently, measurements of the coercive force of individual particles were also made using Lorenz microscopy with maghemite powders prepared from directly precipitated spindle-type hematite particles [46]. The isolated particles showed a very large Hc of about 1200 Oe; however, this was somewhat below thevalue predicted by the cOherent rotation model, which may be partly caused by imperfections in the particles studied. Magnetic particlesareembeddedinplastic binder for useas magnetic recording media. Therefore, particles must have good dispersibility in organic solvents to have high orientability and smoothly coated surfaces. However, it is especially difficult to disperse magnetic particles, compared with nonmagneticparticles, because of strong magneticinteractions. The surface of the magnetic particle is usually coated with a polymertoprevent aggregation. Inoue et al. investigated the effect of an epoxy resin-adsorbed layer on the stability of y-Fe203particles dispersed in organic solvents and found thatalthough both the height of the maximum and the depth of the secondary minimum in the total potential energy of the colloidal interaction strongly affected the dispersion stability, the former was more effective than the latter [47]. The experimentally obtained surface roughness of the tape film produced by using polymer resin-adsorbed magnetic particles increased with the depth of the calculated total potential energy minimum rather than with the decrease in the height of the maximum energy. Homola and Lorenz showed that the recording performance of rigid disks was considerably improved by using magnetic particles coated with small colloidal silica particles [48]. The dispersion properties of the coated particles were enhanced by control of the separation distance between the magnetic particles, resulting in good particle orientability, which led to improved recording performance.

VI.

SMALL MAGNETIC PARTICLES AND MAGNETIC FLUIDS

Historically, dispersions of small magnetic particles with a superparamagnetic nature were used for observing magnetic,domains. When a concentrated suspension of magnetic particles is placedin the gradient of a magnetic field,

SYNTHESIS OF MAGNETIC PARTICLES

433

forces act on the particles and the magnetic interaction between these is enhanced. The magnetic field attracts particles, and the dispersing liquid moves together with the particles. Therefore, the liquid behaves as if it is magnetized. Such behavior is typical of a stable dispersion containing magnetic particles. In the 1960s, after finding a use for concentrated dispersions of magnetic particles in space technology, this magnetic dispersion, called a magnetic fluid or ferro fluid, was introduced as a new material. Since then, it has been finding a variety of applications [6]. At present, magnetic fluids are employed in many industrial technologies. For example, lubricant oil in which magnetic particles are dispersed works as a good sealing material when it is suspended by a magnetic field. The magnetic particles remain in the magnetic field together with the lubricant oil, which works as a sealing material. Nowadays, magnetic fluid is employed for sealing computer disk units. The constituent of the magnetic fluid is the concentrated dispersion of superparamagneticparticles protected by surfactants in an appropriate liquid. Since the magnetic fluid appeared, techniques for the production of a variety of magnetic fluids, including water-based and oil-based magnetic fluids, weredeveloped. The first magnetic fluid was prepared by milling magnetic material in a nonpolar organic liquid in the presence of oleic acid. Concentrated dispersions of magnetic particles were also independently prepared by Shimoiizaka et al. usingprecipitated magnetite particles [49]. According to their method, the magnetite particles were obtained by adding an alkaline solution to an aqueous mixture containing Fe2+ and Fe3+ ions, followed by the adsorption of oleic acid. The magnetic particles covered by oleic acid were filtered and dispersed in an organic liquid to form a stabledispersion of magnetic particles. Thomas succeeded informing 2-3Onm particles of cobalt with a narrow size distribution by decomposing cobalt carbonyl compounds inthe presence of suitable surface-active reagents [50]. By the same procedure, iron and nickel particles were prepared. Papirer et al. studied the decomposition of a toluene solution of c02(co)8in the presence of a surface-active reagent and found that at least two factors are responsible for the formation of particles with an extremely narrow size distribution: the division of the system into microreactors and a diffusion-controlled growth mechanism of the individual particles [51]. Konno et al. prepared magnetite particles using a microemulsion method [52]. The procedure was the mixing of a watedisooctane or watedcyclohexane microemulsion with aqueous FeCl3 and aqueous NH3, followed by the addition of aqueous FeC12 with vigorous stirring. In this procedure, Aerosol OT was necessary as a surfactant to solubilize adequate amounts of FeC12in the hydrocarbon used. In the microemulsion method, the size of the magnetite particle couldbe controlled by adjusting the size of waterdropletsinthe microemulsion and the content of ferric and ferrous ions.

OZAKI

434

Nanometer-sized iron oxide particles were prepared using unilamellar vesicles [53]. Adding an alkaline solution to vesicles containing intravesicular solutions of Fez+, Fe3+, and Fe2+/Fe3+ resulted in the formation of membranebound discrete particles of goethite and magnetite, and ferrihydrite. The particles had very small dimensions, in the range 1.5-12 nm. These results, together with particle formation in a microemulsion, are not only of interest in colloid chemistry but also have significance in mineralization in biosystems such as magnetotactic bacteria, in which particles are formed within enclosed organic compartments. Small metal particles can also be obtained by vacuum evaporation in lowpressure inert gas [54]. Magnetic particles of metals,such as iron, cobalt, nickel, and alloys of these metals, can be prepared by this method. Although the amount of particles obtainable by this method is limited, particles are clean compared with particles precipitated from solution. They are mainly used for studies of the physical properties of fine particles.

VII.

METAL AND METALOXIDEPARTICLES

In the poly01 process developed by Figlarz et al., a powdered inorganic compound, such as Co(OH)2 or Ni(OH)2, is suspended in a liquid polyol, such as ethyleneglycol [7]. The suspension is then heated to the boiling point or near the boiling point of the poly01 with continuous stimng. A complete reduction of these compounds can be achieved within a few hours. In this reaction the poly01 acts as a solvent for the starting compound; subsequently, for example, ethyleneglycol reduces the cobalt(II) or nickel(II) species in the liquid phase to metallic states, in which nucleation and growth of the particles occur. Metal particles obtained by the poly01 process are corilposed of equiaxial particles with a mean size of 1-10 pand narrow size distributions. Coprecipitation has been used for the production of ferrite particles. For example, according to Tang et al., MnFe04 particles of relatively small size (5-25 nm) can be obtained through the reactions MnCl,

+ 2FeC1, +xNaOH +MnFeO, + (x - 8)NaOH +4H,O + 8NaCl

for x > 8 ([Me]/[OH-I) and MnxFe3-x04 (0.2 < x < 0.7) particles of greater size, up to 180 nm, are produced from ferrous salts, where [Me] is the concentration of metal ions [55]. They found that in either case, the particle size appeared to be a unique function of the ratio of metal ion concentration to hydroxide concentration. A variety of ferrites can be obtained by similar reactions. Many metal oxides and hydroxide particles with narrow size distributions were prepared byMatijeviC et al. by forced hydrolysis at elevated temperatures. Hematite particles with a narrow size distributions can be prepared as spheri-

SYNTHESIS OF MAGNETIC PARTICLES

435

cal, cubic, spindle, or platelet types. Spherical hematite particles were prepared from ferric chloride solution [56], and spindle-type hematite particles of an extremely narrow size distribution were obtained by heating a ferric chloride solution containing small amounts of phosphate ions [57]. A transmission electron micrograph of the spindle-type particles is shown in Fig. 5. Some magnetic particles are produced by phase transformation in aqueous solutions. In this procedure, particles are recrystallized from other kinds of precipitates into the final forms. Spherical magnetite particles were formed by phase transformation of precipitated Fe(OH)2 under mild oxidizing conditions in the presence of m03 [58]. Ferrite particles were obtained by the same procedures. Platelike barium ferrite particles can also be prepared by this method. Large magnetite particles were also obtained in the presence of chelating agents and reducing agents [59]. Large cubiclike hematite particles were produced through the conversion of previously deposited p-FeOOH [60] in a acid solution of HC1 at 100°C. In Fig. 6, photomicrographs of loose and ordered agglomerates of the coarse cubiclike hematite particles observed in a stable dis-

Figure 5 TEM ofspindle-typehamatiteparticlesobtainedbyforcedhydrolysis of ferric chloride solution at 100°C in the presence of small amounts of phosphate ions.

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OZAKI

a

b

20um Figure 6 Photomicrographs of (a) loose andordered agglomerates of coarse cubiclike hematite particles and (b) spindle-type hematite particles.

SYNTHESIS OF MAGNETIC PARTICLES

437

persion of the solids by a metallurgical microscope are shown together with the agglomerates observed with coarse spindle-type hematite particles [61]. Hematite is known to show parasitic ferromagnetism, which has weak magnetic properties [62]. Thus, the formation of the loose and ordered agglomerationwas explained by the magnetic interactionsbetween hematite particles caused by the weak magnetism [63]. Elongated hematite particles were produced under basic conditions by aging freshly precipitated femc hydroxides at temperatures of 100-200°C in the presence of small amounts of organic compounds, such as sulfonic acids or hydrocarboxyl acids or salts of these compounds, as crystal growth control reagents. The elongated hematite particles thus produced can be used for the production of maghemite particles[27]. Platelike hematite particles were produced from a-FeOOH or Fe(OH)3 under strong basic conditions at elevated temperatures [64,65]. There was a critical temperature for each alkaline concentration in the hydrothermal transformation [M].In Fig. 7, a scanning electron micrograph of platelike hematite particles obtained by this method is shown. The complex reactions in metal solutions are affected by minor changes in experimental parameters, such as temperature, the species and concentration of anions, and the pH of the solution. In view of the complexity of the processes

Figure 7 Scanningelectronmicrograph (SEM) of platelikehematiteparticles duced by transformation from precipitated Fe(OH)3 at 200°C.

pro-

OZAKI

438

involved in the formation and transformation of iron oxides in aqueous media, special attention to well-defined magnetic iron oxides are described in the paper by Blesa and Matijevi'c[66].

VIII.

COMPOSITE MAGNETIC PARTICLES

Hirano et al. reported that spherical carbon particles containing highly dispersed cementite (Fe3C) particles were formed by heating copolymers of divinylbenze and vinylferrocene in a high-pressure bomb at 125 MPa and 650°C [67]. The cementite particles thus formed could be transformed into a-Fe by further heating at 850°C for 6 h. When the same copolymers were heated with water, magnetite particles were formed. A scanning electron micrograph of the spherical carbon particles containing dispersed magnetite particles is shown in Fig. 8. Using this technique, spherical carbon particles containing metal particles, such as cobalt or cobalt alloys, can also be prepared. Attempts have been made to improve its essential properties or to introduce new magnetic functions to particles by coating with nonmagnetic or magnetic

Figure 8 SEM ofsphericalcarbonparticlescontainingdispersedFe304particles prepared by pyrolysis of divinylbenzene and vinylferrocene copolymers. (Courtesy Hirano, Nagoya University, Japan.)

of

SYNTHESIS OF MAGNETIC PARTICLES

439

material. Iron particles covered with polystyrene were prepared for making the core of a high-frequency transformer [8]. Ishikawa et al. [9] coated large polymer particles with a thin magnetic film using a ferrite plating method developed by Abe and Tamura [68]. The particles thus obtained were found to be useful for toners and carriers for copying tools. Ajay and Matijevii: developed procedures for the preparation of well-defined, uniformly coated inorganic particles by forced hydrolysis of metal salt solutions involving coated spindle-type hematite particles with chromium hydrous oxide [69]. Recently, a so-called dry mixing or dry blending method was used for the modification of particles in powder technology [70]. Inthis method, the surface modification of coarse particles is carried out by mixing fine particles and coarse particles with an automatic ceramic mortar or with a centrifugal rotating mixer. This procedure can be applicable to the production of a variety of composite magnetic particles. Inada et al. succeeded in combining enzymes on synthesized magnetic particles, which enabled the easy reuse of enzymes [lo]. Furthermore, magnetic polymeric microsphere particles are employed for the labeling and separation of biocells [ 1 l].

IX.

MAGNETIC PARTICLES OF MAGNETOTACTIC BACTERIA

Magnetotactic bacteria were discovered by Blakemore in 1975. The magnetic particles obtained from the magnetotactic bacteria were confirmed as well-crystallized magnetite particles 50-100 nm [71]. The particles found in some bacteria were cubiclike, with a narrow size distribution, and were aligned in a single or multiple chains more or less parallel to the axis of the cell. The magnetic moment of each particle is small, but the total magnetic moments of the aligned particles are large enough to orient along the geomagnetic field. It is believed that magnetotactic bacteria navigate using these magnets as a direction detector toward north in the northern hemisphere and toward south in the southern hemisphere. The magnetite particles isolated from the bacteria showed magnetic properties similar to those of synthesized particles, and the aligned magnetic particles in their bodies were found to be a good model for the chain of spheres theory for the rotation of the magnetic moment of a magnetic particle [72]. It was also found that the surfaces of the magnetic particles isolated form magnetotactic bacteria were covered witha strong organic membrane. Matsunaga and Kamiya succeeded in immobilizing glucose oxidase and uricase on the organic membranes attached to the magnetic particles [73]. The enzymesbound on the particles were more active than enzymes combinedon synthesized magnetic particles. Such particles could be easily separated from the reactant solution. Introduction of the magnetic particles into bioparticles, such as blood cells and microphages, was also successful. Such small biopar-

440

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ticles carrying magnetic particles could easily be moved to desired places by applying a magnetic field.

REFERENCES 1. Bate, G., Magnetic recording materials since 1975, J. Magn. Magn. Mater., 100, 413 (1991). 2. Hibst, H., Magnetic pigment for recording information, J. Mugn. Magn. Mater., 74, 193 (1988). 3. Bate, G., Present and future of magnetic recording media, in Ferrires: proceeding of the ICF3, '80, (H. Watanabe, S. Iida and M. Sugimoto, eds.), Center for Academic Publications Japan, Tokyo, Japan, (1981), pp. 509-515. 4. Imaoka,Y.,Takada,K.,Hamabata,T.,andMaruta, F., Advances in magnetic recordingmediafrommaghemiteandchromiumdioxidetocobaltadsorbed gammaferricoxide,in Ferrites: Proceeding of theICF3, '80, (H.Watanabe, S. Iida and M. Sugimoto, eds.), Center for Academic Publications Japan, Tokyo, Japan,1981,pp.516-520. 5. Taketomi, S., and Chikazumi, S., Zisei Ryutai (Magnetic Fluidr), Nikkan Kogyo Shinbun, Tokyo, (1988), p. 3. 6. Rosensweig, R. E., Magnetic fluids, Sci. Am., 247, 136 (1982). 7. Fievet, F., Lagier, J. P., and Figlarz, M. Preparing of metal powders in micrometer and submicrometer sizes by the poly01 process, MRS Bull., 5, 29 (1989). 8. Ochiai, K., Hone, H., Kamohara, H., and Morita, M., An encapsulation process Nipfor magnetic metal powder and its application to powder core manufacturing, pon Kagaku Kaishi, 233 (1987). 9. Ishikawa, K., Ohishi, M., Saitho, T., Abe, M., and Tamura, Y., Newmagnetic capsule toner and carrier with high imaging-quality and low fusing temperature prepared by ferrite plating, in Abstracts of 6th International Conference on Ferrites, '87, Tokyo, EB-04 (1987). 10. Inada, Y., Takahashi, K., Yoshimoto, T., Kodera, Y., Matsushima, A., and Saito, Y.,ApplicationofPEG-enzymeandmagnetite-PEG-enzymeconjugatesfor biotechnical process, Trends Biotechnol., 6, 131 (1988). 11. Molday,R. S., Yen, S. P. S., andRembaum, A., Applicationofmagneticmicrospheres in labelling and separationof cells, Nature, 268, 437 (1977). 12. LaMer, V. K., and Dinegar, R. H., Theory, production and mechanismof the formation of monodispersed hydrosols, J. Am. Chem Soc., 72,4847 (1950). 13. Sugimoto, T., Preparation of monodispersed colloidal particles, Adv. Colloid Interfac. Sci., 28, 65 (1987). 14. MatijeviC, E., Monodispersed colloids: Art and science, Lungmuir, 2, 12 (1986). 15. MatijeviC, E., The art and science of colloids, Roy. Inst. Proc., 62, 161 (1990). 16. Blakemore, R. P., Magnetotactic bacteria, Science, 190, 377 (1975). for magnetic field 17. Kirshvink, J. L., and Gould, J. L., Biogenic magnetite as a basis detection in animals, Biosystems, 13, 181 (1981). 18. Chikazumi, S., Physics of Magnetism, Wiley, New York, (1965), p. 5.

441

SYNTHESIS OF MAGNETIC PARTICLES

19. Scholten, P. C., and Tjaden,D. L. A., Mutual attraction of superparamagnetic particles, J. Colloid Intelfac. Sci., 73, 254 (1980). 20. Scholten, P. C., How magnetic can a magnetic fluid be? J. Magn. Magn. Mater., 39, 99 (1983). Magnetic Oxides, Part II J. Craik, 21. Bate, G., Oxides for magnetic recording, in d.), Wiley, New York, (1975), p. 689. femc oxide, Crys22. Monish, A. H., Morphology and physical properties of gamma tals: GrowthandApplications, Vol. 2 (H.C.Freyhard,ed.),Springer,Berlin, (1979), p. 172. 23. Chubachi, R., and Tamagawa, N., Characteristics and Application of metal tape, IEEE Trans. Magn., Mag., 20, 45 (1984). 24. Fugiwara, T., Isshiki, M., Koike, Y., and Oguchi, T., Recording performances of Ba-ferritecoatedperpendicularmagnetictapes, IEEE Trans.Magn.,Mag., 18, 1200 (1982). 25. Yashiro, T., Kikuchi, Y., Matsubayashi, Y., and Morizumi, H., The effects of barium femte particlesaddedtoVHStapes, IEEE Trans.Magn.,Mag., 23, 100 (1987). 26. Kiyama, K., and Takada, T., Iron compounds formed by the aerial oxidation of ferrous salt solutions, Bull. Chem Soc. Japan, 45, 1923 (1972). 27. Matsumoto, M., Koga, T., Fukai, K., and Nakatani, S., Production of acicular ferric oxide, U.S. Patent 4, 202, 871 (1980). 28. Corradi, A. R., Andress, S. J., French, J. E., Bottoni, G., Candoflo, D., Cecchetti,

(D.

A., and Masoli, F., Magnetic properties of new(NP)hydrothermal particles,IEEE Trans. Magn., Mag., 20, 33 (1984). E., Preparation and magnetic properties of monodispersed spindle-type y-Fe203 particles, J. Colloid Intelfac. Sci., 107, 199 (1985).

29. Ozaki,M.,andMatijevie,

30. Amdt, V., y-Fe203 of improved quality from direct synthesis of acicular a-Fe203 particles, IEEE Trans. Magn. Mag., 24, 1796 (1988). 31. Yada, Y., Miyamoto, S., and Kawagoe, H., A new high Hc gamma ferric oxide exhibiting coercive force as high as 450-470 Oe, IEEE Trans. Magn., Mag., 9, 185 (1973). 32. Umeki, S., Saitoh, S., and Imaoka, Y.,A new high coercive magnetic particle for recording tape, IEEE Trans. Magn., Mag., IO, 655 (1974). 33. Asada, S., Preparation of fine acicular iron particles with high coercivity by reduction method, Nippon Kagaku Kaishi, 19, 22 (1985). 34. Asada, S., Surface-stabilation of fine acicular iron particles for magnetic recording media, Nippon Kagaku Kaishi, 1372 (1984). 35. Van der Giessen, A. A., and Klomp, C. J., The preparation of iron powders con-

sisting of submicroscopic elongated particles by pseudomorphic reduction of iron oxides, IEEE Trans. Magn., Mag., 15, 317 (1969). 36. Chen, H. Y., Hiller, D. M., Hudson, J. E., and Westenbroek, C. J. A., Advances in properties and manufacturing of chromium dioxide,IEEE Trans. Magn. Mag.,

20, 24 (1984). 37. Tanaka, T., Tagawa,K., and Tazaki, A., Synthesis and magnetic properties of iron nitride particles, Nippon Kagaku Kaishi, 930 (1984).

OZ4KI 38. Kim, T. K., and Takahashi, M., New magnetic material having ultrahigh magnetic moment, J. Appl. Phys. LRft., 20, 492 (1972). 39. Komuro,M.,Kozono, Y.,Hanazono, M., and Sugita, Y.,Epitaxial growth and

magnetic properties of FenN2 films with high saturation magnetic flux density,

J. Appl. Phys., 67 (part IIA), 5126 (1990).

40. Iwasaki, S., Perpendicular magnetic recording,IEEE Trans. Mugn., Mug., 16, 71 (1980). 41. Kubo O., Ido, T., and Hidehira, Y., Barium ferrite super-fine particles produced by glass crystallization method, Toshibu Rev., 43, 897 (1988). 42. MatijeviC, E., Uniform colloidal barium ferrite particles,J. Colloid Interfac. Sci., 117, 593 (1987). 43. Lemke, J. U., Ultra high density recording with new heads and tapes, IEEE Trans. Mugn. Mug., 15, 1561(1979).

44. Knowels, J. E., Magnetic properties of individual acicular particles, IEEE Trans.

Mugn. Mug., 17, 3008 (1981). 45. Aharoni, A., Angular dependence of nucleation field in magnetic recording media, IEEE Trans. Mugn. Mug., 22, 149 (1986). 46. Salling, C., Schultz, S., McFadyen, I., and Ozaki, M., Measuring the coercivity of

individualsub-micronferromagneticparticlesbyLorentzmicroscopy, Trans. Mugn., 27, 1 (1991).

IEEE

47. Inoue,

H.,Fukke,H.,andKatsumoto,H.,Effectofpolymeradsorbedlayeron magneticparticledispersion, IEEE Trans. Mugn., 26, 75(1990). , 48. Homola, A. M., and Lorenz, M. R., Novel magnetic dispersions using silica stabilized particles, IEEE Trans. Mugn., Mug., 22, 716 (1986). 49. Satoh, T., Higuchi,S., and Shimoiizaka,J., Dispersibility and van der Waals force attractive energy of magnetite colloid in cyclohexane, Abstracts of the 19th Annual meeting of The Chemical Society of Jupun I, p. 293 (1966). 50. Thomas, J. R., Preparation and magnetic properties of colloidal cobalt particles, J. Appl. Phys., 37, 2914 (1966). 51. Papirer, E., Homy, P., Balard, H., Anthore, H., Petipas, C., and Martinet, A., The

preparation of ferrofluid by decomposition of dicobalt octacarbonyl, J. Colloid In-

terfac. Sci., 94, 220 (1983). 52. Gobe, M., Konno, K., Kandori, K., and Kitahara,

A., Preparation and characterization of monodisperse magnetite sols in W/O microemulsion,J. Colloid Inter-

face. Sci., 93, 293 (1983). 53. Mann, S., and Hannington, J. P., Formation of iron oxide in unilamellar vesicles, J. Colloid Interfac. Sci., 122, 326 (1988). 54. Granqvist, C. G., and Buhrman, R. A., Ultra fine metal particles, J. Appl. Phys., 47, 2200 (1976). 55. Tang, Z. X.,Sorensen, C. M., Klabunde, K. J., and Hadjipanayis, G. C., Prepara-

tion of manganese ferrite fine particles from aqueous solution, J. Colloid Inter-

fuc. Sci., 145, 38 (1991).

56. MatijeviC, E.,and Scheiner, P., Ferric hydrous oxidesols,J. Colloid Intetjiuc. Sci., 63, 509 (1978). 57. Ozaki, M.,Krathovil, S., andMatijeviC,E.,Preparationformationofmonodispersed spindle-type hematite particles, J. Colloid Inter&. Sci., 102, 145 (1984).

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58. Sugimoto, T., and MatijeviC, E., Formation of uniform spherical magnetite particles by crystallization from ferrous hydroxide gels, J . Colloid Interjac. Sci., 74, 227 (1980). 59. Sapieszko, R. S., and MatijeviC, E., Preparation of well defined colloidal particles by thermaldecomposition of metalchelates, J . Colloid Interjac. Sci., 74,405 (1980). 60. Hamada, S., and MatijeviC, E., Ferric hydrous oxide sols. IV. Preparation of uni-

form cubic hematite particlesby hydrolysis of ferric chloridein alcohol water solutions, J . Colloid Interjac. Sci., 84, 274 (1981). 61. Ozaki, M., Suzuki, H., Takahashi, K., and MatijeviC, E., Reversible ordered agglomeration of hematite particles due to weak magnetic interactions, J . Colloid Interjac. Sci., 113, 76 (1988). 62. Dunlop,D. J., Magneticproperties of fine-hematite, Ann. Geophys. 27,269 (1971). 63. Ozaki, M., Egami,T.,Sugiyama,N.,andMatijeviC,

E., Agglomerationincolloidal hematite particles due to weak magnetic interactions, J . Colloid Znterjac. Sci., 126, 212 (1988). 64. Nobuoka, S., and Ado, K., Studies on thin iron oxide platelets I. Formation of airon oxide by hydrothermal reactions, Shikizai, 60, 265 (1987). 65. Ozaki, M., Ookoshi, N., and MatijeviC, E., Preparation and magnetic propertiesof uniform hematite platelets, J . Colloid Interjac. Sci., 137, 546 (1990). 66. Blesa, M. A., and MatijeviC, E., Phase transformation of iron oxides, oxohydroxides, and hydrous oxides in aqueous media, Adv. Colloid Interfac. Sci., 29, 173 (1989). 67. Hirano, S., Yogo, T., Suzuki, H., and Naka, S., Synthesis of iron dispersed car-

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coatedmicrospongeprepared bydryimpactblendingmethod, Chem. Pharm. Bull., 37, 3367 (1989). 71. Masuda, T., Endo, J., Osakabe, N., Tonomura, A., and Ani, T., Morphology and structure of biogenic magnetite particles, Nature, 302, 411 (1983). M., Frankel, R. B., Flanders,P. J., Blakemore, R. P., and 72. Moskowitz, B. Schwartz, B. B., Magnetic properties of magnetotactic bacteria, J . Magn. Magn. Mater., 73, 273 (1988). 73. Matsunaga, J., and Kamiya, S., Use of magnetic particles isolated from magnetotactic bacteria for enzyme immobilization, Appl. Microbiol. Biotechnol., 26, 328 (1987).

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l9 Chemistry and Processing of HighTemperature Superconductors Shin-Pei Matsuda Hitachi, Ltd. Ibaraki, Japan

1.

HIGH-TEMPERATURESUPERCONDUCTORS

In 1986, Bednon and Muller of IBM Zurich discovered an oxide superconductor based on La-Ba-Cu-0 with a critical temperature Tc of 30K [l]. The crystal structure was determined by the Tanaka and Kitazawa group at Tokyo University to be La2-xBaxCu04 [2]. In the following years oxide superconductors with higher Tc have been found [3-71: La2-SrxCul04 YlBa2Cu307 Bi2SnCa2Cu30x T12Ba2Ca2Cu30x Tl1Sr2Ca2Cu3Ox

[La (214, TPO K1 [Y (123), Tc93 K] [Bi (2223), Tc105 K] [TI (2223), Tc122 K] [Tl (1223), Tc120 K]

All the compounds contain copper oxide, which plays a fundamental role in the origin of superconductivity. The high-temperature superconductors have a crystal structure of so-called perovskite, whose primary structure is ABO3, where A is a di- or trivalent metal atom with a larger ion size and B is a tetra- or trivalent metal atom with a smaller ion size. In the preceding high-temperature superconductors, the A site is occupied by La3+, Y3+, Ba2+, Sr*+, and Ca2+ andthe B-site is occupied by C$+, B?+, and TP+. If the copper valency is 2, the cation valency is insufficient to form the nominal ABO3. This insufficiency is compensated for by 445

446

MATSUDA

the unusual Cu3+ and/or by oxygen deficiency. The occurrence of Cu3+ is apparently related to the origin of superconductivity in the copper oxide perovskite. In the synthesis of copper-containing oxide superconductors, therefore, they are fired in a highly oxidizing atmosphere. The mechanism of the metal superconductors, such as Nb-Ti and NbsSn, is explained byBCS theory, which emphasizes phonon-electron interaction to form Cooper pairs. Taking into consideration the coupling strength and the mobility of electrons, a conceptual comparison between the metal and the oxide superconductors is schematically illustrated in Fig. 1 [8]. Inmetal superconductors electrons are free and mobile in the normal state. The superconducting state arises at low temperatures when two electrons are coupled to form Cooper pairsmediatedthrough the electron-phonon interaction. Since the electronphonon interaction is relatively weak, the critical temperature is limited to lower than 30 K. In the oxide superconductors, paired electrons exist in the bonding molecular orbital with their spins antiparallel. When the paired electrons are delocalized and become mobile, they form Cooper pairs. In this case the coupling force is relatively strong and thus a critical temperature as high as 120 K is obtained for the oxide superconductors. The bonding between Cu and 0 is ionic, with some covalent nature, in the copper oxide superconductors. The parent compounds are intrinsically insulating, and the superconducting state exists adjacent to the insulating phase (anInsulator

Conductor

Electron

% Clonic Crvsta1s)l

””“.

i r t-l [Metals] “““C.

Figure 1 Origin ofoxide and metal superconductors. Shadowed region on left corresponds to oxide superconductors and that on the right to metal superconductors.

HIGH-TEMPERATURE SUPERCONDUCTORS

447

tiferromagnetic phase). Since the breaking of the superconducting state is caused by the scattering of paired electrons by a phonon, there naturally exists the isotope effect to some extent in both the metal and oxide superconductors. The carrier concentration of the oxide superconductors is in the range of -1022 ml-1, which is one order of magnitude smaller than that of the metal superconductors. The critical current density Jc is therefore smaller in the oxide superconductors.

II. CRYSTALLOCHEMICALCLASSIFICATION The high-temperature superconductors based on copper oxides have a principally perovskite structure. The crystal structures of four typical perovskite superconductors, BaPbxBil-& (BPB), La (214), Y (123), and T1 (1223) are shown in Fig. 2. BPB is a simple perovskite, and the others are layered perovskites. The perovskite superconductors are classified in Table 1 from the viewpoint of crystal structure and constituent elements. Class I is the simple cubic perovskite with AB03 composition. Typical of class I, BaPbxBi1-& was first synthesized by Sleight et al. in 1975 [9]. The parent compound is BaBiO3, whose Bi3+ site is partially substituted by Pb4+. After the discovery of the hightemperature superconductors, it was found that the replacement of Ba by K also gives rise to a superconductor, B a ~ ~ K x B i 0 with 3 , Tc = 28 K [lo]. The superconductors in class JI have the crystal structure A g 0 4 , which is conventionally called a K2NiF4 compound. A typical class II superconductor is ( L a ~ ~ B a ~ ) 2 C uwith O q Tc = 30 K, discovered by Bednorz and Muller. As seen in Fig. 2, A 3 0 4 is composed of an AB03 perovskite unit and an A 0 unit

Ba Pb,-xBixO3

(Lal-xBax)rCu O4 Y( BazCu30rT~lSrzCazCu30, TI cu cu cu TI

30K

Tc=13K

-

1 Layered1.5

105K -Layered

3 -Layered

4 -Layered

Figure 2 Crystal structure of four types of layered perovskite superconductors.

pical

448

MATSUDA

Table 1 Classification of SuperconductivePerovsKtes Class Ba-Pb-Bi-0 I AB03 II A 9 0 4

Derivatives La-Ba-Cu-0 Nd-Ce-Cu-0

ILI A 1A'2B30x Y-Ba-Cu-0 Y-Sr-Cu-0 IV (AIA')-(BIB')-Ox Bi-Sr-Ca-Cu-0 T1-Sr-Ca-Cu-0 T1-Ba-Ca-Cu-0

Tc (K) 13 Ba(K, Rb)-Bi-0 30 Na)-Cu-0 Ca, La-(Sr. 24 93 (La,Nd, Sm,Eu,Gd, Dy,

80

Ho, Er, Tm, Yb)-Ba-Cu-0 (Y/Ca)-Sr-(Cu/Pb)-O

(2223), (2212), (2201) 105 105 (1212) (1223), (2223), 122 (2212), (2201), (1223). (1212)

stacked alternately. Thus, Ago4 is a layered perovskite. Crystallographically it is a three-layered perovskite, but the repetition unit is deemed to be a 1.5 perovskite unit. The parent compound of ( I A ~ - ~ B ~ ~ ) is ~ CLa2CuO4. U O ~ The hole carriers are introduced by partly replacing La3+ by BP+. The oxygen content should be written as 4 - y instead of 4, where x > y under normal oxidizing preparation conditions. The partial substitution of the La site by Sr, Ca, and Na also leads to the formation of a superconductive state, the highest Tc of 40 K being obtained in the (Lal-xSrx)2Cu04 system at x = 0.075 [3]. (Ndl-xCex)2Cu04, whichhas a slightly different crystal structurefrom K2NiF4, was found to be a superconductor with electron carriers when prepared under mild reducing conditions, for example, a nitrogen atmosphere [1l].In this case Nd3+ is partly replaced by C e b , introducing electron carriers. Superconductors inclass 111 have a nominal composition ofA1A'2B30x, which is characterized by the existence of two kinds of A-site ions (A and A') and one B-site ion. Typical of class 111 is Y1Ba2Cu30x with Tc = 90 K, first synthesized byWu et al. The A-site ions are Y3+ and Ba2+ [4].The crystal structure is a three-layered perovskite; that is, the repetition unit is three blocks of the primary perovskite. The Y site can be partly or completely replaced by 10 other rare earth elements, such as La, Nd, Sm, Eu, and Ho, giving rise to a 90 K class of superconductors in all cases. From a chemical point of view the constancy of the critical temperature is surprising, since the rare earthelements change their valencies (Sm3+, Sm4+, E d + , and Eu3+) and ion sizes. This suggests that Tc depends only slightly on the nature of the rare earth elements as long as the crystal structure is preserved. The only exception is PrBa2Cu30x, which cannot be made superconductive although PrBa2CusOx is prepared in the (123) structure because of its rather high electrical resistance in the normal state. It is supposed, however, that there must be a synthetic condition under whichPrBa2Cu30xbecomes superconductive. The oxygen content 7 in

HIGH-TEMPERATURE SUPERCONDUCTORS

449

YlBa2Cu307 deviates largely from the nominal 9 in (AB03)3, indicating that two of the nine of oxygens are deficient in the formal perovskite structure. It may not be appropriate to call these high-temperature superconductors perovskite compounds. The Ba site in Y1Ba2Cu307 can be replaced by Sr, leading to the formation of a superconductor with Tc = 80 K in a metastable state. A superconducting Y1SnCu30x can be synthesized only under high oxygen pressure. When Pb is incorporated into the Y-Sr-Cu-0 system, it forms a stable superconductor (Y/Ca)-Sr-(Cu/Pb)-0 with Tc = 70-80 K [12]. Under certain conditions Y1Ba2Cu40x withTc = 80 K, a derivative of Y (123), is formed. Superconductors in class N have the general formula (A/A')-(BIB')-0, which is characterized by having two kinds of A-site ions and two kinds of Bsite ions. Typical of class N are Bi2SnCa2Cu30x [5] and T12Ba2Ca2Cu30x [6],which are four-layered perovskites. The AIA' is Sr/Ca or BdCa and BIB' is Bi/Cu or TVCu. In the Bi-Sr-Ca-Cu-0 system there exist three variations, namely, the Bi(2223) phase with Tc = 105 K,Bi(2212) phase with Tc = 80 K, and the Bi(2201) phase with Tc = 20 K.These Bi superconductors contain double Bi-0 planes in their crystal structures. In the T1-Ba-Ca-Cu-0 system there arethree superconducting phaseswiththe double 1-0 plane, namely, T1 (2223) with Tc = 120 K, Tl (2212) with Tc = 110 K, and T1 (2201) with Tc = 80 K. The Tl-Sr-Ca-Cu-0 system forms two superconducting phases with the single Tl-0 plane, that is Tl (1223) with Tc = 105 K and T1 (1212) with Tc = 80 K, in a metastable state [7]. The metastable superconductors in the Tl-SrCa-Cu-0 system become stable when a partial substitution of T1 by Pb, Sr by Ba, and/or Y by Ca is performed [13]. The substitution means the adjustment of carrier concentrations, accompanied byan increase in Tc. An appropriate concentration of carrier is essential to attain the stable superconductive states. The Tl-Sr-Ca-Cu-0 system has a potential for use in magnets operable in liquid nitrogen for its high Tc and high critical magnetic field Hc [14].

111.

SYNTHESIS OF SUPERCONDUCTING POWDER AND BULK

A.

Y(123) Phase of Y-Ba-Cu-0 System

The most common method of synthesizing the high-temperature superconductors in bulk form is shown in Fig. 3. Appropriate amounts of Y203, BaC03, and CuO are mixed and ground thoroughly. The mixed powder is precalcined in an electric furnace at 900-980°C for about 10 h in an oxygen or air atmosphere. The calcined powder is ground and then pelletized intoa disk form. The pellets are sintered at 975°C for 10 h under oxidizing atmosphere andthen cooled slowly to room temperature. It is very important to cool slowly in an oxygen or airatmosphere, especially passing the 400-600"C region, for at least 2 h. The resultant bulk has a chemical formula ofYlBa2Cu306.7-7, which

450

MATSUDA

I

Mixing Grinding

I

Calcination

I

Grin,ding

950°C,10h

I

I Sintering I 975OC,lOh 1 I Slow Cooling I -10h Figure 3 Preparationprocedure for bulk Y-Ba-Cu-0 superconductors.Calcination must be done in an oxidizing atmosphere.

shows the superconducting transition at 90-93"C, as shown in Fig. 4. For the proofof superconductivity it is required to show both zero resistance and a Meissner effect (diamagnetism). It is seen in Fig. 4 that both the resistivity and susceptibilitydrop sharply near90°C. The critical temperature depends strongly on the oxygen stoichiometry, the value of 6.7-6.95 normally obtained being smaller than the ideal value 7.0. When the oxygen stoichiometry is less than 6.4, the sample shows no superconductivity [15]. The long annealing at 400-600"C in oxidizing atmosphere is necessary for the bulk materials to absorb oxygen. The absorption and desorption of oxygen are foundtobereversible.

B.

Bi (2223) and Bi (2212) Phases of Bi-Sr-Ca-Cu-0 System

Preparation of the Bi-Sr-Ca-Cu-0 superconductors is essentially the same as in the Y-Ba-Cu-0 system, except thecalcination temperature. The Bi (2223) phase is formed by firng at 850°C for 200 h in air atmosphere. To shorten the firing time, theconventional practice is to replace 10-20 atom% Bi by Pb.The replacement induces the formation of liquid-phase Ca2PbO4, which accelerates the solid-state reaction. Lead atoms are incorporated in the Bi sites. The Bi (2212) phase is more stable than the Bi (2223) phase and is formed by calcin-

HIGH-TEMPERATURE SUPERCONDUCTORS 1

E 10

0

-

I

q

T

E

v

iz2

7

5: 3

v)

W

U

0

, \,, 0

50

3W

Tc 90K

U,

Tc 90K

E

5 -

l-

451

v)

100

150 250 200

0

(b)

50

100

TEMPERATURE (K)

Figure 4 Superconductingpropertiesof Y-Ba-Cu-0 bulksample: change, (b) susceptibility change (diamagnetism, Meissner effect).

(a) resistivity

ing at 880°C for 2-10 h. At 880°C the Bi (2212) phase is partially melted, and therefore the reaction is considerably faster [16]. The Bi system contains a double Bi-0 plane in the crystal structure whose interplane strength is relatively weak and easily laminated. Crystallites in thick films ofBi superconductors prepared by the doctor blade or screen printing method are easily oriented by rolling or pressing because of the laminating property [17]. A high critical current density, for example, 105 Ncm2 at 4.2 K and 1-30 T of applied magnetic field, is obtained in wire and tape forms. Although the critical temperature of the Bi (2223) phase is much higher than the liquid nitrogen temperature 77 K, the Jc at 77 K is very low in the presence of an external magnetic field of 1 T because of the low critical field Hc. The application of the Bi system to magnets is limited to temperatures below 20 K.

C.

TI (2223) Phase of TI-Ba-Ca-Cu-0 System

The Tl(2223)phase of the Tl-Ba-Ca-Cu-0 system with Tc = 122 K is prepared by calcining the corresponding mixture at 845-860°C for 5-10 h in air. Special care must be paid, since the mixture loses a few percent T1 through vaporization during calcination. Recalcined pellets are calcined in a crucible together with Tl (2223) powder to compensate for the T1 losses. At temperatures higher than 870°C Tl (2223) tends to decompose to Tl (2212) and CaCuOz. The Tl(2223) superconductor is much more easily formed than Bi (2223). The interplane bonding between the T1-0 planes is considerably stronger than that in Bi (2223). The dependence of Jc on the magnetic field has been studied using a polycrystalline Tl(2223) thin film on a single crystal of MgO prepared by the laser ablation method [18]. Platelike crystallites are aligned with the c plane on the MgO surface. As shown in Fig. 5, the Jc is 2 x 105 Ncm2 at 77 K in the absence of an external magnetic field but decreases very rapidly in a magnetic field applied perpendicular to the c plane. The influence of the magnetic field parallel to the c plane is rather small. The two dimensionality or

452

MATSUDA

Figure 5 Criticalcurrentdensity of polycrystalline "l (2223) thin film inmagnetic fields. Crystallite c planes are parallel to the MgO surface. Fields are applied parallel and perpendicular to the c plane. (From Ref. 14.)

anisotropy of the superconducting properties is clearly seen. The facts suggest that the Tl (2223) has no effective pinning centers at 77 K when a magnetic field is applied perpendicular to the c plane. The Hc of Tl (2223) is low at 77 K, although slightly higher than that of Bi (2223).

2.

TI (1223) Phase of TI-Sr-Ca-Cu-0 System

The Tl (1223) phase of the pure Tl-Sr-Ca-Cu-0 system is a metastable superconductor. Among many samples prepared, only a few samples show a weak superconductivity or Meissner signal and many others show no superconductivity [7]. This is a reason the T1-Sr-Ca-Cu-0 system is called a metastable superconductor. When Tl is partially substituted by Pb, Sr by Ba, or Ca by Y, the system becomes a stable superconductor [13]. It is thought that the pure T1Sr-Ca-Cu-0 system is too doped to become superconductive.Reducing the carrier concentration by the substitution thereforeleads to a superconductive state. Examples of an optimum composition are as follows:

A high-quality superconductor with Tc = 120 K is formed when a mixture of composition 1 is calcined at 875°C or composition 2 at 900°C for 8-30 h. The substitution of Sr by Ba leads to the formation of a liquid phase comprised

453

HIGH-TEMPERATURE SUPERCONDUCTORS

of Tl-Pb-Ba-Cu-0 at 877 K, reducing the temperature of sintering [18]. At temperatures above 890"C, Tl (1223) decomposes according to Tl(1223)+ Tl(1212)+CaCu02 The bulk samples prepared thus show crystallite size of2-3 pm. When samples with larger grain sizes are required, the heat treatment must be done at 930-980°C with partial melting [19].

IV. PREPARATION OF LARGE-GRAINCRYSTALS A.

Single Crystals of La-Sr-Cu-0 System

Single crystals of high quality are required to study the fundamental physical and chemical properties of a superconductor. Single crystals of La2-xSrEu04 and YBa2Cu307 are prepared by a conventional traveling solvent floating zone (TSFZ)method [20]. A schematic diagram of a TSFZ apparatus is shown in Fig. 6. A focused infrared light is irradiated on a rod-shaped sample prepared by an ordinary sintering or a cold isostatic pressing method. The sample rod is

Vacuum

Computer

f?

Figure 6 Traveling solvent floating zone apparatus. (From Ref. 20.)

454

MATSUDA

moved slowly downward at a speed of0.2-5mm/h. The irradiated region is melted. To obtain better homogeneity in the liquid, the upper part of the rod is counterrotated against the lower part. A single crystal of La1.8sSro.1sCu04 of 8 mm diameter and 100 mm length has been successfully prepared. Magnetization hysteresis of La (214) is shown in Fig. 7 [21]. It is seen that AM has two peaks, the first at zero field and the second at higher fields. The phenomenon is called a peak effect, which is explained by the introduction of pinning centers as weak superconductive regions are converted to nonsuperconductive by the external magnetic fields. Crystal defects, such as oxygen deficiencies, cation concentration inhomogeneity, and inclusions, may play a role of pinning centers. The second peak shifts to higher fields as the temperature is lowered. The AM corresponds to the intragrain Jc as formulated in Bean’s model,

AM

J, =2Mm2 d

where d is the thickness of a sample perpendicular to the magnetic fields.

B.

Large-Grain Crystals of Y-Ba-Cu-0 System

Several methods of growing a large single-grain crystal of Y (123) have been developed based on partial melting. They are called melt textured growth [22], QMG(quenchandmelt growth) [23],and melt powdermeltgrowth[24], whose principal reaction is a peritectic reaction between a liquid-phase Ba-Cu0 and a solid-phase YzBaCuOs in a very slow cooling process. A pseudobinary phase diagram of Y2CuOs-Ba3Cuso8 is shown in Fig. 8 [25]. A liquid phase is formed above 980°Cin a mixture of Y203,BaO, and CuO,and YzBaCuOs decomposes at 1200°C.The preparation procedure of QMG process is showninFig.9. A typical composition of20%richin Y203 from the YBa2Cu307 stoichiometry is heated and melted at 1500°C in a Pt crucible and quenched to room temperature very rapidly using two copper plates Y2BaCu05, BaO, and CuO + melt + Y203 The quenched melt is quickly heated to 1200°C for partial melting for 30 minutes: Melt -+ Y2BaCu05+ liquid (Ba,Cu,08) The peritectic reaction takes place during the slow cooling from lo00 to 940°C for 10 h. Y2BaCu05 +liquid(Ba,Cu,O,)

+ YBa,Cu,O,

+excess Y203

In the process it is crucially important to disperse the Y2BaCuOs (or Y203) particlesfinelyand to use a YzBaCuOs-(orY203)rich composition, which

HIGH-TEMPERATURE SUPERCONDUCTORS

455

c

-10

-5

0

5

10

Bex [TI Figure 7 Magnetizationhysteresis of La-Sr-Cu-0single-crystal La1.82Sro.18CuO4: overdoped sample, magnetic fields parallel to c axis. (From Ref. 21.)

1400

-

L Y*O3 +L

1200 -.............

I

0 .

Y

2 a

-

211+L

c.

2

g, E

F

1OOo-

123i-L

-

2n+123

8005YsOa

+2Bao

211

3BaCuol 123

+2cuo

Figure 8 Pseudobinaryphasediagramalongtieline liquid; (211), Y2BaCuOs; and (123), YBa2Cu30x.

of Y2BaCuOs-YBa2Cu30x. L,

456

MATSUDA

v203

R.T.

+L

I Time (hrs)

Figure 9 Quench and melt process. Peritectic reaction takes place between IO00 and

940°C for 10 h.

eliminates liquid residues between grain boundaries. The QMG is later modified so that the initial melting is omitted when 2 wt%R02 is added to the mixture [26]. Staiting from Y2BaCuOs and Ba3CU508, a large single-grain YBa2Cu307 with the dimensions 4 cm diameter and 2 cm thickness is successfully prepared. To grow a single-crystal YBa2Cu307, it is practical to seeda small SmBa2Cu307 crystal that has a higher melting point than Y (123) in the partial melt. The intragrain Jc of a single-crystal YBa2Cu307 prepared by QMG, shown in Fig. 10, is in the range of2-5 x 104 A/cm* at 77 K and 1 T [27]. The intragrain Jc is calculated from AM using Bean's equation. A large single grain can also be prepared by the unidirectional solidification method. The superconducting pellet behaves as a permanent magnet, which could be applied to the magnetic bearing and flywheel. The intergrain or transport Jc in polycrystalline Y (123) has been found to be poor. The insulating phases between Y (123) grains are yet to be identified to improve the transport Jc.

C.

Single Crystals of TI-Sr-Ca-Cu-0 System

Single crystals of the Tl (1223) phase of the Tl-Sr-Ca-Cu-0 system are prepared by a conventional flux method. An example of the starting composition is

457

HIGH-TEMPERATURE SUPERCONDUCTORS

I 0

I

I

I

0.5

1

1.0

1

Magnetic Field (T) Figure 10 Criticalcurrentdensity of Y-Ba-Cu-0largesingle-grainsamplesprepared by melt powder melt growth method. Three samples contain different amountsof y203.

where the last compound acts as a flux. The precalcined powder is loosely packed in an alumina crucible 3-4 cm in height, heated to 940°C, and cooled slowly to 850°C for 100 h. A sintered bulk sample thus obtained is composed of T1 (1223), nonsuperconducting grains, such as Sr-Ca-Cu-0, and residues of liquid phases. Breaking the bulk by a hammer, single crystals with 0.5-1 mm size are found and collected under microscope. A photograph of a single crysSeven square sheets ofT1 (1223) tal of Tl (1223) is shown in Fig. 11 [28]. crystals are stacked with c planes parallel. The thickness of each sheet was 0.05 mm. The critical temperature was found to be 114 K. An anisotropy of magnetization was measured using the single crystal andis shown in Fig. 12. Even at 77 K and 1 T, there exists a large magnetization hyteresis, indicating a large intragrain Jc. The intragrain Jc in the c plane is calculated as >5 x 104 Akm2 at 77 K and 1 T, which is much greater than in T1 (2223) and Bi (2223) and in the same range as Y (123). The linkage between Tl (1223) grains has been found to be stronger than that of Y (123) [29]. Since the magnetic hysteresis extends to 8 T, Tl (1223) could find applications in magnets generating 1-5 T and operated in liquid nitrogen.

Figure 11 Scanningelectronmicroscopeimageof

T1 (1223) singlecrystal.

B B

20

300

200 100

0

5

l0

2a,

0

W

-1 00

x

-10

-200 "mn -"I"

-6

-4

-2

0

2

Magnetic field (T)

4

6

-20 -6

-4

-2

0

2

4

6

Magnetic field (T)

Figure 12 Magnetization hysteresis of TI (1223) single-crystal. Anisotropy is clearly seen for the applied magnetic field direction. (From Ref. 28.) 458

459

SUPERCONDUCTORS HIGH-TEMPERATURE

V.

FABRICATIONOF SUPERCONDUCTINGWIRES

Several methods, such as the drawing, screen printing, plasma spray, and dip coating, have been tried to fabricate a wire from the oxide superconductors. The drawing method is widely used to fabricate a long wire (-1 km) from NbTi and Nb3Sn superconductors. It is not practical to make a magnet unless a wire whose transport Jc exceeds 104 Ncm2 in the presence of 1-5 T magnetic fields is developed. To attain the high Jc the following four points, as illustrated in Fig. 13, are crucially important: 1. Densificationof superconducting core 2. Clean grain boundaries 3. Alignmentof crystallites 4. Superconducting core withhigh intragrain Jc

Among these it is most difficult to eliminate weak links at the grain boundaries. The drawing method combined with the rolling process, which is one of the mostpromisingmethods to fabricate tape-shaped wire, is showninFig. 14 [30]. For the sheath material, a Ag tube that allows oxygen diffusion into the core is generally used. Powders of the oxide superconductor are packed into a Ag tube with 4-6 mm diameter and 0.5-1 mm wall thickness. The diameter of

7

" "

I

I

: I

I I

Impurities,

a

/ Oxygen Deficient Phase I IClean Boundary1

i

J

Figure 13 Threeimportantstepstoincreasethecriticalcurrentdensity conducting wire from oxide superconductors.

of super-

TPSr-Ca-Cu-O Powder

list9 Ag Tube f

l

Q

W

Outer Diameter

Outer Diameter

6mm

2.8mm

Thickness 0.3mm

Packing

Drawing

Rolling

Figure 14 Fabricationmethodfortape-shapedwirewith

an oxidesuperconductor

core.

loo000

10000 n

"E 0

h

I

W

0 T

l000

5T :600A/cm2

CI L

0

% C 100

I

I

I

I

4

6

8

10

Magnetic Field (T) Figure 15 Transportcriticalcurrentdensity of T1 (1223). TI (2223), and Bi (2223) tape in magnetic fields applied perpendicular to the tape surface. Critical field H*:TI (1223) >> T1 (2223) > Bi (2223).

461

HIGH-TEMPERATURE SUPERCONDUCTORS

the tube is reduced to 1-3 mm in successive drawings and then rolled to a tape with 0.2-0.5 mm thickness. Taking into considerations the melting point of Ag (960"C), the tapes are heat treated at temperatures below 910°C for 8-50 h in air atmosphere. For the Bi (2212) and Bi (2223) phases, crystallites are automatically aligned in the rollingprocess because of their laminating property between c planes [31]. On the other hand, randomly oriented crystallites are observed in the Y (123), Tl (1223), and Tl (2223) tapes [32]. The transport Jc at 77 K of the Bi (2223), Tl(1223), and Tl (2223) tapes in magnetic fields is shown in Fig. 15 [33]. Magnetic fields are applied perpendicular to the tape surface. In a zero magnetic field the Jc is in a range of 2-5 x 104 Ncm2 at 77 K. It is seen that the Jc drops very sharply, anorder of magnitude, at fields less than 0.1 T in all three tapes. At fields more than 1 T the Jc of Tl (1223) stays relatively constant; those of Bi (2223) and Tl (2223) further decrease. The sharp drop at low fields is supposedly caused by the weak links between superconducting grains. Since T1 (1223) has a higher irre-

NbTi

lo8

I

0.01

/

I

0.1

I 1

MAGNETIC FIELD (T)

."

I

10

-

I

20

Figure 16 State-of-the-art presentation of critical current density of metal and oxide superconductingwirein 1992. MetalsuperconductorsareusedonlyinliquidHeand oxide superconductors in liquid NZ and He.

MATSUDA versibility magnetic field H*,the Jc value extends over 8 T. The transport Jc should approach the intragrain Jc if crystallites are aligned with the c plane parallel to the tape surface. The status of Jc in 1992 is shown in Fig. 16, which compares the oxide superconductors with the metal superconductors. The transport Jc of the oxide su. perconductors at 77 K is more than one order of magnitude lower than that of the metal superconductors at 4.2 K. The Bi (2212) wire will find an application in the very high field magnets over 20 T cooled by liquid helium and Tl (1223) in magnets between 1and 5 T cooled by liquid nitrogen. Superconducting magnets operable in liquid nitrogen will be used widely in manyfields of industry whenever a high magnetic field is required.

REFERENCES 1. Bednorz, J. G., and Miiller, K. A., Z Phys., BM,189 (1986). 2. Uchida, S., Takagi, H., Kitazawa, K., and Tanaka, S., Jpn. J. Appl. Phys., 26, L1 (1987). S., Jpn. J. Appl. Phys., 26, 3. Uchida, S., Takagi,H.,Kitazawa,K.,andTanaka, L151 (1987). 4. Wu, M. K., Ashburn, J. R., Torng, C. J., Hor, P. H., Meng, R.L., Gao, L., Huang, Z. J., Wang, Y. Z., and Chu, C. W., Phys. Rev. Lett., 58; 908 (1987). 5. Maeda,H.,Tanaka,T.,Fukutomi,M.,andAsano,T., Jpn. J. Appl. Phys., 27, L209 (1988). 6. Sheng, Z. Z, and Herman, A. M., Nature, 332, 138 (1988). Sheng, Z. Z., Kiehl, W., Bennett, J., El Ali, A., Marsh, D., Mooney, G. D., Arammash, F., Smith,J.,Viar, D., andHerman,A.M., Appl. Phys. Lett., 52, 1738 (1988). 7. Matsuda, S., Takeuchi, S., Soeta, A., Suzuki, T., Aihara, K., and Kamo, T., Jpn. J. Appl. Phys., 27, 2062 (1988). Matsuda, S., Takeuchi, S., Soeta, A., Suzuki, T., Aihara, K., and Kamo, T.,in Advances in Superconductiviv, (K. Ishiguro, ed.) Springer-Verlag,Tokyo (Proc. Int. Symp. Superconductivity, August28-31, 1988, Nagoya, Japan), 1988, p. 804. 8. Matsuda, S. P., Shokubai (Jap. J. Cataly.), 31, 279 (1989). 9. Sleight, A. W., Gillson, J. L., and Bierstedt, P. E.,Solid State Commun., 17, 27 (1975). 10. Cava, R. J., Batlogg, B., Krajewski, J. J., Farrow, R.,Rupp, L. W., White, A. E., Short, K., Peck, W. F., and Kometani, T., Nature, 332, 814 (1988). 11. Tokura, Y., Takagi, H., and Uchida, S., Nature, 337, 345 (1989). 12. Oda,M.,Murakami,T.,Enomoto,Y.,and Suzuki, M., Jpn. J. Appl. Phys., 26, L804 (1987). 13. Soeta, A., Suzuki, T., Takeuchi,S., Kamo, T.,Usami, K., and Matsuda,S. P.,Jpn. J. Appl. Phys., 28, L 1186 (1989). 14. Kamo, T., Doi,T., Soeta, A.,Yuasa,T.,Inoue, N., Aihara,K.,andMatsuda, S. P., Appl. Phys. Lett., 59, 3186 (1991).

HIGH-TEMPERATURE SUPERCONDUCTORS

15. 16. 17. 18. 19. 20. 21. 22. 23. 24. 25.

26. 27. 28. 29. 30. 31. 32. 33.

463

Doi, T., Nabatame, T., Kamo, T., and Matsuda, S. P.,Supercond. Sci. Technol., 4, 488 (1991). Cava,R. J.,Batlogg, B., Chen,C.H.,Rietman, E. A., Zahurak, S. M., and Werder, D., Phys. Rev., B36, 5719 (1987). Morgan, P.E.D., Piche, J. D., and Housley, R. M., Physicu C, 191, 179 (1992). Kase, J. I., Togano, K., Kumakura, H., Dietderich, D. R., Irisawa, N., Morimoto, T., and Maeda, H., Jpn. J. Appl. Phys., 29, L1096 (1990). Nabatame, T., Watanabe, K., Awaji, S., Saito, Y., Aihara, K., Kamo, T., and Matsuda, S. P.,Jpn. J. Appl. Phys., 31, L1041 (1992). Doi, T., Okada,M.,Soeta,A.,Yuasa, T., Aihara, K.,Kamo,T.,andMatsuda, S. P., Physicu C, 183, 67 (1991). Ozkan, N., Glowski, B. A., Robinson, E. A., and Freeman, P. A., J. Muter. Res., 6, 1829 (1991). Kimura, T., Kishio, K., Kobayashi, T., Nakayama,Y.,Motohira,N.,Kitazawa, K., and Yamafuji, K., Physicu C, 192, 247 (1992). Jin, S., Tiefel, T. H., Sherwood, R. C., Davis, M. E., van Dover, R. B., Kammlot, G. W., Fastnacht, R. A., and Keith, H. D., Appl. Phys. Lett., 52, 2047 (1988). Murakami, M., Morita, M., Doi, K., Miyamoto, K., and Hamada, H., Jpn. J. Appl. Phys., 28, L399 (1989). Murakami, M., Mod. Phys. Lett., B4, 285 (1990). Murakami, M., Flux pinning of melt textured processed YBCO superconductors and their applications, in Studies of High Temperature Superconductors, Vol. 9 (A. V. Narlikar, ed.), Nova Science, New York, 1991, pp. 1. Ogawa, N., Hirabayashi, I., and Tanaka, S., Physicu C, 177, 101 (1991). Morita, M., Tanaka, M., Takebayashi, S., Kimura, K.,Miyamoto,K.,and Sawano, K., Jpn, J. Appl. Phys., 30, L813 (1991). Matsuda, S. P.,Soeta, A., Doi, T., Aihara, K., and Kamo, T., Jpn. J. Appl. Phys., 31, L1229 (1992). Sasaoka, T., Nomoto, A., Seido, M., Doi, T., and Kamo, T., Jpn. J. Appl. Phys., 30, L1868 (1991). Matsuda, S., Okada, M., Morimoto, T., Matsumoto, T., and Aihara, K., Muter. Res. Soc. Symp. Proc., 99, 695-698 (1988). Wilhelm, M., Neumuller, H. W., and Ries, G., Physicu C, 185-189, 2399 (1991). (Proc. Int. Conf. M2S-HTSC 111, Kanazawa, Japan, July 1991). Okada, M., Nabatame, T., Yuasa, T., Aihara, K., Seido, M., and Matsuda, S. P., Jpn. J. Appl. Phys., 30, L2747 (1991). Matsuda, S. P.,Doi, T., Soeta, A., Yuasa, T., Inoue, N., Aihara, K., and Kamo, T., Physicu C, 185-189,2281 (1991) (Proc. Int. Conf. M2S-HTSC III, Kanazawa, Japan, July 1991).

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Preparation and Properties of Tantalum Oxide Thin Films by Sol-Gel T. Ohishi Hitachi, Ltd. Ibaraki, Japan

1.

INTRODUCTION

Dielectric thin films are used in a variety of modern electronic devices, such as thin-film condensers and large-scale integrated (LSI) and electroluminescent (EL) devices [l]. Demands are strong for thin films having high dielectric constants and good breakdown-voltage characteristics for use in electron devices with superior performance and high degrees of integration [2].There are several ways of preparing these films. Sputtering as a physical method of preparation and CVD (chemical vapor deposition) as a chemical method are well known. However, these methods have drawbacks; they require expensive vacuum equipment, and manufacturing of films with a large surface area or multicomponent films is difficult. The sol-gel method, which is a chemical method of film preparation, has received much attention for its advantages over the sputtering and CVD methods [3 & 61; it uses simple equipment and can produce large area and multicomponent films with high homogeneity at comparatively low temperatures. This chapter discusses the preparation of tantalum oxide thin films using the sol-gel method, which we believe is promising as a practical method for producing dielectric thin films whose dielectric constants and breakdown voltages are both comparatively high, Tantalum oxide films are good candidates for insulators in LSI devices, EL devices, and film capacitors [4].Reports on the for465

466

OHISHI

mation of fine particles of tantalum oxide prepared by the sol-gel method are avalable, but a detailed report on thin-film preparation is lacking.

II. PREPARATION OF TANTALUM OXIDE THIN FILMS The basic reactions in the sol-gel method are the hydrolysis and condensation reactions shown here, which should proceed at a moderate rate.

+

Ta(OR)5 H,O

+ (R0)4Ta- OH + ROH

(R0)4Ta- OH + Ta(OR),

hydrolysis (R = alkyl group)

+ (RO),Ta - 0 - Ta(OR), + ROH

condensation

Factors influencing these reactions include the type of alkoxide reagents, the types of solvents and catalysts, the concentration, and the reaction temperature [6]. In practice, these factors affect each other in complicated ways, and ensuring that these reactions proceed ideally is not easy. It is necessary to identify the most influential factors and to optimize the reaction conditions. Dense transparent films are essential for insulation in electron devices. To produce high-quality films, a transparent sol suitable for coating must be prepared. The process for preparing these thin films is shown in Fig. 1. Ta (OC2Hs)s wasused as starting material. To the ethanol solution of Ta(OC2H5)5, the ethanol solution of H20 necessary for hydrolysis wasaddedslowly. After about 15 minutes a white spherical precipitate formed. Spectral and chemical

No HCI I

Proper Amount of HCI

I Excess HCI

Spin Coating

t i & Heat Treatment

Transparent Films

Figure 1 Syntheticprocedure for Ta2oS thin films.

TANTALUM OXIDE THIN FILMS BY SOL-GEL

467

analysis revealed that it had a structure similar to that of Ta(0H)s. The hydrolysis reaction between Ta(OC2Hs)s and H20 proceeded faster than the condensation reaction, and the alkoxy groups bonded to the Ta all appeared to be replaced with hydroxyl groups. When the correct amount of HC1 was added, however, no precipitates formed, and a uniform sol solution was obtained. The addition of HCl balanced the consecutive processes of the hydrolysis and condensation reactions. When excess HC1 was added, however, the sol solution became a turbid white. The reason for this is thought to be particle growth at the high ionic strength. The factors with the greatest influence on these reactions are clearly the ratios of Ta(OC2Hs)s, H20, and HC1. Examining the ratios of these three ingredients in detail revealed that a clear sol solution was best achieved with proportions of5-18 mol% Ta (OC2H6)s,82-95mol% H20, and 0.54.0 mol% HC1. Spin coating of the substrate surface with any of the three solutions (the one with the white precipitate, the clear sol solution, and the turbid white sol) produces a thin film. The films could then be densified through heat treatment. The thickness of the film can be controlled through the concentration of the sol, the number of revolutions per minute during spin coating, and the number of layers of spin coating before firing.

111.

RESULTS AND DISCUSSION

A.

Scanning Electron Microscope Observation of Surface and Cross-sectional Structure of Taz05 Thin Films

Figure 2 shows scanning electron microscope (SEM) photographs of the surface of thin films obtained from the three solutions. Spherical particles 1 4 pm in diameter can be observed in film (Fig. 2a) obtained from the solution with the white precipitates and no HCl. Film prepared from the clear sol with 1 mol% HCl added (Fig. 2b) is smooth and transparent. Excess HCl produced the white turbid sol, which resulted in the film in Fig. 2c. The smoothness is preserved here, but the film is opaque. The surfaces of the three films after firing at 800°C are shown in Fig. 3. The particles in Fig. 3a have shrunken visibly. No change is apparent in Fig. 3b, but the film in Fig. 3c is peeling. The reason for this peeling is not clear, but the cracks may be caused by vaporization of the excess HCl during firing. Thus, the thin film obtained from the clear sol prepared with precisely controlled quantities of Ta(OC2H5)s, H20, and HC1 exhibits favorable structural properties between the room temperature and 800°C. Figure 4 shows SEM micrographs of the cross sections of a thin film prepared using the sol-gel method. For comparison, a film prepared using the sputtering method is also shown. A fine film prepared by the sol-gel method was

468

n 0 U

c

4 0

470

cN

0“ c-

U

OHISHI

TANTALUM OXIDE THIN FILMS BY SOL-GEL

471

formed on the IT0 substrate, but there is no definite border between the IT0 film and the Ta2O5 film prepared by the sputtering method and the film thickness is not uniform. This is because, during sputtering, the IT0 film itself was sputtered off. The conductivity of the IT0 film thus decreases to two-thirds o€ its originalvalue after the Ta205 film is formed. This phenomenon is well known. It is not seen in films prepared by the sol-gel method, so there is no adverse influence on the first coating film by Ta205 film.

B.

Characteristics of TasOs Thin Film

We investigated how the properties of Ta205 films prepared by the clear sol vary according to the firing temperature. 1. Infrared Spectra Figure 5 shows the infrared absorption spectra of Ta205 films dried at room temperature and fired at 200, 400, and 600°C. The room temperature film displays a wide variety of absorptions attributed toTa-0 stretching vibrations and Ta-O-Ta bending vibrations in the range of 400-1000 cm-1, indicating the formation of Ta205. In addition, -CH stretching vibrations and -CH2- rocking vibrations for ethoxy groups (-0C2H5) around 2900 and 1400 cm-1 suggest that

I

4000

l

l

3000

1600 1200

800

400

WAVENUMBER (cm-’) Figure 5 Infraredspectra of TazOs thin films after firing at

various temperatures.

OHISHI

472

unreacted organic materials remain in the film. Also, absorptions at 3500 and 1600 cm-1 indicate the presence of water. As the firing temperature increases, absorption caused by organic materials and water decreases and is not observed in films fired at temperatures over 400°C. 2.

X-ray Diffraction Analysis The x-ray diffraction patterns of Ta20s films fired at temperatures from room temperature to 800°C are shown in Fig. 6. At room temperature and 600"C, no distinct peaks are observed; the material is amorphous. Peaks appear beginning at 700°C. indicating that the film has begun to crystallize. At 800"C, the peaks attributable to orthorhombic p-Ta205 are visible, and the film is completely crystallized. This crystallization temperature is much higher than the 600°C for Ta2O5 films prepared using the sputtering and CVD methods [7]. 3. Differential Thermal Analysis Figure 7 shows the results of differential thermal analysis (DTA)

of the film A wide exothermic peak around 280°C and a

prepared at room temperature.

i r.t. I

20

I

I

40

60

I J

80

28 Figure 6 X-ray diffraction patterns of Ta205 thin films after firing at various temperatures.

TANTALUM OXIDE THIN FILMS BY SOL-GEL

473

sharp peak at 720°C can be seen. The infrared spectrum and x-ray diffraction results indicate that organic materials remain in the Ta2O5 film and that crystallization begins at over 700°C. These facts suggest that the wide exothermic peak around 280°C in the DTA curve is attributable to the combustion of the organic materials and the sharp peak at 720°C is caused by crystallization of the amorphous film. 4. Refractive Index Figure 8 shows the relationship between firing temperature and the refractive indices. The refractive index for the room temperature film is 1.73, but this value increases with the firing temperature, reaching 1.96 at 400°C. For firing temperatures above 500"C, however, the value decreases, and at 800°C is 1.90. The initial increase in the refractive index is presumed to correspond to the increase in the film's densification as the unreacted organic materials bum off. The decrease in the refractive index at higher temperatures is thought to be caused by the lower density as the film crystallizes. This is discussed in detail .with the films' microstructure. 5. ElectricalProperties and Microstructure Figure 9 shows the relationship between firing temperature and the dielectric constant of the Ta205films. For the film prepared at room temperature, the dielectric constant is e = 12. This value increases with the firing temperature, for example, E = 28 at 400°C. At temperatures above 400"C, the value decreases slightly. The rise in the dielectric constant between 200 and 400°C is striking. This is attributed tothe elimination of organic materials with low dielectric constants from the Ta205 films. According to the infrared spectra, organic materi-

2

U W T

O b 0 0 200 400 600 800 1000 z W

TEMPERATURE("C1

Figure 7 DTA curve for Ta205 thin film.

474

OHISHI

L 0

I

200

I

400

I

600

1

800

1

TEMPERATURE("C) Figure 8 Refractive index of Ta205 thin film versus firing temperature.

i i5

0

U

c i, W -l

W n

I 0

I

200

I

400

I

600

I

800

J

TEMPERATURE("C) Figure 9 Dielectric constant of Ta205 thin film versus firing temperature.

TANTALUM OXIDE THIN FILMS BY SOL-GEL

475

als are present in the film fired at 200"C, but no absorption by organic materials is apparent in the 400°C film. The DTA results also show heat generation at around 280°C from the combustion of organic materials. The dielectric constant & = 28 of the 400°C film is equal to that of films prepared by sputtering. The relationshipbetween firing temperature andthebreakdown voltage characteristics of Ta205 films is shown in Fig. 10. The breakdown voltage of the room temperature film is 2.5 MV/cm, and this value increases with the firing temperature. It is about 3.0 MV/cm at 400 and 500°C. Above this, however, it drops sharply. For a firing temperature of 600°C it is 1.6 MV/cm, and for 800°C is only 0.6 MV/cm. This abrupt drop in the breakdown voltage is thought to be caused by changes in the microstructure of the films during heating. No notable change in the surface of the Ta205 film isvisible by SEM observation. Thus the changes must be very small. We therefore examined the microstructure of the Ta2O5 films using transmission electronmicroscopy (TEM). Figure 11 showshowtheTa2O5 film microstructure changes withfiring temperature, as observed through TEM. The thin film prepared at room temperature is a conglomeration of fine spherical particles. The 400°C film has the same spherical structure, but the size of particles is slightly larger. It is pre-

a

(3

50 >

S0

n Y

a

W U

m

I

0

200

I

400

I

600

I

J

800

TEMPERATURE("C) Figure 10 Breakdown voltage of Ta205 thin film versus firing temperature.

Figure 11 E M micrographs of Ta205thin films (16.7 nm)heatedat various temperatures: (a) room temperature, (b) 400"C, (c) 650°C (d) 750°C.

TANTALUM OXIDE THIN FILMS BY SOLGEL

i

477

478

OHISHI

sumed that the particles grow during firing. The electron diffraction patterns for both show halo patterns; both are amorphous. The structure of the film fired at 650°C is quite different from the spherical structure. Furthermore, its selective area electron diffraction pattern shows a clear diffraction grating pattern, indicatingthatcrystallization is taking place.Some places inthis film still show halo patterns, however, depending on the area examined by TEM. Thus this is apparently a mixture of amorphous and crystalline materials. The 750°C film consists of thin flakes and thereare pinholes. The finestructure of the film is not preserved. The electron diffraction pattern showed a perfect diffraction grating pattern, indicating that the film has entirely crystallized. The microstructure of Ta205 films thus varies greatly depending on the firing temperature. The fine structure is preserved through 400°C but disappears with crystallization. Therefore, the sudden drop in the breakdown voltage for firing temperatures above 600°C is caused by degradation of the fine structure and the appearance of pinholes with rearrangement of the microstructure during crystallization. A firing temperature of 400°C is clearly best for preparing Ta205 film with both high dielectric constants and favorable breakdown voltage characteristics.

IV.

CONCLUSIONS

Tantalum oxide thin films were prepared using the sol-gel method, and their spectralandelectricalpropertieswere examined. The following conclusions were derived: 1. To obtain a fine, transparent film, HCl should be added during the hydrolysis of tantalum &oxide. An amount between 0.5 and 4.0 mol% is

2.

3. 4.

5.

6.

best. Tantalum oxide films prepared by the sol-gel method are amorphous from room temperature to 600°C and are completely crystallized at 720°C. Thin films prepared at room temperature contain organic materials,but these can be eliminated by firing at 400°C or above. The dielectric constant of the thin films increases as the organic materials are eliminated. For films processed at 400"C, E = 28. The breakdown voltage of the film increases with the firing temperature and is about 3.0 MVkm for films at 400°C. This value drops sharply for films fired above 600°C. This is attributed to degradation of the fine structure as a result of crystallization. To obtain a tantalum oxide thin film with a high dielectric constant and favorable breakdown voltage characteristics, it is bestto fire the film at 400°C.

TANTALUM OXIDE THIN FILMS BY SOLGEL

479

REFERENCES 1. (a) Ohta, K., Yamada, K., Shimizu, K., andTarui,Y.,Quadruplyself-aligned

VLSI dynamic stackedhigh-capacitance RAM usingTa2Oshigh-density memory, IEEE Trans. Electron Devices, 29, 368 (1982). (b) Tiku, S. K., Choice of dielectrics for " E L displays, IEEE Trans. Electron

Devices, 31, 105 (1984). 2. Asai, S., Trends in megabit DRAMS, IEDM Tech. Dig., I , 6 (1984). Melnick,B.M.,Araujo,C.A.,Mcmillan, L. M.,Carver,D.A..andScott.

J.F.,Recentresultsonswitching,fatigueandelectricalcharacterizationof sol-gelbased PZT capacitors,in hoc. SecondSymp. on IntegratedFerroelectrics, 1990, pp. 79-93. Sheppard, L.M., Advances in processing of ferroelentric thin films, Am. Ceramic Soc. Bull., 71, 85, (1992).

Yamagishi, K., and Tarui, Y., Photo-CVD of tantalum oxide film from pentamethoxytantalumfor VLSI dynamicmemories, Jpn. J. Appl.Phys., 25, L306 (1986). Shinriki, H., Kisu, T., Kimura, S., Nishioka,Y.,Kawamoto,Y.,andMukai, K., Capacitortechnologycompatiblewithadvanced VLSI process, in Symp. VLSI Technology Tech. Dig., San Diego, 1988, pp. 29-30. Ogihara, T., Ikemoto, T., Mizutani,N., and Kato, M., Formation of monodispersed Ta20s powers, J. Mater. Sci., 21, 2771 (1986). Jean, J., Synthesis of monodispersed Ta20s powers, J. Mater. Sci., 25, 1013 (1990). 6. Brinker, C. J., and Scherer, G.W., Sol-Gel Science, Academic Press, San Diego, (1990). 7. Okada, M., Preparation and properties of TazOs thin film prepared by CVD and sputteringmethod(Japanease), Denki Kagaku (Electrochemistry), 53, 109 (1985).

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21 Crystalline and Amorphous Thin Films of Ferroelectric Oxides Ren Xu University of Utah Salt Lake City, Utah

1.

INTRODUCTION

Chemical and physical processing techniques for ferroelectric thin films have undergone explosive advancement in the past few years (see Ref. 1, for example). The use of PZT (PbZr1-~TixO3)family ferroelectrics in the nonvolatile and dynamic random access memory applications present potentially large markets [2]. Other thin-film devices based on a wide variety of ferroelectrics have also been explored. These include multilayer thin-film capacitors [3], piezoelectric or electroacoustic transducer and piezoelectric actuators [M], piezoelectric ultrasonic micromotors [7], high-frequency surface acoustic devices [8,9], pyroelectric intrared (IR) detectors [lo-121, ferroelectric/photoconductive displays [13], electrooptic waveguide devices or optical modulators [14], and ferroelectric gate and rneWinsulatodsemiconductortransistor (MIST) devices [15,16]. Recently, efforts have been devoted to the fabrication and characterization of PbZr~~TixO3 family thin films for their potential applications in nonvolatile memory devices (See Ref. 17, for example). Partly because of the convenient stoichiometry control during processing, it was found that chemical methods, such as sol-gel and metal organic decomposition (MOD), are superior to physical means in many aspects. To appreciate better the science and technology of ferroelectric thin-film fabrication, it is important to give a brief account of the past efforts and the presentstatus and, it is hoped, shed some light on the future. 481

482

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The first report of a wet chemical processing of ferroelectric thin film was by Fukushima et al. in 1975 [18]. They reported the use of a mixed alkoxide and organic salt precursors in the fabrication of BaTiOs film. Application of sol-gel processing for the PZT thin films was started in 1984 by Wu et al. [l91 and Fukushima et al. [20] and followed by Budd et al. in 1985 [21]. More recently, continuing efforts in the processing of PZT family thin films by sol-gel and MOD methods can also be found in the literature [22-271. Meanwhile, chemical processing of thin films of other ferroelectric oxides resulted in remarkable progresses. One now can make ferroelectric thin films with crystallinities ranging from polycrystalline to texture-oriented polycrystalline to epitaxial in nature. These include (1) polycrystalline films by Hirano and Kat0 [28] for LiNbOs, Xu et al. [29] for Sr1-xBaxNb206 (SBN), Francis and Payne [30] for PMN-PT, Swartz et al. [31] for Pbl-,, LayZr~-~TixOs, and Nazeri-Eshghi et al. [32] for KNbO3; (2) texture-oriented films byXu et al. [33,34] for SBN and LiNbOs and by Hagberg et al. [35] for LiNbOs; and (3) epitaxial films by Partlow and Greggi [36], Hirano and Kat0 [37], and Xu et al. [34] for LiNbO3, Cheng et al. [38] for KNbO3, and Hirano et al. [39] for PbTiOs. In the following sections, we attempt to discuss some important aspects of sol-gel processing technique through a few model systems. Examples are given to illustrate the formation of various morphologies.Enoughdetails are discussed to aid those readers interested in sol-gel processing of other multicomponent oxide thin films. Finally, we devote one section to the fabrication and characteristic of amorphous thin films of a few ferroelectric compositions made by sol-gel. We further present some phenomenological observations and our preliminary attempts to understand them.

II. SOLUTIONPREPARATION It is now appropriate to differentiate the sol-gel method from MOD. In general, we consider a process sol-gel if it involves these two characteristic chemical reactions: (1) hydrolysis of at least one precursor compound and (2) polycondensation of the hydrolyzed precursor compound. A MOD process involves the thermal decompositions of all precursor compounds physically condensed onto substrates through rapid solvent evaporation. Therefore, methods using all alkoxide andpartially alkoxide precursors are considered sol-gel processing, and those using precursors other than alkoxides are MOD. One last interesting point to emphasize is that one can useall alkoxide precursors ina MOD process. This is possible when neither of the two characteristic reactions for sol-gel is involved and the reactions used to form the oxide is simply thermal decomposition of the metal alkoxides.

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THIN FILMS

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Like any other thin-film deposition techniques, sol-gel processing is essentially a mass transport process. The transformation of a liquid solution to a solid crystalline film is accomplished through three steps: 1. Precursor materials are dissolved in a homogeneous solution, thus assuring

molecular-level mixing of different precursor compounds. 2. Mass transport is completed upon spin or dip coating of a thin layer of the solution onto the substrate surface. Because of a combined process of hydrolysis-evaporation-polycondensation,a thin layer of amorphous gel film is formed on the substrate. 3. The as-deposited thin film together withthesubstrate is thenheatedto cause densification and crystallization of the film. A simple schematic diagram is shown in Fig. 1,where a similar diagram for MOD is also included for comparison. Note that for sol-gel and MOD steps 1 and 3 are similar, but step 2 for MOD is more physical than chemical in nature in that the formation of the amorphous film on the substrate is due solely to the evaporation of solvent and the disordered precipitation of solutes from the supersaturated solution. Ferroelectrics of interest are often multicomponent oxides. The metal elementsin concern do not always show enough solubilityingivensolvent. Therefore, the choice of precursor compounds and the dissolution procedures, as well as their behaviors toward moisture and heat, are all important aspects to consider. We select a few representative ferroelectric systems to illustrate the practices reported on the solution preparations.

A.

PbZr~-~TixOs and Pb1-yLa~Zrl-~TixO3 Systems

Homogeneous solutions containing Pb, Zr, Ti, and La in the correct proportion can be made with soluble precursors in alcoholic solvents. Table 1 lists examples of precursors and the corresponding references. The choices of precursor compounds for Zr, Ti, and La [40] in these examples are mainly for solubility reasons. The use of methoxyethanol as solvent has two effects: first, it is a chelating agent, which prompts the dissolution of lead acetate, and second, it is an ideal solvent with enough volatility to be used in practical coating procedures.

B.

S r i - ~ a d b n O sSystem

SBN is a solid solution system containing strontium, barium, and niobium oxides. The sol-gel solution preparation requires special precautions. This is because the alkaline-earth metal alkoxides dissolve in alcohols slowly and they are extremely moisture sensitive. Although alkoxides of both barium and stron-

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Solution

1

Step 2.

Step 3.

I

Spin’Dip Coating

I

Hydrolysis-Polycondensation

Treatment Heat

I

Hydrolysis-Polycondensation

MOD

Step 2.

I 7:z:iI

Step3.1

Heat Treatment

Evaporation

l

Thermal Decomposition

Figure 1 Sol-gel and MOD methods.

tium can be made easily by allowing reactions between the metal powders directly with the corresponding anhydrous alcohol, the solution thus made is difficult to preserve even under a dry atmosphere. Precipitation problems may arise because of either slow dissolution of Ba(OR)2 and Sr(OR)2 in thesolvent or the hydrolysis of the alkoxide to form less soluble hydroxides. On the other hand,reacting these metal powderswith an alcoholicsolution of Nb(OR)5 helps the dissolution of the alkaline-earth alkoxides caused by the formation of double-alkoxide complexes, for example BaNb2(OC2H5)nr which are readily soluble in the parent alcohol [41]. Xu et al. [29] prepared an ethanol solution of 1 mol of 60% Sr(OC2H5)2 + 40% Ba(OC2H5)2 and 2 mol Nb(OCzH5)5 by reacting strontium and barium

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Table 1 Common Precursors Used for PZT and PLZT Thin Films

ecursor

Metal

232-Ethylhexanoate 22, 20,n-Propanol 19, Acetate Ethoxide Neodecanoate zr n-Propanol n-Propoxide Acetylacetonate 20 26 n-Butoxide Ti n-Propanol i-Propoxide n-Butoxide ZEthylhexanoate 40 La

Pb

Methoxyethanol 2521, 26 40 19,21-23,25,40 19.21-23,25,40 20,26

metals with anhydrous ethanol and subsequently mixing with Nb(OC2&)5 lution in ethanol.

SO-

C. LiNbOs System This is one of cleanest and simplest systems with which to work. Lithium and niobium ethoxides are readily soluble in ethanol. These same precursors and solvents are used widely by several research laboratories. However, the sol-gel processing procedure deviates markedly. Two types of solutions are made for LiNbOs thin-film deposition. The first contains stoichiometric LiOC2H5and Nb(OC2H5)5 witha certain amount of H20 addedin the form of dilute HOC2H5 solution and gels within a few days [36,37,42]. The second type contains no preadded water and is stable for a year [34]. It was argued that the &oxides involved here are quickly hydrolyzed within a few minutes by the ambient moisture [43]. Therefore, no water addition is necessary, except the moisture in the ambient during spin and dip coating. Further, we show in the latter part of this chapter that such an arrangement is a better way of controlling hydrolysis.

D.

KNb03 System

Nazeri-Eshghi et al. [32] were the first to report sol-gel processing of KNbO3 ceramic powders. Their method was later adopted for the deposition of epitaxial thin films by Cheng et al. [38] The precursors usedwere KOC2H5 and Nb(OC2H5)s. The preparation procedure is fairly simple. Potassium metal is allowed to react with anhydrous ethanol, and the solution is then mixed with Nb(OC2H5)s. The solutions are usually stable for a few weeks. It was found that the addition of small amount of 2-ethylhexanoic acid helps to modify the

ca

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hydrolysis behavior of the solution, and subsequently better coating was obtained [38].

E. Other Ferroelectric Systems Other reported ferroelectric systems include PbTiOs [39], BaTi03 [18], and PMN-PT [30]. Solution preparation for these systems is similar to the procedures already discussed. The reader should easily find precursors and appropriate procedures according to these examples.

111.

CRYSTALLIZATION

After the homogeneous solutions are prepared, simple coating procedures are used to deposit thin gel films by either dip or spin coating on substrates. Alcoholic solutions are generally capable of wetting all oxide substrates or metalsemiconductor substrates with a thin oxide layer on the surface. A simple capillary tube experiment was performed [43] to provide semiquantitative data on therelative surface tensionbetween alcoholic solutionsandcommon substrates. Table 2 lists the results. We use the product of the surface tension with the cosine of the contact angle ycose, to represent the afinity of the solution to the corresponding substrates. By normalizing all data against the ycose value for ethanol on a VWR Microslide, it was found that the relative surface tensions of ethanol on all substrates tested are comparable in value. Therefore, it was concluded that alcoholic solutions are a good solution system for coating on oxide substrates. After coating of the gel films on substrates, the thin films together with the substrates are usually slowly heated to a few hundred degrees Celsius and maintained at that temperature for a few hours to cause crystallization of the ferroelectric crystalline phase. Depending on the typeof substrate and the processing procedures, crystalline films with different morphologies can be obtained. The

Table 2 Relative Surface Tension y cos 8 of HOCfl5 on Various Substrates

Substrates

VWR Microslides Silicon (1 11) Fused Sapphire (0.12) LiNbOs (001)

y COS ely COS e 1.m

1.14

1.53

1.14 0.97

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487

most common product is polycrystalline film with no apparent preferred crystallographic orientations. We present a few examples in the following section.

A.

PolycrystallineFilms

Polycrystalline films are usually obtained when any one of the following substrate types is used polycrystalline substrates, amorphous substrates, and single-crystal substrates with large lattice mismatch. For example, LiNbO3 [34] and SBN [44] grownon silicon (1 11)and(100) are polycrystalline. SBN grown on fused silica was also polycrystalline [29]. PZT thin films were grown on a large numbers of substrates, nearly all of which were polycrystalline in nature (see Ref. 45, for example). When polycrystalline films are the inevitable result, problems associated with the formation of grain boundaries, impurities, second phases, and so on, are the central concern. For PZT family ferroelectrics, one of the main problems is the fatigue of the materials after large numbers of repetitive switchings, which is attributed to the domain wall pinning on defect structures, second phases, and grain boundaries [46]. A special heat treatment technique, rapid thermal annealing, is commonly used to reduce extensive grain growth, which was believed to improve the fatigue behavior of PZT thin films [45]. When optical application is the target for these polycrystalline films, scattering of light by the grain boundaries, second phases, and other defect structures are to be addressed. In general, the problems are far from being resolved at present.

B. Texture-OrientedFilms Texture-oriented films may be obtained under several special conditions. First, when single-crystal substrates with fairly large lattice mismatch were used, the film grown over the surface may be highly preferentially oriented.For example, for LiNbO3 grown on a sapphire (012)face, the x-ray diffraction patternsof the film indicated two principal directions only, the (012) and (300) [34]. Alternatively, when an amorphous substrate was used, it was found that because of the sensitivity of the ferroelectric crystallites toward the electrical field, the films can be grown with a preferred orientation if small a dc bias field is applied along the substrate surface direction during the heat treatment. This was reported by Xu et al. [33] for SBN grown on fused silica. Figure 2 shows this result. The x-ray diffraction pattern shows significant enhancement of peak intensities for (130), (121), (131), and (620)for the tungston bronze structure.These represent oxygen close-packing planes preferentially lined up along the external field direction. A third type of textural orientation was obtained in the arrangement in which LiNbO3 film was grown on platinum [35]. It was found that with a rapid heattreatmentprocedure,LiNbO3preferentiallyorientedwiththe(006) face parallel to the substrate surface. The cause of this orientation effect is not clear.

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(Sr0.60Ba0.40)Nb206 t h i n film on fused silica, 850 "C. 1 h; w i t h E = I kVlmm. thickness: 1300 A

240 200 v)

E

160

1

6 120 80

40

0

I60

(Sr0.60Ba0.40)Nb206 thin film on fused silica, 850 "C. I h; without E . thickness: 1300 A

I20 80

40

- .

I

I

30

10

I

Diffractionangle

1

50

I

70

2 0 (&g)

Figure 2 X-raydiffractionpatterns ofSro.6Bao.4Nb206thin films grownonfused silica. Sample heated in the presence of a dc field showed preferential orientation, and peaks with indices (130), (121), (131), and (620) were significantly enhanced. (AfterXu et al. [33].)

C.

Epitaxial Films

When single-crystal substrates with a small lattice mismatch are used, sol-gel produces epitaxial films for a few ferroelectric systems.Although epitaxial growth of crystalline films from an amorphous layer has been observed in the amorphous silicon to silicon transformation, sol-gel epitaxy onlybegan to emerge as a possible fabrication technique in the last few years. Hirano and Kat0 were the first to observe the epitaxial growth of LiNbO3 on the sapphire (1 10) face [37]. Xu et al. [34,43] found the epitaxial growth of L i m o 3 on the LiTa03 (1 10) face and the LiNbOs (006) face. Epitaxial KNbO3 was reported

489

FERROELECTRIC OXIDE THIN FILMS

on a SrTiO3 substrate [38] and PbNbo.02 (Zro.52 Ti0.48)0.9803on sapphire (012) [47]. Figure 3 shows a set of x-ray diffraction patterns of LiNbO3 grown on various single-crystal substrates. LiNbO3 on silicon (1 11) and (400) showed polycrystalline diffraction patterns but onsapphire (012) showed only two principal diffraction peaks, (012) and (300). On LiTaOs (110) and LiNbOs (006) the films grown were clearly epitaxial. The electron diffraction pattern and the energy-dispersive spectroscopic (EDS) analysis of composition at the interface LiNb03LiTa03 is showninFig. 4. The diffractionpatternshowstheview along the [221] zone axis; EDS shows the elemental distribution of Nb5+ and

LN/Si(lII)

5

=

Y

en

W O

LN/LT( I I O )

I

LN/LN(006)

10

50

30

70

TWO THETA

Figure 3 X-ray diffraction patterns of Limo3 films grown on various single-crystalsubstrates.PolycrystallinediffractionpatternswerefoundforLiNbo3onsilicon (1 11) and silicon (100). Highly oriented LiNbog was found on sapphire (012). Epitaxial LiNb03 was found on LiTaOs (110) and Limo3 (006).

0.5

0.4

0.3

0.2 W

0. l

0.5

0.0 0.0

2.0

0.0

2.0

ENERGY (IO3eV)

Figure 4 (a) Electron diffraction pattern of [221] zone at the interfacial area between LiNbO3 and LiTaO3 substrate. (b) EDS peaks of Nb(La, Le), Ta(Ma, Mp), and O(&) across the interface with a lo00 8, beam size. On the substrate side only Ta(Ma, Me) and O(&) were found; on thefilm side the Nb(La, LP) peak is predominant, with only a small peak corresponding to Ta(Ma, Mp).

FERROELECTRIC OXIDE THIN FILMS

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Tas+ across the interface line. The lattice mismatch between the film being grown and the substrates are probably responsible for the various degrees of orientation. Table 3 summarizes the percentage of lattice mismatches for L i m o 3 with various single-crystal substrates. In addition to a lattice-matched substrate, sol-gel epitaxy also requires excellent stoichiometry control of the solution. Most of the demonstrated sol-gel epitaxies were performed at relatively low temperatures. Therefore, molecularlevel homogeneity is essential. To appreciate fully the delicate features of solgel epitaxy, it is now appropriate to compare sol-gel processing with other thinfilm deposition techniques. In the attempt to achieve optical signal processing, modulation, amplification, and memory functions in integrated circuits similar to those on electrical signals by semiconductor devices, integration of ferroelectric devices is the ultimate goal. However, to achieve integration of microscopic devices based on materials as complex as oxide ferroelectrics, which are predominantly multicomponent metal oxide compounds, reliable thin-film deposition techniques are critically needed. One of the most important aspects of multicomponent oxide thin-film deposition is the control of stoichiometry. We can now try to understand the existing deposition techniques by analyzing the precursors. First, we introduce the terms compositionally “proper” and“improper”precursors.A compositionally proper precursor has on the molecular level precisely the same metal-metal ratio ( M ” ’ ) as the oxide compound one wishes to prepare. An improper precursor does not possess the correct ”/”‘ ratio on the molecular level. It may or may not possess the correct M’/”’ ratio on a larger scale, for instance, microscopic (or nanoscopic) scale, as in a micrometer-sized powder mixture. We immediately conclude that most of the existing deposition techniques use improper precursors. For example, such techniques as metal oxide chemical vapor deposition, multiple-target sputtering, and multiple-target evaporation use improper precursors. In these cases, single metal organic compounds or single metal oxide targets are used for the deposition, and thus arise the associated difficulties in stoichiometry control. Table 3 LatticeParametersandMismatchesBetween LiNbOs and Various Single-Crystal Substrates Crystal

CH (A)

aH (A)

Mismatch (%)

Limo3 LiTa03 Sapphire Silicon (111) Silicon (100)

13.864 13.785 12.991

5.150 5.155 4.758 3.84 5.43

0.00 -0.57 -7.61 -25.4

-

-

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xu

These are often caused by the variation in evaporation rate, sticking coefficients, and inhomogeneous spatial distribution of components in the deposition apparatus. Single-target techniques, such as sputtering, evaporation, and laser ablation, use a composite target with overall correct stoichiometry or sometimes the compound itself in polycrystalline form. However, if we consider the precursors after the liquid vapor or plasma has left the target, the liquid, vapor or plasma in these cases becomes improper before the moment of arrival on the substrate. In other words, the precursors lose the correct M'/"' ratio on the molecular level. Therefore, these depositions are in nature similar to other processes using improper precursors. One mightconsider these deposition techniques as having an improper deposition mode. Sol-gel processing, on the other hand, uses proper precursors and can facilitate a completely proper deposition, for example for the epitaxial growth of LiNbO3 thin films. Precursors used in this case are LiOC2Hs and Nb(OC2Hs)s. A double-alkoxide LiNb(OC2Hs)6 is prepared in alcoholic solution by extensive refluxing of a mixture of LiOC2Hs and Nb(OC2Hs)s [28,48]. The solution ofLiNb(OC2Hs)fj is a proper precursor. In addition, the double-alkoxide LiNb(OC2Hs)6 has nearest neighbor atomic arrangement remarkably similar to that of a crystalline LiNbO3, as shown separately in Fig. 5a and 5c. This solution is then used to coat thin films directly onto substrates. When the hydrolysis is controlled, the complex structure can be preserved during the solgel processing as a result of the steric hindrance of the bridging -0C2Hs groups to hydrolysis. In particular, if no water was added to the solution before the

Figure 5 Nearest neighbor shucturesof (a) double-ethoxide LiNb(OC&)6 in solution; (b) amorphous LiNbo3; and (c) crystallineLimos.

FERROELECTRIC OXIDE THIN FILMS

493

deposition, the hydrolysis occurs only during and after the spin or dip coating. Therefore, the preservation of such nearest neighbor structures is accomplished through immediate polycondensation after the double alkoxides are partially hydrolyzed by the ambient moisture. The as-deposited amorphous gelfilm, having a nominal composition LiNb0~x(OC2H5)2~, presents the nearest neighbor structure shown in Fig. 5b. Uponheating to a few hundred degrees Celsius L i m o 3 is crystallized through a local relaxation rather than extensive longdistance diffusion. Therefore, such sol-gel processing constitutes proper deposition with proper precursors. The mechanism of sol-gel epitaxy has not yet been explored in detail. The thermodynamic driving force for epitaxy is probably different from that for common solid-state epitaxy. Miller and Lange [49] studied a simpler system, zro;? epitaxial growth, and proposed a simplistic model for the epitaxy. It is generally considered that the nucleation in a gel film is three rather than two dimensional according to the Avrami model of thermal analysis. This suggests that the epitaxy occurs after nucleation. The Miller and Lange model is consistent with such an argument. A common problem associated with all singlecrystal films derived from sol-gelepitaxy, including LiNbO3[34,36,37], KNbO3 [38], PbTi03 [39], PNZT [47], and ZrO2 [49], is the porous and defective nature of the films. No effective solution to this problem has yet been reported.

W. AMORPHOUSFILMSOF FERROELECTRICS One of the direct outcomes of the proper deposition of a proper precursor by sol-gel processing is the possibility of preparing a structurally controlled amorphous thin film. Ferroelectricityhas been commonly associated with crystalline materials, although it has been postulated that this phenomenon can occur in an amorphous solid [51]. In the past,attempts were reported to demonstrate experimentally thepossibility of amorphous ferroelectricity.For example, a glassy L i m o 3 was produced by rapidly quenching molten LiNbO3 [52]. A dielectricanomalywasobserved in theglassyLiNbO3,attributed to possible amorphous ferroelectricity. Similarly, glassy LiNbO3, when doped with iron, showed an anomalous temperature dependence of resonance, quadruple splitting,and center shift, characteristic of a ferroelectricphasetransition[53]. Radio frequency sputtered amorphous LiNbO3 [54,55] and PbTiO3 [56] also showed anomalous dielectric behaviors below the crystallization temperature. However, the issue of amorphous ferroelectricity remains highly controversial because of the lack of concrete physical evidence for the existence of such a phenomenon. A more realistic approach to this controversy is probably to look at the materials from a strictly phenomenological point of view. We know that sol-gel processing is particularly convenient for the preparation of amorphous

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films on a wide variety of substrates. It is worthwhile to compare systematically the properties of an amorphous film of ferroelectric composition with those of a crystalline film.

A.

Experimental Results

Amorphous LiNbO3 films made by sol-gel processing were subjected to a series of characterizations [57]. It was found that an amorphous LiNbO3 film obtained by heating the gel film at 100°C for 2 h showed P-E hysteresis with remnant polarization PT = 10 $/cm2 and coercive field Ec = 110 kV/cm. Electron diffraction of such film showed a diffuse ring pattern characteristic of an amorphous nature. These are shown in Fig. 6 in which the scale for E is 147 kV/cm division and that for P is 5.6 pC/cm2 division. Further measurement showed a pyroelectric coefficient of 8 pC/cm2 K at 28°C. Note that for singlecrystal LiNbO3, PT = 50 pC/cm2 and the pyroelectric coefficient was reported to be 20 pC/cm* K [l]. Further, a piezoelectric resonance was observed at similar frequency range for both amorphous and crystalline LiNbO3, characteristic of a ferroelectric material [57]. A systematic comparison of the electrical and optical properties of amorphous and crystalline Pb(Zr0.52Tio.48)03 was reported by Xu et al. [57,58] It is noticeable that the amorphous film has lower values than the crystalline film for all parameters, consistent with the observations in a LiNbO3 system. The most remarkable characteristic of amorphous Pb(Zro.52Ti0.48)03 film is the exceptionally high pyroelectric coefficient. Figure 7 shows the pyroelectric current of an amorphous Pb(Zr0.52Tio.48)03 film measured 2, 44, and 162 h after poling at 150°C. The film remained highly pyroelectric several months after the original poling. It is thus possible to use these amorphous films in some applications, in this case pyroelectric heat sensors, for which crystalline ferroelectrics were the typical materials of choice. We have shown that although amorphous ferroelectricity as a physical phenomenon is not fully demonstrated, amorphous films of ferroelectric oxides have shown a number of useful properties that may warrant further studies of their structure and properties. It is possible to use them in limited areas in place of crystalline ferroelectrics.

B.

Structural Model

A model was proposed by Xu et al. [57] to elucidate the preceding observations in an amorphous material. As we have discussed, the nearest neighbor structure of the amorphous LiNbO3, as shown in Fig. 5b, consists of building blocksoflithium-andniobium-centeredoxygen octahedra sharing a face. When many of these building blocks come together, there are clusters of octahedral pairs in the amorphous film in the form shown in Fig. 8a. The schematic

FERROELECTRIC OXIDE THIN FILMS

495

Figure 6 (a) P-E hysteresis loop of anamorphous LiNbo3 film coated on a goldplated silicon wafer with a platinum top electrode (at 60 Hz; scale x axis; 147 kV/cm division, and y axis, 5.6 pC/cm2 division. (b) Electron diffraction pattern of the amorphous LiNbO3 film; the diffuse ring indicated the amorphous nature of the film.

xu

496 60

o 2hafterpoling o 44hafterpoling A 162hafterpoling heating after (24O0C/O

Q v

, !

E

40

3 0

.0

L

c. 0

P)

e5.

P)

20

n

Sample: Au I amorphous PZT(52/48) film I Au. Electrode area= 3.14 X 10-%m2. PZT thickness= 0.3pm. Poling condition: d.c. field= irkVImm, T = 150°C. 0 20

40

60

80

100

Temperature (%)

Figure 7 Pyroelectriccurrent of amorphous PZT (52/48) film on gold. Filmwas heated at 350°C for 1 h, and the thickness was 3000 A. (After Xu et al. [58].)

representation of these clusters shows oriented octahedral pairs within the cluster; adjacent clusters may possess a different overall orientation. By comparing this with a typical crystalline LiNbO3 structure (Fig. 8b), we make the following observations: (1) each individual cluster may possess a unique polarization throughtheinteraction of neighboring octahedral pairs; (2) thepolarization within a cluster may be smaller than in the crystallinestructure of comparable size; (3) the order within a cluster need not be that of a crystalline order, since the sequence neednot be that of oxygen octahedra centered with Nb-Li-vacancy-Nb-Li-vacancy; (4) the adjacent clusters may interact to favor one polarization direction to allow projection of polarizations along a unique macroscopic direction;and (5) whensubjecttoanexternalelectricalfield,the polarizations of individual clusters (referred to as ferrons in Ref. 57) may be reversed to allow macroscopic polarization along the external field direction. These observations are preliminary: the model is also our first attempt to understand the observations. Although ferroelectricity can be consistent with an amorphous structure in theory, to be able to demonstrate such a phenomenon unequivocally is by no means an easy task. However, the preceding discussion may be helpful in shedding light on future efforts in the sense that it suggests a possible avenue to prepare “structurally controlled” amorphous materials, which may be essential to the preparation of any amorphous material with “locally dialectically soft” structural units, as proposed by Lines [51]. After all,

(b) Figure 8 (a) Two clusters of orientedoctahedralpairsand of the octahedral pairs in crystalline LiNbO3.

(b) typicalarrangement

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ferroelectricity relies on the existence of reversible electric dipoles, which by andlarge are constructed through a certain degree of shortto intermediate range order.

V.

SUMMARY

It has been demonstrated that sol-gel processing can be used to grow ferroelectric thin films with a wide range of qualities: they are polycrystalline, oriented, epitaxial, and amorphous. The advantages of sol-gel over other methods include convenient stoichiometry control, low heat treatment temperature, and simple processing procedures. One could easily grow and study ferroelectric thin films of all varieties with simple laboratory equipment. However, it is important to be aware that for the sol-gel films to compete with those from other deposition techniques, especially integrated ferroelectric device applications, a systematic analysis of current deposition techniques, which may provide insight into the future development of the sol-gel method, is critically needed. Areas in need of improvement include (1) the nanopores and associated defect structures in the sol-gel epitaxial films; (2) better understanding and systematic study of the solution structure and its effect onsol-gel processing; and (3) more rigorous structural and property investigationsof the amorphous films made by sol-gel and the correlation between these. Sol-gel processing is a very promising method for the fabrication of multicomponent oxide thin films. There are still many unanswered questions and many obstacles to overcome. Nonetheless, there is little doubt that a brighter future lies ahead for this technique given the vast numbers of areas of demonstrated applicability for this method and the degree of success reported to date.

REFERENCES 1. Xu, Y . H.Ferroelectric Materials and Their Applications, North-Holland, Amsterdam, 1991. Ferroelectric Thin Films (E. R. Myers and A. Kingon, eds.), Materials Research Society, Pittsburgh, (1990). 2. Chapman,D. W., J. Vac. Sci.Technol., 9, 425(1972).Pazde Araujo, C. A., McMillan, L. D., Melnick, B. M., Cuchiaro, J. D., and Scott, J. F., Ferroelectrics, 104, 241 (1990). 3. Feuersanger, A. E., in Thin Film Eielectrics (F.Vratny, e d . ) , Electrochem. Soc., New York, (1969), p. 209. 4. Destefanis, G. L.,Gaillaiard, J. P.,Ligeon, E. L.,Valette, S., Farmery, B. W., Townsend P. E., and Perez, A., J. Appl. Phys., 40, 420 (1979). 5. Foster, N. F.,J. Appl. Phys., 50, 7898 (1979). 6. Mansingh, A., Ferroelectrics, 102, 69 (1990). 7. Udajakumar, K. R., Bart, S. F., Flynn, A. M., Chen, J., Tavrow,L. S., Cross,

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

20. 21. 22.

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28. 29. 30. 31.

32. 33. 34.

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L. E., Brooks, R. A., and Ehrlich, D. J., Proc. IEEE Micro Electro Mechanical Systems (Nara, Japan), (1991), p. 109. Castellans, R. N., and Feinstein, L. G., J. Appl. Phys., 50, 4406 (1979). Adachi,H.,Mitsuyu,T.,Yamazaki, O., andWasa, K., J. Appl. Phys., 60, 736 (1986). Takayama, R., Tomita, Y., Lijima,K.,andUeda,I., J. Appl.Phys., 61, 411 (1987). Glass, A. M., and Abrams, R. L., J. Appl. Phys., 41, 4455 (1970). Takayama, R., Tomita, Y., Lijima, K., and Ueda, I., J. Appl. Phys., 61,411 1987. Chapman, D. R., and Mehta, P. R., Ferroelectrics, 3, 101 (1972). Webster, J. C., and Zernike, F., Ferroelectrics, IO, 249 (1976). Wu, S. Y., Ferroelectric, 11, 376 (1976). Wu, S. Y., IEEE Trans. on Electron. Devices, ED-21, 499 (1974). Symp.Proc. lst, 2ndand3rdInt.Symp.Integr.Ferroelect.,ColoradoSprings, Colorado, 1989,1990,1991. Fukushima, J., Kodaira, K., Tsunashima, A., and Matsushita, T.,Yogyo Kyoaishi, 83, 204 (1975). Wu, E., Chen, K. C., and Mackenzie, J. D., in Better Ceramics Through Chemistry (C.J. Brinker, D. E. Clark, and D. R. Ulrich, eds.), North-Holland,New York, 1984, p. 169. Fukushima, J., Kodaira, K., and Matsushita, T., J. Mater. Sci., 19, 595 (1984). Budd, K. D., Dey, S. K., and Payne, D. A., Br. Ceram. Soc. Proc., 36,107 (1985). Chen, K. C., Janah, A., and Mackenzie, J. D., in Better Ceramics Through Chemistry Il, (C. J. Brinker, D. E. Clark, and D. R. Ulrich, eds.), Materials Research Society, Pittsburgh, (1986), p. 731. Lipeles, R. A., Coleman, D.J.,andLeung,M. S., in Better Ceramics Through Chemistry, II, (C. J. Brinker, D. E. Clark, and D. R. Ulrich, eds.), Materials Research Society, Pittsburgh, (1986), p. 665. Dey, S. K., and Zuleeg, R., Ferroelectrics, 108, 37 (1989). Dana, S. S., Etzold, K. F., and Clabes, J., J. Appl. Phys., 69, 4398 (1991). Tohge, N., Takahashi, S., and Minami, T.,J. Am Ceram. Soc., 74, 67 (1991). Sayer, M.,inSymp. Proc. 3rd Int. Symp. Integr. Ferroelect., Colorado Springs, Colorado, (1991), pp. 1-9. Hirano, S., and Kato, K., Adv. Ceram. Mater., 2, 142 (1987). Xu,R.,Xu,Y.H.,Chen, C. J., andMackenzie,J. D., J. Mater. Res., 5, 916 (1990). Francis, L. F., and Payne, D. A., in Ferroelectric Thin Films (E. R. Myers and A. Kingon, eds.), Materials Research Society, Pittsburgh, (1990), p. 173. SW&, S. L., Bright, S. J., Melling, P. J., and Shrout, T. R., Ferroelectrics, 108, 71 (1990). Nazeri-Eshghi, A., Kuang, A.X., and Mackenzie, J. D., J. Mater Sci., 25, 3333 ( 1990). Xu, Y. H., Chen, C. J., Xu,R., and Mackenzie,J. D., Phys. Rev. B., 41,35 (1991). Xu, R., Xu, Y. H., and Mackenzie, J. D., in Symp. Proc. 3rd Intern. Symp. Integr. Ferroelect., Colorado Springs, Colorado, (1991). p. 561.

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S., Eichorst,D. J.. andPayne,D.A.,in Sol-gel Optics (J. D. MackenzieandD. R. Urich, eds.),SPIE Roc. 1328, SanDiego,California, (1990), p. 466. 36. Partlow, D. P., and Greggi, J., J. Muter. Res., 2, 595 (1987). 37. Hiano, S., and Kato, K., Muter. Res. Symp. Proc., 155, 181 (1989). H., and Mackenzie, J. D., Muter. Res. Symp. Proc., 271 38. Cheng,C.H.,Xu,Y.

35. Hagberg,D.

(1992). 39. Hirano, S., Yogo, T., Kikuta, K., Kato, K., Sakamoto, W., and Ogasahara, S., in 93rd Annual Meeting 8~ Exposition of the American Ceramic Society, 43-SV-91, Cincinnati, Ohio, 1991. 40. Vest, R. W., and Xu, J., Ferroelectrics, 93, 817 (1989). Merul Albxides, Academic 41. Bradley,D.C.,Mehrotra,R.C.,andGaur,D.P., Press, London, (1978). p. 306. 42. Eichorst, E., and Payne, D. A., in Sol-Gel Optics, (J. D. Mackenzie and D. R. Wrich, eds.), SPIE Roc. 1328, San Diego, California, (1990), p. 456. 43. Xu, R., Ph.D. Dissertation, University of California, 1992. 4 4 . Chen, C. J., Xu, Y. H., Xu, R., and Mackenzie,J. D., J. Appl. Phys., 1763 (1991). 45. Proceedings of lst, 2nd. 3rd and 4th International Symposium on Integrated Ferroelectrics, Colorado Springs, Colorado, 1989-1992. 46. Ramesh,R.,Chan,W. K., Wilkens, B., Sands, T., Tarascon, J. M.,Keramidas, V. G., and Evans, J. T.,Integrated Ferroelectrics, I , 1 (1992). 47. Barlingay, C. K., and Dey, S. K.,Appl. Phys. Lett. 61, 1278 (1992). 48. Hirano, S., and Kato, K.,J. Non-Cryst. Solid., 100,538 (1988). 49. Miller, K. T., and Lange, F. F., J. Muter. Res., 5, 151 (1990). 50. Miller, K. T., and Lange, F. F., J. Muter. Res., 6, 2387 (1991). 51. Lines, M. E., Phys. Rev. B, 15, 388 (1976). 52. Glass, A. M., Lines, M. E., Nassau, K., and Shiever, J. W., Appl. Phys. Lett., 31, 249 (1977). 53. Engelmann, H., Kraemer, N., and Gonser, U., Ferroelectrics, 100, 127 (1989). 54. Varma, K. R. B., Harshavardhan, K. S., Rao, K. J., and Rao, C. N. J., Muter. Res. Bull., 20, 315 (1985). 55. Kitabatake, M., Mitsuyu, T., and Wasa, K., J. Appl. Phys., 56, 1780 (1984). 56. Kitabatake, M., Mitsuyu, T., and Wasa, K., J. Non-Cryst. Solid., 53, 1 (1982). 57. Xu, R., Xu, Y. H., and Mackenzie, J. D., in Sol-Gel Optics I1 (J. D. Mackenzie, ed.), SPIE Roc. 1758, San Diego, California, 1992. 58. Xu, Y. H., Cheng, C. H., Xu, R., and Mackenzie, J. D., Muter. Res. Symp. Proc., 271, (1992).

22 Ceramic Membrane Processing C. Guizard, A. Julbe, A. Larbot, and L. Cot Centre National de la Recherche Scientifique Montpellier, France

I.

,

INTRODUCTION

The use of inorganic membranes in separation technology is relatively new and has given rise to much interest in recent years. This is a result of the inherent properties of inorganic membrane materials, which are generally more stable chemically, structurally, and thermally than organic materials. Ceramic membranes represent a distinct class of inorganic membranes. Other classes consist of such membrane materials as glasses [1,2], carbon and metals [3,4], and organic-inorganic polymers [5]. This chapter focuses on the chemical processing of ceramic membranes, which has to date constituted the major part of inorganic membrane development. Before going further into the ceramic aspect, it is important to understand the requirements for ceramic membrane materials intermsofporous structure, chemical composition, and shape. In separation technologies based on permselective membranes, the difference in filtered species ranges from micrometer-sized particles to nanometer-sized species, such as molecular solutes or gas molecules. One can see that the connected porosity of the membrane must be adapted to the class of products to be separated. For this reason, ceramic membrane manufacture is concerned with macropores above 0.1 pm in diameter for microfiltration, mesopores ranging from 0.1 pm to 2 nm for ultrafiltration, and nanopores less than 2 nm in diameter for nanofiltration, pervaporation, or gas separation. Dense membranes are also of interest for gas 501

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separation. These membranes must work in liquid or gaseous media, usually under harsh conditions. For this reason the major type of ceramics used in ceramic membrane manufacturing consists of refractory oxides-alumina, zirconia, or titania. Nevertheless, manyother ceramic materials, including cordierite, mullite, silicon carbide, silicon nitride, and silica, and also borosilicate glasses, have been mentioned as suitable materials for inorganic membrane preparation [6]. Concerning the shape of ceramic membranes, two main geometries have largely been investigated: tubular and flat membranes. Recently, alumina hollow fibers were proposed by DuPont in the microfiltration range. Membranes designed for the separation and concentration of various species must work in a cross-flow mode in which the liquid to be filtered circulates across the surface of the membrane while the permeatedliquid passes through the membrane perpendicular to the feed flow direction. Provided that the interactions between the feed liquid and the membrane can be minimized, the cross-flow mode prevents the accumulation of the retained products at the membrane surface and slows fouling phenomena. Returning to the ceramic aspect of inorganic membranes, different structural levels are concerned with the expected selectivity of the different categories of a ceramic membrane. Macropores in ceramic membrane supports are derived from micrometer-sized particles processed through conventional shaping and sintering methods, leading to porous bulk materials. Submicrometer-sized pores in the microfiltration range are obtained through submicrometer particle sintering. At this stage the system can no longer be considered a bulk material, and coating processes must be involved in membrane preparations [7,8]. With respect to pore size, “1 m’’constitutes the borderline between bulk materials and thin-film shaped materials to obtain high flow rates through the membranes.Finally, ceramic nonfilters are included in the newtechnologyof nanophase ceramic development [9,10]. After rapidly reviewing the historical aspects of inorganic membranes and outlining the basic principles in the preparation of porous ceramic materials, conventional methods adapted to the preparation of porous ceramic supports are first developed. The second part considers the new ceramic processing techniques that can be used advantageously in ceramic membrane preparation. Processing of submicrometer powders, sol-gel technologies, and chemical vapor deposition are some of the many advanced technologies that can be applied to ceramic membrane processing. The most adapted characterization techniques for each class of porous ceramic membrane are also described.

II.

HISTORICAL BACKGROUND

Workon inorganic membranes started withVycor glass membranes, which were studied in the mid-1940s[l]. In fact, two periods preceded the current de-

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velopment of inorganic membranes. The first was related to the separation of uranium isotopes by the gaseous diffusion process applied to UF6. The challenge was to select membrane materials able to work in avery aggressive chemical medium. Moreover, the radioactive environment required very reliable membranes. After the oil crisis in 1973, ceramic oxide-based supports were proposed and produced by two companies, Ceraver (a subsidiary of the CGE group) and Euroceral (a joint venture between Lafarge and Norton), with the view to providing the Eurodif plant in France with 4,000,000 m2 membranes. The membranes were coated by SFEC (a subsidiary of CEA in France). The nature of the membranes still operating in Eurodif remains classified, but until today very reliable behavior has been observed. A few years later (1982), nuclear energy programs were considerably decreased and the project for a second diffusion plant was abandoned. This was the end of the first period, and the Euroceral, Ceraver, and SFEC production plants were shut down. The second period consisted of the development of microfiltration and ultrafiltration inorganic membranes as a consequence of the know-how accumulated by the companies that built the gaseous diffusion plants. The inorganic membrane concept for the filtration of liquids was first developed by Carre (a subsidiary of DuPont) in the 1960s, with a dynamic (nonpermanent) zirconium hydroxide membrane on a stainless steel support, and by Union Carbide in the 1970s, with ceramic oxide layers coated on carbon supports. The first commercial wholly cross-flow filtration system equipped with an inorganic membrane was manufactured by SFEC (now TECH-SEP) in 1978. The Carbosep membranes from SFEC were made up of zirconia porous layers coated on a macroporous carbon support. In the 1980s, Ceraver (now SCT) developed a range of alumina microfiltration membranes (Membralox), introducing the concept of multichannel support. This was the starting point for the important development scheduled for inorganic membranes in the 1990s. A comprehensive review of the current advance in inorganic membranes is given in arecent book by Bhave [l l].

The previous chapters in this book present current developments in ceramic material processing with emphasis on new technologies. Most of these concepts are of interest to the preparation of both ceramic porous supports and supported ceramic membranes, showing that these kinds of materials enter the class of high-technology ceramics. An additional point is that, compared with other ceramic materials, the porous structure is the main aspect. In high-technology ceramic processing, a porous material usually appears as an intermediate structural stage that must be further densified, unlike ceramic membranes,

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in which a tailored porous structure appears to be key characteristicof the final material. The connected porosity in ceramic supports and membranes results from a process in which no pressure assistance is provided during sintering. Focusing on the voids, not on the solid part of the ceramic, the size, volume, and shape of the pores become evident characteristics of the ceramic material. The basic idea is that the porosity is directly related to the size and arrangement of individual particles in the fired material. Consequently, each filtration field (micro-, ultra-, and nanofiltration) consists of a different class of porous ceramic membranes in which the particle size must be adjusted. As a general rule, an ideal packing of monosized quasi-spherical particles generates interparticle voids for whichthesize, shape, andporousvolume (porosity) depend on the chosen arrangement model. Figure l a shows two different ways of packing spherical particles. This generates a connected porosity. Figure l b shows the corresponding shapes of the pores. With some ideal packings, such as the hexagonal compact mode, straight channels (tortuosity of 1) can be observed. All deviation from these ideal packings broadens the pore size distribution. In practice, these ideal arrangements are not observed in ceramic membranes and particle packing can be assumed to be more or less randomly arranged with a tortuous porosity in the 3040% range. When particles deviate from the spherical shape, an oriented porous structure can be obtained. This is the case for alumina ulafiltration membranes obtained from boehmite plate crystals [12]. Another phenomenon affecting the porous characteristics of ceramic membranes is the particle size evolution during sintering as a function of temperature and time. A general tendency is forthe individual ceramic grains and pores to increase in size with increasing temperature and duration of the sintering process. An example is given in Sec. V.A. Figure 10 shows the evolution of pore size as a function of sintering temperature for an alumina membrane [13]. Ceramic membranes are not prepared with ideal particles, and so this phenomenon is exacerbated when the starting particles exhibit a broad size distribution, because smaller particles are swallowed up by larger particles during the sintering process. The flux performance of a ceramic membrane is directly related to the porosity, which must be as high as possible without sacrificing mechanical strength. Porosity can be improved by a partial agglomeration of the particles at the initial stage of the process, but pore size distribution is increased as well. This is a contradictory approach because a good flux performance needs high porosity but good separation selectivity requires a narrow pore size distribution. Thus, a good ceramic membrane is a balance between these two requirements. Considering porous supports, the starting materials are ceramic powders a few micrometer in size. With these, one can generate a bulk structure withpore diameters larger than 1 pm. Because of the low surface area of these powders, high temperatures (up to 1600°C) are required for sintering. This kind of ce-

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505

Figure 1 (a) Two different packing modes for spherical particles. (b) Corresponding pore shapes.

ramic material obeys the general rules of conventional ceramic processing. On the contrary, the design of supported membranes must be considered in a different way. Supported membranes differ from bulk supports by being thin-film shaped structures obtained through an appropriate coating process. Moreover, because the physical characteristics (including pore size, porosity, and surface area) define the application area for these membranes, different starting materials and processing techniques must be considered. Concerning microfiltration, in which pores of less than 1 pm are needed (typically 0.1-1 pm), suspensions of submicrometer powders must be processed using slip-casting or tape-casting techniques, depending on the support shape (flat or tubular). Smaller pore sizes are required for ultrafiltration. In this case pores result from the packing of colloidal particles, which cannot

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behandled as drypowders. Instead they are formed in aqueous media and maintained as stable suspensions throughout the process. For this purpose, new techniques like the sol-gel process are of vital importance. Membranes exhibiting pore diameters down to 3 nm (present limit for commercial membranes) can be created. New developments of this process are underway on the laboratory scale with the objective of nanoporous membrane production. The development of ceramic nanofilters is closely related to the very innovative work on nanophase material. Sections V.B and C are specially devoted to recent work in this field.

W.

PREPARATION AND CHARACTERISTICS OF CERAMIC SUPPORTS

Most inorganic membrane supports exhibit a tubular shape. This is awelladapted geometry for cross-flow filtration in which the feed stream is circulated across the surface of the membraneand the permeated flux passes through the membrane in a perpendicular direction. Stainless steel, carbon, and ceramic are the most frequently used materials in the preparation of supports. A s shown in Fig. 2, tubes or multichannel substrates can act as membrane supports. A well-designed support must be mechanically strong, and its resistance to fluid flow must be very low. Aiming at enhancing flux performances, multilayered substrates have been prepared that exhibit an asymmetric structure

permeate

multichannel porous ceramic Figure 2 Porous ceramicsdesigned for cross-flowfiltration: (a) tube-shaped ceramic; (b) multichannel ceramic.

CERAMIC MEMBRAIVE PROCESSING

507

with an increasing pore size gradient, generally from inside to outside the tube. An example of a ceramic support made of several alumina porous layers is given in Fig. 3. The bulk ceramic substrates are formed by the extrusion of a ceramic paste derived from a ceramic powder with a regulated grain size. This is a wellknown shaping method used in conventional ceramic processing. The shape of the extruded green body depends, as shown in Fig. 4, on the geometry of the die used. Tubular or multichannel supports can be produced in this way. To obtain an asymmetric support in a one-step process, two layers can be successfully coextruded in the manufacturing stage. The general formulation of ceramic pastes used in support manufacturing can be defined as a mixture of four basic compounds: a ceramic powder, an organic and/or an inorganic binder, a lubricant, and water. Because the ceramic porous structure depends on the shape of individual grains and the way they are packed, different factors can affect the two major characteristics of membrane supports, mechanical strength and porosity. The pseudoplastic behavior of the paste during extrusion is responsible for an exponential dependence of extrusion velocity versus applied pressure. High pressures to increase support permeability and strength have been emphasized in the literature [14]. These can be linked to a better (more dense) particle pack-

l

Figure 3 Cross-sectionalimage of anasymmetricceramic SCT (France).

porous substrate from

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c

c

tubular

support

Figure 4 Two different dies used for ceramic support extrusion. ing at high extrusion pressures and velocities as a result of high shear forces that are able to break up agglomerates [15]. High extrusion velocities can be achieved provided that the rheological properties of the paste have been carefully adjusted during the formulating step by adding the proper quantities of water and organic additives. Thermal treatment of the extruded green bodies can be carried out in two stages: (1) the drying step, with elimination of water at temperatures lower than lOo"C, under a controlled atmosphere, and (2) the firing step, with decomposition of organic additives and sintering of the ceramic up to 1200-1600°C. Concerning plate-shaped ceramics, tape casting (or thedoctor blade process) is an important forming operation commonly used to prepare multilayer capacitors and packages in the electronic industry [16]. This technique was recently investigated for the preparation of flat ceramic filtration supports [17]. Tape casting is a low-cost process for the manufacture of large-area thin ceramic sheets of controlled thickness andhighquality.Mosttape-casting processes start with a milling procedure in which ceramic powders are mixed with a solvent and a dispersant in a ball mill. A mixing and homogenization step then occurs in which binders and plasticizers are added to the low-viscosityslurry. The homogeneous, well-dispersed,and concentrated slurry is degassed and spread on a flat movingcarrier surface using the doctor blade process. Figure 5 is a schematic representation of the doctor blade apparatus. Thickness control isa function of several parameters, including viscosity (which canbe adjusted by solvent content or temperature), castingcarrier

CERAMIC MEMBRANE PROCESSING Micrometer screws

0

cm

509 5

L ” J

,Substrate

film

-Aluminium

Heating elements

Figure 5 Doctor bladeapparatus. (From Ref. 17)

speed, doctor blade gap setting, and reservoir depth behind the doctor blade [17]. After drying, the film is removed from the support andthenusually stamped to shape and further stacked into laminated configurations. Finally, burnout of organic polymers and sintering are performed under controlled temperatures. Three characterization methods are usually applied to ceramic supports: 1. Related to the ceramic structural properties, the porous texture can be imaged by scanning electron microscopy. 2. Mechanical strength can be measured using pressure tests. Highly pressure resistant supports (up to 100 bar) with a multichannel geometry are now commercially available. The integrity of the support is checked by controlling air (or N2) bubble emission through the porous structure of the support immersed in a liquid, such as water or alcohol. The gas pressure at which bubbles are detected can be related to the size of cracks (or capillary-shaped pores) using a calibration method according to the Laplace equation, d*Ap=4y

COS

0

(1)

where d is the mean diameter of the crack (or pore), Ap is the overpressure with respect to the atmospheric pressure, y is the surface tension at the airniquid interface, and 0 is the wetting angle of the liquid on the ceramic material. An air bubble penetrates the support through a void when the bubble radius is equal to (or smaller) than that of the void; this means that the contact angle is zero. This bubble-point method has become a standard technique used by suppliers to measure the largest active pores within

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membranes. Table 1gives an indication ofthe pressure required for a given pore radius (using water as the wetting medium). If small pores are present, it is necessary to apply such high pressures because the surface tension of water/air is relatively high. Water must often be replaced by alcohol (i-propanol is often used as a standard liquid). This method has been approved as an American Society for Testing and Materials procedure (F316). 3. Finally, liquid permeability can bemeasured to evaluate the flux performance of the support, the aim being to obtain a porous material with a very low resistance to the filtered liquid. Once the supports exhibit suitable properties, they can be used as the base porous material in supported membrane processing. The successive layers acting as active membranes can be deposited on either the inner or the outer surface of tubes or supports. Starting from the same support, microfiltration, ultrafiltration, and nanofiltration performance is achieved by superimposing successive layers, yielding an asymmetric structure with porosity percolating from one layer to another. The main condition for a successful process is that the firing temperature of the final layer must be sensibly lower than the firing temperature of the supporting layer.

V.

SUPPORTEDCERAMIC MEMBRANES:SYNTHESIS AND CHARACTERIZATION

The effectiveness of a membrane depends both on its ability to separate particles or molecules in a selective way and on the flux that can be achieved across this membrane. Although some membranes, such as glass membranes, exhibit a symmetrical structure, in most cases they are asymmetric, consisting of several layers coated on a bulk porous support and with a gradual decrease in pore size (Fig. 3). The main advantages of an asymmetric structure are high fluxes and the ability to have tailored membranes made of materials different from the

Table 1 RelationBetweenPressureand Pore Radius Using Wateras the Wetting Medium in the Bubble-Point Method (pm)

radius Pore 10 1.o

0.1 0.01

Pressure (MPa) 0.014

0.145 1.45

PROCESSING MEMBRANE CERAMIC

51I

support material. Ceramic supports are not the only category of material on which ceramic membranes can be deposited. Porous carbon or metal supports can also be used for this purpose. Depending on the mean pore size to be obtained for the membrane, an appropriate preparation method must be used.

A.

Ceramic Powder Method for Microfiltration Membranes

Microfiltration layers with pore diameters larger than 1 pm can be made by coextrusion with the support in a one-step process. For this purpose a special die geometry is needed in which two concentric shells of paste (one for the support and the other for the microfiltration layer) are simultaneously extruded. Ceramic membranes with pore diameters in the 0.1-1 pm range can be prepared through aslip-casting method using asubmicrometer powder as the starting material. Membrane processing consists of preparing a slurry of the powder in water andthen slip casting the suspension onto the surface of the support. Organic additives like binders or plasticizers are added to the slurry to adjust the rheological behavior of the coated layer. An example of a zirconia microfiltration membrane coated on a porous stainless steel support is shown in Fig. 6. Two different microfiltration mem-

Figure 6 Cross-sectional image of a zirconia microfiltration membrane coated on a stainless steel macroporous support.

a

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branes have been coated on this metal support using submicrometer zirconia powders exhibiting surface areas of 20 and 40 m2/g. These powders were dispersed in water, with polyvinyl alcohol and ethyleneglycol chosen as a binder and a plasticizer,respectively, to control the rheological properties of the slurry. When slip cast on porous stainless steel tubes, the suction force created on leaving this slurry in contact with the support macropores (5 pm in diameter) forced the dispersed particles to concentrate at the macropore entrance and form a layer of packed particles. Once this layer was dried at a controlled atmosphere and temperature to prevent crack formation, the sintering stage was performed at 1OOO"C under hydrogen atmosphere to preserve the metal support from oxidation. Two membranes with mean pore diameters of 0.1 and 0.2 pm were obtained, depending on the surface area of the original powder (Fig. 7).

B.

Sol-GelApproachfor Ultrafiltration Membranes

When, to satisfy ultr&iltration, nanofiltration, or gas separation requirements, the required pore size is under 0.1 pm, the ceramic powder approach is no longer viable. Indeed, individual particles yielding pore diameters smaller than 0.1 pm cannot be handled by powder processing. In fact, particles of this type enter the category of colloids and must be maintained as a stable suspension during the process. A s indicated in the introduction, the sol-gel method is a very suitable way to produce mesoporous and nanoporous membranes. The lat&r is elaborated in the next section.

diameter (urn)

l'"

m

diameter(pm)

Figure 7 Pore size distribution of zirconiamicrofiltrationmembranesmeasuredby mercury porosimetry.

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There are two main routes for the sol-gel process, which are described in Fig. 8. The route on the left side is known as the colloidal(or particulate) route. It involves the reaction of a metal salt or hydrated oxide with excess water. This yields a precipitate of gelatinous hydroxide colloidal particles, which can be redispersed through a peptization reaction using an electrolyte. The particle sizes in the resulting sol are typically a few tens of nanometers and are suitable to create mesopores down to 2-3 nm in diameter. The basic principles of soldestabilizationwiththeaim of producing these mesoporousmembranes were described in one of our previous papers [S] and widely exploited by a number of authors [18-201. The porous structure and pore size distribution are governed by particle aggregation at the sol stage. The peptization phenomenon used to prepare a colloidal sol is based on electrostatic interactions according

I

controlled hydrolysis

peptization

W (colloidal species)

4-

WDERS

-..lc

\

I

I

Sol (polymeric species)

/

1

Sol-Gel transition

GEL LAYER

1

Drying

QRGANIC-MORGANIC MEMBRANE

1

Firing

PURE INORGANIC MEMBRANE

Figure 8 The two sol-gel routesusedto produce inorganicmembranes.

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to DLVO theory [21,22]. The repulsion forces that prevent particle aggregation in the sol are a result of the electrical double layer caused by the amphoteric behavior of most of the oxide surfaces. Electrophoresis allows us to determine the particle stability range through particle mobility measurement as a function of pH and electrolyte concentration. When the pH value is decreased or increased away from the isoelectric point (EP), a maximum stability for the sol is observed with high repulsive forces between particles. Near the IEP a flocculation phenomenon occurs, resulting in gel or precipitate formation. A s shown in Fig. 9, a change in particle mobility versus pH results in an evolution of particle stacking, which influences pore size and the structure of the membrane. Based on the control of these phenomena, tailored pore structures have been produced for titania and zirconia membranes [23]. This concept was also applied to alumina membranes, which were the first to be described in the literature. The pore sizes were in the 2.5-6 nm range. Drying and firing steps can also affect the porous structure in different ways. The influence of drying conditions was revealed to be of prime importance to silica membranes [24]. Concerning the firing step, Fig. 10 illustrates the influence of temperature on pore size evolution for an alumina membrane. Recent attempts in the literature [20] to correlate the characteristics of bulk xerogels with those of coated membranes are questionable, mainly because the

+ Figure 9 Influence of particlechargeandmobilityin a colloidal sol on theporosity of the resulting material. V = potential energy, H = interparticle distance.

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IL ’

1200°C 55 nm I

l I I

10

I l I l

’

l

l

A

E

l #

*

v

a

500

1000

Figure 10 Influence of the tiring temperature on the pore size evolution

of an alu-

mina membrane.

environment of the drying and firing steps of a coated layer differs markedly from that of a bulk material. Because of the influence of the capillary forces exerted by the support, the coated sol coalesces in a very dramatic way, yielding a gel that cannot be compared to the corresponding bulk gel in terms of pore structure and porous volume. Moreover, during the drying step capillary forces that result from solvent evaporation are oriented in a direction perpendicular to the surface of the film, resulting in an important diminution of the layer thickness.Normally these capillary forces are responsible for the destruction of a bulk xerogel. Basic mechanisms involved in sol-gel derived coating formation (during aging, drying, and firing treatments) have been extensively developed by Brinker and Sherrer [25].

C.

Ceramic Nanofilters by the Sol-Gel Process

Recentlymuch attention has beenpaid to ceramic membranes exhibiting a nanoporous structure with the aim of new membrane processes for the nanofiltration of liquids [26], pervaporation [27], gas separation [27,28], or catalysis

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[29]. These membranes are achievable using the concept of nanophase ceramics. According to literature, this new class of materials can result from the emphasis of some new ceramic processes, such as the condensation of gaseous atomic clusters [30] or the sol-gel process [31]. This last method, which has beensuccessfully applied toultrafiltrationmembranes,wasusedrecently to prepare ceramic nanofilters. Nanophase materials deal both with the nanometer-sized particle and with the nanometer pore size aspects. The nanopore aspect is central to membrane technologies because of the need for selective separation processes at the molecular level. To understand the importance of nanostructures in microsieving membranes, the basic structure of nanophased ceramics must be briefly described. Because the particles are extremely small, one to a few tens of nanometers, an important fraction of the atoms is found in or very near the interface between grains, as reported in Table 2 [32]. Figure 11 isa schematic representation of a nanophase material. One can see that individual grains in the 5 nm range induce a biphasic material with an interfacial phase between the grains and a residual nanoporosity, evidenced by positron lifetime spectroscopy [33]. Transmissionelectronmicroscopy is also a well-adapted technique for nanoscale structure characterization, as illustrated later. Concerning membranes, new separation capabilities are expected for these materials. The molecular sieving effect caused by connected nanopores can be applied tothe separation of molecules with molecular weights smaller than 1000. The key properties of such membranes are based on the preponderant effect of activated diffusion in nanopores, however. This phase transport phenomenon derives from the nanophased ceramic concept and classes these membranes among those materials expected to be crucial in the areas of modern technology, such as environmental protection, biotechnology, and the production of effect chemical. Recent advances in sol-gel science allowed us to go further in controlling the individual particle size at the sol stage, resulting in individual nanoscale grains in the ceramic. This new development is based on the chemical modification of both metal salt and metal alkoxide precursors to modify their reactivity and to achieve a nanopore structure for the membranes [34]. Some ex-

Table 2

Atom Repartition in aMaterial as a Function

of the Individual Grain Size

Meangraindiameters(nm) 5 10 100

% Atomsinthegrainboundary

30-60 15-30 1-3

CERAMIC MEMBRANE PROCESSING

Figure 11 A nanophasematerialaccordingto

517

Ref. 48.

amples are given to show the interest of this method in ceramic nanofilter preparation. 1. Zirconia Nanofiltration Membranes Prepared by the Colloidal Route. In this case zirconia sols obtained from zirconium oxychloride have been used for the preparation of inorganic membranes active in nanofiltration. Previously described conditions [31] to achieve high selectivity and flux during nanofiltration can be expressed in terms of the structural requirements of the membrane material. Here the two main parameters are the phase stability of zirconia and crystallite size. If zirconia exhibits suitable basic properties .as a membrane material, different polymorphs .(tetragonal, monoclinic, and cubic) can be encountered starting from the amorphous state obtained at room temperature. A neutron diffraction study performed on this material showed the evolution of crystallization kinetics and crystallite growth along the transformation sequences under different firing atmospheres (air, nitrogen, and a hydrogen-nitrogen mixture). Several phenomena have been evidenced that can affect the basic properties of the membrane. Tetragonal crystallites of less thm 6 nm, obtained from the amorphous state at 360"C, are very suitable to generate nanopores of less than 2 nm after sintering at 500°C. When raising the firing temperature to obtain a microporous sintered layer, the tetragonaVmonoclinic phase transition occurs, changing the crystallite size abruptly from 6 nm to more than 12 nm. Depending on the firing atmosphere, this transition temperature can be shifted from 550°C under air to 600°C under nitrogen or hydrogen and nitrogen atmosphere. In each case individual tetragonal particle size can be maintained under 6 nm provided that the firing temperature does not exceed the transition temperature, as shown in Fig. 12.

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30 300

I

I

I

I

400

500

600

700

TEMPERATURE ("C)

Figure 12 Influence of the f ~ n atmosphere g on the evolution of zirconia crystallite size versus temperature for the tetragonal phase.

A transmission electron microscope (TEM) observation of the ZrO2 nanocrystallites obtained at 500°C is given in Fig.13. Homogenous, quite spherical, consolidated particles of about 5-45 nm in size can be observed. This clearly shows that, to maintain nanoporosity inside the membrane, an upper limit for membrane sintering exists. Because these results, a zirconia nanofilter has been obtained by coating a 1 pm thick layer on a microfiltration zirconia layer. The separation performance of this membrane, characterized with model solutes in aqueous media, is in the nanofiltration range [26].

2. Lanthanum Oxychloride Catalytic Membranes These membranes have been prepared from homogeneous sols based on lanthanum chloride aqueous solutions. Two routes have been compared for the sol preparation. In the first, ammonia was added to a stirred aqueous solution of lanthanum chloride; a complex mixture of lanthanum hydroxide and/or basic salts intermediate species was obtained in the form of an opalescent sol. In the second route, an organic complexant (acetic acid) was used to create soluble species around the working pH (pH = 8). Ammonia addition to the modified solution led to clear sols containing probably smaller particles than those obtained by the first preparation method. Sol formation required a binder and a plasticizer to achieve the desired behavior of the cast films during the drying and firing treatments. Here polyvinylic alcohol fulfilled these functions and led to crack-free ceramic layers. Typical processing consisted of a drying step at 110°C to avoid the crystallization of stable carbonate species and then heat

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Figure 13 TEM image of zirconia nanoparticles forming a nanofiltration membrane obtained from a zirconium oxychloride precursor.

treatment at 800°C for 5 h in an N2 atmosphere to obtain a lanthanum oxychloride porous thin film. LaOCl crystallizes at about 400°C in the tetragonal system and is still stable at 800°C. The first LaOCl layer (prepared without acetic acid) was used as a support for the deposition of a second finer membrane from a modified sol. Figure 14 is a scanning electron microscope observation of the cross section of a supported double layer. Very small particles (around 6 m in size) have been evidenced by E M ; N2 adsorption and desorption measurements have revealed a pore diameter distribution with a maximum at around 1.5 m. The catalytic performance of such a material for oxidative coupling of methane has been described elsewhere [35] as a function of preparation conditions. 3. Silica Membranes from Polymeric Sols The preparation of submicrometer monodisperse particles by the controlled hydrolysis of metal alkoxides has been widely investigated for a number of ceramic materials. In particular, sol-gel processes involving the formation of particulate materials or glass-precursor gels from silicon alkoxides, among them tetraethylorthosilicate (TEOS) and tetramethoxysilane, have received considerable attention in the literature. The potential of TEOS sols to form crack-free nanoporous thin films suitable for use as membranes has been investigated and is used as an example in this section. Starting from TEOS in ethanolic solution, the alkoxide hydrolysis was performed with water under acid conditions. Indeed, compared withbase-catalyzed hydrolysis reaction,acid catalysis is known to lead to smaller polymers with a low degree of cross-linking [36] and

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Figure 14 Cross-sectionalimage of a double-layer lanthanum oxychloride membrane.

is thus a most interesting starting material for the preparation of microporous materials. An organic binder (e.g., polyvinylic alcohol or polyoxyethylene) was added to the sol before coating to prevent crack formation during drying and firng treatments. After elimination of the organic products at 450°C, a supported crack-free amorphous silica thin film (0.2 pm thick) was obtained. This membrane was deposited on a mesoporous silica membrane, described in Ref 37. A TEM observation of the top layer is presented in Fig. 15 and shows nonspherical grains with a length around 6 nm and 2 nm wide. The corresponding NZ adsorption isotherms carried out on this top layer are characteristic of microporous materials. The Dubinin-Astakhov pore volume plot in Fig. 16 shows a pore radius distribution in the nanometer range. The specific surface area measured for this membrane is about 700 m2/g. 4. Titania and Zirconia Membranes Prepared by the Polymeric Route Titanium and zirconium propoxides can be used as precursors for the preparation of nanoporous titania and zirconia membranes. To avoid the precipitation of inhomogeneous hydroxide particles during the hydrolysis step, the alkoxide reactivity can be modified with acetylacetone (acacH). This chelating agent reacts readily with transition metal alkoxides, as follows [34]: M(OR), + acacH + M(OR),(acac) + ROH with M = Ti or Zr

CERAMIC MEMBRANE PROCESSING

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Figure 15 TEM image of silica nanoparticles in a TEOS-derived membrane.

This ligand then acts as a functionality locker when substoichiometric hydrolysis ratios are used. A ratio acacH/M greater than 1 prevents precipitation and leads to stable colloids or gels. Consequently, with a good formulation choice, sols can be prepared in air without precipitation. Either titania or zirconia ultrafiltration layers have been used as supports for these membranes. Figure 17 is a cross-sectional image of a zirconia microporous membrane that is 0.2 pn thick. The supported layers obtained after treating at 500°C exhibited crystallized structures (anatase for titania and a tetragonal metastable form for zirconia) and revealed a very fine texture from TEM observation. For Ti with a Ti/acac ratio = 1, grains with defined faces can be observed whose size is around 20 nm. For Zr with a Zr/acac ratio = 2, a finer texture was obtained than with titania, with a mean grain size of around 4 nm. In the two cases, powder x-ray diffraction (Sherrer formula) was used to measure an individual crystal size, which is in good accordance with TEM observations. For the zirconia membrane the N2 adsorption and desorption isotherm in Fig. 18a shows a microporous domain with a 1.7 nm centered pore size distribution (Fig. 18b).

D. Characterization of Mesopores and Micropores in Ceramic Membranes Many characterization techniques developed for the characterization of mesoporous and microporous materials have been adapted to membrane characterization (e.g., mercury porosimetry, adsorption and desorption isotherms, and thermoporometry). These techniques are related to morphological parameters

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0.028

2 0.024

0 &I

0.008 0.004

0.000 5

6

7

8

9

10

11

12

13

EQUIVALENT PORE RADIUS

14

15

(A)

Figure 16 Dubinin-Astakhov differential pore volume plot obtained from N2 adsorption-desorption measurements on a microporous silica membrane.

[38] and can be defined as static characterization techniques in which the sample has no active role. A series of specific techniques for the characterization of active pores in membranes (e.g., permeability of gases and liquids) have been also developed that involve the entire filtration element, not only samples. These dynamic techniques, in which the membrane plays an active role, are usually nondestructive and yield permeation-related parameters. A very precise characterization of the porous texture of materials is often difficult because of interactions between the tested material and the method but also because this approach requires the use of a model for the pore system (idealized pore shape assumption). Basic principles of the main characterization techniques of membranes are described here. 1. Mercury Porosimetry Mercury porosimetry is a well-adapted method to characterize the pore size and pore size distributions of nlicrofiltration ceramic membranes. In this method,

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Figure 17 Cross-sectional image of a zirconia microporous membrane obtained from a modified alkoxide precursor and deposited on a zirconia ultrafiltration layer.

which is a variation of the bubble-point method, mercury is forced into the pores of a dry sample. The quantity of mercury forced into the pores is determined very accurately at each pressure; it can be measured by means of the capacity variation induced by the reduction in a Hg column height connected to the measuring cell. The relationship between the applied pressure P and the pore rp is given by the modified Laplace equation (Washburn equation),

Because mercury does not wet the membrane, the contact angle 8 is greater than 90" and cose has a negative value. A widely accepted value for oxides is eHg/oxide = 140" and mg/air = 0.48 N/m. Very high pressures are needed for pores in the nanometer range. A pore diameter of 4 nm corresponds to a pressure of about 400 MPa, which may damage the ceramic layer. 2. Gas Adsorption and Desorption Technique [39] This technique can be considered a standard method in the science of porous ceramics and catalysts. It is based on the principle that inside a small pore a gas can condense to a liquid at a relative pressure lower than unity; this introduces the capillary condensation theory. The adsorption and desorption isotherms of an inert gas are determined as a function of the relative pressure @re1 = p/po, i.e., the ratio between the applied pressure and the saturation pressure). N2 is often used as adsorption gas, and the experiments are carried out at the boiling liquid nitrogen temperature (at 1 bar). The adsorption isotherm

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is determined by measuring the quantity of N2 adsorbed for each value of PR]. Adsorption starts at a low relative pressure. Ata certain minimum pressure, the smallest pores are filled with liquid nitrogen. As the pressure is increased further, 1arger.pores are filled, and near the saturation pressure, all the pores are filled. The total pore volume is determined by the quantity of gas adsorbed near the saturation pressure. Desorption occurs when the pressure is decreased from the saturation pressure.Manymesoporous systems exibit distinct adsorption-desorption behaviors, which lead to a characteristic hysteresis loop. The curve shape is linked to different geometrical factorsthat rulethe adsorption and desorption processes. The reason for this hysteresis is that capillary condensation occurs differently in adsorption and desorption. Because of the concave meniscus of the liquid in the pores, N2 evaporates at a lower relative pressure because the vapor pressure of the liquid is reduced. The lowering of the vapor pressure for a capillary of radius rk is given by the Kelvin relationship

the contact angle 8 being assumed to be zero. The pore radius rp may be calculated from rp = rk + t, where t is the thickness of the adsorbed layer of vapor in the pores (estimated from calibration curves). This method is generally not very accurate in layers with a large pore size distribution and without a definite pore geometry. The Kelvin equation is useful to explain the hysteresis phenomenonand to calculatethe pore volume and pore area distributionsas a function of pore diameters (e.g., BJH method). Nevertheless, this Kelvin equation is based on thermodynamic considerations and can only be applied to a pore radius between 15 and 300A. In microporous systems, various methods have been developed for the determination of microporous volume and pore size distribution, but the validity of pore sizes is often ambiguous.

3. Thermoporometry This method is based on the microcalorimetric analysis of solid-liquid transformations in porous materials [40]. Since the system of water-filled pores has the closest resemblance to the practical situation of membrane filtration, the solid-liquid transition of water is often used for membrane pore size analysis. Because of the strong curvature of the solid-liquidinterface present within small pores, a freezing (or melting) point depression of the water (or ice) occurs. A full thermodynamic description of this phenomenon is given in Ref. 40. According to this concept, the size of a confined ice crystal (which is set by the size of the pore) is inversely proportional to the degree of undercooling, whereas the pore volume is directly related to the apparent transition energy.

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A differential scanning calorimeter is used to obtain a solidification (or melting) thermogram, from whicha pore size distribution can be extracted. By comparing the solidification and melting processes, thermoporometry can also be used to determine a thermodynamic pore shape factor, which varies generally from 1 (spherical pores) to 2 (cylindrical pores). 4. Permeabilityand Retention Measurements Typically, liquid permeabilities are obtained with water as permeate and expressed in terms of Lhm2-bar. Gas permeabilities are often expressed in terms of air or nitrogen permeability. For gases, membrane pore size also affects the transport mechanisms through the pores [27]. Manufacturers tend to characterize membranes by means of rejection measurements with reference molecules, such as dextrans, proteins, or polyglycols. A parameter extensively used for membranes characterization is the cutof value, which is defined as the lower limit of solute molecular weight for which the rejection is at least 90%. We must keep in mind, however, that rejection measurements always depend on the type of solute (shape and flexibility of the macromolecular solute), the membrane (its interaction with the solute), and the process parameters used (pressure, cross-flowvelocity,geometryof the test cell, and concentration and type of solute). In particular, concentration or polarization, pore blocking, and fouling phenomena affect rejection measurements very significantly. Some other techniques involving membrane permeability have been developed that have not yet been extensively used. Concerning gas permeability, a method called permeametry has been developed [41], based on Adzumi equation. It consists of measuring the variation in membrane permeability when favoring either molecular (Knudsen) or laminar (Poiseuille) flow regime. A mean pore radius is obtained with this method. Another more recent technique, called permporometry (or liquid displacement porosimetry), is based on the controlled blocking of pores by condensation of a vapor, present as a component of a gas mixture, and the simultaneous measurement of the gas flux through the membrane [42,43]. By measuring the gas transport through the membrane upon decreasing relative vapor pressure, the size distribution of the active pores can be found in the limit of validity of the Kelvin equation. The calculation of the number of pores can be performed by assuming a Knudsen type of transport regime. A liquid-liquid displacement technique (biliquid permporometry) was also found to be well adapted to the characterization of meso- and macroporous membranes [M]. It is based on the combined principles of bubble pressure and solvent permeability. In this case the applied pressure P and the flux J through the membrane are measured simultaneously. The recorded P and J values, introduced in the Laplace equation directly, give the pore equivalent radius and

PROCESSING MEMBRANE CERAMIC

52 7

the distribution of permeability versus pore radius. Pore number and pore area distributions versus pore size can be obtained by using Hagen-Poiseuille (or Kozeny-Carman) relationships. 5. Techniques Under Development A number of techniques were developed recently to solve some of the several problems linked to classic technique principles (low sensitivity, necessary pore shape assumptions, lack of validity for pores smaller than 1.5 nm, and others). A nuclear magnetic resonance spin-lattice relaxation technique was recently successfully demonstrated on a number of types of porous media. The basic principle is that the portion of pore fluid near a pore wall undergoes spin-lattice relaxation in a magnetic field faster than pore fluid removed from the pore wall [45]. Small-angle scattering techniques (small-angle x-ray and neutron scattering) have recently become established as versatile probes for the determination of structures of porous solids [46,47]. The size, shape, concentration, and surface area of inhomogeneities, such as pores, can be studied. Positron lifetime spectroscopy has been shown to be a good means of investigating the structural levels of nanocrystalline materials [48]. Different annihilation sites (dislocations, micropores, and mesopores) have been attributed to the different measured positron lifetimes.

VI.

NEW TECHNIQUESUNDERDEVELOPMENT IN CERAMIC MEMBRANE PROCESSING

Separation processes using membranes arise as a key technology in many industries (e.g., chemical, biotechnological, and water and wastewater treatment). For achieving the goal of more efficient processes, higher performance and more cost effective membrane materials are needed. Because of their excellent inherent properties, a rapid development of inorganic membranes is expected over the next 10 years, andnewthin-film deposition technologies are in progress with the aim of producing these new inorganic membranes. Electrophoresis was recently described as a promising method to produce inorganic membranes [49]. An electrical field is applied to deposit charged particles of dispersed mineral solids in water. Negatively or positively charged resins are used to coat the surface of the particles so they are transported to the electrode surface. Electrophoresis is therefore limited to conductive substrates that can serve as an electrode, but complicated shapes are easily accommodated. As mentioned earlier, new developments in gas separation, catalysis, and petroleum industries have stimulated interest in microporous inorganic membranes suitable for high-temperature applications and with the property of high

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permselectivity. Physical vapor deposition methods that have been extensively developed for the electronic industries can be transposed to membrane preparationwith the objective of yielding anew generation ofmetal or ceramic membranes [50]. Several methods are of interest to produce these thin inorganic films active in separation processes. In fhennal evaporation,solid materials are first vaporized by heating at sufficiently high temperature in vacuum, and then a thin film is formed on a cooler substrate by condensation of the vapor. This technique can be used for the deposition of thin layers made of such metals as Ag or Pt that are active in gas separation. Cathodic sputtering can also be avantageously used to deposit thin films of Ag and Pd on Ag supports. In this case materials are ejected atomistically from a target material by bombarding high-energy positive ions, such as Ar ions, and then condensed on a substrate to form a thin solid film. In chemical vapor deposition (CVD), chemically reacting gases or gaseous precursors are vaporized and decomposed under temperature to produce nonvolatile reaction products that deposit on the substrate. The low-temperature plasma-enhanced chemical vapor deposition (PECVD) process is also of interest to deposit inorganic thin layers on substrates that cannot be heated at high temperatures. Instead of temperature, plasma or microwaves are used to perform decompositionof the volatile precursors undervacuum. For example, nonporous silica glass, which is known for its excellent selectivity for hydrogen, is a good candidate for gas separation (28). Deposition of this kind of thin solid films using PECVD of Si02 can be envisaged on various porous supports such as ceramics, glasses, or polymers. These new developments show that ceramic membrane processing is fully involved in recent progress in ceramic materials and can be considered a hightechnology ceramic manufacturing industry.

REFERENCES 1. Nordberg, M. E., Properties of some Vycor-Brand glasses, J. Am. Ceram. Soc., 27, 299 (1944). 2. Schnabel, R.,and Vaulout, W., High pressure techniques with porous glass membranes, Desalination, 24, 249 (1978). 3. Bauer, J. M., Elyassini, J., Moncorge, G., Nodari, T., and Totino, E., New developments applicationsof carbon membranes, Key Eng. Mater.. 61/62,207 (1991). 4. McCaffrey, R. R.,McAtee, R. E., Grey, A. E., Allen, C. A., Cumming, D. G., Appelhans, A. D., Wright, R. B.,andJolley, J. G.,Inorganicmembrane technologies. Separ. Sei. Technol., 22, (2/3), 873 (1987). 5. Guizard, C., Larbot, A., and Cot, L., A new generation of membranes based on organic-inorganic polymers in TCIM '89 Proceedings, Montpellier, France, 1989, pp. 55-64.

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6. Chan, K. K., and Browstein, A. M., Ceramic membranes: growth, prospect and opportunities, Ceram. Bull., 70 (4), 703r (1991). 7. Charpin, J., Bergez, P., Valin, F., Barnier, H., Maurel, A., and Martinet, J. M., Inorganicmembranes:Preparation,characterization,specificapplications. High Tech Ceramics (P. Vincenzini, ed.), Elsevier, Amsterdam, 1987, p. 2211. 8. Larbot, A., Fabre, J. P., Guizard, C., and Cot, L., Inorganic membranes obtained by sol-gel techniques, J. Memb. Sci., 39, 203 (1988). 9. Burggraaf, A. J., Keizer, K., and Van Hassel, B. A., Ceramic nanostructure materials, membranes and composite layers, Solid Stare Ionics, 32/33, 771 (1989). 10. Guizard, C., Julbe, A., Larbot, A., and Cot, L., Nanostructures in sol-gel derived materials. Application to the elaboration of nanofiltration membranes, Key Eng. Mater., 61/62, 47 (1991). 11. Bhave, R. R., (ed.), Inorganic membranes. Synthesis, Characteristics and Applications, Van Nostrand-Reinhold, New York (1991). 12. Leenaars, A. F. M., and Burggraaf, A. J., The preparation and characterization of

alumina membranes with ultra-fine pores. 2. The formation of supported membranes, J. Coll. Interfac. Si., 105(1), 27 (1985). 13. Larbot, A., Alary, J. A., Guizard, C., and Cot, L., New inorganic ultrafiltration membranes: Preparation and characterisation,Int. J. High Technol. Ceram, 3, 143 (1987). 14. Terpstra, R. A., Bonekamp, B. C., and Veringa,

H. J., Preparation, characterizamition and some properties of tubular alpha alumina ceramic membranes for crofiltration and as a support for ultrafiltration and gas separation membranes, De-

salination, 70, 395 (1988). 15. Alfotd, N. M., Birchall, J. D., and Kendall, K., High-strength ceramics through colloidal control to remove defects, Nature, 330, 51 (1987). 16. Mistler, R. E., Tape casting: The basic process for meeting the needs of the electronic industry, Ceram. Bull., 69(6), 1022 (1990). 17. Simon, C., Bredesen, R., Raeder, H., Seiersten, M., Julbe, A., Monteil, C., Laaziz, Key Eng. I., Etienne, J., and Cot, L. Tape casting of flat ceramic membranes, Mater. 61/62, 65, (1991). 18. Anderson,M.A.,Gieselmann,M.J.,andXu, Q., Titania and alumina ceramic membranes, J. Memb. Sci., 39, 243 (1988). 19. Zaharescu M., Pirlog, C., Vasilescu, A., Crisan, D., and Braileanu, A., Inorganic membranes by the sol-gel method, Rev. Roum. Chim., 35(7-g), 909 (1990). 20. Xu, Q., andAnderson,M.A.,Synthesisofporositycontrolledceramicmembranes, J. Mat. Res., 6(5), 1073 (1991). 21. Parks, G. A., The isoelectric points of solid oxides, solid hydroxides, and aqueous hydroxo complex systems, Chem. Rev., 65(2), 177 (1965). 22. Overbeek, J. T. G., The rule of Schulze and Hardy, Pure Appl. Chem., 52, 1151 (1980). 23. Larbot, A., Fabre, J. P., Guizard, C., Cot, L., and Gillot, J., New inorganic ultrafiltration membranes: Titania and zirconia membranes;J. Am. Ceram. Soc., 72(2), 257 (1989). 24. Larbot, A., Julbe, A., Guizard, C., and Cot, L., Silica membranes by the sol-gel process, J. Memb. Sci., 44, 289 (1989).

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25. Brinker, C. J., and Scherer, G. W., Sol-Gel Science, Academic Press, New York, 1990. 26. Guizard, C., Ajaka, N., Garcia, F., Larbot, A., and Cot, L., New membranes for

the hyperfiltration of small molecules. Influence of the mesoporous structure on Proc. of V FiltrationCongress, separationandfractionationperformances,in Nice, France, 1990, pp. 143-145. 27. Ulhorn, R. J. R., and Burggraaf, A. J., Gas separation with inorganic membranes, in Inorganic Membranes (R. R. Bhave, ed.), Van Nostrand-Reinhold, New York, 1991, p. 155. 28. Klein, L. C., and Giszpenc, N., Sol-gel processing for gas separation membranes, Ceram. Bull., 69(11), 1821(1990). 29. Armor, J. N., Catalysis with permselective membranes, Appl. Catal., 49, 1(1989). 30. Siegel, R. W., Nanophase ultrafine grained materials, Mater. Sci. Forum, 37, 299 (1989). 31. Guizard, C., Julbe, A., Larbot, A., and Cot, L., Nanostructures in sol-gel derived

materials. Application to the elaboration of nanofiltration membranes; Key Eng.

Mater., 61/62, 47 (1991).

MRS Bull., 32. Siegel, R. W. Nanophase materials assembled from atomic clusters, 15(10), 60 (1990). t H., Structureof 33. Schaefer, H. E., Wiiurschum,R.,Birringer,R.,andGleiter,

nanometer-sizedpolycristallineironinvestigatedbypositronlifetimespectroscopy, Phys. Rev., B38(14), 9545 (1988). 34. Sanchez,C.,andLivage,J.,Sol-gelchemistryfrommetalalkoxideprecursors.

New J. Chem., 14, 513 (1990). P., Larbot, L., Guizard,C.,Cot,L.,Mirodatos,C.,and 35. Julbe,A.,Chanaud,

Borges, H., Lanthanum oxychloride catalytic membranes,Key Eng. Mater., 61/62,

65 (1991). D., Schaeffer, D. W., Assink, T. A.,Kay,B. D., and 36. Brinker,C.J.,Keefer,K. Ashley, C. S., Sol-gel transition in simple silicatesII,J. Non-Cryst. Solids, 63, 45 (1 984). 37. Larbot,A.,Julbe,A.,Guizard,C.,Cot,L.Silicamembranesbythesol-gel process, J. Memb. Sci., 44, 289 (1989). 38. Cuperus, F. P.,Bargeman, D., and Smolders, C. A., Characterizationof UF mem-

branes,membranecharacteristicsandcharacterizationtechniques,

Adv.Colloid Interfac. Sci., 34, 135 (1991). 39. Gregg, S. J., and Sing, K. S. W., Adsorption, Surface Area and Porosity, Acade-

mic Press, London, (1982). 40. Eyraud,C.,Quinson,J.F.,andBrun,M.,Theroleofthermoporometryinthe

study of porous solids, in Characterization of Porous Solids, Elsevier, Amsterdam, 1988, p. 295. 41. Charpin, J., and Rasneur, B., Caract6risation de la texture poreuse des matCriaux, Tech. Ing., 1050-1(France). 42. Eyraud, C., Application of gas-liquid permporometry to characterization of inorganic ultrafilters, inProc. Europe-Japan Cong. Memb. Mernb. Processes (E.Dridi and M. Nakagaki, eds.), 1984, p. 629.

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43. Mey-Marom, A., and Katz, M. J., Measurement of active pore size distribution of microporous membranes-a new approach, J. Memb. Sci., 27, 119 (1986). 44. Capannelli, G.,Becchi, I., Bottino, A., Moretti, P., and Munari, S., Computer driven porosimeter for ultraf3tration membranes;Characterisation of Porous Solids, Elsevier, Amsterdam, 1988. p. 283. 45. Glaves, C. L., and Smith, D. M., Membrane pore analysis via NMR spin lattice relaxation experiments, J. Memb. Sci., 46, 167 (1989). 46. Ramsay, J. D. F., and Avery, R. G.,Neutron scattering investigation of adsorption processes in modelporous systems, in Characterisation of Porous Solids II, Elsevier, Amsterdam, 1991, p. 235. 41. Long, G.G.,Krueger, S., Gerhardt, R. A., and Page, R. A., Small-angle neutron scattering characterizationof pmcessinghicrostructure relationships in the sintering of crystalline and glassy ceramics, J. Mater. Res., 6(12), 2706 (1991). H., Structureof 48. Schaefer, H. E., Wiiurschum,R.,Birringer,R.,andGleiter, nanometersizedpolycrystallineironinvestigatedbypositronlifetimespectroscopy, Phys. Rev. B, 38(14), 9545 (1988). 49. Alary, J. A., Bauer, J. M., and Boudier, G.,Synthesis of new inorganic membrane byelectrocoatingprocess,in Proc. International Congress on Membranesand Membrane Processes, Chicago, 1990, pp. 555-58. 50. Ilias, S., and Govind, R., Development of high temperature membranes for membrane reactor: An overview, AIChE Symp. Ser., 85(268), 18 (1989).

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Index

Acetic acid, 9, 37 Acetyl acetone, 11, 37, 520 ABO3,68,445,447 AnBn03n-l,63,445 An-lBn03n+29 68 Acid-base reactions Lewis, 41 Acrylic acid, 19 Aerogels, 311, 315, 331 acoustic attenuation in, 322 acoustic properties of, 323, 330 alumina, 3 11 applications of, 327 catalysts, 331 Cherenkove detectors, 327 gas filters, 331 laser fus,ion targets, 332 photoluminescent light sources, 33 1 solar panels, 3 18 thermal insulation, 328

[Aerogels] thermal insulation windows, 327 tiles, 327 base-catalyzed, 323 Bragg scattering, 324 Brillouin scattering, 321 carbon, 312,314,320,330 as supercapacitors, 330 carbon black-doped, 320 characterization: Bragg peaks, 318 light scattering, 317, 324, 344 NMR, 316 SAXS, 316, 323 skeletal, 316 structure, 315 density, 322,324 fractal range, 323 gas filters, 331 Guinier region, 317 insulating foams, 3 11, 319, 328

533

534

[Aerogels] iron oxide, 311 loss tangent, 324 mass fractal dimension, 317 mechanical properties of, 320 acoustic attenuation methods, 321, 324 pulse-echo method, 321 melamine-formaldehyde, 3 14 mixed oxide, 312 modulus of rupture (MOR), 322 nanostructure, 325 one-step silica, 318 optical properties, 319, 325 interferometric method, 327 Rayleigh scattering, 325, 344 organic, 313, 332 resorcinol-formaldehyde,3 14, 318 phonon mean free paths, 322 photon mean free path, 326 porosity, 322 quality factor, 324 Raman-Nath diffraction, 322 refractive index, 326, 344 silica, 266, 3 11 skeletal densities of silica, 316 sound velocities, 321, 322, 330 Kelvin-Voigt model, 324 thermal insulation of, 265, 328 thermal properties, 3 19 stationary hot-plate, 320 thermal conductivities, 319, 329 titania, 3 11 trap, 332 two-step silica, 3 18 Young's modulus of, 324 Alcohols alcoholysis, 35

INDEX [Alcohols] exchange reactions, 30,75, 77, 216 polydentate, 32 Alkali flux method, 67 Alkaline-earth metal nitrates, 65 Alkanolamines, 31, 35 Alkoxides chemically modified, 3 1 double, 40, 79 heteroleptic, 3 1 heterometallic, 41 homoleptic, 3 1 metal, 24, 25, 30,76,405 mixed metal, 75, 77 [M(OC2H40Me)nlm, 34 niobium, 35 precursor, 516 silicon, 3 transition metal, 5 16 Alkoxy alcohols, 35 groups, secondary, 5 tertiary, 5 ligands, 4 Allyl function, 19 &Alumina, 67 Aluminum alkyls, 117, 365 Aluminum chloride, 365 Aluminum etac, 11 Aluminum hydride, 116 Aluminum hydroxide, 184 Al(OC3H7)3, 79, 80, 81 [A11304(oH)24(H20)17+, 47 Aluminum nitride, 10, 115, 122, 361,365,381 Al(Et)3, 365 N-alkylamino alanes, 365 cationic impurities, 115 formation of, 121

INDEX [Aluminum nitride] oxygen content, 123 powder, 197 sintering aids, 115 TGA profiles, 121 Aluminum oxide boehmite, 190 EVA mix, 248 polystyrene with, 258 rheology, 199 S i c whisker composite, 187, 190 suspension, 176, 180, 181 ZrO2-SiC composites, 190 Aluminum sec-butoxide, 11,278 Ammonia, 24, 108 ammonolysis of dichlorosilane, 36 1 method, 76 Aryloxide, 3 1 A, WO3,66 Azeotropic distillation, 35, 77 Barium &oxides, 217,484 hydroxide, 217 metal, 217 BaCo3,71 Barium ferrites hexagonal 70 particles, 426 . platelet, 430 Bag(OH)(OR)g, 25 BaPb03,65 Ba2PbO4,65 BaPbXBil_,03 (BPB), 44 Ba phenoxide, 32 Bal-xKxBi03, 67,447 BaTi03.627 231 -based dielectric, 401 -based materials, 397

535

[BaTi03] crystal structure, 139 powder, 90 S~2Og-d0ped,93 BaTi system, 41, 217 sol-gel derived, 218 BaTi[(OiPr)4 iPrOH14, 8

Ba4Ti404(0iPr)16(iPrOH)x,44, 46 B ~ ~ ~ ( c I - O R ) ~ ( O45R ) ~ I ~ , Benzoyl peroxide, 19 Binder amount of, 250 ethylene ethylacrylate (EEA), 243 ethylene vinylacetate, 247 GPC of, 252 molten organic, 250 oxidation of, 252 polymeric, 243 stability of, 251 thermoplastic, 246, 257 waxes, 257 Bi2An_1Bn03n+3,69 Bi2CaSr2Cu208,73 Bismuth cuprates, 73,449, 450 Bi(2212), 449,450,451 Bi(2223), 449,450 B(N,C) ceramics, 364 Borazoles, 118 Boron nitride (see also Nonoxide ceramics), 118 borazine, 363 boric acid, 363 ceramic yields, 118 fibers, 371, 381 hexagonal, 118 trichloroborazine, 363 urea, 363 Boron polymers, 112, 118, 363, 364

536 Boronylpyridine, 112 Brownmillerite, 63

Ca2Co204.63 Ca2Co205,63 CaCu02,453 ca2Fe2-xMnx(co3)4, 64 Ca2Fe204,62 Ca2Fe205,62 Ca2Mn205,63 Ca2Nb207,68 Cag.5Ti2p30 12,61 Capacitor dielectric, 398, 465 lamination defects, 401 MLC, 398,416 Capillaries, 257 force, 515 tortuosity constant, 257 Carbides, transition metal 120 Carbon content, 122 excess, 110 precursor, 106 Carbothermal reduction, 105, 106 Carboxylic acids, 9, 19, 35 Catalysis or catalyst, 24, 515 acid, 3 base, 3 Ceramic membranes, 501 alumina (Membralox), 503,514 biphasic materials, 516 calibration, 509 Carbosep, 503 catalyic, oxidative coupling of methane, 519 charaterization, 521 Adzumi equation, 526 bubble-point method, 523 gas adsorptioddesorption technique, 523 Kelvin equation, 525

INDEX [Ceramic membranes] Laplace equation, 523 microcaloric analysis, 525 mercury porosimetry, 521, 522 N M R , 527 permeability, 526 positron life time spectroscopy, 527 rejection measurements, 526 small angle scattering, 527 texture, 522 thermoporometry, 521,526 Washburn equation, 523 fabrication, 504 acetyl acetone, 520 by colloidal, 513 by sol-gel, 512 cathodic sputtering, 528 CVD, 528 electrophoresis, 527 extrusion, 507 firing atmosphere, 5 17 inorganic binders, 507 organic additives, 508, 5 11 organic complexants, 518 physical vapor deposition, 528 plasticizers, 508 rheology, 508 slip casting, 505, 511 tape casting, 505,508 thermal evaporation, 528 Tilacac ratio, 521 Zrlacac ratio, 521 zirconium oxychloride, 5 17 flux performance, 504 macropores, 501 mesopores, 501, 512 nanofilters, 502, 506, 516 nanopores, 501,512 porosity, 501

INDEX [Ceramic membranes] supports, 503 carbon, 503 ceramic porous, 503, 504 multichannel, 503, 506 multilayered, 506 stainless steel, 503 Vycor glass, 502 zirconia, 503, 514, 517, 520 Ceramic powders, 129, 143 binder mixture, 239 characterization, 129 density, 131 hafnia, 94,96,97 morphology, 131, 136 nonsilicate, 75, 81 oxide, 75 physical properties of, 131 PLZT,85,88 polydisperse, 131 porosity, 131 silicate, 75, 78 silicon nitride, 129 size distribution, 131 specific surface area, 131 synthesis, 75, 359 YAG, 98 zirconia, 93, 97 Ceramic fibers, 106,359,382 BN, 371,381 graphite, 360, 371 green, 362 Sic, 375,385 silicon nitride, 375, 381 silicon oxynitride, 375, 390 Si-Ti-C-O,375 Tyranno, 375,387 Ceramic yields, 109, 119, 362 of boron carbide, 119 of boron nitride, 118 of polycarbosilane, 382

537 Cerium alkoxide, 5 isopropoxide, 15 Ce2(0ifi)8(mH)2,15 Ce604(0W4(acac)l2, 16 CFC in aerogel, 31 1, 329 blown polyurethane, 320, 328 Chemical modification, 3, 34 of alkoxide, 3 in ceramic membranes, 516 by complexation, 3 Chemical vapor deposition (see also CVD), 24,40, 48, 108, 118 Chevrel phases, 61 B-Chloroborazoles, 118 Chlorosilanes, 108, 114 Chromium-containing precursors, 122 CrN, 122 Cinnamic acid, 19 Coatings heat resistant, 391 infrared, 392 methods dip, 467 spin, 467 perhydropolysilazane, 392 tyranno, 391,392 C-0 bond cleavage, 33 Coercive force, 423 remnant, 423 coercivity, 425 Colloidal ceramics acid-base interaction, 200 gels from, 278 for membranes, 5 13 particles, 198 processing, 189 seeds, 187 Colloidal suspensions, 133, 164, 174,425

538

[Ceramic membranes] Doppler shift, 165 electrophoretic mobility, 182 electrostatic stabilization, 172 electrosteric stabilization, 186 ESA, 165,205 flocculation, 144 lyophobic sol, 172, 174 maghemite, 425 magnetite, 425 microemulsion, 434 polyelectrolytes, 186 stability, 144, 157, 172, 224 steric stabilization, 185, 186 zeta potential, 146, 164, 172, 176, 199,202 Columbite, 405 MgNb206,408 route, 406 Combustion method glycine, 73 urea, 73 Complexation step, 41 of zirconium akoxide, 21 Complexing additives, 12 Complexing reagent, 9 Composites alumina, 190 carbon-carbon, 364 ceramic-ceramic, 190 cordierite-ZrO2, 192 lithium aluminosilicate matrix, 192 molecular, 300 mullite matrix, 190 nano-, 300 Sic whisker, 190 zirconia, 190 Condensation, 3 CuFeO2,67 CU,MO&, 61

INDEX

Coprecipitation route, 62 Cordierite, 11, 178, 181, 275 Cracks aging liquid, 268 in gels, 266, 288, 314 pressure gradient, 269 routes to avoid, 272 solvent exchange, 269, 273 supercritical liquid, 269 Cross-linking, 362, 371 agents, 17 degree of, 115 of polycarbosilane, 383 Crystalline silicates, 79 Crystallite sizes, 123 Curing autoclave, 377 kinetics, 366 methods, 372 by oxidation, 385, 391 Raman spectra, 385 ring opening, 377 CVD, 23,40,48, 108, 295, 359, 465,528 plasma-enhanced, 360, 528

Darcy’s law, 270 liquid flux, 270 pressure gradient, 270 Density, 134 packing, 179, 199 relative, 181 skeletal, 3 18 Derjaguin equation, 197 Design chemical, 3 molecular, 3 Dielectric microwave, 400 multilayer capacitors, 481 thin films, 465

INDEX Dielectric constant of lead dielectrics, 41l of medium, 146, 163 of organic solvent, 12 of Ta205,473 B-Diketones, 9, 11, 35, 37 Diimid precipitation, 112 Dimethylchlorosilane, 108 Dip coating, 459,493 wetting for, 486 Direct pyrolysis, 77 Direct synthesis, 106 Disintercalation, 67 of alkali metal ions, 67 Dispersants, 252 acrylic acid, 217 amino groups, 205 amino polyisobutylene (APIB), 203,208 aminosilane (AHAS), 201, 206, 208 Darvan C,41 1 linolenic aid, 200, 206 mixed, 212 PMMA, 230,232,234 DLVO theory, 172, 174, 187,210, 256,514 Double alkoxides, 40,79 Ba-Nb, 484 Li-Nb, 492 Mg-AI, 79,81 Y-AI, 98 Double layer, 157, 160, 256, 514 adsorption, 158 adsorption density, 159 adsorption energy, 163 Coulombic potential, 160 Debye-Huckel length, 161, 172, 198 diffuse layer, 160,' 178 diffuse layer potential, 169

539

[Double layer] DLVO theory, 172, 174, 187, 210 electroneutrality, 160 electrostatc repulsion, 163 GCSG model, 160, 163 Helmholtz plane, 161 hydrodynamic slipping plane, 169 ionization, 158 Nernst equation, 165, 166 OHF' potential, 164 Poisson-Boltzmann equation, 160 relaxation, 178 shear plane, 164 specifically adsorbed ions, 163 Stem plane, 161 Stem potential, 172 surface charge, 158, 164 thickness, 146, 198 zeta potential, 172 Drying autoclave, 291, 314, 315 capillary condensation, 272 capillary stress, 270, 288 control chemical additives, 273 formamide, 273,279 glycerol, 273 oxalic acid, 273, 279 chlorination, 294 curvature, 270 Darcy's law, 270 dehydroxylating agent, 294 differential strain, 270 of gels, 269, 275, 288 hypercritical or supercritical conditions, 272, 315, 318 Laplace's law, 269, 275 shrinkage, 266,279

540

Electrochemical method, 73 Electroluminescent devices, 465 Electrooptical ceramics, 23 Kerr effect, 299 Electrosynthesis, 30 of polymeric precursors, 123 Enolic form, 11 Epitaxial thin films, 48,488 Epitaxy, 488,492 sol-gel, 488, 491,493 solid-state, 493 Ester exchange, 75,77 Esterification, 10 2-(2-Ethoxy-ethoxy)ethoxide,6 Ethylacetoacetate, 11 EuBa$u307,73 EXAFS spectra, 6 , 9 frequency shift, 10

Ferroelectrics, 23,48 1 amorphous, 493 devices, 481 electric dipoles, 498 materials, 482 KNbO3,482,485,488 LiNbO3, 482,485 LiTa03,489 phase transition, 493 PLZT, 483 PMN-PT, 482,486 PNZT, 489 PT, 482,486 SBN, 482,483,487 oxides, 482 thin film preparation, 481 Ferric oxides cobalt modified, 428,429 B-FeOOH, 426 y-FeOOH, 426 &Fe203 (see also Hematite), 168

INDEX Flocculation, 144, 256, 258 Fluoride compound NH4F, 277 HF, 277 Fourier transform infrared spectroscopy (FTIR,see also Infrared), 99, 147,320,367,368 adsorption spectra, 220 diffuse reflectance, 144 IR active, 144 photoacoustic, 144 spectra of boron nitride, 118 powders, 143 Fractal, 337 chain entnanglements, 347 cluster-cluster aggregation, 347, 350 cluster nucleation, 337 cluster size, 346, 354 density, 339 diffusion limited growth, 338 dimension, 339 Eden growth model, 337 gel network, 354 growth, 337 models, 350 rate factor, 348 number density, 341 post gelation reaction, 345 radial growth rate, 346 radius, 340 radius of gyration, 339 shape factor, 342 surface, 337 transmission intensity, 348 Freeze drying, 217,219

Gas-phase synthesis, 106, 108 GC/mass spectroscopy, 367

INDEX Gelation, 5 Gel optical composites: inorganic dopants, 298 laser dyes, 298 NLO,299 perylene-silica-PMMA, 299 silica-PMMA, 298 Gels amorphous films, 483 aging, 268,288 alkoxide-derived silica, 289 bloating, 281 bulk xerogels, 5 14 capillary forces, 275 chlorination treatment, 281 colloidal, 278,289 connectivity, 274 cracking of, 266 dehydroxylating agent, 281 densification of, 266 drying, 288 filler phases, 279 fluorinated, 277 fracture of, 268 helium atmosphere, 28 1 He-02 mixture, 28 1 HF catalyzed silica, 347 hydrothermal aging, 291, 292 liquids formamide, 272,273,279 dimethyl formamide, 272,279 mechanical strength, 266 microbubbles in, 281 microporosity, 274 monolithicity in, 265, 270, 273 network of, 268, 279 organic, 3 14 permeability, 266, 268, 270, 273,274 polymeric, 279 network, 346 polymerization, 268

541

[Gels] powder, 275 silica, 288 sintering of, 280 foaming during, 281 syneresis, 268 thermal expansion coefficient of liquid, 268 wettability, 274 Gel permeation chromatography, 252,368 Gel time, 9 of silicon alkoxide, 9 Glycols, 35 Gradient index (GRIN) by ion exchange, 303 lenses, 287, 303 titanium aluminosilicate glasses, 305 Grain growth: homogeneous, 2 16 inhomogeneous, 215 Gravity sedimentation, 132 Stoke’s law, 132 Green density, 266 Hafnia (HfO2),94,96 Hamaker constant, 172, 198,210, 256 polarizability, 21 1, 256 Hematite, 168, 426,437 Heteroalkoxide, 8 High Tc superconducting ceramics, 6, 197, 445 Homometallic alkoxides, 23 Hydrolysis, 3 acid, 81 base catalyzed, 5 19 of organometallic polymers, 363 polycondensation, 37 rate, 5, 12, 35, 217 reaction, 78

542

Hybrid compounds organic-inorganic, 18 Hydrolytic decomposition, 78 of metal alkoxides, 78 Hydrothermal transformation, 428, 437 Hyponitrite precursors, 73 Imino alanes, 116,366 poly-, 117, 118, 119 Indium tin oxide (ITO) film, 471 Infrared absorption spectra, 220 diffuse reflectance, 144 hydrogen bond, intensity of, 228 spectra of Bag 2Pb0 gTiO3,221,228 binders, 252 boron nitride, 118 silica, 28l Ta2O5 films, 471 Injection molding, 197, 246 Inorganic dopants, 298 Inorganic membranes, 501 fabrication by sol-gel, 502 supports, 503 supported, 505 tortuosity of channel, 504 Interface analysis calorimetry, 147 complexed ions, 145 in concentrated suspension, 165 counter ions, 145 diffuse layer, 145 of dispersed particles, 144 Doppler shift, 165 electrolyte concentration, 146 extended X-ray absorption fine structure, 147 potentiometric titration, 167, 168 streaming potential, 158 surface species, 147 in suspension, 144

INDEX Intercalation, 62, 67 Interfacial phenomena, 157, 187 electrochemical, 189 free energy, 174 Hamaker constant, 172 ion pairs, 169 at metauwater, 157 negative surface sites, 158 positive surface sites, 158 Intergrowth structure, 67 Ion exchange, 62,67 Isoelectric point (IEP),133, 146, 164, 182,200 pH, 146,200,225,228,230 point of zero charge (PZC), 147, 159, 167, 174,200,224 PZR, 164, 182, 183 Lability of alcohol, 25 of alkoxide bond, 40 La2-xBaxCu04, 70,445,447,448 La2Cal-xSrxCu06, 72 La3(0tBu)gk, 32 LaCoOg, 62 L ~ ~ C U O67~ + ~ , LaFeO3,62 L*e0.5co0.5(cN)6.5 H209 65 H209 65 %.5Ndo.’jCo(CN)6.5 LaNiO3,65 LNI-~M,(OH)~,64 Laser ablation, 492 composite target, 492 single target, 492 Laser dyes, 298 coumarin-314T, 380 fluorescein, 300 poly(p-phenylene vinylene), 300 rhodamine B, 300 rhodamine 6G, 300 Lattice parameter, 136 Lead acetate, 85 Lead alkoxides, 40

INDEX Lead-based electronic ceramics, 397 Curie temperature, 41 1 dissipation factor, 41 1 Goldschmidt tolerance factor, 403 PCN, 398,403,412 P m , 398,402 PEW, 398,403 phase stability, 401 PMN, 398,402,411,412,482 PNN, 398,403 PT, 398,403,411,412 relaxor, 403,416 Lead cuprates, 72 pbl-xbx(zr Ti,)l-x/403, 85 Lead niobates, $3 PbgNb404(0Et)24,44,45,47 Pb/Nb stoichiometry, 8 Pb4o(oEt)6, 8 [PB60(0Et)4lINb(OEt)gl4,8 Lead oxide (PbO): columbite method, 410 excess route, 416 MgO, 410 rich grain boundary, 415 vaporization, 216,404,408 Pb2Ru2-xPbx07-y.67 Lead titanates, 23, 217, 397 Ligand, 38 double-bridging, 3 1 triple-bridging, 3 1 Light scattering Fraunhofer and Mietheories, 133 Rayleigh equation, 133 LiCro2, 67 LiM02, 67 Lithium niobate, 47,482 EDAX of, 489 XRD of, 489 LiNb(OR)6, 41, 45 LiTa03,47

543

LiVO2,67 Low-temperature method, 72,216

Maghemite, 425 aspect ratio, 43 1 elongated, 425 spindle-type, 43 1 Magnetic bacteria, 434,439 critical field, 432 dipole, 424 moment, 424 orientation of, 423 dispersion, 433 domain, 423 ferro fluid, 433 fluid, 433 interaction energy, 424 medium, 430 particles, 421 properties, 422 recording materials, 424 recording media, 425 Magnetic particles bioparticles, 439 composite, 438 vinylferrocene, 438 chelating agent, 435 chromium dioxide ((2109,429 coated with silica, 432 dispersion, 433 oleic acid, 433 elongated, 43 1 longitudinal recording system, 43 1 FeOOH, 426 goethite, 426,428 hematite, 426,437 plastic ferromagnetism, 437 spindle-type, 43 1,435 iron powder, 428 iron nitride, 430

544

[Magnetic particles] Fe@, 430 Fe16N2 fih,430 Fe/N ratios, 430 maghemite, 425 magnetite, 425,433 spherical, 435 magnetotactic bacteria, 434,439 geomagnetic field, 439 MnFeO4,434 microemulsion, 434 polymeric microsphere, 439 unilamellar vesicles, 434 Magnesium aluminosilicate (MAS), 78 MgM2(0R)g, 7 9 , ~ Magnesium ethoxide, 79, 81 MgNb2(0Ac)2(0iW 1 0 , 4 4 4 6 Magnetic tapes audio, 425 video, 425 Magnetization domain, 423 remnant 423 saturation, 423,425 Melt processing, 369 Melt spinning, 371, 383 Membrane filtration gas separation, 512 microfiltration, 501,504 layers, 5 11 zirconia, 511, 518 nanofiltration, 501,504 ultrafWation, 501, 504, 512, 516 alumina, 504 Membrane permeability, 526 Adzumi equation, 526 biliquid permporomerty, 526 Hagen-Poiseuille relationship, 527

INDEX [Membrane permeability] Kozeny-Carman relationship, 527 laminar (Poiseuille), 526 molecular (Knudsen) flow regime, 526 permeability vs pore radius, 27 permporometry, 526 rejection measurements, 526 Mercury, 24 HgC12, 76 HgI2, 76 porosimetry, 136 Metal acetates, 405 Metal carbonates, 62 Metal halidehlcohol, 75 Metal hydrides, 30 Metal hydroxides, 405 Metal organic chemical vapor deposition (MOCVD), (see also CVD), 23,48, 108 Metal-organic compounds, 75 decomposition (MOD), 481,482 Metal oxides, 61,75 rare earth, 93 Metal silylamides, 24 Methacrylic acid, 19 Methacryl amido salicylic acid, 19, 20 Methanol, 3 12 2-methoxyethanol, 34.49 Microscopy optical, 132 scanning electron, 132 techniques, 136 transmission electron, 132 Microwave dielectric, 400 Mixed oxide method, calcined 217, 218,221 Mixing, 250 analysis of mixedness, 260

INDEX Mixing] degree of mixedness, 260 mechanism, 250 critical stress, 251 dispersive, 250 laminar, 250 molecular diffusion, 250 orientation of agglemerates, 25 1 process tolerance capability, 252 rate of, 261 simulation of, 261 Monte Carlo technique, 261 MoO3,66 M01-~W,03,65,66 Molten salt synthesis method, 217, 218,221 Monolithic ceramics aerogels, 270 silica, 265 xerogels, 266, 270, 271 Mullite (3Al203 2Si02), 79, 80, 190,275 Multicomponent ceramics, 8, 216 glass-ceramics, 78 oxides, 483 silicate glass, 78 thin films, 498 Nanoscopic cluster, 3 13 Nasicon, 61 Near-net-shape, 287 Nd2-xCexCu04, 72 NdNiOg, 65 Nd(OiPr)3(iPrOH)4, 32 Niobium alkoxides, 484 ethoxide, 13 oxide (Nb2O5), 406,408,409 polymers, 120

545

ml0028(NMe4)6 6H20, 13 m8010(oEt)20, 13 Nicalon", 375 ESR spectra of, 385 excess carbon in, 385 fiber, 385 Nitridation of polycarbosilane, 390 rate, 107 temperature, 107 Nitrides, transition metal, 120 Nonaqueous medium, 185, 197 Nonoxide ceramics, 105 aluminum nitride, 105, 115,361, 363,365 boron carbide, 105, 112, 361, 364 boron nitride, 118, 361, 363 composites, 359 elemental analysis, 370 microstructure, 369 silicon carbide, 105, 115, 361 silicon nitride, 105, 112, 115, 129,361 surface analysis, 370 titanium carbide, 363 titanium diboride, 361, 363, 366 titanium nitride, 361,363 Nuclear magnetic resonance (NMR), 171,368,527 2 7 ~ 1 , ~ for aerogels, 316 Knight shift, 139 magic angle spinning (MAS), 139 nuclear magnetogyric ratio, 138 nuclear spin number, 138 31P, 139 powders, 130, 138 ulse Fourier transform, 139 MAS, 139

546

INDEX

muclear magnetic resonance] spin relaxation times, 139, 171 spin-spin interaction, 138 tetramethyl silane (TMS), 138 Nucleophilic addition, 4 attack, 35 ligands, 5 substitution, 4 Oligomerization, 5 Oligomers, 3, 7, 30 Organic amines, 72 Ormosil, optical metallic quatum dots, 303 photochromic molecules, 303 photochromism, 303 semiconductor quantum dots, 303

Oxidation curing, 385 high temperature, 364 Raman spectra for polycarbosilane, 390 Oxophilic, 37 Oxopolymers, 3, 12 alkoxides, 5, 8 , 23,25 bridges, 8, 10 ligands, 8 Oxygen stoichiometry, 72 Paratungstate, 38 Particles acoustic field, 146 dispersion of, 157, 197 electrokinetics, 146 sonic amplitude (BA), 146, 147,164

electrolyte concentration, 146 electrophoresis, 146,164

[Particles] electrophoretic dispersion, 202 hydrodynamic shear plane, 146 in liquids, 145 magnetic, 421 monodispersed, 422 quasi-spherical, 504 sedimentation potential, 146 spherical, 244 ultrasonic wave, 146, 147 zeta potential of, 146, 164, 172, 176

Particle surface, 157 acid-base equilibria, 169 interaction, 200 adsorption density, 159 amphoteric dissociation, 158, 167

charge density, 159,163,164 complex, 169,170 counter ions, 160 DLVO theory, 172, 174,210 metal-hydroxo complexes, 158 potential, 164 potential determining ions, 166, 168

site-binding model, 169, 170 surface sites, 158 zeta potential, 146, 164, 172, 176,224

Particulate acicular, 80,426 prismatic, 80 Peptization, 217,277,513 Percolation, 190 theory, 337 Perovskite, 63,400,445 phase, 397 structure, 68 Pervaporation, 515 bonding, 34

INDEX

547

PLZT, 47,85,483 alkoxide derived, 88 DTA of, 88 powders, 85, 88 PMN ceramics, 8,47, 398,402, 408,412,482

Phenolic resins, 106 Photochromic properties, 19 Piezoelectric actuators, 481 resonance, 494 tranducers, 481 Plastic forming, 239 Polarization coercive field in ferroelectric, 494

ferrous, 496 remnant, 494 Polyacrylonitrile (PAN), 366 Polyborazylene, 118, 364 Polyborosilazanes, 364 Polycondensation via dechlorination, 379 in gels, 265 rate of, 267 Polymers (see also Polysilanes) acrylic acid, 217 atactic polypropylene, 257 branched-ring, 362,376 carbon-based, 362 chain, 245 mobility, 362 cross-linking, 17, 115, 362 entanglements in, 245 ethylene ethylacrylate (EEA), 243

ethylene vinylacetate (EVA), 247,248

grafted chains, 256 high-melting, 257 linear, 362

[Polymers] melts, 245 metal-organic, 376 organo-silicon, 377 PMMA, 230,232,234 polyacrylonitrile (PAN), 366 polyborosilazane, 364 polyethylene, 252 polysilazane, 113, 114, 361, 380 polysiloxanes, 113, 362, 380 polystyrene, 258 polytitanosilane (PTC), 379 preceramics, 360,361 pyrolysis, 362, 375 ring, 114,362,376 sterically hindered, 362 structure, 362 thermoplastic resin, 257 Poly01 process, 421,434 Polyoxo anions, 38 Polyoxometallates, 18 Polysilanes, 108,377 borodiphenyl siloxane, 377, 379, 392

carbosilanes, 108, 109,375, 381, 382

oxygen cured, 391 dimethyl silane, 377 dimethylsiloxane (PDMS), 303, 361

silastyrene, 377 silazanes, 113, 114, 361, 380 siloxanes, 113, 362, 380 titanosilane (PTC), 379,387 Polyvinyl borazine, 364 Pore BET adsorption, 315 macro- ,meso- , and micro- , 274

mercury porosimetry, 316 porosity, 266

548

Pore1 radius, 18 size, 269, 273, 315 Porous structure;503 bubble-point method, 509 capillary shaped, 509 concentric shell, 511 diameter, 506 Dubinin-Astakov pore volume, 520 layers, 507 oriented, 504 total pore volume, 525 Portland cement, 199,208 slumes, 209 surface, 209 tensile strength of cast, 209 Potassium hydride, 114, 115,381 [K(OtBu)lq, 32 K4SiW11039, 18 Powder characterization, 129 A E S , 140, 141, 142 atomic absorption spectroscopy, 136 bulk composition, 136 density, 134 inductively coupled plasma, 136 milling, 405 morphology, 136 NMR and EPR, 137, 140 size, 131 surface area, 134 X P S , 140, 141 XFS, 137 XRD, 140 Precursors improper, 49 1 method, 62 proper, 491 organosilazane, 361 Primary alcohols, 24 Primary particles, 131

INDEX Processing aids, 200, 217, 243, 230,252,256 Pyrochlore phase, 398,402 structure, 67 Pyroelectric coefficient, 49 1 current, 494 heat sensors, 494 Pyrolysis, 77, 109, 361, 362, 371-376 or organometallic polymers, 360 of polysilanes, 110 of polysilazanes, 114, 115 of polyvinyl pentaborane, 364 rate of gas evolution, 363 of B-triamino-Ntriphenyborazine, 363 PZT, 48,481,482

Raman spectroscopy, 143 frequency, 144 Rayleigh scattering, 144 surface silica, 144 Rapid thermal annealing, 487 Rayleigh scattering, 144, 298 ReO3,66 Rheology, 144,239,369 Bingham solid, 240 Brookfield shear rate, 208 for ceramic membranes, 508 continuous-phase, 247 dilatant, 240, 243 drag force, 250 Einstein relation, 249 constant, 249 elastic deformation, 241 fluidity, 250 frictional voids, 240 of gels, 274 inverted plasticity, 243 Newtonian flow, 240

INDEX [Rheology] nonNewtonian, 240,242 properties, 188,226, 360 pseudo plastic, 240 and related behavior, 199 Reynolds number, 250 shear rate, 240 thixotropic, 241 threshold shear stress, 240 viscoelastic, 241 viscosity, 240 of ZR02lAl203, 188 RbxP8W3201 129 61 Rock salt structure, 62 Schlenk vacuum line, 363 Screen printing, 45 1,459 Sialon, 61 Silanols, 35 Silica aerosil, 277 Cab-0-Sil, 276 with carbon black, 320 coating, 190 colloidal, 276, 288 condensed, 3 13 electrokinetics, 192 Eu-doped, 294 fluorescence spectra, 296 fiber-optic amplifiers, 293 fumed, 273 gel, 288 residual silanols, 294 glasses, 287 infrared spectra of, 281 La-dopd, 293 laser glasses, 293 Ludox", 273 modified ormocer, 279 ormosil, 279,287,300,302 Nd-doped, 293

549

[Silica] quartz transformation, 273 rheology, 199 Stober, 273 surface, 205 uranium fluorescence, 296 Silicates grain boundary phase, 129 GRIN glasses, 305 potassium, 276 sodium, 276,312 surface silanol, 182 Silicon &oxide, 5, 80, 190 sodium coupling reaction, 365 TEOS, 519 Silicon carbide, 361 auger spectrum of, 142 fibers, 375, 385 ESR spectra of, 385 excess carbon in, 385 filler, 245 Nicalona, 375 polytypes, hexagonal 108 powder, 243 from rice shell, 391 13-silicon carbide, 108, 385 silicon nitride composites, 115 Tic composit fibers, 387 whiskers, 391 composites, 190, 192 Si(CH3)3Cl, 109 Si(CH3)3CH=CH2, 109 Silicon monomers dichlorosilane, 377 dimethyldichlorosilane,377 HSiCH3C12,380 dodecamethylcyclohexasilane, 379 phenylmethyl dichlorosilane, 377 methyltrichlorosilane, 380

550

Silicon nitride, 105, 112, 115, 129, 258,261 auger spectrum of, 142 coatings, 59 fibers, 375,381 a and B mixture, 361 substrates, 257 XRD of, 390 Si-0-C, 383 Si-0-Si, 383 Si-0-Ti, bonds, 17 Siloxanes, 17 Si-Ti-C-0 (Tyranno) fibers, 375, 387 Silylamid derivatives, 25 Sintering additives, 105, 115, 320 of B% 2Pb0 gTiO3,230 of gels; 272,'280 B-AgA102,67 Slip casting, 505,511 NaCaqNb5017,68 Na2Fe60(CH3)18 6CH30H, 8 NaMo406, 61 Sodium sulfur battery, 329 NaZn2Mo308.61 Sol-gel chemistry, 3 derived B% 2Pb0 gTiO3, 225-230 derived powders, 216 epitaxy, 488,491,493 in ferroelectrics, 48 1,482 homogeneity, 405 hydrolysis, 482 under basic conditions, 267 for membranes, 502,512 method, 62.7 1,215,465 multiphasic, 187 polycondensation, 265,482 rate of, 267 process, 41,75, 265,287,492

INDEX [Sol-gel] routes, 121 synthesis, 3 thin film, 465 Solid-state chemistry, 61 Solvation, 6 Solvents, 24 acetone, 204,205 enolization of, 206 basicity scale, 201 benzene, 23 1,232 chloroform, 202,278 l-decanol, 278 ethanol, 204,205 ethyl acetite,'2W good, 230 hexane, 204,23 1,232 methylene chloride, 203 mixed, 2 12 nonpolar, 41 organic media, 199,432 for PMMA, 231 0-point, 231,232 solubility parameter, 230 THF, 200 toluene, 200 Sonochemistry, 24 Spin coating, 467, 493 contact angle, 486 oxide substrates, 486 wetting, 486 Spinel (MgA1204), 81 Stability colloidal, 224 dispersion, 144,432 electrostatic, 172, 200 electrosteric, 86 ratio, 176 of solvates, 25 steric, 185, 186, 198 barrier, 278 hindrance, 5

INDEX

[Stability] repulsion, 185 stabilization, 185, 186, 198 stabilizer, 204 Steric stabilization, 185, 186, 198 chain length of stabilizer, 198 entropic repulsion, 256 Storage modulus vs. oscillation frequency, 244 Sr2PbO4,65 Sr phenoxide, 32 sr4(oph)~(phoH)2(THF)6. 32. srTi2(oiPr)8(iProH)3,45 Substitution reactions, 25 Superconducting cuprates, 62, 70, 197,445 Bi(2212), 449 Bi(2223), 449 copper oxide, 447 grain boundaries, 459 La(214), 447,448 La2xSrxCu4, 453,454 preparation, 453 melt textured growth, 454 quench and melt growth (QMG),454 traveling solvent floating zone (TSFZ), 453 PrBa2Cu30x, 448 SmBa2Cu307,456 superconducting grains, 46 1 T1(1223), 447,452 weak links, 459 Y(123), 65,70-72, 139, 447-449 Y2BaCuOg,454 Superconducting wires, 459 intergrain Jc, 459 transport Jc, 459 weak link, 459 Superconductivity,446 BCS theory, 446

551

[Superconductivity] Bean’s equation, 456 critical current density, 447 diamagnetism, 450 electron-phonon interaction, 446 isotope effect, 447 magnetization hysteresis, 457 Meissner effect, 450 transport Jc, 456 Surface acidic sites, 200 acidity-basicity, 209 amphoteric, 199 area, 134 basic sites, 199 of cement, 209 charge, 158, 164 contact angle, 257 force apparatus, 188 induced coating, 189 of mica, 187 -OHgroup, 205 of silica, 205 sites, 157 species, 147 tension, 257, 510 Surface area of B% 2Pb0 gTiO3,219 BET, i34,2i9 internal and external, 134 of lanthanum oxycloride membrane, 520 permeametry, 134 Surfactants for stabilization, 133 Suspensions filled, 247 interparticle forces, 247 mixed, 191 stability, 144, 157, 172, 197 viscosity, 240 Suspension stability, 144, 157, 172, 197,258

552

[Suspension stability] attractive energy, 198 Born repulsion, 173 Derjaguin equation, 197 electrostatic stabilization, 172 Eliers & Korff rule, 175 flocculation, 144 Hamaker constant, 172 kinetically stable mixed, 191 potential energy barrier, 176 primary maximum, 173 rapid coagulation, 176 secondary minimum, 173, 189 slow coagulation, 175 Smoluchowski limit, 176 stability ratio, 176

Tantalum alkoxide, 466 Tantalum oxide, 465 breakdown voltage, 478 dielectric constant, 473 electron diffraction pattern of, 478 films, 471 infrared spectra of, 472 microstructure, film, 475 refractive index of, 473 &Ta205,472 Tape casting, 45 1,508 for ceramic membranes, 505 doctor blade, 45 1, 508 Teratogenic, 49 Tertiary butoxide, 77 Tetraethylorthosilicate or tetraethoxysilane (TEOS), 80,278,519 Tetramethoxysilane (TMOS), 278, 312,519 Texture-oriented films, 487 epitaxial growth, 488 of LiNbO3,492

INDEX [Texture oriented films] nearest atomic arrangement, 492 lattice mismatch, 488,491 KNBO3,488 substrates, 488,489 zro2,493 Thallium cuprates, 72,447, 449, 45 1-453,456,457 Tl2O3,72 Thermal anlysis DSc, 369 DTA, 368,472 TGA, 367 TMA, 368 Thermal decomposition, 72,78, 383 of Ba[TiO(C204)2], 62 Of LaCO(CN)6 5H20,62 of Li[Cr(C204)2(H20)21, 62 of metal alkoxides, 482 Thermalysis reactions, 40 Thin films epitaxial, 488 gel, 486 . ITO, 471 multilayer capacitors, 481 preferred crystallographic orientation, 487 preparations, 467 dip coating, 459 sol-gel, :482 spin coating, 467 sputtering, 467 substrates, 486 tantalum oxide, 465 thickness, 467 transparent, 466 Titanates, powder, 90 Titanium albxides, 5, 14, 19, 32, 278

553

INDEX Ti alkoxy oxoacetates, 10 Titanium compounds, 366 acetilides, 120 carbide, 120 diboide, 361, 363, 366 nitride, 120 precursors, 122 tetrachloride, 120 tetrakis-dialkylamine, 120 TiNmiC solid solutions, 122 Ti containing monomers, 366 Ti02 gels, 10 powder, monodispersed, 6 TiN, 107 Ti20(0Ac)6, 11 Ti(acac)2(0R)2, 14,48 Ti-glycolate bond, 37 Ti-O-Ba bonds, 8 Titration conductometric, 147 potentiometric, 147, 167 Topochemical dehydration, 65 methods, 62 reactions, 65 Trialkylsi1oxo.n donor, 34

S-Triamino-N-triphenylbonine, 363 B-Trichloroborazine, 363 Triphenyl boron, 112 W(OEt)6, 11 WO3,66 bronze, 66 hexagonal, 70 intergrowth, 70

Uranium hexafluoride, 503 isotopes, 503

Video recorder, high-definition, 430 Viscosity, 240 of alumina suspension, 200 of supension, 342

Wetting, 257 adhesional, 257 contact angle, 257 immersional, 257 spreading, 257 surface tension, 257 Wolframite, 405 AlNb04, 408 CrNbO4,408 W-0-Si-C bonds, 18

XANES spectra, 6 X-ray photoelectron spectroscopy, 130, 140-143,370 binding energy, 141 photoelectrons, 141 Scofield cross-sectional values, 141 Sic, 142 Si3N4, 142 Si02, 142, 143 X-ray powder diffraction, 140,367 pattern of Ta20.3472 X-ray Sedigraph, 132

YAG (Y3Al5O12) (see also Yttrium aluminum garnet), 98,100 Yttrium aluminum garnet, 98,100 YBa2Cu307 (YBCO), 65,70-72, 139 447,449

554

Zero-expansion ceramic, 61 Zinc alkoxides, 7, 30 ZnNb2O6,408 Zirconia alumindsilicon carbide composites, 190 dispersion, 192 green bodies, 180 powder, monodispersed, 15 submicron, 93, 97 toughened alumina with (ZTA), 189

INDEX [Zirconia] Y203-stabilized. 94, 97 Zirconium borides, 366 Zirconium organ0 complex, 15, 37 &oxides, 14 polymers, 120 Zirconium-oxopoly-PA, 20, 21 -methacry1 amidosilicylated copolymers, 20