Progress in Nanotechnology: Processing (Progress in Ceramic Technology)

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progress in Nanotechnology

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Progress in Nanotechnology Processing

A Progress in Ceramic Technology series publication

@WILEY A John Wiley & Sons, Inc., Publication

Copyright 02010 by The American Ceramic Society. All rights reserved. Published by John Wiley & Sons, Inc., Hoboken, New Jersey. Published simultaneously in Canada. No part of this publication may be reproduced, stored in a retrieval system, or transmitted in any form or by any means, electronic, mechanical, photocopying, recording, scanning, or otherwise, except as permitted under Section 107 or 108 of the 1976 United States Copyright Act, without either the prior written permission of the Publisher, or authorization through payment of the appropriate per-copy fee to the Copyright Clearance Center, Inc., 222 Rosewood Drive, Danvers, MA 01923, (978) 750-8400, fax (978) 750-4470, or on the web at www.copyright.com. Requests to the Publisher for permission should be addressed to the Permissions Department, John Wiley & Sons, Inc., I 1 1 River Street, Hoboken, NJ 07030, (201) 748-601 1, fax (201) 748-6008, or online at http://www.wiley.com/go/permission. Limit of Liability/Disclaimer of Warranty: While the publisher and author have used their best efforts in preparing this book, they make no representations or warranties with respect to the accuracy or completeness of the contents of this book and specifically disclaim any implied warranties of merchantability or fitness for a particular purpose. No warranty may be created or extended by sales representatives or written sales materials. The advice and strategies contained herein may not be suitable for your situation. You should consult with a professional where appropriate. Neither the publisher nor author shall be liable for any loss of profit or any other commercial damages, including but not limited to special, incidental, consequential, or other damages. For general information on our other products and services or for technical support, please contact our Customer Care Department within the United States at (800) 762-2974, outside the United States at (31 7) 572-3993 or fax (3 17) 572-4002.

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Progress in nanotechnology : processing. p. cm. - (Progress in ceramic technology) Includes index. ISBN 978-0-470-40839- 1 (cloth) I . Nanostructured materials. 2. Nanotechnology. TA4 18.9.N35P765 2010 620'.56c22 2009035822 Printed in the United States of America. 10 9 8 7 6 5 4 3 2 1

Contents

xi

Introduction Synthesis Methods for Powders

Freeze Casting as a Nanoparticle Material-Forming Method K. Lu and X. Zhu Int. J. of Appl. Ceram. Techno/., Vol. 5, Is. 3, p. 219-227

Preparation of a Nanoscale/SOFC-GradeYttria-Stabilized Zirconia Material: A Quasi-Optimization of the Hydrothermal Coprecipitation Process

3 13

Y-C Chang, M-C Lee. W-X Kao, and T-N Lin Int. J. of Appl. Ceram. Techno/., Vol. 5, Is. 6, p. 557-567

Synthesis of NanosizeTin Dioxide by a Novel Liquid-Phase Process

25

Y. Zhou, N. Dasgupta, and A. Virkar J. Am. Ceram. SOC., Vol. 91, No. 3, p. 1009-1012,2008

Fabrication of Nanocomposite Powders of Carbon Nanotubes and Montmorillonite

29

J. Feng and Q. Wang J. Am. Ceram. SOC.,Vol. 91, No. 3, p. 975-978,2008

Synthesis of Highly Dispersed Barium Titanate Nanoparticles by a Novel Solvothermal Method X. Wei, G.Xu, Z. Ren, Y. Wang, G. Shen, and G.Han

33

J. Am. Ceram. SOC., Vol. 91, No. 1, p. 315-318,2008

Continuous Production and Harvesting of Inorganic-Ceramic Nanoparticles S.A.E. Abdulla, P.A. Sermon, M. Worsley, and I.R. Collins CESP, Vol. 28, No. 6, p. 131-141,2008

Nanocrystalline Scandia Powders Via Oxalate Precipitation: The Effects of Solvent and Solution pH Z. Xiu, J-G. Li, X. Li, D. Huo, X. Sun, T. Ikegami, and T. lshigaki J. Am. Ceram. SOC.,Vol. 91, No. 2, p. 603-606,2008

A Pulse Combustion-Spray Pyrolysis Process for the Preparation of Nano- and Submicrometer-Sized Oxide Particles

37

49 53

W. Widiyastuti, Wei-Ning Wang, Agus Purwanto, 1. Wuled Lenggoro, and Kikuo Okuyama J. Am. Ceram. SOC., Vol. 90, No. 12, p. 3779-3785,2007

One-Step Synthesis of Luminescent Nanoparticles of Complex Oxide, Strontium Aluminate C. Li, Y. Imai, Y. Adachi, H. Yamada, K. Nishikubo, and C-N Xu

61

J. Am. Ceram. SOC.,Vol. 90, No. 7, p. 2273-2275,2007

Nan0 a-Al,O,

Powder Preparation by Calcining an Emulsion Precursor

Y-C Lee, S-B Wen, L. Wenglin, and C-P Lin J. Am. Ceram. SOC., Vol. 90, No. 6, p. 1723-1727,2007 Contents

65

V

Lanthanum Strontium Manganite Powders Synthesized by Gel-Casting for Solid Oxide Fuel Cell Cathode Materials

71

L. Zhang, Y. Zhang, Y. Zhen, and S. Jiang J. Am. Ceram. SOC.,Vol. 90, No. 5, p. 1406-1411,2007

Preparation of Matrix-Type Nickel Oxide/Samarium-Doped Ceria Composite Particles by Spray Pyrolysis

77

S. Suda, K. Kawahara, M. Kawano, H. Yoshida, and T. lnagaki J. Am. Ceram. SOC.,Vol. 90, No. 4, p. 1094-1100, 2007

Novel Low-Temperature Synthesis of Ferroelectric Neodymium-Doped Bismuth Titanate Nanoparticles

85

P. Prakash, A. Garg, M. Roy, and H. Verma J. Am. Ceram. SOC., Vol. 90, No. 4, p. 1295-1298, 2007

Hydrothermal Synthesis of CdMoO, Nano-Particles

89

X. Jiang, J. Ma, B. Lin, Y. Ren, J. Liu, X. Zhu, J. Tao, Y. Wang, and L. Xie J. Am. Ceram. SOC.,Vol. 90, No. 3, p. 977-979, 2007

Chromium-Doped Forsterite Nanoparticle Synthesis by Flame Spray Pyrolysis

93

T. Tani, S. Saeki, T. Susuki, and Y. Ohishi J. Am. Ceram. SOC.,Vol. 90, No. 3, p. 805-808,2007

Formation of AI,O,-Tic

Composite Nano-Particles Synthesized from Carbon-Coated Precursors

97

H. Kaga and R. Koc J. Am. Ceram. SOC.,Vol. 90, No. 2, p. 407-411,2007

Synthesis of Sm,~,Sr,,CoO,~,

and La,~,Sr,,4Co0,~,

Nanopowders by Solution Combustion Process

103

N. Bansal and Z. Zhong CT; Vol. 195, p. 23-32,2006

Colloidal Processing and Sintering of Nano-ZrO, Powders Using Polyethylenimine

113

Y. Hotta, C. Duran, K. Sato, and K. Watari Cx Vol. 190, p. 85-93,2006

Synthesis of High Purity p-SiAION Nanopowder from a Zeolite by Gas-Reduction-Nitridation

123

T. Yamakawaa, T. Wakihara, J. Tatami, K. Komeya, and T. Meguro CT Vol. 190. P.3-8.2006

A Novel Supercritical CO, Synthesis of Amorphous Hydrous Zirconia Nanoparticles, and Their Calcination to Zirconia

129

M-H Lee, H-Y Lin, and J. L. Thomas J. Am. Ceram. SOC.,Vol. 89, No. 12, p. 3624-3630,2006

Praseodymium-Doped Photo-Luminescent Strontium lndate Nanoparticles by Ultrasonic Spray Pyrolysis

137

S. E. Lin, K. Borgohain, and W. C. J. Wei J. Am. Ceram. SOC., Vol. 89, No. 10, p. 3266-3269, 2006

Nano-Blast Synthesis of Nano-size Ce0,-Gd,O,

Powders

141

Oleg Vasylkiv, Yoshio Sakka and Valeriy V. Skorokhod J. Am. Ceram. SOC., Vol. 89, No. 6, p. 1822-1826,2006

Sol-Gel Processing and Characterization of Phase-Pure Lead Zirconate Titanate Nano-Powders

147

Yasir Faheem and M. Shoaib J. Am. Ceram. SOC.,Vol. 89, No. 6, p. 2034-2037, 2006

Synthesis of AIN Nanopowder from -y-Al,O,

by Reduction-Nitridation in a Mixture of NH,-C,H,

151

Tomohiro Yamakawa, Junichi Tatami, Toru Wakihara, Katsutoshi Komeya, Takeshi Meguro, Kenneth J. D. MacKenzie, Shinichi Takagi, and Masahiro Yokouchi J. Am. Ceram. SOC.,Vol. 89, No. 1, p. 171-1 75, 2006

Membranes, Films, and Coatings

Microporous ZrO, Membrane Preparation by Liquid-Injection MOCVD

159

S. Mathur, E. Hemmer, S. Barth, J. Altmayer, N. Donia, 1. Kumakiri, N. Lecerf, and R. Bredesen CESP, Vol. 28, No. 6, p. 165-173, 2008 vi

Progress in Nanotechnology:Processing

Growth of Barium Hexaferrite Nanoparticle Coatings by Laser-Assisted Spray Pyrolysis

169

G. Dedigamuwa, P. Mukherjee, H. Srikanth, and S. Witanachchi CESP, Vol. 28, No. 6, p. 73-81, 2008

Two Phase MonaziteKenotime 3OLaPO4-7OYPO, Coating of Ceramic Fiber Tows

179

E. Boakye, R. Hay, P. Mogilevsky, and M. Cinibulk J. Am. Ceram. Soc., Vol. 91, No. 1, p. 17-25,2008

Template-Free Self-Assembly of a Nanoporous TiO, Thin Film

189

Y. Gao, M. Nagai, W-S Seo, and K. Koumoto J. Am. Ceram. SOC.,Vol. 90, No. 3, p. 831-837,2007

Nano-Sized Hydroxyapatite Coatings on Ti Substrate with TiO, Buffer Layer by E-beam Deposition

197

S-H Lee, H-E Kim, and H-W Kim J. Am. Ceram. SOC.,Vol. 90, No. 1, p. 50-56,2007

Sol-Gel Routes to Nanostructured Patterned FerroelectricThin Films with Novel Electronic and Optical Functions

205

M. Kuwabara, Y. J. Wu, J. Li, and T. Koga Ceramic Transactions, Vol. 196, p. 371-380, 2006

Preparation and Properties of Hydrothermally Stable y-Alumina-Based Composite Mesoporous Membranes

215

Md. Hasan Zahir, Koji Sato, Hiroshi Mori, Yuji Iwarnoto, Mikihiro Nornura, and Shin-ichi Nakao J. Am. Ceram. Soc., Vol. 89, No. 9, p. 2874-2880,2006

Synthesis and Tribological Behavior of Silicon Oxycarbonitride Thin Films Derived from Poly(Urea)Methyl Vinyl Silazane

223

T. Cross, R. Raj, T. Cross, S. Prasad, and D. Tallant Int. J. of Appl. Ceram. Techno/., Vol. 3 No. 2, p. 113-1 26, 2006

Synthesis and Tribology of Carbide-Derived Carbon Films

237

A. Erdemir, A. Kovalchenko,C. White, R. Zhu, A. Lee, M. J. McNallan, B. Carroll and Y. Gogotsi Int. J. of Appl. Ceram. Techno/.,Vol. 3, No. 3, p. 236-244,2006

Nanotubes, Nanorods, and Nanowires

Design, Fabrication and Electronic Structure of Oriented Metal Oxide Nanorod-Arrays

249

L. Vayssieres CESP, Vol. 28, No. 6, p. 187-193,2008

Electrospinningof Alumina Nanofibers

257

K. Lindqvist, E. Carlstrorn, A. Nelvig, and B. Hagstrorn CESP, Vol. 28, No. 6, p. 41-51,2008

ZnO Nanofiber and Nanoparticle Synthesized Through Electrospinningand Their Photocatalytic Activity Under Visible Light

269

H. Liu, J. Yang, J. Liang, Y. Huang, and C. Tang J. Am. Ceram. SOC.,Vol. 91, No. 4, p. 1287-1291,2008

Synthesis of Carbon Nanotubes and Silicon Carbide Nanofibers as Composite Reinforcing Materials

275

H. Li, A. Kothari, and B. W. Sheldon CESP, Vol. 27, No. 8, p. 41-48,2007

Polymer Fiber Assisted Processing of Ceramic Oxide Nan0 and Submicron Fibers

283

S. Shukla, E. Brinley, H. J. Cho, and S. Seal, CESP, Vol. 27, NO. 8, p. 57-68,2007

Growth of Quasi-Aligned AIN Nanofibers by Nitriding Combustion Synthesis

295

M. Radwan and Y. Miyamoto J. Am. Ceram. SOC.,Vol. 90, No. 8, p. 2347-2351, .2007

Contents

vii

Synthesis and Optical Properties of Mullite Nanowires

301

H-K Seong, U. Kim, M-H Kim, H-J Choi, Y. Lee, and W-S Seo J. Am. Ceram. SOC.,Vol. 90, No. 6, p. 1937-1939, 2007

(Nao,,Ko,2)o,,Bi,,5Ti03 Nanowires: Low-Temperature Sol-Gel-Hydrothermal Synthesis and Densification Y-D Hou, L. Hou, T-T Zhang, M-K Zhu, H. Wang, and H. Yan

305

Synthesis and Characterization of Ce,Gd,O,_,

311

J. Am. Ceram. SOC.,Vol. 90, No. 6, p. 1738-1743,2007

Nanorods

J. S. Lee and S. Kim

J. Am. Ceram. SOC.,Vol. 90, No. 2, p. 661-663,2007

Synthesis and Characterization of Cubic Silicon Carbide (p-Sic) and Trigonal Silicon Nitride (a-Si,N,) Nanowires

315

K. Saulig-Wenger, M. Bechelany, D. Cornu, S. Bernard, F. Chassagneux, I? Miele, and T. Epiciers CESP, VOl. 27, NO. 8, p. 81-88,2007

Synthesis of Boron Nitride Nanotubes for Engineering Applications

323

J. Hurst, D. Hull, and D. Gorican CESP, Vol. 27, NO. 8, p. 95-102,2007

Novel Process of Submicron-Scale Ceramic Rod Array Formation on Metallic Substrate

331

K. Okamoto, S. Hayakawa, K. Tsuru, and A. Osaka CT; VOl. 195, p. 133-138,2006

Tin Oxide Nanoparticle-Functionalized Multi-Walled Carbon Nanotubes by the Vapor Phase Method

337

W. Fan, L. Gao, and J. Sun J. Am. Ceram. SOC.,Vol. 89, No. 8. p. 2671-2673, 2006

Electrospinning: A Simple and Versatile Technique for Producing Ceramic Nanofibers and Nanotubes

341

D. Li, J. McCann, Y. Xia, and M. Marquez J. Am. Ceram. SOC.,Vol. 89, No. 6, p. 1861-1 869, 2006

Nanocomposites and Nanostructures

Chemical Precipitation Synthesis and Optical Properties of ZnO/SiO, Nanocomposites

353

H. Yang, Y. Xiao, K. Liu, and Q. Feng J. Am. Ceram. SOC.,Vol. 91, No. 5, p. 1591-1596, 2008

Low-Temperature Processing of Dense Hydroxyapatite-Zirconia Composites

359

Y. Nayak, R. Rana, S. Pratihar, and S. Bhattacharyya Int. J. Appl. Ceram. Techno/.,Vol. 5, No. 1, p. 29-36, 2008

Synthesis and Characterization of Chalcogenide Nanocomposites

367

J. Martin and G. Nolas ESP, VOl. 28, NO.8, p. 221-226, 2008

Self Assembled Functional Nanostructures and Devices

373

C. S. Ozkan CESP, VOl. 28, NO.6, p. 91-94, 2008

Carbon Nanotube (CNT) and Carbon Fiber Reinforced High Toughness Reaction Bonded Composites

377

I?Karandikar, G. Evans, and M. Aghajanian CESP, Vol. 28, No. 6, 2008, p. 53-63, 2007

Syntheis, Characterization and Measurements of Electrical Properties of Alumina-Titania Nanocomposites

389

V. Somani and S. Kalita ESP, Vol. 27, No. 8, p. 11-22, 2007

A New Ternary Nanolaminate Carbide: Ti3SnC,

401

S. Dubois, T. Cabioc'h, P. Chartier, V. Gauthier, and M. Jaouen J. Am. Ceram. SOC.,Vol. 90, No. 8, p. 2642-2644, 2007

viii

Progress in Nanotechnology: Processing

Fabrication of a Nano-Si,N,/Nano-C Composite by High-Energy Ball Milling and Spark Plasma Sintering

405

X. Xu, T. Nishirnura, N. Hirosaki, R-J Xie, and H. Tanaka J. Am. Ceram. SOC., Vol. 90, No. 4, p. 1058-1062, 2007

Conversion of Bamboo to Biomorphic Composites Containing Silica and Silicon Carbide Nanowires

41 1

T. L. Y. Cheung and D. H. L. Ng J. Am. Ceram. SOC., Vol. 90, No. 2, p. 559464,2007

Novel Processing to Produce Polymer/Ceramic Nanocomposites by Atomic Layer Deposition

41 7

X. Liang, L. Hakirn, G-D Zhan, J. McCorrnick, S. George, A. Weirner, J. Spencer II, K. Buechler, J. Blackson, C. Wood,

and J. Dorgan J. Am. Ceram. SOC.,Vol. 90, No. 1, p. 57-63,2007

Intra-Type Nanocompositesfor Strengthened and Toughened Ceramic Materials

425

S. Choi, S. Honda, S. Hashirnoto, and H. Awaji C7; VOl. 190, p. 173-180,2006

Prepation and Properties of Mullite-Based Iron Multi-Functional Nanocomposites

433

H. Wang, W. Wang, Z. Fu, T. Sekino, and K. Niihara C7; Vol. 190, p. 203-211,2006

Electrospinning of Ceramic Nanofibers and Nanofiber Composites

443

J. Yuh, H. Park, and W. Sigmund Ceramic Transactions, Vol. 190, p. 9-19, 2006

Microstructure and Properties of Spark Plasma-Sintered Zr0,-ZrB,

Nanoceramic Composites

455

B. Basu, T. Venkateswaran, and D-Y Kim J. Am. Ceram. SOC., Vol. 89, No. 8, p. 2405-2412,2006

Homogeneous Zr0,-Al,O,

Composite Prepared by Nano-ZrO, Particle Multilayer-CoatedAI,O,

Particles

463

Y. Jia, Y. Hotta, K. Sato, and K. Watari J. Am. Ceram. SOC.,Vol. 89, No. 3, p. 1103-1106, 2006

Preparation of a Highly Conductive AI,O,/TiN Growth

lnterlayer Nanocomposite through Selective Matrix Grain

467

X. Jin and L. Gao J. Am. Ceram. SOC., Vol. 89, No. 3, p. 1129-1132, 2006

Preparation and Microstructure of Multi-Wall Carbon Nanotubes-ToughenedAI,O,

Composite

471

J. Fan, D. Zhao, M. Wu, Z. Xu, and J. Song J. Am. Ceram. SOC., Vol. 89, No. 2, p. 750-753,2006

Three-DimensionalAssemblies of Zirconia Nanocrystals Via Shape-Preserving Reactive Conversion of Diatom Microshells

475

S. Shian, Y. Cai, M. Weatherspoon, S. Allan, and K. Sandhage J. Am. Ceram. SOC., Vol. 89, No. 2, p. 694-698,2006

Contents

ix

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Introduction

Although nanotechnology is still an emerging industry, it represents a huge potential in a variety of markets that include biomedical, electronics, and energy totaling billions of dollars. However, before these markets are realized, processing methods must be developed that can produce quality nanomaterials and structures. Whether the material is a powder, thin film, wire, or composite, an optimal processing method is needed. Powders of various compositions can be made by a wide range of methods, including freeze casting, chemical, hydrothermal synthesis, and solution combustion, among others. Each method has its limitations and advantages. The methods to make thin films and coatings include chemical vapor deposition, spray pyrolysis, and sol gel. Wires can be made by electrospinning or hydrothermal synthesis. Other methods are under development for making composites and other structures. This edition of Progress in Ceramic Technology series contains a select compilation of articles on the topic of nanomaterials processing of powders; thin films, wires and tubes; and composites that were previously published in The American Ceramic Society Bulletin, Journal of the American Ceramic Society, International Journal ofApplied Ceramic Technology, Ceramic Engineering and Science Proceedings (CESP) and Ceramic Transactions (CT). The American Ceramic Society contributes to the progress of nanotechnology by providing forums for information exchange during its various meetings and by publishing articles in its various journals and proceedings. For other books on nanotechnology, including Progress in Nanotechnology: Applications, visit the ACerS bookstore at www.ceramics.org or the ACerS-Wiley webpage at www.wi1ey.comJgolceramic.s.

Introduction

xi

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Synthesis Methods for Powders

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Freeze Casting as a Nanoparticle Material-Forming

Method

Kathy Lu*’+and Xiaojing Zhu* Materials Science and Engineering Department, Virginia Polytechnic Institute and State Universiy, Blacksburg, Virginia 24061

Nanoparticle material forming is challenging because of loose packing and agglomeration issues intrinsic to nanoparticles. Liquid processing shows great potential to overcome such hurdles. This study is focused on nanoparticle colloidal processing and freeze-casting forming. A l 2 O 3 nanoparticle suspensions are examined, and microstructure evolution of A203nanoparticle suspension during freeze casting is discussed. The “Fines” effect influences nanoparticle packing on freeze-cast sample surfaces. Trapped air bubbles in the suspension lead to a porous bulk microstructure. Prerest is necessary for dense and homogeneous green microstructure formation. The green strength, fracture mode, and ability to form fine features by freeze casting are also evaluated.

Introduction The huge surface area intrinsic to nanoparticles serves as one of the most striking advantages as well as disadvantages for nanoparticle material forming. Because of the natural tendency of nanoparticles in forming agglomerates, wet forming utilizing a colloidal suspension has become the most active research area in nanoceramics. For almost all the nanoparticle wet forming processes, the first step is to produce a stable colloidal suspension. A green body with a uniform

Supported by Petroleum Research Fund, administered through American Chemical Society. Presented at the 31st International Conference on Advanced Ceramics and Composites, Dayrona Beach, FL, January 21-26, 2007. ‘Member, The American Ceramic Society. ‘Wu&t.edu 02008 The American Ceramic Society

Synthesis Methods for Powders

microstructure can then be produced from the h l l y stabilized colloidal suspension of nanoparticles. Freeze casting is a process that pours the suspension into a nonporous mold, freezes the suspension, demolds the frozen sample, and then dries the sample under vacuum. Liquid-state to solid-state conversion is realized through phase transformation of the dispersing medium such as water. The process has the potential to form near net-shape complex geometry parts with low pressure and often environmentally benign advantages. l T 2 The key requirements are that the suspension is stable and ~nagglomerated.~ When the freeze-casting condition is properly controlled, water separates from solid phases through sublimation and no capillary force exists to cause hard agglomerates or cracks.* Freeze casting has been practiced for sometime. 1,2,5,6 However, the studies were mainly focused on micrometer-sized particles and large-sized samples. For example, A 1 2 0 3 particles of 0.4 pm were freeze cast

3

and > 98% sintered density was achieved.’ Enclosed shells of Al2O3 bodies encapsulating steel parts were fabricated.8 Porous and layered-hybrid materials were freeze cast.’ However, application of freeze casting to nanoparticle systems and understanding the freezecasting behaviors of nanoparticle suspensions have not been explored. One challenge is that nanoparticle research in forming bulk components is still evolving; considerable effort is still needed in achieving high solids loading suspensions. The other challenge is the drastic capillary pressure increase with decreasing particle size; it takes a prolonged duration of time and much improved control to remove water completely by sublimation. Our past work has addressed the first challenge to a certain extent. 10-12 To address the second challenge successfully, it is essential to understand the nanoparticle microstructure evolution from the colloidal state to the freeze-dried state. High particle packing efficiency is preferred for postfreeze casting, handling, and densification purposes. Freeze-cast sample strength and failure mode can be used to understand and evaluate this aspect. With the decrease of particle size to nanoscale, the feature sizes of nanoparticle samples can also be correspondingly decreased. The potential in forming fine features by freeze casting should be explored. This study is focused on A 1 2 0 3 nanoparticle suspension viscosity evaluation and the corresponding freeze casting process for solid sample formation. Microstructural evolution of nanoparticle suspensions from a colloidal state to a freeze-dried state is examined. The surface and bulk microstructures under different prerest conditions are compared. Special nanoparticle phenomena during the suspension to solid-state transformation are explained. Equibiaxial strength and fracture mode of the freeze-cast samples are examined. Fine features produced by the freeze-casting process are presented.

Experimental Procedure A1203 dry nanoparticles with a specific surface area of 45 m2/g were used in this study (Nanophase Technologies, Romeoville, IL). The particles can be redispersed in water by ball milling as reported before by transmission electron microscopy (TEM) analysis. l o In this study, the particle size distribution was further measured by a dynamic light-scattering measurement

4

Fig. I . A1203 nanoparticle size distribution porn dynamic light scattering analysis: (a) weight basis, (b) number basis. as shown in Fig. 1 (Zetasizer Nan0 ZS, Malvern Instruments Inc., Southborough, MA). Agglomerates are absent. However, A203 nanoparticles are polydispersed, consistent with the TEM analysis. Even though on the weight basis there seems to be 10 wt% of > 100 nm particles, on the number basis, large particles are negligible (<0.5%). For the A 1 2 0 3 nanoparticle suspension preparation, poly(acry1ic acid) (PAA, Mw 1800, Aldrich, St. Louis, MO) was used as the dispersant with the segment as [-CH2CH(C02H)-] and glycerol (C3H803, Fisher Chemicals, Fairlawn, NJ) was used as the freeze casting aid with the molecular formula CH20H-CHOH-CH20H. A water-glycerol mixture at a weight ratio of 10:1 (water:glycerol) was used as the dispersing medium. The mixture was homogenized for 5 min using a ball mill (755 RMV, US Stoneware, East Palestine, OH) before use, and the ball mill rotation speed was 75 rpm. A203 nanoparticles were added to the dispersing medium in 10 g increments along with a PAA dispersant at 2.0wt% of A1203. Because low pH

-

Progress in Nanotechnology: Processing

promotes PAA dispersant adsorption onto A 1 2 0 3 nanoparticle^,'^ HCl solution was added to lower the pH to 1.5. The suspension was ball milled for 12 h with periodic adjustment of pH to 1.5. The purpose of the ball milling was to break the soft agglomerates of the dry A1203 nanoparticles and thoroughly homogenize water, glycerol, PAA, and A 1 2 0 3 nanoparticles. The milling media were A1203 rods with 6 mm diameter and 6 mm height. Breaking down of A 1 2 0 3 nanoparticles into smaller sizes was believed to be absent because of the low energy that ball milling provided. However, air bubbles can be introduced into the suspension during mixing. Suspensions of approximately 20 vol% A 1 2 0 3 solids loading were made by this procedure. After this step, A 1 2 0 3 nanoparticles were again added in l o g increments, along with a PAA dispersant at 2.Owt% of A 1 2 0 3 in order to make a 4 0 ~ 0 1 %solids loading suspension. The final PAA amount was thus maintained at 2.0wt% of A l 2 O 3 . N H 4 0 H (28-30wt%, W R International, West Chester, PA) was used to adjust the suspension to p H 9.5 in order to achieve the highest suspension flowability for freeze casting.l o It should be pointed out that the liquid volume of N H 4 0 H and HC1 was considered for the 4 0 ~ 0 1 %solids loading reported. The suspension was then mixed for 24 h for complete homogenization. Suspension viscosities without and with glycerol were characterized using a rheometer with a cone-plate geometry (AR 2000, TA Instruments, New Castle, DE). All the viscosity measurements were performed with a controlled shear rate in a decreasing shear rate order. Freeze-casting molds were developed using poly(dimethylsiloxane) epoxy (RTV 664, General Electric Company, Waterford, NY). An A1203 suspension was filled into the mold immediately after the suspension preparation. Some of the filled molds were kept under ambient conditions for 1 h; other filled molds were frozen in a freeze dryer (Labconco Stoppering Tray Dryer, Labconco, Kansas City, MO) immediately. This 1-h resting period after the mold filling but before the freeze-casting step is termed as prerest in the following discussion. During the freezing process, the suspension-filled molds were kept at -35°C for 2 h . Then, the samples were demolded and exposed to vacuum (< 1 x lo4Pa) to allow ice to sublimate for 36h. The microstructural differences in the freeze-cast samples with and without prerest were systematicallyanalyzed using a LEO550 field emission scanning electron microscope (SEM) (Carl Zeiss MicroImaging Inc., Thornwood, NY).

Synthesis Methods for Powders

For the freeze-dried sample green strength measurement, circular disks 25.4mm in diameter and 1.58 mm in thickness were freeze cast. The A 1 2 0 3 suspension was placed in the mold cavity using a disposable pipette. Care was taken to fill the molds completely with the A1203 suspension. After the mold filling, the samples were allowed to rest for 1 h under ambient conditions. The fracture load was measured using a Texture Analyzer test console (Stable Micro Systems, Surrey, U.K.), equipped with a 5 kN load cell. The equibiaxial flexural strength was then calculated for each sample, using the following equation:

where ofis the fracture strength in units of MPa, F is the fracture load in units of N, h is the specimen thickness in units of mm, v is Poisson's ratio, D is the sample diameter in mm, Ds is the diameter of the support ring in mm, and DL is the diameter of the load ring in mm.

Results and Discussion

A1203 Nanoparticle Suspensions In the previous study," it was identified that for the nanoparticle suspension studied, the PAA amount adsorbed onto A 1 2 0 3 particles increases with PAA concentration increase up to 2.00 wt% of A l 2 O 3 . As more PAA is added to the suspension, the PAA amount adsorbed reaches a plateau of 0.31 mg/m2. The adsorption plateau represents the adsorption saturation at which the A 1 2 0 3 particle surfaces are fully covered by the PAA dispersant. Also, the suspension has the lowest viscosity from pH 7.5 to 9.5; this pH range is the most desirable p H range for stabilizing the A 1 2 0 3 nanoparticles. Based on these results, the PAA concentration is 2.00wt% of A 1 2 0 3 and the pH is 9.5 for the A 1 2 0 3 nanoparticle suspensions in this study. Figure 2 shows the viscosities for the 20-45vol% A1203 solids loading suspensions. The viscosities are a strong function of A 1 2 0 3 nanoparticle solids loading. With an increase in A 1 2 0 3 solids loading, the viscosity increases at all shear rates. Approximately, the suspension viscosity increases by two orders of magnitude from 20 to 45 vol% solids loading. As the shear rate increases,

A1203

5

Prerest Efect

Fig. 2. Kscosity change versus shear ratefor dtfferent AlZO, solids loading suspensions. Each curve is an average of three measurements and the standzrd deviation is much smaller than the markers and cannot be shown on thejgure.

the viscosity decreases monotonically but the Al2O3 nanoparticle suspension does not obey Newtonian flow. Instead, the suspension displays a shear-thinning behavior, reflecting the broken polymeric link between the nanoparticles provided by the PAA dispersant. At 20 vol% solids loading, the viscosity changes from 0.03 to 0.02 Pas when the shear rate is changed from 30 to 200 s-’. At 45 vol% solids loading, the viscosity changes from 4.35 to 1.60 Pas when the shear rate is changed from 30 to 200 s-’. As the A1203 solids loading increases, the viscosity difference becomes much larger at high solids loading levels. This indicates that the suspension undergoes a more drastic “link breakage” and that the suspension conditions should be controlled more rigorously and optimized under high solids loading conditions. There are two states for a polydispersed nanoparticle suspension like the A1203 suspension studied here: fluid and glass, depending on solids loading.’* For freeze casting, especially when forming parts with complex shapes, the suspension should be kept at reasonably high solids loading levels. When the solids loading is too low, the freeze-cast sample will have a very low particle packing density and present a challenge for densification. However, when the solids loading is too high, the suspension will be too viscous to flow into the fine features of the mold. Ideally, the solids loading should be chosen at the boundary of fluid and glass suspensions, slightly on the fluid-state side. Our prior study has identified that 45% solids loading is approximately near the fluid-state to glass-state boundary region.15 Based on this consideration, the suspension of 40 vol% A 1 2 0 3 solids loading was used for the freeze-casting studies reported here.

6

In the previous work,” the critical role of prerest in forming homogenous and crack-free freeze-cast samples was reported. However, the microstructural changes were not understood. In this study, the prerest effects on the freeze-cast sample microstructure evolution are examined from two aspects: surface microstructure and bulk microstructure. The free surface microstructures of the freeze-cast samples without and with prerest are shown in Fig. 3. For the sample without prerest (Fig. 3a), the surface is rather flat and the A1203 particles are densely packed. For the sample with prerest (Fig. 3b), the surface is rough and the particles are loosely packed. This surface smoothness difference can be observed indirectly from

Fig. 3. &$ace microstructure differencefor the fieeze-cast samples: (a) without prerest, (6) with prerest.

Progress in Nanotechnology: Processing

the surface luster difference. For the sample in Fig. 3a, the surface is shiny. For the sample in Fig. 3b, the surface is dull. The unique nature of the freeze casting process is able to capture this evolution that otherwise would be missed unless in situ analysis is used.16 Figure 3 reveals phenomena that are in sharp contrast with the well-established understanding at the micrometer particle size level. Even though the gravity effect seems to be a convenient explanation, it can be eliminated because the suspension is stabilized to prevent sedimentation and the particle size is too small for gravity to play an essential role within the timeframe studied. Lue and Woodcock” generalized this particle self-arrangement process as the “fines” effect and showed that the large to small particle size ratio for the “fines” effect to be effective is 10:l. The “fines” effect refers to the natural tendency of the small particles migrating to fill the space between the large particles under Brownian motion when the suspension is kept undisturbed. From Fig. 1, the d203 nanoparticle suspension certainly meets this particle size criterion. During the undisturbed prerest time period, Brownian motion causes nanoparticles to move within the suspension. When the small nanoparticles migrate to fill the voids between the large nanoparticles, the process causes the large particles on the suspension surface to open up the network to accommodate the highly diffusive small particles. The net result is a more open and less smooth surface microstructure. This observation clearly indicates that Brownian motion is an important phenomenon affecting nanoparticle material surface microstructures. From the freeze-cast surface microstructure comparison, it can be concluded that nanoparticle size distribution plays an important role when the suspension is directly frozen, even if the large particles only take up a small percent of the overall distribution. The particle size distribution should be kept at the narrowest range for suspension-related forming processes if a uniform and consistent surface microstructure is desired. The fracture surface microstructures of the freezecast A 1 2 0 3 samples are shown in Fig. 4. Figure 4a shows the sample without prerest. Figure 4b shows the sample with prerest. Both samples are formed under the same freeze-casting conditions. For the sample without prerest (Fig. 4a), a porous bulk microstructure develops and the pores are 3 pm in size through SEM image measurement of about 50 pores. For the sample with prerest, no micrometer-sized pores are present as shown in Fig. 4b. For the A 1 2 0 3 nanoparticle suspension preparation,

-

Synthesis Methods for Powders

Fig. 4. Bulk microstructure difference fDr the freeze-cast samples: (a) without prerest, (b) with prerest.

a high shear process, ball milling, was used to homogenize A 1 2 0 3 nanoparticles, dispersant PAA, glycerol, and water. A shear force was constantly present to break the nanoparticle suspension and encapsulate air bubbles while homogenizing the suspension. When the shear ceases, the d203 nanoparticle suspension should be given time to reorganize for the air bubbles to migrate out of the suspension. However, it may take a long time to eliminate the air bubbles completely for the high solids loading suspensions due to the relatively higher suspension viscosity and more particle movement hindrance. If the nanoparticle suspension is frozen before such an equilibrium is reached, then observations are often made for a metastable suspension with air bubbles dispersed in the suspension.

7

Because the samples without and with prerest are freeze cast under the same conditions, the freezing process of the water-glycerol dispersing medium and the subsequent impact on the nanoparticles should not play a role in the microstructure difference observed. More revealing evidence can be found in Fig. 5. For the sample without prerest, the pore surfaces are very smooth. The elongated pores show no fracture on some ligaments (Fig. 5a). This means that these elongated pores are not dendrites, which frequently occur during an aqueous suspension freezing (phase transformation) process, even though the shape has some resemblance to dendrites. The left edge of Fig. 5b shows the smooth pores at a high magnification. This work offers the first evidence of the importance of microstructure evolution

before freeze casting. Figures 4 and 5 indicate that the porous structure is a consequence of lack of prerest; the air bubbles trapped in the suspension are the cause for the micrometer-sized pores. After nanoparticle suspension mixing and freeze-cast mold filling, enough time should be given for the trapped air to escape and for the A 1 2 0 3 nanoparticles to rebuild the 3D network. The more open surface microstructure due to the "fines" effect (Fig. 3) seems to be beneficial for the air bubbles to escape. In essence, this means that the surface and bulk microstructures evolve simultaneously to affect the microstructure of the freeze-cast samples.

Equibiaxial strength, Fracture Mode, and Fine Features The freeze-cast disk samples with prerest were fractured using the equibiaxial strength ring test, and the peak load was recorded to determine the strength." Compared with the four-point bending test, the equibiaxial strength test has the advantage of isotropic test conditions, whereas the four-point bending test exerts high stress in the direction perpendicular to its supports. l 9 The equibiaxial strength test was also conceived to eliminate the effect of edge (corner) defects that might be present for a four-point bending test specimen.20 Accordingly, the equibiaxial flexural strength test should offer better evaluation of the freeze-cast samples, especially in light of the brittle failure nature. The equibiaxial flexural strength of the freeze-cast samples was calculated via Eq. (1). The average freeze-cast sample strength is 0.3-1.6MPa. It should be mentioned that this green strength variation is related to the sample thickness. When the sample is thick, the green strength tends to be lower. This is a surprising observation for the freeze-cast sample because the green strength should be dimension independent. Even though Poisson's ratio v is not directly available and is assumed to be 0.22 in this study, a handbook value for fully dense A1203r21a deviation of 25% in the Poisson's ratio results in only an 2% change in the equibiaxial flexural strength and the calculation is made with this knowledge. We believe that the green strength variation is closely related to the microstructure evolution discussed earlier. Complete removal of air bubbles from the A 1 2 0 3 nanoparticle suspension is essential to obtain consistent and highstrength freeze-cast samples. If air bubbles are not removed completely, which is more likely for the thicker samples, the resulting pores can act as defects and

-

Fig.5. Smooth pore surfaces offieeze-cast sample without prerest: (a) elongated pores, (b) spherical pores.

a

-

Progress in Nanotechnology: Processing

cause low green strength. More problematically, these defects can cause inhomogeneous densification and permanent flaws in the sintering process. It is possible that an optimal prerest time exists for each sample thickness. Future efforts will be made to evaluate this aspect. Even though the green strength is low compared with that of a fully dense sample, the result is very encouraging. The freeze-cast samples are fairly strong compared with green samples from other forming processes and can be handled easily for sintering densification. As to the fracture characteristics, the freeze-cast A 1 2 0 3 samples conform to the medium energy fracture mode as shown in Fig. 6. The marker indicates the origin of the cracks. It has been known that the green sample fracture strength and the fracture mode are correlated. The medium fracture mode means that the freeze-cast sample has desirable integrity and consumes a reasonable amount of energy during fracture. If the samples are too weak, the fracture mode will change correspondingly to the low energy fracture mode. Nevertheless, the freeze-cast samples do not have high strength, such as for sintered ceramic or glass samples. It should also be pointed out that the fracture edges of the freeze-cast samples are very smooth. The fractured pieces can be re-assembled easily into its original shape.

Nanoparticle materials have great potential in producing decreasing size features and parts because of the size decrease in the basic constituent species-particles. However, a substantial challenge exists in producing well-defined edges and dimensions.22 Freeze casting seems to be a good candidate in addressing this issue and forming complex shapes. This advantage is demonstrated by freeze casting of a dime sample as shown in Fig. 7. Freeze casting has high fidelity in feature shape replication. The fine features on the dime face and the sharp edges of the dime are precisely reproduced. This represents a very promising opportunity of forming complex-shaped nanoparticle materials with micrometer-sized features. T o further explore the freeze-casting capability in forming small-size features, the 40 vol% A 1 2 0 3 solids loading suspension was freeze cast into small molds with micrometer-sized features. As shown in Fig. 8a, the micrometer-sized features are precisely aligned with a well-defined shape and size. Figure 8b shows an even smaller feature size with 100 pm bottom diameter and < 40 pm height. It should be mentioned that this length bridges the macroscopic scale and the nanometer scale and falls in the range of many component sizes in device fabrication. Freeze-casting forming can be potentially

F&. I; Fracture modes: (a) ideal fiacture modes, and (6)fieeze-cast A1203 samplefiacture images. Actual sample crack initiation point is indicated a star.

Synthesis Methods for Powders

9

Fig. 8. Micrometer-sized features produced by fieeze casting of A1203 nanoparticle suspension: (a) aligned features, (6) individual

feature. Fig. 7. The dime and the fine features produced by A1203 nanoparticle suspension I;.... casting: (a) fill dime, (6) lower left corner of the dime face.

used to form complex surface features for direct device fabrication purposes. To fully explore the potential of freeze casting in forming components in the micrometer length scale, two issues need to be solved. One is creating dimensionally well-controlled molds for fine feature formation during nanoparticle freeze casting. As shown in Fig. 8b, less than desired smoothness at the mold cavity bottom produces micrometer-sized features with an uneven top surface. Second, fine features tend to create stress, concentrating at the feature surrounding areas. It requires great care to control the freeze drying process to obtain intact samples. Even though the cracks shown in Fig. 8a

10

are not observed after freeze drying and is believed to be from SEM sample preparation, freeze casting of intricate features is still at its infancy. Future work will be continued to improve the freeze casting process for complex-shaped micrometer-sized sample forming.

Conclusions Nanoparticle suspension and freeze-casting forming are interesting new areas in nanoparticle processing. This study focuses on an A 1 2 0 3 nanoparticle suspension and the subsequent freeze casting processes. It shows that suspension viscosities increase with A1203 nanoparticle solids loading. After the suspension is cast into a mold, prerest should be given for the nanoparticles to

Progress in Nanotechnology: Processing

rearrange. Otherwise, the air bubbles will be trapped in the bulk and evolve into pores after freeze casting. The “fines” effect was found to be beneficial for the surface nanoparticle packing to open up and for the trapped air to escape from the A 1 2 0 3 nanoparticle suspension. Dense nanostructures can be obtained when the freeze casting conditions are controlled properly. The freezecast samples have good green strength and show a medium energy fracture mode during the equibiaxial strength testing. The freeze-casting process itself shows to be a promising approach in forming complex shapes with fine features. Acknowledgments

The authors thank Prof. Brian Love for help with the equipment use to conduct the equibiaxial flexural strength testing. References 1.

2. 3.

4.

5.

D. Koch, L. Andresen, T. Schmedders, and G. Grathwohl, “Evolution of Porosity by Freeze Casting and Sintering of Sol-Gel Derived Ceramics,” j. Sol-MSci. Technol., 26 11-31 149-152 (2003). R. W. Jones, “Near Net Shape Ceramics by Freeze Casting,” Ind. Ceram., 20 [2] 117-120 (2000). J. Laurie, C. M. Bagnall, B. Harris, R. W. Jones, R. G. Cooke, R. S. RussellFloyd, T. H. Wang, and F. W. Hammett, “Colloidal Suspensions for the Preparation of Ceramics by a Freeze Casting Route,” j. Non-Cyst. Solih, 147 and 148 320-325 (1992). B. E. Novich, C. A. Sundback, and R. W. Adams, “Quickset Injection Molding of High-Performance Ceramics,” Ceramic Transactions, Vol. 26, Forming Science and Technology f i r Ceramics. ed., M. J. Cima. American Ceramic Society, Westerville, OH, 157-164, 1992. S. W. Sofie and F. Dogan, “Freeze Casting ofAqueous Alumina Slurries with Glycerol,”j. Am. Ceram. Soc., 84 [7] 1459-1464 (2001).

Synthesis Methods for Powders

6. Z. S. Rak, “Advanced Shaping Techniques in Advanced Ceramics,” Ceram. Forum Int., 77 [7] E25-E26 (2000). 7. K. Araki and J. W. Halloran, “Room-Temperature Freeze Casting for Ceramics with Nonaqueous Sublimable Vehicles in the NaphthaleneCamphor Eutectic System,”]. Am. Ceram. Soc., 87 (111 2014-2019 (2004). 8. T. Moria and H.-J. Richter, “Ceramic Bodies with Complex Geometries and Ceramic Shells by Freeze Casting Using Ice as Mold Material,” j . Am. Ceram. Soc., 89 [8] 2394-2398 (2006). 9. S. Deville, E. Saiz, R. K. Nalla, and A. P. Tomsia, “Freezing as a Path to Build Complex Composites,” Science, 31 1 [5760] 515-5 18 (2006). 10 K. Lu, C. S. Kessler, and R. M. Davis, “Optimization of a Nanoparticle Suspension for Freeze Casting,” /. Am. Ceram. Soc., 89 [8] 2459-2465 (2006). 11 K. Lu and C. S. Kessler, “Nanoparticle Colloidal Suspension Optimization and Frem-Cast Forming,” Ceramic Engineering and Science Proceedingr, Vol. 27. American Ceramic Society, Westerville, OH, 1-10, 2006. 12 K. Lu, “Microstructural Evolution of Nanoparticle Aqueous Colloidal Suspensions During Freeze Casting,” j . Am. Ceram. Soc., 90 3753-3758 (2007). 13. J. Cesarano and 1. A. Aksay, “Processing of Highly Concentrated Aqueous Alpha-Alumina SuspensionsStabilized with Poly-Electrolytes,”j. Am. Ceram. SOC.,71 (121 1062-1067 (1988). 14. F. Sciortino and P. Tarraglip, ‘‘Glassy Colloidal Systems,” Adv. Phys., 54 [ 6 7 ] 471-524 (2005). 15. K. Lu and C. S. Kessler, “Colloidal Dispersion and Rheology Study of Nanoparticles,” j . Mater. Sci., 41 [17] 5613-5618 (2006). 16. S. Hachisu, “Segregation Phenomena in Monodisperse Colloids,” j. Colloid Interface Sci., 55 [3] 499-509 (1976). 17. L. Lue and L. V. Woodcock, “Depletion and “Fines” Effects on Transport in Hard-Sphere Colloids,” Int. J Thermophys., 23 [4] 937-947 (2002). 18. ASTM Designation C 1499-04, Standard Test Methodfbr Monotonic Equibiaria1 Fluural Strength ofAduanced Ceramics at Ambient Temperature, American Society for Testing and Materials International, West Conshocken, PA, 767-777, 2005. 19. ASTM Designation C1161-02c. Standard Test Methodfir Flexuralstrength of Advanced Ceramicsat Ambient Temperature,American Society for Testing and Materials International, Wesr Conshocken, PA, 223-236, 2003. 20. U. Soltesz, H. Richter, and R. Kienzler, “The Concentric-Ring Test and its Application for Determining the Surface Strength of Ceramics,” Ceramics in Clinical Applications (High Tech Ceramics), ed., P. Vincenzini. Elsevier, Amsterdam, 149-158, 1987. 21. N. Dowling, Mechanical Echauior of Materiak, 2nd Edition, Prentice-Hall, Upper Saddle River, NJ, 178, 1999. 22. G. Hayes, N. Antolino, M. Frecker, C. Muhlstein, E. Mockensturm, and J. Adair, “Processing of Sub-500 Micron Scale Ceramic Parts for Surgical Applications,” Proceedings of Materiak Science and Technology 2006 ed., E. Medvedovski. Cincinnati, OH, 901-910, 2006.

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Preparation of a Nanoscale/SOFC-Grade Yttria-Stabilized Zirconia Material: A Quasi-Optimization of the Hydrothermal Coprecipitation Process Yang-Chuang Chang, Maw-Chwain Lee,* Wei-Xin Kao, and Tai-Nan Lin The Institute of Nuclear Energy Research, W C , Lung-Tan, 325 Taiwan, ROC

Yttria-stabilized zirconia (YSZ) is the key material for an electrolyte of a solid oxide fuel cell (SOFC). In order to prepare a nanoscalelSOFC-grade 8YSZ material, the hydrothermal coprecipitation process is intensively investigated for process improvement and product identification. From the characterization results of the 8YSZ product, the operation conditions were selected for process optimization. The criteria used in optimization condition determination were the reproducibility of the 8YSZ product and the particle size, the crystallinity, as well as the operation simplicity. Experimental results showed that a nanoscale/SOFC-grade 8YSZ powder was directly prepared. The quasi-optimum conditions of the process are proposed.

Introduction

8YSZ is a well-recognized material used as an electrolyte of a solid oxide fuel cell (SOFC). It consists of 8 mol% of Y203 and 92 mol% of Zr02 (zirconia). The basic requirements of yttria-stabilized zirconia (YSZ) for a ood electrolyte are (1) to have good media for pure 0 transport, (2) to have zero electronic conduction, (3) to be porosity free to prevent gas from permeating, and (4) to be thin to minimize ohmic 1 0 ~ s . The ’ ~ ~ typical specifications of a commercial product of YSZ material and SOFC development at the Institute of Nuclear Energy Research (INER) were well summari~ed.~’* It shows that the SOFC-grade YSZ material has characteristics of low particle size (nanoscale for electrolyte and submicrometer scale for an anode) and high thermal stability (100% cubic crystal structure). The

5-

‘[email protected]

0 2008 The American Ceramic Society

Synthesis Methods for Powders

cost of such a YSZ material is expensive. Seabaugh et reported the cost analysis for the manufacturing cost of YSZ material in comparison with SOFC. The material cost is approximately 60% of the stack cost. Hence, selfdevelopment of the SOFC-related materials, especially for high-quality material (e.g., YSZ), is essential for SOFC fabrication. This tendency provides INER with the objective of preparing SOFC-grade YSZ efficiently. In addition to the factor of fabrication cost, the flexibility for the fabrication process and product quality can be tuned by varying the powder characteristic/property. Hence, a study of the preparation of SOFC-grade YSZ powders is critical for SOFC development. Bagger reported that SOFC cells based on a 180-pm-thick cubic YSZ electrolyte were produced by a combination of tape casting and spray coating at sintering temperatures around 1300°C; an area-specific internal cell resistance of 0.3Rcm2 at 1000°C was obtained with controlled electrode structures; and anodes were based on Ni-YSZ cermets with only 40 vol% Ni.‘ Composite cathodes

13

were developed from Sr-doped LaMn03 and YSZ with 3 mol% yttria. YSZ is the key material for SOFC membrane electrode assembly (SOFC-MEA). Both physical Feed Stock Prepwdtion Prccipitants and chemical processes are generally used for prepara1 f a (UrcdAniinonin) Reactor tion of SOFC-grade YSZ material. For the physical .Batchwise reactor method, 8YSZ powders are well mixed and blended in a *Pressure(<22.0 MPa=Pc,, closed system (e.g., a jar or a container) via operation *Residence time (20-80 inii (either in solid or in liquid phases) of ball mills using Teinperdture(2s "C) Cooling zirconia media to prepare 8YSZ powders. A homogeFiltration nous mixture of Y 2 0 3 + Z r 0 2 powder is obtained after long-time blending. The final mixture is subjected to calcination/sintering at a suitable temperature (near 1500°C) to obtain 8YSZ powder with a 100% cubic crystal structure, which is suitable for the preparation of an SOFC electrolyte. This is the simplest way to achieve Fig. 1. Block diagram of the hydrothermal coprecipitation/ the target of production of SOFC-grade YSZ p ~ w d e r . ~ ' ~ synthesis process for yttria-stabilized zirconia (XSZ) preparation. YSZ powder obtained from this process possess a larger particle size, and a further milling process (either in solid or in liquid phase) is required to obtain a fine YSZ powder. The limiting value of particle size is around reports were provided by Swartz et aL2' and Seabaugh 0.5 pm via a solid-phase milling process using a 5 mm et aL5 in this regard. The hydrothermal coprecipitation/ synthesis process developed at INER is shown in Fig. 1 . YSZ ball.9 The processes to prepare ceramic powders have been cited widely.'-' Advanced milling equipA precursor solution based on a mixture of zirconium ments for preparation of nanoscale ceramic powder have and yttrium salts (e.g., ZrO(N03)2/ZrOC12 and been announced by Hsin Fang Co. of Taiwan, NetzschY(N03)3/YC13) was prepared according to the stoiFinmahltechnik Gmbh, and other organizations. chiometry of the required (Zr02)0.92(Y203)0.08 powA lower sintering temperature not only improves der. A suitable dispersant (e.g., polyethyene glycol, citric the flatness of anode-supported SOFCs but also reduces acid, phosphate ester, etc.) may be added to the precursor solution for prevention of particle coagulation the grain growth of Ni phase in the NiNSZ anode subduring the coprecipitation process. pH adjustment is strate. Chen etal.16 found that pellets made from nanosize YSZ powders can achieve 94% of the theoretical also one of the key methods to keep particles apart. density at 1200°C. Many studies have been proposed on Hence, the amount of precipitant (base) must be limited. Dispersants are designed to produce fine (either diverse manufacturing techniques for preparation of nanosize YSZ powders, e. . gas-phase combustion," nanoscale or submicrometer scale) YSZ powders. After completing feedstock preparation, the precipitant (e.g., spray drying,18 ~ol-gel,'~,~'hydr~thermal;~-~~ solvourea (CO(NHJ2) or ammonia (NH3/NH40H)) is thermal,26 and so on. Among these techniques, the hydrothermal process has many advantages such as there added to the reactor for a coprecipitation reaction. being no solvent, no sintering process, and being good The temperature of the reactor was first increased up chemical composition control. In addition, continuous to 80°C for coprecipiration and held for a specified peprocessing systems of hydrothermal reaction are very riod of time to facilitate completion of the reaction. The promising.27 In this paper, the hydrothermal coprecipreactor temperature was further increased up to 180itation process is developed and the experimental results 330°C and held for another specified period of time for are presented for process optimization. synthesis of YSZ oxide. After completion of YSZ oxide synthesis, the temperature and pressure were lowered via cooling through the jacket outside the reactor to the Description of the Hydrothermal Coprecipitation ambient condition. Then, the product was subjected to Process settling, filtration, washing, and drying. The YSZ powder thus obtained can be a nanoscale particle. EssentialDawson studied the hydrothermal synthesis process ly, the mechanisms of the hydrothermal coprecipitation/ for preparation of advanced ceramic powders.28 Many

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14

Progress in Nanotechnology: Processing

synthesis process for YSZ preparation are deduced. It involves the following chemical reaction^:^-^^,^' (1) Dissociation of NH40Hlurea and metal salts:

NH40H(aq>+rNH:(aq)+ OH-(aq)

+

ZrO(N03),(aq) +rZr02+(aq) 2NO;(aq)

(1)

(3)

Experimental Procedure Preparation of Feedstock- Y(N03)3 Yttria (Y203) was slowly dissolved in concentrated nitric acid ( H N 0 3 ) solution under a stirring condition. The free acid in the final stock is kept as low as possible. The concentrations of yttrium and free acid were determined by chemical analysis. The typical concentrations of feedstock are 218 g/L and 0.002 M for yttrium and free acid, respectively. Preparation of Feedstock- YSZ Precursor

(2) Precipitation reactions:

+

Zr02+(aq) 4NH3 . H20(aq)(orNH40H(aq)) +Zr(OH),(s) -1 +4NH:(ag)

+O2-(q) (5)

+

Y3+(aq) 3NH3 . HzO(aq)(orNH40H(aq)) (7) +rY(OH),(s) 1 +3NHf(aq)

(3) Thermal formation:

decomposition

Zr(OH)4(s)

+

reactions9/crystal

ZrOz(s) + 2H20(g) 1

(9)

A YSZ precursor solution based on a mixture of zirconium and yttrium salts (i.e., ZrO(N03)2 and Y(N03)3 ) was prepared according to the stoichiometry of the required (Zr02)0.92(Y203)0.08 powder. Solid Z I - O ( N O ~ ) ~was ( ~ )added to the Y(N03)3(aq) solution. Hydrothermal Coprecipitation/Synthesis The experimental procedure for the hydrothermal coprecipitation/synthesis of YSZ is briefly described in “Description of Hydrothermal Coprecipitation Process.” Temperature control and pressure detection are key factors in the whole process. Industrial safety must be well handled. The capacity of a Pressure Product (Warminster, PA) agitator reactor is 4L with a magnetic drive mixer. The maximum operating (design) temperature and pressure are 350°C (650°F) and 41.4MPa, respectively. Other equipments are a solid-liquid separator (filtration system via centrifuge), a high-temperature furnace for calcinations and sintering, and a ball-milling system (either in liquid or in solid phase milling). After filtration, drying, and calcinations of the YSZ precipitate from the hydrothermal reactor, 8YSZ powders with nanoscale or submicrometer scale particles were obtained. Powder Characterization

(4) Hydrothermal synthesis of the YSZ Zr(OH),(s) + Y(OH),(s) - dissolution Zr(OH)*(aq)+ Y ( O H ) 3 ( 4 - precipitation +

+

ysz Or [ ( Zr 0 2 ) ~.~2 ( Y2 0 3 >o .o sl

(12)

Calcination and sintering of the YSZ precipitate are not required to obtain SOFC-grade 8YSZ powder for SOFC-MEA fabrication, while the process parameters are approximately optimized.

Synthesis Methods for Powders

The characterization of powder properties is very important for process development and evaluation. The crystal structure of the prepared powder was characterized using an X-ray diffractometer (XRD, Bruker AXS D8 Advance; Bruker AXS, Karlsruhe, Germany). The thermal analysis of the prepared powder was carried oat using a Linseis L81/1750 (Linseis, Princeton-Jct, NJ). The surface area of the YSZ powder was analyzed by a Micromeritics ASPS 2020 analyzer (Micromeritics Instrument, Norcross, GA). Tap and bulk densities were obtained via Autotap

15

and Ultra pycnometer 1000 (Quantachrome, Boynton Beach, FL), respectively. The mean particle size was measured using a Malvern Zetasizer Nan0 ZS (Malvern Instruments, Worcestershire, U.K.). Furthermore, the composition of the prepared powder was analyzed using a Horiba 7539H energy-dispersive X-ray spectrometer (EDX; Hitachi, Tokyo, Japan). The EDX was attached to a field-emission scanning electron microscope (Hitachi S-4800; Hitachi, Tokyo, Japan). Results and Discussion

YSZ Preparation without Dispersants The experimental data for preparation of YSZ powder via the hydrothermal coprecipitation process are Table I.

Preparation Conditions and Name of Yttria-Stabilized Zirconia (YSZ) Cakes

Amount used (gmol) of reactants Zr

Y

D1

0.3499

0.06084

D2

0.3499

0.06084

Test name

shown in Tables I and 11. The mole ratio of YlZr in the feed precursor is 115.747. Both N H 4 0 H and urea (CO(NH&) were used as precipitants in the tests. The reaction rate of coprecipitation between Y3+(aq)IZr4+(aq)and NH40H(aq) is so fast that the reaction time and temperature are not critical for quality control of the precipitates. Hence, the reactant concentrations and the pH value of the system have a significant influence on the particle size of the precipitates. The structures of Y(OH),(s) and Zr(OH)4(s) are subject to change at an elevated temperature (i.e., near the critical point of the water) and fine powders are expected to be produced. However, dispersion of nanoscale oxides in an aqueous solvent is complicated by the strong tendency of nanoscale crys-

Name of YSZ cake'

YlZr Mole ratio

Total volume (ml)

P1 (pH = 9.21)

115.747

2000

c1 H T1

YSZ-D 1

P1 (pH = 9.17)

115.747

2000

c1

YSZ-D2 YSZ-D3 YSZ-D4

Precipitanta

Reactionb

D3

0.3499

0.06084

P2 (pH = 7.39)

115.747

2000

HT1 c1 HT 1

D4

1.3996

0.2434

P2 (pH = 7.39)

115.747

2500

c1

HT2 D5

1.3996

0.2434

P2 (pH = 7.39)

115.747

2500

YSZ-D5

c1

HT2 D6 and D7

1.3996

0.2434

P1 (pH = 7.39)

115.747

2,000

c1

D8

1.3996

0.2434

P1 (pH = 7.39)

115.747

2,000

HT3 c1 HT4

YSZ-D8

D9

1.3996

0.2434

P1 (pH = 7.39)

115.747

2,000

c1

YSZ-D9

2500 2500 2500 2500 2500 2500 2500

HT5 HT6 HT6 HT6 HT6 HT6 HT6 HT6

D10 D11 D12 D13 D14 D15 D16

1.3996 1.3996 1.3996 1.3996 1.3996 1.3996 1.3996

0.2434 0.2434 0.2434 0.2434 0.2434 0.2434 0.2434

P1 P1 P1 P1 P1 P1 P1

(pH = 9.28) (pH = 9.45) (pH = 9.36) (pH = 9.35) (pH = 9.29) (pH = 9.26) (PH = 9.44)

115.747 115.747 115.747 115.747 115.747 115.747 115.747

YSZ-D6 & YSZ-D7

(1 h) (2h) (4 h) (8h) (12 h) (20 h) (24h) . .

YSZ-D 10 YSZ-D 11 YSZ-D 12 YSZ-D 13 YSZ-D 14 YSZ-D 15 YSZ-D 16

'Pl: NH40H (conc.: 28-30%, p = 0.899); P2: CO(NH,),. bC1: 8 0 T , 1.5h; HTl: 3OO0C, 1 h (PI11.5MPa); HT2: 270°C, l h (PI28.8MPa); HT3: 330"C, 1 h (PS14.9MPa); HT4: 330"C, 2 h (PI14.9 MPa); HT5: 330"C, 3 h (PS14.9MPa); HT6: 180°C (PI1.0MPa). 'Cake drying: lOO"C, 48 h.

16

Progress in Nanotechnology: Processing

Table 11. Characteristics of Yttria-Stabilized Zirconia (YSZ) CakedPowders and Preparation of YSZ Powders Name of YSZ cake YSZ-D 1 YSZ-D2

Particle size (nm) 302.3 (85.8%) 91.4 (14.2%)

-

Calcination" CA1 cA2

CA1 CA2

Name of YSZ powder D 1A D1B D2A D2B

Particle size (nm) -

3000 3000 (94.3%) 139.9 (5.7%)

N C N C

-

YSZ-D3

1104 (55.5%) 172.4 (44.5%)

CAI CA2

D3A D3B

2993.3 (96.4%) 177.3 (3.6%)

N C

2403.9 (712%) 169.2 (28.8%) 2648.2 (80.0%) 292.2 (20.0%) 297.5 (100%) 582.8 (100%)

N C

YSZ-D4

-

CA3 CA2

D4A D4B

YSZ-D5

-

YSZ-D6

1924.1 (88.1 Yo) 174.8(17.9 %)

CA3 CA2 CA2

D5A D5B D6A

YSZ-D7

-

CA3 cA3-cA2 CA3-CA4 CA3-CA5 CA3-CA6 cA3 CA3-CA2 CA3-CA4 CA3-CA5 CA3-CA6 DY DY DY DY DY DY DY

D7A D7B D7C D7D D7E D8A D8B D8C D8D D8E D6 D7 D8 D9 D10 D11 D12

YSZ-D8

YSZ-D10 YSZ-D 11 YSZ-D 12 YSZ-D13 YSZ-D 14 YSZ-D 15 YSZ-D16

-

-

-

XRD patternb

-

-

-

-

195.5 (100%) 190.2 (100%) 259.8 (100%) 271.7 (100%) 253.7 (100%) 227.9 (100%) 278.4 (100%)

N C C N C C C C N C C C C N N N C C C C

'CAI: 300"C, 2 h; CA2: 1450"C, 3 h; CA3:llO"C, 50 h; CA4: 1450"C, 0 min; CA5: 1450°C,30 min; CAG: 1450°C 1 h; DY: drying at 7 5 T , 12 h. bN: impurity phase or incomplete cubic structure; C: 100% cubic crystal structure. XRD, X-ray diffraction.

tallites to agglomerate and form much larger particles (due to van der Waals and other forces). If no effort is made to control agglomeration, the measured particle size distribution of hydrothermally prepared YSZ suspensions extends well into the micrometer range.29 Hence, it is critical to eliminate agglomeration during the synthesis process. In these tests, the concentrations

Synthesis Methods for Powders

of Zr4' and Y3' were 0.1749 and 0.03042A4, respectively. In tests D1 and D2, NH40H was used as a precipitant, which was added to the feedstock of the YSZ precursor, and the p H values were adjusted to 9.21 and 9.17 for coprecipitation of Zr4' and Y3+, respectively. Then, the reactor was closed and its temperature was increased to 300°C and held for 1 h. After completion

17

Table 111. Influence of Calcination Time on Particle Size and Surface Area for Yttria-Stabilized Zirconia (YSZ) Powder Particle sizeb (pm) Name of YSZ powder

Calcination timea (min)

YSZ-D7

YSZ-D8

YSZ-D9

0 30 60 0 30 60 0 30 60

Mean

d,, (standard deviation)

4.777 4.684 6.553 2.919 5.891

3.642 3.379 5.375 2.072 4.433

(3.898) (4.349) (5.266) (2.783) (5.137)

-

-

2.439 5.043 6.023

1.694 (2.399) 3.670 (4.590) 4.186 (5.583)

Surface

XRD patternc

0.0865 0.0849 0.0915 0.0845 0.0709 0.0733 0.0935 0.0979 0.0658

C C C C C C C C C

area (m2/g)

"Calcination temperature: 1450°C. bVolume meadmedian (&). 'C: 100% cubic crystal structure. XRD, X-ray diffraction.

of the hydrothermal synthesis reaction, the system was cooled down to ambient condition via tap water through the jacket of the reactor. In tests D3, D4, and D5, urea (CO(NH2)2) was used as a precipitant. The coprecipitation reaction was first carried out at 80°C for 1.5 h. Then, the system was heated up to 270300°C and held for 1 h. (pressure > 11.6 MPa) to carry out the hydrothermal synthesis reaction. The reaction rate of coprecipitation among Y3'(aq), Zr4+(aq), and CO(NH2);! is controlled by the dissociation rate of CO(NH2)2 (i.e., reaction temperature). Hence, coagulation of precipitation can be moderately controlled and fine precipitates can be obtained easily. The average particle size of a YSZ cake directly obtained from the reactor is in the range of 91.4 nm to 1.1 pm. The flowability of the precipitate is good. A bimodal particle size is observed in the test system. It is generally accepted to be positive for pellet fabrication to promote the sintering density. Tables I1 and I11 show that the heat treatment of the YSZ cake/powder enhances the growth of the particle size. The calcination process (1450"C/3 h) for the production of a 100% cubic crystal YSZ caused the particle size to increase from 91.4-302.3 nm to around 6pm. It is obvious that the SOFC-grade YSZ powder must be subject to further processing so that a fine YSZ powder can be obtained in this case. Table I11 shows that the calcinations at high temperatures always cause a

-

decrease in the surface area of the powder. This trend may not be well consistent for all the cases due to the short time of calcinations. Table IV shows the typical relationship between milling time versus particle size of YSZ powders under dry milling. The particle size of YSZ powder decreases down to a submicrometer scale if the milling time is longer than I week. Hence, liquidphase milling with surfactants to disperse the agglomerate of the YSZ powder is a promising process to obtain fine YSZ powders. Liquid-phase ball milling and adjustment of the pH range are investigated in this study. Both the milling time and the pH value are key parameters for reduction of the particle size. Ball milling in a water system causes the reduction of the particle size of YSZ powder from 3 to 0.3pm at a fixed pH value. However, the influence of pH reduction (from 4.6 to 2.0) on the particle size of YSZ powder is not obvious or drastic. More experiments must be conducted to study

Table IV. Typical Relationship Between Milling Time Versus Particle Size of Yttria-StabilizedZirconia (YSZ) Powdera Milling time (h)

0

21.0

Particle size (nm)

3000 (94.3%) 139.9 (5.7%)

2487.2 (75.4%) 297.5 (24.6%)

aBy dry milling using a YSZ ball.

Progress in Nanotechnology: Processing

this aspect. However, dispersion of YSZ at a low p H may result in significant leaching of yttrium from YSZ into the solvent." This will limit the range of p H variation to control the particle size. After calcinations, an SOFC-grade YSZ powder with a 100% cubic crystal structure is obtained. The XRD pattern is shown in Fig. 2. At low-temperature calcinations (1 10°C, 50 h), the XRD patterns of D4A and D5A do not show a very good structure of crystal but have a fine particle size. However, the XRD patterns of D4B and D5B show well-crystallized characteristics after high-temperature calcinations ( 1450°C, 3 h). Hence, YSZ powder with a 100% cubic crystal may not be obtained via low-temperature calcinations. It is obvious that the phase separation between the precipitate and the solution is clear and sharp. The YSZ precipitate is well separated from the solution via settling. The morphology of the YSZ powder is shown in Fig. 3. Figure 4 shows the TGA and DTA curves for YSZ powders of YSZ-D2 and YSZ-D7. Table V shows the typical characteristics of a YSZ cake from TGA results. These results imply that the powders collected are hydroxide forms (Zr(OH)* and Y(OH),). A thermal de-

Fig. 3. Scanning electron micrographfor powder D3A.

composition reaction did not occur under these reaction conditions from D 1 to D9. In addition, a YSZ membrane is prepared by tape casting or hot-press methods. The typical powder properties of YSZ and the sintered density of a YSZ membrane are listed in Table VI.

Preparation of YSZ with NanoscaleLSOFC-Grade Cbaracteristics In order to directly prepare a YSZ powder with nanoscale/SOFC-grade characteristics via the hydro20

0

15 -10

10

-20

5

5 -z -30 n v)

2M .-

n

-

0

%

16

3

12

$

-10

8

-20

4

n

0

-30

0

400

800

1200

-A

Temperature ("C)

Fig. 2. x-ray diffaction pattern f i r powder YsZ-D4 and y.zD5 series. USZ, yttria-stabilized zirconia.

Synthesis Methods for Powders

Fig. 4. TGA and DTA curvesforyttria-stabilized zirconia PSZ) powders: (a) YSZ-D2 and (6) YSZ-D9.

19

Table V. Typical characteristics of Yttria-Stabilized Zirconia (YSZ) Cakes from TGA Results Name of YSZ cake YSZ-D2 YSZ-D7 YSZ-D8 YSZ-D9

Mass change

Mass change

(wtO/o)

(WtYO)

(Temperature (Temperature range, "C) range, "c) -22.39 (150-1600) -22.57 (150-1600) -23.18 (150-1600) -23.33 ( 150- 1600)

-28.95 (RT-1600) -29.40 (RT-1600) -26.85 (RT-1600) -26.98 (RT- 1600)

Notesa See Fig. 4a

See Fig. 4b

"1. Mass change for a reaction of Zr(OH)4-+Zr02+2H20to produce 1 gmol of ZrOz is -22.61 wt%. 2. Mass change for a reaction ofY(OH)3+ 1/2Y,0,+3/2 HzO to produce 0.5 gmol of Y 2 0 3 is -19.3Owt%. 3. Mass change for formation of 1 gmol 8YSZ from Zr(OH)* and Y(OH)3 is estimated to be -23.89wt%. 4. Zr(OH), (white amorphous powder) is converted to Z r 0 2 by evolving 2H20. 5. Y(OH), (white-yellow gelatinous material or powder) is decomposed to Y203. RT, room temperature (25°C).

with a mole ratio of ZrO,:Y,O, = 0.92:0.08. The total volume of the mixing solution was 2 L. Then, the proper amount of concentrated ammonia (about 400mL) was added under mechanical stirring, and a white precipitate was obtained via the coprecipitation reaction. The final p H value of the solution was above 9.5. The reaction time was the key parameter in tests D 10-D 16. The reaction temperature and pressure were selected to be 180°C and 965 kPa, respectively. The reaction time was in the range of 1-24 h. After the reaction, white and ultra fine YSZ powder was produced. The powder was washed twice by 0.1M nitric acid and deionized water, and then dried at 75°C in the oven. Figure 5 shows the XRD patterns of the D10-Dl6 powders at different reaction times and there are compared with those of the commercial YSZ powder made by NexTech Materials (Lewis Center, OH). Four mainly characteristic peaks were observed (20 = 30.1, 34.9, 50.2, and 59.6") and the crystal faces were marked as (1 1l), (2OO), (220), and (31 l), respectively. These peaks matched well with those of NexTech YSZ powder and were consistent with the database in JCPDS file (PDF No. 30-1468). However, the intensity of (200) increased with an increase in the reaction time. This suggests that the reaction time of the hydrothermal process was too short to produce a stable crystallite structure when the reaction time was < 8 h.

thermal coprecipitation process, other tests (Dl O-D 16) were performed. The concentrations of feedstock were 199 g/L and 0.265M for yttrium and free acid, respectively. ZrO(N03)2 xH20 was dissolved in deionized water and the yttrium stock solution was added. The mole ratio of Zr/Y was adjusted to produce a powder

-

Table VI. Typical Powder Properties of Yttria-Stabilized Zirconia (YSZ) and Sintered Density of a YSZ Membrane Name of YSZ powder

Density (g/cm3) Tap"

Trueb

Specific surface area

Sintered densityb (g/cm3)

-

5.42' 5.55" 5.31d

(m2/g)

D6A 3.640 aBy Autotap. b ~ Ultra y pycnometer. V i a hot press. dvia lamination.

20

5.844

Fig. 5. X-ray difiaction patterns of the powder prepared at different reaction times (tests Dl O-Dl6J

Progress in Nanotechnology: Processing

The crystallized size of YSZ powder was estimated from XRD data and calculated using the Schemer formula

where 8 is the Bragg angle of the diffraction peak; Kis a geometric factor (usually equal to 0.9); h is the X-ray wavelength using CUKEradiation (0.1542 nm); and P is the half-width of the diffraction peak. Following this equation, the broader peak indicated a smaller particle size. The estimated crystallite size of YSZ powder was about 4 nm. The composition of the prepared powder was measured by EDX, and the results are shown in Fig. 6 . It shows that the mole ratio of Y203 in the YSZ powder increases with an increase of the reaction time, but it retains a constant value when the reaction time is over 8 h. These results were consistent with that of the XRD pattern. Hence, the minimum reaction time for the direct production of a nanoscalelSOFC-grade YSZ powder from the hydrothermal coprecipitation process is 8 h. The surface area and density of YSZ powder at different reaction times were analyzed and are shown in Fig. 7. With increasing reaction time, the density exhibits a decreasing trend, but the surface area decreases rapidly from 182 to 105 m2/g. The results suggest that a

Fig. 6. Effect of reaction time on the yttria-stabilized zirconia powder prepared via the hydrothermal process (tests DI 0-Dl6).

Synthesis Methods for Powders

Fig. 7. Surface area (*) and density (A) ofyttria-stabilized zirconia powder at dzfferent reaction times (tests 010-DIG).

longer reaction time leads to the growth and agglomeration of YSZ particle size. A comparison of particle sizes of YSZ powder obtained from various methods analyzed at different reaction times is shown in Fig. 8. The crystallite size of YSZ powder decreases with increasing reaction time, but the

Fig. 8. Comparison ofparticle sizes ofyttria-stabilized zirconia powder at different reaction times, calcuhtedfiom various analyzed methods: , X-ray dzfiaction; 0, B E E A, dynamic light scattering (tests DIO-DI G).

+

21

Fig. 9. TGA curve f i r a nanoscaleyttria-stabilized zirconia (YSZ) powder of YSZ-Dl6

particle size of YSZ powder increases with an increase in the reaction time. In addition, the data analyzed by dynamic light scattering showed that the particle size of YSZ powder under ultrasonic treatment was about 250 nm. This implied that YSZ particles dispersed in the water solution caused particle agglomeration and then formed larger clusters. The nanoscale YSZ product was analyzed via TGA, and the curve is shown in Fig. 9. The 8 wt% loss is quite different from the results shown in Fig. 4. The weight loss of YSZ-D 16 may be caused by removing the adsorbed nitrate and the surface hydroxide group of the YSZ nanoparticles. In this study, a YSZ ceramic powder has been prepared successfdly via the hydrothermal coprecipitation process. The characteristics of YSZ powder at different reaction times were analyzed. The quasi-optimum reaction condition for YSZ powder preparation is determined.

Conclusion

The following conclusions can be drawn from the experiment: (a) SOFC-grade YSZ powders/materials with a 100% cubic crystal structure are successfully obtained using the hydrothermal coprecipitation process. The YSZ powder is also characterized to show its good flowability and low particle size (submicrometer scale). (b) Calcination/sintering processes lead to an increase of the particle size of YSZ powder. The temperature of calcinations is the key factor influencing the particle size of the powder. The sintering/calcinations always cause particle agglomeration. Hence, the final

22

SOFC-grade YSZ powder has particle size up to 6 pm if the reaction time of hydrothermal coprecipitation is < 1 h. (c) The particle size of YSZ powder may be reduced from 3 to 0.3 pm via ball milling in a water system. Hence, liquid-phase milling is one of the methods to obtain fine particles of YSZ powder. (d) pH adjustment is not a critical factor for particle size reduction of the YSZ powder. This is because the dispersion of YSZ at a low pH may result in significant leaching of yttrium from YSZ into the solvent. Hence, the range for p H adjustment is limited. (e) In closed hydrothermal systems, the system pressure increases with increasing reaction temperature. The precipitant of urea results in a quite higher system pressure than that of the N H 4 0 H case. Hence, the N H 4 0 H precipitant is better recommended for use in a hydrothermal coprecipitation system. (f) The temperature suitable for the direct preparation of nanoscale YSZ powder from a hydrothermal reaction system is in the range of 180”-200°C with a saturated vapor pressure of around 1-1.5 MPa. (g) In order to directly produce nanoscale and cubic crystal 8YSZ powders, the reaction time must be long enough. Reaction time is a critical factor in these tests. The quasi-optimum reaction time is around 8 h, which is the key factor to prepare a nanoscale 8YSZ powder directly. This means that the time required for completion of reaction (12) is longer than 8 h. Acknowledgments

The authors are extremely grateful to the Chemical Analysis Division and Nuclear Fuel and Material Division for carrying out the required work for chemical analysis, powder characterization, and SEM micrograph preparation. References M. C. Lee “Research and Development of Solid Oxide Fuel Cell (SOFC),

I. Activities and Tasks Accomplished in Fiscal Year of 2003 for Preparation/ Fabrication of the Membrane Electrode Assembly (MEA) of SOFC,” INER2682R, INER, Taiwan, January 2004. J. H. Hirschenhofer, D. B. Stauffer, R. R. Engeman, and M. G. Klett, Fuel Ccil Handbook, DOE/FETC-99/1076, 4th edition, 1-36. Parsons Corpora-

tion, Reading, PA, 1998. M. C. Lee “The Study o n the Preparation of Nano-scale Yttria Stabilized Zirconia (YSZ) Via Wet Chemical Coprecipitation Process,” INER-3345R, INER, Taiwan, January 2005.

Progress in Nanotechnology:Processing

4. C. H. Lee, M. C. Lee, C. S. Hwang, R. Y. Lee, and L. F. Lin, “SOFC Development at INER,” Int. /. Appl. Ceram. Technol., 4 [4] 339-349 (2007). 5. M. M. Seabaugh, S. L. Swarm, W. J. Dawson, K. Hasinska, and B. E. McCormick “Low-Cost Fabrication Processes for Solid Oxide Fuel Cells,” Fuel Celh 2000 Confcrence, Portland, October 30-November 2, 2000. 6. C. Bagger, S. Linderoth, M. Mogensen, P. V. Hendriksen, B. Kindl, S. Primdahl, P. H. Larsen, F. W. Poulsen, N. Bonanos, and M. J. Jorgensen, “Status of Danish Solid Oxide Fuel Cell RBrD,” Solid Oxide Fuei Celh VI, PV 99-19. eds., S. C. Singhal and M. Dokiya. The Electrochemical Society, Pennington, NJ, 28, 2001. 7. C. H. Hueh, R. D. Carneim, P. F. Becher, and T. R. Armstrong, “Constrained Sintering of Electrolyte Films on Porous Electrodes,” Solid Omye Fuel Celh VII, PV 2001-16. eds., H. Yokokawa and S. C. Singhal. The Electrochemical Society, Pennington, NJ, 293, 2001. 8. D. Perednis, M. B. Joerger, K. Honegger, and L. J. Gauckler, “Fabrication of Thin YSZ Electrolyte Films Using Spray Pyrolysis Technique,” Solid Oxide Fuel Celh VII, PV 2001-16. eds., H. Yokokawa and S . C. Singhal. The Electrochemical Society, Pennington, NJ, 989, 2001. 9. M. C. Lee and W. X. Kao “Preparation of Solid Oxide Fuel Cell (SOFC) Grade Ynria-Stabilized Zirconia (YSZ) Powder Via Hydrothermal Copredpitation Process and the Fabrication Process of Electrode Substrates for the Membrane Electrode Assembly ( M U ) of SOFC,” Proceeding of2005 Taiwa*/apan /oint Conjrmcc on SOFC, Taiwan, December 19-20,2005, 59-68. 10. C. Hayashi, “Ultrafine Particles,” Pbyr. Today, 40 [12] 44-51 (1987). 1 1 . W. R. Camon, S. C. Danforth, J. H. Flint, J. S. Haggerry, and R. A. Marra, “Sinterable Ceramic Powders from Laser-Driven Reactions: I. Process Description and Modeling,”J Am. Ceram. Soc,, 65 [7] 324-330 (1982). 12. J. N. Armor, A. J. Fanelli, G. M. Marsh, and P. M. Zambri, “Nonaqueous Spray-Drying as a Route to Ultrafine Ceramic Powders,”/. Am. Ceram. Soc., 71 [ l l ] 938-942 (1988). 13. H. J. Lee, K. Eguchi, and T. Yoshida, “Preparation of Ultrafine Silicon Nitride, and Silicon Nitride and Silicon Carbide Mixed Powders in a Hybrid Plasma,”]. Am. Ceram. Soc., 73 [ l l ] 33563362 (1990). 14. J. A. ONeill, M. Horsburgh, J. Tam, K. J. Grant, and J. A. Paul, “Production of Fine Ceramic Powders from Chloromethylsilanes Using Pulsed Excimer Radiation,”/. Am. Ceram. Soc., 72 [7] 1130-1135 (1989). 15. Y. C. Chang, M. C. Lee, C. H. Chiang, T. N. Lin, W. X. Kao, and C. H. Wang “Preparation of Nano-Scale YSZ Powder via Hydrothermal Coprecipitation Process,” INER-4563R, INER, Taiwan, March 2007. 16. X. J. Chen, K. A. Khor, S. H. Chan, and L. G. Yu, “Influence of Microstructure on the lonic Conductivity of Yttria-StabilizedZirconia Electrolyte,” Mater. Sci. Ens A, 335 [ 1-21 246252 (2002).

Synthesis Methods for Powders

17. A. U. Limaye and J. J. Helble, “Effect of Precursor and Solvent on Morphology of Zirconia Nanoparticles Produced by Combustion Aerosol Synthesis,”/. Am. Ceram. Soc., 86 [l] 273 (2003). 18. K. Prabhakaran, M. 0. Beigh, J. Lakra, N. M. Gokhale, and S. C. Sharma, “Characteristics of 8 mol% Yttria Stabilized Zirconia Powder Prepared by Spray Drying Process,” J Mater. Process. Techml, 189 [l-31 178-181 (2007). 19. H. L. Chang, P. Shady, and W. H. Shih, “Effect of Sodium on Crysrallite Size and Surface Area of Zirconia Powders at Elevated Temperatures,”/. Am. Cerum. Soc., 83 [8] 2055-2061 (2000). 20. K. Matsui and M. Ohgai, “Formation Mechanism of Hydrous Zirconia Particles Produced by Hydrolysis of ZrOCI, Solutions: IV. Effects of ZrOCI? Concentration and Reaction Temperature,” /. Am. Ceram. Soc., 85 [3] 545-553 (2002). 21. G. D. Dell’Agli, S. Fsposito, G. Mascolo, M. C. Mascolo, and C. Pagliuca, “Films by Slurry Coating of Nanometric YSZ (8 mol% Y,03) Powders Synthesized by Low-Temperature Hydrorhermal Treatment,” J Eur. Ceram. SOC.,25 [12] 2017-2021 (2005). 22. G. Dell’Agli and G. Mascolo, “Zirconia-Yttria (8 mol’Yo) Powders Hydrothermally Synthesized from Different Y-Based Precursors,” /. Eur. Ceram. Soc., 24 [6] 915-918 (2004). 23. R. R. Piricescu, C. Monty, D. Taloi, A. Motoc, and S . Axinte, “Hydrothermal Synthesis of Zirconia Nanomaterials,” /. Eur. Ceram. Soc., 21 [10-11] 2057-2060 (2001). 24. Y. B. Khollam, A. S. Deshpande, A. J. Patil, H. S. Potdar, S. B. Deshpande, and S. K. Date, “Synthesis ofYttria Stabilized Cubic Zirconia (YSZ) Powders by Microwave-Hydrothermal Route,” Mater. Chon. Pbys.. 71 [3] 235-241 (2001). 25. S. Somiya and R. Roy, “Hydrothermal Synthesis of Fine Oxide Powders,” Bull. Mater. Sci., 23 [GI 453-460 (2000). 26. X. M. Wang, P. Xiao, and G. Lorimer, “Highly Dispersed Suspension of YSZ and Zirconia Nanoparticles by Solvothermal Processing,” Adv. Appl. Ceram., 104 [3] 135-141 (2005). 27. C. Aymonier, A. Loppinet-Serani, H. Reverbn, Y. Garrabos, and F. Cansell, “Review of Supercritical Fluids in Inorganic Materials Science,”/. Supercrit. Fluids, 38 [2] 242-251 (2006). 28. W. J. Dawson, “Hydrothermal Synthesis of Advanced Ceramic Powers,” Ceram. Bull., 67 [lo] 1673-1678 (1988). 29. S. L. Swara, M. M. Seabaugh, and W. J. Dawson, “Low-Cost Fabrication Process of YSZ Electrolyte Films,” Solid Oxide Fuel Celh VI, PV 99-19. eds., S. C. Singhal and M. Dokiya. The Electrochemical Society, Pennington, NJ, 135, 1999.

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Synthesis of Nanosize Tin Dioxide by a Novel Liquid-Phase Process Yingke Zhou, Niladri Dasgupta, and Anil V. Virkar*.+ Department of Materials Science and Engineering, University of Utah, Salt Lake City, Utah 841 12

Welldispersed tin(IV) oxide (SnOz) nanosize crystallites were prepared by a novel liquid-phase process using NazSnO3 precursor of several micrometers in size. The product was characterized using powder X-ray diffraction, scanning electron microscopy, transmission electron microscopy, and BET. The nanosize SnOz powder formed has a narrow particle size distribution ranging between 1 and 7 nm, a very high surface area, and crystallizes in the tetragonal structure. A possible mechanism for the liquid-phase process, which is a combined effect of coprecipitation, hydrolization, and molecular decomposition, is proposed. I. Introduction

N

and nanoscale structures are of considerable interest due to their unique chemical and physical properties resulting from low dimensionality and the quantum confinement effect.' As an n-type semiconducting material with a large band gap, tin(1V) oxide (Sn02) has been widely used for various applications, such as solid-state gas sensors, optical devices, transparent electrodes, and rechargeable lithium ion batteries2-' Nanocrystalline tin oxide can be successfully synthesized by various hysical and chemical processes, including vacuum sputtering,!laser a b l a t i ~ n thermal ,~ evaporation: hydrothermal decomposition: sol-gel," and chemical coprecip itation." l 4 Among these, chemical coprecipitation is extensively used to synthesize nanosize Sn02 powder"-" because of its potentially low cost and amenability to mass production. In a conventional coprecipitation process, tin chloride is used as the precursor and NH40H or NaOH as the precipitant to form precipitate, which is subsequently calcined in the temperature range 4OO0-100O0C to form crystalline SnOz powder. The chemical reactions in the conventional coprecipitation/calcination process are as follows: ANOCRYSTALLITES

SnCl4t4NaOH -+ Sn(OH),

.+

Sn(OH),+4NaCI

Sn02+2H20

(1) (2)

A careful control of pH is often required during the precipitation process.1s17 In addition, the required high-temperature calcination process results in crystal growth, wider grain size distribution, and particle agglomeration, which is undesirable from the standpoint of subsequent processing. A case in point specific to Sn02-based sensors is the difficulty in uniformly distributing fine Pd catalyst particle^.^' In this communication, we report a simple liquid-phase process for the synthesis of nanosize Sn02 powder, which is based on a combined effect of coprecip itation, hydrolization, and molecular decomposition. An inorganic precursor powder, Na2Sn03,of several micrometers was S. Bhandarkar-ontrihuting

editor

Manuscript No. 23628. Received August 18, 2007; approved November 2. 2007. The financial s w w r t hv the Centers of Excellence Program of the State of Utah is gratefully acknowledged. . 'Member, the American Ceramic Society. 'Author to whom correspondence should be addressed. e-mail anil.virkarym.cc.utah. edu

Synthesis Methods for Powders

first synthesized. The precursor was treated in dilute nitric acid (HN03) to directly obtain well-dispersed Sn02 nanocrystallites without the need for subsequent calcination. Sn02nanoparticles with narrow size distribution and large surface area were successfully prepared. 11. Experimental Procedure

For the synthesis of Sn02 nanoparticles, Na2Sn03 was chosen as the precursor, which was prepared by a conventional solidstate reaction. Required amounts of SnO and Na2C03 were mixed by wet ball milling in alcohol for 24 h and then dried. The mixture was calcined in air at 1150°C to form Na2Sn03. The precursor was then boiled in dilute H N 0 3 (0.07M) for half an hour. The resulting powder was thoroughly washed to pH = 7 with deionized water and dried to obtain nanosize Sn02powder. In order to obtain greater insight into the reaction mechanism, the following experiments were conducted: (I) Na2Sn03 powder was mixed with water. The mixture was stirred for half an hour at three different temperatures: 25", 80", and 100°C. In one experiment at 100"C, the mixture was boiled for 4 h. After the reaction, the undissolved powder was recovered by filtration for subsequent characterization; (11) dilute H N 0 3 solution (0.07M) was added to filtrates from the mixtures heated at 80" and 100°C for half an hour, and the resultant solutions were boiled for half an hour. This led to the formation of a precipitate, which was washed until the pH was 7 and was heated to dryness. In all experiments, the ratio of the precursor solid to the liquid (either H 2 0 or HN03) was set at 8 g:100 mL. Negligible loss occurred during subsequent heating for half an hour. Water was, however, replenished at regular intervals to offset the loss due to evaporation in the experiment in which the precursor was boiled in water for 4 h. The structure and morphological properties of the precursor and nanosize SnO2 powder were characterized by several techniques. Powder X-ray diffraction (XRD) was carried out on a Phillips X'Pert PW3040 diffractometer (Lelyweg, Almelo, The Netherlands) with CUKE radiation and a graphite monochromator at 45 kV and 40 mA. Scanning electron microscopy (SEM) was conducted using a Phillips XL30 ESEM FEG. Transmission electron microscopy (TEM) was conducted on a Tecnai F30 TEM (FEI Company, Hillsboro, OR). BET surface area measurements were made on a TriStar 3000 (Micromeritics Instrument Corporation, Norcross, GA). 111. Results and Discussion

Figure 1 shows an XRD pattern of the Na2Sn03precursor. The diffraction peaks correspond to the Na2Sn03anorthic structure (JCPDS 30-1252), except for the two small peaks that are attributed to the hydrous phase Na2Sn03. 3H20 (JCPDS 24-1 143). Na2Sn03 is very sensitive to water vapor in the ambient, and readily absorbs water vapor to form the hydrous compound. Both the anhydrous and hydrous phases readily react with HN03. Figure 2 shows an SEM image of the as-prepared Na2Sn03 precursor. The precursor exhibits an irregular morphology and its size ranges between 1 and 20 pm.

-

25

Fig. 1. An X-ray diffraction pattern of the as-synthesized Na2Sn03 precursor.

After reacting with HNO,, the residual powder was washed and dried. An XRD pattern of the powder is shown in Fig. 3(a). As a comparison, an XRD pattern of commercial Sn02 powder is also shown in Fig. 3(c), which can be indexed as a tetragonal cassiterite structure (JCPDS 77-0447). Note that all major peaks of the cassiterite structure are present in the nanosize powder synthesized. The low-intensity peaks cannot be resolved because of peak broadening and overlapping. Figure 3(a) shows that the XRD peaks of powders synthesized by the liquid-phase process are very broad, which is indicative of a very fine crystallite size. The presence of all major peaks in the as-synthesized nanosize Sn02 powders also suggests that the powder is crystalline and not amorphous. This indicates that reaction of Na2Sn03 with HNO, leads to the formation of crystalline nanosize Sn02 in a single step, i.e. without the subsequent calcination process as required in the typical coprecipitation process. The as-synthesized nanosize Sn02 powder was then subjected to a thermal treatment at 800°C for 5 h. The corresponding XRD pattern is shown in Fig. 3(b). Note that the peak positions in Fig. 3(b) are almost the same as those in Fig. 3(c), although the peaks are still wider. The average crystal size was calculated using the Schemer equation. The peaks of commercial coarse Sn02 powder (around 10 pm, Fig. 3(c)) also show some broadening, attributed to the XRD instrument itself. Thus, for calculating the crystallite size, the selected peak width (110) was corrected by subtracting the corresponding peak width in Fig. 3(c). The results show that the crystal size of the as-prepared sample (Fig. 3(a)) is -2.5 nm (from 110 peak), and after thermal treatment the particle size increased to -41 nm (Fig. 3(b), from 110 peak). This shows that very fine nanoparticles of Sn02 can be formed, starting with a coarse Na2Sn03 precursor several micrometers in size, by reacting with HNO,.

Fig.3. An X-ray diffraction pattern of (a) nanosize tin(1V) oxide (Sn02) powder made by the liquid-phase process; (b) nanosize Sn02 powder made by liquid-phase process, and subjected to thermal treatment at 800°C for 5 h; (c) commercial SnO2 powder.

The specific surface area of the nanosize Sn02 powder formed was measured to be -220 m2/g. Table I compares the specific surface area of the nanosize Sn02 powder synthesized in the present work with nanosize Sn02 powder synthesized by other techniques described in the literature. Note that the Sn02 powder synthesized in the present work exhibits the highest surface area of all the techniques reported. Assuming particles to be spherical for simplicity, the particle diameter (D) may be estimated from the following relation?

D= 6/Sp

(3)

where p is the densit of Sn02 and S is the specific surface area. With p = 6.95 g/cmY for Sn02, and the measured BET surface area of S = 220 m2/g, the particle diameter is estimated t o be -3.9 nm. The estimated particle size from surface area measurement is not much larger than that obtained from XRD peak broadening, indicating that the Sn02 nanoparticles are well separated from each other and exhibit very little agglomeration. The morphology and the structure of Sn02 nanoparticles were further characterized by TEM and high-resolution TEM (HRTEM). Figure 4(a) is a bright-filed TEM image of the assynthesized Sn02, which shows that the sample consists of welldispersed, near-spherical nanoparticles with minimal agglomeration. Typical HRTEM image in Fig. 4(b) reveals the fine structure of the Sn02 powder. The clear lattice fringes in the HRTEM image confirm the high degree of crystallinity of the nanosize Sn02 powder formed. The spacing between two adjacent lattice planes is 0.34 nm, which corresponds to the “d” spacing for (1 10) planes of tetragonal Sn02. Figure 4(c) shows a histogram of the Sn02 nanoparticle size distribution. The particle size ranges from 1 to 7 nm, of which the majority is between 2 and 4 nm (85%); indicating a very narrow particle size distribution.

Table I. ~~~

Comparison of Typical BET Surface Areas of tin(1V) oxide (Sn02)Powder by Different Processes ~

Process

Fig. 2. An scanning electron microscopic image of the as-synthesized Na2Sn03 precursor.

26

Sol-gel Combustion Coprecipitation Redox coprecipitation Liquid phase

~

Surface area (m2/g)

References

143

Zhang and Gaol8 Zhang and Gao” Xie and Varadan” Wang and XieI4 This work

55

42 29 220

Progress in Nanotechnology: Processing

Fig.4. (a) A transmission electron microscopic (TEM) image of nancsize tin(1V) oxide (SnO2) powder; (b) high-resolution TEM images of nanosize Sn02 powder. Inset shows a higher magnification image of one particle; (c) size distribution of nanosize Sn02 powder.

All powders obtained by reacting the precursor with water at various temperatures (Section 11) were characterized by XRD and the results are shown in Fig. 5. After stirring in water at 2 5 T , very little Na2Sn03dissolved (5.8%), and the structure of the residual powder was the same as the precursor except for a small increase in the amount of Na2Sn03. 3H20. With increasing temperature, more and more Na2Sn03could be dissolved in water. After boiling in water (lOO"C), most of the precursor was dissolved (82.2%), and the residual powder was a mixture of the precursor and Sn02. This suggests that when boiled in water for half an hour, a small amount of Na2Sn03can be hydrolyzed to form Sn02. XRD trace of the residual powder after boiling in water for 4 h is shown in Fig. 5(d). As seen in Fig. 5(d), most of the precursor converted to Sn02, with a small amount of Na2Sn03. 3H20, the hydrous precursor left unconverted. From Fig. 5(d), the crystal size of Sn02 was calculated to be 13 nm. This suggests that when boiled in water, the conversion efficiency from precursor to Sn02 can be increased by increasing the reaction time. However, the crystal size of the product also increases. While further work is required, this observation suggests that by varying the boiling time, it may be possible to vary the particle size, yet maintaining it in the nanosize regime. When boiled in dilute HNO, for half an hour, the entire precursor was converted to nanosize Sn02 by the following reaction:

-

The corresponding XRD trace is shown in Fig. 5(e). XRD patterns of precipitates obtained by reacting dissolved precursor with HNO, are shown in Fig. 6. Figures 6(a) and (b) show that the formed powders are identical to those obtained from the liquid-phase process (Fig. 6(c)) described earlier. The corresponding crystal sizes from Figs. 6(a) and (b) are -2.3 and 1.8 nm, respectively, similar to the crystal size of 2.5 nm in Fig. 6(c). The samples were also examined by TEM. The size distributions of the two samples are very similar to each other; but narrower (1-5 nm) than that obtained for the sample prepared by the liquid-phase process as shown in Fig. 4(c) (1-7 nm).

-

Synthesis Methods for Powders

-

Fig.5. X-ray diffraction patterns of the powders obtained as described in Section I 1 (a) Na2Sn03precursor; (b) Na2Sn03 precursor after stirring in water at 25°C for half an hour; (c) Na2Sn03precursor after stirring in water at 80°C for half an hour; (d) Na2Sn03 precursor after stirring in water at 100°C for 4 h; (e) Na2SnO3 precursor after boiling in 0.07M nitric acid solution at 100°C for half an hour.

The preceding suggests that there exist two possible parallel mechanisms for the formation of nanosize Sn02 starting with Na2Sn03. The overall process (consisting of the two parallel mechanisms) is termed here as the liquid-phase process. (1) Na2Sn03dissolved in water exists as Na+ and SnOf ions. In the dissolved state, either a reaction with H2O or H N 0 3 occurs to form nanosize Sn02 by hydrolization precipitation or precipitation by an acid. (2) The other parallel process consists of a direct reaction between undissolved Na2Sn03 and H20 or HN03. Either of the parallel mechanisms directly yields nanosize, crystalline Sn02 powders. The latter mechanism involves molecular decomposition in which the fugitive constituent (here Na20) is leached away from the solid precursor (Na2Sn03)to form soluble products (either NaOH in the case of H2O or NaN03 in the case of HN03) and leaving behind an insoluble oxide (nanosize Sn02). The formation of nanosize powders by molecular decomposition has been previously demonstrated in which nanosize Zr02 was formed by reacting Na2Zr03 with H20.'6The exact assignment ratio of these parallel reactions will depend on specific reaction conditions, such as the ratio of the precursor solid to the liquid, presence or absence of HNO,, reaction temperature and time, etc. When acid is used, most of the Na2Sn03 (dissolved and undissolved) will preferentially react with HN03. Also it may be possible that under some specific condition that suppresses the molecular decomposition reaction, the formation of Sn02 occurs by precipitation from the dissolved Na2SnO3, as the undissolved precursor continuously dissolves and replenishes Sn0:- ions in the solution. 27

decomposition, and is based on a liquid reacting with a solid precursor to leach away fugitive constituent (here Na20) and forming insoluble nanosize Sn02. (b) Dissolution of the precursor in the liquid followed by acid precipitation or hydrolization precipitation resulting in nanosize SnOz. Which of the two processes dominates may depend upon various parameters such as temperature and composition. The resulting particle size is on the order of 2 4 nm with surface area as large as -220 m2/g. The obtained Sn02 is crystalline of a tetragonal structure.

-

!I .-c

References

C u)

c

-

I

40

I

60

I

I

80

C

20

60 2 Theta [deg.]

40

80

Fig. 6. X-ray diffraction patterns of the powders obtained as described in Section 11: (a) nanosize tin(1V) oxide (Sn02) produced by first boiling the precursor in water at 100°C for half an hour, and then boiling the filtrate in dilute nitric acid (HNO,); (b) nanosize Sn02 produced by first boiling the precursor in water at 80°C for half an hour, and then boiling the filtrate in dilute HNO,; (c) nanosize Sn02 powder made by the liquid-phase process.

In a conventional coprecipitation process, typically amorphous metal hydroxides are first obtained from the solution. Subsequently, the hydroxides are calcined to remove water and form crystalline oxides. This, however, often leads to particle growth and the formation of hard agglomerates. The liquidphase process described here directly leads to the formation of nanosize oxide powders requiring no further thermal treatment, unlike conventional coprecipitation processes. We believe this to be the primary benefit of this novel process, as this leads to soft agglomerates which can be easily separated into individual crystallites during subsequent processing. The near congruence of crystallite size obtained from X-ray line broadening and particle sizes obtained by TEM and from surface area data is evidence for the absence of hard agglomerates.

IV. Conclusions In summary, we introduce a new liquid-phase process for the synthesis of nanosize Sn02 powder using Na2Sn03as a precursor. The formation of nanosize powder by reaction with either H N 0 3 or H 2 0 involves two parallel processes: (a) a direct reaction with solid Na2Sn03. This process is termed molecular

28

‘J. Hu, T. W. Odom, and C. M. Lieber, “Chemistry and Physics in One Dimension: Synthesis and Properties of Nanowires and Nanotuhes.” Arc. Chem. Res., 32, 43545 (1999). ’A. M. Azad, S. A. Akbar, S. G. Mhaisalkar, L. D. Birkefeld, and K. S. Goto, “Solid State Gas Sensors: A Review,” J. Electrochem. Soc., 139,369&704 (1992). ’S. R. Vishwakarma, P. Rahmatullah, and H. C. Prasad, “Low Cost Sn02:P/ Si02/n-Si (Textured) Heterojunction Solar Cells,” J. Phys. D,26, 95942 (1993). 4R.Lalauze. J. C. L. Thiesse, C. Pijolat, and M. Soustelle, “Sn02 Gas Sensor: Effect of SO2Treatment on the Electrical Properties of Sn02.” Solid State Ionics, 12,453-7 (1984). ’Y. Idota, T. Kubotd, A. Matsufuji, Y. Maekawa, and T. Miyasaka, “TinBased Amorphous Oxide: A High-Capacity Lithium-Ion-Storage Material,” Science, 276, 1395-7 (1997). ‘G. B. Barhi, J. P. Santos, P. Serrini, P. N. Gibson, M. C. Horrillo, and L. Manes, “Ultrafine Grain-Size Tin-Oxide Films for Carbon Monoxide Monitoring in Urban Environments,” Sensors Actuators B, 24/25, 559-63 (1995). ’G. Williams and G. S. V. Coles, “Gas Sensing Properties of Nanocrystalline Metal Oxide Powders Produced by a Laser Evaporation Technique,” J. Muter. Chem., 8, 1657-64 (1998). *H. Ogawa, A. A. M. Nishikawa, and S. Hayakawa, “Electrical Properties of Tin Oxide Ultrafine Particle Films,” J. Electrochem. Soc., 128, 2020-5 (1981). 9H. M. Cheng, J. M. Ma, Z. G. Zhao, and L. M. Qi. “Hydrothermal Synthesis of S n 0 2 Nanocrystals,” Chem. J. Chin. Univ., 17, 833-7 (1996). “S. S. Park and J. D. Mackenzie, “Thickness and Microstructure Effects on Alcohol Sensing of Tin Oxide Thin Films,” Thin Solid Film, 274, 15&9 (1996). “J. Wang, M. Yang, Y.Li, L. Chen. Y. Zhang, and B. Ding, “Synthesis of FeDoped Nanosized SnOz Powders by Chemical Co-Precipitation Method,” J. NonC r st Solidy, 351, 228-32 (2005). ‘H. Yang, X. Song, X. Zhang, W. Ao, and G. Qiu, “Synthesis of VanadiumDoped Sn02 Nanoparticles by Chemical Co-Precipitation Method.” Muter. Lett., 57, 3124-7 (2003). ”N. C. Pramanik, S. Das, and P. K. Biswas, “The Effect of Sn(1V) on Transformation of Co-Precipitated Hydrated In(II1) and Sn(lV) Hydroxides to Indium Tin Oxide (ITO) Powder,” M a w . Lett., 56. 671-9 (2002). I4X. Wang and Y. C. Xie, “CHI Deep Oxidation on SnO,-Bdsed Composite Oxides: Novel, Active, and Stable Catalysts,” Chem. Lett.. 3, 216-7 (2001). “C. B. Lim and S. Oh, “Microstructure Evolution and Gas Sensitivities of Pd-Doped Sn02-Based Sensor Prepared.” Sensors Actuators B: Chem., 30,223-3 I (1996). ’‘Y. Jiang, S. V. Bhide. and A. V. Virkar. “Synthesis of Nanosize Yttria-Stabilized Zirconia by a Molecular Decomposition Process.” J. Solid State Chem.. 157, 14%59 (2001). ”A. V. Virkar and S. V. Bhide, “Molecular Decomposition Process for the Synthesis of Nanosize Ceramic and Metallic Powders”; US Patent 6803027, October, 2004. ‘*J. R. Zhang and L. Gao, “Synthesis and Characterization of Nanocrystalline Tin Oxide by SoCGel Method,” J . Solid State Chem.. 177, 1425-30 (2004). I9J. Zhang and L. Gao, “Synthesis of Crystalline A T 0 Nanoparticles by Combustion Method,” Wuji Huaxue Xuehuo, 20, 801-1 (2004). %J. Xie and V. K. Varddan. “Synthesis and Characterization of High Surface Area Tin Oxide/Functionalized Carbon Nanotubes Composite as Anode Materials,” Mater. Chem. Phys., 91, 274-80 (2005).

1 I

20

Progress in Nanotechnology: Processing

Fabrication of Nanocomposite Powders of Carbon Nanotubes and Montmorillonite Jingdong Feng and Qingwei Wang' Institute of Applied Chemistry, College of Chemistry, Jilin Normal University, Jilin, China 136000

Nanocomposite powders of carbon nanotubes (CNTs) and montmorillonite (MMT) were successfully synthesized by heating the mixed composites of polypropylene, nickel (Ni) catalyst, and organic-modified montmorillonite in large quantitiesafter the Ni catalysts were removed. Transmission electron microscopy, X-ray diffraction,thermogravimetric analysis, Raman spectroscopy, and scanning electron microscopy show the co-presence of CNTs and MMT. The interlayer distance of MMT was enlarged because of the CNTs, which result in MMT being partly exfoliated and buried in the CNTs. These current nanocomposite powders of CNTs and MMT with superstructure are therefore expected to be ideal fillers for high-performance polymer nanocomposites.

S

1. Introduction

the discovery of carbon nanotubes (CNTs) in 1991 by Iijima,' there has been great interest in developing new methods for the synthesis of CNTs because of their outstanding mechanical, electrical, and chemical properties. Theoretical and experimental studies have shown that single-walled carbon nanotubes (SWCNTs) possess Young's moduli of the order of 1-5 TPa and tensile strengths as high as 60 GPa. Depending on their diameters and orientations of their hexagons with respect to the tube axis, CNTs can be either semiconducting or metallic. The axial thermal conductivity of CNTs is higher than even that of diamond. These unique properties suggest that CNTs hold great promise for synthesizing new kinds of multifunctional nanocomposites. As for their use in ceramic-matrix composites, at least two different architectures can be conceived: ( I ) CNTs can be used as reinforcements to improve mechanical properties and add new functionality, and (2) ceramic materials can be used to protect functional CNT structures from harsh environments. The ceramic-matrix composites are potentially important for those applications where requirements such as mechanical strength or resistance to corrosive environments prevent the use of other matrices such as polymers and metals. Additionally, CNT-ceramic composites have been the subject of several previous investigations. Ma and Liu' synthesized Sic composites containing 10 vol% of multiwalled carbon nanotubes (MWCNTs) by hot pressing the mixture of nanometer-sized powders and CNTs. Peigney and colleagues3" synthesized CNT-Fe/ A1203 composites by growing either SWCNTs or MWCNTs in situ. However, only slight improvements in mechanical properties were observed in these previous studies. More recently, Zhan et ~ 1 . ' reported ~~ the synthesis of nano-phase A1203 composites containing SWCNTs by a spark-plasma sintering technique; the composites possessed significantly improved mechanical properties and electrical conductivity over the pure INCE

J. Ferreird-contributing editor

Manuscript No. 23208. Received May 15, 2007; approved October 28, 2007. This work was supported by the Youth Fund of Jilin Normal University. 'Author to whom correspondence should be addressed. e-mail: qwwdngllnu@yahoo. com

Synthesis Methods for Powders

matrix materials. Montmorillonite (MMT) is another very promising ceramic for the fabrication of nanocomposites. The dispersion of nanometer-size MMT nanoplatelets within a polymer matrix leads to nanocomposites exhibiting markedly improved physicochemical properties, such as a higher Young's modulus and storage modulus, higher thermal stability and flame retardancy, and more efficient gas barrier properties, com ared with pure polymers or conventional microcomposites?-' Recently, nanocomposites of CNTs and MMT also attracted increasin interest because of their structure and potential applications.12-15 The growth of CNTs on MMT layers could provide another interesting class of materials for potential applications. CNTs rooted on smectite layers are particularly attractive for polymer reinforcing applications. One would expect that the combination of CNTs and clay in polymer matrixes would provide outstanding functionalities to the resulting composite materials. Toward this aim, Zhang et al.19 demonstrated that the incorporation of the CNTGclay hybrid system into nylon-6 by simple melt blending led to exfoliated nanocomposites with excellent mechanical properties. Compared with neat polymer, the tensile modulus and the tensile strength of the composite were greatly improved, by about 290% and 150%, respectively, by the incorporation of only 1 wt% CNT-clay hybrid reinforcement. Although this method is very meaningful, the cost to produce the composites is relatively high. Therefore, the development of an easy and continuous low-cost method for the production of these nanocomposite powders enabling them to be used in commercial fields would be welcome. Recently, Tang et aLzo have synthesized CNTs from the composites of polypropylene (PP)/organicmodified montmorillonite (OMMT)/nickel (Ni) on flames in large quantities. In their experimental procedure, MMT was removed as waste product using hydrofluoric acid. This method was employed in the present work to synthesize nanocomposite powders of CNTs and MMT. Herein, we kept MMT in the system but only disposed of Ni catalyst. Thus, large quantities of nanocomposites of CNTs and MMT were obtained. Besides being easy, this method has other potential applications. Owing to the high aspect ratio of these nanocomposite powders, a small addition of these nanocomposites to the polymer matrix would enhance the mechanical, thermal, flame retardant, and barrier properties of polymers without a significant loss of clarity, toughness, or impact strength, enabling one to recycle the carbon sources and use the energy sources more efficiently.

11. Experimental Procedure

( 1 ) Materials and Sample Preparation PP used in this work was a commercially available extrusion grade powder with 0.8 g/min of melt flow index, from Yanshan Chemical Company (Beijing, China). The OMMT (I.44P, composed of units made up of two silica tetrahedral sheets centered with an alumina octahedral sheet) was bought from Nanocor (Arlington Heights, IL) and the Ni catalyst (Ni, on silicaalumina, catalyst) was obtained from Alfa Aesar (Ward Hill, MA). Nitric acid and hydrofluoric acid were purchased from Beihua Chemicals (Beijing, China). 29

The experimental procedure is as follows. In a typical procedure, Ni catalyst ( 5 wt%), PP (85 wtX), and OMMT (15 wt%) were mixed together in a Brabender mixer at 70 rpm at 170°C for 10 min. Then a piece of the above-mixed composite was placed in a crucible and heated in muffle at 800°C for about 5 min (the furnace atmosphere was air), after which the crucible was taken out and rapidly cooled to room temperature. The as-obtained black powder was then purified with nitric acid (3M) to partially dispose of Ni catalyst and obtain purified nanocomposite powders of CNTs and MMT. Pure CNTs can also be obtained by using hydrofluoric acid and nitric acid to remove MMT and Ni catalyst.

( 2 ) Characterization Scanning electron microscopy (SEM) measurements were carried out with field emission scanning electron microscopy (FESEM; Hitachi S-4700, JEOL, Tokyo, Japan). Samples, fixed onto adhesive tape, were used to examine the morphological variation after sputter coating with gold. An accelerating voltage of 20 kV was used. Transmission electron microscopy (TEM) was performed using a JEOL 200CX transmission electron microscope operated at 150 kV. Phase analyses of the products were obtained by X-ray diffraction (XRD) using a Rigaku D/MAX-IIB (Rigaku Ltd., Tokyo, Japan) powder X-ray diffractometer with CuKa radiation. Raman spectroscopy was acquired using a JYHR800 Raman spectrometer (JY Co., Paris, France) with an excitation beam wavelength at 532 nm. Thermogravimetric analysis (TGA) was carried out in air at a heating rate of 10"C/min in a Perkin-Elmer (Waltham, MA) TGA 7 thermal analyser.

tion process, which means that the content of CNTs in the nanocomposite powders of CNTs and MMT increased. On the other hand, a characteristic diffraction peak of MMT at 20" 28 also appeared intensified after the purification process. The two peaks appearing together proved the presence of CNTs and MMT in the final product. A very weak and broad diffraction band at 44.5" 28 still existed after purification using nitric acid, which suggests that the metallic Ni has mostly been removed. However, the diffraction at 20" and 26.5" 28 still exists and even intensified.

( 2 ) SEM Studies To further prove the presence of CNTs and MMT, studies were made by FESEM. Figure 2(a) shows that a large amount of CNTs appear in the SEM image and a relatively small amount of MMT, which are sheet-like particles having a length of about 1 pm and are randomly distributed. Some stack together and others bury in the CNTs. The CNTs exist among the MMT or tangle together. Figure 2(b) shows the CNTs purified using hydrofluoric acid and nitric acid. Only CNTs entangled together can be observed in this photo, suggesting that the MMT has been removed by the acid leaching treatment. The diameter of the CNTs is about 3&50 nm and the length is up to several micrometers. The white dots at the head of CNT are the remainings of Ni nanoparticles, which are protected by the graphic carbon coating on them. So nitric acid could not remove it and this is the reason for the existence of a weak diffraction peak at 44.5 28.

111. Results and Discussion

( I ) X-Ray Diffvaction Analysis Figure 1 shows the XRD patterns of a series of nanocomposites. The pure MMT obtained from heating the OMMT at 800°C for 5 min does not show any diffraction peak in the range of 5'-20" 28. Only a very weak diffraction effect can be observed at about 20" 28, which is ascribed to the remains of MMT structure. At 26.5" 28, there is a very weak diffraction peak that belongs to carbon. Composite-I is the mixture of PP/OMMT/Ni obtained in a Brabender mixer at 170°C. From 10" to 25" 28, there are some diffraction peaks, which are ascribed to PP. Composite-3 is the as-obtained product heated in the muffle at 800°C without purification. In this graph, the diffraction peaks from 10" to 25" 28 disappeared but the diffraction peak at 26" 28 appeared, which means that the graphitic carbon appeared. So it can be concluded that PP has been partially transformed into carbon. The diffraction peak at 44.5" 28 is attributed to metallic Ni. Composite-2 is the prepared nanocomposite powder of CNTs and MMT, which was purified using nitric acid after heating in the muffle. The diffraction peak of graphitic carbon intensified after the purificananocompositepowders of MMT and CNTs

I

750

.b

Y)

-

600 -

-

450-

300 150 -

l 0

.

~ 10

.

, 20

. 30 2 theta

r

.

F 40

'

t

'

50

Fig. 1. X-ray diffraction patterns of produced nanocomposites of carbon nanotubes and montmorillonite.

30

Fig.2. (a) Scanning electron microscopy (SEM) image of nitric acidtreated composites (after purification). (b) SEM image of nitric acid-treated carbon nanotubes (after purification).

Progress in Nanotechnology: Processing

90 -

-g Irn

60 -

.-

s

30 -

1

Fig. 3. Transmission electron microscopy image of montmorillonite and carbon nanotubes.

(3) TEM Analysis To characterize the nanocomposite powders of CNTs and MMT more deeply; TEM analyses were performed to examine the morphology of CNTs and MMT nanocomposites. Figure 3 shows the TEM image of nanocomposite powders of CNTs and MMT after purification with nitric acid. From this image, the coexistence of CNTs and MMT can be observed more clearly. The gray areas in the images correspond mainly to clay platelets. The arrow points to the CNTs buried in the MMT. The sue of MMT is about micrometer scale and the diameter of CNT

/

3.2%

is about 10-50 nm. Because the small CNTs could not be seen in the SEM image, only the relatively big CNTs can be seen. Figures 4(a) and (b) show images of purified CNTs with lengths of about 1-2 pm and diameters of about 2G50 nm. The hollow tube structure of the as-treated CNTs can be seen. Some CNT walls are very thick while others exhibit thin walls, with the difference being more distinctly observed in Fig. 4(b). The hollow tube structure can be more clearly observed for the big tube with very thin walls.

( 4 ) TGA Analysis TGA measurements from the final MMT-CNTs composites and CNTs can survey the carbon content of the obtained composites and also the quality of the CNTs produced, because their oxidation temperature is strongly influenced by the presence of wall defects. Ideal graphite starts to oxidize above 600"C, while CNTs would be oxidized above 500°C. In Fig. 5, pure CNTs began to decompose at 585°C in air; when the temperature was further increased, the weight loss increased rapidly until all the CNTs were exhausted at about 630°C. The nanocomposites of CNTs and MMT systems had similar TGA behavior under these experimental conditions. However, it could be seen that higher amounts of residue were obtained for the nanocomposites of CNTs and MMT systems than for pure CNTs at temperatures ranging from 580" to 900°C. The change in TGA residue at 900°C is also apparent in Fig. 5, with the average listed for pure CNTs (3.2% w/w) and nanocomposites of CNTs CNT

1

1000

Fig.4. (a and b) Transmission electron microscopy image of obtained carbon nanotubes.

Synthesis Methods for Powders

1200

1400 1600 Wave number (cm-')

1800

2000

Fig.6. Raman spectrum of purified carbon nanotubes obtained from polypropylene/organic modified montmorillonite/nickel.

31

Fig. 7. Schematics of the formation process of nanocomposites

and MMT (20% w/w). Hence the content of MMT in nanocomposites of CNTs and MMT was about 17%.

( 5 ) Raman Analysis In order to characterize the nature of obtained CNTs after purification with hydrofluoric acid and nitric acid, Raman spectroscopy was used to further characterize graphitization of the MWCNTs (Fig. 6). As shown in Fig. 6, two strong peaks appear at 1590 and 1346 cm-I, respectively, in the Raman spectrum. The former (so-called G-band) is attributed to an E2g mode of hexagonal graphite and is associated with the vibration of sp2 hybridized carbon atoms in a graphite layer. The latter (so-called D-band) is ascribed to the vibration of carbon atoms with dangling bonds in the plane terminations of disordered graphite. Such disordered structures might originate from the purification process. Thus, the Raman spectroscopic results agree with the TEM observations.

(6) Illustration Formation Process of Nanocomposites of CNTs and MMT Figure 7 shows the formation process of nanocomposites of CNTs and MMT. Firstly, PP, OMMT, and Ni catalyst were mixed together to obtain composites of PP/OMMT/Ni, and then PP was transformed into CNTs in the presence of Ni catalyst at 800”C.20Secondly, nitric acid was used to release the remaining Ni catalyst and the nanocomposite powders of CNTs and MMT were obtained. It is important to discuss the formation mechanism of CNTs from the current PP/OMMT/Ni composites. On the basis of the previous discussion, a formation mechanism was proposed, as shown in Fig. 7. PP and OMMT are known to form spherulites of PP/OMMT under heating conditions. In the process of Brabender mixing, PP and OMMT may form spherulites that contain Ni. When the PP/OMMT/Ni composite is placed in an oven at 900°C, OMMT platelets migrate to the surface of spherulites because of temperature, viscosity gradients, gas formation, interfacial tension between OMMT and PP, and surface free energy of the polymer.16 PP pyrolyzes to give pyrolytic species that are sealed in an autoclave-like microreactor formed by OMMT platelets.*’ Within the autoclave-like microreactor, the pyrolytic species are catalyzed to form CNTs on these Ni nanoparticles. The confinement of the microreactor results in curled CNTs and eventually in chestnut-like spheres.

IV. Conclusions Nanocomposite powders of CNTs and MMT have been prepared in mass production using PP as carbon source and Ni nanoparticles as catalyst after the Ni catalysts were removed. The MMT with partially exfoliated structure appears distributed in the CNTs. This current CNT/MMT hybrid with this superstructure

32

is therefore expected to be an ideal filler for high-performance polymer nanocomposites. Furthermore, for its easy production it can be used in many ways such as catalyst carrier.

References IS. lijima, “Helical Microtubules of Graphitic Carbon,” Nature, 354, 5 6 8 (1991). ’R. Z. Ma and J. Wu, “Processing and Properties of Carbon Nanotube-NanoS i c Ceramic,” J . Mater. Sci., 33, 5243-6 (1998). ’A. Peigney, C. Laurent, 0. Dumortier, and A. Rousset. “Carbon NanotuksFe-Alumina Nanocomposites. Part I:Influence of the Fe Content on the Synthesis of Powders,” J . Eur. Cerum. Suc., 18, 1995-2004 (1998). 4C. Laurent, A Peigney, 0. Dumortier, and A. Rousset. “Carbon NanotubeFc-Alumina Nanocomposites. Part 11: Microstructure and Mechanical Properties of the Hot-Pressed Composites.” J. Eur. C e r m . Soc., 18, 2005-13 (1998). Zhan, J. D. Kuntz, J. Wan, and A. K. Mukherjee. “Single-Wall Carbon Nanotubes as Attractive Toughening Agents in Alumina-Based Nanocomposites,” Nut. Muter., 2, 38 4 2 (2003). ‘G. Zhan, J. D. Kuntz, J. E. Garay, and A. K. Mukherjee, “Electrical Properties of Nanoceramics Reinforced with Ropes of Single-Walled Carbon Nanotubes,” Appl. Phys. Lett.. 83, 1228--30 (2003). B. K. Kuila and A. K. Nandi, “Physical, Mechanical, and Conductivity Properties of Poly(3-HexylthiophenejMontmorillonite Clay Nanocomposites Produced by the Solvent Casting Method,” Macromoleculer, 37,8577 -84 (2004). ‘5. W. Gilman, C. L. Jackson, A. B. Morgan, R. J. Hams, E. Manias, E. P. Giannelis, M. Wuthenow. D. Hilton, and S. H. Phillips, “Flammability Properties of Polymer-Layered-Silicate Nanocomposites. Polypropylene and Polystyrene Nanocomposites,” Chem. Mater., 12, 186673 (2000). 9W. Xie, R. Xie, W. P. Pan, D. Hunter, B. Koene, L. S. Tan, and R. Vaia, “Thermal Stability of Quaternary Phosphonium Modified Montmorillonites,” Chem. Muter., 14, 483745 (2002). ‘OW. Xie, 2. Gao, W. P. Pan, D. Hunter, A. Singh, and R. Vaia, “Thermal Degradation Chemistry of Alkyl Quaternary Ammonium Montmorillonite,” Chem. Mater., 13, 2979-90 (2001). ”T. Tang, X.Chen, H. Chen, X. Meng, Z. Jiang, and W. Bi, “Catalyzing Carbonization of Polypropylene Itself by Supported Nickel Catalyst During Combustion of Polypropylene/Clay Nanocomposite for Improving Fire Retardancy,” Chem. Mater., 17. 279942 (2005). ”R. H. Baughman, A. A. Zakhidov. and W. A. de Heer, “Carbon Nanotubethe Route Toward Applications,” Science, 297, 787-92 (2002). I3A. B. Dalton, S. Collins, E. Mufioz. J. M. Razal, V. H. Ebron, J. P. Ferraris, I. N. Coleman, B. G. Kim, and R. H. Baughman, “Super-Tough Carbon-Nanotube Fibres,” Nature, 423, 703 (2003). l4J. M. Adams. ‘Synthetic Organic Chemistry using Pillared, Cation-Exchanged and Acid-Treated Montmorillonite Catalysts-A Review,” Appl. Clay. Sci.. 2, 3 0 ‘ 4 2 (1987). I5Z. Ding, J. T. Kloprogge, R. L. Frost. G. Q. Lu, and H. Y. Zhu, “Porous Clays and Pillared Clays-Based Catalysts, Part 2: A Review of the Catalytic and Molecular Sieve Applications,” J. Porous Mater., 8, 273-93 (2001). I6Y. Li. J. Chen, Y. Qin, and L. Chang, “Simultaneous Production of Hydrogen and Nanocarbon from Decomposition of Methane on a Nickel-Based Catalyst,” Energy Furls, 14, I188 (2000). I’D. Gournis. M. A . Karakassides, T. Bakas, N. Boukos, and D. Petridis, “Catalytic Synthesis of Carbon Nanotubes on Clay Minerals,” Carbon, 40,2641 (2002). ‘*A.Bakandritsos, A. Simopoulos, and D. Petridis, “Carbon Nanotube Growth on a Swellable Clay Matrix,” Chem. Mater.. 17. 3468-74 (2005). I9W. D. Zhang, 1. Y. Phang, and T. X. Liu, “Growth of Carbon Nanotubes on Clay: Unique Nanostructured Filler for High-Performance Polymer Nanocomposites,” Adv. Mater., 18, 73-7 (2006). ’9 Tang, . X. C. Chen, X. Y. Meng. H. Chen, and Y. P. Ding, “Synthesis of Multiwalled Carbon Nanotubes by Catalytic Combustion of Polypropylene.” Angew. Chem. Inr. Ed., 44, 1517-20 (2005).

Progress in Nanotechnology: Processing

Synthesis of Highly Dispersed Barium Titanate Nanoparticles by a Novel Solvothermal Method Xiao Wei, Gang Xu, Zhaohui Ren, Yonggang Wang, Ge Shen, and Gaorong Hant Department of Materials Science and Engineering, Zhejiang University, Hangzhou 3 10027, China

A novel solvothermal route has been developed to synthesize highly dispersed nanocrystalline barium titanate (BaTiOJ, using a mixture of ethylenediamine and ethanolamine as a solvent. The as-synthesized BaTiOJ nanoparticles were characterized by X-ray powder diffraction, transmission electron microscopy (TEM), high-resolution TEM, Fourier transform infrared spectroscopy, and thermal analysis. Based on the results of characterizations, the organic solvent was found to influence strongly the crystal growth and dispersibility of BaTi03. The BaTi03 nanoparticles obtained were highly dispersed and crystalline with a cubic perovskite structure. The particle size derived from the TEM ranged from 5 to 20 nm.

I. Introduction A R I U M TITANATE

B

(BaTi03) is one of the most extensively investigated and widely used perovskite oxides as a material for thermistors,' multilayer ceramic capacitors (MLCCs)? electro-optic device^,^ and dynamic random access m e m ~ r i e s be,~ cause of its high dielectric constant536 and ferroelectric property.' In view of the current trend of miniaturization in the fabrication of electronic devices with improved properties and high reliability, nanoscale BaTiO3 materials are desirable. For MLCCs, in particular, improved performance can be obtained by decreasing the thickness of the ceramic layer with a decrease of the particle size.* Furthermore, ceramics produced from agglomerate-free nanosized powders with a narrow size distribution have shown unique consolidation and compaction properties.' Another important reason for synthesizing nanosized BaTi03 is as an attempt to understand the finite size effect on the ferroelectricity in nanoscale materials, which is much debated in the literature.'&12 In this case, the highly dispersed nanocrystals have the benefit of precise investigation into the dependence of ferroelectricity on the particle size for an individual nan0partic1e.l~ The traditional solid-state reactionI4 and wet-chemical processing, including sol-gel rocessing,15 oxalate route,16 and homogeneous precipitation,' have been used to produce BaTi03 materials. To form a pure crystalline phase, however, a calcination step above 700°C was necessary in all these methods, which led to coarse and chemically bonded aggregates. Although the synthesis of BaTi03 by an alkoxidehydroxide sol-precipitation process did not require further calcination at an elevated temperature, it also produced agglomerated and irregular-shaped powders with a broad size d i s t r i b ~ t i o n . ' ~The ~ ~ 'hydrothermal method has been used to prepare highly dispersed BaTi03 powders, whereas it was found to be inadequate to further downsize E. Slamovich-ontributing

editor

Manuscript No. 23171. Received May 6, 2007; approved August 6, 2007. This work was supported by the National Science Foundation of China under grant number 50452003. 'Author to whom correspondence should be addressed. e-mail: [email protected]

Synthesis Methods for Powders

the particle size to become smaller than several tens of nanometer.2022 Crystalline BaTi03 particles with diameters smaller than 10 nm have been synthesized by the decom osition of bimetallic alkoxide precursors in organic solventj3 and the "benzyl alcohol Nevertheless, the nonaqueous reaction approaches have shown advantages in the subtle control of the crystallization and particle Herein, we report the solvothermal synthesis of highly dispersed and crystalline BaTi03 nanoparticles with a particle size of 5-20 nm through the reaction between barium acetate (Ba(CH3C00)2)and tetrabutyl titanate (Ti(OC4H9)4)in a mixture of ethylenediamine (NH2CH2CH20H)and ethanolamine (NH2CH2CH2NH2) at a relatively low temperature of 200°C. This method exhibited better control over particle size, crystallinity, and dispersibility of the as-synthesized BaTi03 nanoparticles than that in a hydrothermal system. 11. Experimental Procedure

Ba(CH3C00)2and T ~ ( O C A Hwere ~ ) ~used as the starting materials to introduce the Ba and Ti cations into the solution in a 1:l molar ratio. A mixture of NH2CH2CH20H and NH2CHz CH2NH2 in a 1: 1 ratio of volume was used as the reaction medium. In a typical procedure of preparing BaTi03 nanoparticles, Ti(OC4H9)4 (5 mmol) was dissolved into lOmL of ethanol. Then, 2 mL of 0.15Mammonia solution was added to the above solution to obtain a precipitated titanium hydroxide. The titanium hydroxide precipitate was filtered and washed with distilled water three times, and then dispersed in a 30mL mixture of NH2CH2CH20Hand NH2CH2CH2NH2with vigorous stirring, followed by addition of 2 mL of 2.5M Ba(CH3C00)2 aqueous solution and 2.74 g of KOH. Finally, the final suspension above was transferred into a 50 mL homemade stainless steel Teflon-lined autoclave. The autoclave was sealed and maintained at 200 "C for 12 h, and then cooled to room temperature. The resultant product was filtered and washed with distilled water and ethanol five times. The final product was dried at 90°C for 24 h. The phase purity of the as-synthesized products was measured by X-ray powder diffraction (XRD) using a Rigaku D/ max-RA X-ray diffractometer (Rigaku, Tokyo, Japan) with CuKa radiation (h = 1.5406 A). Transmission electron microscopy (TEM), high-resolution TEM (HRTEM) images, and the corresponding selected area electron diffraction (SAED) pattern were determined on a JEM 2010CX TEM at an accelerating voltage of 200 kV. The TEM sample was prepared by dispersing a small drop of the ultrasonicated suspension in ethanol onto a copper grid precoated with an amorphous carbon film. The Fourier transform infrared spectroscopy (FTIR) spectrum was taken with a Thermo Nicolet (Waltham, MA) Model 47-FTIR in the wave ranging from 4000 to 400 cm-l using the KBr method. Differential thermal analysis (DTA) and thermogravimetric analysis (TGA) were carried out at a heating rate of 10"C/min in air using a VL-VAC-7000 thermal analyzer (Candy Precision Instruments Co. Ltd., Shanghai, China). 33

111.

Results and Discussion

Figure 1 shows the XRD pattern of the as-synthesized product via the solvothermal route at 200°C for 12 h. All diffraction peaks in Fig. 1 can be indexed to the cubic perovskite structure [space group: Pm3m] of BaTiO3 (JPCDS No. 31-174) with very few crystalline byproducts of BaC03 (marked with asterisk in Fig. 1) that can be removed by an acid washing process.26The BaTi03 average crystallite size was calculated to be approximately 14.6 nm using line broadening (of the (200) reflection) with Scherrer's equation

R = 0.9 h / ~ c o eB s

(A),

where R is the average particle size B is the width of the peak at half the peak height (radians), h is the X-ray wavelength (nm), and 0~ is the Bragg angle ("). Figure 2 displays the representative TEM images of the as-prepared BaTiO, nanocrystals in the mixture of NH2CHzCH20H and NH2CH2CH2NH2. An overview image (Fig. 2(a)) at low magnification shows that spherical nanoparticles with particle sizes of 5-20 nm were present in the solvothermally synthesized sample. The SAED pattern taken from this sample is shown in Fig. 2(b), which exhibits polycrystalline diffraction rings. The diffraction rings correspond well to the cubic perovskite structure and the first five diffraction rings are indexed as (IOO), (1 lo), (1 1I), (200), and (210), respectively. This is in agreement with the XRD results. HRTEM images of isolated nanocrystals are shown in Figs. 2(c) and (d). Clear lattice fringes with dl = 2.83 A give evidence that the particles are single crystalline with high crystallinity. Figure 3 shows the FTIR spectrum of the as-synthesized BaTi03 nanoparticles. In the spectrum, the 1636 cm-' band may be attributed to the bending vibration of the -NH2 groups and the 1427 cm-I band may be attributed to the scissoring vibration of -CH2 groups. The weak peaks at 1060 cm-l may be due to the stretching vibration of C-N groups. The broad band at 3436 cm-' may be attributed to the -OH group (or the OH-'). All the above observations confirmed the existence of some organic solvent molecules that were possibly adsorbed on the crystal surface. It was also found that some acetate anions were present and identified at the 1557 cm-' band, which possibly resulted from the recursor of Ba(CH3C00)2.The broad adsorption at 569 cm- may be due to the vibration of Ti06 octahedra. The sharp and weak peaks at 1382 cm- may be assigned to the C0;2vibration.27 BaTi03 nanoparticles solvothermally synthesized in the mixture of NH2CH2CH20H and NH2CH2CH2NH2were characterized for thermal behavior. The TGiDTA results are shown in

P

'

Fig. 1. X-ray powder diffraction pattern of the product synthesized by the solvothermal route at 200°C for 12 h in a mixture of ethylenediamine and ethanolamine.

34

Fig. 2. Representative transmission electron microscope (TEM) images of barium titanate (BaTiO,) nanoparticles solvothermally synthesized in a mixture of ethylenediamineand ethanolamine. (a) An overview image at low magnification; (b) selected area electron diffraction; (c) and (d) high-resolution TEM images of isolated BaTiO, nanocrystals.

Fig. 4. The DTA curve shows that there is an endothermal peak at 90.7"C, which may be due to desorption of water molecules adsorbed on the crystal surface. A further increase in temperature changes the thermal behavior from endothermal to exothermal, and induces two exothermal peaks at 304.9" and 367.6"C, which may be attributed to the combustion of the NH2CH2CH20Hand NH2CH2CH2NH2.The TG curve exhibits a drastic weight loss of -7% up to 600°C, which is due to desorption of adsorbed water molecules and the combustion of the adsorbed NH2CH2CH20H and NH2CH2CH2NH2molecules. This is consistent with the DTA analysis. The small weight loss in the range of 600-1000"C is probably due to the decomposition of a tiny BaCO, byproduct. As revealed by some successful synthesis works, a nonaqueous route has been proved to be a robust way to synthesize crystalline oxides,28in which the crystal growth could be well controlled. In this study, the mixture of NH2CH2CH20Hand NH2CH2CH2NH2is found to be a crucial factor in preparing Progress in Nanotechnology: Processing

100

..

367.6

1

'

,

.

)

.

98

. &?

96 94 92

W Q,

m

4000

I

I

3000

2000

Wavenumber Icm-1

I

1000

Fig. 3. Fourier transform infrared spectroscopy spectrum of barium titanate nanoparticles solvothermally synthesized in a mixture of ethylenediamine and ethanolamine.

901

,

200

.

I

.

400 600 TemperaturePC

,

800

.

,

1000

Fig. 4. Thermo gravimetric (TG)/differential thermal analysis (DTA) of barium titanate nanoparticles solvothermally synthesized in a mixture of ethylenediamine and ethanolamine.

that slowed the mass transport of ions; however, its high dielectric constant led to the formation of few nuclei in the initial BaTi03 nanoparticles with a narrow size distribution. Among stage, which also finally resulted in large crystals (shown in Fig. 5(b)). Accordingly, the mixture of NHzCH2CHzOH and many physical and chemical properties, the dielectric constant and viscosity of solvents play a significantly important role in NH2CH2CH2NH2,which had intermediate dielectric constant and viscosity compared with pure NH2CH2CH20H and the size-control growth. On the one hand, in a solvent with low dielectric constant ions, saturation easily occurs and small crysNH2CH2CH2NH2,was considered to be the optimum solvent to synthesize BaTi03 nanoparticles due to the optimum dielectallites may be more steady than that in a solvent with a high tric constant and viscosity for the nucleation process and crystal dielectric c o n ~ t a n t . ~ This ~ ' ~ ' accelerates the nucleation process growth. and more nuclei are formed in the initial stage in the solvent with In addition, the solvent molecules would bind to the growing a low dielectric constant, which favors the preparation of small nanocrystal surface, and the capping effect might be another crystals. On the other hand, during the formation of nanocrysreason for the weakening of the coarsening growth.32The surtals, the competition between the nucleation and the crystal growth is partially controlled by diffusion. The mass transport face-adsorbed solvent molecules may also act as a dispersant to of the ions of oxide is more difficult in a solvent of high viscosity. prevent the particle agglomeration b a steric stabilization in the This favors the formation and cr stallization of new nuclei, subsequent suspension preparation!' The presence of organic resulting in a smaller crystal size.3r At room temperature, the molecules probably adsorbed on the crystal surface was condielectric constant and viscosity are 12.9, 1.6 mPa.s for firmed by the FTIR spectrum and the thermal analysis shown in NH2CH2CH20H and 37.7,24.1 mPa .s for N H ~ C H Z C H ~ N H ~ , Figs. 3 and 4, respectively. respectively. It was noticeable that, although the dielectric constant of NH2CH2CH20H was low, the particle size of the assynthesized sample from the solvent of pure NH2CH2CH20H IV. Conclusions was large (shown in Fig. 5(a)), which was mainly ascribed to the easy mass transport of ions due to its low viscosity. Compared In summary, highly dispersed single-crystalline BaTi03 nanoparticles with a narrow size distribution in the range of 5-20 nm with NH2CH2CH20H, NH2CH2CH2NH2had high viscosity

Fig.5. Transmission electron microscope images of barium titanate nanoparticles synthesized at 200°C for 12 h in different pure solvents of (a) ethylenediamine and (b) ethanolamine. Synthesis Methods for Powders

35

have been produced with a novel solvothermal reaction using the mixture of NH2CHzCHzOH and NH2CHzCHrNHz as a solvent. The characterization results confirmed that the mixed solvent played a key role in confining the crystal growth by providing a proper dielectric constant, viscosity, and capping effect, and that the solvent molecules adsorbed on the crystal surface were responsible for the high dispersibility in the subsequent suspension preparation through the steric effect. The high dispersibility of the as-synthesized BaTi03 nanoparticles would facilitate investigations into the size effects of ferroelectricity for isolated particle. Furthermore, the agglomerate-free and nanocrystalline BaTiO, would be suitable for lower temperature sintering so as to produce a denser bulk material.

References ‘J. Nowotny and M. Rekas. ”Defect Chemistry of BaTi03.” SolMSfufe 1onic.s. 49. 13>54 (1991). ’D. Hennings, M. Klee. and R. Waser. “Advanced Dielectrics: Bulk Ceramics and Thin Films.” Adv. Muter.. 3. 33440 (1991). ‘M. Mori. T. Kineri, K. Kadono, T. Sakaguchi, M. Miya. H. WakabdydShi, and T. Tsuchiya. “ E f f ~ tof the Atomic Ratio of Ba to Ti on Optical Properties of Gold-Dispersed BaTiO, Thin Films.” J. Am. Ceruni. Soc., 78, 23914 (1995). 4A. 1. Kingon, S. K. Streiffer. C. Basceri, and S. R. Summerfelt. ”High-Permittivity Perovskite Thin Films for Dynamic Random-Access Memories,” MRS Bull., 21, 46-52 (1996). ‘P. K. Dutta. R. Asiaie. S. A. Akbar. and W. Zhu, “Hydrothermal Synthesis and Dielectric Properties of Tetragonal BaTiO,.” Cheni. Murer.. 6, 1542-8 (1994). ‘Y. S. Her, E. Matijevi. and M. C. Chon, ”Preparation of Well-Defined Colloidal Barium Titanate Crystals by the Controlled Double-Jet Precipitation.’‘ J. Muter. Res.. 10, 3106-14 (1995). ’T. Lee and 1. A. Aksay, “Hierarchical Structure-Ferroelectricity Relationships of Barium Titanate Particles,” Cryst. Grosfh Des.. 1. 40-19 (2001). “H. Kishi. Y. Mizuno, and H. Chazono, “Base-Metal Electrode-Multilayer Ceramic Capacitors: Past. Present and Future Perspectives.” Jup. J. Appl. Phw., 42. 1-15 (2003). ‘A. Arora, “Ceramics in Nanotech Revolution.” Adv. Eng. Muter., 6, 244-7 (2004). lop. Ghosez and K. M. Rabe. “Microscopic Model of Ferroelectricity in StressFree PbTiO, Ultrathin Films.” Appl. Ph.w L p f f . ,76. 2767-9 (2000). “R. Bottcher. C. Klimm. D. Michel, H. C. Semmelhack, and G . Viilkel, “Size Effect in Mn’+-Doped BaTiOl Nanopowders Observed by Electron Paramagnetic Resonance.” Phys. Rev. B. 62. 2085-95 (2000). ”X. Deng. X. Wang. H. Wen, A. Kang. Z. Gui. and L. Li. “Phase Transitions in Nanocrystalline Barium Titanate Ceramics Prepared by Spark Plasma Sinte‘’ J. Am. Cerani. Sot., 89. 1059-64 (2006). Ray. Y. V. Kolen’Ko. D. Fu. R. Gallage. N. Sakamoto, T. Watanabe, M. Yoshimura. and M. Itoh. “Direct Observation of Ferroelectricity in Quasi-ZeroDimensional Barium Titanate Nanoparticles.” Srnoll. 2. 1427-31 (2006).

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I4L. K. Templeton and J. A. Pask, “Formation of BaTiO, from BaCO, and TiOl in Air and in C02,” J. A m . Ceruni. Soc.. 42. 212-6 (1959). ”H. Shimooka and M. Kuwabara, “Crystallinity and Stoichiometry of NanoStructured Sol-Gel-Derived BaTi03 Monolithic Gels.” J. Ani. Ceruni. Soc., 79. 2983-5 (1996). “H. I. Hsiang and F. S. Yen, “Effect of Crystalline Size on the Ferroelectric Domain Growth of Ultrafine BaTiO, Powders.” J. h i . Cerum. Soc., 79, I 0 5 3 4 (1 996). I7M. Stockenhuber. H. Mayer, and J. A. Lercher. “Preparation of Barium Titanates from Oxalates.” J. Am. Ceruni. Sot.., 76, 1185-90 (1993). ‘“F. Chaput and J. P. Boilo. “Alkoxide-Hydroxide Route to Synthetize BaTi03-Based Powders,” J. Am. Cerurn. SOC..73. 942-8 (1990). ”S. Yoon and S. Baik, “Formation Mechanism of Tetragonal Barium Titanate Nanoparticles in Alkoxide-Hydroxide Sol-Precipitation Synthesis,” J. Am. Cerum. Soc., 89, 1816-21 (2006). 201. J. Clark. T. Takeuchi. N. Ohtori. and D. C. Sinclair, “Hydrothermal Synthesis and Characterization of BaTiO, Fine Powders: Precursors, Polymorphism and Properties.” J. Mufer. Chon., 9. 83-91 (1999). ’IS. W. Lu, B. 1. Lee. 2. L. Wang, and W. D. Samuels, “Hydrothermal Synthesis and Structure Characterization of BaTiO, Nanocrystals.” J. Crjsf. Growth. 219, 269-76 (2000). 22X. Zhu, J. Zhu, S. Zhou, Z. Liu, N . Ming. and D. Hesse, “BaTiO, Nanocrystals: Hydrothermal Synthesis and Structural Characterization,” J. Crysf. Growrh. 284, 486-94 (2005). 23 S. OBrien, L. Brus. and C. B. Murray, “Synthesis of Monodisperse Nanoparticles of Barium Titanate: Toward a Generalized Strategy of Oxide Nanoparticles Synthesis,’’ J. Am. Cheni. Soc.. 123. 12085-6 (2003). 24M. Niederberger. N. Pinna, J. Polleux. and M. Antonietti. “A General Soft-Chemistry Route to Perovskites and Related Materials: Synthesis of BaTiO3. BaZr03, and LiNO, Nanoparticles.” Angen. Cheni. I n / . E d , 43. 2270-3 (2p). M. Niederberger, G. Garnweitner. N. Pinna. and M. Antonietti. “Nonaqueous and Halide-Free Route to Crystalline BaTi03. SrTiO,, and (Ba, Sr)TiOl Nanoparticles via a Mechanism Involving C-C Bond Formation.” J. Am. Chon. Soc.. 126. 9 I 2 W (2004). 26 J. 0. Eckert, C. C. Hung-Houston, 8 . L. Gersten. M. M. Lencka, and R. E. Riman. “Kinetics and Mechanisms of Hydrothermal Synthesis of Barium Titanate,” J. Am. Ceruni. Soc.. 79, 2929-39 (1996). 27 S. Ghosh. S. Dasgupta. A. Sen. and H. S. Maiti, “Synthesis of Barium Titanate Nanopowder by a Soft Chemical Process.” M u m . Left., 61. 53841 (2007). ’“G. Garnweitner and M. Niederberger. “Nonaqueous and Surfactant-Free Synthesis Routes to Metal Oxide Nanoparticles,” J. Am. Cerum. Soc.. 89. 1801-8 (2006). ”Q. Q. Wang. G. Xu, and G. R. Han. “Solvothermal Synthesis and Characterization of Uniform CdS Nanowires in High Yield.” J. Solid State Chon.. 178. 268W5 (2005). ”X. B. Zhao, X. H. Ji, Y. H. Zhang. and B. H. Lu, “Effect of Solvent on the Microstructures of Nanostructured BizTe3Prepared by Solvothermal Synthesis.” J. A//o.v.s Conipd.. 368, 349-52 (2004). ”J. C. Yu. L. Zhang. Q. Li. K. W. Kwong, A. W. Xu.and I. Lin, “Sonochemical Preparation of Nanoporous Composites of Titanium Oxide and Size-Tunable Strontium Titanate Crystals,’’ Lungniuir, 19. 7673-5 (2003). ’2T. He, D. Chen. and X. Jiao, “Controlled Synthesis of Co304 Nanoparticles Through Oriented Aggregation.” Chem. Muter.. 16. 73743 (2004).

Progress in Nanotechnology: Processing

CONTINUOUS PRODUCTION NANOPARTICLES

AND

HARVESTING

OF

INORGANIC-CERAMIC

S.A.E. Abdulla”, P.A. Sermona’*,M.Worsley“ and I.R. Collinsb “Chemistry, SBMS, University of Surrey, Guildford, Surrey, GU2 7XH. UK bBP. Chertsey Road, Sunbury-on-Thames., TW 16 7LN, UK. ABSTRACT The authors demonstrate here that the continuous synthesis of BaS04, Ca3(PO4)2, LaCo03, BaCe03 etc can be achieved to give a range of particle sizes and morphologies in microemulsions and can then be continuously water-in-TX 100/cyclohexane/2-methyl-2-propanol harvested by cooling the microemulsion to 268 K. Surprisingly, the nanoparticles then appear in the organic-rich upper phase. INTRODUCTION Nanomaterials can be nm sized building blocks for a whole range of applications, e.g. synthetic bone, bone cements, pigments and ceramics provided that their preparation can be scaled up. The last ten years has seen steady development in the uses of microemulsions in the preparation of metal’, semiconductor’ phosphate2, mixed oxide314(e.g. manganate4 or ferrite4) and ceramic5 nanoparticles of controlled size, composition, structure and morphology. Here the authors consider how this may be scaled up using four examples: BaS04, tricalcium phosphate/hydroxyapatite, BaCe03 and LaCo03. Microemulsions are of interest in materials preparation as the reactions that lead to particle formation of a characteristic size and morphology6 occur within dynamic droplets of nm (colloidal) size in which reactant concentrations define the size of solid particles formed. The microemulsion consists of three components: water, oil and surfactant (and possibly a cosurfactant). At certain compositions the system is optically clear and spontaneously formed. Generally water-in-oil microemulsions have the water phase dispersed in the hydrocarbon phase’, but stabilised by the added surfactant.’ The properties of microemulsions has been explained’ in terms of the continual coalescence/re-separation of droplets. The novel aim of this work was to continuously synthesise and harvest nanoparticles of Bas04 (where nanoribbon~’~ have been continuously harvested), tricalcium phosphate/ hydroxyapatite, BaCe03 (recently prepared by a microemulsion route’* as 5-50nm nanoparticles) and Lac003 by mixing separate water-in-oil microemulsions containing the reactive ions (e.g. Ba2+and (S04)2- lo,”) and then thermally cycling these around the phase inversion temperature. The nanoparticles formed were to be compared with those produced in free solution. EXPERIMENTAL ), (2,2,4-trimethylpentane; Decane (p = 0.73 g.cm”), dodecane (p = 0.75 g . ~ m - ~isooctane p = 0.692 g.cm”),p-xylene (1,4-dimethylbenzene; p = 0.86 g.cm”), 2-propanol (p = 0.785 g.cm3 ), cyclohexane (p = 0.78 g.cm”; Fisher Scientific) and 2-methyl-2propanol (p = 0.78 g.cm”) were used as hydrocarbons. Ethanol (p = 0.789 g.cm”) was used with distilled water. The

Synthesis Methods for Powders

37

following surface-active agents have been used: (i) SPlO (Sisterna; HLB2 sucrose ester of stearate/palmitate (10% monoester)), (ii) AOT (sodium bis-(2-ethylhexylsulfosuccinate;sodium dioctyl sulfosuccinate; RMM = 444.59, (iii) C12EO4 or Brij30 (poly(oxyethy1ene-4-dodecyl ether), RMM = 362.56; density = 0.95 g . ~ m - ~(iv) ) , DDAB (didodecyldimethylammonium bromide; RMM = 462.65) and (v) Triton XlOO or octyl phenol ethoxylate (RMM = 625; p = 1.07 g . ~ m - ~In ) . the aqueous phases the reagents used were BaC12, Na2S04 (Fisher Chemicals), Na3P04 (BDH), Ca(NO3)2 (BDH), Ba(NO3)2 (Sigma Aldrich), Ce(N03)3 (Acr6s Organics), La(N03)3, Co(N03)3, CuC12.2H20 and ammonium oxalate monohydrate (RRM = 142.1 1; Sigma Aldrich). Unless otherwise specified all reagents were of analytical grade and obtained from Aldrich. Densities, conductivities (Wayneker bridge (Automatic LCR 7330) and FTIR (Perkin Elmer 2000) were measured on various faces. Harvested nanoparticles were characterised here by transmission electron microscopy (TEM Philips CM200) using holey carbon grids. Preparations in Bulk Free Solution (a) BaS04: Using BaC12 and Na2S04 (2: 1) in ethanol-water (25:75) Bas04 was prepared by mixing 10 cm3 of BaC12 solution (1.2214 g/dm3, 0.005M) in ethanol water 25:75 and 5 cm3 of a Na2S04 solution (0.7102 g/dm3, 0.005M). (b) Ca3(P04)2/hydroxyapatite: Calcium phosphate in ethanol water 25:75; Ca3(PO4)2 (2:3 mol ratio) [Na3P04 and Ca(N0)2] was prepared by mixing 10 cm3 sodium Na3P04 solution (32.69 g/dm3, 0.086M) in ethanol+water (25:75) with 10 cm3 Ca(N03)~ solution (21.17 g/dm3, 0.129M) in ethanol+water (25:75). Calcium phosphate with 1% copper dopin was prepared as above except that 0.16 g CuC12.2H20 (1% with relation to molar Ca concentration) was added to the Ca(N0)2 (21.17 g/dm3, 0.129M) predissolved in water. (c) BaCe03 and Lac003 were prepared by dissolving the nitrates of Ba and Ce (1 :1) (or nitrates of La and Co) in water and then adding excess ammonium oxalate solution with agitation. The product was a mixed oxalate that was washed well with water, dried and calcined to 873 K.

a

Microemulsion Preparations of Bas04 (a) Bas04 in AOT/isooctane/water: 0.25 cm30f 0.01M BaC12 solution (2.4 g/dm3) and 0.25 cm3 of 0.01M (1.42 g/dm3) Na2S04 solution were added drop-wise separately to 2 cm3 0.1M AOT (45 g/dm3 in isoctane) each under vigorous stirring. The two microemulsions were combined to form Bas04 microemulsion. Bas04 in DDAB/dodecane/water and Cl2EOJdecane/water microemulsions were prepared as described for AOT. Surfactant concentration used was 46.3 g/dm3 in dodecane for DDAB and 36.3 g/dm3 in decane for C12E04 surfactant. Bas04 in HLB2/xylene/propanol was prepared using 0.2 cm3 of each BaC12 solution (O.OlM, 2.4 g/dm3) and NaS04 solution were added dropwise to 2 cm3 SE HLB 2 (10 g/dm3 in xylene) separately. The initial suspension was transformed to a transparent microemulsion by adding 2 cm3 2-propanol to each emulsion.

38

Progress in Nanotechnology: Processing

(b) TX 100/cyclohexane/2-methyl-2-propanol/water: 0.07 cm3 of an aqueous solution of 0.1M BaC12 (24.4 g/dm3) was added dropwise to 5 cm3 of 0.2M TXlOO (387.6 g/dm3) (in cyclohexane and 2-methyl-2-propanol 1O:l ratio). 0.07 cm3 of an aqueous 0.1M Na2S04 (14.2 g/dm3) solution was added dropwise to another 5 cm30f 0.2M TXlOO in cyclohexane and 2-methyl-2-propanol ( 10:1). Mixing these formed water-in-oil microemulsions. On cooling to 268 K two layers were again formed. Densities of two layers were measured (upper layer: 0.78 g.cm”; bottom layer: 0.84 g . ~ m - ~Since ) . the densities of both 2-methyl-2propanol and cyclohexane are 0.78 g . ~ m -the ~ , water and surfactant were assumed to have been in the bottom layer. All Bas04 microemulsions formed were colourless and transparent. Conductivities (see Table I) suggested that these microemulsions were water-in-oil. The presence of water in the bottom layer after lowering the temperature of the TXlOO microemulsions was proved by both conductance and IR. Conductivity of Bas04 focussed TXlOO microemulsion (after cooling to 268 K) the bottom layer was approximately 67 times higher than the upper layer (0.105 mS and 6.700 mS for upper and bottom layers respectively). IR (see Figure 1) showed that the intensity of the 3400 cm-’ H - 0 was 25 times higher in the lower than in upper layers. There was a strong C-H bend at 2922 - 2852 cm-’ in both layers and another strong C-H bend at 1450 cm” at upper layer and a weak one in the bottom layer (i.e it was 0.3 intensity ratio of the upper phase value (see Figure la,b)). Weak peaks of TX-100 were found in the bottom layer at 1524 cm-’ and 1512 cm-’ (see Figure 1).

Table I. Conductivities of microemulsions used in Bas04 preparation Surfactant used PH

HLB 2 7

AOT 6.8

c1 2 E 0 4 6

DDAB 6

TX 100 7.2

Other TX 100 Microemulsion-based Preparations (a) Ca3(PO4)2 : The above TXlOO microemulsion route was repeated using 0.1M of each Ca(N03)2(23.6 g/dm3) and Na3P04 (14.2 g/dm3) (1 :1). Densities for both upper layer (0.79 g.cm3, and bottom layer (0.89 g . ~ m ’ were ~ ) measured. Here again most water was in the bottom layer (see Figure 1). (b) LaCo03: Here three microemulsions were produced: (i) using 0.07 cm30f 0.1M La(N03)3 (32 g/dm3), (ii) 0.1M Co(NO3)3 (29 g/dm3) was added dropwise to 5 cm3 of 0.2M TXlOO (in cyclohexane and 2-methyl-2-propanol 1O:l ratio) and (iii) 0.14 cm3 of 0.1M NH40H was added dropwise to 10 cm3 of 0.2M TXlOO (in cyclohexane and 2-methyl-2-propanol 1O:l ratio). The three microemulsions were mixed together and then cooled at 268 K. Again two layers were formed. The densities of the upper (0.77 g . ~ m -and ~ ) bottom (0.80 g.cm”) layers were measured. FTIR again (see Figure 1) suggests water was mainly in the bottom layer. (c) BaCe03: This preparation used the same method as Lac003 except that 0.lM Ba(N03)~(26 g.dm-3)and O.lM Ce(N03)3 ( 43.4 g.dm-3jwere used. Densities for upper and bottom layers were

Synthesis Methods for Powders

39

found to be 0.79 g . ~ mand - ~ 0.84 g . ~ mrespectively. -~ Again FTIR suggested (see Figure 1) that water was mainly in the bottom layer.

4000

3750

3500 3250

3003 2750

2500

2250

2000

1750

1500

v (ern-')

Figure. 1 FTlR spectra of top and bottom phases in TXlOO microemulsions (after phase inversion at 268 K) when these were used for BaS04, Ca3(PO4)2, Lac003 and BaCe03 preparation

~

RESULTS AND DISCUSSION Similar conductivity and FTIR evidence was seen with the TXlOO microemulsions route to Ca3(PO4)2 (see Figure 1). The conductivity of bottom layer (i.e. 7 mS) was much higher than that of the upper layer (0.12 mS) (see Table 111). Furthermore IR again showed the same water (0-H) stretch at 3500 cm" where the upper:bottom intensity ratio was 0.23 upper to bottom layers. The bottom:upper intensity ratio of the C-H stretch at 1450 cm-' was found to be 0.57. Again similar data were available from the Lac003 and BaCe03 TXlOO microemulsions (see Table 11) after phase inversion at 268 K.

40

Progress in Nanotechnology: Processing

4-BaSO-

Figure 2 shows a micrograph of the Bas04 10-20nm nanoparticles formed in a free aqueous + ethanolic solution. These are clearly aggregated, but show a uniformity of primary

Table 11. FTIR evidence of phase-separated TXlOO microemulsions used for Lac03 and BaCeO3 preparation L~coo~/c~-' 3424 bottom 3357 upper

Stretc h/cm-' 0-H

I

C-H

2928-2853 bottom 2932-2852 upper

%T 2.425

BaCe03/cm-' 3442 bottom 3357 upper

similar

1

%T 3.9

2926-2854 bottom 2922-2852 upper

I

(6

1450 bottom 1451 upper

C-H

0.1

bb

1450 bottom 1451 upper

0.3 1

Table 111. Conductivities of Ca3(PO4)2, BaCe03 and Lac003 emulsion phases before and after phase separation Microemulsion Before phase inversion

Bas04 /m S 3.OOO

Ca3(P04)2/ mS 4.000

BaCe03 / mS 1.300

Lac003 / mS 3.300

After phase inversion

0.104 upper 6.700 bottom

0.1 15 upper 7.400 bottom

0.100 upper 5.600 bottom

0.095 upper 8.100 bottom

I

Synthesis Methods for Powders

1

41

Figure 2. Bas04 nanoparticles produced in free ethanol:H20 solution (scale=SOnm)

Figure 3. Bas04 nanoparticles derived with a) HLB2 microemulsion, b) C12E04 microemulsion, c) AOT microemulsion, d) DDAB microemulsion and e) AOT microemulsion (Scale =100nm). EDX or ED confirmed that these were only BaS04.

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Progress in Nanotechnology: Processing

Figure 3a-e shows TEM micrographs of Bas04 of very different morphologies formed in various transparent microemulsions. HLB2 (see Figure 3a) produced baguette particles (12-28nm x 125300nm), C12EO4 (see Figure 3b) baguettes and thin platelets (25-60nm), AOT (see Figures 3c, 2e) platelets (1 8-37nm) and nanoribbons (21 pm) and DDAB (see Figure 3d) cylindrical bundles of nanoribbons (11-60nm in length and 7-26nm width). Elongated particles with AOT are mentioned in the literature. 14,15 Ellipsoidal particles have been reported before for ZnS with Triton XlOO in the presence of cosurfactant n-butanol.I6 EELS confirmed that these were Bas04 are not artefacts. However, all of those Bas04 nanoparticles were very difficult to harvest. Hence more attention was now paid to TXlOO based microemulsions, where phase separation could be based on phase inversion temperatures. Bas04 from TX100 microemulsions after Phase Inversion Figure 4 shows that a large number of Bas04 nanoparticles were found extensively in the TXlOO microemulsion upper phase (50%) after thermal separation. This suggests that the phase inversion method has delivered a usable method of nanoparticle harvesting.

Figure 4. Bas04 nanoparticles (1 1-29nm) produced in upper layer. This upper phase was water-depleted but produced a high density of harvestable nanoparticles similar to those found in free ethanol-water solutions (see Figure 2). CA(PO&

Results and Discussion

Gelatinous Ca3(PO4)2 formed in free solution contained 25-62nm fibrons (see Figure 5a). Densities and conductivities (see Table 111) of both upper and bottom layers of the TXlOO microemulsion containing Ca(NO3)2 and Na3P04 after phase inversion at 268 K proved that the particles were again in the upper organic layer (see Figure 4b). TEM shows that dispersed 27107nm cubic and 80-2 13nm hexagonal Ca3(PO4)2 nanoparticles are formed and harvested from the microemulsion. This is promising for injectable bone.

Synthesis Methods for Powders

43

Figure 5. Ca3(PO& a) formed in free solution, b) in TXlOO microemulsion cubic and hexagonal (scales=50 nm).

BaCe03 And Lac003 Results and Discussion Figures 6a and b shows micrographs of BaCe03 and LaCoOs at 873 K formed by free bulk solution; rods were much larger than nonoparticles. Figures 7a and 7b clearly show that for both Lac003 and BaCe03 very highly dispersed nanoparticles and nanochains were harvested in the upper phases of TXlOO microemulsions on phase segregation at 268 K. Since the bottom layers had conductivities 8.1 mS and 5.6 mS for Lac003 and BaCe03 respectively (see Table 111) that were 85 and 56 times higher than corresponding upper layers, the upper layer was again thought to be organic rich. A comparison of Figures 6 and 7 looks very promising for ceramic production nano-particles.

44

Progress in Nanotechnology: Processing

Figure 6 SEM of a) BaCe03 particles and b) Lac003 formed in free solution (scale 30ym) CONCLUSIONS There is much interest in the design of materials using nanoparticles (e.g. Ca3(P04)2hydroxyapatite 17*18 with and without doping 19, Bas04 20,2 I ,22 , Lac003 23 and BaCe03 24725 and nanocomposites for a whole range of applications e.g. from synthetic bone 26 to ceramics 27.

Figure 7 TEM of a) Lac003 in a TXlOO microemulsion after phase inversion nanoparticles (size 10-21nm) and nanothreads (size 43-108nm) were formed and b) BaCe03 in a TXlOO microemulsion after phase inversion nanparticles (size 10-20nm) and nanothreads (size 40- 121nm) were formed.

Synthesis Methods for Powders

45

This paper shows that the continuous synthesis of well dispersed nanoparticles of BaS04, Ca3(PO4)2, Lac003 and BaCe03 etc can be achieved using water-in-TX 100/cyclohexane/2methyl-2-propanol microemulsions with harvesting after phase separation of an organic-rich top phase on cooling to 268 K. The benefits over free solution routes is most clear for Lac003 and BaCe03 and this bodes well for ceramic production. A range of particle sizes and morphologies (particles, platelets, ribbons, threads and baguettes (as previously seen22)) can be produced for inclusion in products in a range of applications. The different morphologies of particles is determined by the relative rate of growth along different crystallographic directions. Mann28729 reported that faces perpendicular to the fast direction of growth have smaller surface areas and slow growing faces. Therefore fast growth along one specific axis produces a needle-shaped morphology, whereas platletlikes were produced while growth along two directions. BaCe03 and Lac003 nanoparticles formed are primary nuclei which are involved in the growth of the n a n ~ c h a i n s . ~The ~ ' ~ 'formation of Bas04 elongated nanoribbons was related to interweaving of surfactant bound to the long axis in adjacent nanochains (see Figure 2e).30 ACKNOWLEDGEMENTS The authors thank EPSRC and BP/Viaton for funding for SA via a studentship, L. Courtney (UniS) for SEM and A. Reynolds (Brunel University) for some TEM micrographs. REFERENCES 'J. H.Clint, I. R.Collins, J.A.Williams, B.H.Robinson, T.F.Towey, P.Cajean and A. Khanlodhi Farad. Discuss. 9 5 2 19-233,(1993) 2M.Cao, X.Y.He, X.L.Wu and C.W.Hu, Nanotechnology, 16,2129-2133,(2005) 3V.Uskokovic, D.Makovec and M.Drofenik Curr.Res.Adv. Mater.Processes MuterSci. Forum 494,155- 160,(2005) 4V.Uskokovic and M.Drofenik Prog.Adv.Muter.Processes Materials Science, Forum, 453454,225-230,(2004) 5C.Pithan, Y.Shiratori, A.Magrez, S.B.Mi, J.Dornseiffer and R.Waser J. CeramSoc. Japan 114(1335),995-1000,(2006) 6L. Qi, J.Ma, H.Cheng and Z.Zhao, Coll. Surf: 108A, 117-126,(1996) 7B.Niemann, F.Rauscher, D.Adityaawarman, A.Voigt and KSundmacher, Chem. Eng. Proc. 45,917-935,(2006) *T.Nishimi and C.A.Miller Langmuir 16,9233-9241,(2000) 'S.Clark, D.I.Fletcher and Y.Xilin Lungmuir 6,1301-1309,(1990) 10 J.D.Hopwood and S.Mann Chem.Muter. 9,1819-1828,( 1997) "N.I.Ivanova, D.S.Rudelev, B.D.Summ and A.A.Chalykh Co1l.J. 63,7 14-71742001) 12J.Cai,KLaubernds, F.S.Galasso, S.L.Suib, J.Liu, X.F.Shen, H.R.Kunz and J.M.Fenton J.Amer. Ceram.Soc. 88,2729-2735 ,(2005) 13P.A.Sermon,N.Mason McLellan and 1.R.Collins Cryst.Eng.Comm. 6,469-473, (2004) 14M.P.Pileniand J.Tanori, Adv.Mater. 7, 862-864,( 1995) "M. Li and S. Mann, Langmuir 16,7088-7094,(2000) '6T.Charinpanitkul, A.Chanagu1 and J.Dutta, Sci. Tech. Adv. Mat. 6,266-27 1,(2005) 17 S.C.Lee, H.W.Choi, H.J.Lee, K.J.Kim, J.H.Chang, S.Y.Kim, J.Choi, K.S.Oh and Y.K.Jeong J.Muter.Chem. 17,174-180,(2007)

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Progress in Nanotechnology: Processing

"N.Pramanik, P.Bhargava, S.Alam and P.Pramanik Polymer Compos. 27,633-64 1 ,(2006); A.Monkawa, T.Ikoma, S.Yunoki, K.Ohta and J.Tanaka Mater.Lett. 60,3647-3650,(2006); Y.M.Mao, D.X.Li, H.S.Fan, X.D.Li, Z.W.Gu and X.D.Zhang Mater.Lett. 61,59-62,(2007) ''Z.Evis J. CeramSoc.Japan 114,1001 - 1004,(2006) 20 G.G.Chen, G.S.Luo, J.H.Xu and J.D.Wang Powd.Techno1. 153,90-94,(2005); J.K.Liu, Q.S.Wu and Y.P.Ding Chem.Res.Chinese Univ. 21,243-245,(2005) 21 M.H.Qu, Y.Z.Wang, Y.Liu, X.G.Ge, D.Y.Wang and C.WangJ.Polymer Sci. 102,564570,(2006) 22 H. Bala, W. Fu, J. Zhao, X. Ding, Y. Jiang, K. Yu and Z. Wang, Colloids and Surfaces 252A,129-134,(2005) 23X.L.Cui,Q.H.Yang and X.X.Fu J.Rare Earths 21(Suppl S), 124- 126,(2003); M.Chen, Y.W.Wang and X.M.Zheng Chin.J.Inorg.Chem. 19,1145-1149,(2003) 24 D.W.Lee, J.H.Won and K.B.Shim, Mater.Lett. 57,3346-335 1,(2003) 25K. Ouzaouit, A. Benlhachemi, H. Benyaich, J. P. Dallas, S. Villian, J. A. Musso and J. R. Gavarri, M J. Condensed Mater. 7,94-97,(2006) 26J. A. Wihmhurst, R. A. Brooks and N. Rushton J. Bone Joint Surgery (British) 83B,278282,(2001) 27M.F. M. Zawrah and N. M. Khalil Ceramics Intn. 27,309-3 14,(2001) 28S.Mann,Angew. Chem. Int. Ed. 39,3392-3406,(2000) 29 S. Sadasivan, D. Khushalani and S. Mann, Chem. Mater. 17,2765-2770,(2005) 30 G. D. Rees, R. Evans-Gowing, S. J. Hammond and B. H. Robinson, Langmuir 15,19932002,(2002)

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Nanocrystalline Scandia Powders Via Oxalate Precipitation: The Effects of Solvent and Solution pH Zhimeng Xiu,' Ji-Guang Li,',t*5Xiaodong Li,$ Di Hue,$ Xudong Sun,STakayasu Ikegami? and Takamasa Ishigaki5 'Key Laboratory for Anisotropy and Texture of Materials, School of Materials and Metallurgy, Northeastern University, Shenyang 110004, China h a n o Ceramics Center, National Institute for Materials Science, Ibaraki 305-0044, Japan

Nanocrystalline Sc2O3 was synthesized via precipitation with scandium nitrate as the scandium source and oxalic acid as the precipitant, with investigations into the influences of solution pH and solvent type on powder properties. Three types of scandium precursors (normal oxalate, double oxalate, and oxyhydroxide) have been obtained from aqueous media by adjusting pH of the precipitant solution with ammonia water. AU the precursors, however, yield severely aggregated Sc2O3 powders of low surface areas. Replacing water with pure ethanol as the solvent was shown to be able to ield Sc203 nanopowders of high specific surface area (35.7 mY/g) and good dispersion. The mechanisms of solution pH and solvent type on precursor properties were also discussed. I. Introduction

(Sc203)belongs to the cubic space group Ia3 with a reported lattice constant of 0.979 nm, a theoretical density of 3.86 g/cm3, and a melting point of -2430°C.' It is well known

S

CANDIA

as a stabilizer for Zr02, an important ceramic material for structural and electrical applications. It has also been used as a sintering aid for super-strong Si3N4ceramics2and as a dopant in gadolinium gallium garnet (GGG) solid-state laser^.^ Recently, Sc203was shown to hold great potential as a damage-resistant and high-reflection material in light-emitting diodes and in high-power pulsed UV lasers due to its high bulk refractive index (-2.0) and wide bandgap (-5.7 eV).4 Previous work on the reparation of Sc203nanopowders via carbonate precipitation! h droxyl precipitation: urea-based homogeneous precipitation, Y and sol-gel* proved that the Sc3+ ions undergo much more extensive hydrolysis in an aqueous solution than any other rare-earth elements do, exhibiting chemical properties close to those of A13+. It has also been shown that complex anions, such as SO:-, affect appreciably the precipitation mechanism, chemical composition, and morphologies of the resultant particles7 Among the above-mentioned synthetic routes, the carbonate precipitation method' with carefully controlled Sc3+/ammonium bicarbonate (precipitant) molar ratio and the hydroxyl precipitation methodology6 in the presence of particle-morphology engineering SO:- yield the best results. Oxalate precipitation is the most widely used technique in the commercial production of scandia as well V. Thangadurai-ontributing

editor

Manuscript No. 23378. Received June 26, 2007; approved August 27, 2007. This work was partially supported by the Program for New Century Excellent Talents in University (NCET-250290), the National Science Fund for Distinguished Young Scholars (50425413). and the National Natural Science Foundation of China (50672014). 'Author to whom correspondence should be addressed. e-mail: LI.Jiguang~nims.go,jp

Synthesis Methods for Powders

as other rare-earth oxides. The method has the advantage of the nongelatinous nature of the oxalate precursor, which makes the washing procedure readily to perform, but frequently has the shortcomings of large particle size, irregular particle shape, and extensive aggregation of the primary particles. It has been widely observed that the oxides processed wet chemically show properties (especially aggregation and particle morphology) generic to their precursors. Careful control of precursor properties is therefore indispensable to make oxide powders of desired characteristics. In this work the oxalate precipitation technique has been used to prepare nanocrystalline Sc2O3 powders. The effects of solvent for reactants and solution pH on powder properties were investigated, and the optimal conditions that yield dispersed Sc203 nanopowders were outlined. 11. Experimental Procedure

The precursors for Sc203were produced via precipitation using oxalic acid dihydrate (H2C204.2H20, ultrahigh purity, Wako Pure Chemical Industries Ltd., Osaka, Japan) as the precipitant. The scandium source, scandium nitrate, was made by dissolving a commercially available Sc203 powder ( > 99.9% pure, Soekawa Chemicals, Tokyo, Japan) in concentrated nitric acid (analytical grade, Wako Pure Chemical Industries). A watersoluble dry scandium nitrate powder is then obtained by evaporating the salt solution to dryness at 80°C. Precipitation was performed at room temperature on a magnetic stirrer by dripping at 5 mL/min, 300 mL of the scandium nitrate solution (0.1M) into an equivolume of oxalic acid solution (0.3M) under mild stirring. The solvent was changed from pure water to pure ethanol to investigate its effect on powder properties, and pH adjustment was made with ammonia solution (27%, analytical grade, Wako Pure Chemical Industries). After the completion of precipitation, the resultant suspension is homogenized at room temperature for 2 h and is then filtered via suction filtration. After repeated washing via ultrasonication and filtration with ethanol (two times), pure water (two times), and finally ethanol (one time), the precursor is dried at room temperature with flowing nitrogen (200 mL/min) and is then thermally decomposed in a tube furnace to yield Sc203powders. Chemical compositions of the precursors were determined via elemental analysis following standard procedures, with the scandium content determined via inductively coupled plasma spectrophotometry, carbon (from C20:-) assayed on a simultaneous carbon/sulfur determinator (Model CS-444LS, LECO, St. Joseph, MI), NH4f by the distillation-titrimetric method, and NO; via spectrophotometry (Model Ubest-35, Japan Spectroscopic Co. Ltd., Tokyo, Japan). Thermogravimetry (TG, Model TAS-200, Rigaku, Tok 0,Japan) of the dried precursor is made in Rowing air (100 cmY/min) with a heating rate of 10"C/min; phase identification is performed via X-ray diffractometry

49

Table I. Elemental Contents (wt%) of the Precipitates and their Chemical Formula Solution pH

1.4 3.0 4.5 7.0 8.0

sc

C

NO;

N Ha

Approximate composition

20.8fO.l 16.4+0.l 16.3k0.l 16.4fO.l 50.7 Ifr0.l

17.3f0.1 17.3 fO.l 17.5 fO.l 17.3 kO.1 N.D.

N.D. N.D. N.D. N.D. N.D.

N.D. 6.4k0.1 6.3f0.1 6.4 f 0.2 N.D.

sC2(c204)3 .4HZ0 (NH4)Sc(C204)2 ' 2H20 (NWSc(C204)2 '2H20 (NH4)Sc(C204)2' 2H20 ScOOH . 0.6H20

(XRD, Model PWI 800, Philips Research Laboratories, Eindhoven, the Netherlands) operating at 40 kV/50 mA using nickel-filtered CuKu radiation; specific surface area of the Sc2O3 powder is measured by the Brunauer-Emmett-Teller (BET) method on an automatic surface area analyzer (Model 4201, Beta Scientific Corporation, Albertson, NY) via nitrogen adsorption at 77 K; particle morphology is observed via fieldemission scanning electron microscopy (Model s-5000, Hitachi, Tokyo, Japan). 111. Results and Discussion

( I ) The Effects of Solution p H The aqueous solution of oxalic acid (0.3M) used as the precipitant has a pH of 1.4, which can be modified by neutralization with ammonia water. It was observed that the precipitates obtained at pH values up to 7.0 are very fluffy and readily washable while those synthesized at pH = 8.0 or above are gelatinous, indicating that solution pH affects significantly precursor properties. Elemental contents of the dried precursors obtained under some typical pH are summarized in Table I, from which it can be seen that chemical composition of the precursor changes from hydrated normal oxalate to hydrated double oxalate of ammonium and scandium and then to scandium oxyhydroxide. Precursor composition is the result of competition among the ionic species in the solution during combining with Sc3+,and the final composition should assure that the precipitate has the smallest solubility product under given conditions. This explains the varied precursor types under different solution pH. NH4f is known to complex with rare-earth (RE3') ions and the tendency increases with decreasing RE3' size.' As Sc3+ has the tiniest size

-

Fig. 1. X-ray diffractometry (XRD) patterns of the precursors precipitated from aqueous media at various pH values and that of the Sc20z powder pyrolyzed at 650°C from the precursor obtained with pure ethanol as the solvent. The standard XRD pattern for cubic Sc203is included for comparison.

50

among RE3+ (e.g., ri+ = 0.0745 nm while r:: = 0.0861 nm for sixfold coordination)," extensive formation of Sc3+/NHl complexes is expected when the precipitant solution was adjusted with ammonia water to pH = 3.C7.0, leading to the formation of double oxalates. Under higher pH of > 8.0, the Sc3+ions are extensively hydrolyzed to [SC(H~O).,(OH)-~]~-~" species (x+y = 6):' from which ScOOH is formed via olation and oxolation.' The formation of hydrated ScOOH rather than oxalate or double oxalate at pH > 8.0 suggests that oxyhydroxide is the least soluble within this pH range. Figure I shows XRD patterns of the resultant precursors after drying. The powders prepared at pH up to 7.0 are all well crystallized while those obtained at pH = 8.0 or above are amorphous. In addition, the powders synthesized in the pH range 3.0-7.0 show the same XRD patterns, suggesting they are identical in composition and crystal structure. JPCDS data files were found for only two forms of scandium normal oxalates of S C ~ ( C ~ O 6H20 ~ ) ~ .(No. 33-1 156, No. 27-0592) and SC~(C~O 8H20 ~ ) ~ .(No. 28-101 I), while those for double oxalates of scandium and ammonium are not available. XRD patterns of the crystalline phases obtained in this work, however, do not match the data file of any known scandium compound. R E oxalates are known to be formed with various numbers of hydration water and their crystal structures vary accordingly." This might be the reason as to why the normal oxalate S C ~ ( C ~ O 4H20 ~ ) ~obtained . in this work shows an XRD pattern different from those of S C ~ ( C ~ O 6H20 ~ ) ~ .and Sc2(C204)3. 8H2O. Thermal decomposition processes of the normal oxalate (pH = 1.4) and the double oxalate (pH = 4.5) are compared in Fig. 2. That of the scandium oxyhydroxide is not given here as it has been studied in detail previously and was shown to yield Sc2O3 upon calcination at temperatures up to -600"C-' Both the oxalate precursors undergo multistep decompositions to yield ScZO3 at temperatures up to -620°C. The final weight losses observed from the TG curves are close to the values cal4H20 culated from the complete decomposition of S C ~ ( C ~. O ~)~ (67.7%) and (NH4)S~(C204)2. 2H20 (74.9%).

Fig.2. Thermal decomposition behaviors of the normal oxalate (dashed line) and the double oxalate (solid line) precursors. Data were taken in flowing air (100 mL/min) with a constant heating rate of 10"C/min. Progress in Nanotechnology: Processing

Fig. 3. Fieldemission scanning electron micrographs showing particle morphologies of the prccursors, with (a) the normal oxalate made at pH = 1.4, (b) the double oxalate synthesized at pH = 4.5, (c) the oxyhydroxide obtained at pH = 8.0, and (d) the normal oxalate made with ethanol as the solvent. The precursor powders shown in (aHc) were all obtained from aqueous media.

Morphologies of the three types of precursors are compared in Figs. 3(aHc). Both the normal oxalate (Fig. 3(a)) and the double oxalate (Fig. 3(b)) precursors are composed of elongated particles of irregular shapes with sizes measuring up to tens of micrometers. The observed particles are actually aggregates of plate-like primary crystallites, as revealed by high-magnification observations (the insets). However, the two types of precursors exhibit different characteristics: the primary crystallites in the normal oxalate are larger (up to 2 pm) and densely packed while those in the double oxalate are significantly finer (up t o 300 nm) and much less aggregated. The oxyhydroxide precursor (Fig. 3(c)) shows irregularly shaped particles of a wide size distribution, which is directly resulted from hand crushing the dried glassy precursor with mortar and pestle. Morphologies of the precursor particles retain to Sc203after calcining the precursors at 650°C for 2 h, showing the “generic” relationship mentioned earlier (Figs. 4(at-(c)). Obviously, the ScZO3 particles are densely sintered nanocrystallites (Fig. 4(c) and the inset of Fig. 4(b)) with average sizes of -23, 19, and 53 nm, as assayed from the Scherrer equation, for the powders shown in Figs. 4(aHc), respectively. BET analysis found specific surface areas of 0.43, 0.77, and 1.12 m2/g, corresponding to average particle sizes of 3.62, 2.02, and 1.39 pm for the three Sc2O3 powders demonstrated in Figs. 4(a), (b), and (c), respectively.

-

( 2 ) The Effects of Solvent The oxalic acid precipitant and the scandium nitrate both show high solubility toward ethanol, allowing a convenient study of Synthesis Methods for Powders

Fig. 4. Fieldemission scanning electron micrographs showing morphologies of the SczO3 powders obtained by calcining at 650°C for 2 h the precursors shown in Figs. 3(a), (b), (c). and (d), respectively.

the solvent effects on precursor properties. It was found that under identical precipitation conditions, a crystalline normal oxalate of scandium, whose XRD pattern and particle morphology resemble those of the powder synthesized at pH = 1.4 from aqueous solutions, was resulted from an ethanol solution of oxalic acid while an aqueous solution of scandium nitrate (in this case the ethanol/water volume ratio is 50/50 in the resultant suspension). An amorphous precipitate was obtained when pure ethanol was used for both the reactants. Elemental analysis found 22.2k0.1 wt% of Sc and 17.9k0.1 wt% of C for this amorphous precursor, suggesting an approximate chemical formulae of S C ~ ( C ~ O ~ ) ~ . That ~ . ~ is, H ~the O .precursor is still a normal oxalate but has less hydration water. The precursor particles derived from ethanol solvent (Fig. 3(d)) differ dramatically from those synthesized with water, in that they are significantly finer (up to 500 nm) and are much less aggregated. Calcining this ethanol-derived amorphous precursor at 650°C for 2 h yields Sc203particles of much better dispersion (Fig. 4(d)). Although agglomerates are still present, they have much smaller sizes of up to -400 nm. The XRD pattern of this resultant Sc2O3 powder is compared in Fig. 1 with the standard data file (JCPDS: No. 5-629), from which it can be seen that all the characteristic diffraction peaks corresponding to cubic-structured Sc203 have appeared. X-ray line broadening analysis on this ScZO3 powder yields an average crystallite size of 21.7 nm, which is close to the values determined for the Sc203powders synthesized from aqueous solutions. The ethanol-derived powder, however, has a much higher specific surface area of 35.7 m2/g (equivalent particle size: 43.5 nm), suggesting greatly improved dispersion of the Sc203particles with pure ethanol as the solvent for precursor synthesis. The differed properties are direct reflections of the changed solubilities of the oxalate precursors in different media.

-

51

The higher solubility of scandium oxalate in water allows fast particle growth via Ostwald ripening, leading to well-crystallized primary particles of large size and severe aggregation (Figs. 3(a) and (b)). On the other hand, the solubility of scandium oxalate is expected to dramatically reduce in ethanol, greatly retarding the mass transfer, and thus yielding much finer particles of an amorphous nature.

IV. Conclusions The synthesis of nanocrystalline Sc203 powders have been made via precipitation with scandium nitrate as the scandium source and oxalic acid as the precipitant. With water as the solvent, chemical composition of the precursor was found to change with increasing pH of the precipitant solution (adjusted with ammonia water) from normal oxalate (pH = 1.4) to double oxalate (pH = 3.0-7.0) and then to oxyhydroxide (pH 2 8.0). Ethanol was shown to be a better solvent than water, leading to Sc203 nanopowders of high specific surface area (35.7 m2/g, equivalent particle size of 43.5 nm) and good dispersion.

References ‘C. Chen. B. M. Wanklyn, and P. Ramasamy, “The Flux Growth of Scandium Oxide Crystals,” J. Cryst. Growrh, 104, 672-6 (1990).

52

‘D. S. Cheong and W. A. Sanders, “High Temperature Deformation and Microstructural Analysis for Silicon Nitride-Scandium (111) Oxide,” J. Am. Ceram. Soc., 75 [I21 33314 (1992). ’8. Keszei. J. Paitz, J. Vandlik. and A. Suveges, “Control of Nd and Cr Concentrations in Nd,Cr:Gd3Ga3012 Single Crystal Grown by Czochlaski Method,” J. Cryst. Growth, 226, 95-100 (2001). and references therein. 4F. Rainer, W.H. Lowdermilk. D. Milam, T. T. Hart, T. L. Lichtenstein, and C. K. Carniglia, “Scandium Oxide Coatings for High-Power UV Laser Applications,” Appl. Opt., 21 1201 3685-8 (1982). 5J.-G. Li, T. Ikegami, T. Mori. and Y. Yajima, “Wet-Chemical Routes Leading to Scandia Nanopowders,” J. Am. Cerum. Soc., 86 [9] 1493-9 (2003). 6J.-G. Li, T. Ikegami, T. Mori, and Y. Yajima, “Scz03Nanopowders Via Hydroxyl Precipitation: Effects of Sulfate Ions on Powder Properties,” J. Am. Ceram. Soc., 87 [6] 1008-13 (2004). 7J.-G. Li. T. Ikegami, T. Mori, and Y. Yajima, “Monodispered Sc,O, Precursor Particles Via Homogeneous Precipitation: Synthesis, Thermal Decomposition, and the Effects of Supporting Anions on Powder Properties,” J. Muter. Re.s., 18 [S] 1149-56 (2003). 8D. Grosso and P. A. Sermon, “Scandium Oxide Nanoparticles from Sol-Gel Chemistry,” J. Muter. Chem.. 10, 35943 (2000). ’D. 1. Ryahchikov and E. A. Terentyeva. “Recent Soviet Research on the Chemistry of Rare Earth Complexes”; pp. 139-SO in Progress in the Science and Technology ofthe Rare Earth I, Edited by L. Eyring. Pergamon Press, New York, 1964. ‘OR. D. Shannon, “Revised Effective Ionic Radii and Systematic Studies of Interatomic Distances in Halides and Chalcogenides,” Acta Crystollogr., A32, 75147 (1976). “L. C. Thompson, “Complexes”; pp. 209-97 in Handbook on the Physics and Chemistry of Rare Earths. Vd. 3, Edited by K. A. Gschneidner Jr., and L. Eyring. North-Holland Physics Publishing, Amsterdam, the Netherlands, 1979. ”G. Adachi, The Sciences of Rare-Earths. Kagakudojin Press, Kyoto, Japan, I999 (in Japanese). 0

Progress in Nanotechnology: Processing

A Pulse Combustion-Spray Pyrolysis Process for the Preparation of Nano- and Submicrometer-Sized Oxide Particles W. Widiyastuti,' Wei-Ning Wang,' Agus Purwanto,$ I. Wuled Lenggoro,$ and Kikuo Okuyamat3$ tDepartment of Chemical Engineering, Graduate School of Engineering, Hiroshima University, Higashi-Hiroshima 739-8527, Japan hstitute of Symbiotic Science and Technology, Tokyo University of Agriculture and Technology, Tokyo 184-8588, Japan

The preparation of nano- and submicrometer particles using an integrated pulse combustionspray pyrolysis process was investigated for the first time. Zinc oxide nanoparticles with average sizes ranging from 20 to 40 nm were synthesized using a salt precursor with a lower decomposition temperature than the pulse combustion (PC) temperature. Particles of a doped-type oxide, gadolinium oxidedoped with europium, were also produced from a precursor, whose decomposition temperature was higher than the P C temperature, and were observed to be ellipsoidaltoroidal in shape and submicrometer in size. The high-intensity fluctuating pressures generated by P C and the precursor decomposition temperature are proposed as the reasons for the formation of either nanoparticles or aspherical particles. In addition, hot gas fluctuating flow in P C enhanced energy transfer, resulting in more crystalline particles, as compared with particles produced by conventional spray pyrolysis.

I. Introduction

T

HE

spray pyrolysis (SP) method can be used for the production of high-purity particles in the nano- to micrometersized ranges with a continuous operation and a high production rate.' Conventional SP generally produces spherical particles in the submicrometer and micrometer-sized ranges. The spraypyrolyzed particles consist of aggregated nano-sized crystallites derived from nucleation and growth from dissolved precursors inside droplets. Modifications of the SP method to prepare nanoparticles have been used by our group, such as SP under low-pressure conditions: flame SP,' salt-assisted SP; and polymer-assisted SP.' In this study, a new method is introduced that integrates pulse combustion (PC) and SP, called pulse combustion-spray pyrolysis (PCSP). In PC, the momentum, energy, and mass transfer rates are enhanced by the generation of high-intensity acoustic waves and unsteady gas flow rates.637PC has been used in drying processes because it can increase the drying rates significantly. The advantages of PC in spray drying have been reported by Zbicinski et d 8and Kuts et d9It has also been used for drug production, as reported by Liang et al."SP differs from spray drying, in that SP involves evaporation, drying, and thermal decomposition or reaction, while spray drying only involves evaporation and drying. The use of PC for oxide particle synthesis has not been S. Bhandharkar-contributing editor

Manuscript No. 23248. Received May 23,2007; approved July 31, 2007. This work was supported in part by Grants-in-Aid from the Ministry of Education. Culture, Sports, Science and Technology (MEXT) of Japan, Japan Society for the Promotion of Science (JSPS) and Japan Science and Technology (JST) Agency. 'Author to whom correspondence should be addressed. e-mil: okuyama@hiroshima-

u.ac.jp

Synthesis Methods for Powders

studied previously, and the effect of pressure oscillations on the particle synthesis process needs to be studied further. A pulse engine mainly consists of three parts, i.e. the inlet, the combustor, and the tail or exhaust part, as shown in Fig. 1, inset." A combustible mixture is introduced into the combustor, which is subsequently ignited by a spark plug. During the combustion process, the pressure inside the combustor increases and leads to the movement of hot gas toward the tail. This movement produces a negative pressure in the combustion chamber; thus, the hot exhaust gas from the tail flows back into the combustor, resulting in precompression of the mixture inside the combustor. Hot and high-pressure gas comes in contact with the combustible mixture and leads to ignition. The process is repeated, producing a fluctuating flow with a high acoustic wave intensity. The pressure and temperature in the combustor therefore vary with time. To evaluate the effect of PC on the particle characteristics produced, particles of two different materials were prepared: zinc oxide (ZnO) and gadolinium oxide-doped with europium (Gd203:Eu3+).ZnO was selected because it has a lower decomposition temperature than that of the average flow temperature generated by PC," and because it is an often-used semiconductor.I3 ZnO nanoparticles were prepared by many methods, including SP methods, such as salt-assisted SP? flame SP,I4 and low-pressure SP." However, salt-assisted SP needs an additional treatment after spraying to remove the salts. Flame spray and low-pressure SP are relatively energy consumption processes. In this work, a PCSP was introduced, which is considered to be a promising method because it is a one-step process with a relatively low energy consumption. Gd203:Eu3+, which is well known as a red-emitting phosphor, was selected for comparison. It has a higher thermal decomposition temperature than that of the average flow temperature generated by PC.I6 The particle size, morphology, and crystallinity obtained for both ZnO and Gd203:Eu3+ particles produced by PC, PCSP, and conventional SP are presented and discussed in this work.

11. Experimental Procedures The experimental setup is schematically shown in Fig. 1. It consisted of a pulse generator, a nebulizer, a furnace, and a particlecollecting system. The precursor was sprayed using an ultrasonic nebulizer (1.7 MHz, NE-U17, Omron Healthcare Co. Ltd., Tokyo, Japan), with air as the carrier gas. The droplet distribution and the number concentration were measured by a laser diffraction technique (Spraytec, Malvern Instruments Ltd., Malvern, U.K.). The average volume diameters of the droplets of 0.1M aqueous zinc acetate and 0.1M aqueous gadolinium nitrate sprayed by a nebulizer were 4.22 and 4.38 pm, respectively. The droplets carried by an airflow of 10 L/min were prevented from evaporating and precipitating using a double pipe cooling system before entering the furnace. An electrostatic precipitator was used to collect the particles. 53

Fig. 1. Experimental setup of pulse combustion (PC) integrated with spray pyrolysis. Inset: Operating principle of the PC system

A PC system was provided by Pultech Co. (Kobe, Japan). 0.67 L/min of propane and 20 L/min of air were used as a fuel in the PC system, which was recommended as an optimum condition. This system produced 155 dB at a frequency of 700 Hz. To maintain a similar residence time in the furnace, a vacuum pump equipped with a flow meter controller was used to maintain the flow at 35 L/min (GCD-l36X, Ulvac Kiko Inc., Kanagawa, Japan). The temperature at the outlet region of the pulse combustor tail was recorded using a data logger (midi logger GL200, Graphtec Corp., Yokohama, Japan) every second. Carried by combustion flow, the droplets vertically passed through the furnace. A commercial computational fluid dynamics code (FLUENT 6.2, Fluent Asia Pacific Co., Ltd., Tokyo, Japan) was also used to analyze the steady-state temperature profile inside the furnace and the combustion process. Three systems were used for comparison: PC, in which the energy source only arose from pulse combustion; SP, in which the energy only arose from a heated furnace; and PC integrated with SP (PCSP). For precursor preparation, zinc acetate dihydrate (99%, Wako Pure Chemical, Tokyo, Japan) was dissolved in ultrapure water, resulting in 0.1M of zinc acetate solution as a precursor. 0.1M gadolinium nitrate precursor was also prepared by diluting gadolinium nitrate hexahydrate (99%, Wako Pure Chemical) with ultrapure water. Europium nitrate with a concentration of 10 at.% of Gd(NO& was mixed with the gadolinium solution to determine the PC effect on the luminescence characteristics. Thermogravimetric (TG) and differential thermogravimetric (DTG) analyses (TG/DTA 6200, SII Exstar 600, Seiko Instrumentation Inc., Tokyo, Japan) were used to analyze the precursors. Fifteen miiligrams of zinc acetate dihydrate and gadolinium nitrate hexahydrate without pretreatment were heated at a 10"C/min heating rate with 200 mL/min of air flow for thermal analysis measurements. Zinc acetate dihydrate and gadolinium nitrate hexahydrate were placed in an Al cell and a Pt cell in the temperature ranges of 30"400"C and 30"-1000°C, respectively. The crystallinity of the prepared particles was char54

acterized by an X-ray diffractrometer (XRD, RINT 2200 V, Rigaku, Tokyo, Japan) using nickel-filtered CuKi radiation (h = 1.54 A) at 40 kV and 30 mA. Their morphology was characterized by a field-emission scanning electron microscope (FE-SEM S5000, Hitachi, Tokyo, Japan) operated at 20 kV. A field emission transmission electronic microscope (FE-TEM, JEM-3000F, JEOL, Tokyo, Japan) operated at 300 kV was also used to analyze the crystal structure and morphology. The photoluminescence (PL) intensity and surface area were measured to characterize Gd203:Eu3+.PL spectra were measured using a luminescence spectrometer (RF-5300PC, Shimadzu Corp., Kyoto, Japan) with a xenon lamp at an excitation wavelength of 254 nm. The optical spectrum was measured at a resolution of 0.2 nm at room temperature. The specific surface area of Gd203 powder was measured by a nitrogen adsorption apparatus using the Brunauer-Emmett-Teller (BET) method (Belsorp 28SA, Be1 Japan Inc., Osaka, Japan). 111. Results and Discussion

The velocity fluctuations are the main factors that enhance the heat transfer coefficient, which leads to an increase in the evaporation rate. Increasing the instantaneous velocity leads to an increase in the heat transfer coefficient. Furthermore, the presence of an oscillation flow of dry and hot gas of air due to PC led to an enhancement of the evaporation rate. Dubey et al." conducted an experiment wherein the sinusoidal acoustic field influenced the instantaneous droplet velocity, which is an important variable for determination of droplet evaporation. PC also caused fluctuations in temperature. The temperature measurement at the outlet of the PC tail showed that the temperature fluctuated between 300" and 350"C, as is shown in Fig. 2. The steady-state simulation results of the temperature distribution of SP and PC are depicted in Fig. 3. In SP, a parabolic temperature profile in the radial direction was present, with the minimum temperature along the furnace axis. On the Progress in Nanotechnology: Processing

(Zn04(CH3C02)& as measured previously by Ghule et ~ 1 . ' ~ Therefore, the total weight loss due to dehydration and thermal decomposition processes of zinc acetate dihydrate up to 290°C was 76.8%. Figure 4(b) shows the TG-DTG analysis of a gadolinium nitrate precursor. The dehydration of the precursor began at 9 8 T , indicated by the first peak of the DTG curve. The subsequent three peaks observed indicated phase transitions to form Gd203 up to a constant weight at around 700°C. The total weight loss due to Gd203formation was 59.8%, which agreed well with the predictions based on the dehydration process (23.9%) and the decomposition reaction (35.9%). TG-DTG analysis showed that the zinc acetate precursor to form ZnO had a lower decomposition temperature (at 290°C) than that of the average temperature of PC (around 300"350°C). On the other hand, the gadolinium nitrate precursor to form GdzO3 had a higher decomposition temperature (at 700°C) than that of the average temperature of PC.

F 500 1000 2000 2500

200 0

1500

Time [s]

Fig.2. Temperature fluctuation of combustion flow generated in a pulse combustion system at the outlet of the combustor tail.

other hand, hot gas flow released by PC led to a temperature distribution in the furnace almost similar to that at the furnace inlet. This indicated that energy transfer was enhanced by the PC system.

( I ) Thermal Analysis of Precursors Figure 4(a) shows the TG and DTG results corresponding to the percentage weight loss of zinc acetate dihydrate due to heat treatment. The TG curve shows that the dehydration leading to the formation of anhydrous zinc acetate occurs around 93°C. A 16.4% weight loss was observed, which agreed well with the dehydration calculation based on stoichiometry. On increasing the temperature up to 251"C,a 46.5% weight loss occurred corresponding to ZnO formation according to stoichiometric calculation. Increasing the temperature further up to 290°C resulted in a 13.9% weight loss. This weight loss was due to additional sublimation of the decomposed precursor

( 2 ) ZnO Particle Preparation and their Characteristics Conventional SP with an isothermal furnace wall temperature of 500°C (SP500) generated particles of submicrometer size. The particles consisted of polycrystals, which aggregated and could not be separated from the original drying particle. On the other hand, PCSP without and with a heated furnace of 300", 400", and 600°C (PC, PCSP300, PCSP400, and PCSP600) generated nanoparticles. Figure 5 shows FE-SEM images of SP500 and PCSP600 to compare the effect of PC on the generated particles. PC plays a role in producing ZnO nanoparticles. XRD patterns are also depicted in Fig. 5. The patterns show that ZnO particles had a zincite structure, which agrees well with the reference (JCPDS 36-1451), and that the particles synthesized by PC have a larger crystal size than those synthesized from SP. The crystalline size was determined using the well-established Scherrer's method by measuring the full-width at half-maximum of the diffraction peak. Based on [IOI] crystal orientation at 20 = 36.251' as the highest peak of the XRD pattern, the calculated crystal sizes were 14.0, 24.8, 25.2, 26.4, and 32.8 nm for SP500, PC, PCSP300, PCSP400, and PCSP600, respectively.

Fig. 3. Temperature profile in the spray pyrolysis system with an isothermal furnace wall temperature of 500°C and in the pulse combustion region as predicted by computational fluid dynamics.

Synthesis Methods for Powders

55

0.2

.......

23.9% weighl loss .....................................................

h&r&

0.0 -0.2 -0.4

g d

2 35

100

200

300

400

500

'"I

I 1

0.2

0.0

zG

-0.2 p.

-s 3

*s 5

L t 10

-0.8 600

Ih\

-0.4

-0.6 2%

30

f

1

980c

II

v

5

-0.6k

Temperature ["C]

Fig. 4. Thermogravimetric-differential thermogravimetric analysis of precursors: (a) zinc acetate dihydrate (b) gadolinium nitrate hexahydrate.

The crystallinity of the particles generated by PC integrated by a heated furnace (PCSP) increased with increasing furnace temperature. The energy of PC alone was sufficient to form ZnO nanoparticles, and additional heat from heated furnace (PCSP) was used for crystal growth. Figure 6 shows TEM images of particles produced by PC, PCSP, and SP. Using PC, disperse nanoparticles were obtained. The primary particle size generated by PC increased with increasing furnace wall temperature as shown in Fig. 6(a-c). Agglomerated, polycrystalline particles were produced by SP alone (SPSOO) as shown in Fig. 6(d). The insets of Fig. 6 show the corresponding selected-area electron diffraction (SAED) patterns of particles. Spotty ring patterns without any additional diffraction spots and rings of second phases reveal that particles were of high crystalline zincite ZnO. Clearer spotted patterns indicated that an increase in crystallinity was also found with XRD measurements. Unlike PC and PCSP particles, the ZnO particles prepared from SP (SPSOO) were polycrystalline, as indicated by circular line diffraction patterns (Fig. 6(d) inset). A high-resolution TEM (HRTEM) image of a particle generated by PC is shown in Fig. 7(a). The particle had single direction lattice fringes with a lattice spacing of 0.28 nm, which is the [ 1001 orientation. Conversely, conventional SP particles had multi-directional lattice fringes and some defect regions. Amorphous regions were also found as shown in Fig. 7(b).

Fig. 5. X-ray diffractrometer patterns of zinc oxide particles synthesized by spray pyrolysis (SP) at 500°C (SPSOO), pulse combustion (PC), and PC integrated with SP at 300", 400", and 600°C (PCSP300, PCSP400, and PCSP6OO). Field-emission scanning electron microscopic images show submicrometer particle and nanoparticles generated by conventional SP (SPSOO) and PC integrated with SP at 600°C (PCSP600), respectively.

56

(3) Gd20J:Eu3+Particle Preparation and their Characteristics The FE-SEM images in Fig. 8(a) show that spherical Gd203:Eu3+ particles were typically generated by SP.I8 In contrast with ZnO particles, PC produced ellipsoidal-toroidal Gd203:Eu3+ particles as shown by the FE-SEM images in Fig. 8(b). The particles were also ellipsoidal-toroidal in morphology as shown in Fig. 8(c) when PC integrated with SP at a temperature of 800°C was used. Figure S(d-0 show high-magnification FE-SEM images of particles generated by SP (SPSOO), PC, and PC integrated with a heated furnace (PCSPSOO), respectively. PC affects the morphology of the Gdz03:Eu3+particles generated, giving rise to the generation of nonspherical particles. The change in morphology resulted in an increased particle surface area. The specific surface area measured by the BET method showed that particles synthesized by SP800, PC, and PCSP800 had surface areas of 12.606, 15.660, and 17.061 m2/g, respectively. XRD analysis of Gd2O3:Eu3+particles generated with PC, SP, and PC integrated with SP are shown in Fig. 9(a). The XRD patterns of particles generated by PC and SP at a temperature of 800°C agree with the Gd203crystal reference (JCPDS 12-0797).

Fig.6. Transmission electronic microscopic images of particles generated by (a) pulse combustion (PC), (b) PC integrated with spray pyrolysis (SP) at 400°C (PCSP400), (c) 600°C (PCSP600),and (d) conventional SP. The insets correspond to selected analysis electron diffraction.

Progress in Nanotechnology: Processing

crystal size by PCSP. Overall, the results show that the PL intensity correlates with both the crystal size and the distribution of the activator in the host lattice. From our result, PC contributes to an increase in both the activation distribution and the crystal size.

Fig. 7. High-resolution transmission electronic microscopic images of zinc oxide particles generated by (a) pulse combustion and (b) spray pyrolysis at 500°C (SPSOO).

Amorphous phase particles were observed on particles generated by SP at temperatures of 300" and 500"C, as was expected from TG-DTG analysis of the precursor. Surprisingly, the XRD pattern of particles synthesized by PC without using a heated furnace (PC) is in good agreement with the Gd203 reference even though they had low crystallinity. The crystallinity of the particles synthesized by PC combined with SP increased with increasing furnace temperature. PL intensity emission at excitation of Gd203:Eu3+ at 254 nm as shown in Fig. 9(b) correlates with crystallinity and crystal size." Particles synthesized by PC integrated with SP at 800°C (PCSP800) had the largest peak PL intensity, followed by particles generated by SP only with a furnace wall at 800°C (SPSOO). PC and PC integrated with a heated furnace at 300" and 500°C (PC, PCSP300, and PCSP500) generated particles with almost similar PL intensities. They showed a significant luminescence y a k of Eu3+ at approximately 612 nm, indicating a red D ~ - + ' Ftransition ~ at c2sites within europium." Particles generated by conventional SP at 300" and 500°C had the lowest PL intensity. Presumably, PC has a significant role in increasing the PL intensity because PC increases the distribution of Eu3+ activators in the Gd203 crystals and chemical homogeneity increases the efficiency of the activator Eu3+ cations in the host lattice.20 However, more energy is needed to obtain a certain

( 4 ) Proposed Particle Formation Mechanism Experimental results show that a material with a decomposition temperature lower than the pulse average temperature, such as ZnO, can produce nanoparticles. Presumably, the primary particles of nucleated ZnO were formed just before droplets entering the furnace. The initial submicrometer size of ZnO particles tends to undergo fragmentation due to velocity and pressure fluctuations generated by PC. In the case of a higher decomposition reaction such as Gd203:Eu3+,ellipsoidal-toroidal particles in submicrometer size were generated. PC affects on the surface droplet-like particles that deformed easily to an ellipsoidal-toroidal morphology. In this case, fluctuating velocity and pressure are not strong enough to break the molecular chain in a surface drying precursor. There are two types of interactions between two different phases, e.g. fluctuating flow generated by PC (gas phase) and fast precipitated droplets (solid/liquid). These gas-liquid interface interactions can lead to the fragmentation of agglomerate crystals in a precipitated droplet in the case of ZnO synthesis and deformable particles acting as bubbles in the case of Gd203:Eu3+synthesis depending on their characterisTwo characteristic dimensionless parameters are bond number and Weber number. The bond number represents the ratio of the apparent gravity of an impacting particle to p ,dd, and o the surface tension force (Bo = ( A p u d ; ) / ~ ) . ~ ~ Aa, are the difference in the densities of the droplet and the surrounding fluid, the acceleration due to a change in the gas velocity, the droplet size, and the surface tension, respectively. The Weber number defines the square root of the ratio of the kinetic ener y of the impacting particle to the surface ) ) and . ~ ~ U are the gas density tension (We= (p8U ~ i ~ ) / ( 2 0pg and the maximum speed in the acoustic disturbance, respectively. High fluctuating oscillation generated by PC yield the high bond and Weber numbers. They lead to droplet-like particles' deformation. However, the ability to disperse primary particles in an agglomerate particle formed by evaporation, precipitation, and nucleation of a precursor in a droplet is

3

Fig. 8. Field-emission scanning electron microscopic images of gadolinium oxide-doped with europium particles synthesized by (a) spray pyrolysis (SP) at 800°C (SPSOO), (b) pulse combustion (PC), and (c) PC integrated with SP at 800°C (PCSP800). (d-f) show high-magnification images.

Synthesis Methods for Powders

57

Fig. 9. (a) X-ray diffractrometer patterns and (b) photoluminescence spectra of gadolinium oxidedoped with europium particles generated by pulse combustion and conventional spray pyrolysis.

the average gas temperature generated by PC were used to generate nano ZnO and submicrometer-aspherical Gd203. This study showed that the precursor decomposition temperature has a significant role in determining the final particle morphology. In addition, PC contributes toward increasing the characteristics of the materials prepared such as crystallinity, luminescence, and surface area. Acknowledgments The authors wish to thank M. Hazata and T. Ozaki for their assistance with the experiments, and T. Ogi for TEM analysis. Mr. K. Kato, M. Sabi, and N. Ike from Dai-Ichi Kogyo Seiyaku Co. Ltd.. Kyoto, Japan, as well as T. Kubotani and M. Wada from Pultech Corporation, Kobe. Japan, are thanked for valuable discussions. The Ministry of Education, Culture, Sports, Science and Technology (MEXT) of Japan and the Japan Society for the Promotion of Science (JSPS) are acknowledged for providing doctoral scholarships (W. W. and A. P.) and a postdoctoral fellowship (W. N. W.), respectively.

References Fig.10. Proposed particle formation mechanism in the case of zinc oxide and gadolinium oxide-doped with europium.

influenced by its agglomeration behavior such as primary particle size and number, density, and interparticle bond strength.22 The fragmentation of dried submicrometer nanoparticle aggregates into dispersed nanoparticles under low pressure conditions was also investigated by our Pressure differences between liquid and solid nanoparticles inside droplets induced by rapid drying and evaporation rates are considered to be the main reasons for particle fragmentation or droplet distortion. A simple proposed mechanism of ZnO and Gd2O3:Eu3+ particle synthesis by PC is illustrated in Fig. 10. The PC is actually a complex process and will become a specific topic to be investigated.

IV.

Conclusions

A new method for synthesizing nanoparticles and submicrometer-aspherical particles was developed using an integrated PCSP process. Zinc acetate and gadolinium nitrate having lower and higher decomposition temperatures, respectively, than 58

‘K. Okuyama and I. W. Lenggoro, “Preparation of Nanoparticles Via Spray Route,” Chem. Eng. Sci.,58 [ 3 4 ] 53747 (2003). 2W. N. Wang, Y.Itoh, 1. W.Lenggoro, and K. Okuyama, “Nickel and Nickel Oxide Nanoparticles Prepared from Nickel Nitrate Hexahydrate by a Low Pressure Spray Pyrolysis,” Mar. Sci. Eng. B-Solid,111 [ I ] 6%76 (2004). ’A. Purwanto, I. W. Lenggoro, H. W. Chang, and K. Okuyama, “Preparation of Submicron- and Nanometer-Sized Particles of Y203:Eu3+ by Flame Spray Pyrolysis Using Ultrasonic and Two-Fluid Atomizers,’’ J. Chem. Eng. Jpn., 39 [I] 68-76 (2006). 4C. Panatarani, 1. W. Lenggoro, and K. Okuyama, “Synthesis of Single Crystalline ZnO Nanoparticles by Salt-Assisted Spray Pyrolysis,” J. Nunopart. Res., 5 [1-2] 47-53 (2003). ’W. N. Wang, S. G. Kim, I. W. Lenggoro, and K. Okuyama. “Polymer-Assisted Annealing of Spray-Pyrolyzed Powders for Formation of Luminescent Particles with Submicrometer and Nanometer Sizes,” J. Am. Ceram. Soc., 90 [2) 425-32 (2007). %. L. Fraenkel, L. A. H. Nogueira, J. A. Carvalho, and F. S. Costa. “Heat Transfer Coefficients for Drying in Pulsating Flows,” I n f . Commun. Hear Muss, 25 [4] 471-80 (1998). ’I. Zhicinski. A. Delag, C. Strumillo, and J. Adamiec, “Advanced Experimental Analysis of Drying Kinetics in Spray Drying,” Chem. Eng. J., 86 [1-2] 207-16 (2002). ‘1. Zbicinski, C. Strumillo, M. Kwapinska, and 1. Smucerowicz, “Calculations of the Pulse Combustion Drying System,” Energ. Convers. Munage.. 42 [IS171 1909-18 (2001). 9.S. Kuts, P. V. Akulich, N. N. Grinchik. C. Strumillo, 1. Zbicinski, and E. F. Nogotov, “Modeling ofGas Dynamics in a Pulse Combustion Chamber to Predict Initial Drying Process Parameters,” Chem. Eng. J.,86 [I-21 25-31 (2002). ‘“W. Liang, H. Sunada, and F. D. Cui, “Improvement of the Dissolution Rate of the Insoluble Drug-A Pulse Combustion Dryer System.” Pharm. Tech. Jpn., 22 [7] 47-51 (2006).

Progress in Nanotechnology: Processing

‘IF. J. Weinberg. Advunced Combustion Methods, pp. 115-81. Academic Press Inc, London, 1986. ”A. V. Ghule, K. Ghule, C. Y. Chen, W. Y. Chen, S. H. Tzing, H. Chang, and Y. C. Ling, “In Situ Thermo-TOF-SIMS Study of Thermal Decomposition of Zinc Acetate Dehydrate,” J . Muss Spectrum., 39 [lo]1202-8 (2004). I3U. Ozgur, Y. I. Alivov, C. Liu, A. Teke, M. A. Reshchikov. S. Dogan, V. Avrutin, S. J. Cho. and H. Morkoc, “A Comprehensive Review of ZnO Materials and Devices.” J . Appl. Phys.. 98 [4]1-103(2005). ‘9.Tani, A. Kato, and H. Morisaka, “Effects of Solvent on Powder Characteristics of Zinc Oxide and Magnesia Prepared by Flame Spray Pyrolysis,” J . C e r m . Soc. Jpn., 113 [3]2 5 M (2005). ”Y. C. Kang and S. B. Park, “Preparation of Nanometre Size Oxide Particles Using Filter Expansion Aerosol Generator,” J. Murer. Sci., 31 [9]2409-16 (1996). I6Y. C. Kang, S. B. Park, 1. W. Lenggoro, and K. Okuyama, “Gd203:Eu Phosphor Particles with Sphericity, Submicron Size and Non-Aggregation Characteristics,” J . Phys. Chem. Solids, 60 [3]37944 (1999). ”R. K. Dubey, D. L. Black, M. Q. McQuay, and J. A. Carvalho, “The Effect of Acoustics on an Ethanol Spray Flame in a Propane-Fired Pulse Combustor,” Combust. Flume, 110 [I-21 25-38 (1997). “W. N. Wang, W. Widiyastuti, T. Ogi, 1. W. Lenggoro, and K. Okuyama, “Correlation between Crystallite/Particle Sizes and Photo-Luminescent Properties of Submicron Phosphors,” Chem. Muter., 19 [7l 1723-30 (2007).

Synthesis Methods for Powders

I9W. N. Wang, W. Widiyastuti, 1. W. Lenggoro, T. 0. Kim, and K. Okuyama, “Photoluminescence Optimization of Luminescent Nanocomposites Fabricated by Spray Pyrolysis of a Colloid-Solution Precursor,” J . Elecfrochem. Soc., 154 [4] JI2I-8 (2007). *‘C. Louis, R. Bazzi, M. A. Flores, W. Zheng, K. Lehbou, 0. Tillement, B. Mercier, C. Dujardin, and P. Perriat, “Synthesis and Characterization of Gd201:EuZ+Phosphor Nanoparticles by a Sol-Lyophilization Technique,” J . Solid State Chem., 173 [2]33541 (2003). ”A. V. Nguyen, H. J. Schulze, H. Stechemesser, and G. Z o k l , “Contact Time During Impact of a Spherical Particle against a Plane Gas-Liquid Interface: Theo ” Int. J . Miner. Process,50 [I-21 97-1 I I (1997). “R. Moreno-Atanasio and M. Ghadiri, “Mechanistic Analysis and Computer Simulation of Impact Breakage of Agglomerates: Effect of Surface Energy,” Chem. Eng. Sci., 61 [XI247681 (2006). 23F. Iskandar, L. Gradon, and K. Okuyama, “Control of the Morphology of Nanostructured Particles Prepared by the Spray Drying of a Nanoparticle Sol,” J. Colloid fnterf. Sci., 265 [2]296303 (2003). “I. F. Murray and S. D. Heister, “On a Droplet’s Response to Acoustic Excitation,” Inr. J . Multiphas. Flow, 25 [3]531-50 (1999). 25W.N. Wang, I. W. Lenggoro, and K. Okuyama, “Dispersion and Aggregation of Nanoparticles Derived from Colloidal Droplets Under Low-Pressure Conditions,” J. Colloid InterJ Sci., 288 [2]423-31 (2005). 0

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One-Step Synthesis of Luminescent Nanoparticles of Complex Oxide, Strontium Aluminate Chengzhou Li, Yusuke Imai, Yoshio Adachi, Hiroshi Yamada, Keiko Nishikubo, and Chao-Nan Xut On-Site Sensing and Diagnosis Research Laboratory, National Institute of Advanced Industrial Science and Technology, Saga 841-0052, Japan

An organic aqueous solution of metal acetylacetonate precursors was subjected to spray pyrolysis in order to fabricate SrA1204:Eu (SAO) nanoparticles. Non-agglomerated luminescent SAO nanoparticles, having a spherical shape with a size of 10-30 nm, were achieved in a single step, while only submicrometer-sized SAO particles were obtained from the conventional ultrasonic pyrolysis of the metal nitrates. Without any post-annealing process, the as-prepared SAO nanoparticles were observed to exhibit a strong photoluminescence, which is comparable with that of the submicrometer-sized SAO particles. A mechanism for the formation of the nanoparticles is also discussed. I.

Introduction

R

ARE-earth-doped alkaline earth ahminates are very wellknown phosphors because of their long persisting luminescence in the visible region. The europium-doped strontium aluminate has also been proposed for use as a stress indicator or crack indicator because it can emit visible light under mechanical stress.”2 Recently, we have realized the visualization of the stress distribution of a subject by utilizing an intense visible elastico-luminescenceemitted by the europium-doped strontium aluminate, E~.01Sr0,99A1204 (SAO)?’ Non-agglomerated SAO nanoparticles are highly desired from the viewpoints of displaying the stress distribution of a tiny or microscopic object as well as fabricating other mechano-optical nano-devices. Although the elastico-luminescence is easily observed, the underlying mechanism remains obscure.’ The SAO nanoparticle will also be considerably helpful in understanding the luminescence behavior of a single nanoparticle as it may be excited by a mechanic stress imposed using an atomic force microscope.6 Despite the extensive studies on fabricating nanocrystals of single-cation oxides or sulfide^^-'^ as well as nanostructures of non-luminescent complex oxide^,'^-'^ a limited amount of work has contributed to the manufacture of luminescent nanoparticles of complex oxides such as SA0.I’ Nevertheless, good-quality non-agglomerated SAO nanoparticles have never been achieved. For SAO, a high temperature and a reducing atmosphere are necessary for forming luminescent centers during crystallization. However, high-temperature calcination often results in a rapid crystal growth, forming necks, and consequently, aggregates or hard agglomerates. To prepare non-agglomerated particles, ultrasonic spray pyrolysis is a well-established process. Owing to the ineffective generation of tiny droplets, however, the process can only produce submicrometer (0.1-1 pm) and micrometer-

sized particles accordin to the one-particle-per-droplet mechanism.” Xia et d.” demonstrated a unique ultrasonic spray pyrolysis process to obtain BaTi03 nanoparticles after post-washing the as-prepared particles in water that removes a large amount of salt, which was added to prevent any agglomeration. Unfortunately, SAO is vulnerable to water due to the hydrolysis of aluminate, and even a small amount of an inorganic impurity will reduce its performance. In this context, we aim to develop a novel process for preparing non-agglomerated SAO nanoparticles. Such a process will also offer an alternative route to producing other complex oxide nanoparticles. Instead of the aqueous inorganic precursor solution that is often used in a conventional ultrasonic spray pyrolysis, in the present study, metal organic compounds and an organic solvent were used to form the starting solution. It was then atomized by an ultrasonic atomizer and the generated mist was carried into the furnace tube to prepare the SAO nanoparticles.

11.

Experimental Procedure

Strontium acetylacetonate, aluminum acetylacetonate, and europium acetylacetonate were purchased from Wako Pure Chemicals Industries Ltd. (Osaka, Japan) and used without further purification. All other chemicals used in this study were of analytical grade. The ultrasonic spray pyrolysis system used in the present study consisted of an ultrasonic atomizer with a 2.4 MHz transducer, a horizontal reaction furnace, and an electrostatic precipitator for collecting the particles. Strontium acetylacetonate, aluminum acetylacetonate, and europium acetylacetonate were dissolved in deionized water containing the desired amount of acetylacetone to prepare a 0.0025M starting solution in terms of the SAO. The starting solution was atomized and the mist containing the precursors, the organic solvent, and water was carried into the tube furnace by a 5% H2/Ar carrier gas. The carrier gas speed was maintained at 1 L/min. A conventional ultrasonic spray pyrolysis was also carried out using an aqueous solution of metal nitrates as the starting solution, and acetylacetone was supplemented if necessary. The morphology and particle size of the sample were observed using field emission-scanning electron microscopy (FE-SEM, Hitachi 4500, Tokyo, Japan). The crystal structures were determined by X-ray diffraction (XRD, Rigaku 2000, Tokyo, Japan) with CuKx radiation. The photoluminescence measurements were performed using a fluorospectrometer (Jasco FP6600, Tokyo, Japan). 111. Results and Discussion

J. Ballat-ntrihutlng

editoi

Manuscript No. 22550. Received December 6,2006; approved March 5,2007. Partly supported by the Industrial Technology Research Grant Program in 2003 from the New Energy and Industrial Technology Development Organization (NEDO) of Japan. ‘Author to whom correspondence should be addressed. e-mail: [email protected]

Synthesis Methods for Powders

Figure I(a) shows the FE-SEM image of the nanoparticles prepared by the present process in which an aqueous-organic solution of the metal organic precursors was used as the starting solution. For comparison, Fig. l(b) shows the morphology of the submicrometer-sized particles prepared by a conventional ultrasonic spray pyrolysis of the nitrate salt solution. The con-

61

Fig. 1. Field emission-scanning electron microscopic observation of the strontium aluminate particles prepared by various processes at 1400°C. (A) Aqueous acetylacetonate solution with added acetylacetone; (B) aqueous nitrate solution; and (C) aqueous nitrate solution with added acetylacetone.

centrations of the two starting solutions were made equal in terms of the final product. It can be noted in the figure that both the former and the latter are non-agglomerated spherical particles. The former consisted of nanoparticles with a size range of 1&30 nm, whereas the latter was composed of 100-200 nm spherical particles. To understand the role of acetylacetone in the process, a small amount of acetylacetone was added to the nitrate solution to carry out the ultrasonic spray pyrolysis. Figure I(c) shows a typical FE-SEM image of the resulting particles, and it can be seen in the figure to be mainly composed of submicrometer-sized particles. Thus, the formation of the nanoparticles should not be simply attributed to the effect of the added organic solvent. It could be a result of the synergistic effect of the organic precursors and the solvent. To investigate the effect of the processing temperature on the particle morphology, we prepared the particles at various temperatures from 1300" to 1500°C. Figure 2 depicts the different morphologies of the as-prepared particles obtained at the various temperatures. The particles obtained at 1300°C show a bimodal size distribution, which is composed of several tens of nanometers and several hundreds of nanometers particles. In contrast, the particles obtained at 1400" and 1500°C were found to be non-agglomerated nanoparticles as shown in Figs. 2(b) and (c). To our surprise, the particles obtained at 1500"C, the highest temperature, were found to be the best quality non-agglomerated nanoparticles in the present study. This result implies that the highest temperature gradient is favorable for producing monodispersed nanoparticles. It is considered that rapidly heating the mist droplets of the aqueous organic solution may induce fragmentation of the droplets into many smaller subdroplets, thus resulting in the nanoparticles. Therefore, the present novel process did not follow the one-particle-per-droplet mechanism. As it is known that acetylacetone is only partially miscible in water, a tentative mechanism for the formation of the nanoparticles was proposed as follows. When the mist droplets passed through the furnace tube, numerous acetylacetone subunits were first separated from the

bulk aqueous phase while the water escaped as a vapor. For the present system, the acetylacetone would spontaneously extract the precursor from the aqueous phase. Therefore, it is favorable for the precursor to diffuse into the organic phase, and almost no precursor precipitated into the aqueous phase. Each droplet was then broken into many subdroplets composed of the subunits as the water vigorously evaporated due to the high temperature gradient in the furnace tube. These subdroplets were decomposed and finally converted into nanoparticles during their passage through the furnace. The crystal structures of the SAO particles obtained at the different temperatures were determined by the XRD technique, and the results are shown in Fig. 3. The crystal structure of the nanoparticles prepared under the various temperatures was determined as a monoclinic phase (PDF #334379) according to the ICDD database. It was pointed out in our previous study that only the monoclinic phase SAO could exhibit a practical elastico-luminescenceperformance that was much stronger than the hexagonal phase SAO as well as the other strontium alumhates! Thus, the SAO nanoparticles are required to have a monoclinic phase for exploring their applications with respect to the elastico-luminescence. The XRD peak intensity shown in Fig. 3 steadily increased with increasing calcination temperature from 1300" to 1500°C. The higher peak intensity indicates a more well-developed crystalline structure formed at the higher temperature, which may lead to a better performance. Without any post-annealing process, the as-prepared SAO nanoparticles prepared at the different temperatures were observed to be able to emit a green light that is visible to the naked eye in a dim room. Their excitation and emission spectra are shown in Fig. 4. Each excitation spectrum is seen to be a broad band from 270 to 400 nm and the peak shifted to the lower wavelength side compared with a standard sample prepared by a solid-state reaction process (data not shown). The emission spectra of these nanoparticles show a symmetrical broad band that peaked around 515 nm, which could be attributed to the 4f65d'4f7transition of the divalent europium. The peak slightly

Fig. 2. Different morphologies of the strontium aluminate particles prepared by the present process at (A) 1300"C, (B) 1400"C, and (C) 1500°C.

62

Progress in Nanotechnology: Processing

nescent strontium aluminate nanoparticles without postannealing, although the elucidation of the formation mechanism of the nanoparticles in the present process requires further experimental and theoretical studies. The as-prepared nanoparticles, having a spherical shape with a size of 10-30 nm, were observed to exhibit strong photoluminescence that was comparable with the performance of the submicrometer-sizedparticles. This approach is considered to be highly applicable for preparing other complex oxide nanoparticles, such as rare-earth-activated alkaline earth aluminate nanophosphors.

1300 “C

Acknowledgment

200

C. L. wishes to thank the Japan Society for Promotion of Science (JSPS) for providing the JSPS fellowships.

n 10

20

30

40

50

60

70

2theta I degree

Fig. 3. X-ray diffraction patterns of the strontium aluminate nanoparticles prepared by the present process at various temperatures.

shifted to the blue side compared with the standard sample, althou h it was not as remarkable as claimed by some researchers.” For the SAO nanoparticles obtained at 130O0C, in addition to the main emission peak, a small shoulder peak was also observed at about 610 nm, which could have originated from the remaining trivalent europium. No red emission was observed for the other two samples under the present measurement conditions. As shown in Fig. 4, the photoluminescence of the SAO nanoparticles was enhanced with increasing processing temperature. This is in good agreement with the deduction from the XRD results. Among the three SAO samples, the one prepared at 1500°C showed the highest photoluminescence intensity, which is comparable with that of the submicrometer-sized particles prepared by the conventional spray pyrolysis depicted by the dotted line in Fig. 4.

IV. Summary The processing of nanoparticles with the desired crystal phase and uniform morphology is always important for realizing highperformance devices on a nanoscale. The present study demonstrated an effective method to prepare non-agglomerated lumi-

Fig.4. Excitation and emission spectra of the strontium aluminate luminescent nanoparticles achieved at various temperatures. The dotted line shows the emission spectrum of the submicrometer-sized particles prepared by the conventional spray pyrolysis process at 1400°C.

Synthesis Methods for Powders

References ‘C. N. Xu, T. Watanabe, M. Akiyama, and X. G. Zheng, “Direct View of Stress Distribution in Solid by Mechanoluminescence,” Appl. Phys. Lett., 74 [I71 2414-6 (1999). ’K. S. Sohn, S . Y. Seo, Y. N. Kwon, and H. D. Park, “Direct Observation of Crack Tip Stress Field Using the Mechanoluminescence of SrAI2O4:(Eu.Dy,Nd),” J. Am. Ceram. Soc., 85 [3] 7 1 2 4 (2002). ’C. N. Xu, X. G. Zheng, M. Akiyama, K. Nonaka, and T. Watanabe, “Dynamic Visualization of Stress Distribution by Mechanoluminescence Image,” A p t / . Phys. Lett., 76 [2] 179-81 (2000). C. N. Xu. H. Yamada, X. Wang, and X. G. Zheng, “Strong Elasticoluminescence from Monoclinic-Structure SrAI2O4,” Appl. Phys. Lett., 84 [I61 3040-2

(2004).

’C. N. Xu, “Smart Coating With Mechanoluminescence”; pp. 190-201 in Encyclopedia o/Smart Materials. Vol. I,Edited by M. Schwartr. John Wiley & Sons lnc., New York, 2002. 6K. Sakai, T. Koga, Y. Imai, S . Maehara, and C. N. Xu, “Observation of Mechanically Induced Luminescence from Microparticles,” Phys. Chem. Chem.

P h y , 8 [24] 2819-22 (2006). F. X. Redl, K.4. Cho, C. B. Murray, and S. OBrien, “Three-Dimensional Binary Superlattices of Magnetic Nanocrystals and Semiconductor Quantum Dots,” Nature, 423 [6943] 968-71 (2003). ‘T. Masui, H. Hirai, R. Hamada, N. Imanaka, G. Adachi, T. Sakata, and H. Mori, “Synthesis and Characterization of Cerium Oxide Nanoparticles Coated With Turbostratic Boron Nitride,” J . Mater. Chem., 13 [3] 622-7 (2003). 91. W. Lenggoro, Y. Ito, K. Okuyama, and T. 0. Kim, “Nanoparticles of a Doped Oxide Phosphor Prepared by Direct-Spray Pyrolysis,” J. Mater. Res., 19 [l:L3534-9 (2004). R. M. Laine, J. C. Marchal. H. P. Sun, and X. Q . Pan. “Nano-Alpha-AI20~by Liquid-Feed Flame Spray Pyrolysis,” Nar. Mater., 5 [9] 710-2 (2006). “T. Torimoto, S. Y. Murakami, M. Sakuraoka, K. Iwasaki, K. Okazaki, T. Shibayama, and B. Ohtani, “Photochemical Fine-Tuning of Luminescent Color of Cadmium Selenide Nanoparticles: Fabricating a Single-Source Multicolor Luminophore,” J. Phys. Chem. B., 110 [27] 13314-8 (2006). I2Y. W. Zhao, Y. Zhang, H. Zhu, G. C. Hadjipianayis. and J. Q. Xiao, “Low-Temperature Synthesis of Hexagonal (Wurtzite) ZnS Nanocrystals,” J. Am. Chem. Sac., 126 [22] 6874-5 (2004). ”0. Vasylkiv, Y. Sakka, and V. V. Skorokhod, “Nano-Blast Synthesis of Nano-Size CeO+3d203 Powders,” J . Am. Ceram. Soc., 89 [6] 1 8 2 2 4 (2006). I4A. J. Zarur and J. Y. Ying, “Reverse Microemulsion Synthesis of Nanostructured Complex Oxides for Catalytic Combustion,” Nature, 403 [6765] 65-7 (2000). ”J. C. Yu, L. Zhang, Q. Li. K. W. Kwong, A. W. Xu, and J. Lin, “Sonochemical Preparation of Nanoporous Composites of Titanium Oxide and Size-Tunable Strontium Titanate Crystals,” Langmuir, 19 [I81 7673-5 (2003). “B. Xia, I. W. Lenggoro, and K. Okuyama, “Novel Route to Nanoparticle Synthesis by Salt-Assisted Aerosol Decomposition,” Adv. M a w . , 13 [20] 1579-82 (29: 1). H. Liu, C. Hu, and Z. L. Wang, “Composite-Hydroxide-Mediated Approach for the Synthesis of Nanostmctures of Complex Functional-Oxides,” Nano L e y , M) [7] 1 5 3 M (2006). 9.Y. Peng, H. Liu, H. P. Yang, and C. H. Yan, “Synthesis of SrAI2O4:Eu.Dy Phosphor Nanometer Powders by Sol-Gel Processes and its Optical Properties,” Mater. Chem. Phys., 85 [I] 68-72 (2004). I9G. L. Messing, S. C. Zhang, and G. V. Jayanthi, “Ceramic Powder Synthesis 0 by Spray Pyrolysis,” J . Am. Ceram. Soc., 76 [ I11 2707-26 (193).

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Nan0 a-A1203Powder Preparation by Calcining an Emulsion Precursor Yu-Chen Lee, Shaw-Bing Wen,+ and Liang Wenglin Department of Resources Engineering, National Cheng Kung University, Tainan 701, Taiwan

Chih-Peng Lin Department of Information Management, Diwan College of Management, Madou, Tainan 721, Taiwan

An experimental study has been conducted to evaluate the formation of nano a-AI2O3 under various conditions, such as different calcining temperatures and emulsion ratios of aqueous aluminum nitrate solutions and oleic acid with a high-speed stirring mixer. Four batches of the precursor powders were calcined at three different temperatures of lOOO",1050",and 1100°C for 2 h and a terminal product of nano a-AI2O3 powders was obtained. The products have been identified by X-ray diffraction (XRD), specific surface area measurement scanning electron microscope, and transmission electron microscope (TEM). The XRD results show that the phase of powders is determined to be a-Alz03, indicating that the overall process has been effective. The optimum calcination temperature of the precursor powder for crystallization of nano a-AI203 was found to be 1000°C for 2 h. The TEM image indicates that the particle grains have a subspherical shape with a mean size of 50-100 nm.

A

I. Introduction

LUMINA is generally considered to show desirable performance for modern industry, such as mechanical strength for wear resistance, thermal conductivity, and electrical insulation for electronic manufacturing, and is also noted to be favorably cost-effective.'32 Alumina powders have played an important role as abrasives in the planarization of silicon wafers in the semiconductor industry3; however, such applications demand a particular nano-sized powder. It seems that nanosized alumina powders will be indispensable for the processing needs of various electronic materials and devices in the future. The characteristics of nano-alumina particles have been noted to be highly variant, primarily due to the aging process as well as the phase transformation effect of thermal treatment of the feedstock powders so that the morphologies and particle size distribution are distinctively n~ticeable.~ The crystal growth and particle agglomeration are very difficult to avoid during the preparation of alumina powders. Sarikaya and colleag~es'~~ prepared alumina microshells by calcining the precursor material derived from water-in-oil-type emulsions using an aqueous AI(N03)3solution as the water phase, mineral oil as the organic phase, and a non-ionic surfactant, Arlacel 83, as the emulsifier. This particular method has been named the emulsion evaporation technique, and the micrometer products prepared by Sarikaya et al. had a comparatively larger surface area, being suitable for different applications. On the other hand, Lin ef used the chemical precipitation method to prepare an alumina precursor via an emulsified boehmite gel with oleic acid, subsequently calcined to 1300°C for 20 min under a controlled atmosphere, the final product being a-A1203nano-sized powder.

J. Blendell-contributing editor

Manuscript No. 22398. Received October 26, 2006, approved January 17, 2007. This paper was supported by National Cheng-Kung University. 'Author to whom correspondence should be addressed. e-mail [email protected]

Synthesis Methods for Powders

It is noteworthy that the above procedure provides a unique and desirable situation for limiting the growth of the crystal sizes from unnecessary enlargement, seemingly caused by the incomplete combustion of the organics and the presence of char particles. This method could obtain very fine nano-sized a-alumina with a favorable range of particle size near 30 nm. The objectives of this study are twofold, namely to evaluate: (1) the alumina precursor phase transformation by various ratios of an aqueous aluminum nitrate solution and oleic acid and (2) the alterations of crystal structure and size by the calcination temperature. The products were evaluated by X-ray diffraction (XRD) analysis and transmission electron microscopy (TEM). Figure 1 illustrates both the concept of this research and also the sequence of the formation of various products. Our laborious experimentation seems to provide some useful information of the subject matter; however, the scope of this study is limited, and it seems further experimentation and evaluation are necessary for a full understanding.

11. Experimental Procedure The procedure for this study consisted of three steps: the preparation of an emulsion, the formation of an alumina precursor, and the calcination process. Figure 1 shows a flow chart of the experimental procedure. The emulsions were obtained by mixing oleic acid (Ci7H&OOH, Katayama Chemical Co. Ltd., Japan) and 1M aqueous aluminum nitrate (Al(N03)3.9H20, Merck, Germany) and stirring the solution by using a propeller mixer operated at 1200 rpm. The volume ratio (ol/an) of oleic acid and aqueous aluminum nitrate was in the range of 8:2, 7:3, 6:4, and 55. The mixed emulsion was stirred and heated with a plate heater at 180°C to remove water, and at the end of the mixingheating operation, the emulsion micelles were found to have yielded dark brownish alumina precursor particles and also showed high viscosity. These alumina precursors were then placed in a crucible, and calcination treatment was performed from room temperature to lOOO", 1050°,and 1100°C for 2 h. The final products of nano a-A1203powders were obtained. The thermal characteristics of alumina precursors were determined by thermogravimetry (TG) and differential thermal analysis (DTA, Netzsch STA 409 PC). The crystalline phase of calcined powder and average crystallite size were determined by X-ray powder diffractometry (XRD, Siemens D5000) with a Cu target. The specific surface area measurement (BET, Micromeritics Gemini 2360) of alumina powders was performed by adsorption on nitrogen (BET technique) on the powders at 77 K. A scanning electron microscope (SEM, Hitachi S4100) and a transmission electron microscope (TEM, Hitachi HF-2000) were used to observe the morphology and also the sizes of the crystallites. The average crystal size (0,) and full-width at half-maximum (FWHM) peak height in terms of radians are given as follows by the Scherrer formula':

D, = 0.9h/(FWHM) C O S ~

(1) 65

Fig. 1. Flow sheet and conceptual diagram of this study: (A) mixing of aqueous aluminum nitrate solutions with oleic acid in a high-speed mixer to form dispersed uniform emulsions. (B) The emulsions were heated to remove water until a brownish alumina precursorappeared. (C) Calcination, (D) loose particles of nano a-AI20, powders.

where h is the X-ray wavelength of the X rays (ku = 0.154 nm) and 0 is the diffraction angle (rd). 111. Results and Discussion (1) Thermal Decomposition of the Precursor The TG and DTA curves of the precursors that were obtained by heating from 25" to 1500°C in an air atmosphere are given in Fig. 2, an emulsion product having been derived from a mixture of oleic acid and aluminum nitrate with a (ol/al) ratio of 8:2. An endothermic peak was noted near 120°C on the DTA curve, representing the elimination of adsorptive water, and an exothermic peak between 120" and 700"C, seemingly due to the decomposition of oleic acid. The endothermic peak was considered to be due to the decomposition of boehmite's water of crystallization. It is relatively difficult to identify the exothermic peak of the precursor for its transformation into a-A1203.Therefore, we preheated the precursor product to burn out oleic acid at 900°C for 2 h, and Fig. 3 shows the DTA and TG curves, indicating explicitly an exothermic peak between 1070" and 1 190"C, being responsible for its transformation into a-A1203. According to Wefers and Misra," the &A1203 phase is present before the reconstructive phase transformation to a-A1203.The possible phase transformation of the alumina route is reckoned to have followed the sequence: gibbsite, bayerite, boehmite, and diaspore. The maximum formation temperature of the gibbsite

66

and bayerite is I S O T , and that for boehmite and diaspore are 300" and 400°C, Therefore, the precursor obtained at 180°Cfor this study could not have been either gibbsite or bayerite.

(2) Alumina Precursor Phase Transformation The XRD patterns of the alumina precursor separately calcined at 900", 950", and 1000°Cfor 2 h are shown in Fig. 4. It is noted that at 900°C, the predominant phase is 6-A1203,at 950"C, both 6- and a-A1203,and at 1000"C, a-AI2O3becomes the only phase present. The phase transformation observed in this study follows the sequence: 6-A1203-+ a-A1203.Thus, our results seem to contradict the findings of Wefers and Misra" in that an intermediate phase of e-AI2O3has not been identified for the phase transformation route of aluminum hydroxide to alumina. Generally speaking, the emergence of different phases may be controlled by the specific conditions of the operating process. Therefore, the difference in the phase evolution may be henceforth explained in the following. Our processing procedure, as shown in Fig. 1, has been unique in that due to furnace type and atmosphere condition, the burn-out settings may be different from that of Wefers and Misra." In our case, there was an incomplete burn-out of the oleic acid, yielding some residual particles of carbon,'.' which served to block and obstruct the emergence of the 0 phase of Progress in Nanotechnology: Processing

1 " " " " " " " I

t- xo 1

1070OC 114ooc

0-

20 -0

-2

.

0

,

200

.

,

400

.

I

.

,

.

,

.

600 800 1MX) Temperature ("C)

,

.

1200

,

?

1400

Fig. 2. Thermogravimetry (TG) and differential thermal analysis (DTA) curves of the alumina precursor at an (ol/an) emulsion ratio of (8:2) prepared by mixing aqueous aluminum nitrate solution and oleic acid. 0.43.,

. , .

I

.

I

.

,

. , . ,

,

102

-0.4

200

400

600 800 loo0 Temperature ("C)

1200

1400

Fig. 3. Thermogravimetry (TG) and differential thermal analysis (DTA) curves for the initial burn-out of oleic acid at 900°C for 2 h.

A1203.There is a further effect on the crystallite-size distribution and morphologies as given below. (3) Crystal Size and Structure The XRD patterns of the alumina powder prepared from the emulsion precursor at an (olian) ratio of 5:s are given in Fig. 5. The average size of U-AlzO, crystals at 1000°C may be calcu-

Fig.4.

X-ray diffraction patterns of the alumina precursor calcined at No",950", and IO00"Cfor 2 h.

Synthesis Methods for Powders

Fig. 5. X-ray diffraction patterns of alumina precursor for an emulsion (olian) ratio of 5 : s calcined at 1000", 1050", and 1100°C for 2 h.

lated as &re50 nm by using 0 = 25.646" and FWHM = 0.0033rd in the XRD patterns. The average size of the crystallites may be calculated by (1) the Scherrer formula9 and (2) BET-specific surface area, and the results are shown in Figs. 6 and 7, respectively. The average crystal size has been calculated to be in the range of 4 5 4 0 nm by the Scherrer formula, whereas the BET result has been shown to be variable as: (a) a range of nearly 100 nm at lOOO"C, (b) a range of 150-450 nm at 105O"C,and (c) 15&1500 nm at 1100°C. It is noted that the crystalline sizes of a-AI2O3powders are larger both in the increasing concentration of AI3+ ion in the feedstock emulsion and the stage of a higher calcination temperature. The average size of a-AI2O3calculated from BET and XRD measurements has been very different. Wen and Yen4 suggested that the average size of the BET-specific surface area increases very slowly at the nucleation stage and drastically at the growth stage with an increase in the amount of a-AI2O3. The difference between these two results indicates that in the agglomeration during the formation of a-A1203,especially in the high-temperature range, such as IOOOT, the BET average sizes of a-AI2O3powders by BET are near 100 nm, and XRD crystalline sizes by XRD are near 50 nm in spite of the variation in the concentration of A13+ ion. This particular calcination temperature seems to represent the overlap of a-AI2O3formation and the burn-out of oleic acid. As the calcination temperature was set at 950°C, the a-A1203phase had not been completely

Fig.6. Average size of the crystallite size calculated from the peak broadening of X-ray diffraction (XRD) results for products of u-AlzO,, prepared at different calcination temperatures and emulsion ratios.

67

formed, as shown in Fig. 4. These indications seem to point out that the agglomeration of a-A1203formation from the emulsion precursor may have been highly dependent on the calcination temperature. At a higher calcination temperature, the burn-out of oleic acid is faster and has less residual char. Thus, it may be concluded that the optimum calcination temperature of nano aAI,O3 powders prepared by calcining an (ol/an) emulsion is 1000°C.

Fig.7. The mean sizes of the BET-specific surface area vary with the calcination temperaturesofthe a-alumina powders prepared by different emulsion ratios.

( 4 ) Emulsion Ratios It is important to note that the emulsions obtained from different ratios of oleic acid/aluminum nitrate in the feedstock were found to be a highly significant factor in the preparation of nano-alumina powders. The particle size for an (olian) ratio of 8:2 is smaller than that of the other proportions, such as 5:5,6:4, and 7:3, given in Figs. 6 and 7. There are two possible reasons: firstly, the concentration of A13+ ion of the (8:2) precursor is lower during the course of phase transformation. As the calcined temperature increases, the material is considered to be deficient in terms of A13+ content so that the crystallite grain renders a retarded mode of growth. Secondly, the oleic acid content is relatively abundant and is thus capable of cutting off the grain contact among the nucleated a-A1203and restraining the

Fig. 8. Scanning electron microscope images of a-AI203 powders obtained from calcination at (a) lOOo"C, (b) IOSOT, and (c) 1IOo"C for 2 h at (ol/an) emulsion ratios 8:2 and 5 5 , respectively.

68

Progress in Nanotechnology:Processing

oleic acid. The above results are indicative that the overall process represents an effective methodology; however, further experiments are needed to provide more comprehensive results for a full understanding of the preparative processes.

IV. Conclusions The volumetric emulsion ratio (oleic acid/aluminum nitrate) of the starting feedstock has been shown to be a critical factor for the formation of nano a-AI2O3powders, notably the optimum ratio being 8:2. The optimum calcination temperature of the precursor powder for crystallization of nano a-AI2O3was found to be I O O O T for 2 h. The average size of crystallite particles is 5&100 nm, as given by SEM/TEM. Apparently, this size control may be related to calcination in a muffle furnace under an ambient air atmosphere.

Acknowledgments We appreciate Profasors F. S . Yen and H. S. Liu for their assistance with the preparation of the manuscript.

Fig. 9. Transmission electron microscope image of a-A1203powders at an emulsion (olian) ratio of 5 5 calcined at 1000°C for 2 h.

crystallite growth. The schematic diagram as given in Fig. 1 may be considered to provide a clear demonstration of the above inference on the different modes of crystallite growth. ( 5 ) Morphology From the SEM micrographs as given in Fig. 8, the sizes of the aAI2O3 powders are found to be near 100 nm at 1000°C and illustrate the formation of sintering necks among grains. As the temperature increases, the crystalline grain grows and more sintering necks are noted at the higher operating temperatures of 1050”and 1lOo”C, and also the BET data seem to show grain growth, possibly due to agglomeration (Fig. 7). A typical TEM image of a-A1203 powders ( 5 5 ) at 1000°C for 2 h is given in Fig. 9. The image represents a small portion of a-Al2O3,having an average size of near 50 nm. The size of nano a-A1203 products is sensitive to the calcination temperature and the amount of oleic acid surfactant in the feedstock for preparation. This research has shown the possibility of preparing nano u-AI2O3powders by calcining an emulsion consisting of an aqueous aluminum nitrate solution and

Synthesis Methods for Powders

References ‘W.H. Gitzen, Alumina as a Ceramic Murerial. American Ceramic Society,

Columbus, OH, 1970. ’E. Dorre and H. Hiibner, Alumina: Processing, Properries, and Applications. Springer-Verlag, Berlin, 1984. ’Advanced Ceramic Powders and Nan0 Ceramic Powders, Business Communications Co. Inc. Report, Thomas Abraham, 2003. 4H. L. Wen and F. S . Yen, “Growth Characteristics of Boehrnite Derived UItrafine Theta and Alpha-Alumina Particles During Phase Transformation,” J . Crysr. Growrh, 208, 696-708 (2000). ’Y. Sarikaya and M. Akinc, “Preparation of Alumina Microshells by the Emulsion Evaporation Technique,” Ceram. In(., 14 [4] 239-44 (1988). 6Y. Sarikaya, T. Aiemdaroglu, and M. Onal, “Determination of the Shape, Size and Porosity of Fine a-A1203 Powders Prepared by Emulsion Evaporation,” J . Eur. Ceram. Soc., 22 [3] 305-9 (2002). 7P. Lin, S. B. Wen, and T. T. Lee, “Preparation of Nanometer Sized u-Alumina Powders by Calcining an Emulsion of Boehmite and Oleic Acid,” J. Am. Ceram. Soc., 85 [I] 129-33 (2002). ‘C. P. Lin and S . 9. Wen, “Variation of Boehmite and Oleic Acid Emulsion in the Calcining Process of Alumina Powders,” J . Am. Ceram. Soc., 85 [6] 1467-72 (2002). ’D. M. Moore and R. C. Jr. Reynolds, X-Ray Diffracrion and [he Idenr@carion and Analysis ofClay Minerob, 2nd edition, p. 87. Oxford University Press, Oxford, 1997. ‘OK.Wefers and C. Misra, “Oxides and Hydroxides of Aluminum”; Alcoa Technical Paper, No. 19, Revised Alcoa Laboratories, Pittsburgh, PA, 1987. “Y.Sarikaya, 1. Sevinc, and M. Akinc, “The Effect of Calcination Temperature on Some of the Adsorptive Properties of Fine Alumina Powders Obtained by Emulsion Evaporation Technique,” Powder Techno/., 116 [I]109-14 (2001).

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Lanthanum Strontium Manganite Powders Synthesized by Gel-Casting for Solid Oxide Fuel Cell Cathode Materials Lan Zhang and Yujun Zhang Key Lab for Liquid Structures, School of Materials Science and Engineering, Shandong University, Jinan 25006 1, China

Y. D. Zhen and San Ping Jiangt School of Mechanical and Aerospace Engineering, Nanyang Technological University, Singapore 639798, Singapore

Lanthanum strontium manganite ((La,,$r,&.gMn03; LSM) powder was successfully synthesized by an aqueous gel-casting technique, using carbonaceous precursors. Both thermal and Xray diffraction analysis confirmed that the gel-casting LSM powder formed a single perovskite phase at 850°C, which is 1W-150"C lower than that of the LSM powder prepared by the conventional solid-state reaction route. The significantly reduced phase formation temperature of the gel-casting LSM powder is most likely due to the homogeneously distributed and immobilized precursor particles in a polymeric network, promoting the sintering and crystallization process. The LSM electrode prepared by the gel-cast LSM powder showed good electrocatalytic activity for the O2 reduction reaction for solid oxide fuel cells.

I. Introduction

S

OLID oxide

fuel cells (SOFCs) technology has gained considerable attention recently as the most efficient, versatile, and less polluting power-generating system.'32The efficiency of an SOFC stack is critically dependent on the electrical conductivity, chemical and thermal stability at the SOFC operating conditions, and electrochemical performance of the cell components. Among the different cell components, Sr-doped lanthanum manganite ((Lao.8Sro.2)o.yMn03; LSM) is considered as the preferred cathode material for the high-temperature SOFCs due to its high electrical conductivity, high thermal and chemical stability, and good electrocatalytic activity for the oxygen reduction reaction.'" However, the electrochemical catalytic activities of a LSM cathode depend significantly on the microstructure and the interfacial characteristics of the three phase boundary, which in turn depends on the ro rties of the starting powders and the preparation process.7,9 pe There are various processin techniques available for the preparation of LSM p o ~ d e r s . ~ ~Recently, ~ ~ * ~ *the ~ -gel-cast'~ ing technique has been used to prepare cell components of SOFCs such as NiOjSDC composite and yttria-stabilized zirconia (YSZ) and Ceo.yGdo.102electrolyte powder~.",'~ Gel-casting was first reported by Omatet et a1.17 in 1991. In this technique, slurry of high solid-state loading is obtained by dispersing the ceramic powders in the premixed monomer and cross-link solution and is cast into a mold of the desired shape. With heating or addition of catalyst, cross-linking and solidification occur to form a three-dimensional polymeric network structure. Polymerization solidifies and immobilizes the powder into the desired shape. Gel-casting is a novel forming K. Zum Gdhr--contrihuting editor

Manuscript No. 22075. Received August 3, 2006; approved December 21, 2006. 'Author to whom correspondenceshould be addressed. e-mail: [email protected]

Synthesis Methods for Powders

method in fabricating complex three-dimensional ceramic parts and has significant advantages in comparison to the conventional forming methods.'871yOwing to the advantage of the high homogeneity, the gel-casting technique was employed to synthesize various ceramic powders." '6q20*2' In this paper, LSM powders were synthesized by the aqueous gel-casting technique, using carbonaceous precursors. The phase formation, microstructure, particle size, and electrochemical performance of the gel-cast LSM powders were investigated and characterized. For the purpose of comparison, the phase formation and sintering of LSM powder prepared by the conventional solid-state reaction method were also investigated. 11. Experimental Procedure

(1) LSM Powder Synthesis A-site nonstoichiometric LSM powders were synthesized by an aqueous gel-casting method using La203 (99.9%), SrC03 (99.9%), and MnC03 (990/,) (all from Sigma-Aldrich,St. Louis, MO) as precursors. The raw materials were weighed, and ball-milled in propanol for 24 h. A premix solution of monomers was prepared in deionized water with acrylamide, C2H3CONH2, (AM, Sigma-Aldrich) and N,K-methylenebisacrylamide, (C2H3CONH)2CH2, (MBAM, Sigma-Aldrich). The AM to MBAM ratio was 15. The precursors were mixed in the premix solution with SD-03 (ammonium poly-(methacrylate), PMAA) as a dispersant, followed by a planetary ball milling for 30 min. Ammonium persulfate ((NH&S208) (APS, Sigma-Aldrich) was added as the initiator and the slurry was poured into a container, and subsequently heated in an oven at 80°C for 24 h. After being dried, the green body was calcined at a temperature range of 700"-90O0C for 20 h in air. Figure 1 shows the procedures of the gel-casting method for the synthesis of LSM powders. For the purpose of comparison, LSM powders were also prepared by the conventional solid-state reaction technique. La2O3, SrC03, and MnC03 with desired composition were mixed in propanol and ball-milled for 24 h, followed by calcination at a temperature range of700"-900"C for 20 h in air.

( 2 ) Powder Characterization Differential thermal analysis (DTA) and thermogravimetric analysis (TGA) of the gel-cast and solid-state LSM powders were carried out on a 2950 thermal analyzer (TA Instruments Inc., New Castle, DE) with a heating rate of 10"C/min. Phase development of the gel-cast and solid-state LSM powders was determined by X-ray diffraction (XRD, Philips MPD 1880, Oxford, UK) using CuKu radiation at room temperature. Xray scans were run over a 28 spectrum of 20"-80" at a scan rate of 0.02"/s. The morphology of the LSM powder and LSM electrodes was examined by scanning electron microscopy (SEM, Leica S360, Wetzlar, Germany) after the electrochemical measurement. 71

Raw powders

I

Organic Monomer Cross-Link agent (AM, MBAM)

Dispersant (Water, PMAA)

4 (Slurry Formation)

Gelation

Fig. 2.

Fig.1. Procedures of the gel-casting method for the synthesis of

lanthanum strontium manganite powders.

( 3 ) Electrochemical Characterization The electrolyte substrates were prepared by sintering die-pressed disks of 8 mol% YSZ (Tosoh, Shiba, Japan) at 1550°C for 2 h. The sintered substrate was 19 mm in diameter and 1 mm in thickness. Five wt% of graphite was added to the LSM slurry as a poreformer. LSM electrodes were applied to the YSZ disks by slurry painting and sintered at 1150°C in air for 2 h. Platinum paste was painted onto the other side of the electrolyte disk to make the counter and reference electrodes. The counter electrode was symmetrically positioned opposite of the working electrode and the reference electrode was painted as a ring at the edge of the electrolyte substrate. The gap between the counter and ring reference electrodes was 4 mm. Pt mesh was used as the current collector for both working and counter electrodes. The cell was mounted inside the furnace with LSM electrode exposed to an airflow at 100 mL/ min and reference electrode was exposed to open air. The initial electrode behavior of freshly prepared LSM electrodes with cathodic polarization time was studied against the Pt air reference electrode by galvanostatic current interruption (GCI) and electrochemical impedance spectroscopy (EIS) techniques. The polarization performance was carried out under a constant cathodic current of 200 mA/cm2 at 800°C and the cathodic polarization potential (Ecathode) was recorded with a data-logging device. The polarization was imterrupted from time to time to make GCI and EIS measurements. Overpotential (q) and electrode ohmic resistance (Rn) between the LSM electrode and the Pt reference electrode were directly measured from the GCI transition curves. EIS was performed on a Solartron 1260 frequency response analyzer (Farnborough, UK) in conjunction with a 1287 electrochemical interface at a frequency range of 0.01 Hz to 1 MHz with the signal amplitude of 10 mV under open circuit. The electrode was equilibrated at open circuit for 10 min before the EIS measurement. The electrode interface (polarization) resistance (RE) was directly measured by the differences between the low- and high-frequency intercepts on the impedance axis. The details of the electrochemical measurements of LSM electrodes for the O2 reduction reaction under SOFC operating conditions have been given elsewhere.22

Decomposition characteristics of lanthanum strontium manganite precursor samples of (a) thermogravimetricanalysis (TGA) of the gel-casting precursor, (b) TGA of solid-state reaction precursor, (c) differential thermal analysis (DTA) of the gel-casting precursor, and (d) DTA of the solid-state reaction precursor. residual water in the polymer network and raw powder mixtures. As shown in Fig. 2, the major decomposition of the gel occurred in the temperature range of 300"-5OO0C, giving a rise of a very sharp endothermic peak in the DTA spectra around 400°C and a weight loss of 20.05% in the TGA. Such decomposition step indicates the burnout of the polymer network, removal of bounded molecular water and decomposition of carbonaceous precursors. Similar DTA and TGA behavior for the powder mixture before the solid-state reaction were also observed in the temperature range of 300"-500"C. The weight loss of 16.59% for the solid-state reaction specimens is lower, in comparison with that of the gel-casting specimens, indicating the decomposition of carbonaceous precursors and endothermic removal of the bounded molecular water. For the gel-casting specimens in the temperature range of 650"-850"C, there was a weight change of 8.27%, accompanied by two endothermic DTA peaks at 750" and SWC, respectively. The shallow endothermic peak at 750°C may be due to the decomposition of SrC03.23The rather sharp endothermic peak at 800°C could be originated from the formation of the LSM perovskite phase. Further increase of the temperature did not lead to the weight change in the TGA curve and no endothermic DTA peaks were observed (Figs. 2(a) and (c)). This indicates that the perovskite phase may be formed for the gel-casting specimens at temperatures around 800"-850"C. In the case of the solid-state reaction specimens, the weight loss was 5.09% in the TGA plot and there were two endothermic DTA peaks at 722" and 800°C in the same temperature range of 650"-850"C. However, both weight change and the endothermic peaks are smaller than that for the gel-casting specimens. Different to that of the gel-casting specimens, there was an additional endothermic peak in the DTA curve at 944°C. This may indicate that the formation of perovskite phase for the solid-state sample may not be completed at temperatures of 8O0"-85O0C. The significantly separated endothermic peaks appearing at 800" and 944°C show that the formation of LSM perovskite phase of the powder prepared by the solid-state method would occur in two stages.

111. Results and Discussion

( I ) Thermal Analysis of the Gel and Solid-State Reaction Precursors Figure 2 shows the DTA and TGA curves of the gel after heat treatment at 80°C for 24 h and the powder mixture before the solid-state reaction. At temperatures below 150"C, there is a p proximately 1.89% and 1.03% mass loss for the gel-casting specimens and the solid-state reaction powder mixture, respectively. This is most likely caused by endothermic removal of the 72

(2) Effect of Temperature on Phase Formation In order to investigate the phase evolution of gel-casting and solid-state reaction powder, XRD measurements were carried out on samples heat treated at different temperatures (Fig. 3). In comparison, XRD patterns of raw materials, SrC03, MnC03, and La2O3, are shown in Fig. 3(c). As evidenced from Fig. 3, for the gel-casting and solid-state reaction powders heat treated at 700"C, SrC03 phase was detected. This confirms that the Progress in Nanotechnology: Processing

Fig. 3. X-ray diffraction (XRD) patterns of the lanthanum strontium manganite powder synthesized by (a) gel-casting and (b) solid-state reaction techniques at different temperatures. The XRD patterns of raw materials are shown in (c). The solid symbols indicate SrCO, and the empty symbols indicate Lalo,. The insets are the enlarged XRD patterns at 28 ranges from 25" to 31".

endothermic DTA peaks around 700-750"C (Fig. 2) are due to the decomposition of SrCO3. All the high-intensity reflections in the gel-casting sample calcined at 850°C were identified to the reported pseudocubic phase of LSM, indicating the formation of single perovskite phase LSM (Fig. 3(a)). On the other hand, the presence of La203 was clearly detected in the powder prepared by the conventional solid-state reaction and calcined a t 900°C (see the inset in Fig. 3(b)). This shows that the phase was much more pure in the case of gel-casting LSM powder in comparison with that of the solid-state reaction LSM powder. The exact process associated with the endothermic peak above 900°C needs further study. However, we have shown before that the formation of a single perovskite phase for LSM powder prepared by a wet chemical synthesis method such as co-precipitation takes place at 1000"C.24This shows that the endothermic DTA peak at 944°C for the solid-state reaction precursor sample may be related to the formation of the LSM perovskite phase. The results in the present study indicate that the phase formation for the gel-cast LSM powder occurs at 850"C, 100°-1500C lower than that of the LSM powder prepared by the solid-state reaction route. The significantly reduced phase formation temperature in the LSM synthesized by the gel-casting technique is most likely due to the much better homogeneity, achieved through the immobilization of the precursor particles by the

-

Synthesis Methods for Powders

cross-linked polymeric network formed during the cross-linking and gel formation steps. The homogeneously distributed and immobilized particles in a polymeric network are in close vicinities of each other, enhancing the sintering and crystallization process. On the other hand, the solid-state reaction powders are loosely contacted and thus, higher calcination temperature would be needed to enhance the solid-state diffusion and sintering of the solid-state reaction powder for the crystallization and phase formation of the LSM perovskite. ( 3 ) Polarization Behavior Figure 4 shows the initial impedance and polarization responses of a LSM electrode prepared by gel-casting powder under a cathodic current density of 200 mA/cm2 at 800°C in air. Before passing cathodic current, the impedance response was characterized by a large and depressed arc and the initial RE was 7.76 f2.cm2. The impedance arcs were reduced significantly with the current passage. After current passage for 15 min, RE was reduced to 1.8 R.cm2, much smaller than the initial RE value of 7.76 R . cm2. The initial polarization potential of the LSM electrode behaved similarly to the impedance behavior. The initial Ecathode was very high, 673 mV at 200 mA/cm2 before current passage.

73

Fig. 4. Initial impedance and polarization responses of a lanthanum strontium manganite electrode prepared by the gel-casting powder as a function of the cathodic current passage at 200 mA/cm* and 800°C in air.

Fig. 5. Initial impedance and polarization responses of a lanthanum strontium manganite electrode prepared by the solid-state reaction ow der as a function of the cathodic current passage at 200 mA/cmPand800°C in air.

With a cathodic current passage, the polarization potential decreased very rapidly initially, followed by a region where the decrease in b&thode was much smaller. After cathodic current treatment for 15 min, the &tho& started at 407 mV and reduced to 344 mV in 10 min of the current passage. However, after current passage treatment for 30 min, the change in Ecathode with the current passage time became relatively small. The change in Rn was negligible, indicating that the reduction in &tho& was mainly due to the reduction in overpotentials, q. Such impedance and polarization behavior with the cathodic current passage are the well-known activation rocesses for the oxygen reduction on the LSM The initial impedance and polarization responses of a LSM electrode prepared by the solid-state reaction powders under a cathodic current density of 200 mA/cm2 at 800°C in air are very similar to that of the gelcasting LSM powder (Fig. 5). The impedance arc decreased significantly with the cathodic current passage. After the current passage for 15 min, RE was reduced to 1.77 0 .cm2, almost the same as that of the LSM electrode prepared by the gel-casting powder. The polarization performance of the LSM electrode prepared by the solid-state reaction also behaved similarly to that of the electrode prepared by the gelcasting LSM powder. The similar electrochemical performance of the LSM prepared from the gel-casting and solid-state reaction powder is most likely due to the high electrode sintering temperature of 1150°C. High electrode sintering temperature would diminish the benefits of the low formation temperature of

the LSM powder prepared by the gel-casting technique in comparison with that prepared by the solid-state reaction technique. The significant reduced formation temperature for the gelcasting LSM powder implies that LSM electrode may be fired at reduced temperatures, forming finer microstructure and thus enhancing the electrocatalytic activities of the LSM. To confirm this, a LSM electrode prepared from the gel-casting powder was sintered at IOO0"C. The impedance and polarization behaved similarly to that of the LSM electrodes sintered at 1150°C. However, the gel-casting LSM electrode sintered at 1000°C shows significantly high activity as compared with that sintered at 1150°C. Figure 6 compares the impedance responses of the LSM electrodes sintered at 1150" and 1000°C.At 800°C, R E was 1.03 Q . cm2 observed 0.27 Q cm2, substantially lower than for the LSM electrode sintered at 1150°C. The result indicates that the gel-casting LSM powder reduces the electrode sintering temperature and enhances the electrocatalytic activity for the oxygen reduction reaction under SOFC operating conditions.

electrode^^^^^^^

74

-

( 4 ) Movphology and Microstructure of LSM Powder and Electrode Figure 7 shows the SEM micrograph of the LSM powder synthesized by the gel-casting technique and calcined at 850°C. The LSM powder was characterized by irregular shapes and the particle size was in the range of 7&300 nm. The irregular shape and submicrometer size of the LSM powders indicate the high Progress in Nanotechnology: Processing

Fig. 6. Impedance responses of the lanthanum strontium manganite (LSM) electrodes prepared by the gel-casting and solid-state reaction powder and sintered at 1150" and 1000°C. The impedance spectra were measured at 800°C in air after the cathodic current passage at 200 mA/cm2 for 240 min.

activity of the gel-cast LSM powder. Figure 8 shows the SEM micrographs of the LSM electrodes prepared by the gel-cast and solid-state reaction powders after the electrochemical testing. The average LSM grain of the electrode prepared by the gelcasting powder was 0.6+0.2 pm, smaller than 1.1 k0.4pm of the LSM electrode prepared by the solid-state reaction powder. The smaller grain size of the LSM electrode prepared by the gelcasting LSM powder is most likely due to the submicrometer and very fine LSM powder synthesized by the gel-casting technique (Fig. 7).' However, in the case of LSM electrode sintered at high temperatures (1 1WC), the difference in the microstructure is small and might not have substantial effect on the electrochemical catalytic activity of the LSM electrodes for the oxygen reduction reaction (see Figs. 4 and 5). Nevertheless, the significant reduced phase formation temperature of the gelcasting LSM powder shows that it is possible to reduce the electrode sintering temperature, thus forming finer microstructure and enhancing the electrocatalytic activities of the LSM cathode, as shown in Fig. 6. Gel-casting has been widely used in the forming of complex ceramic parts owing to the advanta es of high homogeneity, simplicity, and relatively low ~ o s t . ~ The ~ *flexural ~ ' ~ strength ~ ~ ~ ~ and Weibull modulus can be improved by optimizing the

Fig. 7. Scanning electron micrograph of lanthanum strontium manganite powder synthesized by the gelcasting and calcined at 850°C.

Synthesis Methods for Powders

Fig. 8. Scanning electron micrographs of the lanthanum strontium rnanganite (LSM) electrodes prepared by the gel-casting and solid-state reaction LSM powders after the electrochemical testing. The LSM electrodes were sintered at 1150°C for 2 h in air.

rheological properties of the suspension in the gel-casting process and subsequent sintering The drawback of gelcasting is that the solidification of the suspension is usually induced by the mold temperature; thus, the solidification is not uniform in large complex parts due to the existence of temperature gradient, resulting in inner stress at the forming stage and microcracks on subsequent drying and debinding processes. These defects will become fracture origins and hence reduce the strength and reliability of the product.30 However, such drawback of the gel-casting technique in the forming of complex ceramic parts is not a concern in the powder synthesis. In the present study, it has been shown clearly that the aqueous gelcasting technique can be used to significantly reduce the phase formation temperature of the oxide compound powders. In the case of LSM, the formation of the perovskite phase occurs at 850"C, which is 100°-150"Clower than that synthesized by the sold-state reaction and co-precipitation methods.24The low sintering and phase formation temperature of the LSM powder show that it is possible to significantly reduce the electrode sintering temperature, thus reducing the grain growth and sintering of the LSM particles. This in turn leads to the formation of much finer microstructure, significantly enhancing the electrochemical activities of the LSM cathodes for the 0 2 reduction reaction. This is just the case as shown in Fig. 6. The results in the present study demonstrate the great potential of the aqueous gel-casting technique in the development of nano-sized electrode and electrolyte powders for SOFCs with significantly reduced phase formation temperature. 75

IV. Conclusions LSM powder was successfully synthesized by an aqueous gelcasting technique using carbonaceous precursors. Both thermal analysis and XRD analysis confirmed the phase formation of the gelcasting LSM powder at 850°C. The LSM powder is characterized by irregular shape with particle size in the range of 70-300 nm. On the other hand, in the case of the LSM powder prepared by the conventional solid-state reaction route, La203 was clearly detected for the powder calcined at 900°C. An additional endothermic peak at 944°C indicates that the complete formation of a single perovskite phase for the LSM powder prepared by the solid-state method would be around 950’1000”C?4 The significantly reduced phase formation temperature of the gel-casting LSM powder is explained by the homogeneously distributed and immobilized precursor particles in a polymeric network, in contrast to the loosely contacted particles in the powder mixture prepared by the solid-state reaction technique. The results in the present study demonstrate the potential of the gel-casting technique in the development of nano-sized powders with significantly reduced sintering and phase formation temperature for SOFCs.

References IT. L. Wen, D. Wang, M. Chen, H. Tu, Z. Zhang, H. Nie, and W. Huang, “Material Research for Planar SOFC Stack,” Solid Slate lonics, 148, 513-9 (2002). *N.Q. Minh, “Ceramic Fuel Cells;’ J. Am. Cerom. Soc.. 76, 563-88 (1993). ’S. P. S. Badwal and K. Foger, “Materials for Solid Oxide Fuel Cells,” Marer. Forum, 21, 187-224 (1997). 4H. Yokokawa, “Understanding Materials Compatibility,” Annu. Rev. Morer. Res., 33. 581610 (2003).

5N.P. Brandon, S. Skinner, and B. C. H. Steele, “Recent Advances in Materials

for Fuel Cells,” Annu. Rev. Marer. Res., 33, 183-213 (2003). 6S. P. Jiang and W. Wang, “Sintering and Grain Growth of (LaSr)MnOl Electrodes of Solid Oxide Fuel Cells Under Polarization,” Solid Srare lonics, 176, 1185-91 (2005).

’S. P. Jiang, “Issue on Development of (La,Sr)MnOl Cathode for Solid Oxide Fuel Cells,” J. Power Sources, 124, 390-402 (2003). ‘S. B. Adler, “Factors Governing Oxygen Reduction in Solid Oxide Fuel Cell Cathodes,” Chem. Rev., 104, 4791-843 (2004). 9S. Charojrochkul, K.-L. Choy, and B. C. H. Steele, “Cathode/Electrolyte Systems for Solid Oxide Fuel Cells Fabricated Using Flame Assisted Vapour Deposition Technique,” Solid State lonics, 121, 107-13 (1999). “A. Kumar, P. S. Devi, A. D. Sharma, and H. S. Maiti, “A Novel Spray-Pyrolysis Technique to Produce Nanocrystalline Lanthanum Strontium Manganite Powder,” J. Am. Cerum. Sor., 88,971-3 (2005).

76

“T. Fukui, T. Oobuchi, Y. Ikuhara. S. Ohara, and K. Kodera, “Synthesis of (La,Sr)MnO,-YSZ Composite Particles by Spray Pyrolysis,” J. Am. Cerom. Soc.,

80. 261-3 (1997).

I2 M. Gaudon, C. Laberty-Robert. F. Ansart. P. Stevens. and A. Rousset. “Preparation and Characterization of Lal_,Sr,MnO,+s. ( 0 ~ ~ ~ Powder 0 . 6 ) by SoCGel Processing,” Solid Stare Sci., 4, 125-33 (2002). ”Y. H. Yin, S. Y. Li, C. R. Xia, and G. Meng, “Electrochemical Performance of Gel-Cast NiO-SDC Composite Anodes in Low-Temperature SOFCs,” Elecfrochim. Arm, 51, 2 5 9 M (2005). I4Y. Yin, W. Zhu, C. Xia, and G . Meng, “Gel-Cast NiO-SDC Composites as Anodes for Solid Oxide Fuel Cells,” J. Power Sources, 132, 36 41 (2004). I5X. Liu, G . Li, J. Tong, and D. Chen. “Low Cost Fabrication and Characterization of Thin-Substrate YSZ Solid Electrolyte,” Ceram. h t . , 30, 2057-9 (2004). I6J.-G. Cheng, S.-W. Zha, J. Huang, X.-Q. Liu, and G.-Y. Meng, “Sintering Powder Prepared by the Behavior and Electrical Conductivity of Ce,,&do Gel-Casting Process,” Marer. Chem. Phys., 78, 791-5 (2003). ”0. 0. Omatete, M. A. Janney, and R. A. Strehlow, “Gelcasting-A New Ceramic Forming Process,” Am. Cerum. Soc. Bull., 70, 1641-9 (1991). “A. C. Young, 0. 0. Omatete, M. A. Janney, and P. A. Menchhofer, “Gelcasting of Alumina,” J. Am. Ceram. Soc., 74. 612-8 (1991). I9M. Wei, D. Zhi, and D. G . Brandon, “Microstructure and Texture Evolution in Gel-Cast Alpha-Alumina/Alumina Platelet Ceramic Composites,” Scripra Muter., 53. 1327-32 (2005). 20 Y. Yanhong, Y. Shuting, X. Changrong, and M. Guangyao. “Performance of C q 8Sw,20, (SDC) Prepared from Different Method,” J. Henan Norm. Liniv. (Nar. Sci.), 33, 5 6 9 (2005). ”H. Varma, S. P. Vijaya, and S. S. Babu,’“Transparent Hydroxydpatite Ceramics Through Gelcasting and Low-Temperature Sintering.” J. Am. Ceram. Soc.,

85,493-5 (2002).

22S. P. Jiang, “A Comparison of 0 2 Reduction Reactions on Porous (La.Sr) MnOl and (La,Sr)(Co,Fe)Ol Electrodes,’’ Solid Stare lonics, 146, 1-22 (2002). 23 S. W. Zha, C. Xia, X. Fang, H. Wan& and G . Meng, “Processing and Electrical Properties of Doped -LaGaO, by Gelcasting,” Cerum. lnr., 27,649-54 (2001). 24 S. P. Jiang, J. P. Zhang, Y. Ramprakash, D. Milosevic, and K. Wilshier, “An Investigation of Shelf-Life of Strontium Doped LaMn03 Materials,” J. Muter.

Sci., 35, 2 7 3 H 1 (2000).

*’S. P. Jiang and I. G. Love, “Origin of the Initial Polarization Behavior of SrDoped LaMnO] for O2 Reduction in Solid Oxide Fuel Cells,” Solid Stare lonics, 138, 183-90 (2001).

26 S. P. Jiang, I. G. Love, J. P. Zhang, M. Hoang, Y. Ramprakash, A. E. Hughes, and S. P. S. Badwal, “The Electrochemical Performance of LSM/ZirconiaYttria Interface as a Function of A-Site Non-Stoichiometry and Cathodic Current Treatment,” Solid State lonics, 121, 1-10 (1999). 27W.Wang and S. P. Jiang, “A Mechanistic Study on the Activation Process of (La,Sr)MnO, Electrodes of Solid Oxide Fuel Cells,” SolidSrare lonics. 177, 1361-9

(2006).

2RX.Liu, Y. Huang, and J. Yang, “EtTect of Rheological Properties of the Suspension on the Mechanical Strength of Al20,-ZrO2 Composites Prepared by Gel Casting,” Ceram. lnf., 28, 15944 (2002). 29L.Zhou, Y. Huang, Z. Xie, A. Zimmermann, and F. Aldinger, “Preparation of S I N 4 Ceramics With High Strength and High Reliability Via a Processing Strategy,” J. Eur. Cerum. Soc., 22, 1347-55 (2002). ’OF. F. Lange, “Powder Processing Science and Technology for Increased Reliability,” J. Am. Cerum. Soc., 72, 3-15 (1989).

Progress in Nanotechnology: Processing

Preparation of Matrix-Type Nickel Oxide/Samarium-Doped Ceria Composite Particles by Spray Pyrolysis Seiichi Sudat and Koichi Kawahara Materials R&D Laboratory, Japan Fine Ceramics Center, Nagoya 456-8587, Japan

Mitsunobu Kawano, Hiroyuki Yoshida, and Tom Inagaki Energy Use R&D Center, The Kansai Electric Power Company Inc., Hyogo 661-0974, Japan

Matrix-type nickel oxide (NiO)/samarium-doped ceria (SDC) composite particles, in which NiO and S D C nano-particles were homogeneously dispersed, were synthesized by spray pyrolysis (SP) for an anode precursor of intermediate-temperaturesolid oxide fuel cells (IT-SOFCs). SP of an aqueous solution containing Ni, Ce, and Sm salts resulted in capsule-type composite particles that had NiO enveloped with SDC. The capsule-type composite particles actually prevent Ni aggregation between particles, but they cannot have a large contact area between nickel (Ni) and SDC. A matrix-type composite particle is expected to have a large contact area because the matrix-type composite is comprised of nanometer-sized Ni and S D C particles. An adequate addition of ethylene glycol successfully resulted in matrix-type NiO/SDC composite particles. The matrix-type composite particles also showed higher anode performance than the capsule-type composite particles in these experiments and they were effective as precursors of highperformance IT-SOFC anodes.

I. Introduction

I

NTERMEDIATE-temperatUre solid oxide fuel Cells (IT-SOFCS) have been drawing considerable attention because IT-SOFCs improve cost performance, which is one of the biggest issues for practical application of fuel cells. Lowering the operating temperature would make it possible to use inexpensive refractory metals or ceramics and reduce heat damages caused by holding SOFC cells at high temperatures for several months or years. SOFC performance, such as power densities and fuel efficiencies, are extremely decreased with a decrease in operating temperature. Many researchers have developed highly efficient ITSOFCs mainly from two perspectives: (1) use of highly conductive solid electrolytes such as doped cerial-'" and lanthanum gallates,"- l6 and (2) application of anode- or cathode-supported cell design using a thin e l e ~ t r o l y t e . l ~Recently, -~~ the cell performance of IT-SOFC has been improved markedly by attentive researchers, but it has not shown sufficient power density to afford practical applications yet. One of the reasons for the insufficient power densities is the relatively high electrode polarizations. The polarizations of both the anode and cathode increase steeply with a decrease in operating temperature, and the temperature dependence of electrode polarization is much larger than that of a solid electrolyte. Electrodes with sufficiently low polarizations at relatively low temperatures are required to be applied to IT-SOFC for practical applications, and we have R. Cutler--contributing editor

Manuscript No. 21776. Received May 10, 2006; approved November 20, 2006. 'Author to whom correspondence should be addressed. e-mail: [email protected]

Synthesis Methods for Powders

Y

develo d SOFC anodes for IT-SOFC using composite partic1es."SOFC anodes mainly consist of a metal and an oxygen-conductive material. Nickel (Ni) is generally used as a metal component because Ni has high electronic conductivity larger than lo5 S / C ~excellent , ~ ~ hydrogen oxidation-catalytic properties, and a high-cost performance. More than 40 vol% Ni content in an SOFC anode would result in sufficient electrical conduction and catalytic properties. Samarium-doped ceria (SDC) is also one of the most promising materials as an oxygen conductor in an SOFC anode because SDC has a high ionic conductivity larger than 3 x lo-' S/cm at 700°C25depending on samarium composition. SDC also has oxidation-catalytic properties26and might exert high anodic activities due to rapid electrochemical oxidation of hydrocarbon gas. Therefore, cermets composed of Ni and SDC would be one of the most promising candidates as an IT-SOFC anode. However, Ni has some drawbacks as an SOFC anode material. Ni has a high thermal expansion coefficient as compared with SDC, and an interface between Ni and SDC is fragile on not only redox cycles but also thermal cycles owing to thermal expansion mismatch. Ni is also easy to sinter at low temperatures and the sintering would cut off the electrical skeleton network of an SOFC anode, and shorten three-phase boundaries (TPBs). Thus, the Ni sintering causes degradation of anode performance and would prevent long-term SOFC operation. Moreover, Ni reacts with some high-conductivity electrolytes.16To resolve these issues, we have studied nickel oxide (NiO)/SDC composite particles as a precursor of SOFC anodes. The microor nano-composition can enlarge contact areas between Ni and SDC and suppress the Ni sintering. As precursor particles of NiO/SDC composite anodes, we have prepared NiOjSDC composite particles by the spray pyrolysis (SP) method. An aqueous solution containing Ni, Ce, and Sm salts was atomized ultrasonically and the atomized mists were transferred to an electrical furnace zone with dry air as a carrier gas. The furnace zone was composed of several furnaces arranged in series, and the mists were fired by passing through the furnaces continuously. After drying the mists, decomposition, solid reaction, and sintering, NiOjSDC composite particles were obtained by SP. The composite particles had a capsule-type morphology that had NiO enveloped with SDC. The SOFC anode derived from the capsule-type composite particles sup ressed aggregation of Ni and showed low anodic polarization?'The formation mechanism of capsule-type NiO/SDC composite particles during pyrolysis of the mists and the formation processes from the composite particles to Ni and SDC skeleton networks in the SOFC anode are still vague, but the composite was rather effective in obtaining a high-performance anode. Capsule-type composite particles actually prevent the particles from aggregation because the SDC shells had a relatively high sintering temperature and the SDC would suppress Ni aggregation. But a capsule-type composite particle does not pro77

gated in order to design a composite conformation suitable for high-performance IT-SOFC anodes.

11. Experimental Procedure Samarium oxide (99.9%, Shin-Etsu Chemical Co. Ltd., Tokyo, Japan), cerium nitrate (98.0%, Wako Pure Chemical Ind. Ltd., Osaka, Japan) and nickel acetate (98.0%, Nacalai Tesque Inc., Kyoto, Japan) were dissolved in a nitric acid solution. The batch composition of SDC was (SmOl.5)o2(CeOz)o8 , and the molar ratio of NiOjSDC was adjusted to be 5.6. We then defined the amount of EG (99.5%, Wako Pure Chemical Ind. Ltd.) as the molar ratio against all cations of Ce, Sm, and Ni in the solution. The chelation by addition of EG is known as the Pechini process (Fig 2). The Pechini process allows chelation by heating a concentrated nitric acid solution at temperatures < IOO"C, and each EG molecular can chelate a cation. A ratio of EG to total cations > 1.0 would be necessary to chelate all cations in the solution. But SP allows a non-equilibrium rapid thermal process and will lead to some different results from the synthesis under mild conditions. In this study, the ratio of EG/cation was adjusted to be 0.25, 0.50, 1.0, and 5.0 in order to obtain some information on the relationship between the degree of chelation and the morphology of NiO/SDC composite particles. As for heating for the chelation with the solution containing EG, we investigated two processes. One process aimed at the chelation during SP. The solution without heating was diluted with deionized water until the NiOjSDC concentration was 0.1 mol/ dm3 and the EG-mixed solution was spray pyrolyzed directly. The other aimed at sufficient chelating before the pyrolysis by heating the solution. The solution containing EG was heated to 80°C for 5 h in air to allow sufficient chelation. Several minutes after the solution was heated to 8 0 T , it generated a heat by the

Fig. 1. Morphology of Ni/SDC composite particles capsule type versus matrix type. SDC, samarium-doped ceria; Ni, nickel.

vide a large contact area between Ni and SDC. As compared with capsule-type composites, a matrix-type composite particle is expected to have a large contact area because the matrix-type composite is comprised of nanometer-sized Ni and SDC particles (Fig. I). SOFC anodes derived from the matrix-type composite particles can have improved anodic polarization properties relative to the capsule-type ones. Matrix-type NiOjSDC composite particles were thus synthesized by SP using a chelating reagent. The use of an aqueous solution composed of Ni, Sm, and Ce salts resulted in capsuletype composites. The formation mechanism of capsule-type composites is still unclear, but the difference in mass transfer rates between NiO and SDC during firing could possibly lead to separate inner NiO and outer SDC. Therefore, the addition of a binding agent among Ni, Sm, and Ce can result in a matrix-type composite. In this study, ethylene glycol (EG) was used as a binding agent, and the effects of EG on the morphology of NiO/ SDC composite particles were examined by SP. The synthesis conditions and electrochemical performance were also investi-

H

H

I

1

-

HO - C - C - O H

I

I

H

H

A

+

M+2 = > I HNO3

..*...* M .* .*

0

No

\\ /c

HO

- \ OH

+

H

H

I

1

HO-C-C-OH

I

I

H

H

c-c + H

nH20

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n Fig. 2. Chelation scheme by the Pechini process.

Fig. 3. SEM images of NiOjSDC composite particles. The particles were synthesized by SP using EG-containing solutions at an EG ratio to the cations: (a) 0.25, (b) 0.50, and (c) 1.00. SDC, samarium-doped ceria; SEM, scanning electron microscope; EG, ethylene glycol; SP, Spray pyrolysis.

78

Progress in Nanotechnology: Processing

Fig.4. SEM image of Nio/sDC composite Particles. The Particles were synthesized by spray pyrolyzing diluted chelated solutions at an EG/cation ratio of 5.0. SDC, samarium,oped ceria; SEM, scanning electron microscope; EG, ethylene glycol.

exothermaic chelating reaction, and the temperature of the solution temporarily reached 120°C. After cooling to room temperature, the chelated solution was then diluted with deionized water and the diluted solution was spray pyrolyzed. In the SP instrument, an aqueous solution was atomized with an ultrasonic mist generator with an oscillation frequency of 1.7 MHz. The solution was placed in an atomizing vessel and the atomizer transformed the solution into a few micrometer-sized mists. Ultrasonic atomizers increased the temperature of the solutions. Cold water was then circulated around the atomizing vessel because atomization at a relatively high temperature led to inhomogeneous mists.23 The mists were transferred into a firing zone using dry air as a carrier gas. Four furnaces were arranged in series, and the temperatures of the furnaces were set at 200", 400", SOO", and 1000°C. The mists were dried, decomposed, reacted, and sintered during passage through the various furnaces. The composite particles were thus obtained continuously. The particles were collected using a membrane filter with 0.2 pm pores, and calcined at 1000°C for 24 h in air. The morphologies of the particles and anode microstructures obtained with the particles were observed with a scanning electron microscope (SEM, Hitachi S4500, Tokyo, Japan) and a transmission electron microscope (TEM, Topcon 002EB, Tokyo, Japan), The compositions of the composite particles were investigated with an energy-dispersive X-ray spectrometer equipped with TEM (EDS/TEM).

Fig.5. TEM image and EDSFEM analysis of a NiOjSDC composite particle. The particle was synthesized with chelated solutions at an EG/cation ratio of 5.0. TEM, transmission electron microscope; SDC, samarium-doped ceria; EG, ethylene glycol; EDS, energy dispersive X-ray spectrometer.

Synthesis Methods for Powders

79

Fig. 6. TEM image and EDS/TEM analysisof NiOjSDC flakes. The NiOjSDC flake was synthesized by spray pyrolyzing an EG-containingsolution at an EG/cation ratio of 1.OO.TEM, transmission electron microscope; SDC, samarium-doped ceria; EG, ethylene glycol; EDS, energy dispersive X-ray

spectrometer.

SOFC single cells were prepared by attaching the NiOjSDC composite particles on an electrolyte and a cathode on the other side of the electrolyte, and SOFC performances were examined for the single cells. Lao.sSro.,Gao.8Mgo.z03-~ (LSGM) disks with a thickness of 0.2 or 0.1 mm were prepared as a solid electrolyte.

Fig. 7. Power densities and cell voltages of SOFC single cells composed of the anodes synthesized using EG. The thickness of the LSGM electrolyte was adjusted at 200 pm. The measurement temperature was 750°C. NiOjSDC composite particles were synthesized with either diluted-chelated solutions at an EG/cation ratio of 5 . 0 ( 0 , 0 ) or EGmixed solutions at an EG/cation ratio of 1.0 (0,+), 0.5( O, ), and 0.25 (A, A). SOFC, solid oxide fuel cells; LSGM, LaoSSrO ~ M g o@ - s ; SDC, samariumdoped ceria; EG, ethylene

glycol.

The anode paste was composed of polyethylene glycol (PEG, average molecular weight: 400, Wako Pure Chemical Ind. Ltd.) and the NiOjSDC composite particles were screen printed on the LSGM disk and sintered at 1250°Cfor 2 h in air. The anode paste was prepared by wet blending with the particles, PEG, and ethanol for 15 min using an alumina mortar. SEM observations of the composite particles before and after the wet blending indicated that this wet blending did not break up the composite morphology. After the anode paste was attached on the LSGM disk, the cathode paste composed of PEG and L%.zSro.8C003 (LSC, Seimi Chemical Co. Ltd., Kanagawa, Japan) was also screen printed on the other side of the LSGM disk and sintered at 1000°Cfor 4 h in air.*'When SOFC performance was investigated using LSGM disks with a thickness of 0.2 mm, a reference electrode was attached on the side of the disk with a Pt wire and glass flit-free Pt paste (Tanaka Precious Metals). Hydrogen humidified at 23°C (3%) was used as a fuel gas, whereas dry air was used an oxidation gas. After the complete reduction of NiO to Ni metal with 3% humidified hydrogen at 800°C for 30 min, SOFC performances and anode polarizations were measured at 650"-750"C. The anode polarizations were measured using the reference electrodes and the anodes by the current interruption method. Therefore, the ohmic loss estimated by the anode polarization measurement contains almost half the resistance of the LSGM electrolyte.

111.

Results and Discussion

EG was added in the nitric acid solution of Ni, Sm, and Ce salts at the EG/cation molar ratio of 0.25, 0.50, and 1.00, and the EG-mixed solutions were directly spray pyrolyzed to allow the partial or full chelation and the pyrolysis to occur almost simProgress in Nanotechnology: Processing

Fig.8. Ohmic loss and overpotential properties of SOFC anodes. The anodes were obtained NiOjSDC composite particles prepared with a dilutedchelated solution at an EG/cation ratio of 5.0 (@), and EG-mixed solutions at an EG/cation ratio of 1.0 (+),0.5 ( ), and 0.25 (A).SOFC, solid oxide fuel cells; SDC, samarium-doped ceria; EG, ethylene glycol.

ultaneously by firing into the series of furnaces. The NiOjSDC composite particles obtained were observed with SEM. Figure 3 shows the particles synthesized with EG containing mixed solutions at the various ratios. The mixed solution at the ratio of 0.25 resulted in round composite particles and the particle sizes were 0.3-3 pm. This particle size distribution is large as compared with capsule-type NiOjSDC composite particles synthesized with EG-free solutions. The composite particle size slightly increased with an increase in the EG ratio. The particles obtained with the mixed solutions at the EG ratio of 0.5 and 1.0 contained flakes and their fragments. The addition of E C would bind the cations by chelating, but the addition at the EG/cation ratio larger than 0.5 resulted in the formation of flakes. The EG/ cation ratio larger than 1.0 would lead to full chelation but formed a large amount of flakes. The amount of flakes increased with an increase in the EG/cation ratio, but SEM and TEM observations did not clarify the relationship between the EG/ cation ratio and the dispersion of NiO or SDC in the composite particles. The flakes had a large volume, which prevented the composite particles from forming a homogeneous anode paste for screen printing. Therefore, simultaneous chelation during SP without pre-heating would not be suitable for synthesis of matrix-type composite particles. The EG solutions were then heated to 80°C and the chelated viscous solutions were diluted with deionized water. The solution was then spray pyrolyzed as soon as possible to avoid hydrolysis of the solution with added water. Figure 4 shows an SEM image of NiOjSDC composite particles using the chelated solution at an EG ratio of 5.0. The composite particles had many dimples and some particles were broken into pieces. Small voids were observed in the particles, but solid particles were obtained by SP with the solution. The particle size distribution was relatively small and these particles would be applicable as anode precursor particles. In order to clarify whether the composite particles had matrix-type morphology or not, TEM observation was carried out. The composite particles obtained with the fully chelated solution at an EG/cation ratio of 5.0 were calcined at 1000°C for 24 h in air. The calcined particles were observed with TEM, and detailed compositions were investigated with EDS/TEM (Fig. 5). This TEM and EDS/ TEM images reveal a composite particle comprised of 50-100 nanometer-sized particles, and both Ni and Ce were dispersed together in the composite particle. The images in Fig. 5 are relatively unclear because the calcined particles were too large or thick to observe detailed composite morphologies with TEM or EDS/TEM. The solid particles obtained with the chelated solution are supposed to have a composite morphology similar to the flakes synthesized with the solution at an EG/cation ratio larger than 1.0. The flakes were then observed with TEM and Synthesis Methods for Powders

detailed compositions in the flakes were analyzed with EDS/ TEM because the flakes were thin enough to investigate with TEM. Figure 6 shows a TEM image and the results of EDS/ TEM for the flakes. These flakes were obtained by spray pyrolyzing the EG-mixed solution at an EG/cation ratio of 1.0. The flake also comprised 50-100 nanometer-sized particles. EDS/ TEM analysis showed that the small particles were either NiO or SDC, and that NiO and SDC particles were dispersed together in the flake. Apparently, this flake and the composite particle in Fig. 5 were matrix-type composite particles and flake-free matrix-type composite particles were obtained with the chelated solution. Therefore, adequate addition of EG in the solution for SP converted capsule-type to matrix-type composition particles. SOFC performance was then measured using the matrix-type NiO/SDC composite particles. Four kinds of matrix-type particles were prepared by the SP process with or without pre-heating. Solid particles were obtained by pre-heating the chelated solution at an EG/cation ratio of 5.0, whereas the composite particles containing the flakes were obtained by spray pyrolyzing directly the EG-mixed solutions at an EG/cation ratios of 0.25,0.50, and 1.0. SOFC single cells were fabricated using these NiOjSDC composite particles as anode precursor particles. Power densities and cell voltages for the single cells prepared with LSGM electrolytes with a thickness of 200 pm were measured at 750°C. Figure 7 shows the electrochemical performance of the single cells with anodes derived from various matrix-type composite precursors. The NiOjSDC composite particles obtained using the chelated solution at an EG/cation ratio of 5.0 showed the highest power densities as compared with other matrix-type composite particles. The relationship between SOFC performances and the morphology of the composite particles is still vague. For example, whether the flakes decreased SOFC cell performance or not, and whether pre-heating changed micro- or nano-dispersion in the matrix-type composite particles are under consideration. The flakes, which were not formed using the chelated solutions, had a large volume. The large volume has some advantages such as an easy diffusion of a fuel gas in the anode, but it may have a disadvantage related to the difficult formation of Ni or SDC skeleton network during reduction from NiO to Ni. If the volume of the flakes is too much, the flakes would prevent Ni or SDC from forming a skeleton network. Electrode polarizations for the anode derived from matrix-type composite particles were also measured at 750°C as shown in Fig. 8. The use of the chelated solution at an EG/cation ratio of 5.0 lowered ohmic loss and overpotential voltage as compared with other matrix-type composite anodes. The solid matrix-type composite particles with a relatively small particle size distribution appear to produce low anode polariza81

Fig. 9. Power densities and cell voltages of an SOFC single cell composed of anodes prepared with matrix type and capsule-type NiOjSDC composite particles. The thickness of the LSGM electrolyte was adjusted at 100 pm. The measurement temperatures were 650°C (A, A), 700°C (0, m), and 750°C (0,0 ) . SOFC, solid oxide fuel cells; SDC, samarium-doped ceria; LSGM, Lao.sSro.,Gao.sMgo.zO~_s.

tions. The NiOjSDC composite particles prepared with the EGmixed solution at an EG/cation ratio of 0.25 did not contain flakes, but the particles showed higher ohmic losses as compared with the particles prepared with the mixed solution at an EG/ cation ratio of 0.50 and 1.O. Backscattered SEM images of the

anode derived from various matrix-type particles did not show any apparent difference in the anode microstructure depending on the EG/cation ratios and volume of the flakes. The flakes may have some advantages as the formation of an Ni or SDC network. But the flakes also resulted in inhomogeneous anode

(a) SOFC' mode dcrived from matrix-t> pe paTtic1es.

Fig. 10. Backscattered SEM images of SOFC anodes derived from matrix-type and capsule-type NiOjSDC composite particles. SOFC, solid oxide fuel cells; SDC, samarium-doped ceria; SEM, scanning electron microscope.

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Progress in Nanotechnology: Processing

pastes and the inhomogeneous pastes would affect anode performance. Rheological investigation of anode pastes for screen printing would be additionally required to make good use of the flakes and to clarify the relationship between EG/cation ratios of the matrix-type particles prepared with the EG-mixed solutions and anode performance. It has been noted that NiO reacts with the LSGM electrolyte to form an insulating layer of LaNiO, at the interface between LSGM electrolytes and Ni-containing anodes.“ However, severe degradation owing to this reaction was not observed using the NiOjSDC composite particles obtained by SP. Micro- or nano-composition may suppress this reaction. Finally, the electrochemical performances of the SOFC single cells composed of the anodes derived from matrix-type composite particles and capsule-type particles were measured using an LSGM electrolyte with a thickness of 100 pm at 650”, 700”, and 750°C. Figure 9 shows SOFC power densities and cell voltages for both the single cells. Flake-free matrix-type composite particles were synthesized by SP with the chelated solutions at EG/ cation ratios of 5.0, whereas capsule-type composite particles were synthesized by SP without EG. Both the composite particles had the same NiOjSDC compositions. Matrix-type composite particles resulted in high-performance SOFC anodes as compared with capsule-type composite particles under these measurement conditions, and the maximum power densities of 1.69 W/cm2 at 750°C, 1.05 W/cm2 at 700”C, and 0.53W/cm2 at 650°C were found. Figure 10 shows backscattered SEM images of the anodes prepared with matrix-type and capsule-type composite particles. Partly collapsed composite particles in which some SDC small particles surrounded inner spherical Ni particles remained in the anode derived from the capsule-type composites, whereas nanometer-sized SDC or Ni particles were connected together in the anode derived from matrix-type composites. The matrix-type composite-derived anode showed poor porosity but preferable contact between SDC and Ni as compared with the capsule-type composite-derived anode. The capsule-type. composite anode would suppress Ni sintering as compared with the matrix-type one. However, the Ni sintering would be reduced by decreasing SOFC working temperatures. Matrix-type composite particles were then effective as precursors of high-performance IT-SOFC anodes.

IV. Conclusions Matrix-type NiOjSDC composite particles were synthesized by SP. An adequate addition of EG in the solution for SP converted capsule-type to matrix-type composition particles. The direct SP of the mixed solution containing EG resulted in NiO/SDC composite particles and flakes, whereas the pyrolysis of the solution chelated with EG by pre-heating resulted in solid particles with many dimples on the surface and relatively small particle size distributions. The solid matrix-type composite particles also showed a higher anode performance than the capsule-type composite particles. The matrix-type composite-derived anode showed poor porosity but preferable contact between SDC and Ni as compared with the capsule-type composite-derived anode. The matrix-type solid composite particles were effective as precursors of high-performance IT-SOFC anodes. References ‘V. V. Kharton, A. V. Kovalevsky, A. P. Viskup, A. L. Shaula. F. M. Figueiredo, E. N. Naumovich, and F. M. B. Marques, “Oxygen Transport in

Synthesis Methods for Powders

C Q , ~ G ~202-s-Based ,, Composite Membranes,” Solid Slate tonics. 160. 247-8 (2003). 2J. C. C. Abrantes, D. Perez-Coll, P. Nunez, and J. R. Frade, “Electronic y-6 Ceramics Under Reducing Conditions,” ElecfroTransport in Ce,,.8Sm,, 201 chem. Acfa, 48, 27614 (2003). ’F. Y. Wang. S. Chen, and S. Cheng, “Gd’+ and Sm3+Co-Doped Ceria Based Electrolytes for Intermediate Temperature Solid Oxide Fuel Cells,” Elecfrochem. Commun., 6. 7 4 3 4 (2004). 4V. Agdrwal and M. Liu, “Colloidal Processing of BaCeO,-Based Electrolyte Films,” J. Elecfrochem. Soc., 143, 323944 (1996). 5Z. Wu and M. Liu, “Stability of BaCe,,8Gdn203in a H20-Containing Atmosphere at Intermediate Temperatures,” J. Electrochem. Soc., 144, 217C-5 (1997). 6N. Sammes, G. Tompsett, Y. Zhang, A. Cartner, and R.Torrens, “The Structural and Mechanical Properties of (Ce02)l_,(GdO, .5)x Electrolytes.” Denki Kagaku, 64,674-80 (1996). 7Y. lkuma, “Nonstoichiometry and Diffusion in Ceria-Based Solid Solutions,” J. Soc. Inorg. Mafer. Jpn., 12, 213-20 (2005). ‘K. Higashi, K. Sonoda, H. Ono. S. Sameshima, and Y. Hirata. “Synthesis and Sintering of Rare-Earth-Doped Ceria Powder by the Oxalate Coprecipitation Method,” J. Maier. Res.. 14, 957-6 (1999). ’C. Milliken, S. Guruswamy, and A. Khandkar, “Evaluation of Ceria Electrolytes in Solid Oxide Fuel Cells Electric Power Generation,” J. Elecfrochem. Soc.. 146, 872-2 (1999). ‘OK. Yashira, S. Onuma, A. Kaimai, Y. Nigara, T. Kawdda, J. Mizusaki, K. Kawamura, T. Horita, and H. Yokokawa. ”Mass Transport Properties of Ce,,&do at the Surface and in the Bulk,” Solid Sfare tonics, 152-153, 469-76 (2002). ‘IT. lshihara, H. Minami, H. Matsuda, H. Nishiguchi, and Y. Takita, “Decreased Operating Temperature of Solid Oxide Fuel Cells (SOFCs) by the Application of LaGaO,-Based Oxide as Electrolyte,” Chem. Commun., 929-30 (1996). ”P. R. Slater, J. T. S. Irvine, T. Ishihara. and Y. Takita, “The Structure of the Oxide Ion Conductor L+,ySro lG+ 8M& 202.8s by Powder Neutron Diffraction.” Solid Stare tonics, 107, 319-23 (1998). ”I. Taniguchi, R. C. van Landshoot, and J. Schoonman, “Electrostatic Spray Deposition of Gdo.iC~yO1.g5and LaoySrg.lGa08Mgo202.87 Thin Films,” Solid Stare lonics, 160, 271-9 (2003). I4A. C. Tas, P. J. Majewski, and F. Aldinger, “Chemical Preparation of Pure and Strontium- and/or Magnesiumdoped Lanthanum Gallate Powder,” J . Am. (2000). Ceram. Sac., 83. 295”M. Feng, J. 9. Goodenough, K. Huang, and C. Milliken, “Fuel Cells With Do d Lanthanum Gallate Electrolyte,” J . Power Sources, 63,47-51 (19%). Huang, R. Tichy. J. 9. Goodenough, and C. Milliken, “Superior Perovskite Oxide-Ion Conductor; Strontium-and Magnesium-Doped LaGaO,: 111. Performance Tests of Single Ceramic Fuel Cells,” J. Am. Ceram. Soc., 81, 2581-5 (1998). ”Y. Matsuzaki, Y. Baba, and T. Sakurai, “High Electric Conversion Efficiency and Electrochemical Properties of Anode-supported SOFCs,” Solid Sfafetonics, 174, 81-6 (2004). “Y. J. Leng, S. H. Chan, K. A. Khor, and S. P. Jiang, “Performance Evaluation of Anode-supported Solid Oxide Fuel Cells with Thin Film YSZ Electrolyte.” Inf. J. Hydrogen-Energ., 29, 102S33 (2004). ”S. de Souza, S. J. Visco, and L. C. De Jonghe, “Thin-Film Solid Oxide Fuel Cell With High Performance at Low Temperature,” Solid State lonics, 98, 57-71 (1997). , Matsuzaki. Y. Baba. and T. Sakurai, “Diffusion Characteristics in AnodeSubstrates for AnodaSupported SOFCs,” Elecfrochemisfry,73, 484-8 (2005). ”S. Suda, M. Itagaki, E. Node, S. Takabashi, M. Kawano, H. Yoshida, and T. Inagaki, “Preparation of SOFC Anode Composites by Spray Pyrolysis,” J . Euro. Ceram. Soc., 26, 593-7 (2006). 22 M. Kawano, K. Hashino, H. Yoshida, H. Ijichi, S. Takahashi. S. Suda, and T. Inagaki, “Synthesis and Characterizations of Composite Particles for Solid Oxide Fuel Cell Anodes by Spray Pyrolysis and Intermediate Temperature Cell Performance,” J. Power Sources, 152, 1969 (2005). 23S.Suda, S. Takahashi, M. Kawano, H. Yoshida, and T. Inagaki, “Effects of Atomization Conditions on Morphology and SOFC Anode Performance of Spray Pyrolyzed N i 0 - S ~ , 2 C e o , 8 0 Composite 1,y Particles,” So/idSrafetonics, 177, 121% 25 (2006). 24R.C. West, M. J. Astle, and W. H. Beyer (eds.) C R C Hondhook of Chemisfry and Physics, pp. E-124, CRC Press Inc. Florida. USA. 1989. 25 H. Yahiro, Y. Eguchi, K. Eguchi, and H. Arai, “Oxygen Ion Conductivity of the Ceria-Samarium Oxide System with Fluorite Structure,” J . Appl. Elecfrmhem.. 18, 527-31 (1988). ’%. Zhao and R. J. Gorte. “A Comparison of Ceria and Sm-doped Ceria for Hydrocarbon Oxidation Reactions,” Appl. Cafal.A: General, 277, 129-36 (2004). 27R.Maric, S. Ohara, T. Fukui, H. Yoshida, M. Nishimura, T. Inagaki, and K. Miura, “Solid Oxide Fuel Cells with Doped Lanthanum Gallate Electrolyte and LaSrCoO, Cathode, and Ni-Samaria-Doped Ceria Cermet Anode.” J. Elecfro0 chem. Soc., 146. 200610 (1999).

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Novel Low-Temperature Synthesis of Ferroelectric Neodymium-Doped Bismuth Titanate Nanoparticles Prem Prakash and Ashish Gargt Department of Materials and Metallurgical Engineering, Indian Institute of Technology, Kanpur 208016, India

Mukesh Kumar Roy and Harish Chandra Verma Department of Physics, Indian Institute of Technology, Kanpur 208016, India

In this study, we report on the synthesis of nanopowders offerroelectric Bi3.5N&.5Ti3012 ceramic at temperatures below 500°C via a simple chemical method using citric acid as a solvent. The calcined powders were characterized using X-ray diffraction (XRD), differential scanning calorimetry (DSC), and transmission electron microscopy (TEM). Heating the as-dried powders in air first leads to crystallization of the BizTi207 phase at -310°C,followed by crystallization of the perovskite Nddoped BhTi3Ol2 phase at -490°C as suggested by the peaks in the DSC analysis and confirmed by the evolution of phases in XRD patterns of the powders calcined at various temperatures. TEM of particles calcined at 550°C for 1 h in air showed an average particle size of 50-60 nm. The temperature dependence of capacitance of nanopowders calcined at 700°C for 1 h in air showed a Curie temperature of 615°C evincing a ferroelectric transition.

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I.

Introduction

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ANTHANIDE-doped ferroelectric bismuth titanate (Bi4Ti3OL2) has been extensively studied from both the technological and fundamental points of view because of its potential applications in the non-volatile ferroelectric random access memory devices as well as sensors and actuators. Bi4Ti3012is one of the well-studied compounds of the Aurivillius family' and is represented by the general formula (Bi202)2+ [Am-lB,03m+ with m = 3 consisting of three perovskite-structured (Bi2Ti3010)2units sandwiched between (Bi202)2+ layers along the c axis of the Bi4Ti30L2unit cell, i.e., one layer is present on each side of the stack of three perovskite layers along the c axis. Upon thermal cycling, it undergoes reversible ferroelectric- araelectric phase transition at a Curie temperature of 675"C, a transition from a paraelectric tetragonal Z4/mmm structure, as proposed by Aurivillus,' to an orthorhombic structure.24 Doping with a lanthanide element such as La or Nd creates a significant amount of increment in the fatigue resistance of Bi4Ti3OI2,especially in the thin-film form.5 The current emphasis on nanotechnology has led to the drive toward the synthesis of ferroelectric materials in nanosized form. Nanosized ferroelectrics are further important to enhance the understanding of size effects in ferroelectrics. Conventionally, Bi4Ti3OI2ceramics have been prepared by solid-state reaction in a stoichiometric oxide mixture of Bi203 and Ti02, which

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Z. Bdrbe--contributing editor

Manuscript No. 22374. Received October 20, 2006; approved November 27, 2006. Supported by Department of Science and Technology, New Delhi (Grant no. SR/FTP/ ETA-032/2002). 'Author to whom correspondence should be addressed. e-mail: [email protected]

Synthesis Methods for Powders

are calcined at temperatures above 800°C to yield a fully crystallized perovskite Bi4Ti3OI2phase.637However, this leads to unwanted coarsening and agglomeration of the powders, which is undesirable in many applications. Efforts have been made to decrease the calcination temperatures, and chemical processes have proved to be successful to some extent. The powders of Bi4Ti301 have been synthesized by various chemical rocesses such as co-pre~ipitation,~.~ sol-gel,'03'1 hydr~thermal,'.'~ molten salt route,14 metalor anic decomposition,15 and mechanically assisted synthesis," but reports are limited in nature. Although some of these processes offer crystallization of the perovskite phase at temperatures below 800"C, nano-crystallinity is not achieved. On the other hand, hydrothermal synthesis leads to c stallized nanosized Bi4Ti3012at temperatures as low as 24OoC,' but the processing route remains complex, making the process unviable for large-scale production demanding high throughput at lower costs. Thus, low-temperature synthesis of Bi4Tt3OI2at a nanoscale using a simple process is important and remains a challenge. In this study, we report on the low-temperature synthesis of nanosized powders of well-crystallized ferroelectric Bi3,5Nd0.5Ti3012r using a chemical route. The powders were characterized using X-ray diffraction (XRD), differential scanning calorimetry and thermo-gravimetric analysis (DSCTGA), and transmission electron microscopy (TEM). Dielectric measurements were made to ascertain the ferroelectric nature of the particles.

r

11. Experimental Procedure (1) Synthesis Neodymium-doped BbTi3012 nanopowders (Bi3.5Nd0.5Ti3012 or BNdT) were synthesized by the wet chemistry route. Initially, a solution was made by dissolving the Bi and Nd precursors (hydrated bismuth and neodymium nitrates) in a citric acid solvent. Ethylene amine was used to stabilize the solution. Titanium isopropoxide was used as a titanium precursor and was stabilized using acetyl acetone before adding to the solvent containing Bi and Nd precursors. Extra 10% Bismuth was added to overcome the loss of Bi during heat treatment. A measured amount of sol (200 mL) was dried at 110°C while vigorously stirring using a hot place cum magnetic stirrer until it became viscous enough so that the sol constituents could not move freely. It was dried further in an ultrasonic bath at the same temperature. The dried powder was calcined at various temperatures, from 200" to 700°C for 1 h in air.

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(2) Characterization Analysis of the crystallinity, phase purity, and particle size of the calcined powders was carried out using a Seifert X-ray diffractometer using CuKu radiation (h = 1.540598 A). DSC and TGA were conducted using a METTLER TOLEDO DSC (822e;

85

Greifensee, Switzerland) equipment to investigate the crystallization of the Bi4Ti3OI2phase in as-synthesized powders at a heating rate of 10"C/min from room temperature to 700°C in air. The particle size and phase purity of the synthesized powders were further confirmed by using JEOL-TEM (Tokyo, Japan), operated at 200 kV. To prepare the samples for TEM investigation, calcined BNdT samples were ultrasonicated in isopropanol for 10 min. A few drops of the suspension were placed on the Formvar-coated 3 mm diameter and 200 mesh copper grid. The concentration of BNdT nanoparticles in the suspension was optimized to avoid agglomeration of the particles. The average particle size was also calculated from the bright-field (BF) images, and a selected area diffraction pattern (SADP) was taken to affirm the phase purity of the particles. The stoichiometry of the crystallized powders was examined using EDAX (energydispersive X-ray analysis) in a scanning electron microscope (FEI Quanta 200, Eindhoven, the Netherlands). Crystallized powders were cold pressed into pellets of 10 mm diameter. The faces of the green pellets were graphite coated to perform capacitance versus temperature measurements to determine ferroelectric transition using a Hewlett Packard 4192A (Melrose, MA) impedance analyzer at a frequency of 500 kHz. The sample was placed inside a horizontal tube furnace and was connected with silver wires on both the surfaces. A chromel/alumel thermocouple was placed beside the pellet to measure the sample temperature accurately. The temperature was controlled using a PID controller and a heating rate of l"C/min was maintained to attain the desired temperature uniformity.

111. Results and Discussion Figure 1 represents the XRD pattern of the BNdT samples calcined for 1 h in air at various temperatures: 300", 350", S O " , 600", and 700°C. The pattern shows that the samples calcined up to 300°C do not exhibit any significant phase formation. Crystallization is observed to begin at 350°C with the appearance of a non-ferroelectric Bi2Ti207phase7.17-19(JCPDS file number: 3118). However, upon heat treating the samples further up to 7 W C , Bi2Ti207was observed to transform into perovskite Nddoped B&Ti3012(JCPDS file number: 35-795a) as suggested by the gradual disappearance of Bi2Ti2O7peaks and enhancement of the peaks of the Nd-doped Bi4Ti3OI2phase. EDAX analysis of powders calcined at 550" and 700°C with the fully crystallized Nd-doped Bi4Ti3OI2phase showed a powder composition of Bi~.~+o.~Ndo.5+o.04Ti~O12, suggesting that the stoichiometry was maintained in all samples. Crystallite sizes of the heat-treated powders were calculated utilizing XRD peak-broadening data and using Scherrer's equation

Fig. 1. &20 patterns of Bi3sN& 5Ti3012nanopowders calcined at various temperatures.

heating and was evident by the broad peak between 100" and 250°C and significant weight loss as suggested by the TGA measurements (not shown here). In this figure, a strong exothermic peak (A) is observed at -31O"C, corresponding to the crystallization of the Bi2Ti207phase as suggested by the XRD results. Upon heating the sample further, another peak (B) is observed at -490°C, which can be ascribed to the formation of perovskite Bi4Ti3012phase, substantiated by the XRD results of powders heat treated above 500°C (Fig. I). Another peak (C) at 625°C could be indicative of a reversible ferroelectric-paraelectric phase transition and could be suggestive of the ferroelectric nature of the synthesized powders. This was further substantiated by the dielectric measurements of the powders as explained later. Previously, studies have been carried out to investigate the crystallization behavior of amorphous B&Ti3012.7*19 These studies also predicted the formation of perovskite Bi4Ti3012via a two-step process where Bi2Ti207was the first phase to form. It was reported that the formation of a metastable Bi2Ti207phase occurs at 580"C19and -608"C,7 followed by crystallization of a perovskite Bi4Ti3OI2phase at -85O"Cl9 and -831"C.7 Our results show the crystallization of perovskite Nd-doped Bi4. Ti3012 below 500"C, more than 300°C lower than the above values. To the best of our knowledge, this is also the lowest crystallization temperature reported for Bi4Ti3OI2 powders made using a simple chemical solution route. It is envisaged that this occurs due to the combined effect of vigorous stirring

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7

8-

where d is the crystallite size, k is the shape factor (-0.9), h is the wavelength of the radiation used, is the full-width at halfmaximum of the diffraction intensity of selected peak, and 8 is the Bragg angle. The instrumental broadening was calculated for standard samples and was subtracted from the measured broadening. Measurements were made using broadening of the (117) peak of Bi4Ti3OI2occurring at 28 = 30.157". The calculated particle sizes for samples calcined for I h at S O " , 600", and 700°C were approximately 18, 35, and 40 nm, respectively. An increase in the particle size upon heating up to 700°C is indicative of particle coarsening due to enhanced diffusion and a reduction in the overall surface energy. In order to confirm and analyze the phase evolution observed from XRD patterns, thermal analysis was performed on uncalcined samples. Figure 2 shows the DSC scan of as-synthesized powders upon heating up to 700°C in air. The removal of the volatile chemical constituents takes place in the earlier stages of 86

z Y

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1 100

200

300 400 500 Temperature "C

600

700

Fig. 2. Differential scanning calorimetry curve for as-prepared Bi,,,N~.sTi301zpowders (air atmosphere).

Progress in Nanotechnology:Processing

nanopowders calcined for 1 h in air at (a) 550"C, (b) W C , and Fig. 3. Bright-field transmission electron microscopy micrographs of Bi3.5N&.STi3012 (c) selected area diffraction pattern of the powders calcined at 600°C.

as well as ultrasonication of the solution during the drying stage, which helps to mix the different constituents of the mixture effectively and results in lower crystallization temperatures of various phases. In addition, vigorous stirring and ultrasonication during drying avoids the gel formation, which could be the key to the lowering of crystallization temperatures, eventually leading to nanosized particles. Formation of gel can lead to larger separation between the constituents, leading to higher crystallization temperatures. Figures 3(aHc) show the BF TEM micrographs of Bi4Ti3OI2 powders calcined at 550" and 600°C for 1 h in air and a selected area diffraction pattern of powders calcined at 600°C. It is evident from the micrograph (Figs. 3(a) and (b)) that the average particle size of the powders is between 50 and 60 nm, with the size ranging between -40 and 90 nm, which is reasonably uniform. A comparison of size measurements with those made using XRD peak widths suggests that some particles consist of more than one crystallite. The rings in the selected area diffraction obtained from powders calcined at 6OO0C, as shown in Fig. 3(c), were indexed and correspond to (012), (1 13), and (002) planes of Nd-doped Bi4Ti3OI2. Figure 4 shows the capacitance versus temperature plot for powders heat treated at 700°C for 1 h in air. It should be noted that we chose the nanopowders calcined at 700°C for the measurements to avoid any ambiguities or effects related to size growth that can occur due to heating the low-temperature-calSynthesis Methods for Powders

cined particles up to 800°C during electrical measurements. It is observed that the capacitance increases slowly until about 500"-550"C, and then starts to increase monotonically until the

0.4

G C

v

a, 0

a C

c

'U 0.2 m a

s

100

200

3M)

400

500

600

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Sample Temperature ("C) Fig. 4. Capacitancetemperature plot for neodymium-doped bismuth titanate nanopowders calcined at 700°C (1 h, air) and measured at a frequency of 500 kHz.

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transition from the ferroelectric to paraelectric phase at 615”C, followed by a sharp decline at the transition temperature. This temperature is comparable with the endothermic peak observed in DSC scan (Fig. 2), approximately at the same temperature. The transition process was found to be repeatable and reversible. This plot clearly shows that the nanopowders fabricated in the present study are ferroelectric in nature. Further studies are being carried out to thoroughly investigate the ferroelectric properties of the nanopowders, especially those calcined at lower temperatures.

IV. Conclusions Bi3.5Nd0.5Ti3012 nanopowders have been successfully synthesized at temperatures as low as -500°C via a simple chemical method using citric acid as a solvent. The crystallite sizes measured from XRD patterns were approximately 18,35, and 40 nm, for powders calcined for 1 h in air at 550”, 600”, and 700°C, respectively. The combination of XRD observations with DSC scans suggests a two-step transformation: formation of a nonferroelectric, non-perovskite Bi2Ti207 phase at 310”C, followed by formation of perovskite ferroelectric Nd-doped Bi4Ti3OI2 at -490°C. The sizes of well-crystallized powders calcined at 550°C (1 h in air) were -40-90 nm as shown by TEM, two to three times larger than those computed from Scherrer’s formula. The presence of a ferroelectric-paraelectric phase transition at -615°C is confirmed from the capacitance versus temperature measurements.

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Acknowledgments The assistance of Mr. S.C. Barthwal (IIT Kanpur) with TEM, and Mr P. Padaikathan, and Dr. U Ramamurthy (both from Department of Metallurgy, Indian Institute of Science, Bangalore) with the DSC measurements is gratefully acknowledged.

References ’B. Aurivillus, “Mixed Bismuth Oxides with Layer Lattices I 1 Structure of Bi4TiZOI2.”Ark. Kemi.,1 [58] 499 (1949).

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’J. F. Dorian, R. E. Newnham, and D. K. Smith, “Crystal Structure of Bi4Ti3OI2,”Ferroelectrics,3, 17 (1971). ’R. E. Newnham, R. W. Wolfe, and J. F. Dorian, “Structural Basis of Ferroelectricity in the Bismuth Titanate Family,” M a w . Res. Bull., 6, 1029 (1971). 4A. D. Rae, J. G. Thompson, R. L. Withers, and A. C. Willis, “Structure Refinement of Commensurately Modulated Bismuth Titanate, Bi4Ti3012.”Acra. Cr s f , 846,474 (1990). ‘B.’H. Park, B. S. Kang, S. D. Bu, T. W. Noh, J. Lee, and W. Jo. “LanthanumSubstituted Bismuth Titanate for Use in Non-Volatile Memories,” Nature, 401. 682 (1999). %. C. Subbarao, “Ferroelectricity in Bi4Ti3OI2and its Solid Solutions.” Phys. Rev., I22 131 804 (1961). 7S. Kojima, A. Hushur. F. Jiang. S. Hamazaki, T. Takashige, M. S. Jang, and S . Shimada, “Crystallization of Amorphous Bismuth Titanate.” J. Non-Crysr. Solids, 293-2295, 250 (2001). *Y. H. Wang, C. P. Huang, and Y. Y. Zhu, “Variable-Temperature Raman Scattering and X-Ray Diffraction Studies of BiZ2sNdo75Ti3012Ceramics,” Sol. State Comm., 139,229 (2006). ’M. Villegas, C. Moure, J. F. Fernandez, and P. Duran, “Low-Temperature Sintering of Submicronic Randomly Oriented Bi4Ti3012Materials,” Ceram. Int.. 22, 15 (1996). ’OX. Q. Chen, H. Y. Qi, Y. J. Qi, and C. J. Lu, “Ferroelectric and Dielectric Properties of Bismuth Neodymium Titanate Ceramics Prepared Using Sol-Gel Derived Fine Powders,” Phys. L p f f .A, 346 [I-31 204 (2005). “A. V. Prasada Rao, A. 1. Robin, and S. Komarneni, “Bismuth Titanate From Nanocomposite and SoCGel Processes,” M a w . Left.. 28, 469 (1996). ”D. Chen and X. Jiao, “Hydrothermal Synthesis and Characterization of Bi4Ti3012Powders From Different Precursors,” Mater. Re$. Bull., 36, 355 (2001). ”Y. Shi, C. Cao, and S. Feng, “Hydrothermal Synthesis and Characterization of Bi.,Ti3012,” Mater. L p f f . , 46, 270 (2000). I4Y. Kan, X. Jin, P. Wang, Y. Lia, Y.-8. Cheng, and D. Yan, “Anisotropic Grain Growth of Bi4Ti3OI2in Molten Salt Fluxes,” Mater. Res. Bull., 38, 567 (2003). ”W. L. Liu, H. R. Xia, H. Han, and X. Q. Wang, “Structural and Dielectrical Properties of Bismuth Titanate Nanoparticles Prepared by Metalorganic Decomposition Method.” J. Crysf. Growth, 269 [ 2 4 ] 499 (2004). I6B. D. Stojanovic, C. 0.Paiva-Santos, C. Jovalekic, A. 2.Simoes, F. M. Filho, Z. Lazarevic, and J. A. Varela. “Mechanically Activating Formation of Layered Structured Bismuth Titanate,” Mater. Chem. Phys., 96. 471 (2006). I7S. Shimada, K. Kodaria, and T. Matsushita, “Crystal Growth of Bismuth Titanates and Titanium Oxide From Melts in the System Bi203-V20rTi02.” J. Crysr. Growrh, 41, 317 (1977). ‘*W.-F. Su and Y.-T. Lu, “Synthesis, Phase Transformation and Dielectric Properties of Sol-gel Derived Bi2Ti207Ceramics,” Mater. Ckem. Phys., 80 [3] 632 (2003). I9Y. Yoneda, J. Mizuki, S. Kohara, S. Hamazaki, and M. Takashige, “Crystallization Process of Ferroelectric Bi4Ti3OI2From Amorphous State,” J. App. Pkys., 99,074108 (2006).

Progress in Nanotechnology:Processing

Hydrothermal Synthesis of CdMo04 Nano-Particles Xiaohui Jiang, Junfeng Ma,+,$Botao Lin, Yang Ren, Jun Liu, Xiaoyi Zhu, and Jiantao Tao Institute of Materials Science and Engineering, Ocean University of China, Qingdao 266003, China

Yonggang Wang and Lijin Xie College of Chemistry and Chemical Engineering, Ocean University of China, Qingdao 266003, China

CdMo04 nano-particles were successfully synthesized by a hydrothermal process at a low temperature of 9WC, and the powders were characterized in detail by X-ray diffraction, transmission electron microscopy, scanning electron microscopy, and photoluminescent spectra techniques (PL), respectively. CdMo04 particles could be obtained under the hydrothermal condition from micrometer to nanometer sizes by varying their precursors. The PL spectra results showed that the optical properties of CdMo04 crystallites obviously relied on their particle sizes.

1. Introduction

R

ECENTLY,

considerable attention has been paid to molybdates due to their electrochemistry properties,' special luminescence, and unique structure^.^-^ Metal molybdates gf relatively large bivalent cations (MMo04, ionic radius >0.99 A, M = Ca, Ba, Sr, Pb) are of a so-called scheelite structure in which the molybdenum atom adopts tetrahedral coordination,6 where the emission spectrum of the metal molybdates is mainly attributed to the charge-transfer transitions within the [Moog-] ~ o m p l e x . ~Calcium - ~ ~ molybdate (CaMo04) as an important functional material has been used in various fields, such as in photoluminescence (PL)" and microwave applications,12 and lead molybdate (PbMo04) single crystals are extensively used in acousto-optical and high-voltage measurement devices.l 3 Ternary oxides with the general formula MMo04 (M = Cu, Zn, Ni, Fe) are also characterized as cathode materials for rechargeable batteries,' the idea being to take advantage of the charge couple Mo6+/Mo4+where in the metal-redox oxidation state change by two units is very attractive for the development of batteries with a high capacity and also with high-energy density.I4-l8Cobalt molybdate (CoMo04) as another attractive compound has exhibited a wide range of applications in industrial catalysis for properties related to its s t r u ~ t u r e . ' ~ * ~ ~ However, there have been few reports on CdMo04, whether in its synthesis process or properties, to the best our knowled e Even though there possibly exist environmental drawbacks2 22 of cadmium compounds such as other heavy metal compounds as Lead Zirconate titanate (PZT), it is necessary to study CdMo04 as a potentially functional material with relatively low toxic Cd(II).23 Therefore, it is of significance both in fundamental and applied fields to synthesize and characterize

F.

G . Pdtzke--contrihuting editor

Manuscript No. 21820. Received May 24, 2006; approved November 3, 2006. 'Author to whom correspondenceshould he addressed. e-mail: [email protected] 'COkge of Chemistry and Chemical Engineering,Ocean University of China, Qingdao, China.

Synthesis Methods for Powders

CdMo04 powders with different particle sizes, especially of nanometer scale. It is well known that molybdate powders can be prepared by many methods, e.g. the pulsed laser ablation process for synthesizi;g C ~ M O O the ~ , citrate ~ ~ complex method for BaMoo4, and the spray pyrolysis route to N ~ M o O ~ . ~ ~ Nevertheless, among these routes, there are still such limitations as needing a higher calcined temperature and a longer holding time to eliminate organic material or requiring complicated instrumentations and manipulations. The hydrothermal method has been considered to be a good synthesis process for some inorganic powders27 because of (1) available synthesis of crystallized products at a low reaction temperature, (2) flexibility in the design of reaction conditions, (3) uniformity of production composition, phase, and microstructure, and (4) simplicity of equipment and processing. Here, we report on the hydrothermal synthesis of CdMo04 crystallites. The particle size of CdMoO? powders was about 3-5 pm by a direct hydrothermal route, while the CdMo04 crystallites could be reduced to 30-50 nm in particle size under the same hydrothermal condition by introducing a microemulsion route to the precursor of CdMo04. The PL properties of CdMo04 crystallites with different particle sizes were also studied for the first time. 11. Experimental Procedure

(1) Synthesis of CdMoOl Precursors Sodium molybdate (Na2Mo04 2H20) and cadmium nitrate (Cd(N03)2 4H20) were used as the starting materials, and both of them were of analytical grade without any further purification. In our work, two methods were used to prepare the precursor of CdMo04 for the subsequent hydrothermal treatment. ( A ) Normal Precursor Preparation: Appropriate amounts of Na2Mo04 2H20 and Cd(N03)2 4H20 were dissolved in distilled water of 250 mL to form 1M aqueous solutions, respectively. Then, both solutions were mixed together with vigorous magnetic stirring at room temperature to form a suspension, which was called the normal precursor. ( B ) Microemulsion Precursor Preparation: Cyclohexane was used as the oil phase, OP was used as the surfactant, and nbutanol as the co-surfactant. A typical process was as follows: two kinds of microemulsions (M I and M 11) with different aqueous phases were obtained (see Table I). The aqueous phase in the M I was a solution of Cd(N03)*,while the aqueous phase in the M I1 contained Na2Mo04. The concentration of both of the aqueous phases was 0.1M . These two microemulsions were mixed together and stirred for 30 min at room temperature. The appearance of a white suspension indicated the formation of microemulsion precursor of CdMo04. The precipitate could be separated by a centrifuge, and washed with acetone and absolute

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89

Table 1.

Aqueous phase Surfactant Co-surfactant Oil ohase

Ingredients of the Microemulsions

0.1 M Cd(N0& OP n-butanol Cvclohexane

0.1 M NazMoO4 OP n-butanol Cvclohexane

8 15 10

I0

ethanol several times to eliminate excess OP. The final product was called the microemulsion precursor.

( 2 ) Synthesis of CdMoO, Crystallites The normal precursor of 500 mL was poured into a 1000 mL stainless-steel autoclave, and hydrothermally treated at 90°C for 20 h. Similarly, the above-obtained microemulsion precursor was also placed in the same autoclave with distilled water, filling up to 500 mL of capacity, and then hydrothermally treated at 90°C for 20 h. The two kinds of CdMo04 powders could be obtained by filtrating, washing, and drying after the hydrothermal process. (3) Characterization XRD analysis was carried out using an X-ray powder diffractometer (XRD, D/max, Rigaku, Tokyo, Japan) with CUKEradiation. The morphology and particle size of the as-prepared powders were observed by using a transmission electron microscope (TEM, JEM-l200EX, JEOL, Tokyo, Japan), and scanning electron microscope (SEM, JSM-840, JEOL). The roomtemperature luminescent spectra were recorded on a spectrofluorometer (PL, Fluorolog-3, Jobin Yvon, Edison, FL). 111. Results and Discussion

( I ) XRD Analysis Figure 1 shows the XRD patterns of the samples prepared at 90°C for 20 h by the hydrothermal method with the normal precursor, and microemulsion precursor, respectively. It indicates that the good crystallization of tetragonal phase CdMo04 can easily be achieved under the hydrothermal condition with the different precursors. Both the XRD patterns are well consistent with the reported data (JCPDS: 88-0182), and no other impurities can be found. On the other hand, the XRD pattern obtained when using the normal precursor is much stronger in intensity than of using the microemulsion precursor. This means that the crystallization and development of the CdMo04 phase

Fig. 2. Transmission electron microscope (TEM) and scanning electron microscope (SEM) images of the CdMo04 crystallites prepared at 90°C for 20 h by the hydrothermal process: (a) SEM, the product from the normal precursor, and (b) TEM, the product from the microemulsion precursor.

I 350

400

450

500

550

600

wavelength (nrn)

Fig. 1. X-ray diffraction patterns of the samples obtained by the hydrothermal process at 90°C for 20 h: (a) with the normal precursor, and (b) with the microemulsion precursor, respectively.

90

Fig. 3. Photoluminescence spectra of the as-prepared CdMo04 crystallites by the hydrothermal process with (a) the normal precursor, and (b) the microemulsion precursor, respectively.

Progress in Nanotechnology: Processing

seem to be sensitive to the precursor used in the hydrothermal process.

(2) Particle Size and Morphology of CdMoOd Crystallites SEM and TEM images of the CdMo04 powders obtained from the different precursors are shown in Fig. 2. The normal precursor resulted in a perfectly and homogeneously spherical crystallization of CdMo04 phase about 3-5 pm in particle size (Fig. 2(a)), while the nanometer scale of CdMo04 powders, which is much smaller than the former, could be prepared by using the microemulsion precursor, and the resultant particles were estimated at about 3&50 nm in size and of nearly homogeneous and polygonal morphology (Fig. 2(b)). Obviously, the particle size and morphology of CdMo04 crystallites strongly rely on the precursor species even though both the products have the same crystalline phase of CdMo04. Here, nano-size CdMo04 crystallites, prepared with the microemulsion precursor, should grow in a controlled manner in much smaller regions, determined by the microemulsion process, whereas the normal precursor, obtained from direct precipitation, would lead to an extensive growth of CdMo04 crystallites to larger particle sizes (Fig. 2(a)) in an uncontrolled manner. It is reasonable that the particle size and morphology of the tetragonal CdMo04 crystallites can be tailored under the hydrothermal condition by varying the precursor used.

( 3 ) PL Analysis Figure 3 presents the PL spectra of the CdMo04 samples corresponding to Fig. 2. Both the samples exhibit the same emission peak position at 390 nm in the PL spectra. Nevertheless, it is noteworthy that the emission peak of 390 nm could be obtained by a longer excitation wavelength (330 nm) in the case of nanoparticle CdMo04 than that used in the micrometer-sized ones (320 nm). Furthermore, as shown in Fig. 3, the nano-particle CdMo04 has a higher PL intensity than that of the micrometersized powder. This shows that the CdMo04 nano-particles should have better PL properties than the micrometer-sized ones. IV.

Conclusions

CdMo04 nano-particles with a good crystallization could be synthesized at 90°C for 20 h by a hydrothermal method, where a microemulsion precursor was used. But under the same hydrothermal condition, only the normal precursor afford led to micrometer-sized CdMo04 crystallites. The particle size and morphology of CdMo04 crystallites can be tailored by varying the precursor species. The improved PL properties of CdMo04 crystallites can be obtained with by decreasing the particle size to a nanometer scale.

References IN. N. Leyzerovich. K. G. Bramnik, T. Buhrmester, H. Ehrenberg, and H. Fuess, “Electrochemical Intercalation of Lithium in Ternary Metal Molybdates MMo04 (M: Cu, Zn. Ni and Fe),” J. Power Sources, 127, 76-84 (2004). ’V. B. Mikhailika, H. Krausa, D. Wahla, H. Ehrenbergb, and M. S. Mykhaylyk, “Optical and Luminescence Studies of ZnMo04 Using Vacuum Ultraviolet Synchrotron Radiation,” Nuclear Instrum. Methods Phys. Res. A , 562, 5 1 3 4 (2006).

Synthesis Methods for Powders

’V. B. Mikhailik, H. Kraus, D. Wahl, and M. S. Mykhaylyk. “Studies of Electronic Excitations in MgMo04, CaMo04 and CdMo04 Crystals Using VUV Synchrotron Radiation,” Phys. Star. Sol. ( h i , 242, R17-9 (ZOOS). 4D. Spassky, S. Ivanov, 1. Kitaeva, V. Kolobanov, V. Mikhailin, L. Ivleva, and I. Voronina, “Optical and Luminescent Properties of a Series of Molybddte Single Crystals of Scheelite Crystal Structure,” Phys. Srar. Sol. ( c ) . 2. 65-9 (ZOOS). ’V. B. Mikhailik, H. Kraus, M. Itoh, D. Iri, and M. Uchida, “Radiative Decay of Self-Trapped Excitons in CaMo04 and MgMo04 Crystals.” J. Phys. Condens. Matrer., 17, 7209-18 (2005). 6A. Karipides and D. A. Hailer, “The Crystal Structure of Tetraphenylgermanium,” Acfa CrystaNogr. B, 28, 288S92 (1972). ’R. Grasser, E. Pitt. A. Scharmann, and G. Zimmerer, “Optical Properties of CaW04 and CaMo04 Crystals in the 4 to 25 eV Region,” Phys. Srar. Sol. ( b ) , 69, 35948 (1975). ‘S. B. Mikhrin, A. N. Mishin, A. S. Potapov, P. A. Rodnyi. and A. S. Voloshinovskii, Nucl. Insrrum. Merh. A.. 486,295-7 (2002). 9D. A. Spassky, S. N. Ivanov, V. N. Kolobanov, V. V. Mikhailin. V. N. Zemskov, B. 1. Zadneprovski, and L. 1. Potkin, “Optical and Luminescent Properties of the Lead and Barium Molybdates,” Rad. Meas., 38, 607-10 (2004). “B. K. Chandrasekhar and W. B. White, “Luminescence of Single Crystal CaMo04,” Marer. Res. Bull., 25, 1 5 1 H (1990). “R. Grasser, E. Pitt, A. Scharmann, and G . Zimmerer, “Optical Properties of CaW04 and CaMo04 Crystals in the 4 to 25 eV Region,” Phys. Star. Sol. ( h ) , 69, 35758 (1975). L. F. Johnson, G. D. Boyd, K. Nassau, and R. R. Soden, “Continuous O p eration of a Solid-state Optical Maser,” Phys. Rev., 126, 14069 (1962). ”A. N. Belsky, V. V. Mikhailin, A. N. Vasil’ev, I. Dafinei. P. Lecoq, C. Pedrini, P. Chevallier, P. Dhez, and P. Martin, “Fast Luminescence of Undoped PbW04 Cr stal,” Chem. Phys. Lett., 243, 552-8 (1995). ‘P. A. Christian, J. N. Carides, F. J. DiSalvo, and J . V. Waszczak. J. Elecfrochem. Soc.. 127 [Ill2315 (1980). I5J. 0. Besenhard, J. Heydecke, E. Wudy, H. P. Fritz, and W. Foag, “Characteristics of Molybdenum Oxide and Chromium Oxide Cathodes in Primary and Secondary Organic Electrolyte Lithium Batteries. Part II. Transport Properties,” Solid Srare Ionics, 8. 61-71 (1983). ‘q.Tsumura and M. Inagaki, “Lithium Insertion/Extraction Reaction on Cr stalline MOO,,” Solid State Ionics, 104, 183-9 (1997). y7A.Yu. N. Kumagaj, Z. Liu, and I. Lee, “Preparation of Sodium Molybdenum Oxides by a Solution Technique and Their Electrochemical Performance in Lithium Intercalation,” Solid Srare lonics, 106, 11-18 (1998). “R. H. Sanchez, L. Trevino, A. F. Fuentes, A. Martinez-de la Cruz, and L. M. Torres-Martinez, “Electrochemical Lithium Insertion in Two Polymorphs of a Reduced Molybdenum Oxide,” J . Solid Srare Elecfrochem., 4, 21CL-5 (2000). I9P. K. Pandey, N. S. Bhave, and R. B. Kharat, “Preparation and Characterization of Spray Deposited NiMoO, Thin Films for Photovoltaic Electrochemical Studies,” Mater. Res. Bull., 41, 1160-9 (2006). 2oJ. L. Brito and A. L. Barhosa, “Effect of Phase Composition of the Oxidic Precursor on the HDS Activity of the Sulfided Molybdates of Fe(ll). Co(11). and Ni(ll),” J. Catal., 171, 467-75 (1997). ’IN. Serpone, P. Marathamuthu, P. Pichat, E. Pelizzetti. and H. Hidaka, “Exploiting the Interparticle Electron Transfer Process in the Photocatalysed Oxidation of Phenol, 2-Chlorophenol and Pentachlorophenol: Chemical Evidence for Electron and Hole Transfer Between Coupled Semiconductors,” J. Phororhem. Phorobiol., 85, 247-55 (1995). 22 H. Fujii. M. Ohtaki, and K. Eguchi, “Preparation and Photocatalytic Activities of a Semiconductor Composite of CdS Embedded in a TiO’ Gel as a Stable Oxide Semiconducting Matrix,” J. Mu/. Catal. A , , 129, 61-68 (1998). 23A. E. Raevskaya, A. L. Stroyuk, and S. Y. Kuchmiy, “Preparation of Colloidal CdSe and CdSjCdSe Nanoparticles from Sodium Selenosulfate in Aqueous Polymers Solutions,” J. Colloid Inrerfuce Sci. 302, 13341 (2006). 24 I. H. Ryu, B. G. Choi, J.-W. Yoon, K. B. Shim, K. Machi. and K. Hamada, “Synthesis of CaMo04 Nanoparticles by Pulsed Laser Ablationin Deionized Water and Optical Properties,” J. Lumin., 124, 67-70 (2006). 25J. H. Ryu, J.-W. Yoon, C. S. Lim, and K. B. Shim, “Microwave-Assisted Synthesis of Barium Molybdate by a Citrate Complex Method and Oriented A gregation,” Marer. Res. Ball., 40, 146&76 (2005). k . Mazzocchia, C. Aboumrad, C. Diagne, E. Tempesti, J. M. Herrmann. and G. Thomas. “On the NiMo04 Oxidative Dehydrogenation of Propane tn Propene: Some Physical Correlation with the Catalytic Activity,” Caral. Lefr., 10, 181-91 I).

X. Huang, J. Ma, P. Wu, Y. Hu, J. Dai, Z. Zhu, H. Chen, and H. Wang,

“Hydrothermal Synthesis of LiCoPO4 Cathode Materials for Rechargeable Lithium Ion Batteries,” Marer. Lert., 59, 578-82. 0

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Chromium-Doped Forsterite Nanoparticle Synthesis by Flame Spray Pyrolysis Takao Tanit and Shu Saeki Inorganic Materials Laboratory, Toyota Central R&D Laboratories Inc., Aichi 480-1 192, Japan

Takenobu Suzuki and Yasutake Ohishi Department of Future Industry-Oriented Basic Science and Materials, Toyota Technological Institute, 468-851 1 Nagoya, Japan

Synthesis of chromium-doped forsterite (Mg2Si04:Cr) nanoparticles by flame spray pyrolysis (FSP) was investigated. The morphology, crystalline phase, and photoluminescence of the products were evaluated. Crystalline Mg2Si04:Cr nanoparticles of several 10 nm in diameter were obtained, although a small amount of the submicrometer-sized particles and the unreacted MgO phase existed. The roduct powder showed electron-spin resonance signals from C& and photoluminescence typical for Cr4' in Mg2SiO4, sufgesting that a part of the Cr4' ions were incorporated into Si4 sites by FSP. On the other hand, the effects of excess Si02addition on the structural and optical characteristics of Mg2Si04:Cr were examined. Addition of excess Si02up to 20 mol% did not influence these characteristics of the products. Further addition of excess Si02 (60-100 mol%) enhanced the formation of the amorphous phase and resulted in the emission from Cr3+ in the amorphous phase in addition to an emission from Cr" in Mg2Si04.

Flame spray pyrolysis (FSP)'- is an elegant process to prepare metal oxide nanoparticles in one step by spraying and combusting precursor solutions where metal salts are dissolved in organic solvents. Product particles are formed by evaporation, condensation, and oxidation of the metal species (e.g., atoms or oxide clusters) in the gas phase. High-flame temperatures and short residence time in the hot zone by FSP can result in highly crystalline nanoparticles, which are appropriate for dispersants of the nanoparticle-polymer (or glass) composites. However, it is not clear whether the short processing time is enough for crystallization of Mg2Si04and/or incorporation of cr4+ into si4+ sites or not. This paper investigates the potential to apply the FSP process to Mg2Si04:Cr nanoparticle synthesis. The structural and optical characteristics of the products are evaluated. In addition, the effects of excess Si02 addition on these properties are examined because it may control Mg2Si04:Cr crystal size as seen in ZnO/SiO2 synthesis by FSP" where the ZnO crystal size was varied by addition of SiOz.

I. Introduction

F

ocus is placed on new-type optical waveguides where functional nanoparticles are dispersed in glass or polymer. Yoshida et al.' prepared Ti02 nanoparticles by a reversed micelle method and dispersed them in polyimide to control the refractive index of the polyimide waveguide. A waveguide where semiconductor nano-crystals are dispersed in quartz glass is proposed for an optical signal amplifier with controllable amplification wavelength by selecting the nano-crystals.2 Chromium-doped forsterite (Mg2Si04:Cr) is a lasting material whose activity was first reported in 1988,3 and Mg2Si04:Cr single crystals are applied commercially to tunable near-infrared (IR) lasers, whereas, if Mg2Si04:Cr nanoparticle-polymer (or glass) composites are obtained, they may be advantageous in terms of flexibility of shape and a low production cost compared with single crystals and can be applied widely to laser, amplifier, and other optical devices. Sol-gel synthesis is a well-known route for preparation of Mg2Si04:Crnanoparticles.4-6Although the sol-gel process seems appropriate to synthesize homogeneous nanoparticles, it primarily needs expensive metal alkoxide precursors and complicated steps for synthesis, increasing a process cost. In addition, its low process temperature can result in low crystallinity and thus low optical properties of the product. J. Ballat-ntrihuting

editor

Manuscnpt No. 21560. Received March 7, 2006; approved October 30. 2006. This work is supported in part by MEXT, the Private University High-Tech Research Center Program (2lX&20IO) and JSPS.KAKENH1 (17560322). 'Author to whom correspondence should be addressed. e-mail: tdnitakd@mosk. tytlahs.co.jp

Synthesis Methods for Powders

11. Experimental Procedure

A stoichiometric composition of MgzSi04:Cr is Mg2(Sio.w. Cro.01)04in this study. Synthesis was carried out for five compositions with various amounts of excess Si02 addition (X= 0, 0.1, 0.2, 0.6, and 1.0 in molar ratio). All powders are labeled by the amounts of excess Si02 addition ( X ) . Here, X = 0.1 corresponds to the composition of Mg2(Si0.99Cro,ol)04+0. 1 SOz. Magnesium acetate tetrahydrate (Wako, S grade), hexamethyldisiloxane (HMDSO, Wako, S grade) and chromium acetate (Wako, chemical grade) were used as Mg, Si, and Cr sources, respectively. These precursors were dissolved in a mixture of methanol (Wako, S grade) and I-butanol (Wako, S grade) with a volume ratio of methanol/butanol = 9/1, resulting in precursor solutions of 1.O mol/L in total metal ion concentration. Synthesis was carried out using a spray flame reactor." A pump supplied the precursor solution at a feed rate of 12 cm'/ min to the nozzle where the solution was dispersed into droplets by oxygen at 11 L/min. The droplets were ignited by supporting methane/oxygen flames surrounding the nozzle. The total flow rates of methane and oxygen were 2 L/min each. Air at 50 L/min was supplied through a porous metal plate ring for excess oxidant. The product particles were collected on a Gore-Tex filter with the aid of a vacuum pump. The particle morphology was observed by transmission electron microscopy (TEM; Nihon Denshi, JEM2000EX, 200 kV, Akishima, Japan). The compositional analysis was carried out by energy-dispersive X-ray spectroscopy (EDS; Hitachi, S3600N, Tokyo, Japan). The specific surface area (SSA) was measured by nitrogen adsorption (Micro Data, 4232). The crystalline phase of the product was identified by X-ray diffraction (XRD: Rigaku, RINT-TTR, 50 kV, 300 mA, Tokyo, Japan) at 93

0 Forsterite Periciase

20

30

40

50

60

70

2 Wdeg. (CuKa) Fig. 3. X-ray diffraction spectrum of the stoichiometric Mg2Si04:Cr powder. Forsterite (Mg2Si04)and periclase (MgO) as well as an amorphous phase (a broad peak at 20 = 20'-40") are observed. Fig. 1. Transmission electron microscopic image of the stoichiometric Mg2SiO4:Cr powder. Polyhedral aggregates of nanoparticles and a small amount of submicrometer-sized spherical particles are observed.

20 (CuKa) = 20"-70" and scan speed = Z"/min. The bonding state of metal ions was evaluated by IR spectroscopy (Nicolet, Avantar 360, Waltham, MA). The state of Cr ions was examined by electron spin resonance (ESR: Bruker, ESP300E, 100 mW, Tsukuba, Japan). The measurement was carried out at 9.6 GHz (X-band) in the field range of &I000 G with modulation of 100 kHz at 25°C. The near-IR photoluminescence (PL) spectrum of the product was measured at I-nm intervals in the Wavelength range of 820-1700 nm. The sample was excited by an 800-nm laser beam which was generated from a CW Tisapphire laser (Coherent, 890, Santa Clara, CA) and chopped mechanically at a frequency of 97 Hz. The emission from the sample was dispersed by a single monochromator (blaze wavelength: 1.0 pm, grating: 600 grooves/mm, resolution: 3 nm) and then detected by an InGaAs photo detector (Hamamatsu Photonics, H58S2-11, Hamamatsu, Japan) and a lock-in amplifier (NF, L15640,Yokohama, Japan).

In. Results and Discussion (I)

Stoichiometric Mg2SiO&r Synthesis ( A ) Morphology: Figure 1 shows a TEM image of the stoichiometric Mg2Si04:Cr powder (A' = 0). Polyhedral aggregates of nanoparticles, observed typically in flame-made powders,13 were obtained, while a small amount of submicro-

0

0.5

1 EnergylkeV

1.5

2

Fig.2. Energy-dispersive X-ray spectra of (a) the aggregates of the nanoparticles and (b) the large particles in the stoichiometric Mg2Si04:Crpowder. The MgjSi ratio is similar for both particles, indicating no segregation of the Zn and Si species.

94

meter-sized spherical particles (large particles) also existed. The nanoparticle diameter observed by TEM was about 20-50 nm, which was roughly consistent with the BET-equivalent average primary particle diameter (dBET:32 nm) estimated from the measured SSA (58 m*/g) and solid density (MgzSi04: 3.21 x lo' kg/m3)14assuming solid spherical Mg2SiO4particles. Such inhomogeneous morphologies have been obtained in the case of ZnO/Si02 synthesis by FSP usin zinc acetate and Si02 sol as Zn and Si precursors, respectiveIy.e2 In that case, the Si precursor did not evaporate in the flame and the particle was formed in the li uid phase in the dispersed droplets, as seen in spray pyrolysis?' resulting in gas phase-made aggregates of ZnO nanoparticles and liquid phase-made large Si02 particles, whose composition was confirmed by EDS. On the other hand, the EDS spectra (spot size = about 100 nm) of the aggregates of the nanoparticles (Fig. 2(a)) and the large particles (Fig. 2(b)) showed that the Mg/Si ratio was similar for both particles in this study, indicating no segregation of the Zn and Si species. Different from the previous study, the large particles may be formed irregularly because of the inhomogeneous thermal history of the aerosol in the spray flame. ( B ) Crystalline Structure: Figure 3 shows an XRD pattern of X=O. The main product was forsterite (Mg2Si04: PDF34-0189), suggesting that crystallization of Mg2Si04 primarily takes place during FSP. On the other hand, periclase (MgO: PDF4S-0946) and an amorphous (a broad peak at

12uo

1100

1000

900

800

700

Wavenumberlcm-'

Fig.4. Infrared (IR) spectrum of the stoichiometric Mg2Si04:Cr powder. The IR spectra of Mg2Si04, MgO, and SiO2 are shown also for comparison. The absorptions at 870 and 1000 cm-' typical for Mg2Si04 are observed. Progress in Nanotechnology: Processing

800

900

lo00 1100 1200 1300 1400 1500 Wavelengthlnm

Fig. 5. Normalized photoluminescence spectrum of the stoichiometric Mg2Si04:Cr powder. Emission at 1200 nm typical for Mg2Si04:Crare observed, suggesting that Cr ions substitute not Mg2+ but Si4+ sites in Mg2Si04as Cr4'.

2e=20°40") phase were also observed. In the FSP process, MgO and Si02nuclei can be formed independently from the gas phase during cooling because of different boiling points (MgO: 3600"C, Si02: 295O0C)," followed by crystallization of MgO with Si02 forming MgzSi04. A part of MgO and Si02 nuclei may be cooled down before crystallization of Mg2Si04because of the inhomogeneous thermal history of the aerosol. Figure 4 shows the IR spectrum of X = 0. The IR spectra of Mg2Si04, MgO, and Si02 are also shown in the figure for comparison. Absorptions at 870 and 1000 cm-l typical for Mg2Si04 (SiLO-Mg bonding)I6 were observed, whereas an absorption at 1050 cm-l typical for S O 2 (Si-0-Si bonding)16 was not seen, indicating that the Si4' and Mg2+ ions are linked or influenced by each other. The result suggests that the amorphous phase (Fig. 3: XRD) is not Si02 but magnesium silicates. (C) Emission: Figure 5 shows the normalized PL spectra of X = 0. Emission at 1200 nm typical for Mg2Si04:Cr4was observed, while there were no emissions at low wavelengths (e.g., 70&1100 nm), which have been seen for Cr3+ substituted with Mg2+ in Mg2Si02." or MgSi03.1s The result suggests that a part of the Cr ions substitute not Mg2' but Si4+ sites in Mg2Si04 as Cr4' by FSP. The ESR spectrum of X = 0 showed an ESR signal at 800 G (Fig. 6) although it was very weak and noisy. Budil et al." studied ESR of a Mg2Si04:Cr single crystal and reported that Cr4' showed broad ESR signals at 600 or 800 G, whereas Cr3+ showed sharp ESR signals at 15W3200 G for X-band measurement at room temperature. Although it is difficult to compare directly the ESR spectra of the single crystal with those of the powder sample, the ESR signal in this study

Fig. 7. Transmission electron microscopic images of the product powders with various amounts of excess SiOz addition (X) to [email protected])04.Addition of excess SiOz primarily does not influence the particle morphology.

may have originated from Cr4' in Mg2Si04particles because of its similar position and shape to those of Cr4' in the single crystal, su gesting that a part of the Cr4' ions were incorporated into Si". sites by FSP, qualitatively consistent with the PL result.

( 2 ) Addition of Excess Silica Figure 7 shows TEM images for X = 0.1, 0.2, 0.6, and 1.0. As seen for X = 0 (Fig. I), polyhedral aggregates of nanoparticles

x=o.1

0 Forsterite 0 Periclase

0

. : x .-I 111

X=0.6

A

20

0

200

400

600

800

1000

Magnetic Field (G) Fig. 6. The electron spin resonance (ESR) spectrum of X = 0. The ESR signal was seen at 800 G although it was very weak and noisy.

Synthesis Methods for Powders

30

40 50 2 Wdeg. (CuKa)

x=1.o

-

60

L

70

Fig.8. X-ray diffraction spectra of the product powders with various amounts of excess SiOz addition (X)to Mg2(Si0.99Cr~ 0 1 ) 0 4 . Compared with that for the stoichiometric MgzSi04:Cr powder (Fig. 3), the peak heights from Forsterite (Mg2SiO4)and periclase (MgO) lower whereas the broad peak is intensified as the amount of excess SiOz increases. While, no silicon-rich crystalline phases such as enstatite (MgSi03) are seen.

95

sized particles as well as an unreacted MgO phase existed. The product powder showed electron spin resonance signals from Cr4’ and PL typical for Cr4’ in Mg2Si04,suggesting that a part of the Cr4’ ions were incorporated into Si4+sites by FSP. On the other hand, addition of excess Si02 up to 20 mol% (X=O.l and 0.2) did not influence the structural and optical characteristics of the products. On the other hand, further addition of excess Si02 ( X = 0.6 and 1.0) enhanced a formation of the amorphous phase and resulted in emission from Cr3+ in the amorphous phase in addition to emission from Cr4+ in Mg2Si04,although the particle morphology did not change so much. The crystal size of Mg2Si04 was not controlled by addition of excess Si02. Future study will be focused on evaluation and improvement of emission efficiency by optimizing the Cr-doping amount for X = 0. Fig. 9. Normalized photoluminescence spectra of the product powders with various amounts of excess SiOzaddition ( X ) to Mg2(Si0.9&r0.01)04. For X=O.1 and 0.2, the emission spectra are similar to that for the stoichiometric Mg2Si04:Crpowder (Fig. 5), whereas further addition of excess Si02 (X= 0.6 and 1.O) results the emission from Cr3+ (broad emission at 850-1 100 nm) in addition to the emission from Cr4’.

and a small amount of submicrometer-sized spherical particles were also observed for all powders. In FSP, particle morphology is influenced mainly by the kinds of precursors and solvents used as well as the combustion conditions (e.g., O2 flow rate), and these are the same for all powder syntheses in this study, which could be the reason why addition of excess SiOz did not influence the particle morphology. Figure 8 shows XRD patterns for X = 0. I, 0.2, 0.6, and 1.0. Compared with that for X = 0, the peak heights from Mg2Si04 and MgO decreased, whereas the broad peak was intensified as the amount of excess SiOz increased. On the other hand, no silicon-rich crystalline phases such as enstatite (MgSi03: PDF190768) were seen for all powders. The phase diagram2” shows that the melting point of Mg2Si04 (1888°C) is higher than that of MgSi03 (1 557°C) and that the phase changes “liquid -+ Mg2Si04+liquid -+ Mg2Si04+MgSi03” during cooling in the composition range of Mg/Si = 2/1 to 1/1, indicating that the Mg2Si04formation is preferred during cooling. A rapid cooling rate of the FSP process can result the formation of the “Mg2Si04+amorphous” phase from the “Mg2Si04+liquid” phase without crystallization of MgSi03 regardless of the amounts of excess SO2. The IR spectra showed Si-O-Mg absorptions at 870 and 1000 cm-l as seen for X = 0, although the peak heights decreased as the excess SiOz increased, suggesting that the amorphous phase is also magnesium silicates, while the crystal sizes (dXm) of Mg2Si04calculated from the full width at half maximum of the (131) peak were 31,40,37,43, and 28 nm for X = 0, 0.1, 0.2, 0.6, and 1.0, respectively. Considering experimental errors, the dXRD of Mg2Si04 did not change so much for all powders, suggesting that the crystal size is not controlled by addition of excess Si02 in this study, different from ZnO/Si02 synthesis.’I Figure 9 shows normalized PL spectra for X = 0.1, 0.2, 0.6, and 1.0. For X=O.l and 0.2, the PL spectra were primarily similar to that for X = 0. On the other hand, broad emission at 850-1100 nm was seen in addition to the emission at 1200 nm for X = 0.6, suggesting existence of Cr3+ ions substituting Mg2+ sites as discussed before (Section III(c)). The broad peak was intense for X = 1.0 where the amorphous phase increased, suggesting that C?’ ions are located mainly in the amorphous magnesium silicate.

IV. Summary The MgzSi04:Cr nanoparticles of several 10 nm in diameter were prepared by FSP, although a small amount of submicrometer-

96

Acknowledgments The authors are grateful to Dr. H. Nozaki (Toyota CRDL) for ESR measurement .

References ‘M. Yoshidd, M. Ldl, N. D. Kumar, and P. N. Prasad, “TiOZ Nano-ParticleDispersed Polyimide Composite Optical Waveguide Materials Through Reverse Micelles,” J. Muter. Sci.,32, 4047-51 (1997). ’Japanese Patent, Kokai Number 199664894. ’V. Petricevic, S. K. Gayen, R. R. Alfano, K. Yamagishi, H. Anzai, and Y. Yamaguchi, “Laser Action in Chromium-Doped Forsterite,” Appl. Phys. Lett., 52, 104lL2 (1988). 4D. G . Park, J. M. Burlitch, R. F. Geray. R. Dieckmann. D. B. Barber, and C. R. Pollock. ‘ % G e l Synthesis of Chromium-Doped Forsterite,” Chem. Muter., 5, 518-24 (1993). ’P. S. Devi, H. D. Gafney. V. Petricevic, and R. R. Alfano, “Synthesis and Spectroscopic Properties of Cr4+ Doped Sol-Gels;’ J. Nan-Cryst. Solids,203, 78x3 f 1996). ~~

6‘ P. S. ’Devi, H. D. Gafney, V. Petricevic, R. R. Alfano, D. He, and K. E. Miyano. “Sol-Gel Svntheses and Soectroscooic Characterization of ChromiumDoped Silicates a i d Germanates,’: Chem. haw.,12, 1378-85 (2000). ’M. Sokolowski, A. Sokolowska, A. Michalski. and B. Gokieli, “The “lnFlame-Reaction’’ Method for A1203Aerosol Formation,’’ J. Aerosol. Sci., 8.2 19-

30 (1977).

‘R. M. Laine, T. Hinklin, G. Williams, and S. C. Rand, “Low-Cost Nanopowders for Phosphor and Laser Applications by Flame Spray Pyrolysis”; pp. 5CKLlO in Metustuhle. Mechanically Alloyed and Nunocrystalline Materiuls, PIS 1 und 2 , Materials Science Forum 343-3, Edited by J. Eckert. H. Schlorh, and L. Schultz. Trans Tech Publications Ltd., Zurich Uetikon. Switzerland, 2000. 9L. Madler, H. K. Kammler, R. Muller, and S. E. Pratsinis, “Controlled Synthesis of Nanostructured Particles by Flame Spray Pyrolysis,” J. Aerosol. Sci., 33, 369-89 (2002). ”A. U. Limaye and J. J. Helble, “Morphological Control of Zirconia Nanoparticles Through Combustion Aerosol Synthesis,” J. Am. Ceram. Sac., 85, 112732 (2002).

“T. Tani. L. Midler. and S. E. Pratsinis, “Zinc Oxide/Silica Composite Nanoparticle Synthesis by Flame Spray Pyrolysis,’’ J. Muter. Sci., 37, 4627-32

(2002).

”T. Tani, K. Takatori, and S. E. Pratsinis, “Evolution of the Morphology of Zinc Oxide/Silica Particles Made by Spray Combustion,” J. Am. Ceram. Soc., 87,

365-70 (2004).

”S. E. Pratsinis, “Flame Aerosol Synthesis of Ceramic Powders.” Prog. Energy Comhu7t. Sci.. 24. 197-219 (1998). I4D. R. Lide, CRC Hundbook of Chemistry und Physics, 81st edition, CRC Press, Boca Raton, FL, 2000. ”G. L. Messing, S. C. Zhang, and G. V. Jayanthi, “Ceramic Powder Synthesis by Spray-Pyrolysis,” J. Am. Cerum. Soc., 76, 2707-26 (1993). 16R. A. Nyquist and R. 0. Kagel, lnfrured Spectra oJ lnorgunic Componud,v. Academic Press Inc., Orlando, FL, 1971. ”R. Moncorge, G. Cormier, D. J. Simkin, and J. A. Capobianco, “Fluorescence Analysis of Chromium-Doped Forsterite (Mg2Si04),”IEEE J. Quantum Electron.. 27, 114-20(1991). “R. Moncorge, M. Bettinelli, Y. Guyot, E. Cavalli, J. A. Capobianco, and S. Girard, “Luminescence of Ni[2]+ and Cr’+ Centres in MgSiO, Enstatite Crystals,” J. Phys.-Condes. Mufrer, 11, 683141 (1999). I9D. E. Budil, D. G . Park, J. M. Burlitch, R. F. Geray, R. Dieckmann, and J. H. Freed, “9.6-Ghz and 34-Ghz Electron-Paramagnetic-Resonance Studies Chromium-Doped Forsterite,” J. Chem. Phys., 101. 353848 (1994). ’OR. S. Roth, Phase Equilibria Diugrams. Vol. 111, Fig. 10369. The American Ceramic Society, Westerville, OH, 2001.

Progress in Nanotechnology: Processing

Formation of AI2O3-TiC Composite Nano-Particles Synthesized from Carbon-Coated Precursors Hisashi Kaga*’+?$ and Rasit Koc* Department of Mechanical Engineering and Energy Processes, Southern Illinois University at Carbondale, Carbondale, Illinois 62901

The formation of nano-sized alumina-titanium carbide ( A 1 2 0 ~ TIC) composite powders from a carbon-coated titanium dioxide-aluminum (TiOIAI) mixture was investigated. The carbon-coated TiOIAl mixture altered the mechanism of the reaction, compared with standard mixtures, to produce highquality nano-sized AI2O3-TiC powders. Data synthesized from intermediate temperatures indicate that these products form via Ti203 and A13Ti. TEM images of the A120TTiC powders showed fine size (50-100 nm), narrow size distribution, and lack of agglomeration. DSC data for the carbon-coated TiOl A1 mixture showed a single endothermic and four successive weak exothermic reactions as the carbon coating moderated the heat release during the reaction.

A

I. Introduction

ceramic materials are used in a great variety of applications. Among these advanced ceramics, aluminatitanium carbide (AI2O3-TiC) composites have been used as cutting tools and magnetic head slider substrates for a hard disk drive. This application is aided by the combination of hightemperature strength, high wear resistance, and good corrosion resistance.’-’ For these applications as with other ceramic materials, A1203-TiC powders with a homogenous chemical composition, fine particle size, narrow particle size distribution, and loose agglomeration are required for the manufacture of the final components.8 Several different methods exist for synthesizing AI2O3-TiC composite powders, each resulting in various particle size distributions, morpholog , state of agglomeration, Y*9 Efficient production of chemical purity, and stoi~hiometry.~high-quality AI2O3-TiC powder begins with low-cost, high-purity raw materials. Titanium dioxide (Ti02), aluminum (Al), and carbon (C) powders are commonly used for an aluminothermic reaction to produce AI2O3-TiC powders in industry. The aluminothermic reaction produces large amounts of powder from inexpensive precursor materials. However, there is presently no commercial process for the production of submicrometer AI2O3-TiC powders. The arrangement and particle sizes of Ti02, Al, and C in the mixture must be controlled in order to produce high-quality, low-cost AI2O3-TiC composite powders. The carbon coating of Ti02 or a TiO2-AI mixture can be a very effective way to control precisely the reaction path and rate of the aluminothermic reaction for the production of superior AI2O3-TiC composite powders.

The carbon coati:f,Rrocess was developed and patented by Koc and Glatzmaier for the production of TIC and silicon carbide. The advantage of the coating process is the intimate contact between reactants. This intimate contact allows reactions to occur closer to the thermodynamic predictions than conventional mixtures. The industrial production of AI2O3-TiC powders generally occurs as per reaction. 3Ti02 + 4A1+ 3C -+ 2A1203

+ 3TiC

This process is widely used because of the inexpensive starting materials, which can easily be obtained, and the reaction does not require high temperatures (1200°C).12-14Synthesis via reaction (1) is generally considered to be divided into two stages, shown in the following reactions”-”:

DVANCED

T. ksmann--wntrihuting editor

Manuscript No. 21949. Received June 28. 2006; approved September 18,2006. Research for this paper was performed at Southern Illinois University at Carbondale, sponsored by the U S . Department of Energy, Assistant Secretary for Energy Efficiency and Renewable Energy, Office of Industrial Technologies, Advanced Industrial Materials Program, under contract DE-AC05-960R22464 with Lockheed Martin Energy Research Corporation/UT-Batelk. ‘Member, American Ceramics Society. ‘Author to whom correspondence should be addressed. E-mail: [email protected] ‘Present address: National Institute of Advanced Industrial Science and Technology (AIST) 2266-98 Anagahora, Shimo-shidami, Moriyama-ku, Nagoya 463-8560, Japan.

Synthesis Methods for Powders

Ti

+ C + TIC

(3)

In reaction (2), molten aluminum reduces the surrounding solid Ti02 particles to titanium. The heat released in reaction (2) provides the required energy for reaction (3), in which the reduced titanium reacts with free C to form TIC. Although the formation of AI2O3-TiC composites generally follows the above sequence, the effect of C-coated starting materials has not been investigated. This paper describes the formation of AI2O3-TiC from Ccoated precursor materials. The effectiveness of the C-coated precursor was investigated in terms of the formation of A1203TIC compared with standard mixtures. The C-coated precursor provides a low-cost method to produce high-purity nano-sized AI2O3-TiC composite powders. 11. Experimental Procedure

(1) Starting Materials The starting Ti02 powder (P-25, Degussa Corp., Ridgefield Park, NJ) has a high specific surface area (51 m2/g), and Xray diffraction (XRD) showed that both anatase and rutile phases were present. Propylene gas (C3H6) was used as the C source for the C coating process. Carbon black (Monarch 880, Cabot, Waltham, MA) and Al (41000, Alfa Aesar, Ward Hill, MA) powders with specific surface areas of 220 and 1 m2/g, respectively, were used as starting powders for mixtures. Three different starting mixtures were used to prepare three precursors: fully coated, partially coated, and fully mixed. Figure 1 shows schematic drawings of the three precursors. ( A ) Preparation of C-Coated Ti021AI Mixtures (Fully Coated): The fully coated precursor was prepared by coating a mixture of Ti02 and Al. First, Ti02 and Al powders at the desired weight percent ratios were thoroughly mixed in ethyl alcohol by an attritor (HD01-A100-9, Union Process, Akron, OH) at 140 r.p.m. for 4 h in air. The ratio of the powders to ethyl 97

Fig. 1. Schematic illustrations of the three precursors investigated in this research.

alcohol was 2.1 by volume. After drying in a vacuum furnace (1400E, VWR Scientific, Philadelphia, PA) for 24 h, approximately 100 g of the Ti02/AI mixture was placed in a stainlesssteel vessel (10 cm inner diameter x 35 cm length) and subjected to the carbon coating process as explained el~ewhere.'*~'~ The steel chamber was purged with argon and evacuated. The chamber was then filled with C3H6 gas to a pressure of 275 kPa and heated to 550"C, the temperature at which C3H6decomposes to C, HZ,and CH4. After reaching 550"C, the system (powder, gas, and steel vessel) was allowed to equilibrate for 60 min. The hydrocarbon gas was then released to atmospheric pressure and refilled to 275 kPa. C, released from the hydrocarbon, coated the surfaces of the Ti02/A1mixture. After approximately 5 min, the remaining CH4 gas was released and the vessel was refilled with fresh hydrocarbon gas. This process was repeated until approximately 9.3 wt% of C was deposited, which required about 20 cycles (each cycle was 10 min). The weight percentage of the C in the precursor was determined using themogravimetric analysis. As the source of C is a hydrocarbon gas, impurities commonly found in C black, e.g., iron, are not introduced into the system. As the starting powders also have a high surface area, these precursors should result in high-purity, high-surface-area AI2O3-TiC powders. ( B ) Preparation of C-Coated Ti02 and A1 Mixtures (Partially Coated): The partially coated precursor was prepared by coating Ti02 powders as above, and then mixing this coated powder with Al powder. The C-coated Ti02 powders were mixed with Al in the attritor for 1 h in air. They were not mixed with ethyl alcohol to prevent removal of the uniform C coating of TiO2, which can occur with wet mixing. The partially coated precursor was prepared in order to prove the effectiveness of the C coating. If the Ti02 is uniformly coated by C, the C coating will act as a barrier and inhibit the reaction between Ti02 and Al, which is the first step of the aluminothermic reaction, and prevent the overall reaction of Eq. (1) from occurring. ( C ) Preparation of Ti02,Al, and C Black Mixtures (Fully Mixed): The fully mixed precursor was prepared using a method similar to that used in industry, with no C coating attempted. The appropriate weight percent ratios of TiO2, Al, and C black were mixed in ethyl alcohol in the attritor for 4 h in air. After drying, the specific surface area of the mixed precursor was measured to be 50 m2/g. The mixed precursor was intended for comparison with the coated precursors.

(2) Synthesis of A120J-TiC Powders All three precursors were reacted in a tube furnace (CTF 17/75/ 300, Carbolite, Sheffield, U.K.) with an inner diameter of 7 cm; 10 g of a given precursor was placed in a graphite crucible (6.5 cm inner diameter x 20 cm long) inside the furnace. The AlzO3T i c powders were then produced by promoting the aluminothermic reaction. The fully coated precursors were heated for 2 h at 600", 800", IOOO", 1200", and 1500°C to investigate the reaction mechanisms. The partially coated and mixed precursors were synthesized at 1000", 1200", and 1500°C for 2 h. All samples were heated in a flowing argon atmosphere; the heating and cooling rates were 4"C/min. 98

(3) Characterization of Synthesized A120TTiC Powders XRD (DMAX-B, Rigaku, Tokyo, Japan) of powders was carried out on a lab X-ray source using CuKa radiation for phase identification of samples. Each powder sample was scanned from 28 = 20"-80" at a scanning speed of 4"/min. The specific surface areas of the powders were measured using a BET gas adsorption surface area analyzer (Gemini 2360, Micromeritics, Norcross, GA) with 0.34.4 g of each sample. Samples were degassed at 175" for 4 h before the BET multipoint specific surface area was measured based on the amounts of gases absorbed with the corresponding pressures. TEM (FA7100, Hitachi Inc., Tokyo, Japan) imaging was used to examine the C coating on fully and partially coated precursors, as well as to examine the morphology of reacted powders. Samples were prepared and imaged on holey C formvar grids. DSC was utilized to describe the formation mechanism of AI2O3-TiC powders. Pellets of each precursor were prepared by single action, uniaxial pressing in a hardened steel die (inner diameter = 31 mm). The pellets were broken to accommodate the size restriction of the DSC apparatus (Labsys TG-DTA/DSC, Setaram Corp., Caluire, France). These samples were heated to 1200°C in a Pt crucible at 20"C/ min in flowing argon gas. Each experiment was repeated to ensure consistent results. III. Results and Discussion

( I ) C Coating Precursors A very uniform coating resulted from the carbon coating process. Figure 2(A) shows the TEM micrograph of the fully coated precursor. It shows two large particles in the lower region, which are Al, while the smaller dark particles are Ti02. Both Al and Ti02 particles are uniformly surrounded by the slightly lighter area, which is the C coating. BET specific surface areas of the uncoated Ti02/AI mixture and Ti02were measured to be 35 and 51 m2/g, respectively. After being coated with 9.3 wt% C, the surface areas were 31 and 45 m2/g, respectively; the reduced specific surface area indicates an increase in particle size. The TEM and BET results indicate that a uniform C coating can be obtained on the Ti02/AI mixture by using a hydrocarbon gas. Figure 2(B) shows the TEM micrograph of the partially coated precursor, and a higher magnification on the coated region is also presented. This image was obtained before Al was mixed with C-coated Ti02, and All Ti02particles were uniformly coated by C. Figure 2(C) shows the TEM micrograph of the fully mixed precursor. The dark particles indicate Al (large) and TiOz (small), as indicated in the image. The lighter particles in the upper middle portion indicate C black. Unlike fully and partially coated precursors, the mixed precursor has very little contact area between the C and TiO2 or Al. ( 2 ) Results of Synthesis ( A ) Fully Coated Precursor: Figure 3 shows the XRD patterns of the fully coated precursor and its products after synthesis at various temperatures. The diffraction pattern for the precursor exhibits peaks corresponding to Ti02 and Al. The sample synthesized at 600°C shows the same XRD pattern as the precursor because the aluminothermic reaction does not take place at this temperature. However, at 800"C, above the melting temperature of Al (660"C), the diffraction pattern shows intermediate phases, such as AI3Ti and TizO3, as well as the desired final product of AI2O3.The initial reaction between Al and Ti02 leads to the formation of A1203,A13Ti, and Ti203according to the following reaction: Eq. (4). 3TiO2 + 5AI + A1203

+ A13Ti + Ti203

(4)

The two intermediate phases Ti203 and A13Ti are considered to form A1203 and TIC in the presence of C according to the following equation: Eq. (5).

I STi203 + A13Ti + 4C -+ 1.5A1203 + TIC

(5)

Progress in Nanotechnology: Processing

Fig.2. TEM micrograph of the three precursors. (A) Fully coated precursor; The carbon coating is visible as a translucent fringe surrounding the A1 and TiOz particles. (B) Partially coated precursor; Ti02 particles are shown coated with carbon and a higher magnification on the coated region was also presented. A1 powder was not yet mixed with the coated TiO2. (C) Fully mixed precursor; Al, TiOz. and carbon black particles are indicated.

However, at IOWC, there is an intermediate-phase Ti203 but not AI3Ti. It is believed that Ti203 reacts with C to form TIC and CO(g) according to: Ti203 + 5C -+ 2TiC

+ 3CO(g)

(6)

This was supported by the fact that the total C content decreased by about 10% after synthesis as COk) was produced. As the total amount of COk) was approximately two orders of magnitude lower than that of the solid phase, reaction (1) is not affected by the gaseous species, as reported by Choi and Rhee13 Fig. 3 also includes the BET surface areas of the fully coated precursor and the resulting products. At 800"C, the surface area is slightly higher than that of the 600°C product and the precursor due to intermediate phases. At 1000°C,the surface area decreased from that at 800°C because most intermediate phases transformed into the final products of A1203 and TIC. At both 1200" and 1500"C, intermediate phases were all transformed to A1203 and TIC. The decrease in the surface area of the 1500°C products relative to that from 1200°C is caused by Synthesis Methods for Powders

sintering or particle growth in the powder at higher reaction temperatures.'5316Figure 4 shows a TEM micrograph of reaction products from the fully coated precursor at 1200°C. Particles range in size from 50 to 100 nm and are mostly spherical with a few edges. ( B ) Partially Coated Precursor: Figure 5 shows the XRD patterns of the partially coated precursor synthesized at 1200" and 1500°C. At both temperatures, there exist unreacted and intermediate phases such as Ti02, Ti305, and unknown Ti,O,,, indicating very little A1203-TiC formation. The C coating acted as a barrier to Ti02-A1 particle contact, preventing the initial reaction of Ti02 and Al. In other words, the Ccoated Ti02inhibited reaction with A1 and the aluminothermic reaction did not take place at these temperatures. The BET surface areas were 30.4 and 2.8 m2/g for 1200" and 15OO"C, respectively. The high surface area at 1200°C is due to the presence of unreacted TiOz and C, while the low surface area at 1500°C is due to reaction between Ti02 and C. The C-coated Ti02 reacted with C to form TIC and CO(g), via Ti407 and Ti305 at higher temperatures, as Al did not reduce it to titanium.16 99

0

r

I

UI

e!

0

20

30

A

.

40

1200°C. 30.4d/g

.l* 50

281'

O A e

60

70

1

80

Fig. 5. X-ray diffraction pattern synthesized from the partially coated precursor along with the specific surface areas of each.

I?

? 20

IJ

p o 00

30

40

-

oyQooy I

-

Precursor, 31m2/g

I

50

28 / "

60

70

80

Fig.3.

X-ray diffraction patterns of fully coated precursor and its synthesized products along with the specific surface areas of each.

( C ) Fully Mixed Precursor:

No extra phases, except Al2O3 and Tic, were observed to be synthesized at IOOO", 1200", and 1500°C. However, the results of the surface area showed that a low specific surface area < 1 mZ/gwas obtained for all samples. The powder obtained was much larger than that obtained from the fully coated precursor. As a result, the proposed processing method is advantageous for the formation of nano-sized A1203-TiC composite powders.

(3) Effect of C Coatings To understand the effect of the carbon coating and the formation mechanism, all three precursors were examined by the DSC. The heat flow as a function of temperature was recorded, and the DSC curves are shown in Fig. 6. All DSC curves show an absorption peak at about 670°C, corresponding to the melting of Al. In the Ti02-AI-C system, it is believed that Ti02 and A1 react initially to form A1203 and Ti as shown in reaction (2)?,1*,13 In other words, TiOz and Al must be in contact for the first step of the aluminothermic reaction to occur. In the case of the partially coated precursor, the C coating inhibits the reaction between Ti02 and Al by preventing their contact. The DSC data for the partially coated precursor had two minor exothermic reactions, with peak maxima at 801 and 866°C. This confirms the XRD results that the aluminothermic reaction

500 600 700 800 900 I000 1100 1200 TemperatureP C

Fig. 6. DSC patterns for A1203-TiC formation from three different

precursors. was incomplete. This is explained by the fact that the Ti02 was uniformly coated by C, which provides a suitable method for controlling the reaction path and rate. The mixed precursor had a single large exothermic reaction with a maximum at 941°C. Clearly, the mixture of Al, Ti02, and C released its heat at once with no control of the reaction rate due to the absence of a C coating. This large exothermic reaction sinters the powder as indicated by a low specific surface area (< 1 m2/g) in the final product. For the fully coated precursor, there were four successive weak exothermic reactions, with peak maxima at 823", 871", 968", and 1092°C. The C coating acts as a barrier to control the reaction path and rate when heat is released by the reaction between Al and Ti02. This leads to low heat-released exothermic reactions and a high surface area of the final product. From these results, coating the Ti02/AI mixture with C provides an effective means to control the reaction path and rate. This control of the reaction leads to the production of AI2O3-TiC composite powders with nano-sized powders.

IV. Conclusion

Fig. 4. TEM micrograph of the reaction product from the fully coated precursor after 2 h at 1200°C in flowing Ar.

100

AI2O3-TiC composite powders were synthesized using three different precursors. The fully coated precursor showed that the products had only two phases, A1203and TIC, which formed via intermediate phases of Ti203, A13Ti, and gaseous CO. The partially coated precursor showed that the reaction did not achieve completion due to the C coating preventing a reaction between the Al and Ti02. The mixed precursor showed that the synthesized powder had a very low surface area, even though A1203and T i c were produced. These results show that there are three advantages of using the fully coated precursor instead of mixed precursors, which are commonly used in industry. First, the C-coated Ti02/AI mixture leads to an intimate contact of the reactants and can be a very effective way to control the reaction path and rate. Second, the C coating inhibits the partial sintering of particles by moderating the magnitude of the exothermic reProgress in Nanotechnology: Processing

action. Third, it produces high-purity A1203-TiC composite powders because the hydrocarbon gas provides a non-contaminated C source. Consequently, the new process results in a very complete reaction, yielding high-purity and nano-sized A1203TIC powders that meet the exact requirements for the manufacture of ceramic materials and composites.

Acknowledgment The authors wish to thank Dr. Geoffrey A. Swift and Zach Crothers for contributions to this manuscript.

References ‘A. G. King, “Ceramics for Cutting Tools,” Am. Ceram. Soc. Bull., 4 3 , 3 9 5 4 3 (1965). ’R. P. Wahi and B. Ilschner, “Fracture Behaviour of Composites Based on A1203and TIC,” J . Marer. Sci., 15, 875-85 (1980). ’R. A. Culter. A. V. Virkar, and J. B. Holt, “Synthesis and Densification of Oxide Carbide Composites,” Ceram. Eng. Sci. Proc., 6, 715-28 (1985). 4S. J . Burden, “Comparison of Hot-Isostatically-Pressed and Uniaxially HotPressed A1203-TiC Cutting Tool,” Am. Ceram. Soc. Bull., 67, 1003-5 (1988). ’ S . Adachi, T. Wada. T. Mihara, Y. Miyamoto, and M. Koizumi, “HighPressure Self-Combustion Sintering of Alumina-Titanium Carbide Ceramic Composite,” J . Am. Ceram. Soc., 73, 1451-2 (1990).

Synthesis Methods for Powders

“R. A. Culter. A. C. Hurford. and A. V. Virkar, “Pressureless-Sintered Al2O3T i c Composite,” Marer. Sci. Eng. A , 105/106, 183-92 (1988). ’A. Krell and P. Blank. “Tic-Strengthened A1203 by Powder Tailoring and Doping Procedures,” Muter. Sci. Eng. A , 161, 295-301 (1993). ‘M. N. Rdhaman. Ceramic Processing and Sintering, ch. 2, pp. 3 8 4 2 . Marcel Dekker Inc, New York. 1995. 9C. R. Bowen and B. Derby, “The Formation of TiC/A1203Microstructures by a Self-propagating High-Temperature Synthesis Reaction,” J . Marer. Sci., 31, 3791-803 (1996). ‘OR. Koc and G. Glatzmaier, “Process for Synthesizing Titanium Carbide, Titanium Nitride, and Titanium Carbonitride, US Pat.5, 417, 952 (1995). “ G . Glatzmaier and R. Koc, “Method for Silicon Carbide Production by Reacting Silica With Hydrocarbon Gas, US Pat.5, 324, 494 (1994). I2T. D. Xia, Z. A. Munir, Y.L. Tang, W. J. Zhao, and T. M. Wang, “Structure Formation in the Combustion Synthesis of AI2O3-TiC Composites,” J . Am. Ceram. Sac.,83, 507-12 (2000). I3Y, Choi and S.-W. Rhee, “Effect of Precursors on the Combustion Synthesis of TiC-AIZ03 Composite,” J . Mater. Res., 9, 17614 (1994). I4M. P. Borom and M. Lee, “Rapid Rate Sintering of AI2O3-TiC Composites for Cutting-Tool Application,” Adv. Ceram. Mafer., 3, 3 8 4 (1988). I5R. Abramovici, “Composite Ceramics in Powder or Sintered Form Obtained by Aluminothermal Reactions,” Muter. Sci. Eng., 71, 313-9 (1985). I6C. R. Bowen and B. Derby, “Modelling of Self-propagating High-Temperature Synthesis Reactions,” Br. Ceram. Proc., 50, 2%38 (1993). I7C. R. Bowen and B. Derby, “Self-Propagating High Temperature Synthesis of Ceramic Materials,” Br. Ceram. Trans., 96, 25-30 (1997). “R. Koc and J. Folmer, “Synthesis of Submicron Titanium Carbide Powders,” J . Am. Ceram. Sue., 80, 952-6 (1997). I9G. A. Swift and R. Koc, “Formation Studies of TIC From Carbon Coated Ti02,” J . Mafer. Sci., 34, 308S93 (1999).

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SYNTHESIS OF Sm0.5Sr0.~Co03-, AND La0&-0.4Co03-,NANOPOWDERS BY SOLUTION COMBUSTION PROCESS Narottam P. Bansal National Aeronautics and Space Administration Glenn Research Center Cleveland, OH 44 135 Zhimin Zhong QSS Group, Inc. NASA Glenn Research Center Group Cleveland, OH 44 135

ABSTRACT Nanopowders of S ~ O . S S ~ O . ~ C (SSC) O O ~and - , Lao.$3-0.&003-, (LSC) compositions, which are being investigated as cathode materials for intermediate temperature solid oxide fuel cells, were synthesized by a solution-combustion method using metal nitrates and glycine as fuel. Development of crystalline phases in the as-synthesized powders after heat treatments at various temperatures was monitored by x-ray diffraction. Perovskite phase in LSC formed more readily than in SSC. Single phase perovskites were obtained after heat treatment of the combustion synthesized LSC and SSC powders at 1000 "C and 1200 "C, respectively. The as-synthesized powders had an average particle size of -12 nm as determined from x-ray line broadening analysis using the Scherrer equation. Average grain size of the powders increased with increase in calcination temperature. Morphological analysis of the powders calcined at various temperatures was done by scanning electron microscopy. 1. INTRODUCTION Solid oxide fuel cells (SOFC) are being considered' as the premium power generation devices in the fbture as they have demonstrated high energy conversion efficiency, high power density, extremely low pollution, in addition to flexibility in using hydrocarbon fuel, A major obstacle for commercial applications of SOFC still is high cost, both in terms of materials and processing. Intermediate Temperature Solid Oxide Fuel Cell (IT-SOFC) operated between 50O-80O0C, which allows utilization of available and inexpensive interconnects and sealing materials, can significantly reduce the cost of SOFC. The IT-SOFC also will have better reliability and portability. To keep up with the performance of traditional SOFC that operates between 900-1000°C, new materials with improved performance have to be used2*'. To enhance the oxygen ion conductivity of the electrolyte at the reduced temperature, La~-,Sr,Ga~-yMg,O, (LSGM), scandium stabilized zirconia or lanthanum (gadolinium, samarium) doped ceria can be used to replace the yttrium stabilized zirconia. Similarly, cathode materials with higher performance at the lower temperature such as Sm&30.~CoO3-, (SSC), L ~ O . ~ S ~ O . ~(LSC), COO~-~ Lao.~Sro.2Coo.2Feo.~03-, (LSCF) will be used to substitute Lal-,SrYMn03-,(LSM), the performance of which decreases rapidly when the operating temperature is below 800°C. The primary objective of this study was to synthesize fine powders of SSC and LSC compositions for applications as SOFC cathodes. A number of approaches such as, solid state reaction, sol-gel, hydrothermal, spray-drying, freeze-drying, co-precipitation, and solution

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combustion have been used for ceramic powders processing. The solution-combustion method is particularly useful in the production of ultrafine ceramic powders of complex oxide compositions in a relatively short time. This approach has been utilized4-" for the synthesis of various oxide powders such as ferrites, chromites, manganites, Ni-YSZ cermet, zirconates, doped ceria, hexaaluminates, pyrochlores, oxide phosphors, spinels, etc. An amino acid such as glycine is commonly used as the fuel in the combustion process. However, urea, citric acid, oxylydihydrazide, and sucrose have also been recently utilized6>l o as complexing agents and fuel in the combustion synthesis. In the present study, SSC and LSC cathode powders were synthesized using the glycinenitrate solution-combustion t e ~ h n i q u ebecause ~-~ of its high energy efficiency, fast heating rates, short reaction times, and high reaction temperatures. This process is also unique as all the reactants are mixed in solution at the molecular level resulting in homogeneous reaction products and faster reaction rates. Development of crystalline phases in the powders, on heat treatments at various temperatures, was followed by powder x-ray diffraction. Morphology of the powders was characterized by scanning electron microscopy (SEM). 2. EXPERIMENTAL METHODS 2.1. Powder Synthesis: The starting materials used in the synthesis were metal nitrates Sm(N03)3.6H20 (99.9 % purity), La(N03)3.6H20 (99.9% purity), Sr(N03)~(98 % purity), Co(N03)2.6H20 (97.7 % purity) and glycine (NH2CH2COOHY99.5 % purity), all from Alfa Aesar. A flow chart showing the various steps involved in the synthesis of powders by the solution-combustion process is shown in Fig. 1. Metal nitrates are employed both as metal precursors and oxidizing agents. Stoichiometric amounts of the metal nitrates, to yield log of the final SSC or LSC oxide powder, were dissolved in deionized water. A calculated amount of the amino acid glycine (0.7 mole per mole of NO3-) was also dissolved in deionized water. The glycine solution was slowly added to the metal nitrate aqueous solution under constant stirring. Glycine acts as a complexing agent for metal cations of varying sizes as it has a carboxylic group at one end and an amino group at the other end. The complexation process increases the solubility of metal ions and helps to maintain homogeneity by preventing their selective precipitation. The resulting clear and transparent red colored solution was heated on a hot plate until concentrated to about 2 mole/liter on metal nitrate basis. While the solution was still hot, it was added drop wise to a 2 liter glass beaker that was preheated between 300-400°C. The water in the solution quickly evaporated, the resulting viscous liquid swelled, auto-ignited and initiated a highly exothermic self-contained combustion process, converting the precursor materials into fine powder of the complex oxides. Glycine acts as a fuel during the combustion reaction, being oxidized by the nitrate ions. Oxygen from air does not play an important role during the combustion process. The overall combustion reactions can be represented as:

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0.6 La(N03)3 + 0.4 Sr(N03)2 + Co(N03)2 + 3.2 H2NCH2COOH + (1.8 - x/2) O2 + La0.6Sr0.4Co03-,+ 6.4 C02 + 8 H 2 0 + 3.9 N2

(1)

0.5 Sm(N03)3 + 0.5 Sr(NO3)2 + Co(NO3)z + 3.2 H2NCH2COOH + (1.95 - x/2) 0 2 + Smo.~Sro,sC003.,+ 6.4 C02 + 8 H20 + 3.85 N2

(2)

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1

lear red solution; heat at -80 "C; concentrate to -2M metal nitrate basis

beaker preheated to 300-400"C

Black powder; heat treat 700-1300 "C,2 h each, in air

(SEM) Figure 1.-Flow chart for solution-combustion synthesis of La0.6Sr0.4CoO~~ and Sm0.5Sr0.5CoO~~ nanopowders

indicating the formation of C02, N2, and H20 as the gaseous products. The evolution of gases during the combustion process helps in the formation of fine ceramic powder by limiting the inter-particle contact. The resulting black powder contained some carbon residue and was further calcined to convert to the desired product. Small portions (-1 g) of this powder were heat treated in air at various temperatures between 700 and 1300°C for two hours to study the development of crystalline phases. 2. 2. Characterization Thermal gravimetric analysis (TGA) of the powders was carried out using a PerkinElmer Thermogravimetric Analyzer 7 system which was interfaced with computerized data acquisition and analysis system at a heating rate of 10 "C/min. Air at 40 ml/min was used as a purge gas. X-ray diffraction (XRD) analysis was carried out on powders heat treated at various temperatures for crystalline phase identification and crystallite size determination. Powder XRD patterns were recorded at room temperature using a step scan procedure (0.02"/28 step, time per step 0.5 or 1 s) in the 28 range 10-70" on a Philips ADP-3600 automated diffractometer equipped with a crystal monochromator employing Cu K, radiation. Microstructural analysis was carried out using a JEOL JSM-840A scanning electron microscope (SEM). Prior to analysis, a thin layer of Pt or carbon was evaporated onto the SEM specimens for electrical conductivity.

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3. RESULTS AND DISCUSSION 3.1. Thermogravimetric Analysis Figure 2 shows the TGA curves recorded at a heating rate of 10"C/min in air from room temperature to 1200°C for the as-synthesized LSC and SSC powders using the solutioncombustion method. For both precursors, about 6% weight loss was observed between 600 to 850°C that was likely due to loss of carbon residue by oxidation and also from decomposition of SrC03. For SSC, there was additional 1% weight loss between 850 to 1000°C for which there is no simple explanation based on the x-ray diffraction results of Figure 4.

3.2. Phase Formation and Microstructure Both the LSC and SSC as-synthesized powders were calcined in air for two hours at various temperatures between 700 to 1300 "C to investigate the evolution of crystalline phases. X-ray diffraction patterns for these heat treated LSC and SSC powders are shown in Figs. 3 and 4, respectively and the results are summarized in Table I. The as-prepared LSC powder shows weak crystallinity of the perovskite phase. SrC03 phase was also observed in the as-synthesized powder and after calcination at 700 "C. An unknown peak at 32" (probably Sr3Co206.13, 83-375) appeared for the powder calcined at 800 and 900 "C. Formation of the perovskite phase, Lao.~SrO.4C003-~, is completed above 1000°C as observed by XRD results in Fig. 3. The asprepared SSC powder showed the presence of Sm203,Co304, and SrC03 phases. The desired Smo.~Sr0.~CoO3., perovskite phase emerged as the major phase after the powder was calcined at

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700 "C. Secondary phases such as Sr3Co206.13remained even after the powder was heat treated at 1100 "C. Perovskite phase-pure Smo.5Sro.~CoO3-,powder was obtained after heat treatment at 1200" C for 2 hours. Earlier investigation7 of SSC synthesis by solid-state reaction method indicated that the perovskite phase was formed after calcination at 1200°C for 6 hours. The products calcined at this temperature will have low porosity and non-ideal microstructure as cathode materials.

Table. I. X-ray diffraction analysis of Smo.5Sro.5Co03.x and Lao.~Sro.4CoO3.,powders made by solution-combustion synthesis after heat treatments at various temperatures in air Heat treatment Crystalline phasesa Temp. Time ("C) (h) As -- L ~ , , ~ S ~ O . ~ CSrC03 OO~-~, L%.6Sr0.4CoO3-x synthesized System

1000

2

1100

2

Average grain size (Wb 12

S ~ O . ~ S ~ O , ~Sr3C0206.13, C O O ~ - low ~ , intensity

38

Sm&ro.5Co03-x,Sr3C0206.13, low intensity

41

peak at 32" 20

peak at 32" 20

1200 smo.5sro.5coo3-x 2 1300 2 Smo.5SrosCoO3-x "Phases in decreasing order of peak intensity bCalculatedfrom Scherrer formula using FWHM of XRD peak in 47-48" range of 20. The SEM micrographs of Lao.~Sro.&o03., and Smo.&-o,5CoO3-,powders made by solution-combustion synthesis after heat treatments at different temperatures for 2 h in air are presented in Figures 5 and 6, respectively. The as prepared powders were highly porous and particles were linked together in agglomerates of different shapes and sizes. Substantial particle growth was observed upon calcination for two hours at 1000°C or higher temperatures. The particle size of samples calcined at 1000°C increased but the structure remained highly porous, which resembled the typical cathode structure for SOFC. Therefore, LSC and SSC powders

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should be sintered around 1000°C for fabrication of cathode structures. After calcination at 1200"C, LSC became dense and lost porosity. SSC powder sintered into a dense pellet following heat treatment at 1200°C.

3.3. Particle Size Analysis After each heat treatment of the as synthesized LSC and SSC powders, the average particle size was evaluated from X-ray line broadening analysis using the Scherrer equation'' :

where t is the average particle size, h the wave length of Cu K, radiation, B is the width (in radian) of the XRD diffraction peak at half its maximum intensity, and 0 B the Bragg diffraction angle of the line. Correction for the line broadening by the instrument was applied using a large particle size silicon standard and the relationship

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B~

=

B~~ - B~~

(4)

where BM and Bs are the measured widths, at half maximum intensity, of the lines from the sample and the standard, respectively. Values of average grain sizes of the as synthesized SSC and LSC powders and of those after heat treatments at various temperatures are given in Table I.

The as synthesized powders had an average grain size of about 10-12 nm. A number of factors are responsible for the nanosize of the resulting powders. Before the reaction, all the reactants are uniformly mixed in solution at atomic or molecular level. So, during combustion, the nucleation process can occur through the rearrangement and short-distance diffusion of nearby atoms and molecules. Also, large volume of the gases evolved during the combustion reactions (1) and (2) limits the inter-particle contact. Moreover, the combustion process occurs at such a

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fast rate that sufficient energy and time are not available for long-distance diffusion or migration of the atoms or molecules which would result in growth of crystallites. Consequently, the initial nanosize of the powders is retained after the combustion reaction. The X-ray line broadening method can be used only for the size determination of small crystallites (< 100 nm). The values obtained are not the true particle size, but the average size of coherently diffracting domains; the latter being usually much smaller than the actual size of the particles. The crystallite size of the as-synthesized powder depends'. on the glycine to nitrate ratio used during the combustion synthesis. Powder made using a fuel-deficient system has the highest surface area. The powder surface area decreases as the glycine to nitrate ratio is increased. This has been attributed to an increase in the flame temperature during combustion which helps in the growth of crystal size. The average grain size of the SSC and LSC powders increased (Table I) with the increase in calcination temperature, as expected.

4.SUMMARY AND CONCLUSIONS Nanopowders of Smo.sSr0.sCo03-~ (SSC) and La0.&,.&003-~(LSC) cathode materials for solid oxide fuel cells have been synthesized by the glycine-nitrate solution-combustion method. Formation of crystalline phases in both the powders started at relatively low temperatures. However, the as-synthesized powders had to be calcined at or above 1000 "C to yield phase pure perovskite products. The high temperature calcination caused significant reduction in surface area, coarsening of the powders, and sintering which is not favorable for forming the cathode structures for SOFC. The investigations of electrochemical activity of these materials and co-sintering with fuel cell electrolytes are being investigated and will be presented in the future. ACKNOWLEDGMENTS Thanks are due to Ralph Garlick for X-ray diffraction analysis. This work was supported by Low Emissions Alternative Power (LEAP) Project of the Vehicle Systems Program at NASA Glenn Research Center. REFERENCES 1. N. Q. Minh, Ceramic Fuel Cells, J. Am. Ceram. Soc., 76[3], 563-588 (1993). 2. D. Stover, H.P. Buchkremer, S. Uhlenbruck, Processing and Properties of the Ceramic Conductive Multilayer Device SOFC, Ceram. Int., 30 [7], 1107-1113 (2004). 3. Y. Liu, S. Zha, M. Liu, A&. Muter., 16 [3], 256-260 (2004). 4. S.-J. Kim, W. Lee, W.-J. Lee, S. D. Park, J. S. Song, and E. G. Lee, Preparation of Nanocrystalline Nickel Oxide-Yttria-Stabilized Zirconia Composite Powder by Solution Combustion with Ignition of Glycine Fuel, J. Muter. Res., 16 [12], 3621-3627 (2001). 5. L. A. Chick, L. R. Pederson, G. D. Maupin, J. L. Bates, L. E. Thomas, and G. J. Exarhos, Glycine-Nitrate Combustion Synthesis of Oxide Ceramic Powders, Muter. Lett., 10, 6- 12 (1990). 6. M. Marinsek, K. Zupan, and J. Maeek, Ni-YSZ Cermet Anodes Prepared by CitrateNitrate Combustion Synthesis, J. Power Sources, 106, 178-188 (2002). 7. T. Ishihara, M. Honda, T. Shibayama, H. Minami, H. Nishiguchi, Y. Takita, Intermediate Temperature SOFCs Using a New LaGa03 Based Oxide Ion Conductor. I. Doped SmCo03 as a New Cathode Material, J. Electrochem. Soc., 145 [9], 3177-3 183 (1998).

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8. R. D. Purohit, S. Saha, and A. K. Tyagi, Nanocrystalline Thoria Powders via GlycineNitrate Combustion, J. Nuclear Mater., 288 [l], 7-10 (2001). 9. T. Ye, Z. Guiwen, Z. Weiping, and X. Shangda, Combustion Synthesis and Photoluminescence of Nanocrystalline Y2Og Eu Phosphors, Mater. Res. Bull., 32, 50 1 (1997). 10. K. Prabhakaran, J. Joseph, N. M. Gokhale, S. C. Sharma, and R. Lal, Sucrose Combustion Synthesis of LaxSql-,$ln03 (x 5 0.2) Powders, Ceram. Int., 31 [2], 327-33 1 (2005). 11. B. D. Cullity, Elements of X-Ray Diffraction, TdEdition, Addison-Wesley, Reading, MA, p. 284 (1978).

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COLLOIDAL PROCESSING AND SINTERING OF NANO-Zr02 POWDERS USING POLYETHYLENIMINE (PEI) Yuji Hotfa'*, Cihangir Duran"2, Kimiyasu Sato' and Koji Watari' 1

National Institute of Advanced Industrial Science and Technology, Advanced Sintering Technology Group, Advanced Manufacturing Research Institute, Anagahora 2266-98, Shimoshidami, Moriyama-ku, Nagoya, Japan 2GebzeInstitute of Technology, Department of Materials Science and Engineering, PK 141,41400, Gebze, Kocaeli, Turkey *Correspondingauthor

ABSTRACT A stable colloidal suspension is important for fabricating dense samples with uniform microstructure using colloidal processing methods. Aqueous nano-Zr02 suspensions were prepared using polyethylenimine (PEI) as a dispersant. PEI adsorption on nano-Zr02 surfaces was influenced by PEI content and suspension pH. The isoelectric point (IEP) shifts from pH 7 at 0 wt% PEI to pH 10.3 at 3 wt% PEI. Stable suspensions had mean particle sizes in the range of 100 to 150 nm and sedimentation rates less than 0.4 &h, as compared to 2-5.5 pm and 10-50 mm/h for unstable suspensions. Electrostatic interactions, hydrogen bonding and PEI conformation were found to be controlling mechanisms on the colloidal stability of the suspensions. The amount of PEI adsorbed on nano-ZrO2 surfaces was characterized using Thermogravimetric analysis (TG) and Fourier transform infrared spectrometer (FT-IR). The densification behavior of samples containing 3 wt% PEI at pH 7.1 sintered for 4 h at 1000 to 1300 "C was characterized. Relative density was found to increase rapidly from 54 % at 1100 "C to 92 % at 1200 "C and finally to 98 % at 1300 "C. Therefore, pellets were sintered at 1300 "C for 4 h to quantitatively correlate the processing conditions such as effect of pH and PEI content with densification. INTRODUCTION The dispersion of ceramic powders in liquid is of importance in the colloidal processing methods such as slip casting. These methods have been shown to be superior to conventional dry pressing in terms of controlling density and uniform microstructure evolution in the green and sintered states." * If aggregates form, suspension stability and subsequently sintered properties are severely degraded. Therefore, aggregation must be prevented during colloidal processing3.

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Nano-particles can form aggregates in solution due to van der Waals attractive forces at short interaction distance^.^. Therefore, sufficiently large stabilizing forces such as electrical double-layer repulsion or steric interactions should be used to create an energy barrier to inhibit aggregation. Polyethylenimine (PEI) has been used as a dispersant for various ceramic powders and has been shown to enhance stability of ceramic powders in ~ a t e r ~When - ~ . PEI is dissolved in a neutral or acidic solution, proton adsorption results in the protonation of the amine groups. Therefore, positively charged PEI easily adsorbs on the negatively charged ceramic surfaces, which provides an electrosteric effect that prevents aggregation of the ceramic powders5. In this study, the effects of PEI content and pH on the dispersion of nano-Zr02 powders in water were characterized. The objectives were to investigate both PEI adsorption mechanisms on nano-ZrO2 particles and to define the pH range in which colloidally stable nano-Zr02 suspensions can be prepared to fabricate dense sintered ceramics. The effects of pH and PEI content on the suspension properties were characterized by measuring zeta potential, particle size and sedimentation rate. PEI adsorption on nano-Zr02 powders was evaluated from thermogravimetric analysis (TG) and FT-IR measurements. Green samples were used to compare densification behavior at 1300°C as a function of pH and PEI content. EXPERIMENTAL PROCEDURE 3 mol% yttria-stabilized nano-ZrO2 powders with average particle size of 50 to 75 nm were purchased from Aldrich. Surface charge of nano-ZrO2 particles was modified with polyethylenimine (PEI) (MW 10000, Anhydrous, Wako Pure Chemical Ind., Ltd., Japan). PEI was added at concentrations ranging from 0 to 3 wt% with respect to the dry weight of nano-ZrO2. 1 vol% nano-ZrO2 suspensions at various PEI contents (0 to 3 wt%) and pH values were prepared. First, PEI was dissolved in distilled water and then nano-ZrOz powder was added. Suspensions were ultrasonicated at 120 Watts for 10 min. pH was adjusted using reagent-grade HC1 and NaOH (Wako Pure Chemical Ind., Ltd., Japan). Then, the suspensions were stirred for 6 h and finally centrifuged at 5000 rpm for 1 second to remove bigger particles from the suspension. Zeta potential was manually measured by applying 50 V (Model 502, Nihon Rufuto Co., Ltd, Japan). Particle size distribution was measured by using a laser particle size analyzer (Horiba LA-920, Japan). Sedimentation behavior of suspensions was characterized with a pulsed near infrared light (Turbiscan ma 2000, Formulaction, France). A clarified region at the top, and the sediment region at the bottom were characterized as a function of time. In sedimentation kinetics analysis, thickness of the clarification region from the top was considered and light transmitted at 3 % intensity was used.

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The amount of PEI adsorption was determined from Thermogravimetric Analysis (TG) (Seiko Instruments SII, SSC/5200). Suspensions were centrifuged at 25000 rpm for 15 min to separate sediment and supernatant. The sediment was washed twice with distilled water to remove any excess (free) PEI. After final centrifugation, the sediment was dried under vacuum at 100°C for 2 h before analysis. TG of as-received ZrO2 powder was chosen as a reference. Weight loss between 200-600°C was used in the adsorption calculations. The surface characterization of nano-ZrO2 and PEI-modified nano-ZrO2 powders was studied by FT-IR (Perkin Elmer, Spectrum GX, USA) after the samples were mixed with KBr powder. IR spectra of as-received PEI and 25 wt% aqueous PEI solutions at various pH values were measured after each solution was sandwiched as a thin layer between two CaF2 plates. Green bodies were fabricated by slip casting using gypsum molds. Densification behavior as a function of sintering temperature was first tested on the samples containing 3 wt% PEI prepared at pH 7.1 from 1000 to 1300 "C with an isothermal hold of 4 h. PEI burn out and sintering were performed at the same heating cycle as the pellets were heated at 100 "Ck. Densities were calculated using the Archimedes principle. The relative density was found to increase sharply from 54 % at 1100 "C to 92 % at 1200 "C and finally to 98 % at 1300 "C. Therefore, pellets were sintered at 1300 "C for 4 h to quantitatively correlate the processing conditions such as effect of pH and PEI content with densification. Microstructure observations were carried out by a scanning electron microscopy (SEM) (Model JSM5600N, JEOL, Japan). RESULTS AND DISCUSSION Figure 1 shows the effect of pH and 80 PEI on zeta potential (5) of the nano-ZrO2 60 suspensions. Electrophoresis properties of 40 > ZrO2 in an aqueous solution determine the E 20 suspension stability. Attractive London-van .-cd z0 o der Waals forces should be overcome by 0 a -20 Q c . repulsive forces such as electrostatic or N -40 polymeric to prepare stable suspensions'. In -60 other words, high zeta potential value induces -80 sufficiently high surface charge, which causes strong repulsive double-layer force6. Figure 1 Zeta potential plots of suspensions as a function of pH and PEI content; ( 0 ) 0 wt% PEI, Nano-ZrOZ suspension with 0 wt% PEI has an (+) 0.5 wt% PEI, (A) 1 wt% PEI, (m) 2 wt% PEI, isoelectric point (pHiep) at nearly pH 7. PEI and ( 0 ) 3 wt% PEI. interacts with nano-ZrO2 surface and shifts the pHiep to more alkaline region as the amount of PEI content is increased. The pHiep increases from pH 7 at 0 wt% PEI to pH 10.4 at 3 wt% PEI. This

4-4

b)

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can be attributed to the fact that PEI is a cationic polyelectrolyte and addition of a strong acid to the PEI-containing solution neutralizes -NH- basic groups, resulting in a positively charged polymer skeleton according to the following reaction';

As a result, positively charged -NH;- groups can adsorb on negatively charged nano-ZrOz surfaces by an

30

0

/

electrostatic attraction, which causes the PHjep to shift to the alkaline region. Furthermore, nano-ZrO2 powders have various negative zeta potentials in the alkaline -72 region due to the different pHiep values such as mV at 0 wt% PEI compared to as low as -9 mV at 3 wt% PEI at pH 11. The zeta potential curves for the suspensions with PEI have a tendency to approximate the original curve of the suspension without PEI at highly acidic (e.g., pH 3) and highly alkaline (e.g., pH

c=

12) solutions. It was reported that the degree of PEI dissociation (a)increases with decreasing pH, that is, fully dissociated PEI at pH 2 (a=1) and undissociated

2000

1800

I600

Wavenumber, cm-'

1400

2000

0 0

25

0-

E 2o

2

2

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0

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10

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0

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0 0

5

'0 8 . 6

0'0 0

o ct

9J.o

5

-7 04.4

0.0

2.0

1.o

3.0

Initial PEI amount, wt%

Figure 2 PEI adsorption on nano Zr02 powders as a function of initial PEI . . content ana p n . iwmDers inaicare suspension pH. 1

1800

.

1

1600

1

TT

1.

1400

Wavenumber, cm-'

Figure 3. (a) FT-IR spectra of (A) as-received PEI and (B-F) aqueous PEI solutions at various pH; (B) pH 11.6, (C) pH 9.2, (D) pH 7.1, (E) pH 3.3 and (F) pH 0.4. (b) Effect of PEI content on FT-IR spectra of ZrO,; (G) PEI-free ZrO,, (H) 0.5 wt% PEI, (I) 1 wt% PEI, (J) 2 wt% PEI, (K) 3 wt% PEI and (L) PEI solution at pH 7.1.

(e.g., free of charge) PEI structure after pH 11 (a=O)can be obtained. Therefore, PEI is strongly

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desorbed due to a strong electrostatic repulsion between the positively charged polymer and solid sites at pH 3, and PEI does not carry charge at pH Figure 2 shows the amount of PEI adsorbed on nano-ZrO2 powders as a function of initial PEI content and pH. The plot represents the amount of PEI adsorbed per gram of nano-ZrO2. The dashed line is given as a reference for the complete adsorption. The amount of PEI adsorbed is incomplete at the as-prepared pH conditions although it increases with increasing polymer content. The amount of PEI adsorbed increases steeply with increasing pH, that is, it is almost nil (e.g., 0.05 mg/g) at pH 4.4 in contrast to nearly full (9.9 mg/g) at pH 10.8. Figures 3a and 3b compare the adsorption curves as a function of initial PEI contents and pH at 1 wt% PEI, respectively. The spectra of PEI-free ZrO2 and the PEI solution at pH 7.1 are shown for reference in both graphs. Consistent with the TG data given in Figure 2, the broad absorption band at 1555 cm", which becomes more dominant with increasing PEI content and pH in both plots, is evidence of the PEI adsorption onto the nano-Zr02 particles. Furthermore, a shift in the band position from 1514 cm-' to 1555 cm" can be attributed to the bonding of PEI on nano-25-02 surface via -NH2+- groups, as compared to the free PEI present in the solution. Furthermore, TG and FT-IR data (Figs. 2, 3b, and 3c) show that the amount of PEI adsorbed increases with increasing initial PEI content and suspension pH. The adsorption

12. 3 .5

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3 3 .

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2.5 .$? v)

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0.5

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0

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6

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10

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PH

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3.5 3 2.5

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PH Figure 4 Mean particle size and sedimentation rate as a function of pH and PEI content; (a) 0 wt% PEI, (b) 1 wt% PEI and (c) 3 wt% PEI. Vertical arrows show PHiep.

process can be explained by considering a and zeta potential (Fig. 1) as a function of pH. The sign and magnitude of electrical surface charge of nano-ZrO2 in the suspension can be modified

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with pH, which, in turn, strongly affects the adsorption process. The main adsorption mechanism was reported to occur by hydrogen bonding between isolated ZrOH groups on the Zr02 surface and the C-N- groups of PE14. Electrostatic repulsion begins to fade towards pHiep due to the decreased numbers of positively charged surface sites on Zr02, as compared to the ones at pH 4.4, which contributes to the adsorption process. After pHiep 7, nano-Zr02 particles have negative charge (Fig. 1) and, therefore, the positively charged PEI can easily and strongly adsorb on the negatively charged nano-Zr02 surfaces. Figure 4 summarizes and compares mean particle size and sedimentation rates as a function of pH and PEI content. The suspensions with various PEI contents exhibit similar behavior with pH, in that, both mean size and sedimentation rate increase sharply on reaching pHiep of the respective suspensions, as indicated by the vertical arrows on each plot. In general, the suspensions whose pH
-

4c). In addition, this suspension is unstable due to its proximity to the pHieP(10.4) and low 5 (-19 mV). 151 < 20 mV is considered to be unstable suspension". These results indicate that suspension

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stability is deteriorated due to the interconnection of nano-ZrOz powders by hydrogen bonding and coiled conformation of PEI molecules in the alkaline region. In other words, the pH range where 6 is negative represents a kinetically unstable region for the current system, which is characterized by coagulation of the polymer-particle system. Figures 4b and c further indicates that the sedimentation rate increases after pH>pHiep even though the particle size decreases slightly. The particle size measurements were carried out right after the sample preparation step and then the suspensions were subjected to the sedimentation tests for 24 h. Aggregation of the particles with time takes place due to the instability of the suspensions. In this case, the gravitation force accelerates the settling of the aggregates. However, the stability range in nano-Zr02 suspensions is successfully extended up to pH 9.1 with 3 wt% PEI addition (Fig. 4c). This can be attributed to the adsorption and conformation of PEI in the region where the 5 is positive. The PEI molecules undergo a transition from a compact coiled to a stretched conformation with increasing a (or, alternatively, decreasing pH) due to an electrostatic repulsion between the neighboring ionized sites. Therefore, there happens a strong electrostatic repulsion between both positively charged PEI sites and nano-Zr02 powders at pH<-7, which in fact limits the adsorption but also prevents the coagulation of nano-ZrOz powders. Figure 5 shows densification plots as a function of pH and PEI content. Arrows indicate the pHiep of the respective suspensions. Densification decreased suddenly on reaching pHiep for

'1 'z

85

8o

15

@Owt%PEI A 1 wt%PEI

'

I

I

I

I

I

4

5

6

7

8

9

I

I

1 0 1 1 1 2

PH Figure 5 Densification plots as a function of pH and PEI content. Samples were sintered at 1300°C for 4 h. Arrows show pHi,,.

each suspension. Irrespective of PEI content, the samples prepared from the suspensions whose pHpHi,,. Figure 6 shows SEM micrographs of

Synthesis Methods for Powders

119

Figure 6 SEM images of fracture surfaces of sintered compacts prepared at various PEI contents and pH values; (a) 0 wt% PEI and pH 7.2, (b) 1 wt% PEI and pH 9.1, (c) 3 wt% PEI and pH 7.1, (d) 3 wt% PEI and pH 8.4,(e) 3 wt% PEI and pH 10.3, respectively.

the fracture surfaces of the sintered samples. Figures 6a and 6b are micrographs of samples containing 0 and 1 wt% PEI prepared at their respective pHi,, values, respectively. These samples are made of dense regions with less-dense regions in which large, micro-scale pores are present. This can be attributed to the formation of aggregates, due to the instability of suspension when the pH changes, that giving rise to differential sintering in the compacts. In contrast, samples containing 3 wt% PEI dispersed at up to pH 9.1 are almost fully dense (Figs. 6c and 6d). Thus, PEI broadens the pH range in which the stable suspensions can be prepared In turn, this leads to the fabrication of a dense sample. In the dense samples, mostly nano-scale pores were concentrated either along the grain boundaries or at the triple grain junctions which indicates an effective pore removal (or densification). Note that final grain size in the dense samples is in the range of 1 to 3 pm. A sample with 3 wt% PEI at pHi,, 10.3 (Fig. 6e) shows microstructure evolution similar to that given in Figures 6a-b. These results strongly suggest that a dense ceramic can be fabricated from stable colloidal suspensions having a finer particle size and a narrow particle size distribution. CONCLUSION A colloidal dispersion of nano-ZrOz powders was successfully produced using polyethylenimine (PEI) as a dispersant. The addition of PEI shifted pHi,, from pH 7 at 0 wt% PEI to pH 10.4 at 3 wt% PEI. Due to the electrosteric stabilization effect of PEI, the range for

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Progress in Nanotechnology: Processing

colloidal stability was extended up to pH 9.1. Electrostatic interactions, hydrogen bonding and PEI conformation were found to be the controlling mechanisms for the colloidal stability of the suspensions. The amount of PEI adsorbed on nano-Zr02 surfaces increased with increasing initial PEI content and suspension pH. Green bodies made from this colloidally stable nano-Zr02 suspension sintered compacts with 98 % of relative density on heating at 1300°C for 4 h. REFERENCES I I. A. Aksay, F.F. Lange and B.I. Davis, “Uniformity of Al2O3-ZrO2 composites by colloidal filtration”, J. Am. Ceram. SOC.,66, C-190 (1983). *F.F. Lange, B.I. Davis and E. Wright, “Processing-related fracture origins: IV, elimination of voids produced by organic inclusions”, J. Am. Cerum. SOC.,69,66-68 (1986). 3V. A. Hackley, “Colloidal processing of silicon nitride with poly(acry1ic acid): I, Adsorption and electrostatic interactions”, J. Am. Cerum. SOC.,80,23 15-25 (1997). 4J. Wang and L. Gao, “Surface properties of polymer adsorbed zirconia nano particles”, NanoStructured Materials, 11,45 1-57 (1999). 5A. Dietrich and A. Neubrand, “Effects of particle size and molecular weight of polyethylenimine on properties of nanoparticulate silicon dispersions”, J. Am. Ceram. SOC.,84, 806-12 (2001). 6X. Zhu, T. Fengqiu, T. S. Suzuki and Y. Sakka, “Role of the initial degree of ionization of polyethylenimine in the dispersion of silicon carbide nanoparticles”, J. Am. Cer. SOC.,86, 189-91 (2003). 7J. Wang and L. Gao, “Adsorption of polyethylenimine on nanosized zirconia particles in aqueous suspensions”, J. Coll. and Int. Sci., 216,436-39 (1999). ‘R. Moreno, “The role of slip additives in tape-casting technology: Part I-solvents and dispersants”, Am. Cerum. SOC.Bull., 71, 1521 (1992) 9

S. Baklouti, C. Pagnoux, T. Chartier and J. F. Baumard, “Processing of aqueous a-&O3, a-SiOz, a-Sic suspensions with polyelectrolytes”, J. Eur. Ceram. SOC.,17, 1387-92 (1997). 10 S. Vallar, D. Houivet, J. El Fallah, D. Kervadec and J.-M. Haussonne, “Oxide slurries stability and powders dispersion: Optimization with zeta potential and rheological measurements”, J. Eu. Cerum. SOC.,19, 1017-21 (1999).

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SYNTHESIS OF HIGH PURITY P-SiAlON NANOPOWDER FROM A ZEOLITE BY GASREDUCTION-NITRIDATION Tomohiro Yamakawa, Tom Wakihara, Junichi Tatami, Katsutoshi Komeya and Takeshi Meguro Graduate School of Environment and Information Sciences, Yokohama National University 79-7, Tokiwadai, Hodogaya-ku Yokohama, 240-8501, Japan

ABSTRACT P-SiAlON powders were synthesized by gas-reduction-nitridation of zeolite using NH3 and C3Hg as reactant gases. Zeolite changed from crystalline into glassy phase around 900 "C. The peak intensity of P-SiAlON increased with increasing firing temperature and soaking time. It was confirmed that products were high purity P-SiAlON. Specific surface areas of the products nitrided at 1400 OC increased with increasing soaking time, and it reached 20 m2/g after firing at 1400 OC for 120 min. FE-SEM observation showed that morphology of the products obtained are composed of nanoparticles (ca. 20-50 nm). INTRODUCTION P-SiAlON ceramics are one of the most promising materials because it has potential for engineering applications owing to having the excellent properties, e.g. high strength and good thermal-shock resistance. P-SiAlON is a solid solution of P-Si3N4; a part of Si and N are replaced by A1 and 0. The chemical formula of P-SiAlON is indicated as follows:

P-SiAlON ceramics are commonly produced by a reaction sintering of a mixture of Si3N4, AlN, and A1203 over 1500OC. However, their corrosion resistance degrades due to glassy phase in the grein boundary Therefore, preparation of P-SiAlON powders followed by their sintering received considerable attention to avoid the formation of glassy phase '. In the present study, we focused on a zeolite as a raw material. Zeolites are hydrated, crystalline tectoaluminosilicates that are constructed from TO4 tetrahedra (T = tetrahedral atom, e.g. Si and Al) 3-5. Typical zeolites are aluminosilicate, which contain Si and A1 uniformly in an atomic scale. In our previous study 6, P-SiAlON was synthesized from a mixture of zeolite and carbon by carbothermal reduction-nitridation (CRN). In recent years, gas-reduction-nitridation (GRN) process, which uses a mixture of NH3 and hydrocarbon as reactant ases, has been proposed as one of the synthetic techniques of high-purity nitride powders 7- 1 GRN is the quite simple process since it dose not require mixing of raw materials with carbon and decarbonizing after nitriation. Moreover, formation of nitride via GRN occurres at lower temperatures than via CRN. Therefore, the objective of the present work is to synthesize high-purity P-SiAlON powder from zeolite via GRN.

'.

!

EXPERIMENTAL PROCEDURE A commercial zeolite powder (HSZ-330HUA, Tohso Chem., Tokyo) was used as a starting material. Its main characteristics are listed in Table 1. The raw zeolite has high specific surface area and is composed of fine particles about 0.6 pm. It was weighed on an alumina boat Synthesis Methods for Powders

123

and placed in an electric fhrnace with a high-purity alumina tube. Then, it was fired to 700 "C at a heating rate of 5 "C/min in Ar gas (99.999% purity) in order to eliminate oxygen in the system and remove the adsorption water in zeolite. It was confirmed that the structure of zeolite did not change during the pre-heat treatment up to 700 "C. Heating was continued in a gas mixture of NH3 (99.999% purity) and 0.5 vol% C3H8 (99.99% purity). The flow rate was 4 l/min. The sample was fired at 1200-1400 "C for 0-120 min before being cooled in N H 3 . Furthermore, CRN process was also applied for the synthesis of SiAlON using the same raw zeolite for comparison Phases present in the products were identified by X-ray diffractometry (RINT2500, Rigaku, Tokyo, Japan) using CuKa radiation operated at 50 kV and 300 mA. Morphological changes were observed by field emission scanning electron microscopy (FE-SEM; S-5200, HITACHI, Tokyo, Japan). Specific surface areas of the powders were measured by a singlepoint Brunauer-Ematt-Teller method (BET; Quantasorb, Quantachrome, Boynton Beach, FL).

Table 1. Characteristics of a raw powder

Zeolite type

Y-type 3 (z = l S )

Si/A1 (mol/mol) Mean particle size

(pm)

Specific surface area (m2/g>

0.6 550

RESULTS AND DISSCUSSIONS Figure. 1 shows XRD patterns of the products prepared by CRN and GRN methods using the same zeolite at 1400°C for 120 min. P-SiAlON was a main phase in both samples, but XRD pattern of the P-SiAlONs phase obtained by CRN method indicated the broad and splited peak in higher diffraction angle. This fact indicates the formation of P-SiAlONs with different compositions. On the other hand, high-purity P-SiAlON with an uniform composition was confirmed in the sample produced by GRN (Fig. 1(b)). XRD patterns of the products synthesized at 1200-1400 "C for 0-60 min, together with the raw zeolite are shown in Fig. 2(a). The raw zeolite was transformed into amorphous phase at 1200 "C (Fig.:! (b)). In the sample fired at 1400 "C without soaking time, 0-SiAlON and XSiAlON were confirmed to be formed (Fig. 2(c)) similarly to a result in a previous study 10. Finally, high-purity P-SiAlON phase was obtained by heat treatment at 1400 "C for 60 min. Preparation of P-SiAlON from zeolite by CRN requires firing at a higher temperature for longer soaking time to form P-SiAlON. Mullite phase, one of the stable phases at high temperatures was formed during nitridation process. In the case of CRN, as a result of heterogeneity of the composition, P-SiAlONs with different composition obtained. On the other hand, since GRN proceedes at lower temperature than CRN, it is possible to prevent the formation of mullite to maintain homogeneity of the chemical composition of starting materials, As a results, high-purity P-SiAlON with uniform compositon was obtained by GRN of zeolite.

I24

Progress in Nanotechnology: Processing

0 1

p

0 P-SiAlON 0

a-Si3N4

0

0

10

20

0

28 I deg.

40

60

Figure 1. XRD patterns of the products (a) carbothermal reduction-nitridation (b) gas-reduction-nitridation

20 I deg. Figure 2. XRD patterns of (a) raw material and the products synthesized at (b)1200 "C for 0 min (c) 1400 OC for 0 min and (d) 14OO0Cfor 60 min.

Synthesis Methods for Powders

125

Fig. 3 shows FE-SEM photographs of raw zeolite and the products synthesized at 12001400 "C for 120 min. Morphology of the raw zeolite (Fig. 3(a)) was approximately maintained after transition to glassy phase (Fig. 3(b)). As the soaking time increases to 1400 "C, equiaxed grains with 30-50 nm (Fig. 3(c)) were obtained in diameter.

Figure 3. FE-SEM photographs of (a) the raw material and the products synthesized at (b)1200 "C for 120min (c) 1400 "C for 60 min.

Fig. 4 shows the specific surface areas of the products synthesized at various reaction temperatures for 0-60 min. Specific surface area of a raw material drastically decreased with the amorphization. However, in the sample prepared at 1400 "C, it was found that the specific surface area gradually increased up to 20 m2/g with the formation of P-SiAlON.

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20

-

Heating

Soaking-

-!

/ \

-

h

w

3

W

1200

1300

1400 'c 0

30

60 min

Figure 4. Specific surface area of the products synthesized at various reaction temperatures at 1400 "Cfor 60 min.

CONCLUSIONS Zeolite was used as a raw material to prepare P-SiAlON by gas-reduction nitridation method. In consequence, high purity P-SiAlON powder was synthesized at 1400 "C for 60 min, and it was found that P-SiAlON obtaind are comosed of nanoparticles with diameter of about 3050 nm. Mullite phase could not be confirmed in the samples synthesized by GRN.BET surface area of the products increased along with the formation of p- SiAlON. These promising results suggest the potentiality of gas-reduction-nitridation for synthesizing high purity p- SiAlON nanoparticles from zeolite.

Synthesis Methods for Powders

127

REFERENCES 1 T. Sato, Y. Tokunaga, T. Endo, M. Shimada, K. Komeya, K. Nishida M. Komatsu and T. Kameda, “Corrosion of Silicon Nitride Ceramics in Aqueous Hydrogen Chloride Solutions,” J.Mater.Sci., 23 3440-3446 (1988). T.Ekstrom, P.O. Kall, M.Nygren and P.O.Olsson, “Dense single-phase 0-SiAlON ceramics by glass-encapsulated hot isostatic pressing” J.Mat.Sci., 24 1853-6 1 (1989). D.W. Breck, “Zeolite Molecular Sives” wiley, New York (1974). 4 R.M. Barrer, “Hydrothermal Chemistry of Zeolites,” Academic Press, London (1982). T. Wakihara and T. Okubo, “Hydrothermal Synthesis and Characterization of Zeolites,” Chem Llett., 34 278-79 (2005). F. Li, J. Tatami, T. Meguro and K. Komeya, “Synthesis of 0-SiAlON powder by Carbothermal Reduction-Nitridation of Zeolite” Key. Eng. Mater., 247 109-112 (2003). T. Suehiro, J. Tatami, T. Megro, S. Matsuo and K. Komeya, “Synthesis of Spherical A1N Particles by Gas-Reduction-Nitridation Method,” J. Eur. Ceram. SOC.,22 52 1-26 (2002). T. Suehiro, J. Tatami, T. Megro, K. Komeya and S. Matsuo, “Aluminum Nitride Fibers Synthesized from Alumina Fibers Using Gas-Reduction-Nitridation Method,” J. Am. Ceram. SOC.,85 [3] 715-17 (2002). T. Suehiro, N. Hirosaki, R. Terao, J. Tatami, T. Megro and K. Komeya, “Synthesis of Aluminum Nitride Nanopowder by Gas-Reduction-Nitridation Method,” J. Am. Ceram. SOC.,86 [6] 1046-48 (2003). lo T.Yamakawa, J.Tatami, K.Komeya and T.Meguro, “Synthesis of A1N Powder from Al(OH)3 by Reduction-Nitridation in a Mixture of NH3-C3Hg,’’ J. Eur. Ceram. SOC.,in press, (2005). l 1 K.J.D. Mackenzie, R.H. Meinhold, G.V. White, C.M. Sheppard and B.L. Sherriff “Carbothermal Formation of P-SiAlON from Kaolinite and Halloysite by 29Si and 27Al Solid State MAS NMR,” J. Mat. Sci. 29 261 1-2619 (1994).

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A Novel Supercritical C 0 2 Synthesis of Amorphous Hydrous Zirconia Nanoparticles, and Their Calcination to Zirconia Mei-Hwa Lee Department of Materials Science and Engineering, I-Shou University, Kaohsiung 840, Taiwan, China

Hung-Yin Lin Department of Chemical and Materials Engineering, National University of Kaohsiung, Kaohsiung 8 11, Taiwan, China

James L. Thomast Department of Physics and Astronomy, University of New Mexico, Albuquerque, New Mexico 87131

A novel synthesis of amorphous hydrous zirconia nanoparticles was performed in a supercriticalcarbon dioxide (scCOz) reverse microemulsion, converting a high concentration of a very inexpensive starting material (zirconyl nitrate hydrate) into a product that was then calcined to yield monoclinic zirconia nanoparticles. The amorphous hydrous zirconia precursor particles were obtained by simply adding a precipitating agent to IZrf.J/perfluoropolyether/scCOz. Calcination converts the amorphous hydrous zirconia precursor into the oxide, and the corresponding phase changes that occur were confirmed by differential thermal analysis. Some control of particle size and shape (ellipticity) could be achieved by selecting the reaction pressure from within the range over which stable microemulsions are obtained (13.9-17.3 MPa): a higher reaction pressure yields smaller and more spherical particles. This novel route for the synthesis of zirconia nanoparticles is both “green” (environmentally friendly) and economical.

I. Introduction UPERCRITICAL CARBON DIOXIDE (scC02) is an attractive replacement for organic solvents in a wide range of chemical processes. It is advantageous for many reasons: it is nontoxic, nonflammable, inexpensive, chemically inert, and has a low viscosity. This last property leads to a high diffusivity for reactants in scCO2, which increases the reaction rates. Moreover, when reactions are carried out in scC02, reaction products can be easily recovered by simply venting the solvent. Finally, the properties of scC02 can be manipulated by temperature and pressure, providing a tunable medium for reactions. Ultra-powders (submicrometer or nanometer-scale particles) with quasi-spherical morphologies and narrow size distributions are useful to enhance the packing density and thus the uniformity of ceramics. In addition, finer ceramic powders can enhance the sinterability, reducing the sintering temperature and giving smaller residual pores. For these reasons, there is great current interest in ceramic ultra-powders. A number of papers have successfully demonstrated that supercritical C02 can replace organic solvents as a hydrophobic phase in microemulsions and reverse microemulsions with water,132(comprehensively reviewed by Liu et uL3 and by Ye and Wai4) Fluorinated rather than hydrogenated surfactants are used to stabilize the microemulsion droplets against van der

S

A. Krell-ontrihuting

editor

Manuscript No. 12540. Received March I, 2006; approved July 25, 2006. ‘Author to whom correspondence should be addressed e-mail [email protected]

Synthesis Methods for Powders

Waals coalescence.596Importantly, microemulsion droplets can serve as effective reaction containment vessels or nanoreactors. This has led to the increasing use of microemulsions for nanoparticle synthesis in the past decade. In general, nanoparticle synthesis by coprecipitation and microemulsion routes requires reducing or precipitating agents. A number of reducing agents have been studied and added to water-in-scC02 microemulsions to prepare various n a n o p a r t i c l e ~ Supercritical .~~~ propane (scC3H8)has also been used: Maston et al. first demonstrated the synthesis in scC3Hxof spherical Al(OH)3 particles with a mean size of 0.5 pm using N ~ B H ( O A Creduction. )~ Monoclinic zirconia powder, partially stabilized with doping such as CaO, MgO, and Y2O3, is used in refractory ceramics to enhance thermal shock resistance and abrasion resistance. These materials are used in demanding thermal applications such as sliding gate plates for pouring steel, insulating fiber, and thermal barrier coatings. In the past several years, fine zirconia particles have been successfully synthesized in reverse macroemulsions,lO,l‘microemi~lsions,’~~’~ and by isothermal decomposition of nitrate precursor^.'^ Although scC02 reverse microemulsions have been successfully used in the synthesis of quasi-spherical semiconductor and metallic nanoparticles, there are very few publications in which an scC02 reaction step has been used in the preparation of metal oxides, such as titanium oxide. I s Previous syntheses of nanoparticles using scC02 have required high pressures (20.2-36.7 MPa)931c’8and have used low concentrations of starting materials (0.33-100 mmole/ ~).7,8,16,17,19In this work, we report on a new, efficient synthesis of nano-sized zirconia (zirconium oxide, Zr02)particles using a water-in-scC02 microemulsion reaction step, using lower pressures and higher concentrations of raw materials than used previously. This new “green” synthesis uses an inexpensive precursor zirconium salt and should prove to be a useful and efficient route to the formation of relatively monodisperse zirconia nanoparticles with spherical morphology for ceramics and other applications. 11. Experimental Procedure

Three milliliters of 1 mol/L aqueous zirconyl nitrate hydrate and 4 g (6.7 mmoles) perfluoropolyether (PFPE; both from Sigma- Aldrich, St. Louis, MO) surfactant were added to a 285 mL stainless-steel high-pressure reaction vessel. The vessel was sealed, and liquid C 0 2 was pumped into the vessel to a pressure of 0.6 MPa, while stirring the reaction mixture with a magnetic stir bar at room temperature. To obtain a supercritical solution, the temperature of the vessel was first heated over ca. 20 min to 67”+ 1°C using a thermostatted water bath. During this heating process, the liquid C 0 2 influx was halted until a pressure of 129

Table I. Effect of Ratio of [Zr?$/[PFPEJ

and Reaction Conditions (Pressure and Temperature) on the Stability of Water/scCOZ Microemulsions

Aqueous solution volume (mL)

Temperature ("C)

Pressure (MPa)

Microemulsion stability

0.65

4.3

3 x 5

No microemulsion formed

0.65

4.3

66-83

8.6 or 12.9

3S83

11.1-17.3 13.9 15.3 17.3 12.2-17.3

[Zr4'(dq)]/[PFPE] molar ratio

1.3 or 1.95

Stable for > 2 h

No microemulsion formed

Boldface indicates conditions used for zirconia synthesis. The aqueous phase was 1M zirconyl nitrate hydrate; the chamber volume is 285 mL.

2.2k0.3 MPa was reached, after which additional liquid C02 was injected into the reactor to further increase the pressure, while keeping the temperature within the stated range. The liquid+vapor CO2 was observed through sapphire windows in the vessel. After ca. 11 min of C 0 2 addition, at a pressure of about 7.6 MPa, the liquid+vapor became supercritical (the critical point of scC02 is at T, = 7.5 MPa and T, = 32°C). Continued addition of liquid C02 increased the pressure at 0.5 MPa/min to achieve experimental pressures of 13.9, 15.3, or 17.3 MPa; the temperature of the reactor reached 70"-72°C. This aqueous/ scC02 microemulsion was stable for > 2 h at pressures above 13.9 MPa at 70°C. For safety reasons, the highest pressure used in the reaction vessel was 17.3 MPa. The precipitating agent, either 0.04 moles of ammonia in a 14 mol/L solution or 3 mL of urea (14 mol/L), was then injected into the vessel via a three port injection. An excess of precipitating agent was used to neutralize the acidic nature of the water-scC02 mixture.20The pressure declined quickly by about 0.3 MPa after the addition. The reaction mixture rapidly turbidified, most likely owing to the rapid formation of nanoparticles within the microemulsion droplets. Over the course of 40 min, the nanoparticles aggregated and/or settled to the bottom of the reaction tank, leaving a clear "supernatant." After 1 h continuous stirring, the vessel was vented. Venting was performed slowly in order to prevent the entrainment of the synthesized powder in the expelled gas. The powder, consisting of a zirconia precursor (amorphous hydrous zirconia),*' was collected from the bottom of the vessel. The powder was repeatedly (>3 x ) washed with 50 mL ethanol and filtered, until the color changed from dark-yellow into pale yellow. Finally, the powder was dried overnight at room temperature (25°C). The zirconium oxide nanoparticles were then synthesized from the amorphous hydrous zirconia precursor particles by calcining the dried powder at 450"C, then at 550"C, and finally at 650"C, for 2 h at each temperature. Thermal analysis of precursor calcination was performed at a heating rate of IO"C/min in air (the same rate and conditions used in the synthesis

9

calcinations above), using thermal ravimetric (TGA) and differential thermal analyses (DTA).' The sizes of the zirconia nanoparticles were characterized by transmission electron microscopy (TEM) analysis, and X-ray diffraction analysis (XRD) was used to determine the crystal structure of Zr02 nanoparticles. Electron Spectroscopy for Chemical Analysis (ESCA) was used to determine the surface elemental composition of the amorphous hydrous zirconia.

III. Results ( 1 ) Microemulsion Stability Table I shows the effect of [Zr4'(aq)]/[PFPE] molar ratio and the reaction conditions (pressure and temperature) on the microemulsion stability of l mol/L aqueous zirconium in scCO2 in our reaction chamber. Temperatures above 65°C were necessary to obtain a stable zirconium hydrate nitrate(aqueous)/PFPE/ scC02microemulsion system at a [Zr4'(aq)]/pFPE] molar ratio of 0.65 (i.e., with 4.3 mL of solution), and stable microemulsions were found at pressures between 13.9 and 17.3 MPa. No stable microemulsions could be obtained at higher ratios of water to surfactant, at any of the temperatures and pressures studied. ( 2 ) Sizes and Morphologies of Amorphous Hydrous Zirconia Precursor Particles and Zirconia Nanoparticles After determining the conditions for microemulsion stability, amorphous hydrous zirconia precursor particles were synthesized by adding ammonia to stable water/ZrO(NO& microemulsions in scCO2. The precursor particles were then calcined in zirconia at S O T , as described above. TEM photographs of typical zirconia nanoparticles calcined from precursors prepared at three different pressures are shown in Fig. 1. Reactions at each pressure were performed at least twice, and different runs gave qualitatively similar results, in terms of particle shape and size. Particles formed from precursors made at the lowest pressure studied, 13.9 MPa, were especially ellipsoidal, with lengths

Fig. 1. Transmission electron microscopy (TEM) micrographs and electron diffraction pattern of zirconia (ZrOz) nanoparticles synthesized in a [Zr4'(aq)]/PFPE/scCOz microemulsion by adding aqueous ammonia, followed by calcination at 550°C. Reaction pressure: (a) 13.9 MPa, (b) 15.3 MPa, and (c) 17.3 MPa.

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Progress in Nanotechnology: Processing

Fig. 2. Long and short axes of zirconia nanoparticles calcined at 550"C, vs. scCOz pressure during precursor synthesis. The trend toward smaller and more spherical particles with increasing synthesis pressure correlates well with changes in microemulsion droplet morphology reported in the literature. The lines shown are "guides to the eye."

Fig. 3. Transmission electron microscopy (TEM) micrograph of precursors and the Zr4' concentration was 1M. The operation conditions are the same as those in Fig. l(a).

of 21-72 nm along the major axis and from 16 to 37 nm along the minor axis. Precursors synthesized at higher pressures gave smaller and more spherical final products, with typical dimensions shrinking to 10-30 nm at 15.3 MPa and to 4-13 nm at 17.3 MPa. Figure 2 summarizes the effect of reaction pressure on the dimensions of the final zirconia nanoparticles. Figure 3 shows a TEM photograph of a sample of the amorphous hydrous zirconia precursor (i.e., before calcination) that was calcined to the zirconia shown in Fig. I(a). The TEM images show that the calcined zirconia particles are slightly smaller than the precursors, as expected. A tabulation of sizes of particles measured and marked in Figs. 1 and 3 is presented in Table 11.

( 3 ) Analysis of Amorphous Hydrous Zirconia Precursor Figure 4 shows the thermal gravimetric/differential thermal analysis (TGA/DTA) of the zirconia precursor particles prepared at 17.3 MPa, using a heating rate of 10"C/min in an air atmosphere. The TGA curve shows two major weight-loss stages. One stage ranges from room temperature to 210°C; there is about a 10% weight loss over this temperature range. In the DTA plot in Fig. 4, there is one broad endothermic peak, at

139"C, and three exothermic peaks at 303", 347", and 449°C. The TGA curve shows that major weight losses were associated with both endothermic and exothermic effects. The endothermic valley over the temperature range from room temperature to 210°C corresponds to the 10% weight decrease in the sample over the same temperature range. The two exothermic peaks at -303" and at 347°C are associated with the decomposition of amorphous hydrous zirconia precursor, or of trapped surfactant or solvent. The third exothermic peak at 449°C indicates the crystallization of tetragonal zirconia, as shown by XRD measurements (vide infra). Zirconia precursor powders synthesized in the supercritical C 0 2 reverse microemulsions were calcined at temperatures of 300", 450", 550", and 650"C, separately. XRD measurements were made at room temperature after slowly cooling each of these powders in air, Fig. 5. The XRD patterns were essentially the same for ZrOl precursors prepared at any of the three pressures studied (13.9, 15.3, and 17.3 MPa). Also evident from Fig. 5: XRD of zirconia precursors did not show any evidence of crystallinity after two hours at the calcination temperature of

Table 11. Sizes of Zirconia Nanoparticles Labeled in Figs. 1 and 2, in Nanometers Fig. I(a)

Fig. 2

No.

Major axis

Minor axis

Fig. l(b)

Fig. l(c)

Major axis

Minor axis

1 2 3 4 5 6 7 8 9 10

63.00 48.00 36.00 46.00 41.00 22.00 41.00 46.00 29.00 51.00

27.00 23.00 29.00 27.00 21.00 20.00 22.00 34.00 27.00 29.00

22.22 11.11 14.81 16.67 29.63 22.22 14.81 18.52 15.74 20.37

34.09 36.36 16.14 29.32 30.68 37.50 24.77 28.86 55.68 25.00

18.18 20.45 13.18 26.59 25.00 25.68 23.86 18.64 32.95 23.64

42.30+ 10.94

25.90k4.13

18.61f4.98

9.26 12.96 9.26 12.96 10.19 6.48 4.63 4.63 14.81 10.74 7.59 8.33 11.11 12.96 5.56 9.43 3.12

3 l.84+9.93

22.82 f5. I9

11

12 13 14 15 Average

Synthesis Methods for Powders

+

131

h

40

P

v

w 20

2

f

E n O

f

e

1L-J-

550'C t

t

t *

I

I

450°C

L

20

c

01

0

i

139°C

$

I

100

I 200

I

I

I

400 500 Temperature ("C) 300

I

600

+

700

Fig. 4. Thermal gravimetric/differential thermal analysis (TGA/DTA) curves of zirconia precursors (17.3 MPa synthesis pressure), calcined in air.

300°C. The tetragonal phase was observed after 2 h at the calcination temperature 450°C (cf. JCPDS card #17-0923) and then the monoclinic phases were found after calcining for 2 h at either 550" or 650°C (cf. JCPDS card #Ol-0750).

( 4 ) Particles Obtained With Urea As the Precipitating Agent At high temperature and pressure, ammonia solution and carbon dioxide can react to form urea. It is not likely that any significant urea was formed in our reactions, however, as urea is not formed at temperatures below 180"C, even at very high C 0 2 Nonetheless, we explored the effects of using urea as a precipitating agent in the synthesis of zirconia precursors, to further rule out the possibility that urea is the active reactant in our synthesis. SEM photographs of zirconia nanoparticles synthesized using urea are shown in Figs. 6(a) and (b). The other reaction conditions were the same as those in Figs. ](a) and (c), respectively. Zr02 particles formed using urea are highly heterogeneous, with sizes from microns to tens of microns, and with irregular shapes. Although urea formation is not an important side reaction in our synthesis, a urea precursor, ammonium carbamate, may be present in the reaction. To determine whether ammonium carbamate or other nitrogen-containing species are present in significant quantities, ESCA was carried out on hydrous zirconia precursors, synthesized using ammonia as the precipitant. The results are shown in Fig. 7. In addition to zirconium and oxygen, fluorine and carbon were found (from the surfactant), but nitrogen was undetectable. ESCA on hydrous zirconia prepared using urea also showed no nitrogen.

IV. Discussion ( I ) Effect of Reaction Pressure on Zirconia Precursors and Zirconia Nanopavticle Products There are many factors, including the pressure and temperature of scC02, solute concentration in aqueous phase, water-tosurfactant molar ratio, and surfactant choice, that can affect the properties of scC02 microemulsions and that would therefore influence the characteristics of the particles synthesized in the microemulsion. In particular, synthesis of uniform nanoparticles requires a stable microemulsion. Microemulsion droplets can take on a variety of sizes and shapes, depending upon the amount of water, surfactant type, and ionic strength, e t ~ In. scCO2 ~ ~ inverse microemulsions, it 132

300°C

E

-20 w

20

I

I

30

40

50 20 (degrees)

60

0

Fig. 5. X-ray diffraction (XRD) analysis of zirconia precursors (either 13.9, 15.3, or 17.3 MPa synthesis pressures)calcined at 300", 450", 550",

and 650"C, bottom to top.

has been shown that decreasing pressure increases droplet size, at a fixed surfactant Concomittant with the increase in droplet size, an increasing deviation from sphericity is expected-nly by forming increasingly elongated droplets can the fixed amount of surfactant be accommodated at the droplet surfaces. These changes in microemulsion droplet morphology parallel the changes that we observed in the final (calcined) particle morphology. As shown in Figs. l(a<), and in Fig. 2, the ellipticities and sizes of zirconia particles after calcining show the same trends with pressure that are expected for the microemulsion droplets. Thus, it appears that droplet morphology determines nanoparticle morphology. We also found almost no differences in the tendencies for ZrO2 nanoparticles synthesized at different pressures to agglomerate, shown in Figs. l(a-c). We suggest that the nucleation rate was very rapid at high pressure, causing a large number of precursor particles to appear suddenly throughout the reaction vessel in a very short time; subsequently, precursors would be easily agglomerated as those particles gradually sedimented to the bottom of the reaction tank. As evidenced in Figs. I(a) and 3, amorphous hydrous zirconia precursor particles synthesized using microemulsion precipitation also exhibit shrinkage on calcination, which has been attributed to loss of surfactant and trapped solvents during calcination.26 Similar losses undoubtedly occur with the amorphous hydrous zirconia precursor formed in the scC02-based synthesis. It is of interest to compare the results from the scCOz synthetic method with a liquid-nitrogen-based synthesis developed ' ~that work, anion exchange was first used by Sinitsky et ~ 1 . In to convert zirconium nitrates into a colloidal solution of hydrated precursors, which were then freeze dried using liquid N2. Calcination at 400°C gave a zirconia product. The focus of that work was primarily on the activation energy for dehydration of the precursors. However, electron micrographs show a much smaller particle size, compared with the synthesis described here, when using the colloidal freeze-drying method.

(2) Characteristics of Amorphous Hydrous Zivconia Precursor Figure 4 depicts the weight loss of the amorphous zirconia precursors during calcination. Part of this weight loss is likely caused by the evaporation of residual water, solvent, and surfactant PFPE (the PFPE boiling point ranges from 160" to 215°C). A second stage of weight loss, from 210" to 380°C, results in a mass decrease of -45%. Wang ef al." studied the calcination of zirconia precursors under atmosphere and found Progress in Nanotechnology: Processing

Fig.6. Scanning electron micrographs of ZrOz powder using urea as the precipitating agent under (a) 13.9 M Pa and (b) 17.3 MPa reaction pressure. The other operating conditions are the same as those in Fig. 1.

a 45% weight loss over the temperature range from 220" to 410"C, which they ascribed to the decomposition of the hydroxide. The similar temperature range and identical fractional weight loss in our TGA lead us to similarly attribute this weight loss to hydroxide decomposition. If this attribution is correct, at 4 10"C, the nanoparticles are an amorphous oxide. The DTA shows three exothermic peaks at 303", 347", and 449°C. Previous studies have shown similar exothermic peaks between 300" and 50O0C,2' showing rapid exothermic reactions when amorphous zirconia precursors are heated at a constant rate in this temperature range. The reactions responsible for these exothermic peaks are not yet identified; however, because TGA shows ongoing weight loss during these exotherms, they may be associated with decomposition of trapped solvent or surfactant. The exotherm at ca. 450°C is the well-established

o.6

,

Amorphous Hydrous Zirconia Pmcursoi

0.0 800

600

400

200

0

Binding Energy (eV) Fig. 7. Electron spectroscopy for chemical analysis survey of the surface of amorphous hydrous zirconia precursor, synthesized using ammonia as a precipitant.

Synthesis Methods for Powders

"glow exotherm" reported in the literature, and is generally associated with the crystallization of z i r ~ o n i a . ~ ~ ~ ~ ' For bulk zirconia, a monoclinic baddeleyite structure is thermodynamically most stable at low temperatures, with a monoclinic to tetragonal phase transition at 1150°C. (A further transition to a cubic phase occurs at 237O0CF7) In contrast, XRD on calcined zirconia nanoparticle precursors showed that the tetragonal phase was formed on crystallization at only 450"C, and the phase persisted on cooling to room temperature. Several explanations have been offered for the unusual phase behavior of the n a n o p a r t i ~ l e s . ~The ~ ~ ~crystal ' formation and transformation behavior of zirconia powders may be affected significantly by various synthesis conditions such as reagent concentration, solvent type, precipitation species, pH and reaction temperature, and thus may not follow the bulk phase transition b e h a ~ i o r ? ' , ~ ' For , ~ ~ example, the pH during precursor synthesis has been shown to play a role in nanoparticle structure, with extremes of pH (below 4 or above 13) giving rise the ~ ;high pH to a metastable tetragonal ZrO2 on ~ a l c i n a t i o n ~ of the synthesis reported here is consistent with that report. The stability of the metastable tetragonal zirconia phase at low temperatures has been attributed to many factors, such as the effect of water vapor,34 a critical particle size effect:' the effect of anionic vacancies, etc.35,36Surface structure and defects are likely to be important, especially for nan~particles.~~ Srinivasan et d3' have suggested that oxygen vacancies present on the surface may be responsible for the low-temperature transformation of the tetragonal phase to the monoclinic phase in the cooling of zironia nanopowders. If so, small amounts of residual reactants (such as surfactant) adsorbed on the nanoparticle surfaces may mask surface defects and thereby inhibit the transformation from t to m on cooling from 450°C. The XRD patterns show that the tetragonal phase is completely converted into the monoclinic one when calcinations are performed at 550" or 650"C, in accord with previous studies on particulate ~ i r c o n i a . ~ 'It . ~is ~ possible that the higher temperature calcinations more completely remove residual surfactant or solvent and restore the low-temperature t -tm phase transformation of zirconia nanoparticles on cooling. 133

(3) Use of Urea As the Precipitating Agent Because ammonia can react with C 0 2 to produce urea, experiments were performed using urea as a precipitating agent, and the resultant nanoparticles were compared with particles synthesized using ammonia. The Zr02 particles formed using urea were highly heterogeneous, with sizes from microns to tens of microns, and with irregular shapes (Figs. 6(a) and (b)). The large differences in particles obtained with urea and ammonia reactants indicate that the precipitating agent is an important factor in the size and morphology of the produced particles; most importantly, these differences demonstrate that urea is not the active precipitant in the ammonia-based synthesis. NH3 and C 0 2 are converted to urea via an ammonium carbamate intermediate. The intermediate could form under our reaction conditions; however, ESCA shows that no significant amounts of ammonium carbamate or other nitrogen-containing species are present in the precursor particles.

5.

V. Conclusions Amorphous hydrous zirconia precursor nanoparticles can be readily synthesized using a water-in-supercritical carbon dioxide reverse microemulsion by adding an aqueous ammonia solution to an inexpensive raw material consisting of an inorganic salt solution. There are many factors that affect the formation of microemulsions with supercritical fluid COP as a hydrophobic phase to replace normal organic solvents. In our work, microemulsion stability was found only at temperatures above 65°C and at pressures between 13.9 and 17.3 MPa. (Higher pressures could not be studied with our apparatus, but may also give stable microemulsions.) Microemulsion droplet size and shape appears to affect nanoparticle size and shape, based on estimates of microemulsion droplet properties from the literature, and on our own measurements of nanoparticle sizes and shapes. As expected, calcined zirconia nanoparticles are slightly smaller than their precursors. The precipitating agent is also an important factor in determining the size and morphology of both precursor and Zr02 particles. This synthesis should prove to be an efficient method for making nanoparticles, because of the ability to manipulate pressure and temperature to control nanoparticle size and morphology. Increasing the precursor reaction pressure from 13.9 to 17.3 MPa caused the size of zirconia particles to decrease from 60 to 10 nm, and led to a more spherical particle morphology, correlating well with expected pressure-induced changes in the microemulsion droplet morphology. As is often observed with nanoparticle synthesis, many primary amorphous hydrous zirconia precursor nanoparticles agglomerated together because of particle interactions during the gradual sedimentation to the bottom of the reaction tank. We observed several exothermic events in thermal analysis of calcinations, including the prominent exotherm at 449°C (the socalled “glow exotherm”) that is characteristic of the formation of tetragonal phase zirconium oxide. Monoclinic phases of zirconia were obtained after higher temperature calcinations, at either 550” or 650°C. Finally, we note that this is the first report of the synthesis of a metal oxide nanoparticle (other than Ti02 using TTIP)using supercritical C 0 2 for a critical synthetic step in which quasispherical or rod-like morphologies could be obtained (depending on reaction pressure). Important practical advantages of the synthesis are the lower reaction pressure (< 20 MPa) and the higher concentration of zirconium salt precursor (>0.1 mol/L) than in previous syntheses.

-

-

Acknowledgments The authors would like to thank the Center for Micro/Nano Technology Research, National Cheng Kung University.

134

References ‘J. Eastoe, 2. Bayazit, S. Martel, D. C. Steytler, and R. K. Heenan, “Droplet Structure in a Water-in-C02 Microemulsion,” Longmuir, 12, 14234 (1996). ’C. T. J. Lee, K. P. Johnston, H. J. Dai, H. D. Cochran, Y. B. Melnichenko, and G. D. Wignall, “Droplet Interactions in Water-in-Carbon Dioxide Microemulsions Near the Critical Point: A Small-Angle Neutron Scattering Study,” J. Phys. Chem. B, 105, 3540-8 (2001). ’J. Liu, Y. Ikushima, and Z. Shervani, “Environmentally Benign Preparation of Metal Nano-Particles by Using Water-in-C02 lcroemulsions Technology.” Curr. Opin. Solid State Mater. Sci., 7, 25541 (2003). 4X. Ye and C. M. Wai, “Making Nanomaterials in Supercritical Fluids: A Review,” J. Chem. Edu., 80, 198-204 (2003). ’J. Eastoe, A. Paul, A. Downer, D. C. Steytler, and E. Rumsey, “Effects of Fluorocarbon Surfactant Chain Structure on Stability of Water-in-Carbon Dioxide Microemulsions. Links Between Aqueous Surface Tension and Microemulsion Stability,” Lnngmuir, 18, 3014-7 (2002). 6J. Eastoe, A. Dupont, and D. C. Steytler, “Fluorinated Surfactants in Supercritical C02,” Curr. Opin. Colloid Interface Sci., 8, 267-73 (2003). ’H. Ohde, F. Hunt. and C. M. Wai, “Synthesis of Silver and Copper Nanoparticles in a Water-in-Supercritical-Carbon Dioxide Microemulsion,” Chem. Muter., 13, 413Ck-5(2001). ‘H. Ohde, M. Ohde, F. Bailey, H. Kim, and C. M. Wai, “Water-in-C02 Microemulsions as Nanoreactors for Synthesizing CdS and ZnS Nanoparticles in Su rcritical C02,” Nano Lett., 2, 7 2 1 4 (2002). W. Matson, J. L. Fulton, and R. D. Smith, “Formation of Fine Particles in Su rcritical Fluid Micelle Systems,” Muter. Lett., 6, 31-3 (1987). K ‘. Kon-no, pp. 3 W in Fine Particles: Synthesis, Characterization, and Mechanisms of Crowfh, Edited by Tadao Sugimoto. Marcel Dekker, New York, NY, 2000. “C. T. Lee, P. A. Psathas, and K. P. Johnston, “Water-in-Carbon Dioxide Emulsions: Formation and Stability,” Langmuir, IS, 6781-91 (1999). I2J. Wang, L. S. Ee, S. C. Ng, C. H. Chew, and L. M. Gan. “Reduced Crystallization Temperature in a Microemulsion-Derived Zirconia Precursor.” Muter. Lett., 30. 119-24 (1997). I’M. H. Lee and C. Y.Tai, “Control of Zirconia Particle Size by Using TwoEmulsion Precipitation Technique,” Chem. Eng. Sci., 56, 2389-98 (2001). 14A.S. Sinitsky, V. A. Ketsko, 1. V. Pentin, G. P. Murav’eva, A. L. ll’inskii,and N. N. Oleynilov, “High-Temperature Dehydration of Hydrophilic Oxides Zr02 and Al2O3.” Russ. J. Inorg. Chem., 48, 406-9 (2003). ”K. T. Lim, H. S. Hwang, M. S. Lee, G . D. Lee, S . 4 . Hong, and K. P. Johnston, “Formation of Ti02 Nanoparticles in Water-in-C02 Microemulsions,” Chem. Commun., 152&9 (2002). 16M.Ji, X. Chen, C. M. Wai, and J. L. Fulton, “Synthesizing and Dispersing Silver Nanoparticles in a Water-in-Supercritical Carbon Dioxide Microemulsion,” J. Am. Chem. Soc., 121,2631-2 (1999). I7H. Ohde, J. M. Rodriguez, X.-R. Yea, and C. M. Wai, “Synthesizing Silver Halide Nanoparticles in Supercritical Carbon Dioxide Utilizing a Water-in-C02 Microemulsion,” Chem. Commun., 23534 (2000). “H. Ohde, C. M. Wai, H. Kim, J. Kim, and M. Ohde. “Hydrogenation of Olefins in Supercritical C 0 2 Catalyzed by Palladium Nanoparticles in a Water-inC 0 2 Microemulsion,” J. Am. Chem. Soc., 124,454Cb-1(2002). I9X. Dong, D. Potter, and C. Erkey, “Synthesis of CuS Nanoparticles in Waterin-Carbon Dioxide Microemulsions,” Indus. Eng. Chem. Res., 41. 4489-93 (2002). 2oK. L. Toews, R. M. Shroll, C. M. Wai. and N. G. Smart, “pH-Defining Equilibrium Between Water and Supercritical C02. Influence on SFE of Organics and Metal Chelates,” Anal. Chem., 67, 404Cb3 (1995). ”G. Stefanic, S. Popovic, and S. Music, “Influence of pH on the Hydrothermal Crystallization Kinetics and Crystal Structure of Zr02,” Thermochim Acta, 303, 31-9 (1997). ”EFMA. “Booklet No. 5 of 8: Production Of Urea and Urea Ammonium Nitrate”; pp. 8-14 in Best Available Techniques. 2nd version, Edited by European Fertilizer Manufacturers’ Association. EFMA, Belgium , 2000. 23Kirk-Othmer, “Amino Resins and Plastics”; pp. 607-21 in Kirk-Othmer Encyclopedia of Chemical Technology, Ffth version, Edited by The Kirk-Othmer Encyclopedia of Chemical Technology. John Wiley & Sons, New York, NY, 1998. 24E. J. Beckman, J. L. Fulton, and R. D. Smith, “Microemulsions in NearCritical and Supercritical Fluids”; pp. 40549 in SupercriticaI FIuid Technology: Reviews in Modern Theory and Applications, Edited by T. J. Bruno, and J. F. Ely. CRC Press, B o a Raton, FL, 1991. 25J. L. Fulton, J. P. Blitz, J. M. Tingey. and R. D. Smith, “Reverse Micelle and Microemulsion Phases in Supercritical Xenon and Ethane-Light-Scattering and Spectroscopic Probe Studies,” J. Phys. Chem., 93. 419&204 (1989). 26M. H. Lee, C. Y. Tai, and C. H. Lu, “Synthesis of Spherical Zirconia by Precipitation Between Two Water/Oil Emulsions,” J . Eur. Ceram. Soc., 19, 2593-603 (1%).

27M. 2. C. Hu, R. D. Hunt, E. A. Payzant, and C. R. Hubbard, “Nanocrystallization and Phase Transformation in Monodispersed Ultrafine Zirconia Particles from Various Homogeneous Precipitation Methods.” J. Am. Ceram. Soc., 825231>20 (1999). B. Djuricic, S. Pickering, D. Mcgarry, P. Glaude, P. Tambuyser, and K. Schuster, “The Properties of Zirconia Powders Produced by Homogeneous Precipitation,” Ceram. lnt., 21, 195-206 (1995). 29 P. E. D. Morgan. “Synthesis of 6-nm Ultrafine Monoclinic Zirconia.” . I Am. . Ceram. Soc., 67, C204-5 (1984). ’OR. C. Garvie, “Stabilization of Tetragonal Structure in Zirconia Microcrystals;” J. Phys. Chem., 82, 218-24 (1978). X. Bokhimi, A. Morales, 0. Novaro, M. Portilla. T. Lopez, F. Tzompantzi, and R. Gomez, “Tetragonal Nanophase Stabilization in Nondoped Sol-Gel Zir-

Progress in Nanotechnology: Processing

conia Prepared with Different Hydrolysis Catalysts,” J. Solid Store Chem., 135, 28-35 (1998).

12J. Denkewicz, P. Raymond. K. S. TenHuisen, and J. H. Adair, “Hydrothermal Crystallization Kinetics of m-ZrOa and t-Zr02,” J. Muter. Res., 5, 2698-705

(1990).

33B. H. Davis, “Effect of pH on Crystal Phase of ZrOz from Solution and Calcined at 600°C.” J. Am. C e r m . Soc.. 67, C-168 (1984). 34Y.Murase and E. Kato. “Role of Water Vapour in Crystallite Growth and Tetragonal-Monoclinic Phase Transformatiom of Zr02,” J. Am. Cerum. Soc., 66, 19G200 (1983).

I5M. I. Osendi, S. S. Moya, J. Serna, and J. Soria, “Metastability of Tetragonal Zirconid Powders,” J . Am. Cerum. Soc., 68, 135-9 (1985).

Synthesis Methods for Powders

16E.Tani, M. Yoshimura, and S. Somia, “Formation of Ultrafine Tetragonal Zr02 Powder Under Hydrothermal Conditions,” J. Am. Cerum. Soc., 66, II4 (1983). ”R. Srinivasan, S. F. Simpson, J. M. Harris, and B. H. Davis, “Discrepancies in the Crystal Structures as Signed to Precipitated Zirconid,” J. Muter. Sci. Lett.. 10, 3 5 2 4 (1991).

”R. Srinivasan, B. H. Davis, and 0. B. Cavin. “Crystallization and Phase Transformation Process in Zirconia: An In Siru High-Temperature X-Ray Diffraction Study,” J. Am. Cerum. Soc., 75, 1217-22 (1992). I9J. R. Combes, L. D. White, and C. P. Tripp, “Chemical Modification of Metal Oxide Surfaces in Supercritical COz: In Situ Infrared Studies of the Adsorption and Reaction of Organosilanes on Silica,” Lungmuir, 15. 7870-5 ( 1999).

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Praseodymium-Doped Photo-Luminescent Strontium lndate Nanoparticles by Ultrasonic Spray Pyrolysis S. E. Lin, K. Borgohain, and W. C. J. Weit Department of Materials Science & Engineering, National Taiwan University, Taiwan 106, ROC

Submicrometer spherical particles of praseodymium (Prkdoped strontium indate were synthesized by spray pyrolysis using ultrasonic nebulization. The morphology, crystalline structure, particle size, and size distribution were studied by X-ray diffraction, scanning, and transmission electron microscopies. The spherical agglomerates were found to be made up of very small nanocrystallites. Upon excitation by UV radiation, sharp luminescence emission bands due to Pr were observed in the blue, green, and red regions in these Prdoped samples. I. Introduction

0

XIDES

doped with rare-earth ions have been extensively studied because of their interesting luminescence properThe rare earth praseodymium (Pr) is a very attractive activator, because it contains several metastable multiplets, which can lead to a variety of emissions right from the blue to the red and infrared wavelength^,^ which is of great importance for luminescent devices, especially in modern display technologies like field-emission displays. Strontium indate (SrIn204)has a wide band gap of 3.6 eV? which makes it possible to incorporate luminescent centers of various rare-earth activator ions. To date, studies on indates as a luminescent host material have been very few in number.536 Rare-earth, Pr-doped indates have the possibility of emitting all the three primary colors of light, which makes it an ideal phosphor for use in modern display panels. We have used the ultrasonic-assisted spray pyrolysis (SP) technique to obtain submicrometer, spherical, luminescent SrIn204 particles doped with Pr. SP is a simple and rapid process for synthesizing submicrometer, spherical agglomerates. It is commonl used for producing a wide variety of materials in powder form,'5 e.g., metal 0xides,8*~non-oxides," ceramics, superconducting materials'' luminescent material~,'~-'~ and so on. The processing steps involve: (1) generation of small-size (usually in sub- or micrometer size droplets from the precursor solution or colloidal suspension1216 by nebulization or ultrasonic atomization; (2) carrying the droplet by flow gas, the atomized droplets evaporate when passing through a temperature-controlled furnace; (3) nucleation and crystallization of solid phase on the droplets due to super-saturation of the solution; and (4) densification of the particles during a transient period at a high temperature as desired. Song et aL8 have reported systematic research on controlling the agglomerate size and size distribution by adjusting precursor concentration, gas flow rate, and furnace temperature. How-

J. Ballato-contributing editor

Manuscript No. 21616. Received March 22,2006; approved May 17,2006. This work was financially supported by the National Science Council (NSC) in Taiwan, under the contract NSC93-2120-M-002-011. 'Author to whom correspondenceshould be addressed. e-mad: [email protected]

Synthesis Methods for Powders

ever, to date, no detailed study on the internal structure of the agglomerates has been carried out. We have made an attempt to analyze the nano-scale structure of the particles and study the photo-luminescent (PL) properties of the Pr-doped SrIn204particles.

11. Experimental Procedure Powder synthesis was carried out by SP using an ultrasonic generator for nebulization of the droplets. III(NO~)~, Sr(N03)2, and Pr(N03)3 solutions were used as the precursors. To study the effect of doping level on the PL properties, samples with three different doping concentrations (5.0%, 0.5%, 0.05°/~)were prepared and denoted as SP-1, SP-2, and SP-3, respectively. The formulation and experimental conditions for the SP process are listed in Table I. The powder samples were characterized by X-ray diffraction (XRD; Philips PW 1710 Instrument, Eindhoren, the Netherlands), scanning electron microscopy (FE-SEM, LEO 1530 Instrument, Cambridge, England), and transmission electron microscopy (TEM 100 CXII, JEOL Co., Akishima, Tokyo, Japan). PL measurements were made under ultraviolet excitation (266 nm) using a frequency-doubled Nd:YAG laser. For XRD measurements, the sample was dispersed in deionized (DI) water and a film was made on a glass substrate. Samples used for SEM were films deposited directly on Si substrates from SP. The SEM was operated at a voltage of 5 kV. For PL measurements, the powder sample was dispersed in DI water and a film was made on an alumina substrate, after which it was calcined at 1000°C for 4 h in a furnace. For TEM observation, samples were prepared by dispersing the particles in 99.5% pure alcohol and placing a drop of the suspension on a carbon-coated TEM copper grid. In order to study the internal structure of the particles by TEM, a cross section of the particles was prepared. A G1 epoxy (Gatan Coronado, CO) was mixed with the sample powder, taking care so as to obtain a homogeneous mixture of the sample particles and the epoxy. The surface tension of the epoxy mixture forms a thin layer on the TEM copper grid, which was then ion milled (PIPS, Gatan) to obtain a nearly transparent film to electron beam. The operating voltage for TEM was 100 kV.

III. Results and Discussion The SEM micrograph of the spherical agglomerates of the asprepared sample SP-2, along with the size distribution obtained by taking into consideration 100 particles from the SEM image, are shown in Fig. 1. The average particle size was about 580 nm. Figure 2 shows the XRD pattern of the as-prepared sample SP-2. The phase of the particles was found to be predominantly that of SrIn204,which crystallizes in the orthorhombic structure (JCPDS File No. 33-1336), but the In203phase was also present in a very minor quantity. The XRD patterns for the different doping concentrations were found to be the same. 137

Table I.

Experimental Conditions for Sample Preparation by Spray Pyrolysis Concentration (mol/L) x

Sample

SP-I SP-2 SP-3

4

6004

Sr2+

pr3+

Furnace temperature ("C)

16 16 16

9.6 9.6 9.6

5 .O 0.5 0.05

800 800 800

Nore: The concentrations of In and Sr ions shown are diluted from the original chemicals obtained from Titanex Cop. (Burgstall, Leiblfing, Germany).

il

2- 300._ 2

Strontium lndate

+ Indium Oxide

ll

Doping (%)

1~3+

.k

1

*

.

5c 200VI

L

-

100-

Centered dark-field (CDF) TEM images of sample SP-2, as prepared and postcalcined at 8OO0C,are shown in Fig. 3. These images reveal that the submicrometer spherical particles are made of small crystallites. The CDF cross-sectional image (Fig. 3(a)) shows that the crystallites are in two size ranges: one is less than 10 nm, and the other ranges from 30 nm to about 65 nm. The smaller crystallites are In203, whereas the large ones are SrInz04. After calcination at 800°C for 4 h, InZO3and SrIn204 crystallites were found to grow to about 17 and > 100 nm, respectively. The sharp diffraction peaks in the XRD pattern would indicate a larger crystallite size than that obtained from TEM results. The small crystallite size shown by the CDF TEM image is most probably that of Inz03, while the large ones belong to SrIn204.The crystallite size of SrIn204being larger than 100 nm (Fig. 3(b)), the electron beam cannot pass through these large crystallites. Hence, the corresponding electron diffraction pat-

Fig. 1. Scanning electron microscopy micrograph and size distribution histogram of as-prepared sample spray pyrolysis-2.

138

t

o

'

I

10

'

I

.

20

I

.

30

I

.

40

I

.

50

I

'

60

I

.

70

1

1

ao

2(e) Fig. 2.

X-ray diffraction spectra of as-prepared sample spray pyrolysis-2.

tern obtained shows a mixture of In203and SrInz04.XRD results, on the other hand, mainly show a SrIn204phase, and very little Inz03phase. It can also be seen from XRD results that the diffraction peaks due to In203 are much broader, indicating a smaller crystallite size, whereas that from SrIn204 are very sharp, indicating a larger crystallite size. For PL measurements, it was noted that the as-prepared, uncalcined samples did not show any PL emission. The samples were therefore postcalcined at 1000°C for 4 h, for better PL efficiency. The sample post-calcined at 800°C for 4 h shows PL emission owing to better crystallinity. PL measurements were carried out under ultraviolet excitation (266 nm) from a frequency-doubled NdYAG laser and detected by a PMT. Figure 4 shows the PL spectra of Pr-doped SrIn204samples with three different doping concentrations 0.05%, 0.5%, and 5.0%. The emission spectra show narrow bands typical of rare-earth ions. The most prominent bands were seen in the blue and red region at 492,605, and 619 nm. Some less intense bands were also seen in the green and red region. The emission features are all due to different transitions from Pr. The blue emission that peaked at 492 nm is due to transition from 3Po-+3H4.The 605 emission line is due to the transition 'D2-+'H4,and the emission at 619 is the 'Po+ 3H6 transition.' Apart from these three main emission bands, the weak band located at 548 nm is due to the transition 3PO+3H5.The emission bands at 654, 691, and 719 nm can be ascribed to transitions from the 3Po+3F2.3,4of Pr. It is to be noted that the feature observed at 532 nm is the second harmonic frequency of the incident laser. The rare-earth elements have partially filled shells off electrons that give rise to narrow localized electronic transitions that occur at wavelengths ranging from the far-infrared to the vacuum-ultraviolet. It is also well known that the emission characteristics of rare-earth elements are relatively weakly influenced by the host material. This is because the 4felectrons of the rareearth ions are highly localized due to shielding by the outer filled shells of 5p and 5s electrons, and hence their optical transitions are atomic-like even when the ion is in a crystalline solid. The intensity of PL emission varied with varying concentrations of Pr. The highest PL intensity was observed in the case of SP-2 with a Pr doping of 0.5%. The PL intensity for SP-3 (0.05% Pr) was much less and it was the weakest for SP-1 (5% Pr). The small PL intensity in the case of SP-3 can be explained by the low doping level of Pr, which leads to fewer Pr ions available for emission. In the case of SP-I, the large number of Pr sites could possibly lead to cross-relaxation between the Pr ions and thus a decrease in the PL intensity." The parameters for PL measurements for the three samples with different doping concentrations were kept identical so that a comparison of the PL intensities could be made. Progress in Nanotechnology: Processing

Fig. 3. Centered dark-field transmission electron microscopy images of sample spray pyrolysis-2 (a) as-prepared in thin section; (b) post-calcined at 800°C for 4 h; (c) shows the diffraction pattern of (a); (d) analysis of the diffraction pattern.

and made up of very small nanocrystallites. CDF TEM images of the agglomerates by TEM show the smallest size of the crystallites to be about 10 nm, which grow to > 100 nm by calcination at 800°C. Sharp, blue, green, and red luminescence bands were observed in these Pr-doped SrInz04 samples. The PL intensity varied with varying concentrations of Pr. The highest PL intensity was observed in the case of samples with a Pr doping of 0.5%.

0.60 0.55 0.50 -

0.45 0.40 4 0.35 5 0.30 L ‘ b j,0.25 -

Q

.-c

5

References

c ; 0.20 -

a

J

0.15 -

0.10 -

0.05,-

0.00 -0.05

300

sp-2 sp-3 sp-1 I

400

.

I

500

.

I

600

.

I

700

.

800

Wavelength (nm)

Fig. 4. Photoluminescence spectra of the praseodymium-doped strontium indate samples with varying doping concentrations.

IV. Conclusion Spherical, luminescent, Pr-doped SrIn204 particles were obtained by an ultrasonic SP process. SP is a simple and effective technique for obtaining phosphor particles at a low cost. The main advantages of this technique are the high purity, spherical shape, and uniform size character of the powder obtained. XRD and electron diffraction results show that SrIn204phase is generated at a furnace temperature of about 800°C. The average powder size as obtained from SEM results was about 580 nm. TEM results show the spherical particles to be polycrystalline, Synthesis Methods for Powders

‘R. Balda, J. Ferna’ndez, I. Sae’z de oca’riz, M. Vcda, A. J. Garci’a, and N. Khaidukov, “Laser Spectroscopy of Pr3+ Ions in LiKYI-,Pr,FS Single Crystals,” Pb s Rev. B, 59, 9972-80 (1999). ‘M. Malinowski, M. F. Joubert, and B. Jacquier, “Dynamics of the IR-to-Blue Wavelength Upconversion in Pr’+-Doped Yttrium Aluminum Garnet and LiYF, Cr stals,” Pbys. Rev. B, SO, 12367-74 (1994). &H. i.Dieke, “Chapter 13: Terbium Tb,” in Spectra and Energy Levels of Rare Earth Ions in Crystals, Edited by H. M. Crosswhite, and H. Crosswhite. Interscience, New York, 1968. 4S. E. Dali, V. V. S. S. Sunder, M. Jayachandra, and M. J. Chockalingan, “Synthesis and Characterization of Aln204 Indates, A = Mg. Ca, Sr, Ba,” 1. Mater. Sci. Lett., 17, 619-23 (1998). ’F. S. Kao and T. M. Chen, “A Study on the Luminescent Properties of RedEmitting Praseodymium-Activated Srlnz04 Phosphon,” 1.Solid Sfare Cbem., 156. 8 4 7 (2001). 6F. S. Kao, “A Study on the Luminescent Properties of New Green-Emitting Srln204:xTbPhosphor,” Muter. Cbem. Pbys., 76. 295-8 (2002).

’S. Jain, D. J. Skamser, and T. T. Kcdas, “Morphology of Single-Component Particles Produced by Spray Pyrolysis,” Aerosol Sci. Tecbnol., 27, 575-90 (1997). *Y. L. Song, S . C. Tsai. C. Y. Chen, T. K. Tesng, C. S. Tasi, J. W. Chen, and Y. D. Yao, “Ultrasonic Spray Pyrolysis for Synthesis Of Sphericdl Zirconia Particles,” J. Am. Ceram. Soc., 87 [lo] 186471 (2004). 9A. Gurav, T. T. Kodas, T. Pluym, and Y. Xiong, “Aerosol Processing of Materials,” Aemyol Sci. Techno/., 19,411-52 (1993). ‘ON. K. Teddy and K. T. R. Reddy, “Growth of Polycrystalline SnS Film by Spray Pyrolysis,” Thin Solid Films, 325, 4-6 (1998). “T. T. Kodas, “Generation of Complex Metal Oxide by Aerosol Processes: Superconducting Ceramic Powders and Films,” Angew. Chem. Int. Ed. Engl. Adv. Muter., 28 [6] 794-806 (1989).

139

I2Y. C. Kdng, H. S. Roh, H. D. Park, and S. B. Park, “Optimization of V U V Characteristics and Morphology of BaMgA110017:Eu2+Phosphor Particles in Spray Pyrolysis,” Ceram. Inf., 29, 41-7 (2003). ”S. M. Abrarov, Sh. U. Yuldashev, B. B. Lee, and T. W. Kang, “Suppression of the Green Photoluiminescence Band in ZnO Embedded into Porous Opal by Spray Pyrolysis,” J . Lwnin., 109, 25-9 (2004). I4Y. C. Kang, H. S. Roh, and S. B. Park, “Preparation of Yz03: Eu Phosphor Particles of Filled Morphology at High Precursor Concentration by Spray Pyrolysis,” Adv. Mater., 12 [6] 451-3 (2000).

140

I5M. Abdullah, F. Iskandar, S. Shibdmoto, T. Ogi, and K. Okuyama, “Preparation of Oxide Particles with Ordered Macropores by Colloidal Templating and Spray Pyrolysis,” Acta Mafer.,52. 5151-6 (2004). 16F. Iskandar, L. Gradon, and K. Okuyama, “Control of the Morphology of Nanostructured Particles Prepared by the Spray Drying of a Nanoparticles Sol,” J . CoNoid Inreface Sci., 265, 29C303 (2003). I7J. Hegarty, D. L. Huber, and Y. M. Yen, “Fluorescence Quenching by Cross Relaxation in LaF3:Pr3+,” Phys. Rev. B, 25, 5638-45 (1982).

Progress in Nanotechnology: Processing

Nano-Blast Synthesis of Nano-size Ce02-Cd203 Powders Oleg Vasylkiv’ ICYS, National Institute for Materials Science 1-1, Namiki, Tsukuba, Ibaraki 305-0044, Japan

Yoshio Sakka MEL. National Institute for Materials Science 1-2-1, Sengen, Tsukuba, Ibaraki 305-0047,Japan

Valeriy V. Skorokhod Institute for Materials Science, NASU, 3, Kiev, 03680, Ukraine

-

)

-

CeC13 7 H z 0 and GdC13 6 H 2 0 that were dissolved in water were precipitated with urea (NHzCONHz) to produce matrix agglomerates for three-component nano-reactors. Mixing hexamethylenetetramine with dilute nitric acid resulted in the formation of well-dispersed nano-particles of cyclotrimetilene & &trinitramine & c ( (RDX) in the solvent. Nano-reactors were produced by impregnating the nano-C&N606 into the matrix agglomerates of an intermediate complex of cerium and gadolinium compounds. Blast initiation of the C&I&O, resulted in extremely rapid detonation and gaseous products formation at temperatures of 2000”-5000”C, which were compressed into a volume nearly equal to the initial volume of each RDX nano-particle. Multiple “nano-blasts” occurred in the volume of each nano-reactor. The impact of the blast waves led to fragmentation of the surrounding matter. The evolution of a large volume of gaseous products dissipated the heat of the process and limited temperature increase, thus reducing the possibility of local sintering among the primary particles. The shortterm high temperature generated during the blasts enhanced the solid solubility of the metal oxides. Uniform aggregates of 22-74 nm consisting of 6-14 nm crystallites of gadolinia in ceria solid solution were synthesized. I. Introduction

M

ULTI-METAL

oxide and composite ceramic nano-size powders enable quality improvement and differentiation of product characteristics at scales currently unachievable with commercially available coarse powders. Fabrication of nano-powders with uniform morphology and precise stoichiometry is the key to realizing high-performance devices based on nano-structured metal oxide ceramics for a wide range of application^.'-^ Particle agglomeration is a natural result of the dominant effect of interparticle forces when the particle size is less than 1 pm. Agglomeration refers to adhesion of particles to each other because of van der Waals forces of attraction, which is significantly larger in nano-particles. If weak forces hold the nano-size particles together, the agglomerates are referred to as soft agglomerates. These agglomerates can be easily redispersed in a suitable liquid medium or in the dry state. In contrast, strong forces due to dissolution-reprecipitation during synthesis and post-synthesis treatment at the contact regions, which form G . Messmg--contributing editor

Manuscript No. 21334. Received January 6, 2006; approved January 23. 2006. Presented at the 9th International Ceramic Processing Science Symposium, Coral Springs, FL, Jan. 8-11.2006. ‘Author to whom correspondence should be addressed. e-mail: oleg.vdsylkiv@nims. wjp

Synthesis Methods for Powders

necks, and (or) subsequent solid necking due to sintering (high-temperature calcination) result in aggregates or hard agglomerates.”16 For these reasons, nano-sized powders readily agglomerate during processing. Once the nano-powder is agglomerated, the strength of the dried agglomerate is too high to realize the benefits of the nano-size primary crystallites. The critical parameter, which had been chosen for controlling and optimization in this study, was the mean secondary aggregate size of the powder after synthesis, post-synthesis treatment, and after calcination. Single- and multi-component nano-powders have been synthesized from aqueous and non-aqueous solutions. A typical procedure involves several sequential steps: (1) preparation of single- or multi-component starting solutions of metal salts (usually aqueous solutions); (2) preparation of aqueous solutions of different precipitants (reductants); (3) reductive decomposition of the starting single- or multi-component solution to obtain the precipitant, colloidal suspension, or gel of the desired end-product phase or intermediate multi-component product; (4) separation of the end (intermediate) product; (5) deagglomeration of the synthesized (precipitated) powder prior to calcination; and (6) synthesis of the end-product powder via calcination, i.e., thermal decomposition of the intermediate products. 1-8*17-24 Processing of multi-component non-agglomerated nano-powders has proven to be extremely challenging and often results in a non-homogeneous multiphase compound with a poor morphol0gy..2-24Typically, nucleation, growth, aggregation, and agglomeration of the first component occur within seconds under very mild conditions. The nucleation of the second component often starts at a higher temperature and requires more time and (very often) different pH. The final product of such “co-precipitation” is a non-homogeneous composite powder, nano-crystalline in nature, but in fact consisting of micrometer-sized agglomerates with very poor morphology and composition h~mogeneity.~ To achieve the desired solid solution, such a multi-component composite powder would require an excessive calcination temperature and unnecessarily prolonged holds with no realistic prospect of achieving a fine-grained s t r ~ c t u r e . ” ” ” ~ ~ The present study was aimed at establishing a new method of preparing nano-size, agglomerate-free, cerium-gadolinium oxide (CGO) ceramic powders with precise morphology and chemical composition. During the past decade, Ce02-based materials have been intensively investigated as catalysts, structural and electronic promoters of heterogeneous catalytic reactions, and oxygen-ion conducting solid electrolytes in electrochemical solid oxide fuel Cyclotrimethylene trinitramine ( R D X b i s a colorless crystalline solid with a density of 1.8 g/cm3. Its structural formula is hexahydro-1,3,5-trinitro1,3,5-triazine or (CH2-NNo2)3.9-’6’2527At temperatures lower than 1IO’C, it is extremely 141

Table I. Characteristics of Cyclotrimethylene Trinitramine (cfi~606)zs-27 Molecule size 0.48 nm Density 1.8 g/cm3 Temperature of thermal detonation 233°C Detonation rate 8,350 m/s Pressure at front of blast wave 33.8 GPa Generated heat during explosion 1.3 kcal/g 0.9 L/g Gas product volume

stable. Decomposition of RDX starts at about 170"C, melting at 2WC, and exploding at 233°C (Table I). Cyclotrimethylene trinitramine is chemically or otherwise energetically unstable. Usually, explosions involve a rapid and violent oxidation reaction with a sudden violent release of mechanical, chemical energy, accompanied by generation of a high temperature and release of extremely hot gases. It causes pressure waves in the local medium in which it occurs.1~1632s27 11. Experimental Procedure

For multiple nano-blast synthesis of CGO, CeCI3. 7H20 and GdC13.6H20 (both 99.9% pure from Wako Pure Chemicals Co., Osaka, Japan) were weighted and separately dissolved in doubly distilled and deionized water at a concentration of 0.1M. The starting materials were used as received without further purification. The initial amount of cerium and gadolinium compounds varied according to the concentrations of both ceramic oxides in the resulting solid solution. To produce complex bi-component intermediate agglomerates, urea (NH2CONH2) (Wako Pure Chemicals Co.) was used as a precipitation agent. NHzCONHz was dissolved in deionized water at a concentration of 2M per I-xM of CeCI3. 7H20 and xM GdCI3.6H20. Two beakers with a total volume of the NH2CONH2aqueous solution of 300 mL (200 mL for cerium chloride and 100 mL for gadolinium chloride solutions) were prepared. Nano-reactors comprising complex threecomponent intermediate agglomerates were synthesized as follows: (I) To produce matrix agglomerates for the threecomponent nano-reactors, nucleation of cerium oxide in the aqueous solution was conducted by spraying 200 mL of urea aqueous solution into the cerium chloride aqueous solution (both solutions were heated to 60°C) under the stirring conditions of 1600 rpm. After precipitation, the suspension was stirred at the prescribed temperature for 10 h. Gadolinium complex was nucleated by spraying gadolinium chloride aqueous solution into a rapidly stirred (I600 rpm) suspension of the as-synthesized cerium oxide. Because of the existence of the residual non-reacted urea, decomposition began within the first minute from the start of the gadolinium chloride solution spraying. After 30 min, 100 mL of urea solution was added by spraying into the stock suspension. Subsequent stirring at 80°C for 5 h was conducted to finalize synthesis and homogenize the suspension. Finally, the product was washed with water, followed by re-dispersion of the agglomerates of cerium and gadolinium intermediate compounds in ethanol (C2H50H, 99.5% reagent grade, Kanto Chemicals, Japan) using an ultrasonic horn (Model USP-600, Shimadzu, Kyoto, Japan). (2) To produce cyclotrimethylene trinitramine (C3H6N606), hexamethylenetetramine was dissolved in deionized water at a concentration of 0.1M. Concentrated (-93%) nitric acid (from Wako Pure Chemicals Co.) was added to urotropin solution. Mixing urotropin with nitric acid led to the formation and re cipitation of cyclotrimethylene trinitramine ( c ~ H ~ N ~ o?5-27 ~ )- . Mixing dissolved hexamethylenetetramine with diluted nitric acid caused the formation of well-dispersed nano-particles of the C&N606 in the Solvent. 142

Fig. 1. TEM micrograph of as-synthesized cyclotrimethylene trinitramine's nano-particles.

(3) Nano-reactors were produced by colloidal impregnation of the C3H6N606 nano-particles (see Fig. 1) into the matrix agglomerates of an intermediate complex of cerium and gadolinium compounds. (4) Subsequently, the threecomponent intermediate agglomerates i.e., nano-reactors, were separated from the supernatant by centrifugation (10 000 rpm for 15 min). The powder was then washed once with distilled and deionized water to remove the residual reaction by-products. Washing with water was followed by two-time ethanol washing with subsequent centrifuging, re-dispersing in fresh ethanol, and then a final slow evaporation of the ethanol ( T = 70°C) using a drying oven. Multiple "nano-blast" calcination/deagglomeration of ceriagadolinia solid solution was conducted by extremely rapid heating of the nano-reactors to the temperature of thermal detonation Of C3H6N606. Cyclotrimethylene trinitramine and nano-reactors were tested for decomposition using thermogravimetric and differential thermal analysis (TG-DTA) (Thermo Plus 2 Rigaku, TG8120, Rigaku, Tokyo, Japan). The reference used for the DTA was an AI2O3sample, the sample containers in the equipment were alumina, and an atmosphere O& was used. Heating rates of 20"C/min were used. The TG-DTA experiments were used to determine the dependence of calcination conditions versus heating rate. The particle-size distribution was analyzed by the dynamic light scattering method (DLS) using a laser particle-size analyzer (Model LSPZ-100, Otsuka Electronics, Osaka, Japan). A very small amount of each powder (15 mg) was dispersed in the distilled water for the analysis. Observation via TEM (Model JEM-2100-F, JEOL, Tokyo, Japan) operated at 200 kV was used to determine the powder morphology. Phase identification of the powders and distribution of the components into each aggregate were determined by nano-area energy dispersion X-ray spectroscopy analyzer (TEM-EDX), and from X-ray diffractometry data (XRD) (Model RINT 2000, Rigaku). The XRD profiles were recorded using CuKu radiation under 40 kV and 300 mA at room temperature.

III. Results and Discussion A general diagram of the processing pathway is shown in Fig. I . A well-dried powder composed of the threecomponent interProgress in Nanotechnology: Processing

Table 11. Powder Aggregale/Agglomerate Size Distribution for Three-Component Intermediate Agglomerates Cerium and Gadolinium Compounds with and without c3H,&o6, as Synthesized after Multiple "Nano-blast" Calcinations/Deagglomeration, and, Finally, after Subsequent Non-Isothermal Calcination up to 450°C with 30 min Holds Composition

Engineered nano-reactors(nm)

After multiple blast treatment (nm)

After calnnation up to 450°C (nm)

28-740 30-1260

1W7

22-74 230-360

GdzoCesoOl95 with C3H6Nb06 GdJ2eanOl0,without C2H,N,0~

mediate agglomerates with a relatively poor non-homogeneous morphology was filled into an alumina container for further thermal treatment. An important factor behind this methodology is to prevent the ignition of the impregnated particles of C3H6N606at 180°C by ultra-rapid heating of nano-reactors through thermal detonation temperature of 233°C (Table I). The thermal detonation temperature is the temperature at which spontaneous multiple ruptures of the N-NOz bonds occur. The blast initiation of cyclotrimethylene trinitramine begins in nanosize regions that are capable of accumulating the heat and transferring it into chemical energy, thus starting the blast reaction. Cleavage of the N-N02 chemical bonds requires less energy for isolated molecules (or clusters) than for molecules located in the s/g) forms bulk of the solid. Extremely rapid detonation ( gaseous products with a temperature of 2000"-5000"C compressed into a volume equaling the initial volume of each cycbtrimethylene trinitramine particle. Multiple nano-explosions start within the volume of each nano-reactor. The instantaneous power of each explosion (i.e., the expansion of compressed gases from the volume equaling the initial volume of each exploded particle of cyclotrimethylene trinitramine) is 500 MW/g.25-27The impacts of the blast waves lead to the fragmentation of the surrounding matter of matrix agglomerates. The rapid evolution of a large volume of gaseous products during combustion dissipates the heat of the process and limits temperature increase, thus reducing the partial sintering among the primary particles. This gas evolution also limits inter-particle contact, resulting in a less-agglomerated product. Multi-blast deagglomeration of the nano-powder occurred due to the highly energetic impact of the blast waves, while the short-term high temperature generated during the explosions enhanced the solid solubility of one component into the other. Nano-size cerium and gadolinium oxide composite and, very soon afterward, a solid solution of the gadolinium oxide in the cerium matrix oxide were synthesized. Utilizing this "nano-blast" calcination

-

-

.

:.

100; 100

700 600

! '. ;

-

0 Time (sec) Fig. 3. Thermogravimetric and differential thermal analysis (TG-DTA) analysis of the thermal detonation and explosion of RDX (heating rate = 20"C/min, dashed-dotted line, temperature; dashed line, DTA; solid line, TG).

technique, we produced a cerium-gadolinium oxide (CGO) powder with an average primary crystallite size of 6 - 14 nm, an aggregate size distribution of 22-74 nm (see Table II), uniform morphology, and precise stoichiometry (Fig. 2). We analyzed the thermal decomposition under rapid heating of both cyclotrimethylene trinitramine itself (Fig. 3) and of cyclotrimethylene trinitramine distributed within the matrix complex precursor agglomerates of cerium and gadolinium compounds (Fig. 4). This enabled us to identify the blast decomposition conditions of the C3H6N606nano-particles synthesized by the described technique and colloidally impregnated into intermediate precursor agglomerates. In addition, this allowed us to confirm them by comparison with the blast decomposition conditions of the C3H6N606itself (Fig. 5 ) The results for TG-DTA of the thermal blast of the C3H6N606 itself at a heating rate of 20"C/min are shown in Fig. 3. Three stages of thermal decomposition could be identified. At the heating rate of 20"C/min, the ignition of the C3H6N606 started at around 180°C. About 202'-205"C is the melting point of cyclotrimetilen trinitramin. A slight endother-

_______..._.------

75

480

Fig. 2. General diagram of the processing pathway.

Synthesis Methods for Powders

-50

520

600 Time (sec)

560

640

150

680

Fig. 4. Thermogravimetric and differential thermal analysis (TG-DTA) analysis of multi-blast calcination of preliminary engineered nano-reactors (heating rate = ZO"C/min, dashed-dotted line, temperature; dashed line, DTA; solid line, TG).

143

Fig. 5. TEM micrograph of porous cerium-gadolinium intermediate

agglomerates.

mic peak could be detected irrespective of the extremely short melting time (2 s for the particles of the C3H6N6O6filled into the container) prior to the multiple blasts. The thermal detonation and multiple blast decomposition of the impregnated C3H6N6O6 particles occurred at -233°C. Just at the beginning of the ignition reaction, the TG analysis detected a significant increase in the mass of the sample (approximately 10.8 weight % at a heating rate of exactly 20”C/min). This is explained by the capturing of external oxygen from the neighboring space by the reacting species. The ignition instantly (within nano-seconds) transformed into thermal detonation, i.e., the C&N606 exploded. Even the temperature as detected by the thermocouple of the TG/DTA system momentarily increased by about 100°C.

Fig. 6. TEM micrograph of cerium-gadolinium oxide nano-aggregates that led to nano-blast synthesis from nano-reactors. 144

Fig.7. TEM micrograph of cerium and gadolinium oxide aggregates calcined with no blast treatment.

Figure 4 shows TG-DTA analysis results for the total summarized explosion of impregnated C3H6N6O6particles during the thermal decomposition of three-component intermediate complex agglomerates i.e., nano-reactors under heating at a rate of 2O”C/min. The reaction results observed here are not, however, limited to this particular rate. The strong exothermal peak detected by differential thermal analysis at a similar temperature and time as for C3H6N606itself (seen in Fig. 3) confirms the Occurrence of multiple blasts of the C3H6N606of nano-particles distributed into the volume of the nano-reactors. Table I1 shows the size distribution of the powders aggregate/ agglomerate for the three-component intermediate agglomerates of cerium and gadolinium compounds with and without C3H6N606,as synthesized, after multiple “nano-blast” calcination/deagglomeration; and, finally, after subsequent non-isothermal calcination up to 450°C. As-synthesized agglomerates with a very wide size distribution of 30-1260 nm were impregnated with separately synthesized particles of cyclotrimetilene trinitramine. Washing of such three-component agglomerates and subsequent ultrasonic deagglomeration reduced the size of the agglomerates to 28-740 nm. Multiple nano-blast calcination reduced the average aggregate size to 1 8 4 7 nm. Nano-aggregates of gadolinia solid solution in ceria matrix synthesized by the multiple “nano-blast” calcinations/deagglomeration technique are shown in the TEM micrograph in Fig. 6. After the nano-blast treatment, the powder was non-isothermally calcined up to 450°C for 30 min to remove the products of explosive decomposition of cyclotrimetilenetrinitramine.

Fig. 8. XRD pattern of ceria-gadolinia solid solution synthesized by multiple nano-blasts of engineered nano-reactors.

Progress in Nanotechnology: Processing

Thus, such treatment acts to preserve both the powder’s compositional homogeneity and morphology. The same calcination treatment was conducted for the agglomerates of cerium and gadolinium oxides that were not embedded with explosive particles, but were preliminarily ultrasonically deagglomerated. From the data listed in Table 11, and TEM micrograph of cerium and gadolinium oxide aggregates calcined without the blast treatment (see Fig. 7), we can conclude that the aggregates of the composite oxide produced were much coarser and the size distribution was wider (230-360 nm). Figure 8 shows XRD patterns of the ceria-gadolinia solid solution produced by multiple nano-blast synthesis. All XRD peaks were attributed to the Gd20Ce8001.95 solid solution. Moreover, these peaks were relatively broad, indicating that the powder was composed of very fine crystallites.

IV. Conclusion The fabrication of nano-powders with uniform morphology and precise stoichiometry is the key to realizing high-performance devices based on nano-structured metal oxide ceramics and metal-ceramic composites for a wide range of applications. Here, we demonstrate a new processing technique that is based on engineering multi-component nano-reactors with subsequent multiple “nano-blast” calcination/deagglomeration. Multiple nano-blasts of impregnated particles of C3H6N6O6deagglomerate the powder due to the highly energetic impacts of the blast waves. The solid solubility of one component into the other is enhanced by the extremely high local temperature generated during the nano-explosions. We produced a nano-size agglomerate-free ceria-gadolinia powder with excellent morphology and an average aggregate size of 48 nm.

Acknowledgments This study was performed through Special Coordination Funds for Promoting Scienceand Technology from the Ministry of Education, Culture, Sports, Science, and Technology of the Japanese Government.

References ’H. Gkiter, “Nanocrystalline Materials: Basic Concept and Microstructure,” Acfa Mafer., 48, 1-29 (2000). ’S. Tjong and H. Chen. “Nanocrystalline Materials and Coatings,” Mafer. Sci. En Res., 45, 1 4 8 (2004). K ‘ . Johnston and P. Shah, “Making Nanoscale Materials with Supercritical Fluids,” Science, 303, 482-3 (2004).

Synthesis Methods for Powders

’T.Zhang, J. Ma, L. Kong. P. Hing, and J. Kilner, “Preparation and Mechanical Properties of Dense Ceo8Gdn.202-a Ceramics,” Solid Stare Ionics, 167. 1 9 1 4 (2~4). 0 . Vasylkiv, T. Kolodiazhni, Y. Sakka, and V. Skorokhcd, “Synthesis and Characterization of Nanosize CeriaGadolinia Powders,” J . Ceram. SOC.Jpn., 113, 1014 (2005). 60.Vasylkiv and Y. Sakka, “Synthesis and Colloidal Processing of Zirconia Nanopowder,” J. Am. Ceram. Soc., 84,2489-94 (2001). ’0. Vasylkiv, Y. Sakka. Y. Maeda, and V. Skorokhod, “Nano-Engineering of Zirconia-Noble Metals Composites.” J. Eur. Cerum. Soc., 24. 469-73 (2004). *N.-L. Wu, S.-Y. Wang, and 1. A. Rusakova, “Inhibition of Crystallite Growth in the Sol-Gel Synthesis of Nanocrystalline Metal Oxides.” Science, 285. 1375-7 (1999). 9G. Ulrich, “Flame Synthesis of Fine Particles,” Chem. Eng. News, 62. 22-9 (1998). “ S . Pratsinis, “Flame Aerosol Synthesis of Ceramic Powders,” Prog. Energy Combust. Sci.,24. 1977219 (1998). “C. Sorensen, W. Hageman, T. Rush, H. Huang, and C. Oh, “Aerogelation in a Flame Soot Aerosol,” Phys. Rev. Left, 80, 1782-5 (1998). ”R. Purohit, B. Sharma, K. Pillai, and A. Tyagi, “Ultrafine Ceria Powders Via Glycine-Nitrate Combustion,” Mafer. Res. Bull., 36, 271 1-21 (2001). ’A. Varma and J.-P. Lebrat, “Combustion Synthesis of Advanced Materials,” Chem Eng. Sci., 41, 2179-% (1992). ‘9.Tillotson, L. Hrubesh, R. Simpson, R. Lee, R. Swansiger, and L. Simpson, “SolCel Processingof Energetic Materials,” J. Non-Crys. Solidr. 225,358-63 (1998). ”T. Tillotson, A. Gash, R. Simpson, J. Hrubesh, L. Jr. Satcher, and J. Poco, “Nanostructured Energetic Materials Using SokGel Methodologies,” J. NonCr s f Solidr. 285, 338-15 (2001). ‘M. Kuklja, “Thermal Decomposition of Solid Cyclotrimethylene Trinitramine,” J . Phys. Chem. B, 105, 1015942 (2001). ”M.-P. Pileni, “The Role of Soft Colloidal Templates in Controlling the Size and Shape of Inorganic Nanocrystals,” Nut. Mater., 2, 145-50 (2000). “J. Millman, K. Bhatt, B. Prevo, and 0. Velev, “Anisotropic Particle Synthesis in Dielectrophoretically Controlled Microdroplet Reactors,” Naf. Mater., 4, 98-102 (2005). I9J.-S.Lee and S.-C. Choi, “Crystallization Behavior of NanoCeria Powders by Hydrothermal Synthesis Using a Mixture of HzOzand NH40H,” Mafer.Left., 58, 39&3 (2004). ”0. Vasylkiv and Y. Sakka, “Synthesis and Sintering of Zirconia Nanopowder by Non-[sothermal Decomposition from Hydroxide.” J . Ceram. SOC.Jpn., 109, 500-5 (2001). ”0. Vasylkiv and Y. Sakka, “Non-Isothermal Synthesis of Yttria-Stabilized Zirconia Nanopowder Through Oxalate Processing: I, Characteristics of (Y-Zr) Oxalate Synthesis and its Decomposition,” J. Am. Ceram. Soc., 83, 2196-202 (2000). 220. Vasylkiv, Y. Sakka, and H. Borodians’ka, “Non-Isothermal Synthesis of Yttria-Stabilized Zirconia Nanopowder Through Oxalate Processing: I, Morpholog Manipulation,” J. Am. Ceram. Soc., 84, 2484-8 (2001). ‘3Z. Tianshu. P. Hing, H. Huang, and J. Kilner. “Ionic Conductivity in the ce02-Gdz0, System (O.OS
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Sol-Gel Processing and Characterization of Phase-Pure Lead Zirconate Titanate Nano-Powders Yasir Faheem+ and M. Shoaib KRL, Rawalpindi 46000, Pakistan The present research describes a modified sol-gel technique used to obtain nano-crystalline lead zirconate titanate (PZT) powders by using triethanolamine (TEA) as a polymerizing agent. The metal ions chemically interacted with TEA in the precursor sol and gel under refluxing conditions. Drying and aging treatments led to development of a precursor-polymeric gel network. A single-phase perovskite structure was formed at 470°C. The PZT gel and powders were characterized by using X-ray diffraction, simultaneous thermal analysis, and Fourier transform infrared spectroscopic techniques. The nano-crystalline morphology was examined by using a transmission electron microscope (TEM). The average particle size determined from TEM was less than 15 nm.

I. Introduction

L

EAD zirconate titanate (PZT) is one of the most widely used piezoelectric ceramic materials because of its excellent piezoelectric properties. PZT-based materials can be found in numerous applications including ultrasonic transducers, hydrophones, speakers, fish finders, sensors and actuators, etc. PZT exhibits a non-centrosymmetric perovskite structure below its Curie temperature (T, = 386°C) and exhibits enhanced piezoelectric properties at the morphotrophic phase boundary (MPB), which occurs at Pb(Zr0.52, Ti0_48)03 composition.' The crystallization behavior of pure PZT can be tailored in many ways by controlling solvent and precursor chemistry, concentration, polymerizing agents as well as proper aging treatments. Important physical properties such as sintered density, chemical homogeneity, and piezoelectric properties are also affected by the final crystallite size.2 The solid-state method is typically used to prepare PZT powders.334Some of the important synthetic routes are hydrothermal, solbgel, and co-precipitation.'-* Particle sizes of the PZT powders, prepared by the above-mentioned methods, are generally in the micrometric to sub-micrometric range. Researchers have focused on the synthesis of nano-crystalline PZT powders to improve certain characteristics such as better sintering at a lower temperature and a higher inherent surface area of the nano-powders. This provides a significantly higher stored energy for solid-state densification, which ultimately results in a lower sintering temperature. This also minimizes the lead loss problem at a high sintering temperat~re.~,'~ The sol-gel method has been extensively used to prepare single-phase nano-crystalline PZT powders and ceramics, thin films, fibers, etc."-I3 In the present research, we report a modified route of PZT nano-powder synthesis from the organic solution of the respective metal ions using triethanolamine (TEA) as the polymerizing agent. TEA also enhances the solubility of J. Nin-ntributing

editor

Manuscript No. 20874. Received August 13, 2005; approved January 3, 2006. 'Author to whom correspondenceshould be addressed. e-mail: [email protected]

Synthesis Methods for Powders

alkoxides and metallic salts such as lead acetate tri-hydrate into alcoholic solutions.

11. Experimental Procedure

Lead acetate trihydrate (Pb(CH3C00)2.3H20; 2 99.5%, Merck, Darmstadt, Germany), zirconium n-butoxide n-butanol complex (Zr(C4H90)d; 80 wt% in n-butanol, Aldrich, Munich, Germany), and titanium isopropoxide (Ti(C,H,O),; 27% Ti, Aldrich) were used as the starting materials to introduce the Pb, Zr, and Ti cations into the solution in a 1:0.52:0.48 ratio, respectively. 2-propanol (CH&CHOH; > 99%, Merck) was introduced as a solvent and stabilizer for zirconium and titanium precursors. Tnethanolamine (HOCH2CH&N; > 98%, Merck) was used to dissolve lead acetate trihydrate in iso-propanol at 80°C under reflux. The processing scheme opted is illustrated in Fig. 1. The PZT sol was maintained at 60"-70°C until a transparent resinous gel was formed. The gel was dried between 100" and 160°C for 10 days continuously. The temperature was increased up to 250"C, and a swollen mass was obtained. This brownish black mass of PZT precursor was ground in an agate mortar to a fine powder and heat treated at various temperatures from 350" to 500°C for 1.5 h to study its crystallizing behavior and other characteristics.

( I ) Characterization The phase chemistry of the PZT gel powder was characterized by X-ray diffraction (XRD; Model JDX-9C, JEOL, Tokyo, Japan) analysis at room temperature, with CuKu radiation and a nickel filter. Differential thermal analysis (DTA) and thermo gravimetric analysis (TGA) were conducted (Model; STA 409C, Netzsch, Selb, Germany) in air atmosphere at a heating rate of S"C/min up to 830°C. The molecular structures of the PZT precursor gel and powders were interpreted by using Fourier transform infrared spectroscopy (FTIR; Model 470-FTIR Nexus, Thermo Nicolet, Madison, WI), in the wave range from 4000 to 500 cm-' and a resolution of 8 cm-I. The particle size and grain morphology were evaluated by using a transmission electron microscope (TEM; Model CX200, JEOL) working at an accelerating voltage of 200 kV. The TEM sample was prepared by dispersing a small drop of the ultrasonicated suspension onto a copper grid precoated with an amorphous carbon film.

1II. Results and Discussion

( I ) XRD Analysis The XRD results of PZT gel powders processed at various temperatures are shown in Fig. 2. The precursor gel exhibited amorphous behavior below 25OoC, which indicates that a polymeric network exists holding the metal ions together. The decomposition of lead acetate started above 300°C and lead oxide was observed to exist in a free state at 35OoC,as shown by the XRD plot. At 430"C, several peaks from perovskite lattice planes were observed, which were accompanied by two small pyrochlore 147

I

I

ZI(OC~H~ complex )~ dissolved in 2-propanol solution RefluxinglNp at 60°C

I

Ti (OC3H7)4 dissolved in 2-propanol solution

I

Refluxing/6O0C Pb(CH3C00)2.3H20dissolved in

Refluxing/BO"C

PZT precursor sol

I

I

Aging and drying

1

I

I

Gelation

1

Calcinations

PZT perovskite nano-powder Fig. 1. Scheme for the preparation of lead zirconate titanate (PZT) nano-powders.

peaks. Thus, it may be concluded that the perovskite phase started to form at 430°C. The XRD plot of the gel powder treated at 500°C identified a single-phase perovskite. But, in order to investigate the exact perovskite formation temperature, further experimentation was carried out by using the FTIR and DSC/TG analytical techniques, which also helped to study the nature of reactions taking place at various temperatures. The XRD pattern of PZT gel powder heat treated at 500°C exhibited broad peaks indicative of fine crystalline particles. The PZT average crystallite size was

A

Perovskite

+ Pyrochlore

500°C c

.-2 -ma,

430°C

c

a

*

250°C I

20

30

I

I

40 50 2theta "degrees"

I

60

70

Fig. 2. X-ray diffraction plots of lead zirconate titanate gel powders calcined at various temperatures. 148

T = 0.9h/Bc ose ~ where T is the average particle size in angstroms, B is the width of the peak at half the peak height in radians, his the wavelength in angstroms, and OB is the Bragg angle in degrees. Das et al.14 and Bose and BanerjeeIs have produced nano PZT powders at 450°C by using inorganic metal precursors. ' ~ reported nano-crystalline PZT perovskite C. Liu et ~ 1 . have formation at 650°C by precipitating the PZT recursor powder from an organic solution, and D. Liu et d.' have produced nano-PZT powders at 450°C by using an ethylene glycol solvent system. The present paper illustrates nano-crystalline phasepure PZT perovskite formation at 470°C by using organic metal precursors and a TEA/iso-propanol solvent system following a specific process, which has not been reported before. The lowtemperature perovskite crystallization can be attributed to the decomposition and combustion reactions associated with the organics and TEA complexes.

(2)

350°C

[r

calculated to be approximately 14.5 nm by using Scherrer's equation

FTIR Analysis

Figure 3 shows the FTIR spectra of the PZT gel and powder heat treated at various temperatures at a rate of 5"C/min. Many absorption bands related to the acetate anions were clearly observed, i.e. the 1407 and 1550 cm-' bands from 80°C gel may be attributed to the symmetric and asymmetric (COO-) vibrations, respectively, while the medium bands at 1339 and 910 cm-l may be attributed to the deformation vibration of CH, groups and the stretching mode of C-C bonds, respectively. The medium Progress in Nanotechnology: Processing

-=!

t

10 Ot-\

910

3318 1550

v

-0 a

2400°C

A f

lii

t

a

1076

I-

n

m +: ._

-'"

5E

!-

-60 -50

0

I 4000

I

3500

I

3000

I

I

I

2500 2000 1500 Wavenumber (cm-')

I

1000

I 1

500

Fig. 3. The Fourier transform infrared spectra of lead zirconate titanate (PZT) gel and powders calcined at different temperatures.

and weak intensity peaks at 1016, 1157, and 1260 cm-' may be attributed to the stretching vibration of the C-N bonds of TEA. A sharp band at 1076 cm-' was assigned to the C-0 stretchin mode of iso-propanol. Two weak bands at 2870 and 2964 cm- $ may be attributed to the CH3 asymmetrical stretching and C-H stretching from various methyl groups present among the precursors and solvents, respectively. The broad shallow band at 3318 cm-', approximately, may be attributed to the OH- group (hydrogen bond) present in isopropyl alcohol, TEA, and chemically combined water in lead acetate.I8 As predicted by FTIR data, heat treatment at a higher temperature was required to decompose the acetate and other organic ligands present in the precursor gel. The gel powder heat treated at 250°C exhibited peaks from acetate ligands, with lower amplitudes indicative of their partial decomposition. The temperature was raised further to remove the residual carbon associated with the acetate, methyl, and amine ligands. Above 400°C the PZT gel powder exhibited no association of acetate ligands which indicated that they were decomposed completely and residual carbon, if any, was oxidized. Traces of 1407 and 1550 cm-' bands were observed up to 430°C, and they disappeared completely at 500°C. A hump at 980 cm-l was observed in the spectrum of 430" and 500°C powders, which is an indication of formation of a new structure, i.e perovskite. The FTIR data exhibited perovskite crystallization onset at 430°C, which continued up to 500°C. The actual crystallization temperature of perovskite was determined from TG/DTA analysis.

( 3 ) TGIDTA Thermal Analysis PZT gel prepared at 80°C was characterized for thermal behavior. The TG/DTA results are shown in Fig. 4. The DTA curve had two main exothermic peaks at 334" and 422"C, respectively, and some small exothermic peaks associated with these. The TG curve had three major weight loss steps at 240", 320", and 350"C, respectively. The DTA curve indicates that the PZT precursor gel decomposed exothermally with a sharp peak at 334°C. This exotherm may be assigned to the major decomposition reaction associated with TEA complexes. Some minor peaks observed just before 334°C peak may be associated with the burning and oxidation of propoxy, butoxy, and acetate ligands associated with metal ions. The formation of free lead oxide occured in this temperature range, as can be observed in the XRD plot. Comparison with FTIR results indicated that the intensity of the peaks from TEA ligands decreased as the temperature increased and at 350°C, the peaks disappeared. Therefore, the weight loss associated with the exotherm at 334°C indicates the TEA decomposition reaction. These decomposition reactions were accompanied by a drastic weight loss in the temperature range between 240" and 320°C. Synthesis Methods for Powders

200

400

600

800

Temperature /"C Fig. 4. Thermo gravimetric (TG)/differential thermal analysis (DTA) of lead zirconate titanate resinous gel prepared at 80°C.

Weight loss continued from 320" to 422"C, which was accompanied by another exothermic peak on the DTA curve. The exotherm may be attributed to the crystallization of the pyrochlore phase, as evidenced by the previous XRD results. The weight loss during this stage resulted from the decomposition reactions of residual organics. Substantial heat energy was released by these reactions and the crystallization of the perovskite phase was finally observed by a small exothermic shoulder close to 470°C on the DTA c ~ r v e . ~ Finally, ~ . * ~ at the completion of decomposition reactions and PZT perovskite formation, no further weight loss or chemical activity was observed.

( 4 ) TEM Observations The PZT gel powder heat treated at 500°C was subjected to TEM microstructural analysis to determine the average particle size. Figure 5 shows the TEM image of the as-dispersed PZT nano-powder. The image reveals that the nano-particles were agglomerated at some locations with an average particle size of less than 15 nm. This is also in good agreement with the XRD line-broadening results (14.5 nm). The individual particles exhibited a spherical morphology. The selected area electron diffraction (SAED) confirmed the crystallinity of the as-obtained PZT nano-particles.

IV. Conclusion Nano-crystalline (< 15 nm) PZT powders have been synthesized successfully by using sol-gel processing. The formation of PZT

Fig. 5. Transmission electron microscope image of lead zirconate titanate nano-powder heat treated at 500°C. Inset is the selected area diffraction pattern.

149

nano-particles was discussed, interpreted, and verified by using FTIR along with XRD and TG/DTA analytical techniques altogether by adopting a specified procedure. Exothermic decomposition of the organics and TEA complexes resulted in the evolution of gases (NH3, C02, H20) and substantial heat energy, which crystallized a single-phase perovskite at 470°C. The PZT particle size was measured by using TEM and compared with XRD line-broadening results. The as-obtained nano-powders may be suitable for lower temperature sintering and to produce a denser bulk material. The volatility of lead oxide that occurs at higher sintering temperatures may be minimized, and enhanced physical and piezoelectric properties may be obtained as well. Acknowledgment The TEM investigation work conducted by Dr. M. Farooque is gratefully acknowledged.

References ‘B. Jaffe, R. S. Roth, and S. Marzullo, “Piezoelectric Properties of Lead Zirconate Titanate Solid-Solution Ceramics,” J. Appl. Phys., 25 [6] 809-10 (1954). ’W. Aiying, P. M. Vilarinho, I. M. M. Salvado, and J. L. Baptista, “SolGel Preparation of Lead Zirconate Titanate Powders and Ceramics: Effect of Alkoxide Stabilizers and Lead Precursors,” J . Am. Ceram. Soc., 83 [6] 137%85 (2000). ’Y. Matsuo and H. Sasaki, “Formation of Lead Zirconate-Lead Titanate Solid Solutions,” J. Am. Cerum. Soc., 48, 289-91 (1965). 4S. S. Chandrateriya, R. M. Fulrath, and J. A. Pask, “Reaction Mechanisms in the Formation of PZT Solid Solutions,” J. Am. Ceram. Soc., 64 171422-5 (1981). ’H.Hirashima, H. Onishi, and M. Nagakowa, “Preparation of PZT Powders from Metal Alkoxides,” J. Non-Crysr. Solids., 121, 4 0 4 4 (1990). 6Q. F. Zhou, H. L. W. Chan, and C. L. Choy, “Nanocrystalline Powders and Fibers of Lead Zirconate Titanate Prepared by the SolGel Process,” J. Mufer. Process. Techno!.,63, 281-5 (1997).

150

’H.M. Cheng, J. M. Ma, B. Zhu, and Y. H. Cui. “Reaction Mechanisms in the Formation of Lead Zirconate Titanate Solid Solutions Under Hydrothermal Conditions,” J. Am. Ceram. Soc., 76 (31 625-9 (1993). *A. M. Bruno and J. A. Eiras, “Preparation of Coprecipitated Ferroelectric Ceramic Powders by Two-Stage Calcinations,” J. Am. Cerum. Soc.. 76 [I I] 27366 (1993). 9W. D. Kingery, H.K. Bowen, and D. R. Uhlmann, In Introduction ro Cerumics, 2nd ed, pp. 469-77. Wiley, New York, 1976. ‘OJ. Ryu, J. J. Choim, and H. E. Kim, “Effect of Heating Rate on the Sintering Behavior and the Piezoelectric Properties of Lead Zirconate Titanate Ceramics,” J. Am. Ceram. Soc., 84 [4] 9 0 2 4 (2001). “I. M. Aiying Wu,M. Salvado, P. M. Vilarinho, and J. L. Baptista, “Lead Zirconate Titanate Prepared from Different Zirconium and Titanium Precursors by Sol-Gel,” J. Am. Ceram. Soc., 81 [lo] 2 6 4 M (1998). ”G. Yi, Z. Wu, and M. Sayer, “Preparation of Ph(Zr,Ti)O, Thin Films by SolGel Processing: Electrical, Optical and Electro-Optic Properties,” J. Appl. Phys., 64 5 2717-24(1988). Jr. Meyer, T. Shrout, and S. Yoshikawa, “Lead Zirconate Titanate Fine Fibers Derived from Alkoxide-Based Sol-Gel Technology,” J. Am. Cerum. Soc., 81 [4] 861-8 (1998). I4R. N. Das, A. Pathak, and P. Pramanik, “Low-Temperature Preparation of Nanocrystalline Lead Zirconate Titanate and Lead Lanthanum Zirconate Titanate Powders Using Triethanolamine,” J. Am. Ceram. Soc., 81 [I21 3357-60 (1998). ”S. Bose and A. Banerjee, “Novel Synthesis Route to Make Nanocrystalline Lead Zirconate Titanate Powder,” J . Am. Ceram. Soc., 87 [3] 487-9 (2004). I6C. Liu, B. Zou, A. J. Rondinone, and Z. J. Zhang, “SolGel Synthesis of FreeStanding Ferroelectric Lead Zirconate Titanate Nanoparticles,” J. Am. Chem. so;? 173,43445 (2001). D. Liu, H. Zhang, W Cai. X. Wu, and L. Zhao, “Synthesis of PZT Nanocrystalline Powder by a Modified SolGel Process Using Zirconium Oxynitrate as Zirconium Source,” J. Mater. Chem. Phys., 51, 186-9 (1997). ‘*K. Kitaoka, H. Kozuka, and T. Yoko, “Prepartion of Lead Lanthanum Zirconate Titanate Fibers by SolGel Method,” J. Am. Ceram. Soc., 81 [5] 1189-96 (1998). I9A.Towata, H. J. Hwang, M. Yasuoka, and M. Sando, “Seeding Effects on the Crystallization and Microstructure of S o l 4 e l Derived PZT Fibers,” J. Mater. Sci.,35, 4009-13 (2000). ’k.Kitaoka, H. Kozuka, and T. Yoko, “Prepartion of Lead Lanthanum Zirconate Titanate Fibers by SolGel Method,” J. Am. Ceram. Sac., 81 [5] 1189-96 0 (1998).

“2.

Progress in Nanotechnology: Processing

Synthesis of AIN Nanopowder from y-AI2O3 by Reduction-Nitridation in a Mixture of NH3-C3H8 Tomohiro Yamakawa,+ Junichi Tatami, Toru Wakihara, Katsutoshi Komeya, and Takeshi Meguro Graduated School of Environment and Information Sciences, Yokohama National University, Yokohama 240-8501, Japan

Kenneth J. D. MacKenzie School of Chemical and Physical Sciences, Victoria University of Wellington, Wellington, New Zealand

Shinichi Takagi and Masahiro Yokouchi Kanagawa Industrial Technology Research Institute, Evina 243-0435, Japan

Aluminum nitride ( A N ) powders were synthesized by gas reduction-nitridation of y-Al2O3 using NH3 and C3Hs as the reactant gases. A N was identified in the products synthesized at 1100"-1400"C for 120 min in the NH&&18 gas flow confirming that AIN can be formed by the gas reduction-nitridation of y-A1203. The products synthesized at 1100°C for 120 min contained unreacted y-A120J. The 27Al MAS NMR spectra show that AI-N bonding in the product increases with increasing reaction temperature, the tetrahedral A104 resonance decreasing prior to the disappearance of the octahedral A106 resonance. This suggests that the tetrahedral A104 sites of the y-A1203 are preferentially nitrided than the A106 sites. A N nanoparticles were directly formed from y-AI203 at low temperature because of this preferred nitridation of A104 sites in the reactant. A N nanoparticles are formed by gas reduction-nitridation of yA1203 not only because the reaction temperature is sufficiently low to restrict grain growth, but also because y-A1203 contains both A104 and A106 sites, by contrast with a-A1203 which contains only A106.

A

I. Introduction

In recent years, the synthesis of AIN nanoparticles has active1 been pursued because of their properties. Plasma syndirect nitridation using NH3 gas?0321vapor phase synthesis,22323 and electron beam heating24have been proposed. More recently, Suehiro et ui.2s,26used gas reduction-nitridation in the AI2O3-NH3-C3H8system to synthesize AIN particles at lower nitridation temperatures which are thermodynamically advantageous compared with the AI2O3-C-N2 system. It has also been reported that AlN nanoparticles can be synthesized from 6, y-A1203.27The formation mechanism of AIN from transition alumina has not yet been elucidated. Studies on the synthesis mechanism of AIN from A1203have been carried out, based on phase analysis and microstructural observations using X-ray diffraction (XRD), transmission electron microscopy (TEM), and scanning electron microscopy (SEM). 27Al MAS NMR has also been used to investigate the formation mechanism of oxynitride ceramics such as A10N28*29 and SiAION.3@-33 NMR spectroscopy should also provide useful information on the formation of AIN by gas reduction-nitridation of y-A1203.The purpose of the present work is to investigate the formation mechanism of nano AlN particles from a transition alumina (y-A1203), by gas reduction-nitridation using several analytical techniques including 27AlMAS NMR.

L U M I N U M NITRIDE (AIN) ceramics

have attracted considerable attention as IC substrates, packages, heat-sinks, and fillers, because of their high intrinsic thermal conductivity (330 W/mK), high electric insulation (> lOI4 R . cm), and low thermal ex ansion coefficient (3.2 x 10-6/K), which is close to that of Si?,' Several investigations have been conducted on the synthesis of AIN by various techni ues, including direct n i t r i d a t i ~ n , ~ , ~ chemical va or deposition,' and carbothermal reduction-nitridation.8-1aIn particular, A1203-C-N2 system has found industrial application, as the morphology of the products can be controlled because of the endothermic reaction. Furthermore, it is well known that the AIN powder synthesized by this technique is easy to be densified the sintered body. However, a longer reaction period and higher temperature is needed to fully nitride aA1203.In order to synthesize AIN at lower temperatures and shorter times, several researchers have investigated the use of various A1203 polymorphs as starting materials. They reported that transition aluminas are easier to be nitrided than aA1203.'2-'4 L. Klein--contributing editor

Manuscript No. 20406. Received April 12, 2005; approved July 19, 2005. 'Author to whom correspondenceshould be addressed. e-mail: [email protected]

Synthesis Methods for Powders

11. Experimental Procedure

Commercial nanocrystalline y-A1203powder (AKPGOI 5, Sumitom0 Chem. Co., Tokyo, Japan) was used as the starting material. The main characteristics of the raw powder are summarized in Table I. The raw AI2O3powder was weighed into an AI2O3 boat, placed in an electric furnace with a high-purity AI2O3work tube, and fired to 700°C at a heating rate of 5"C/ min in Ar gas (99.999% purity) to eliminate oxygen in the system and remove the surface water of the y-A1203.We initially confirmed that the characteristics of the y-AI2O3did not change during the pre-heating treatment up to 700°C. Heating was continued in a flowing gas mixture (4 Ljmin) of NH3 (99.999% purity) and 0.5 vol% C3Hs (99.99% purity). The sample was heated to the reaction temperature of 1100°-1400"C at a rate of 8"C/min and held for &120 min before being cooled in NH3. After the heating, the sample was removed the reactor in shortterm contact with ambient air, it was vacuum-encapsulated the samples. The phases present in the products were identified by X-ray diffractometry (RINT2500, Rigaku, Tokyo, Japan) using CuKa radiation operated at 50 kV and 300 mA. Their morphologies were observed with a transmission electron microscope (JEM2000FX, JEOL, Tokyo, Japan). The specific surface areas of the 151

Table 1. Characteristics of the Raw Powder Characteristics

Value

Purity (%) Crystalline phase Specific surface are (m2/g) XRD crvstallite size (nm) Transition temperature to C L - A I ~("C) O~ \

I

1281'

+Calculated from a diffraction peak y-A1203 (440). 'Determined by differential thermal analysis.

powders were measured by the single-point of Brunauer-EmattTeller (BET) method (Quantasorb, Quantachrome, Boynton Beach, FL). The nitridation ratio was calculated from weight change by oxidation of the resultant powder in air with therm~ g r a v i m e t r y(TGA; ~ ~ TG-8 120, Rigaku, Tokyo, Japan). Oxygen contents were measured by OxygenjNitrogen analyzer based on infrared absorbance technique (EGMA-650, HORIBA, Tokyo, Japan) after correcting calibration curve of standard samples. The solid-state 27Al MAS NMR spectra (Unity 500 spectrometer, Varian, Palo Alto, CA) were acquired at 11.7 T, operating at 130.244 MHz and a 5 mm Doty MAS probe with a spinning speed of 1@12 kHz. A 1 ps x/10pulse for solution was used with a recycle delay time of 1 s, and the spectra were referenced to Al(HzO);+.

IIl. Results and Discussion Figure 1 presents the XRD patterns of the products synthesized at 1100"-1400"C for 120 min. The products formed at 1100°C consist chiefly of AIN and residual unreacted y-A1203. Singlephase A1N can be obtained over 1200°C. Figure 2 shows the XRD patterns of the products synthesized at 1400°C without holding the sample at the maximum temperature, in Ar and NH3-0.5 vol% C3H8 atmospheres. Although most of y-AI203 fired in Ar was transformed into a-A1203 because of a high temperature, u-A1203was not identified in the sample fired at the NH3-C3Hg gas mixture. It was suggested that )'-A1203 should be directly converted to AIN. The nitridation behavior of y-Al O3 is consistent with that reported in a previous study,265 however, its detailed mechanism are, as yet, unexplained. Figure 3 shows the temperature dependence of the nitridation ratio at 1100"-14OO"C for 120 min. Nitridation ratio significantly

-p

0 AlN

0

Fig. 2. XRD pattern of the products synthesized at 1400°C (a) NH3C3H8 gas mixtures and (b) Ar atmosphere.

increased as the reaction temperature increased, consistent with the XRD results. A nitridation ratio of 80% was achieved by firing at 1200°C. Even in samples fired at high temperatures for long soaking times, complete nitridation was not achieved, possibly because of the formation of aluminum hydroxide species by reaction of AIN with high surface area and the moisture in the air. However, the presence of such thin surface oxide layer on the AIN particles could not be confirmed by XRD?7,36 Figure 4 shows TEM photographs of the products synthesized at 1200"-140O0C for 120 min, together a photograph of the initial material (Fig. 5(a)). All the fired products consist of nanoparticles, but in the sample synthesized at 1400"C, these nanoparticles are seen to have sintered during primary particle formation and coalescence (Fig. 5(c)). Figure 5 shows the specific surface area and BET average particle size of the product produced by firing at 1100"-1400"C for 120 min. The size of the nanoparticles increases with increasing temperature. The specific surface area was estimated to be 71.8-24.2 m2/g, the average BET particle size being 2G76 nm, an appreciable change by comparison with the initial powder. This result is in good agreement with the TEM observation.

2o 10

w .

50

2O/ifegree

70

80

Fig.1. XRD pattern of the products synthesized at (a) llOO°C, (b) 1200"C, (c) 1300°C, and (d) 1400°C for 120 min.

152

70

2Wdegrce

t

01 1100

I

I

1200 1300 Temperature ("C)

1

Fig. 3. Nitridation ratio of the products obtained from firing at various temperatures for 120 min.

Progress in Nanotechnology: Processing

Fig. 5. Specific surface area and Brunauer-Ematt-Teller particle size of the products synthesized at various temperatures for 120 min.

Table 11. Oxygen Contents and Lattice Constants of AW Nanopowder Lattice constants Reaction conditions

140OoC-120 min 1400T-60 rnin 14OOoC-120 rnin

Oxygen contents ( ~ 1 % )

2.8 3.2 5.7

a

(A)

3.1110 3.1 110 3.1 109

c

(4

4.9794 4.9792 4.9730

than that prepared in our previous and are comparable with commercial grade of AIN powder. Figure 6 shows the 27Al MAS NMR spectra of samples reacted at 1100"-14OO0C for 120 min, together with the spectrum of the initial y-A1203,which shows the typical tetrahedral A104 resonance at 66 ppm and the octahedral AIOd resonance at 12 pprn (Fig. 4(a)). As nitridation proceeds, the AI-N resonance at 114 ppm appears and increases in intensity. The broadening and asymmetry of this peak to higher fields in the spectra of the products formed at 1100" and 1200°C suggests the formation of a transient aluminum oxynitride species which typically appears at 80-100 ppm, depending on the N:O This result suggests that nitridation proceeds via these intermediate

Fig.4. TEM photographs of (a) the raw material and the products synthesized at (b) 1200°C and (c) 1400°Cfor 120 min.

Oxygen contents and lattice constants are summarized in Table 11. AIN nanoparticles synthesized at 1200°C for 120 rnin and 1400°C for 60 and 120 rnin contain 5.6, 3.2, and 2.8 wt% oxygen contents, respectively. It was found that AIN lattice constants of c-axes were increased along with the synthesis period. It is well known that a change in c-axis lattice parameter is caused by dissolution of oxygen in A1N.37Thus, these results suggest that residual oxygen in AIN nanopowder is reduced by reduction nitridation. XRD analysis shows the presence of AIN only, with no y-A1203detectable in samples fired at 1200°C for 120 min, even though the nitridation ratio indicates that oxygen still remains in these samples. The oxygen contents in samples Prepared at 1400°C for 60-120 min are mainly from the surface oxygen layer because of lattice contents were almost the same. However, resultant AIN possessed lower low oxygen contents

Synthesis Methods for Powders

(4 so 0 -60 Chemical shift (pprn) w r t Al(H20)63t

100

-100

*'A1 MAS NMR spectra of (a) the raw material and the products synthesized at (b) 110o"C, (c) 1200"C, and (d) 1400°C (*denotes

Fig.6.

spinning sidebands).

153

AILCLN products. The tetrahedral A104 sites of the y-alumina appear to be more susceptible to nitridation than the octahedral sites, which persist in the sample fired at 1200°C even after the tetrahedral resonance has disappeared. All the AIL0 resonances claim on have given way to the AIN peak by 1400°C. Min et the basis of quantum mechanical modeling that structural models for both y-A1203 and AlON with vacancies located at the tetrahedral sites are more unstable than those with vacancies located at the octahedral sites. Ydmaguchi and Yanagida4’ have reported that spinel phases such as y-AI2O3are stabilized not by AI3+ but by nitrogen. Thus, it appears that the A104 in the spinel are preferentially converted to AIN4, thereby stabilizing the tetrahedral sites and preventing their thermal conversion to A106 (the u-Al2O3structure contains only octahedral Al). It is well known that the bond length of A 1 4 in the A104 site is shorter that that of A106 site.42743Moreover, the covalent bond of AI-N is stronger than that of A 1 4 in A106 or A104. However, previous research on defective spinel structure model of y-A1203 utilizing the several simulations has demonstrated that cation vacancies are preferentially located at the octahedral Beside, nitrogen atoms preferentially replace oxygen atoms in A104 tetrahedron!’ These results show that the energy cost for a modest occupation of tetrahedral sites by nitrogen is small. From the above consideration, we suggest that results of preferentially nitridation of A104 sites are corresponding to previous model by simulations. Nitridation of y-AI2O3by the gas reduction-nitridation process is thus seen to be because of the particular susceptibility of the tetrahedral A104 site to nitridation, forming the tetrahedral crystal structure of AIN. Reaction at such low temperatures also results in restricted grain growth, producing nanosize AIN particles. When oxygen dissolves into the AIN lattice, vacancies form at Al sites and coalesce to generate A106 stacking faults. The residual oxygen in the initially formed products of this study may therefore occur in conjunction with stacking faults in the AIN. It seems therefore that the formation of AIN from y-AI2O3 progresses initially by the rapid conversion of A104 into AIN4, followed by the gradual nitridation of the residual A106 units.

IV. Conclusion AIN nanopowders were synthesized from y-A1203by gas reduction-nitridation over 1200”C, and that nitridation involved grain growth. Despite the temperature of the reaction, aA1203 was not generated. The tetrahedral A104 sites of the y-AI203 are nitrided to tetrahedral AIN, in preference to the octahedral A106 sites, but eventually all the AIL0 units are converted to nitride. The relatively low reaction temperature of the process enables AIN nanoparticles readily to be formed from y-AI203.

References ’V. A. Virkar, T. B. Jackson, and R. A. Cutler, “Thermodynamic and Kinetic Effects of Oxygen Removal on the Thermal Conductivity of Aluminum Nitride,” J. Am. Ceram. Soc., 72 [ 1 I] 203 1 4 2 (1992). ’Y. Baik and R. A. L. Drew, “Aluminum Nitride: Processing and Applications Key,” Eng. Muter., 122-, 553-70 (1996). ’I. Kimura, K. Ichiya, M. Ishii, N. Hotta, and T. Kitamura, “Synthesis of Fine AIN Powder by a Floating Nitridation Technique Using an N2/NH3Gas Mixture.” J. Mater. Sci. Lett., 8, 3 0 3 4 (1989). 4K. Komeya, N. Matsukaze, and T. Megro, “Synthesis of AIN by Direct Nitridation of Al alloys,” J Ceram. Soc. Jpn., 101, 1319-23 (1993). ’K. G . Nickel, R. Riedel, and G. Petzow, ‘Thermodynamic and Experimental Study of High-Purity Aluminum Nitride Formation from Aluminum Chloride by Chemical Vapor Deposition,” J . Am. Ceram. Soc., 72, 1804-10 (1989). ‘J. R. Park, S.-W. Rhee, and K.-H. Lee, “Gas-Phase Synthesis of AIN Powders from AIC13-NH3-N2.” J . Mater. Sci., 28, 57-64 (1993). ’R. Riedel and K . 4 . Gaudl, “Formation and Characterization of Amorphous Aluminum Nitride Powder and Transparent Aluminum Nitride Film by Chemical Va or Deposition.” J . Am. Ceram. Soc., 74, 13314 (1991). Bachelared and P. Joubert, “Aluminum Nitride by Carbothermal Nitridation,” Marer Sci. Eng., A109, 247-51 (1989).

ak.

154

’P. Lefort and M. Billy, “Mechanism of AIN Formation Through the Carbothermal Reaction of A1203in Flowing N2 Atmosphere,” J . Am. Cerom. Soc., 76. 2295-9 (1993). ”H.-K. Chen, C.4. Lin, and C. Lee, “Kinetics of the Reaction of Carbon/Alumina Powder Mixture in a Flowing Nitrogen Stream,” J . Am Ceram. Soc.. 77. 17534 (1994). “R. G. Odonnell and M. B. Trigg. “The Mechanism of Conversion of A1203to AIN Via Carbothermal Synthesis,” Micron, 25, 575-9 (1994). ”A. Tsuge, H. Inoue, M. Kasori, and K. Shinozaki, “Raw Material Effect on AIN Powder Synthesis from A1203Carbothermal Reaction,” J . Marer. Sci., 25. 235941 (1990). I3Y. W. Cho and J. A. Charles, “Synthesis of Nitrogen Ceramic Powders by Carbothermal Reaction and Nitridation: Part 3 Aluminum Nitride,” Matrr. Sci. Tecknol., 7,495-504 (1991). I4D. Qian, T. Meguro, Z. Xiong, J. Xie, and K. Komeya, “Effect of AIN and Raw Materials on Synthesis of AIN Powder by Carbothermal Reaction-Nitridation,” J. Mater. Sci. Soc. Jpn., 32, 82-6 (1995) (in Japanese). ‘’K. Baba, N. Shohata, and M. Yonezawa, “Synthesis and Properties of UItrafine AIN Powder by R F Plasma,’’ Appl. Pkys. Lett., 54, 2309-1 1 (1989). 16M. Iwata. K. Adach, S. Furukawa, and T. Amakawa, “Synthesis of Purified AIN Nan0 Powder by Transferred Type Arc Plasma,” J . Fhys. D:Appl. Pkys., 37, 1?:-7 (2004). S. Yu, D. Li, H. Sun, H. Li, H. Yang, and G. Zou, “Microanalysis of Singlephase AIN Nanocrystals and AIN-AI Nanocomposites Prepared by DC Arc-Discharge,” J . Crysral Growth, 183, 284-8 (1998). ’*F.J. Moura and R. J. Munz, “Vapor-Phase Synthesis of Nanosize Aluminum Nitride Particles Using a Two-Stage Transferred Arc Reactor.” J. Am. Ceram. Soc., 80 [9] 2425-8 (1997). ”Y. Sakka, H. Okuyama, T. Uchikoshi, and S. Ohno, “Characterization of Degraded Surfaces of Al Ultrafine Powders,” NonoStruct. Mater., 5, 577-88 (1995). ’OY. Qiu and L. Gao, “Nitridation Reaction of Aluminum Powder in Flowing Ammonia,” J . Eur. Ceram. Soc., 23, 2015-22 (2003). ”Y. Qiu and L. Gao, “Novel Way to Synthesis Nanocrystalline Aluminum Nitride from Coarse Aluminum Powder,” J . Am. Ceram. Soc., 86 [q 1214-6 (2003). ”M.-C. Wang, N.-C. Wu, M . 3 Tasi, and H.-S. Liu, “Preparation and Characterization of AIN Powders in the AIClrNH3-N2 System,” J . Cryst. Growth., 216,69-79 (2000). =B. C. D. Lello, F. J. Moura, and I. G. Solorzano, “Synthesis and Characterization of Nano-Scale Aluminum Nitride Produced from Vapor Phase,” Marer. Sci. Eng. C, 15, 67-9 (2001). “S. Iwama, K. Haydkawa, and T. Arizumi, “Ultrafine Powders ofTiN and AIN Produced by a Reactive Gas Evaporation Technique with Electron Beam Heatin ” J . Crystal Growth. 56, 265-9 (1982). ”T. Suehiro, J. Tatami, T. Megro, S. Matsuo, and K. Komeya, “Synthesis of Spherical AIN Particles by Gas-Reduction-Nitridation Method,” J. Eur. Ceram. Soc., 22, 5 2 1 4 (2002). ’q.Suehiro, J. Tatami, T. Megro, K. Komeya, and S. Matsuo, “Aluminum Nitride Fibers Synthesized from Alumina Fibers Using Gas-Reduction-Nitridation Method,” J . Am. Ceram. Soc., 85 [3] 715-7 (2002). 27T.Suehiro, N. Hirosaki, R. Terao, J. Tatami, T. Megro, and K. Komeya, “Synthesis of Aluminum Nitride Nanopowder by Gas-Reduction-Nitridation Method,” J . Am. Ceram. Soc., 86 [6] 10468 (2003). =. Y. I.Kim, M. A. Sriram, P. H. Mcmichael, P. N. Kumta, B. L. Phillips, and S. H. Risbud, “New Molecular Precursors from the Reaction of Hydrazine and Aluminum Alkoxide for the Synthesis of Powders in the A I a N System,’’ J. Pkys. Ckem. B., I01 I241 4689-96 (1997). -M. E. Smith, “Observation of Mixed AI(0, N)4 Structural Units by ”AI M a ic Angle Spinning NMR,” J . Phys. Ckem., 96, 1444-8 (1992). C. Ekstroem, K. J. D. MacKenzie, V. V. White, 1. W. M. Brown, and G . C. Barris, ‘Volatile Products Formed by Carboreduction and Nitridation of Clay Mixtures with Sillica and Elemental Silicon,” J . Mater. Ckem., 96,1225-30 (1996). ”K. J. D. MacKenzie, R. H. Meinhold, G. V. White,C. M. Sheppard, and B. L. Sherriff, “Carbothermal Formation of p’-Sialon from Kaolinite and Halloysite Studied by 29Si and 27AISolid State MAS NMR,” J. Muter. Sci., 29 [lo] 261 1-9 (1 ~ 4 ) . K. J. D. MacKenzie, “Solid State Multinuclear N M R A Versatile Tool for Studying the Reactivity of Solid Systems,” Store lonics, 172, 383-88 (2004). 33T.Brauniger, P. Kemp ens, R. K. Harris, A. P. Howes, K. Liddell, and D. P. Thompson, “A Combined F4N/Z7AlNuclear Magnetic Resonance and Powder Xray Diffraction Study of Impurity Phase in p’-Sialon Ceramics,” Solid Stare Nucl. Ma n Reson., 23 [I-21 62-76 (2003). 4 . . V . Nicolaescu, G. Tardos, and R. E. Riman, “Thermogravimetric Determination of Carbon. Nitrogen, and Oxygen in Aluminum Nitride,” J . Am. Ceram. Soc., 77 [9] 2265-72 (1994). ”N. Hashimoto, Y. Sawada, T. Bando, H. Yoden, and S. Deki. “Preparation of Aluminum Nitride Powder from Aluminum Polynuclear Complexes,” J. Am. Cerum. Soc., 74 [6] 12824 (1991). xE. Ponthieu, P. Grange, B. Delmon. L. Lonnoy, L. Leclercq, R. Bechara. and J. Grimblot, “Proposal of a Composition Model for Commercial AIN Powder,” J . Eur. Ceram. Soc., 8, 23341 (1991). 37G. A. Slack, “Nonmetallic Crystals with High Thermal Conductivity,” J . Ph s Ckem. Solids.. 34. 321-35 (1973). ‘K. J. D. MacKenzie and M. E. Smith, Mulrinuclear Solid Stare N M R <$Inorganic Marerials, Pergamon Series in Materials Science, Vol. 6. Pergamon/Elsevier, Oxford, 2002. ’9J. J. Fitzgerald, S. D. Kohl, and G. Piedra, “Observation of Four-Coordinate Aluminum Oxynitride (A104-yNr) Environments in AlON Solids by MAS ”AI NMR at 14 T,” Ckem. Marer., 6, 191S7 (1994).

4.

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40 X.M. Min, Y. Daiand, and C. W. Nan, “Structures of y-A1203and AlON Studied by ”A1 NMR and Quantum Chemistry Calculations,” Acru Merul. Sin., 12)?] 3 7 2 4 (1999). G . Yamaguchi and H. Yanagida, “Study on the Reductive Spinel-A New Spinel Formula AIN-AIzO, Instead of the Previous One A1,04,” BUN.Chem. Sox. Jpn., 32 [Ill 1264-5 (1959). 42G.Gutierra and B. Johansson, “Molecular Dynamic Study of Structural Properties of Amorphous A1~03,”P h p . Rev. B., 65, 1042021-9 (2002).

Synthesis Methods for Powders

“’H.P. Pinto, R. M. Nieminen, and S . D. Elliott, “Ab Initio Study of y-AIZO, Surfaces,” Phys. Rev. B., 70, 1254021-1 1 2004. 44F. H. Streitz and J. W. Mintmire, “Energetics of Aluminum Vacancies in Gamma Alumina,” Phys. Rev. B., 60.773-7 (1999). 45 S.-D. Mo, Y.-N. Xu.and W.-Y. Ching, “Electronic and Structural Properties of Bulk y-Al20,.” J. Am. Ceram. Soc., 80 [5] 1193-7 (1997). 46G.Gutierrez, A. Taga, and B. Johansson, “Thermal Structure Determination of y-Al,O,,” Phys. Rev. B., 65, 01210114 (2001).

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Membranes, Films, and Coatings

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MICROPOROUS ZrO2 MEMBRANE PREPARATION BY LIQUID-INJECTION MOCVD

S. Mathur, E. Hemmer, S. Barth, J. Altmayer and N. Donia Leibniz Institute for New Materials, Nanocrystalline Materials and Thin Film Systems Saarland University Campus, Building D2 2 D-66123 Saarbriicken, Germany

I. Kumakiri, N. Lecerf and R. Bredesen SINTEF, Materials and Chemistry PB 124 Blindern N-03 14 Oslo, Norway ABSTRACT Inorganic microporous membranes, presenting pores smaller than 2 nm, have an important potential for large-scale application in gas purification and separation. Given its good mechanical strength, thermal stability and chemical inertness, zirconia appears to be a promising material for the preparation of mesoporous and microporous membrane systems. We present here a membrane system based on the chemical vapor deposition of microporous zirconia on aalumina / y-alumina supports prepared by slip casting and sol-gel. For this purpose, Zr(OBu')4 was chosen as precursor for the deposition of nanocrystalline ZrO2 films. The thermal decomposition of the precursor in the liquid-injection CVD process was investigated as a function of the substrate temperature and different precursor injection frequencies. The asdeposited and post-annealed zirconia layers were investigated by EDX, XRD, SEM and AFM. First permporometry tests showed that the additional zirconia layer influenced the gas flux through the membrane in dependence of the CVD process parameters and resulted in a reduction of the total pore volume. INTRODUCTION Inorganic membranes, allowing a selective mass transport of species, have an important potential for industrial applications like purification and separation processes, particularly at high temperatures and under chemically harsh conditions . Recently, microporous silica is the most investigated system for membrane applications. However, the silica system suffers from poor stability to humidity. The limited chemical resistance, especially in alkaline media or in strong electrolyte solutions, turns the choice of adequate ceramic materials out to be complex and restricted. Herein, zirconia appears to be a promising material for the preparation of mesoporous and microporous membrane systems due to its good mechanical strength, thermal stability and chemical inertness. Zirconia can in fact be used in aqueous filtration with much better alkali durability than other ceramic membranes such as alumina or silica *. Currently, slip-casting of suspended colloidal particles and sol-gel processes are the most common techniques for ceramic membrane preparation '. The obtained macroporous supports can be modified by chemical vapor infiltration to deposit solids within the pores. This strategy allows tuning the pore size distribution of the membranes '. In this study a membrane of

'

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159

microporous zirconia on a-alumina / y-alumina supports was prepared by liquid-injection metalorganic chemical vapor deposition (LI-MOCVD). For the deposition of Zr02 by MOCVD the range of suitable precursors includes pdiketonates 6 , fluorinated P-diketonates 63 and alkoxides '. Although the alkoxide zirconiumtert-butoxide, Zr(OBu')4, shows sufficient volatility to be suitable for decomposition and formation of ZrO2 films in the conventional thermal MOCVD process 9, the use of a liquiddelivery system for the precursor implies several advantages: In liquid-injection MOCVD the solvent vapor is the major gas phase species and the concentration of the precursor in the gas phase is directly related to the concentration of the precursor in the liquid solution inducing an easy control of vapor pressure and precursor concentration in the gas flow. Furthermore, flow rate is fixed by the opening frequency and time of the piezovalve between vaporizer and reaction chamber offering the controlled supplies of precursor to reaction zone. This allows a simple control of film thickness by the number of injections. In addition, the precursor reservoir remains cold up to the point where the solution is vaporized providing any unwanted aging of the precursor and being a benefit under up-scalability and economical aspects lo

',

EXPERIMENTAL Membrane Preparation and Liquid-Injection MOCVD Macroporous a-Al2O3 supports were fabricated via slip-casting, followed by firing at 1100 "C for 1 h and machining to the required dimension as well as polishing. Porosity of these flat supports was 35 % and they had an average pore size of 80 nm. In a following step, mesoporous y-Al203 layers were deposited on the a-Al2O3 support by dip-coating under class100 clean-room conditions. After dip-coating, the membranes were dried and fired in air at temperatures between 600 "C and 1000 "C for 3 h with a heating / cooling rate of 1 K / min. The synthesis of zirconium tert-butyloxide (Zr(OBut)4) followed the procedure described in literature " in a modified Schlenk type vacuum assembly, taking stringent precautions against atmospheric moisture. Deposition of thin ZrO2 films on the y-Al2O3 layer of the support was affected by a liquid-injection CVD process using a solution of 2 Vol % Zr(OBu')4 in toluene. Precursor solution was injected for 30 min with 1 Hz, respectively 2 Hz injection frequency by 1 Torr reaction chamber pressure. The ms opening time of the piezovalve at 60 "C and alumina support was inductively heated in a range of 500 to 600 "C using a graphite susceptor. Some samples were additionally annealed. Therefore, post-thermal treatment was affected at 500°C during 2 h under atmospheric conditions. Instrumentation Surface morphology was recorded on a scanning electron microscope JSM-6400F (JEOL). Surface roughness was analyzed by atomic force microscopy (TMX2000 Explorer AFM with standard Si3N4 tip). Phase characterization was done by X-ray diffraction (XRD) analysis at room temperature on a Siemens diffractometer D-5000 fitted with a SEIFERT X-ray tube operating with Cu & (h = 0.154 18 nm).

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Permporometry Since the membranes prepared in this study are asymmetric, the BET method would not be adequate due to additional gas adsorption on the alumina support. Therefore, the mean pore diameter of the membranes has been estimated by means of permporometry based on ethanol / nitrogen permeation. In this technique, gas permeation is measured by increasing vapor pressure from dry (no vapor) conditions to close to saturated vapor conditions. Capillary condensation occurs first in the smaller pores during the measurement and blocks the gas permeation. The simultaneous gas flux through the remaining open pores allows the measurement of the fraction of pores with pore size larger than the pores that are already blocked. Upon desorption the same process occurs in reverse order. Accordingly, gas permeation decreases with increasing vapor pressure in relation with the Kelvin condensation equation (1):

where d k stands for the Kelvin radius, 0 for the surface tension, 6 for the contact angle, V, for the molar volume and where P is the vapor pressure. The pore size distribution of both y-Al203 supports and zirconia membranes were examined. Ethanol was used as vapor. All the gas flux measurements were done at room temperature. RESULTS AND DISCUSSION Film Growth by LI-MOCVD and Characterization The precursor Zr(OBu')4 was chosen for its gas phase stability and for the presence of pre-formed Zr-0 bonds which favor the formation of ZrO2 at low processing temperatures. The use of a single molecular source simplifies the growth procedure and allows a precise control over phase composition and particle size in the deposits 12. The aim of this study was to investigate the thermal decomposition of the molecular precursor as a function of the substrate temperature (500 "C, 500 "C followed by a 2 h annealing step at 500 "C and 600°C) and different precursor injection frequencies (1 Hz, 2 Hz). Morphology of the as-deposited and annealed films was investigated in high-resolution scanning electron microscope. A ZrOz film deposited at 500 "C (1 Hz) showed a very smooth and dense surface morphology without visible defects (Fig. la). Due to higher precursor concentration in the reaction area at enhanced injection frequency (500"C, 2 Hz) grain growth prevails nucleation resulting in a heterogeneous morphology (Fig. lb). In this case, film structure is built up by grains with facetted morphology with an average length of ca. 140 nm and an average width of ca. 60 nm. A similar effect with slightly smaller grains (about 120 nm in length and 45 nm in width) was obtained at 600 "C (1 Hz), however in this case nucleation of clusters that were unevenly distributed over the substrate was observed (Fig. lc).

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161

Figure 1: SEM micrographs of membranes deposited at 500"C, 1 Hz (a), 500"C, 2 Hz (b) and 600"C, 1 Hz (c): Development of the microstructure of the ZrOz membranes. Investigations of surface roughness by atomic force microscopy confirmed the described development of the surface morphology. Films deposited at 500°C (1Hz) showed only slightly increased roughness values independent of an additional post-annealing step (& = 4.6 nm, respectively 4.7 nm) when compared with the uncoated alumina support (Ra = 3.4 nm), while an increase in both injection frequency and substrate temperature led to enhanced roughness of about 6 nm (5OO0C, 2 Hz) and ca. 18 nm (6OO0C, 1 Hz), respectively. X-ray diffraction profiles of the membranes deposited at different process parameters showed good crystallinity and indicated in all cases a mixture of tetragonal (t-phase) and monoclinic (m-phase) ZrOz (Fig. 2).

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Figure 2: XRD patterns (Cu K,) of the ZrO2 deposits obtained at different process parameters: (a) influence of temperature, (b) influence of injection frequency. Referent patterns: t-ZrO2 [50-10891 (dotted line), m-ZrO2 [37- 14841 (drawn line). The recorded patterns show an increase in contribution of m-Zr02 for higher injection frequency (Fig. 2b) and higher substrate temperature (Fig. 2a). This indicates a t-to-rn transformation and / or grain growth mainly of the m-phase in dependence on temperature and precursor concentration in the reaction zone. The observed t-to-rn transition at increased substrate temperature (600"C, 1 Hz, especially the reflexes at 2 8 = 34.160" and 34.383', respectively 49.266' and 50.116' are indicating the m-phase) is conformal with the results of our former investigations for ZrO2 film growth by thermal MOCVD which showed a phase transition already at 450 "C '. However, while the transition from t- to rn-ZrO2 was completed at 550 "C in case of thermal MOCVD, the tetragonal modification in the films prepared by liquid-injection MOCVD remained partially stable till 600 "C. Formation of the high temperature tetragonal phase (transition points of the bulk material: monoclinic to tetragonal: 1170 "C, tetragonal to cubic: 2370 'C) at low tern eratures may be due to the lower surface energy of nanocrystalline tzirconia. While Garvie 13, reported about formation of stable t-Zr02 at low temperatures in case of particles smaller than the critical value of 10 nm in diameter and a t-to-rn transformation of larger grains, Chraska et al. l5 confirmed that even larger grains of the t-phase could be stable in case of direct nucleation in the tetragonal modification. A similar effect is observed by increasing the injection frequency that is directly correlated with a higher precursor concentration in the gas phase inducing faster growth of the first nuclei on the substrate. Furthermore, average grain sizes increased slightly from ca. 12 nm to ca. 15 nm for t-phase and to ca. 19 nm for rn-phase by doubling the injection frequency as calculated from the x-ray data using the Scherrer formula (reflexes used for calculations: V and 0 , Figure 2b). The different

B

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intensities of the reflexes for a-A1203resulting from the alumina support are due to the enhanced film thickness in case of higher injection frequency. The post-annealing of the as deposited membranes at 500 "C provokes strong grain growth up to an average grain size of about 21 nm (reflexes used for calculations: V and 0 , Figure 2a). The strong a-Al2O3 signals in this case may indicate diffusion processes of A13' into the deposited zirconia layer. Further, cross-section analysis of the membrane systems by high-resolution SEM was performed in order to investigate thickness of the zirconia membrane and microstructure within the membrane system. These investigations confirmed the prediction about enhanced film growth at higher injection frequency derived from XRD measurements and surface morphology. The thickness of the ZrO2 layers is uniform for the entire coated surface. The thickness values are directly related to the process parameters. The growth rates at 500 "C were found to be 3 nm / min and 12 nm / min for injection frequency 1 and 2 Hz, respectively. Figure 3 illustrates the development of the pore size in the three layers of the membrane: The macroporous a-A1203 support, the mesoporous y-A1203interlayer and the zirconia top layer.

Figure 3: Cross-section SEM micrograph showing the three compounds of the membrane system (500"C, 1Hz): a-Al203 support base, y-Al203 support coating and Zr02 top layer16. Flux Measurements and Pore Size Distribution Flux measurements and estimation of the mean pore diameter of the membranes by means of permporometry based on ethanol / nitrogen permeation illustrated the effect of deposition process parameters on membrane behavior in gas permeation. Gas flux in the y-A1203 membrane varied in the range of 3 - 6 mol/m2/s/Pa. Figure 4a presents the nitrogen flux measured through the membranes with increasing ethanol vapor pressure. Applying a ZrO2 layer by LI-MOCVD decreases the permeation flux to 50 % to 10 %, depending on the CVD conditions, of the initial value of the pure y-Al2O3 membrane. As expected, the smooth, dense films (5OO0C, 1 Hz) and the thick films with facetted structure (500"C, 2 Hz) provoked minor gas flux, while the membranes deposited at higher substrate

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temperatures (6OO0C, 1 Hz) suffering from pores and clusters induced enhanced gas flux. Further, a sharp flux drop at lower Kelvin radius (0.5 nm) was obtained for membranes deposited at increased temperature (600"C, 1 Hz) and after post-thermal treatment (500°C). In contrast, the membrane deposited at 500 "C with injection frequency of 1 Hz and 2 Hz showed no flux drop which is characteristic for denser layers. Most of the pores in y-Al203 membrane are in the range of 1 to 2.7 nm (see open keys in Figure 4b). Please note that the pore size in the figure represents Kelvin radius which is smaller than the real pore size as the thickness of adsorption layer is not taken into account. The shape of pore size distribution curves was changed after applying ZrO2 layer by CVD (see closed keys in Figure 4b). Approximately 50% of nitrogen permeation through CVD membranes was obtained through pores smaller than lnm Kelvin diameter. This suggested that pore size was narrowed by CVD deposition. Some of the CVD membranes had pores larger than 3 nm. These large pores might be originated from support and got less influence by CVD deposition. 2 0 - " " ' " " " ' " ' " " ' ' " ' 0- AI,O, support

-

.

-0- AI,O, support -*+ 500°C, 1Hz -A- 500°C, 2 Hz -#- 6OO0C, 1 HZ

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(a) (b) Figure 4: (a) Nitrogen flux with increasing ethanol vapor pressure for the uncoated alumina support and the ZrO2 membranes deposited at different process parameters. (b) Pore size distribution (permeance ratio (wet flux / dry flux)) as a function of Kelvin radius. CONCLUSIONS Crack free Zr02 membranes have successfully been synthesized by gas phase deposition using Zr(OBut)q as precursor. Its thermal decomposition in the liquid-injection CVD led to a nanoporous ZrO2 layer on the surface of the y-AlzO3 support layer. Film characterization by Xray diffraction method showed a mixture of t-ZrOz and rn-ZrO2 with further t-to-rn transition at elevated substrate temperatures and injection frequencies. In analogy, changes in morphology from smooth and dense films to thick films with facetted microstructure were obtained. Permporometry showed that the zirconia layer properties depend on the CVD process parameters. The developed procedure based on modification by pore size reduction of macroporous A1203 supports via zirconia deposition through gas phase technique presented the advantage to prepare zirconia membranes in one single step. While the sol-gel process, which usually necessitate clean room condition with careful drying and annealing, holds the high risk of

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agglomerations having automatically a negative effect on the pore size and the membrane quality this drawback does not appear when the membrane layer is deposited from the gas phase as it is achieved here. This study hence emphasized the feasibility of this procedure for the preparation of zirconia membranes which is interesting from an academic and from an industrial point of view. Significant stability under milder steam conditions, as encountered in steam sterilization and pervaporation, is expected from the prepared membranes. ACKNOWLEDGEMENTS Authors are thankful to the German Academic Exchange Service (DAAD) and to the Norwegian Research Council (NFR) for providing financial assistance in the frame of the Project Based Personnel Exchange Programme (PPP) between Germany and Norway. N.L., I.K. and R.B. thank the NFR for the financial support provided in the NANOMAT Programme (Functional Oxides for Energy Technology, Part C). S.M., E.H., S.B., J.A. and N.D. would like to acknowledge the NATO-CLG Program (No. 979756) for supporting this work. REFERENCES I

L. Cot, A. Ayral, J.J. Durand, C. Guizard, N. Hovnanian, A. Julbe and A. Larbot, "Inorganic membranes and solid state sciences", Solid State Sci., 2 , 3 13-334 (2000). 2 T. Yazawa, H. Tanaka, H. Nakamichi and T. Yokohama, "Preparation of water and alkali durable porous glass membrane coated on porous alumina tubing by sol-gel method", J. Membr. Sci., 60,307-317 (1978). 3L. Cot, C. Guizard and A. Larbot, "Novel ceramic material for separation process: present and prospective applications in microfiltration and ultrafiltration", Ind. Ceramic, 8, 143148 (1988). 4C.F. Lin, D.L. Flowers and P.K.T. Liu, "Characterization of ceramic membranes 11. Modified commercial membranes with pore size under 40 A",J.Membr. Sci., 92,45-58 (1994). 'J. Si and S.B. Desu, C.Y. Tsai, "Metal-organic chemical vapor deposition of Zr02 films using Zr(thd)4 as precursors", J. Mater. Rex, 9, 1721- 1727 ( 1994). 6M. Balog, M. Schieber, S. Patai and M. Michman, "Thin films of metal oxides on silicon by chemical vapor deposition with organometallic compounds. I", J. Cryst. Growth, 17,298-301 ( 1972). 7C.S. Hwang and H.J. Kim, "Deposition and characterization of ZrOz thin-films on silicon substrate by MOCVD" J. Mater. Res., 8, 1361- 1367 (1993). 'Z. Xue, B.A. Vaartstra, K.G. Caulton, M.H. Chisholm and D.L. Jones, "Chemical vapordeposition of cubic-zirconia thin-films from zirconium alkoxide complexes", Eur. J. Solid State Inorg. Chem., 29,2 13-225 (1992). 9 S. Mathur, J. Altmayer and H. Shen, "Nanostructured Zr02 and Zr-C-N coatings from chemical vapor deposition of metal-organic precursors" 2. Anorg. Allg. Chem., 630,2042-2048 (2004). "P.J. Wright, M.J. Crosbie, P.A. Lane, D.J. Williams, A.C. Jones, T.J. Leedham and H.O. Davies, "Metal organic chemical vapor deposition (MOCVD) of oxides and ferroelectric materials", J. Mater. Sci.: Materials in Electronics, 13, 67 1-678 (2002).

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"I.M. Thomas, "The preparation of alkoxides and triethylsilanolates of Ti, Zr, V, Nb, Ta and Sn from the dialkylamides", Can. J. Chem., 39, 1386-1388 (1961). 12 S. Mathur and H. Shen, "Inorganic nanomaterials from molecular templates" in "Encyclopediaof nanoscience and nanotechnology",H. Nalwa (Ed.), Amercican Scientific Publisher, 4, 131-191 (2004). I3R.C. Garvie, "The occurrence of metastable tetragonal zirconia as a crystallite size effect" J. Phys. Chem. Solids, 69, 1238-1243 (1965). 14 R.C. Garvie, "Stabilization of tetragonal structure in zirconia microcrystals" J. Phys. Chem. Solids, 82,218-224 (1978). 15 T. Chraska, A.H. King and C.C. Berndt, "On the size-dependent phase transformation in nano articulate zirconia'l Mater. Sci. Eng. A , 286, 169-178 (2000). P,S. Mathur, E. Hemmer, S. Barth, J. Altmayer, N. Donia, N. Lecerf, 1. Kumakiri and R. Bredesen, "Microporous ZrO2 film preparation by chemical vapor deposition", Proc. 9'hInt. Con$ on Inorganic Membranes / ed. R. Bredesen and H. Raeder, Lillehammer - Norway, 524527 (2006).

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GROWTH OF BARIUM HEXAFERRITE NANOPARTICLE COATINGS BY LASERASSISTED SPRAY PYROLYSIS G. Dedigamuwa, P. Mukherjee, H. Srikanth, and S. Witanachchi

Department of Physics, University of South Florida, Tampa, 33620 ABSTRACT Barium Hexaferrite (BaM) nanoparticle coatings have been deposited by a laser assisted spray pyrolysis method. The concentration of the Ba and Fe inorganic salts in the starting aqueous solution determined the final particle size. Films with particle size tunability in the range of 30-100 nm have been grown. Incorporation of C02 laser heating in the process enabled the growth of well-defined smaller particles in comparison to thermal pyrolysis. The use of SF6 as a carrier gas, which has a high absorption at the C02 laser wavelength, heated the nebulized droplets above 150°C. Several post annealing conditions to form the BaFel20-1structure have been studied. Structural and morphological characterization of the films has been carried out by x-ray diffraction, Scanning Electron Microscopy (SEM), and Atomic Force Microscopy (AFM). Magnetic properties of the nanoparticle films have been investigated by using a Physical Property Measurement System (PPMS). INTRODUCTION The formation of films consisting of particles of magnetic materials with size tunability is important for applications that require single magnetic domains. Some of the applications include magneto-sensors, bio-sensors, magneto-electronics, data storage, magnetic heads of computer hard disks, single-electron devices, microwave electronic devices, etc''2. Out of the variety of magnetic materials available, barium hexaferrite3 is one of the most widely studied materials due to its unique properties, that include high uniaxial magnetic anisotropy, high Curie temperature, high coercive force, and large saturation magnetization. Most of the applications require the grains to be small enough to contain only a single magnetic domain while large enough not to be superparamagnetic. As a result, several different techniques have been developed to grow nanoparticles and nanograined films of barium hexaferrite'. One of the simplest methods for the preparation of BaM is firing mixtures of iron oxide and barium carbonate at high temperatures (1200°C) followed by grinding to reduce the particle size4. This method generally yields mixtures which are non-homogeneous. In addition, BaM nanoparticle powders have also been grown by several methods that include chemical c~precipitation~~~, sol-gel method7, citrate precursor method', and thermal spray pyrolysis'. The citrate precursor method has the advantage of forming the BaM phase at a post annealing temperature of about 700C in comparison to 900°C required by other techniques". In this method aqueous solutions of stoichiometric amounts of Ba2+and Fe3+ nitrates are reacted with citric acid in a controlled pH environment to form a citrate complex. Thermal decomposition of the complex around 700°C leads to the formation of BaM. The citrate precursor method has also been combined with spray pyrolysis to form nanoparticle coatings of BaF". The chemical spray pyrolysis technique can be used to form coatings of a variety of different material^'^. The main component of a spray pyrolysis system is an atomizer that generates micro-droplets of a precursor solution dissolved in a relatively volatile solvent. Ultrasonic nebulizers are known to produce a fairly uniform distribution of micrometer size droplets. Generally, the nebulizer is operated at a frequency Membranes, Films, and Coatings

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of 2.4 MHz, where the precursor solution is converted into a mist of particles in a range of 1-2 pm in diameter. These particles lack sufficient inertia and thus have to be transported by a carrier gas. The droplets in the form of a fine spray are carried out of a nozzle onto a heated substrate by the carrier gas that can be inert or reactive. The constituents of the droplets decompose and react on the hot substrate to form the chemical compound. The substrate temperature should be high enough to evaporate the volatile solvents. The precursor is decomposed to form the compound when the droplets impinge upon the heated substrate. The subsequent film formation and morphology of the film is dependant on the velocity of the drop, rate of reaction and the rate of evaporation of the solvent. At high velocities the droplets will flatten on the substrate leading to large particle sizes". If the solvent in a droplet is evaporated so that the salts are condensed into a particle, the particle tends to retain its shape when it strikes the substrate. We have developed a laser-assisted spray pyrolysis method in which the evaporation of the solvent by the laser leads to the deposition of nanoparticle films with particle sizes that are smaller than those deposited without the laser. The solvent in the droplets are evaporated at the nozzle by a CO2 laser that causes the material in the droplets to condense into smaller solid particles. The laser energy is transferred to the droplets by the carrier gas SFs which has a high absorption at the CO2 laser wavelength. The size of the particles deposited on the substrate can be controlled by controlling the concentration of the solvent. The role of the laser in the laser assisted spray pyrolysis process can be seen in the AFM images of films deposited with and without the laser (Fig. 1). Formation of the BaM phase and the particle size distribution of the nanoparticle coatings have been observed to depend on the composition and the concentration of the starting solution as well as the post annealing condition. Several reports have pointed out the

Fig. 1: Films deposited by spray pyrolysis method (a) without laser heating, and (b) with laser heating need for excess Ba to form the BaM phase13. Even though the citrate precursor method facilitates the formation of BaM at a lower annealing temperature, time of annealing should be carefully controlled to prevent grain growth. This paper presents the results of a

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systematic study that was carried out to determine the optimum growth parameters to form nanoparticle BaM coatings with particle size control. EXPERIMENTAL METHOD Aerosols with an average drop size of 1.5 pm were generated by a (Sonaer Ultrasonic Model 241CST) nebulizer. The aerosol was carried through a conical nozzle into the growth chamber by SF6 gas. The schematic diagram of the laser-assisted spray pyrolysis system is shown in Fig. 2. The droplets and the SF6 gas interact with the focused 10.6 pm wavelength CO2 laser beam at the nozzle. The laser power of 3W used in the experiment was sufficient to increase the temperature of the gas and subsequently the droplets to about 350°C. Approximate value of the gas temperature was measured in a calibration experiment where a thermocouple was placed in front of the gas jet, just outside the laser beam-gas interaction volume. Oxygen was introduced into the growth chamber to promote oxidation of the depositing film. The flow rate and thus the speed of the aerosol into the chamber were controlled by gas flow meters. The pressure inside the chamber was about 700 Ton. Films were deposited on quartz or silicon substrates at a temperature of 300°C. The growth rate of the film depended on the concentration of the precursor. The average growth rate was about 50 d m i n .

Fig. 2: Schematic diagram of the laser-assisted spray pyrolysis film growth system.

Two different precursors were studied using the laser-assisted spray pyrolysis method. The first precursor was prepared by dissolving BaI2 and FeC13 in water. In addition to the stoichiometric Ba:Fe ratio of 1:12, films were deposited with Ba:Fe ratios of 1:6, 1:4, and 1:3. All the films were post annealed in air in a furnace using three different schemes: (a) anneal at 1000 "C for 5 minutes, (b) increase the temperature to 1000°C at the rate of 10"C/min followed by 10"C/min decrease to room temperature with no holding time, (c) increase the

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temperature to 1000°C at the rate of 10"C/min, hold for 1 hour, followed by 10"C/min decrease to room temperature. The second precursor was prepared by the citate method. The aqueous solution of 0.25M concentration Ba2+and Fe3+ citrate precursor was made by dissolving Fe(N03)3.9H20 (99% purity, Alfa Asaer), Ba(OH)2.8H20 (98% purity, Alfa Aesar), and anhydrous citric acid 12:19 in de(99.5% purity, Alfa Aesar) in molar ratios of Ba2+:Fe3+:C3H4(OH)(COOH)3=1: ionized water. The solution was then neutralized (pH=7) using an aqueous solution of NH3 (concentration: 28 wt%). The neutralized solution was refluxed at 60 "C for 3 h to allow carboxyl groups of citric acid to completely chelate metallic ions in the solution. Films deposited by using this precursor were post annealed at a temperature of 700°C with a heating and cooling rate of lO"C/min and a holding time of 1 hour. Films were also deposited using solutions with concentrations of 0.2M and 0.1M. The stoichiometry of as-deposited films was determined by EDS analysis. Crystallinity of the films was analyzed by x-ray diffraction while the morphology of the films and the nanoparticle sizes were determined by SEM and AFM. The magnetic properties of the BaM film grown under optimum conditions were investigated by a PPMS with a maximum magnetic field of 7 T. RESULTS AND DISCUSSION EDS analysis of the as deposited films from aqueous solutions of BaI2 and FeCl3 (first precursor) for different Ba:Fe ratios agreed closely with the composition of the starting solution. X-ray diffraction results of films deposited with this solution for the Ba:Fe ratios of 1:12, 1:6, and 1:3 are shown in Fig. 2. All the films have been post annealed in oxygen at 1000°C for 2 hours. X-ray diffraction of the films deposited with stoichometric composition of the starting solution (1:12) showed very weak signals from the BaM phase while preferential formation of Fe2O3 is indicated by the presence of strong peaks corresponding to Fe2O3. With increasing Ba content suppression of the Fe203phase and formation of BaM was observed. For a composition ratio of 1:3, which contained 300% more Ba than what is present in a stoichiometric solution, only the BaM phase was observed. This result is consistent with other reports that point out the requirement of excess Ba to stabilize the BaM phase14. In a typical x-ray powder pattern of BaM the highest intensity peaks are observed for the (107) and (114) orientations. Intensity ratio between the (008) and (107) is about 1:5. However, the highest peak observed for the deposited films was the (008) orientation, which indicates that most of the particles in the film are deposited on the substrate with a c-axis orientation. AFM image of the film (Fig. 4) clearly show the hexagonal grains which are c-axis oriented. It is also clear from this image that grain growth during the annealing process have given rise to grain sizes in the micron range. To prevent grain growth annealing temperature and annealing time have to be lowered. An SEM image of a film annealed at 1000°C for 5 minutes is shown in Fig. 5. The grain sizes are of the order of 200 nm. However the structure appears to be cubic. Formation of the BaM structure under different annealing conditions was studied by x-ray analysis of the films (Fig. 6). According to Fig. 6 (b), the 5 minute annealed film show only the Fez03 structure, which is consistent with the cubic structure observed in Fig. 5. These results suggest that gradual increase of temperature to 1000°C is important for the formation of the BaM Phase. When the films are held at 1000°C for a longer time, larger grains are formed producing high x-ray intensities. Films deposited by the

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citrate precursor were also analyzed by x-ray diffraction. Fig. 7 shows diffraction patterns corresponding to an as deposited film and a film annealed at 7OO0C for 1 hour.

30

40

2-Theta (degrees)

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Fig. 3. X-ray diffraction patterns of films deposited using precursors with Ba:Fe ratios of (a) 1:12, (b) 1:6, (c) 1:4, and (d) 1:3. Films were post annealed at lOOOC for 1 hour.

Figure 4.AFM image of film (d), which has only BaM peaks.

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Figure 5. SEM image of a film annealed for Smin. at 1000 "C

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2-Theta (degrees) Fig. 6 . XRD patterns of the samples that are subjected to different heat treatments. (a) As deposited. (b) Quick annealed (5 mins) at 1000 "C. (c) Ramp up to 1000 "C and ramp down. (d) Ramp up to 1000 OC, soak 1 hrs at 1000 "C and ramp down.

0

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2-Theta (degrees) Fig. 7: X-ray diffraction patterns of films deposited with the citrate precursor (a) as deposited, (b) annealed for 700 "C. All the peaks observed in the annealed film correspond to the polycrystalline structure of BaM. As indicated by the AFM images (Fig. 8) of films deposited by the citrate 174

Progress in Nanotechnology: Processing

Fig. 8 : AFM image of the film deposited with a precursor of 0.25 M concentration and annealed for 700 "C.

Fig. 9: 2D AFM image of the Samples prepared by different concentration at substrate Temp 350 OC by spray pyrolysis with the laser heating. (a) O.lM, estimated average particle size: 50 nm, (b) 0.2M, estimated average particle size: 65 nm. 300K In Plane

IOK In Plane 2.OE-03

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Fig. 1 0 Hyseteresis curves of a BaM film measured at 300 K and 10 K.

precursor method, the particle size remained in the nanoscale after the annealing step. The average particle size was about 50 nm. Fig. 9 shows AFM images of films deposited by the Membranes, Films, and Coatings

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citrate precursor and post annealed at 7OO0C for starting solution concentrations of 0.2M and 0.1M. The average particle sizes for these films were 50 nm and 40 nm, respectively. X-ray diffraction of these films also showed diffraction peaks that are similar to those in Fig. 8. M-H curves measured using a Physical Property Measurement System (PPMS) for films deposited by the citrate precursor and annealed at 7OO0C for 1 hour are shown in Fig. 10. Based on the measured hysteresis curves for magnetic fields applied perpendicular to the film, high coercivities of more than 4000 Oe were observed in both room temperature and at low temperature. The large coercivity indicates that the nanograins are mostly c-axis oriented. CONCLUSION We have compared two different precursors for the growth of BaM nanoparticle film using a laser assisted spray pyrolysis technique. This method has been successful in reproducing the stoichiometry of the starting solution in the deposited films. All the films had to be post annealed to form the BaM phase. When an aqueous solution of a mixture of BaI2 and FeC13 was used as the precursor, stoichiometric solutions led to the formation of FezO3. Excess Ba in the solution, up to a Ba:Fe ratio of 1:3, was required to form the BaM structure. Furthermore, the required high temperature annealing step at 1OOO°C caused grain growth leading to grains in the micrometer range. The use of stoichiometric citrate precursors in laser assisted spray pyrolysis produced the BaM phase at an annealing temperature of 70OoC. Low annealing temperature prevented grain growth leading to nanoparticles whose size was determined only by the concentration of the starting solution. A majority of the nanograins are preferentially oriented along the c-axis. Large coercivities (in excess of 4000 Oe) observed in hysteresis curves also indicate the c-axis orientation of the films. ACKNOWLEDGEMENTS This work was partially supported by the National Science Foundation, under Grant No. DMI0 127939. REFERENCE 1

B.J. Palla, D.O. Shah, P. Garcia-Casillas and J. Matutes-Aquino, “Preparation of Nanoparticles of Barium Ferrite from Precipitation in Microemulsions”. ,Journal of Nanopurticle Research., Vol 1,2,215-221 (1999). ’Kubo, 0. Ido, T.Yokoyama, H., “Properties of Ba ferrite particles for perpendicular magnetic recording media.”, Magnetics IEEE, Vol 18, issue 6, 1122-1124 (1982). 3C. S. Kim, S. W. Lee, S. Y. An, In-Bo Shim, “Magnetic Properties of Barium Ferrite Thin Films on Pt(ll1) by a Sol-Gel Meth0d”phy.s. stat. sol., (a) 189,3,903-906 (2002). 4 G. A. Ward, K. H. Sandhage., “Synthesis of Barium Hexaferrite by the Oxidation of a Metallic Barium-Iron Precursor”., J. American Ceramic Society., 80,6, 1508 (1997). 5T. Ogasawara, M.A.S. Oliveira,.”Microstructureand hysteresis curves of the barium hexaferrite from co-precipitation by organic agent.”, J. Mugn. Magn. Muter., 217,72 (2000). 6 S.R. Janasi, M. Emura, F.J.G. Landgraf, D. Rodrigues,. ”Theeffects of synthesis variables on the magnetic properties of coprecipitated barium ferrite powders.”, J. Mugn. Mugn. Muter., 238,2-3, 168-172 (2002).

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I

W. Zhong, W. Ding, Y. Jiang, N. Zhang, J. Zhang, Y. Du, Q. Yan., “Preparation and Magnetic Properties of Barium Hexaferrite Nanoparticles Produced by the Citrate Process”., J. American Ceramic Society., 80, 12,3258 (1997). 8V. K. Sankaranarayanan, R. P. Pant and A. C. Rastogi, “Spray pyrolytic deposition of barium hexaferrite thin “lms for magnetic recording applications”., J. Magn. Magn. Mater., 220, 72-78 (2000). 9M.V. Cabanas, J.M. Gonzalez-Calbet, M. Labeau, P. Mollard, M. Pernet, M. Vallet-Regi., “Evolution of the microstructure and its influence on the magnetic properties of aerosol synthesized BaFe[sub 12]O[sub 191 particles” J. Solid State Chem., 101,265 (1992). ‘OE. E. Vidal, P. R. Taylor, “Fundamental consideration in the deposition of Barium ferrite films in an inductively coupled plasma reactor”., Annals of the New York Academy of Sciences 891, 152-163 (1999). “Hsuan-Fu Yu and Hsin-Yi Lin,. “Preparation and thermal behavior of aerosol-derived BaFe1201
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Two Phase Monazite/Xenotime 30LaP04-70YP04 Coating of Ceramic Fiber Tows E. E. Boakye,t,S R. S. Hay,§ P. Mogilevsky,Sand M. K. Cinibulk§ %JES Inc., Dayton, Ohio 45432 $Air Force Research Laboratory, Materials and Manufacturing Directorate (AFRL/MLLN), Wright-Patterson Air Force Base, Ohio 45433

Equiaxed yttrium-lanthanum phosphate nanoparticles (Yo.7,Lao.3)P04 0.7H20 were made and used to continuously coat Nextel'" 720 fiber tows. The particles were precipitated from a mixture of yttrium and lanthanum citrate chelate and phosphoric acid (H3P04), and characterized with differential thermal analysis and thermogravimetric analysis, X-ray diffraction, transmission electron microscopy, and scanning electron microscopy. The coated fibers were heat treated at 1000"1300°C for 1, 10, and 100 h. Coating grain growth kinetics and coated fiber strengths were determined and compared with equiaxed La-monazite coatings. The relationships between coating porosity, coating hermeticity, and coated fiber strength are discussed.

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I. Introduction

M

(LaP04) fiber coatings promote crack deflection and fiber ull out in oxide-oxide ceramic-matrix composites (CMCs).l-'Two significant fiber coating issues are the thermomechanical stability of fiber coatings controlled by fiber coating coefficient of thermal expansion (CTE) difference ~ ?and~ the high-temperature microstructural instability of coatings, principally coating spheroidization enabled by grain growth.8- Large coating-fiber CTE differences promote debonding of coating during fiber handling and CMC processing. Spheroidization of thin coating limits coating coverage and therefore mechanical functionality, and exposes fiber surfaces to environmental effects during processing and use that may degrade fiber strength.12-15 The two oxide fibers most suitable for high-temperature structural CMCs are NextelT"610, an all alumina fiber with an average CTE of 7.9 x 10-60C-1,16317 and Nextel'" 720, an aluminamullite fiber with an average CTE of 6.6 x 10-6"C-1.16317 Both have CTEs lower than monazite (9.6 x 10-6"C-1).1 An oxide coating similar to monazite, but with a lower CTE that is more compatible with Nextel" 720 is desirable. Xenotime (YP04)18-20 bonds weakly to many oxides and has an average CTE of 6.2 x 10-6"C-1.21 There are some reports that compared with monazite, grain growth and sintering of xenotime are slow,21,22 so the high-temperature microstructural stability of xenotime coatings could be better than that of monazite. Two-phase mixtures of monazite and xenotime can in principle be tailored to match the CTE of Nextel'" 720. An additional possible benefit of two-phase mixtures is further reduction of grain growth rates,232s which in turn promote the high temperature microstructural stability and limit spheroidization of thin fiber coatings. ONAZITE

W. Mullins--contributing editor

Manuscript No. 23142. Received April 27, 2007; approved July 19, 2007. Work performed under USAF Contract #: F33615.01-C-5214. 'Author to whom correspondence should be addressed. e-mail: emmanuel.boa kye@wpa(b.af.mil

Membranes, Films, and Coatings

The formation, characterization, and possible benefits of twophase monazite/xenotime continuous fiber coatings of (Y0.7,La0.3)P04bulk composition are reported. Rare-earth phosphates such as monazite, can be synthesized by a variety of wet chemical methods.2G37A controlled precipitation method was used to make nano-sized, equiaxed LaP04. (nH20) rhabdophane particles by use of lanthanum citrate chelate (La-Cit) and phosphoric acid (H3P04).7"38.39 The controlled precipitation method was applied to make single phase (Y,La)PO4. (nH20) nanoparticles with a high Y content that convert to a two-phase equiaxed monazite/xenotime mixtures after heat treatment above 950°C during fiber coating. The effect of heat treatment temperature on grain growth, coating hermeticity and coated fiber strength is measured and discussed. 11. Experimental Procedure

(1) Materials Reagent grade lanthanum nitrate hydrate, phosphoric acid, aluminum nitrate nanohydrate (Aldrich Chemical Co., Milwaukee, WI); Yttrium nitrate hexahydrate (Alfa Aesar, Ward Hill MA) and citric acid (CA) (Fisher Scientific Co., Pittsburgh, PA) were used without further purification. The water of hydration of La(N03)3.xH20 was determined as 0.6 and that of Y(N03), . 6H20 was confirmed before (Y,La)-rhabdophane synthesis. Water was purified with the nanopure system (Model D4744, Barnstead/Thermolyne Corp., Dubuque, IA) for all experiments. ( 2 ) Precursor Chenustry Equiaxed hydrated (Y,La)P04 nanoparticles were formed by controlled precipitation from yttrium nitrate, lanthanum-nitrate, CA and phosphoric acid as previously done for La-Cit derived r h a b d ~ p h a n e .A~ ~yttrium-lanthanum citrate mixture (Y,La-Cit) was made by dissolving yttrium nitrate and lanthanum nitrate in ethanol and adding CA to complex the yttrium and lanthanum ions. The yttrium to lanthanum molar ratio was 7:3 and the (Y+La):CA molar ratio was 1:5. Phosphoric acid was added. The P to (Y+La) molar ratios was 1:5. The solution was stirred and equilibrated at -20°C for 15 min. Rhabdophane precipitation was not observed during the equilibration period. The solution was heated to -30°C to precipitate (Y,La)P04.xH20 according to: (Y, La)-Cit + H3P04

-+

(Y,La)P04 . H2O + H3Cit

(1)

The particles were separated from the supernatant by centrifugation, and washed and redispersed in deionized water to remove residual nitrate and citrate. Washing was repeated until the sol acidity decreased from a pH of < 1 to a pH of 3. Particles dispersed at pH = 3 were not stable; they agglomerated and settled within 3 h. The particles were de-agglomerated by

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Fig. 1. Transmission electron microscope image and as energy-dispersiveX-ray spectrum (EDX) of as precipitated (Y,La)FQ. 0.7H20 particles. EDX simulation showed the Y:La molar ratio to be 70:30.

adding NH40H to pH = -9, and were further dispersed with tetramethyl ammonium hydroxide (TMAH) solution. The pH was determined with a pH/ion meter (Coming Inc., Corning, NY). The dispersed particles were separated by centrifugation, washed and redispersed in deionized water. Washing was repeated until the pH was -7. The sol was then used for fiber coating. Transmission electron microscope (TEM) images of the coated fibers showed a thin layer of aluminum phosphate (AlP04) at the fiber coating interface caused by the reaction of the excess phosphorus in the coatings and alumina in the fiber surface (Section IlI(3) below). To scavenge excess phosphorus, some (Y,La)P04.H 2 0 sol was doped with 1 wt% AI2O3in the form of aluminum nitratsitrate. These sols were also used for fiber coating. (3) Characterization Differential thermal analysis (DTA) and thermogravimetric analysis (TGA) (Model STA-409, Netzsch-Geratebau Gmbh, Selb, Germany) were done at a heating rate of IO"C/min from 50" to 1400°C in air. Powder samples for DTA and TGA analysis were dried at 140°C for 18 h. X-ray diffraction (XRD) was done with a diffractometer (Rotaflex, Rigaku Co. Tokyo, Japan) after a 1 h heat treatment at 1000", 1 loo", 12W, 1300", and 1400°C. Particles were separated with Sorvall Super Centrifuge (Model T21, L.P., Newtown, CT). Particle agglomerate size was

Fig. 2. Particle size distribution for La-rhabdopane and (Y,La) rhabdophane particles measured by light scattering.

180

measured with a light scattering Particle Size Analyzer (Model N4 Plus Coulter'", Miami, FL). The coatings were examined with a scanning electron microscope (SEM) (Model 360FE Emst Leitz, Wetzler, Germany) equipped with energy-dispersive X-ray spectrometer (EDX) (ThermoNoran Voyager, Thermo Electron North America LLC, Madison, WI) operating at 15 kV. TEM of the particles and coatings was done with a 200 kV electron microscope (Philips CM200, FEI Co., Hillsboro, OR) equipped with EDS (ThermoNoran Voyager, Thermo Electron North America LLC, Madison, Wl). Particles were prepared for TEM analysis by drying sols on TEM grids. Coated fiber cross-section TEM thin foils were made by established methods?0s4' Single filament tensile strengths and Weibull modulii were measured using a 2.54 cm gauge length. A minimum of 50 tests were done for each data point under the assumption of a uniform fiber diameter!2

( 4 ) Fiber Coating Nextel"" 720 fiber tows were coated with the hydrated yttriumlanthanum phosphate [(Y,La)P04. H20] sol washed to a pH 7.

-

Fig. 3. Thermogravimetric analysis weight loss of rod-shaped LaPO4. XH20, equiaxed LaP04. XH2O and (Y,La)P04. XHzO. The weight loss is normalized to final monazite weight at 1400°C. Rod-shaped LaP04. XHzO bas a negligible weight loss at the coating temperature of I1OO"C. The corresponding weight loss for the equiaxed Lam4 XH2O and (Y. La) P04.XH20 are -0.8% and 0.9%0, respectively.

Progress in Nanotechnology: Processing

Table I. Thermogravimetric Analysis (TGA) Weight Loss for (Y,La)P04 and L a P 0 4 Precursors TGA weight loss% (normalized)

4

I

.-

0 v1

Temperature ("C)

(Y,La)P04

LaP04-equiaxed

LaF'04-rod-shaped ~

25-200 1100

5.2 0.9

950°C

5

A

~

3.8 0.7

3.3 Nedi pible

-. 20

Fibers were desized at 1100°C before coating. A continuous vertical coating apparatus was used in all the coating experiments along with I-octanol for immiscible liquid displacement.12 The coating apparatus was similar to that used previously, except that separate vessels housed the coating precursor and the immiscible liquid. Desized tow is wetted by coating precursor as it passes continuously under a teflon pulley. The wetted tows then go under another teflon pulley into I-octanol, where excess sol within and around the tow is displaced. The wetted tow then moves continuously at 3 cm/s through a vertical tube furnace, where the coatings are heat treated at 1100°C. The furnace hot zone was 8 cm in length and the total furnace length was 30 cm.The coatings are heated very rapidly and held at maximum temperature for -3 s. The supply and take up spools are interchanged and the process is repeated up to 6 times to increase coating coverage and thickness.

-

111. Results

( 1 ) TEM, Light Scattering Particle Size, and EDX Analysis TEM observation shows (Y,La)-rhabdophane equiaxed particles of 10 nm in diameter (Fig. I), similar to La-rhabdophane particles.3xHowever, light scattering measurements show larger particle sizes of 200 nm (Fig. 2). This suggests that there was significant particle agglomeration in the sols, similar to that inferred for La-rhabdophane!' Y to La ratio was analyzed by the standardless EDX method in the TEM using the Voyager software. The analysis yielded an average Y content in the powder of 69.8 at.% (with respect to the total cation content), in good agreement with the 7:3 Y to La molar ratio in which the components were mixed. Also, EDX elemental mapping of the asprecipitated powder revealed no phase separation or compositional inhomogeneity. (2) TGAIDTA and XRD The TGA data for (Y,La PO4 particles are compared with LaP04 particles in Fig. 3?" Data is normalized to the final anhydrous orthophosphate weight at 1400°C. From 50" to 200°C there was a 5.2% weight loss for (Y,La)P04, 3.8% weight

650°C

R

22

24

26

28 20

30

32

34

I

36

Fig.5. X-ray diffraction of (Y,La)P04 precursor heat treated to 860°C and 950°C.

loss for the equiaxed La-rhabdophane, and 3.3% weight loss for the rod-shaped La-rhabdophane. This is consistent with the known water of hydration of 0.7H20 for (Y,La)P04, 0.6 H 2 0 for e uiaxed LaP04, and 0.5 H 2 0 for rod-shaped LaPO)3~15~30~31s'8-M Weight loss above 1100°C was negligible for rod-shaped LaP04, 0.7% for the equiaxed LaP04 and 0.9% for the (Y,La)P04 (Table I). Weight loss for equiaxed Larhabdophane and (Y,La)P04 particles was consistently higher than that for rod-shaped particles above 300°C. This was attributed to the decomposition of absorbed CA and TMAH.38 The exothermic DTA peak for dehydration of rhabdophane to monazite was at -750°C for both equiaxed and rod-shaped LaP04 The corresponding peak for (Y,La)P04 was at -950°C (Fig. 4). This identification was confirmed by XRD. XRD spectrum of the powder heat-treated at 10"C/min to 860°C was consistent with hexagonal rhabdophane, but with diffraction peaks shifted to higher angles compared with pure La-rhabdophane (Fig. 5). The lattice parameters of the rhabdophane powder were consistent with its compo~ition!~ Powders heat treated to 950°C had a mixture of monazite and xenotime (Fig. 5). The positions of the monazite diffraction peaks were shifted to higher angles compared with pure La-monazite. From XRD data, the first monazite phase appearing immediately after the 950°C phase transformation from rhabdophane retains most (66 at.% out of -70 at.% based on the total cation content) of the Y from the parent rhabdophane solid solution.45 The diffraction patterns of the powder heat treated at 1000"-1600"C also showed monazite peak shifts to higher angles compared with pure LaP04 m0nazite.4~The magnitude of the shift increased with temperature, consistent with increasing solid solubility of Y in monazite at higher temperatures from 12 at.% at 1000°C to -42 at.% at 1600°C. The positions of the xenotime diffraction peaks were in good agreement with the PDF data46for YP04, regardless of the heat treatment. This suggests little solid solubility of La in xeno-

-

775°C Rod-Shaped

I

500

.

,

600

.

I

700

.

,

.

800

,

900

.

,

1000

Temperature ("C) Fig. 4. Differential thermal analysis showing exothermic peak temperaturesfor LaP04.XHzOand (Y,La)F'04.XH20at -750"Cand950"C, respectively.

Membranes, Films, and Coatings

Fig. 6. Transmission electron micrograph of as-deposited (Y,La)P04 fiber coatings on Nextel 720, showing a thin film of AIP04 in the coating fiber interphase. 181

Fig. 7. Transmission electron micrograph and electron diffraction patterns of fiber coatings of (Yo.~LQ.~)PO~, DyP04 (xenotime) and Lap04 (moh. Prominent rings are identified as monazite (blue) and xenotime (red). nazite) heat treated at 12OOoC/10

Fig.8. Transmission electron micrograph of [(Y,La)P04+ 1% A1203]coating on Nextel 720 showing a clean fiber coating interface.

182

Progress in Nanotechnology: Processing

Fig.9. Cross-sectionalscanning electron micrographs of [(Y,La)FQd+ I %

time.45The detailed analysis of the effect of heat-treatment temperatures and time on lattice arameters, phases, and compositions is presented elsewhere.4 P

(3)

Fiber Coating

Similar to the spherical La-monazite coatings:' grains in the (Y,La)-monazite/xenotime coating were equiaxed and averaged 13k0.7 nm in diameter (Fig. 6). This grain size was slightly larger than the particle size (10k0.5nm) in the (Y,La)P04. H 2 0 sol, suggesting minor growth during fiber coating. A semi-quantitative evaluation of the Y to La molar ratio in the coatings was done. Selected area electron diffraction pattern of the fiber coat-

A12031

coating on Nextel 720 fibers.

ing heat treated at 1200°C for 10 h (Fig. 7(a), left side) was compared with the diffraction patterns from pure monazite and pure xenotime coatings as standards (Fig. 7(b) and (c), respectively). Because pure YP04 xenotime coatings were not available, a diffraction pattern from a DyP04 xenotime coating obtained using the method48 was used in this analysis. Each pattern in Fig. 7 is a digital summation of eight individual patterns from coatings on different filaments. An enlargement of a strip of each pattern is shown in the upper right hand corner, and the integrated diffraction intensity of the entire pattern is shown in the lower right hand corner over the same d-spacing range (monazite-blue, xenotime-red). A digital summation 5G50 (weighting) of the monazite and xenotime patterns

Fig. 10. Scanning electron micrographs of thick equiaxed LaP04and (Y,La)P04fiber coatings heat treated from 1 to 100 hat 1100"-1300"C, showing higher grain growth rates of monazite.

Membranes, Films, and Coatings

183

Fig. 11. Transmission electron micrographs of equiaxed LaP04 and (Y,La)PO, in cross section after heat treatment at 1 loo", 1200", and 1300°Cfor 10100 h. Monazite coatings are locally hermetic for heat-treatment temperatures > 1100°C and time > 10 h. Monazite/xenotime coatings heat treated at temperatures up to 1100°C and time 100 h are not hermetic.

(the green line) is shown for comparison with the actual (Y,La)monazite/xenotime pattern (the black line). Prominent monazite and xenotime rings are identified with the d-spacings. As expected, some of t h e d-spacings were slightly shifted in the (Y,La)-monazite solid-solution. Both monazite and xenotime are present in the (Y,La) monazite/xenotime coating in roughly equal amounts. A thin film of amorphous AIP04 was present at the fiber/ coating interface (Fig. 6). Trace amorphous Alp04 was previously observed in fiber coatings made from rod-shaped rhabdophane particles, and was inferred to form by reaction of slight excess coating phosphorous with A1203 in the f i b e r ~ . ' ~ In coatings made with 1 wt% alumina doped sols, AIP04 films were not observed at the fiber-coating interface (Fig. 8). SEM observations show significant variations in coating thick184

-

ness 2 pm to (Fig. 9).

-

50 nm along and between fiber filaments

(4) Coating Grain Growth Lap04 and( Y, La) PO41 YPOd Grain growth of (Y,La)-monazite/xenotime coatings heat treated at I 100"-1300°Cfor I , 10, and 100 h is compared with LaP04 coatings made from similar precursors3' (Figs. 1C13). TEM and SEM observations show slower grain growth rates in (Y,La)-monazite/xenotime coatings (Figs. 10 and 1 I). The grain growth kinetics was evaluated using SEM data from Table 11, the results of which are shown in Fig. 12. The growth kinetics was represented by an equation of the form4': d - do = k( T)f'

(2) Progress in Nanotechnology: Processing

Table 11. Average Grain Size (nm) in Monazite (M) and Monazite/Xenotime (M/X) Coatings T 1ooo"c

1 IOWC

1300°C

1200°C

I

M

MIX

M

MIX

M

MIX

M

MIX

Ih 10 h 100h

35 45 48

37 41 45

53 76 93

47 63 75

140 200 256

96 130 156

268 3 14 597

157 213 302

of [In(d-d,)-nIn(t)] vs. l/Tfor monazite and (Y,La)-monazite/ where do is the initial and d is the final grain size, k(T) is a temxenotime grain growth are shown in Fig. 13. Q for monazite perature dependent growth constant, r is time, and n is a growth grain growth was 135 and 109 kJ/mol for monazite/xenotime exponent (Fig. 12). The growth exponents for monazite (0.1@ grain growth. These values were smaller than those found 0.23) were consistently larger than those for monazite-xenotime for CeP04 monazite (192 kJ/mol) and ErP04 xenotime (159 mixtures (0.084.12). All measured growth exponents were kJ/mol). much smaller than those observed for Ce-monazite (0.5)50and Er-xenotime (0.33).5' YP04 xenotime sintering rates are reported to be lower than those. for Lap04 monazite.22 Coating porosity and thin film constraints may be at least partly responsible ( 5 ) Coated Fiber Strength for the low growth exponent^.^^,^^ Also, two-phase. mixtures The cause of fiber strength degradation during fiber coating has may alter the grain growth rates by grain boundary pinning or been inferred to be high temperature stress corrosion from coatZener But in this particular situation the Zener ing precursor decom osition Froducts trapped in the coating model was not applied due to the following reasons: Experi(Table I and Fig. 3). '2-15339*5 s9 For rod-shaped La-rhabdop mental evidence shows that during heat treatment, the YP04 hane, high as-coated fiber strength correlated with low precursor phase forms a solid solution with the LaP04 phase. Powders weight loss at high temperatures. As-coated fibers made with heat treated at 1000°C for 100 h had 16.2 at.wt% of Y dissolved equiaxed La-rhabdophane precursor had high ascoated in LaP04. Also, powders heat treated at 1400°Cfor 1 h had 33.3 strength despite significant precursor weight loss above at.wt% of Y in LaP04.45The volume fraction of LaP04 will 1100°C39(Table I, Figs. 3 and 14). However, unlike rod-shaped therefore vary with the heat treatment time and temperature. In precursor, the fibers coated with equiaxed monazite had a conaddition, the grain growth study show that both the YP04 and siderable strength loss after long-term heat treatment for 100 h LaP04 phase grow whereas in the Zener model the inclusion is at 1100°C. Similar to equiaxed La-rhabdophane coated fibers, assumed to be inert and its size is fixed. Finally the kinetics of monazite/xenotime as-coated fibers made with the equiaxed growth of xenotime are slower than that of m ~ n a z i t e . ~ " ~ ~ ' ~ ' (Y,La)-rhabdophane ~' precursors had high as-coated strength k(7) is expressed by the usual Arrhenius equation: despite significant precursor weight loss above 1100°C (Table I, Figs. 3 and 14). However, their response to long-term heat treatment was rather similar to the rod-shaped pure La-mona(3) zite coatings. These fibers retained their high strength after heat treatment for 100 h at 1100°C (Fig. 14).'3s'4339 Nextelm where Q is an activation energy, R is the gas constant, T is 720 fibers previously coated with rod-shaped LaP04 had the temperature ("K) and c is a rate constant. Arrhenius plots strengths of 1.7550.05 GPa after 110o"C/100 h heat-treatment (Fig. 14). The corresponding values for equiaxed monazite and (Y,La)-monazite/xenotime mixtures were 1.38 kO.06 GPa and 1.76k0.05 GPa. The corresponding strength of uncoated fiber was 1.80&0.07 GPa. SEM and TEM studies of the coated fibers heat treated at 1100"C/100h showed the equiaxed monazite coatings to densify

?

-

Fig.12. Isothermal grain growth in thick equiaxed LaP04 and (Y,La)P04 fiber coatings heat treated at 1000"1300"C for 1-100 h. Growth exponents are shown on the graph for each heat-treatment temperature.

Membranes, Films, and Coatings

Fig.13. Arrhenius plots of equiaxed LaP04 and (Y,La)P04 grain growth for 1, 10, 100 h showing the calculated activation energies. 185

Fig. 14. Tensile strengths of equiaxed Lap04 coated Nextel 720 fibers; (a) as-coated fiber and (h) heat treated at 1 lOO”C/lOOh.

after heat treatment at IlOO°C/lOO h, whereas the monazite/ xenotime coating given a similar heat treatment does not (Fig. 11). In principle dramatic grain growth rate changes are expected as the rate controlling mechanism changes from coarsening by surface diffusion in porous non-hermetic coatings to grain boundary and lattice diffusion in denser coatings. This transition is diagnostic of open porosity transitioning to closed porosity in the coatings. This transition was not observed for either monazite or monazite/xenotime grain growth exponents or activation energies from 1000” to 1300°C (Figs. 12 and 13), although TEM observations of coatings heat treated at 1 100”C/ 100 h showed the monazite coatings to be at least locally hermetic and monazite/xenotime coatings to be non hermitic (Fig. 11). A tentative explanation for the strength loss after long-term heat treatment for coated fiber made from equiaxed Larhabdophane may relate to the high surface area and high sintering rate of equiaxed particles (Figs. 10-12). Equiaxed particles may adsorb more precursor decomposition products than rods of the same diameter. These desorb at higher temperatures. Some decomposition products may be trapped in coating pores adjacent to the fiber, and cause growth of fiber surface flaws by stress corrosion cracking driven by intergranular residual stress. The two-phase monazite/xenotime coatings was non hermetic (Fig. 11). Hence the less rapid densification of the two-phase coatings may trap a smaller amount of decomposition products adjacent to the fiber. Therefore, strength degradation by the two-phase coating was less severe.

IV. Conclusions A single precursor two-phase monazite-xenotime mixture (Y,La)PO4.0.7Hz0 was made and used to coat Nextel 720 fiber tows. The precursor was characterized by TGA/DTA and XRD studies before fiber coatings. The precursor transformed to a mixture of a (Y, La)-monazite and xenotime at -950°C. Coatings heat treated in-line at 1100°C formed a mixture of monazite/xenotime with a 13 nm grain size. Grain growth of two-phase monazite/xenotime coatings was slower than that of pure monazite coatings. The growth exponents for monazite (LaP04) were consistently larger than those for monazite-xenotime [(Y,La)P04/YP04]mixtures. As-coated fiber strength was not degraded by the monazite/xenotime coating derived from nanosized equiaxed particles, despite significant weight loss from the coating precursor at high temperature.

-

186

After heat treatment at I lO0”CjlOO h, the monazite/xenotimecoated fiber retained strength, while the monazite coated fiber was degraded in strength. Coated fiber strength loss may be related to the relative densification rates of the two coatings.

References ‘P. E. D. Morgan and D. B. Marshall, “Ceramic Composites of Monazite and Alumina,” J. Am. Ceram. Soc., 78, 1 5 5 3 4 3 (1995). *K. A. Keller ef a/., “Effectiveness of Monazite Coatings in Oxide/Oxide Composites After Long Term Exposure at High Temperature,” J. Am. Ceram. Sac., 86, 325-32 (2003). ’P.-Y. Lee, M. Imai, and T. Yano, “Effects of Sintering Condition on Mechanical Properties of Monazite-Coated Alumina-Fiber/Alumina-Matrix Composites Fabricated by Hot-Pressing,’’ J. Ceram. Soc. Japan, 112, 29-34 (2004). 4R. J. Kerans, R. S. Hay. T. A. Parthasarathy, and M. K. Cinibulk, “Interface Design for Oxidation Resistant Ceramic Composites,” J. Am. Ceram. Sac., 85, 25!9-632 (2002). F. W. Zok, “Developments in Oxide Fiber Composites,’’ J. Am. Ceram. Soc., 89, 330%24 (2006). 6A. G. Evans and J. W. Hutchinson, “Overview No. 120, The Thermomechanical Integrity of Thin Films and Multilayers.” Acta Mefall. Mafer., 43, 2507-30 c l y . E. E. Boakye, P. Mogilevsky, J. Welter, R. S. Hay, and R. J. Kerans, “Monazite Coatings on SIC Fibers I: Fiber Strength and Thermal Stability,” J. Am. Ceram. Soc. 89 [ I I], 3475-80 (2006). ‘D. J. Srolovitz and S. A. Safran, ‘Capillary Instabilities in Thin Films. I. Ener etics,” J . Appl. Phys., 60, 247-54 (1986). ’D. J. Srolovitz and S. A. Safran, “Capillary Instabilities in Thin Films. 11. Kinetics,” J. Appl. Phys., 60, 255-60 (1986). ‘OK. T. Miller, F. F. Lange, and D. 9. Marshall, “The Instability of Polycrysd i n e Thin Films: Experiment and Theory,” J. M a w . Res., 5, 151-60 (1990). “E. Werner, “The Spheroidization of Thin Plates,” Acfa. Mefall.,37, 2047-53 y9). E. Boakye, R. S. Hay, and M. D. Petry, “Continuous Coating of Oxide Fiber Tows Using Liquid Precursors: Monazite Coatings on Nextel 720,” J. Am. Ceram. So:? 82. 2321-31 (1999). E. E. Boakye, R. S. Hay, P. Mogilevsky, and L. M. Douglas. “Monazite Coatings on Fibers: 11, Coating without Strength Degradation,” J. Am. Ceram. Soc., 84, 2793-801 (2001). I4E. E. Boakye and P. Mogilevsky, “Fiber Strength Retention of La and Ce-P04 Coated Nextel’“ 720,” J. Am. Ceram. Soe., 87,314-6 (2004). ”R. S. Hay and E. Boakye, “Monazite Coatings on Fibers: 1, Effect of Temperature and Alumina Doping on Coated Fiber Tensile Strength,” J. Am. Ceram. Soc., 84, 2783-92 (2001). 16D. M. Wilson, “New High Temperature Oxide Fibers.” pp. 1-12 in High Temperalure Ceramic Matrix Composifes, Edited by W. Krenkel and R.H.N. Schneider. Wiley-VCH, Weinheim, Germany, 200 I . I’D. M. Wilson and L. R. Visser, “High Performance Oxide Fibers for Metal and Ceramic Composites,” Composites A, A32, 1143-53 (2001). “D.-H. Kuo and W. M. Kriven, “Characterization of Yttrium Phosphate and a Yttrium Phosphate/Yttrium Aluminate Laminate,” J. Am. Ceram. Soc., 78, 31214 (1995). I9D. 9 . Marshall, J. 9. Davis, P. E. D. Morgan, and J. R. Porter, pp. 27-36, Key Engineering MareriaLr. Trans Tech Publications. Switzerland, 1997.

Progress in Nanotechnology: Processing

2”P. E. D. Morgan and D. B. Marshall, “Functional Interfaces for Oxide/Oxide Composites,” Muter. Sci. Eng.,A162, 15-25 (1993). ”Y. Hikichi, T. Ota, K. Daimon, T. Hattori, and M. Mizuno, “Thermal, Mechanical, and Chemical Properties of Sintered Xenotime-Type RP04 (R = Y, Er, Yb. or Lu),”J . Am. Ceram. SOC.,81, 2 2 1 H (1998). 22D.Bregiroux, S. Lucas, E. Champion, F. Audubert, and D. Bernache-Assollant, “Sintering and Microstructure of Rare Earth Phosphate Ceramics R D O 4 with R E = La, Ce, or Y,” J. Eur. Cerum. Soc., 26, 2 7 9 4 7 (2006). ”K. B. Alexander, P. F. Becher, S. B. Waters, and A. Bleier, “Grain Growth Kinetics in Alumina-Zirconia (CeZTA) Composites,” J. Am. Cerum. Soc., 77, 93z446 (1994). J. D. French, M. P. Harmer, H. M. Chan, and G . A. Miller, “Coarsening Resistant Dual Phase Interpenetrating Microstructures,” J . Am. Cerum. SOC.,73. 250910 (1990). 2SF. F. Lange and M. M. Hirlinger, “Grain Growth in Two-Phase Ceramics: Al O3 Inclusions in 2102,’’ J. Am. Cerum. Soc., 70, 827-30 (1987). “L. Bo, S. Liya, L. Xiaozhen, Z. Shuihe, W. Chunfang, and L. Wenjing, “Monazite Coatings on Short Alumina Fibers Using Layer-by-Layer Assemby Techni M,”Muter. Sci. Eng., A354,3 2 6 3 2 (2004). q7V. Buissette, M. Moreau, T. Gamin, and J.-P. Boilot, “Luminescent Core/ Shell Nanoparticles with a Rhabdophane LnP04.xH20 Structure: Stabilization of Ce’+-Doped Compositions.” Adv. Funct. Muter., 16, 351-5 (2006). 28V.Buissette, M. Moreau, T. Gamin, and J.-P. Boilot, “Colloidal Synthesis of Luminerent Rhabdophane LaP04:Ln3+.xH20(Ln = Ce, Tb, Eu; x = 0.7) Nanocrystals.” Chem. Muter., 16, 3767-73 (2004). 29 H.-K. Jung, J . 3 . Oh, S. Seok, and T.-H. Lee, “Preparation and Luminescence Properties of LaPO4:Er,Yb Nanoparticles,” J . Lumin., 114, 2005 (2005). ’OS. Lucas, E. Champion, D. Bernache-Assolant, and G. Leroy, “Rare Earth Phosphate Powders ReP04.nH20 (Re = La, Ce, or Y) - Part 11. Thermal Behavior,” J. Solid-SfufeChem., 177, 1312-20 (2004). ”S. Lucas, E. Champion, D. Bregiroux, D. Bernache-Assolant, and F. Audubert, “Rare Earth Phosphate Powders ReP04.nH20 (Re = La, Ce, or Y) - Part 1. Synthesis and Characterization,” J. Solid-Sfare Chem., 177, 1302-1 I ( 2 W ) . ”T. Masui, H. Tategaki, S. Furukawa, and N. Imanaka, “Synthesis and Characterization of New Environmentally-Friendly Pigments Based on Cerium Phosphate,” J. Cerum. Soc. Jupun. 112, 646-9 (2004). ”J. M. Nedelec, D. Avignant, and R. Mahiou, “Soft Chemistry Routes to YPOCBased Phosphors: Dependence of Textural and Optical Properties on Synthesis Pathways,” Chem. Muter., 14,651-5 (2002). ’‘J. M. Nedelec, C. Mansuy, and R. Mahiou. “Sol-Gel Derived YPO4 and LuPO4 Phosphors, a Spectroscopic Study,” J . Mol. Str., 651653, 165-70 (2003). ”H. Oncda, H. Nariai, H. Maki, and I. Motooka, “Syntheses of Various Rare Earth Phosphates from Some Rare Earth Compounds,” Mafer. Chem. Phys., 73, 19-23 (2002). 36K. Rajesh, P. Mukundan, P. K. Pillai, V. R. Nair, and K. G. K. Warner, “High-Surface-Area Nanocrystalline Cerium Phosphate through Aqueous SolGel Route,” Chem. Mater., 16, 2 7 W 5 (2004). 37K. Rajesh e f ul., “Synthesis of Nanocrystalline Lanthanum Phosphate for Low Temperature Densification lo Monazite Ceramics,” Muter. Letf.,58,1687-91 (2004). ”E. E. Boakye, P. Mogilevsky, and R. S. Hay, “Synthesis of Spherical Rhabdophane Particles,” J . Am. Cerum. Soc., 88,274M (2005).

Membranes, Films, and Coatings

39 E. E. Boakye, R. S. Hay, and P. Mogilevsky, “Spherical Rhabdophane Sols 11. Fiber Coating,” J . Am. Ceram. Soc., 90 [5]. 1580-8 (2007). 40R.S. Hay, J. R. Welch, and M. K. Cinibulk. “TEM Specimen Preparation and Characterization of Ceramic Coatings on Fiber Tows,” Thin Solid Filmy, 308-309, 389-92 (1997). 4’M. K. Cinibulk, J. R. Welch, and R. S. Hay, “Method for Preparation of TEM Specimens of Coated Fibers,” J. Am. Cerum. Soc., 79, 2 4 8 1 4 (1996). 42M. D. Petry, T. Mah, and R. J. Kerans, “Validity of Using Average Diameter for Determination of Tensile Strength and Weibull Modulus of Ceramic Filaments,” J . Am. Ceram. Soc., 80, 2 7 4 1 4 (1997). 47E.E. Boakye, R. S. Hay, and P. Mogilevsky, “Spherical Rhabdophane Sols 11. Fiber Coating,” J . Am. Ceram. Soc., 90, 1580.8 (2007). %. Min, K. Daimon, T. Ota, T. Matsubara, and Y.Hikichi, “Synthesis and Thermal Reactions of Rhabdophane-(Yb or Lu),” M u m . Res. Bull., 35,219%205 (2ow. “P. Mogilevsky, E. E. Boakye, and R. S. Hay, “Solid Solubility and Thermal Ex ansion in LaP04-YP04 System,” J. Am. Cerum. Soc., 90, 1899-907 (2006). ‘M. K. Carron, C. R. Naeser, H. J. Jr. Rose,and F. A. Hildebrand, “Fractional Precipitation of Rare Earth with Phosphoric Acid,” U. S . Geol. Surv. Bull., 1036, 253-275 (1958). 41 E. E. Boakye, R. S. Hay, and P. Mogilevsky, “Spherical Rhabdophane Sols II. Fiber Coating,” J . Am. Ceram. Soc., 90 [5], 158C-8 (2007). “G. F. Fair. R. S. Hay, and E. E. Boakye, ”Precipitation Coating of Monazite on Woven Ceramic Fibers: 1. Feasibility,” J. Am. Cerum. Soc., 90, 44-55 (2006). 49J. W. Martin, R. D. Doherty, and B. Cantor, “Microstructural change due to Grain Boundary Energies”, pp. 3 0 7 4 6 in Srubility of Microstrucfure in Metallic Sysfems, Edited by R. W. Cahn. University Press, Cambridge, 1997. 9. Hikichi, T. Nomura, Y. Tanimura, S. Suzuki, and M. Miyamoto. “Sintering and Properties of Monazite-Type CeP04,” J . Am. Cerum. Soc.. 73, 3 5 9 4 6

(1990).

”I. A. Bondar, A. I. Domanski, L. P. Mezentseva, M. G. Degen, and N. E. Kalinina, “A Physicochemical Study of Lanthanide Orthophosphate,” Russ. J. Inorg. Chem., 21, I 1 2 6 9 (1976). ”M. Stech, P. Reynders, and J . Rcdel, “Constrained Film Sintering of Nanocrystalline TiOZ.” J. Am. Cerum. Soc., 83, 1889-96 (2000). ”M. N. Rahaman, Cerumic Processing undSintering. Marcel Dekker Inc., New York, 1995. “N. Maazi and N. Rouag, ‘Consideration of Zener Drag Effect by Introducing a Limiting Radius for Neighbourhocd in Grain Growth Simulation,” J. Crystul Growth, 243, 361-9 (2002). S5E.Nes, N. Ryum, and 0. Hunderi, “On the Zener Drag,” Acra Metall., 33, 11-22 (1985). 56M.N. Rahaman, R. E. Dutton, and S. L. Semiatin, “Effect of Solid Solution Additives on the Densification and Creep of Granular Ceramics,” Acru Muter., 45, 3017-28 (1997). ”N. Moelans, B. Blanpain, and P. Wollans, “Pinning Effect of Second-Phase Particles on Grain Growth in Polycrystalline Films Studied by 3-D Phase Field Simulations,” Acfu Muter., 55, 2173-82 (2007). ’*E. E. Boakye, R. S. Hay, M. D. Petry, and T. A. Parthasarathy, “ZirmniaSilica-Carbon Coatings on Ceramic Fibers,” J Am. Cerum. Soc., 87, 1%7-76 (2004). 59R. S. Hay, E. E. Boakye, and M. D. Petry, “ElTect of Coating Deposition Temperature on Monazite Coated Fiber,” J . Eur. Ceram. Soc., 20, 589-97 (2000). 0

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Template-Free Self-Assembly of a Nanoporous TiOz Thin Film Yanfeng Gaot and Masayuki Nagai Advanced Research Laboratories, Musashi Institute of Technology, Tokyo158-0082, Japan

Won-Seon Seo Korea Institute of Ceramic Engineering & Technology (KICET), Seoul, Korea

Kunihito Koumoto Graduate School of Engineering, Nagoya University, Nagoya 464-8603, Japan

We report a nanoporous TiOl thin film prepared using a supersaturated aqueous solution containing peroxotitanium complex ions. The film morphology can be regulated by chemical kinetics, which was partially controlled by solution conditions such as the concentrations of starting materials, pH values, and the temperatures of the solutions. Porous films with various morphologies from particulate to curved sheet shaped were prepared on different substrates including Si, polymers, and glass. Porous microstructures of films permitted us to prepare a crack-free film 2-3 pm in thickness. These films were still amorphous under the present treatment conditions. Dye-sensitized solar cells using annealed (anatase) films of different morphologies as electrodes yielded conversion efficiency ranging from 1.3% to 3.1%. Improvement in performance may be achieved by either increasing the film thickness or inducing crystallization in solutions. 1. Introduction OROUS Ti02 with

P

a tailored pore size of the order of nano to micrometers has attracted considerable attention because of potential applications including controlled drug delivery, photo catalysis, energy conversion, filtration, and biomedical membranes.' Although a variety of template-based approaches2 have been developed for the production of porous Ti02 thin films, the creation of high-quality structures, preferably over large areas, uniformly and at a low cost, is still a challenging problem.' Inspired by procedures used by living creatures in nature, where high-performance materials (magnets in magnetotactic bacteria, ferritine, teeth, bone, shells, etc.) are produced by self-assembling of highly selective structure^,^ we have developed a template-free, one-step method for the preparation of a nanoporous continuous Ti02 thin film over a large area. One feature of this process is gas generation during film deposition. The release of gas during film deposition leaves voids among self-linked colloid particles and acts as a template for the self-assembled construction of a networked porous film, and the size of the voids can be tailored by controlling the rate of gas generation, which is associated with the reaction conditions, such as temperature and/or pH of the solution.

N. Padture-contributing editor

Manuscript No. 21846. Received May 29,2006 approved November 9,2006. This work was financially supported by Murata Science Foundation and a specific fund from Musashi Institute of Technology. 'Author to whom correspondence should be addressed. e-mail: [email protected]

Membranes, Films, and Coatings

Another feature of this process lies in the control of the film morphology by simply regulating the solution conditions (temperature, pH), which enables us to obtain deposits with various morphologies, from a sphere shape to a nanosheet shape. When used as an electrode in a dye-sensitized solar cell (DSSC), a nanoporous film (-2 pm in thickness) with a unique morphology shows an efficiency of 3.1 YOunder illumination by AM 1.5 (- 100 mW/cm2) simulated sunlight. Our method is different from traditional chemical solution approaches. It is based on the solution chemistry and the solid-liquid interface reaction. We began with the preparation of a transparent peroxotitanium solution. The dissolution of H2Ti03 resulted in the formation of peroxotitanate complexes represented by [TiO(02)(OH)2]2- at high pH (> 10). The existence of hydroxyl (OH) and peroxo groups in the chemical composition of the peroxotitanate complex suggests that low temperatures and high pH values are favorable to stabilize the peroxotitanate complex. On increasing the temperature (to room temperature) or lowering the pH value (to 2.4), OH and peroxo groups were released, and a TiOz-based solid phase was pre~ipitated.~ Under these conditions (room temperature, pH = 2.4), the formation of a sphere-shaped particulate film was confirmed by both SEM and atomic force microscope (AFM) observations! The fundamentals of this process have been developed in our previous research. This work, however, is not a simple extension of the previous reported work, but provides further scientific insights toward understanding the fundamentals of the process. The film deposition process in a supersaturated solution is controlled by several factors, in which the solution physical chemistry, nucleation, and crystal growth are involved. Understanding how the reaction proceeds is a prerequisite for regulation of the deposition conditions, such as pH, temperature, and additives, such that it is possible to control the film properties in terms of morphology, roughness, adhesion of films to substrates, and crystalline phases along with their crystallinity, sizes, and orientation. The examples can be seen in studies involving the crystal growth and design of artificial minerals such as CaC03, BaS04, silica, etc.' Figure 1 shows the schematic mechanism proposed for the formation of a nanoporous Ti02 film. Firstly, the nucleation and growth of TiO2-based films were performed by adsorption of clusters grown in the solution or through the heterogeneous mode under different conditions (a). The bubbles that are evolved from the solution are present among separated islands of deposits during the initial stage of nucleation and growth. The release of these gas bubbles opens up pathways among deposits, acting as a negative template for the construction of foam architecture (b). Oxygen is generated partially from the condensation process as shown in Eq. (4) (see Section III), and partially from the decomposition of hydrogen peroxide (H202).

-

189

Fig. 1. Schematic description of the formation mechanism for a nanoporous TiOz thin film.

In fact, the formation of bubbles is observed during the deposition process, which plays a significant role in the construction of a porous structure. Control of the size and the quantity of pores can be achieved by controlling the chemical kinetics of a series of reactions. When the evolution rate of oxygen gas is slower than the rate of particle growth and densification, the oxygen gas that is generated from the solution and sealed among colloid particles should create voids or a continuous path, permitting gas to escape from the substrate surface to the solution. The rates of gas generation, film deposition, and solid formation in the bulk solution are all related to the solution conditions, such as the concentrations of the reactants, pH values, and the temperatures of the solutions. By controlling these parameters, we can manipulate the size of the pores or even fabricate a dense film. 11. Experimental Procedure

( 1 ) Materials and Method The raw chemicals in this study were mainly titanic acid (H2Ti03, 97%, Mitsuwa Chem., Osaka, Japan), ammonia (NH3, 28% in water, Kishida, Osaka, Japan), hydrogen peroxide (H202,30% in water, Kishida), and deionized water (resistivity > 18.4 MQ . cm, Millipore, Tokyo, Japan). First, we prepared a transparent peroxotitanium complex solution by adding H2Ti03 to a solvent mixture containing an appropriate amount of NH3 and H202. The beaker, along with the chemicals in it, was cooled in an ice bath during this operation. The detailed description of this process has been published previously? The deposition solution contained 5 mM of Ti4+,and the solution pH was adjusted to 1.0-2.0 by adding an appropriate amount of HN03. cm, Shinetsu, Tokyo, Japan) The p-Si (resistivity: 5-10 glass with transparent conductive oxides (resistivity: 10 R cm, Asahi Glass, Tokyo, Japan), or polyethylene terephthalate (film type with a thickness of 50 pm, UNITIKA, Osaka, Japan) of 1 cm x 1 cm was used as a substrate. Before use the substrates were cleaned ultrasonically in acetone (99.5%, Kishida), ethanol (99.5%, Kishida), and deionized water. After drying at 5 0 T , it was exposed to ultra-violet light (184.9 and 253.7 nm; low-pressure mercury lamp, 200 W, PL21-200, 15 mW/cm2 for 254 nm, SEN Lights Co., Osaka, Japan) for 5 min. Deposition was conducted by soaking the substrates in the prepared solution containing 5 mM Ti4' at 20"-95"C for several hours to 48 h. The substrates were hung vertically in the solution with the glass side against the beaker wall. Subsequently, the substrate was rinsed with deionized water three times, and then dried at 50°C in air.

-

( 2 ) Integration into Solar Cells and Characterization In order to increase the adhesive property of the Ti02 film to the substrate, a thin layer of dense anatase Ti02 was grown on the F:Sn02 (FTO, < 12 Q/square, Asahi Glass) substrate by the liquid-phase deposition method before the substrate was treated in a peroxotitanium complex solution. Briefly, the FTO substrate was soaked at 90°C for 2 h in an aqueous solution containing 0.05M of (NH4)2TiF6 (W%, Stella Chemifa, Osaka, Japan) and 0.015M of H3BO3 (99.5%, Showa Chemical, Tokyo, 190

Japan); the starting pH of the solution was regulated to 2.8 by adding an appropriate amount of HN03.5b Porous Ti02 films were then produced repeatedly in fresh peroxotitanium complex solutions (5 mM Ti4+) at 95°C for 12 h on transparent conductive substrates under different pH. The thicknesses of films deposited five times were 2-3 pm. After crystallizing to anatase by annealing at 500"C, these films were immersed subsequently in a 0.5 mM Ruthenium535-bisTBA (N-719, chemical name: cis-bis (isothiocyanato) bis (2,2'-bipyridyI4,4'-dicarboxylato)ruthenium(I1):bis-tetrabutylammonium, produced by Solaronix, Aubonne, Switzerland)-ethanol solution for 12 h at room temperature. A Pt layer coated on a transparent conductive plate was used as a counter-electrode. A small amount of electrolyte solution, Iodolyte TG-50 (Solaronix), was loaded into the gap of two electrodes by the capillary effect. Simulation solar light from a Xenon lamp had an intensity of 100 mW/cm2 (AM1.5) (Solar Simulator: Model YSS-50, Yamashida Denso, Tokyo, Japan). The distance from light exit to cell was 12 cm. The area of the Ti02 electrode was 0.28 cm2. An evaluation system (R0240A, Advantest, Tokyo, Japan) was used to measure the current-voltage (J-V) characteristics of the solar cells.

(3) Characterization Other Properties The film morphology was observed with a field-emission scanning electron microscope (FE-SEM, JSM-6700F, JEOL, Tokyo, Japan) and a scanning probe microscope (SPM, SPI3800N, Seiko, Tokyo, Japan) operated to obtain AFM graphs of the films. The scanning frequency for SPM was 1-2 Hz. The calcined samples were also characterized with a transmission electron microscope (TEM, JEM4010, 400 kV, point-to-point resolution: 0.15 nm, JEOL). The phase composition was characterized using X-ray diffraction (XRD, RINT2100, 40 KV, 30 mA, CuKa, Rigaku, Tokyo, Japan) with a graphite monochromator. The scan was operated in the out-of-plane mode at a rate of I"/min. Transmittance was measured using an ultraviolet visible spectrophotometer (V-570, JASCO, Tokyo, Japan). The chemical compositions were investigated using an X-ray photoelectron spectroscope (Esca3200,8 kV, 30 mA; Pass Energy: 75 eV, Shimazu, Kyoto, Japan). A trace amount of Au was sputtered to the surface as a reference. The X-ray source was MgKa and all spectra were referenced to the Au4f signals. Raman spectra were measured using a JASCO NRS-2100 system (laser power: 10 mW, scan speed: 100 cm-I min-I). III. Results and Discussion Understanding the reaction equations is fundamental to regulate kinetically the reaction for obtaining the desired morphologies of films. The formation of peroxotitanium (IV) from either (titanium) Ti metal powders or Ti4+ has been reported to release protons and to produce a cationic complex ion of peroxotitanium (Eq. (I))?' These Ti species are hydrolyzed immediately in water, resulting in turbid solutions. For obtaining a homogeneous solution for film preparation, an additional amount of acid is usually added, and the solution temperature is usually maintained below or close to 0°C. In the present study, we selected a relatively stable, solid Ti(IV) chemical, titanium acid (H2Ti03), Progress in Nanotechnology: Processing

as the starting material for Ti(1V). Although H2Ti03is difficult to dissolve in water even in 30% H202, the addition of a large amount of ammonia shifts the reaction to the right-hand side (Eq. (l)), promoting the dissolution of H2Ti03 (Eq. (2)). The as-prepared solution is composed of a peroxotitanium complex in the anionic form (Eq. (2)). The component of peroxotitanium (IV) may change with a decrease in the solution pH, as a result of a condensation reaction among the peroxotitanium species. However, these condensation reactions may probably produce either H+ or OH- depending on the charge forms of peroxotitaniums,6b which are cationic or anionic depending on the solution P H . ~ ~ Therefore, the effect of solution pH on the condensation reactions is complex. Previous studies have suggested that protons can play the role of a reactant when the solution pH is decreased from -9 to a pH range of 24.’ At pH below 1, a mononuclear complex with a formula of Ti(02)’+ is formed. Condensation of this peroxotitanium species may release protons. Thus, increasing the solution pH may lower the rates of hydrolysis/condensation reactions toward the formation of Ti02-based precursors (Eq. (4)). The condensation of peroxotitanium complexes may precipitate various titanium oxide hydrates with or without residue of peroxo groups. One of these deposits has been determined to be Ti01.4(02)0.5(OH)0.2 when deposition was carried out close to pH 2.4?

Ti4++ H2Oz + Ti(Oz)(OH),-,

(4-n)+

+ nHt

TiO(OH),+HzO2+2OH-@ [TiO(02)(OH)2]2-+2H20 (2)

Ti(O2)”

+ nH20

@

Ti0,T(02),(OH),

+ 2 H+ + 0 2

(4)

The stability constant of peroxotitanium (IV) in different acidic media was found to decrease with increasing solution temperature,6b and both the decomposition of peroxotitanium to Ti (IV) ions and the hydrolysis reaction of Ti (IV) to solid phases are endothermal reactions.’ Thus, one can hope that treatment in solutions at relatively high temperatures can accelerate the formation of solids. On the basis of these reaction equations, we proposed to optimize the film microstructures by controlling the growth kinetics, i.e. by regulating the solution conditions: pH and temperature. Figure 2 shows scanning electron micrographs for the Ti02 precursor films prepared at pH 1.8 in solutions with different temperatures for various times along with those of anatase films obtained by annealing the as-deposited films. All the films show porous characteristics with pore sizes of tens to about 100 nm and an approximate thickness of 100-200 nm. The pore size increased with increasing solution temperatures. It seems that the solution temperature has significant effects on the film morphology; on increasing the soaking temperatures, typically from 20°C (Fig. 2(B)) to 30°C (Fig. 2(C)) to 50°C (Fig. 2(D)), the number of cracks decreased and a homogeneous porous film was obtained (Fig. 2(D)). Crack generation is associated with stress distribution in the film during annealing. Both substrate properties and film thickness affect the stress distribution. There should be, therefore, a certain critical thickness, beyond which films tend to form cracks during drying or heating.8a It is probable that the constructed porous structure itself can relax the stress and prevent cracks from generating during annealing. In fact, we have obtained films about 2-3 pm in thickness without apparent cracks. Compared with those in the as-deposited film prepared at room temperature (Fig. 2(A)), the cracks became large after annealing at 600°C (Fig. 2(B)). However, the film thickness remained almost unchanged. This finding suggests that the densification in the direction perpendicular to the substrate is not obvious, implying that the porous structure was not destroyed even after annealing, Membranes, Films, and Coatings

Fig. 2.

Scanning electron microscope photographs of TiOz thin films prepared under different conditions: (A) as deposited at room temperature (-20°C) for 48 h; (B) annealed at 600°C of film A; (C) deposited at 30°C for 48 h and annealed at 600°C for 2 h, (D) deposited at 50°C for 48 h and annealed at 600°C for 2 h. All the scale bars are 0.1 pm.

which was also the case for those prepared under other conditions. From the cross-sectional graph in Fig. 2(D)2, it can be seen that pores were present not only in the superficial layer but also throughout the film. At low temperatures (Figs. 2(B) and (C)), the formation of aggregates is not noticeable. However, at a relatively higher temperature (50”C), the adhesion of particles can be clearly observed. These colloid particles 10-30 nm in diameter linked with each other, enclosing voids of various sizes. The difference in deposit microstructures implies the various formation mechanisms. As we have discussed above, the rate for the solid formation was accelerated with increasing reaction temperature. The induction period obviously became short at relatively high temperatures; precipitates were formed in all solutions as deposition progressed on. A large amount of precipitates in the bulk solution enables films to grow by attachment of tiny particles from the solution at relatively high temperatures. To investigate the effects of solution pH on the film morphology, films were prepared at the same temperature of 95°C for 24 h in solutions at pH = 1.0, 1.5, and 2.0, respectively. As shown in Fig. 3, void sizes can be varied by regulating the solution pH, whose length increased with decreasing pH. At pH 1.0, thin films were constructed by distorted thin nanosheets of a TiOz precursor, forming a homogeneous surface with interconnected nanovoids. These nanosheets seem to have grown from a 191

inn

I.

PET film

I - -

300

500

700

900

1100 1300 1500 1700 1900

Wavelength / nm Fig. 5. UV-V spectra of Ti02 films deposited on polyethylene terephthalate (PET) films (conditions: pH 1.0, pH 1.5, pH 2.0; [Ti]5 mM, 95°C at pH 2 for 12 h).

Fig. 3. Scanning electron microscope photographs of TiOz precursor films deposited at 95°C for 24 h in 5 mM peroxotitanium complex solutions with different pH (Al, A 2 pH 1.0; BI, B 2 pH 1.5; C1, C2: pH 2.0).

central point, extending into a flower-like architecture (Fig. 3(A2)), which were also observed in a cross-sectional SEM photograph. This growth behavior may imply that the film was formed by a heterogeneous nucleation mechanism, where preferable growth on pre-existing nuclei resulted in competitive growth from one point to various directions. The aspect ratio of nanosheets decreased obviously on increasing the pH to 1.5 and 2.0. These SEM photographs also suggest the changes in surface roughness, increasing with decreasing pH. Further characterization of the porous film was performed with an AFM. The film surface clearly demonstrated porous characteristics; most of the pores were nanometers in size (Fig. 4). For the measured areas of 2 pm x 2 pm, the quantitative surface roughness (represented using root mean square) was

11, 21, and 51 nm for the film obtained at pH 2.0, pH 1.5, and pH 1.0, respectively. The relatively rough surface of that obtained at pH 1.0 may lead to the problem of poor contact of the film to the electrode when the film is applied to solar cells. It was reported that there were three kinds of pores in the TiOz powders prepared by the hydrolysis of titanium tetraisopropoxide in the presence of HNO, and NH40H under ultrasonic irradiation: fine intra-a gregated pores, larger interaggregated pores, and voids.8' The existence of voids and inter-aggregated pores in both the as-deposited and calcined films is obvious according to the present research results (SEM micrographs), whereas direct observation of intra-aggregated pores such as mesopores is limited to the resolution of SEM. Adsorption edges of ultra-violet spectra became sharp with increasing pH, suggesting that the surfaces of films became smooth and there was an increase in film density, which is in accord with the results of surface roughness and SEM observation. Note that the transmittance of sample pH 1.0 is only half of that of pH 2.0 (Fig. 5). The decrease in transmittance in the visible range may be due to the pores-resulted light scattering. A similar result was also reported for a film synthesized by the sol-gel method.8c For the as-deposited film, the crystallinity was characterized by both XRD (not shown here) and Raman spectroscopy. No crystalline phases such as anatase or rutile were detected by either XRD or Raman for the samples prepared at pH 1.5 and pH 2.0, suggesting that the film deposited under the present conditions was in an amorphous state. For integration into a solar cell, the film needs to be annealed for crystallization. For film prepared under pH 1.O, a small peak at about 51C520 cm-' was observed, which may be assigned to crystallized titaniarelated forms. However, almost no corresponding peaks were detected by XRD. Thus, the crystallization in this case may be very weak.

Fig.4. Atomic force microscope images ofTiOz thin films deposited under different pH values.

192

Progress in Nanotechnology: Processing

Fig.6. Scanning electron microscope graphs of Ti02 powders after annealing of those obtained under different pH; (A) pH 1.0, (B) pH 1.5, (C) pH 2.0; the scale bars are 1 pm.

For integration into solar cells, the samples were crystallized by annealing at different temperatures. Crystallization into anatase Ti02at 25Oo-3WC was confirmed by XRD. In this study, all the samples were calcined at 500°C for 30 min to increase their crystallinities. To collect sufficient amounts, the precipitates in solutions were collected for characterization by SEM and TEM. Figure 6 shows the SEM photographs for the samples after annealing. From the SEM graphs, the morphologies of the deposits changed slightly compared with those observed for Ti02 films on substrates. Highly porous Ti02 aggregates were obtained for that synthesized under pH 1.0 (Fig. 6(A)). Although the difference in sheet dimensions is visible, a single deposit collected from solutions is also sheet shaped (Fig. 6(A)), similar to that in thin films (Figs. 3(A1) and (A2)). The pores obviously became small when the pH increased to 1.5. Isolated particles were formed when the pH was further increased to 2.0. This result suggests that the formation of inter-aggregated and voids is closely related to the solution conditions, which may have effects on both the surface physical chemistry of colloid particles in solutions and the chemical compositions of precursor precipitates. Figure 7 shows the TEM photographs of Ti02 powders after annealing at 500°C. It seems that the crystals prepared under pH 1.0 are thin plate shaped, with particle sizes from several nanometers to about 50 nm in length (Fig. 7(A)). In a magnified graph (Fig. 8), the crystalline lattice fringes of the anatase (101) plane have clearly been observed. The sizes of powders obtained at pH 1.5 appear to be the same as those obtained at pH 1.O, but these powders seem to be aggregated from tiny nanoparticles (Fig. 7(B)). The particles obtained at pH 2.0 were found to be spindle shaped, which were also observed to be formed through aggregation (Fig. 7(C)).

Figure 9 shows the J-V curves of DSSC using the nanoporous Ti02 thin films as electrodes. The best result was obtained for the TiOz film prepared at pH 1.5, which gave a short-circuit current density (J,) of 13 mA/cm2, a fill factor (FF) of 37%, an open-circuit photovoltage (V0J of 662 mV, and a conversion efficiency of 3.1%. A device using a film prepared at pH 2.0 gave a conversion efficiency of 2.2%. Among three sample films, the film prepared at pH 1.0 showed poorer properties than the other two, which may be partially due to the large surface roughness of the film, and partially due to lowered transparency (see Fig. 5). Compared with V , of the DSSCs derived from pH 1.0 to pH 1.5 films, the J, increases largely from 3.9 to 13 mA/ cm2, more than a twofold increase. The V,, fluctuates between 662 and 632 mV, but the F F increased for the low-efficiency curves (0.37 vs 0.56). Accordingly, the increase in the overall conversion efficiency is entirely due to an increase of J,. The film obtained at pH 2.0 showed the largest value of V , = 730 mV. However, F F is obviously lower than those obtained for the film obtained at pH 1.5. In some cases, we also observed a rapid decrease in J, of a closed circuit, which may be caused by the unexpected reverse electron transport that weakened the charge separation or the limited redox reaction rate that reduced electron mobility. Table I lists the results of this study and those from representative studies reported in the l i t e r a t ~ r e . ~A- ' device ~ with the electrode prepared at pH 1.5 demonstrated conversion efficiencies similar to that obtained by electrosynthesis" or hydrothermal crystallization." Note that either electrosynthesis or hydrothermal crystallization was conducted using a suspension containing P-25 Ti02 nanoparticles.lO*" The film obtained by electrosynthesis was annealed at 450°C. Considering the film thickness, the conversion efficiency of the present film ( 2-3 N

Fig. 7. Transmission electron microscope graphs ofTiOz powders after annealing of those obtained under different pH; (A) pH 1.0, (B) pH 1.5, (C) pH 2.0; the scale bars are 50 nm.

Membranes, Films, and Coatings

193

<

0

nH 1.5

200

400

600

800

Biaslm V

Fig. 9. Current-voltage curves of dye-sensitized solar cells using the nanoporous T i 0 2 as electrodes.

in solar cells.14 The preparation of arrayed Ti02 nanotube films with sufficient thickness is considered to be an effective means to obtain high-efficiency solar cells. Fig. 8. High-resolution transmission electron microscope graph of Ti02-derived powders from a solution of pH 1.0.

pm, measured by cross-sectional SEM) is relatively high, although the overall conversion efficiency is still low compared with the P-25-cast Ti02 electrode.“ The high conversion efficiency of a nanoporous thin film is partly because of its high surface areas, which results in a large amount of dye loading. It may also be attributed to the nanosheet structure of the film. Unlike nanoparticles, interconnected nanosheets would serve as current paths to give rise to smooth electron transport through the film. A device using Ti02 prepared at pH 2.0 gave a relatively low conversion efficiency of 2.2%. Further improvement of the conversion efficiency may be achieved by increasing the film thickness and optimizing microstructures such as the arrangement of nanosheets, the dimension of the voids, and the interface between the electrode film and the substrate. Mor et al. have recently reported a method to obtain porous titania films, in which Ti thin films, 400-1000 nm thick, fabricated by radio frequency sputter deposition are anodized in an electrolyte containing acetic acid and hydrofluoric acid to form optically transparent films of highly ordered titania nanotube a r r a y ~ , ’ ~ and this highly porous Ti02 film was found to be promising

Table I. Comparison of Device Performance of Ti02-Based Dye-Sensitized Solar Cells (DSSCs) Fabricated by the Current Method and those Reported Ti02 film

Thickness J, V, (pm) (mA/an2)(mV) F F q (%)

pH 1.0 (PCD)+ 3 3.9 3 13 pH 1.5 (PCD) 2 6.3 pH 2.0 (PCD) Casting of Unknown 1.26 TiOz tubes Electrosynthesis 7.8 10.5 Hydrothermal crystallization Casting of P-25

18

10

632 0.53 662 0.37 730 0.5 704 0.66

Reference

1.3 3.1 2.2 2.9

This study This sudy This study Uchida et a1.9 690 0.57 4.1 Yamamoto et a/.‘’ 9.4 710 0.63 4.2 Zhan et al. 18.2 720 0.72 10 Nazeeruddin et al.“

8

‘PCD refers to peroxotitanium complex deposition, the present process. FF, fill factor.

194

IV.

Conclusion

In summary, we report a simple, controllable process for the preparation of porous Ti02 films. The porous characteristics of our films suggest potential applications to DSSCs, photocatalysts, self-cleaning materials, and bioactive interfaces. DSSCs using Ti02 nanoporous films demonstrated high performance, namely an overall conversion efficiency of 3.1 YO(AM 1.5) for a 3-pm thick film. To further optimize the film properties, efforts are being made to address the following questions, including the in situ crystallization, preparation of one-dimensional arrays, increasing the film thickness, and optimizing the adherence of films to substrates. Moreover, the method proposed in this paper could lead to new approaches for the creation of porous thin films of other ceramic materials through the design of specific gas-generated routes.

References ‘A. Hagfeldt and M. Gratzel, “Light-Induced Redox Reactions in Nanocrystalline Systems,” Chem. Rev., 95, 45 (1995). ’(a) E. L. Crepaldi, G. Soler-Illia, D. Grosso, F. Cagnol, F. Rihot, and C. Sanchez, “Controlled Formation of Highly Organized Mesoporous Titania Thin Films: From Mesostructured Hybrids to Mesoporous Nanoanatase TiO,,” J. Am. Chem. Sac., 125, 977&86 (2003). (b) P. Yang, D. Zhao, D. I. Margolese, B. F. Chmelka, and G. D. Stucky, “Generalized Syntheses of Large-Pore Mesoporous Metal Oxides With Semicrystalline Frameworks,” Nature, 3%, 152-5 (1998). (c) D. Grosso, G. J. A. A. Soler-Mia, F. Bahonneau, C. Sanchez, P.-A. Alhouy, A. Brunet-Bruneau, and A. R. Balkenende, “Highly Organized Mesoporous Titania Thin Films Showing Mono-Oriented 2D Hexagonal Channels,” Adv. Muter., 13, 108S90 (2001). (d) S. M. Liu, L. M. Gan, L. H. Liu, W. D. Zhang, and H. C. Zeng, “Synthesis of SingleCrystalline Ti02 Nanotubes,” Chem. Muter., 14, 1391-7 (2002). (e) P. Hoyer, “Formation of a Titanium Dioxide Nanotube Array,” Langmuir, 12, 1411-13 (1996). (f) Z. R. Tian, J. A. Voigt, J. Liu, B. Mckenzie, and H. Xu, “Large Oriented Arrays and Continuous Films of Ti02-Based Nanotubes,” J. Am. Chem. Soc., 125[41] 123845 (2003). ’S. Mann, Biomimetic Muterials Chemisrry, pp. 140. VCH Publishers, New York, 1996. ‘(a) Y.-F. Gao, Y. Masuda, Z.-F. Peng. T. Yonezawa, and K. Koumoto. “Room Temperature Deposition of a Ti02 Thin Film from Aqueous Peroxotitanate Solution,” J . Muter. Chem., 13,608-13 (2003). (h) Y.-F. Gao, Y. Masuda, and K. Koumoto, “Micropatterning of Ti02 Thin Film in an Aqueous Peroxotitanate Solution,” Chem. Mater., 16, 1062-7 (2004). 5(a) H. Colfen and S.-H. Yu, “Biornimetic Mineralization/Synthesis of Mesoscale Order in Hybrid Inorganic-Organic Materials Via Nanoparticle Self-Assembly,” M R S BUN.,30, 727-35 (2005). (h) Y.F. Gao and K. Koumoto, “Bioinspired Ceramic Thin Film Processing: Present Status and Future Perspectives”. Crys. Growth Des., 5 [5] 1983-2017 (2005). 6(a) M. Kakihana, M. Tada, M. Shiro, V. Petrykin, M. Osada, and Y. Nakamura, “Structure and Stability of Water Soluble (NH4)R[Ti4(C,H,0,)4(02)4]. 8H20,” Inorg. Chem., 40,891-4 (2001). (h) J. M. Muller and G. Schwarzenhach, “The Peroxo Complexes of Titanium,” Inorg. Chem., 9, 2381-90 (1970).

Progress in Nanotechnology: Processing

’R. C. Thompson, “Oxidation of Peroxotitanium(1V) by Chlorine and Cerium(lV) in Acidic Perchlorate Solution,” Inorg. Chem.. 23 [I31 17948 (1984). ‘(a) F. F. Lange, “Chemical Solution Routes to Single-Crystal Thin Films”, Science, 273,903(1996). (b) J. G. Yu. J. C. Yu, M. K.-P. Leung, W. Ho, 9. Cheng, X. Zhao. and J. Zhdo. “Effects of Acidic and Basic Hydrolysis Catalysts on the Photocatalytic Activity and Microstructures of Bimodal Mesoporous Titania,” J. CutuL. 217 [I]69-78 (2003).(c) J. G . Yu. X. Zhao. and Q. Zhao, “Effect of Surface Structure on Photocatalytic Activity of TiOz Thin Films Prepared by Sol-Gel Method,” Thin Solid Films, 379,7-14 (2000). ’S. Uchida, R. Chiba, M. Tomiha, N. Masaki, and M. Shirdi, “Application of Titania Nanotubes to a Dye-Sensitized Solar Cell,” Elecfruchemislry, 70 [6]52 (2002). ‘OJ. Yamamoto. A. Tan, R. Shiratsuchi,S. Hayase, C. R. Chenthamarakshan. and K. Rajeshwar, “A 4% Efficient Dye-Sensitized Solar Cell Fabricated from Cathodically Electrosynthesized Composite Titania Films,” Adv. Muter., 15, 1823 (2003).

Membranes, Films, and Coatings

“D. Zhang, T. Yoshida, and H. Minoura, “Low-Temperature Fabrication of Efficient Porous Titania Photoelectrodes by Hydrothermal Crystallization at the Solid/Gas Interface,” Adv. Muter., 15, 8147 (2003). 12M. K. Nazeeruddin, A. Kay, I. Rodicio, R. Humphry-Baker, E. Muller, P. Liska, N. Vlachopoulos, and M. Gratzel, “Conversion of Light to Electricity by cis-X2bis(2,2‘-bipyridyl-4,4-dicarboxylate)Ruthenium(ll) Charge-Transfer Sensitizers (X=CI-, Br-, I-, CN-, and SCN-) on Nanocrystalline TiOz Electrodes,” J. Am. Chem. Soc., 115 [I416382 (1993). I3G. K. Mor, 0. K. Varghese, M. Paulose, and C. A. Grimes, “Transparent Highly Ordered Ti02 Nanotube Arrays Via Anodizdtion of Titanium Thin Films,” Adv. Func. Mater., 15, 12914 (2005). l4G. K. Mor, K. Shankar, M. Paulose, 0. K. Varghese, and C. A. Grimes, “Enhanced Photocleavage of Water Using Titania Nanotube-Arrays,” Nuno Left., 5, 191-5 (2005).

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Nano-Sized Hydroxyapatite Coatings on Ti Substrate with TiOz Buffer Layer by E-beam Deposition Su-Hee Lee and Hyoun-Ee Kim*,' School of Materials Science and Engineering, Seoul National University, Seoul, Korea

Hae-Won Kim* Department of Dental Biomaterials, School of Dentistry, Dankook University, Cheonan, Korea

A nano-sized hydroxyapatite (HA) layer was coated on a Ti substrate with a titanium oxide (Ti02) buffer layer by the electron-beam deposition method. The morphological features as well as the mechanical and biological properties of the HA/ TiO2-layered coating were noticeably different from those of a conventional HA coating on Ti. The HA on the T i 0 2 layer replicated the fine grain structure of the Ti02 layer, with grain sizes of just a few tens of nanometers. The Ti02 buffer layer was highly effective in preserving the adhesion strength of the coating layer following the heat treatment at 500"C, which was necessary to crystallize the structure. Moreover, in contrast to the HA single coating wherein severe cracking was observed under moist conditions, the HA/Ti02 coating retained its mechanical stability under the same conditions. The dissolution of the HA/Ti02 coating in a physiological saline solution exhibited a more favorable pattern than that of the HA single coating, with a reduced initial burst and a subsequent steady release rate. Preliminary in virro cellular tests showed that osteoblastic cells expressed a significantly higher alkaline phosphatase level on the HA/Ti02 coating than on the HA single coating. Conclusively, the nano-sized HA coating with the T i 0 2 buffer layer holds great promise as a bioactive coating system.

I. Introduction

T

ITANIUM-BASED metals (Ti and its alloys) are currently the most widely used implant materials in orthopedics and dentistry. To improve the biocompatibility of these implant materials, much effort has been made over the last decade to modify the physical and chemical characteristics of their surfaces through various methods, such as sand blasting, chemical etching, oxidation, and by coating them with bioactive materials. Among the coating materials developed thus far, titanium oxide (generally Ti02) has been studied most widely. This is because the biocompatibility of Ti is known to be due to the thin oxide layer ( < 10 nm) that exists under ambient moist conditions. The thickness of this oxide layer is crucial in determining the longterm biocompatibility of the implants, so that in an attempt to produce a satisfactory oxide layer, a large number of studies have been conducted using techniques such as thermal oxidation, anodic oxidation, sol-gel deposition, and chemical vapor deposition.'-3 T. Trouynsk-contrihuting editor

Manuscript No. 21453. Received February 3,2006; approved August 16, 2006. This resedrch was supported by a grant (code no.: OSKIWI-01510) from the Center for Nanostructured Materials Technology under "2 1st Century Frontier R&D Programs" of the Ministry of Science and Technology, Korea. 'Member, the AWncdn Ceramic Society. 'Author lo whom correspondence should be addressed. e-mail: [email protected]

Membranes, Films, and Coatings

Along with Ti02, calcium phosphate (mainly hydroxyapatite (HA)) coatings have been studied extensively over the past decade, in order to improve the direct contact of bone-implants and shorten the healing period."' Conventionally, the plasmaspraying technique has been utilized most commonly in order to create a bioactive HA layer on Ti. Studies of these plasmasprayed HA coatings have demonstrated their excellent performance both in vitro and in vivo,which is mainly attributable to the bioactivity and osteoconductivity of the HA layer and the direct apposition of the coatings to the hard tissues without the formation of fibrous tissues. However, recent studies have unveiled significant clinical problems with plasma-sprayed HA coatings, particularly with regard to their long-term usage under load-bearing conditions, wherein various deficiencies were observed such as the fragmentation and breakdown of the coating layer and its mechanical instability, leading to an acute inflammatory reaction. This was attributed to the fact that the coating layer was too thick (- 5&200 pm thickness) and porous (- 10%-20% porosity), in addition to being inhomogeneous in terms of its phase and structure, due to the extremely high processing temperatures that were used. In an attempt to resolve these difficulties, various coating techniques have been developed to produce much thinner HA layers (less than a few micrometers) with a more homogeneous phase composition and structure, such as the sol-gel depositi0n,8.~ electrochemical deposition, and electron-beam (ebeam) evaporation methods6 The e-beam-deposited apatite films (HA and its fluoridated form) were shown to be very thin < 1 pm) and to have dense and homogeneous structure~!,~ Furthermore, this technique is versatile, allowing the coating composition and properties to be tailored, by modifying the target material and equipment variables. Generally, in the case of e-beam or sol-gel coatings, a subsequent thermal treatment process is necessary to crystallize the films. However, because of the difference in the thermal properties between the coating and the substrate, high thermal stress and consequent severe cracking are generated within the coating layer, resulting in mechanical instability and coating failure. Consequently, a great deal of effort has been made in an attempt either to relieve the thermal stress by changing the composition and thickness of the film or to decrease the crystallization temperature required to confer mechanical stability on the films." In this study, the authors aimed to produce a stable HA thin film on a Ti substrate by using a Ti02 buffer layer between HA and Ti. The present study was motivated by our previous observation of an HA/Ti02 sol-gel-layered coating on Ti, wherein the T i 0 layer improved the bonding strength of HA with respect to Herein, both the HA and Ti02 layers were produced by means of the e-beam deposition technique. The HA film deposited over the Ti02 layer exhibited an ultrafine grain structure (< 100 nm), which has not been previously reported in the literature. We report on the effect of the Ti02 buffer layer on the morphological features, as well as on the mechanical stability and in vitro biological properties of the HA coating on the Ti substrate. 197

11. Experimental Procedure

(1) Preparation of HA/ TiOz-Layered Coatings Commercially pure samples of titanium (cp Ti) with dimensions of 10 mm x 10 mm x 2 mm were prepared as the substrate and then polished with a 1000-grit Sic abrasive paper. The Ti substrates were rinsed ultrasonically in acetone, ethanol, and distilled water successively for 10 min each. The substrates were placed on a holder in a vacuum chamber and then cleaned and etched by Ar ion bombardment with an ion beam at 90 V and 2 A for 30 min. As the target for the Ti02 coating, commercially available Ti02 powder (Sigma-Aldrich Chem. Co., Milwaukee, WI) was used after sintering at 1300°C for 3 h. The Ti02 was then deposited on the pretreated Ti substrate at a rate of 0.1 nm/ s. During the whole process, the substrate holder was rotated at 20 rpm to achieve a uniform coating morphology. After depositing the oxide layer up to a thickness of about 400 nm, it was heat treated in air at 500°C for 2 h to crystallize the oxide layer. The HA layer was subsequently deposited on the Ti02 layer. As the target for the HA coating, commercially available HA powder (Alfa-Aesar Co., Ward Hill, MA) was mixed with 25 wt% of CaO (Sigma-Aldrich Chem. Co.), and the resulting mixture was sintered at 1200°C for 3 h. The composition of the mixture was experimentally determined to obtain a pure phase of HA film after heat The apatite target was evaporated onto the Ti02 precoated Ti substrate by the e-beam process at a deposition rate of 0.1 nm/s. The final thickness of the HA layer was about 700 nm. The HA/Ti02coating obtained by the e-beam process was annealed at different temperatures (450"4WC) for 2 h in air. For the purpose of comparison, an HA film was deposited on a Ti substrate without the Ti02 buffer layer under the same e-beam evaporation conditions and subsequently heat treated at the same temperatures. In particular, the thickness of the HA single film was adjusted so as to be similar to that of the HA layer in the HA/Ti02 film. ( 2 ) Characterizations X-ray diffraction (model MISXHF, MAC Science, Tokyo, Japan) was used to characterize the phase and crystallinity of the films, using a CUKEsource at 40 kV and 200 mA within a scan range of 28 = 20-50" with a step size of 0.5". Scanning electron microscopy (SEM) (using field emission electron sources-JSM-6330F, JEOL,Tokyo, Japan) was used to examine the surface morphology and microstructure of the films. The grain morphology of the films determined by SEM was analyzed to calculate the average size of the grains. The adhesion strength of the coating layer was measured using an adhesion test apparatus, as described previously.' I Before the tensile tests, 42.69 mm x 12 mm aluminum pull studs (Quad Group, Spokane, WA), precoated with a thermal-curing epoxy, were adhered to each coating surface and cured at 150°C for 1 h. The cured studs were pulled until failure occurred and the maximum strength was measured by taking into consideration the failed surface area. SEM and energy-dispersive spectroscopy (EDS) analyses of the failed samples were conducted, in order to determine the failure mechanism of the coating layer. In order to analyze the dissolution behavior of the coating layers, the HA and HA/Ti02 coatings (10 mm x 10 mm x 1 mm) were placed in vials containing 10 mL of physiological saline solution. They were maintained at 37°C for periods of up to 12 days, with the medium being refreshed daily. After predetermined periods of time, the medium was analyzed with inductively coupled plasma atomic emission spectroscopy (ICP-AES; ICPS-7500, Shimadzu, Kyoto, Japan) to determine the concentration of dissolved Ca2+ ions. (3) Cellular Assays Osteoblast-like MG63 cells were cultured in a total medium as described previously. The cells pooled in the medium were seeded on the films at a cell density of 1 x lo4 cells/mL and incubated at 37°C in an atmosphere of 5% co2/95% humidity. The 198

attachment and proliferation of the cells on the films were observed with SEM after the fixation and dehydration processes. The cell proliferation rate was measured using a hemocytometer after exclusion with trypan blue. The osteoblastic functional activity of the cells was determined by measuring the alkaline phosphatase (ALP) expression level, by means of an enzymatic reaction using a p-nitrophenyl phosphate substrate (SigmaAldrich Chem. Co.)." The absorbance was measured at 410 nm using a spectrophotometer and the ALP activity was calculated from a standard curve, after normalizing the results to the total protein content (BioRad, Hercules, CA). Each set of tests was performed in triplicate, and the data were normalized by taking the surface area into consideration. 111. Results

( 1 ) Phase and Microstructure The HA coatings both without and with the Ti02 buffer layer started to crystallize when the heat treatment temperature reached 450°C. The XRD patterns of the HA coatings on Ti without and with the Ti02 buffer layer following their heat treatment at 500°C are representatively shown in Fig. 1. The HA peaks that were observed were similar for both types of films. Moreover, additional Ti02 anatase peaks were also observed, both in the case of the HA/Ti02 coating and the HA single coating on Ti. There were no other peaks except for those corresponding to the HA and Ti02 anatase phase. Based on these XRD results, the Ti02 buffer layer is considered to have little effect on the phase evolution and crystallization temperature of the HA coating. In contrast to the phase evolution, the microstructure of the HA coating was significantly influenced by the presence of the Ti02 buffer layer. The SEM morphologies of the coatings following their heat treatment at 500°C are presented in Fig. 2. The HA single coating on Ti showed the typical grain morphology of apatite with a grain size of approximately 300 nm, as is normally observed in the case of HA coated by the e-beam deposition method (Fig. 2(A)). On the other hand, the HA coating produced above the Ti02 layer did not reveal such a grain morphology at normal magnification (Fig. 2(B)). At a higher magnification ( x 50 OW),however, the ultrafine grain morphology of the HA layer was revealed (Fig. 2(D)). Interestingly, this grain morphology of the HA layer appears to reflect that of the underlying Ti02 layer (Fig. 2(C)). The grain size of the coating layers was measured by conducting an image analysis of the SEM micrographs, as presented in Fig. 3. The grain sizes of the Ti02-buffered HA coating were significantly smaller (from 40 to 100 nm on average) than those

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Fig. 1. X-ray diffraction patterns of the (A)HA and (B) HA/Ti02 coating layers on Ti substrates by the electron-beam process following heat treatment at 500°C for 2 h. HA, hydroxyapatite; TiOz, titanium oxide.

Progress in Nanotechnology: Processing

Fig. 2. Scanning electron microscopy photographs of the electron-beam-depositedcoating layers on Ti substrates: (A) HA, (B) HA/Ti02, (C) TiOz, and (D) high magnification of (B). The films were heat treated at 500°C for 2 h. HA, hydroxyapatite; TiO2. titanium oxide.

of the HA layers directly deposited on Ti (from 250 to 350 nm on average) at all the crystallization temperatures (500"-600°C). In particular, the grain sizes of the Ti02/HA coating layers were similar to (or only slightly higher than) those of the Ti02 layers on Ti, suggesting that the underlying Ti02 layer played a crucial role in determining the grain size of the HA layer deposited on it.

(2)

Adhesion Strength

The adhesion property of the HA/TiO24ayered coating with respect to the Ti substrate was assessed by means of a debonding

400

1

tensile strength test. Figure 4 shows the tensile strength of the coatings with respect to the crystallization temperature. Compared with the HA single coating, the HA/Ti02 layered coating showed a significantly higher strength at all the crystallization temperatures. The bonding strength of the HA single coatings was markedly decreased by the crystallization process, and this reduction became more pronounced as the crystallization temperature increased. This was mainly due to the increase in the thermal mismatch between HA and Ti induced by the crystallization process. On the other hand, the HA/Ti02-layered coatings retained their initial bonding strength values after the heat-treatment.

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Fig.3. Grain sizes of each coating layer represented with respect to the

Fig.4. Adhesion strength of each coating layer without and with heat

heat treatment temperature.

treatment at different temperatures.

Membranes, Films, and Coatings

199

““1

HA/Ti02

Fig.6. Changes in the adhesion strength of the coating layers on Ti substrates following their immersion in physiological saline solution for 1 day. The strength was preserved in the HA/Ti02 coating, whereas it was not measurable in the HA single coating due to severe cracking. HA, hydroxyapatite; Ti02, titanium oxide.

The dissolution behavior of the HA/Ti02 coating was assessed by measuring the concentration of Ca2+ dissolved into a physiological saline medium after incubation at 37°C for periods of up to 12 days, as shown in Fig. 8. In the case of the HA single coating, the Ca2+was released very rapidly within a short period of incubation time (-2 days), and then very slowly thereafter. Fig.5. Scanning electron microscopy images of the (A) HA and (B) HA/TiOZ coating layer after the adhesion test, showing the associated failure modes (T, at HA/Ti interface; TO, at Ti02/Ti interface or within Ti02; H, at HA/Ti02 interface or within HA, and E, epoxy). HA, hydroxyapatite; TiOz, titanium oxide.

The failure surface of the coating layer was observed with SEM, as presented in Fig. 5 , and the composition of the failure surface was analyzed with EDS. The HA single coating on Ti exhibited failure associated mainly with the interface of the HATi substrate (note the large area designated as “T,” wherein only Ti was detected). However, the HA/Ti02-layered coating showed a failure mode associated mainly with the HA/TiO2 interface (note the area designated as “TO,” wherein Ti and a large amount of 0 were detected). Based on these results, the adhesion of the HA layer with respect to the Ti02 layer is considered to be stronger than that of the HA layer with respect to Ti.

( 3 ) Stability Under Moist Condtions The beneficial effect of the Ti02 buffer layer became more manifest in the examination of the stability under moist conditions. As shown in Fig. 6, the HA/Ti02-layered coating heat treated at 500°C retained its initial adhesion strength after immersion in a saline solution for 1 day. However, in the case of the HA single coating heat treated under the same conditions, severe cracks were generated, thus making it impossible to carry out the strength test. The surface morphologies of the coatings after immersion are shown in Fig. 7. While severe cracking was generated on the HA coating (Fig. 7(A)), there were no such cracks on the HA/Ti02 coating (Fig. 7(B)). Interestingly, the initial grain morphology of the HA single Coating disappeared completely after the hmersion test (cf. Fig. 2(A)). Based on these results, the HA/TiO~layeredcoating is considered to have improved mechanical stability under moist conditions. 200

Fig.7. Scanning electron microscopy image of the (A) HA and (B) HA/Ti02 coating layers on Ti substrates after immersion in a physiological saline solution for 4 days, showing severe cracking in the HA single coating. HA, hydroxyapatite; TiO,, titanium oxide.

Progress in Nanotechnology: Processing

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Fig. 10. Proliferation and alkaline phosphate (ALP) activity of the cells cultured on each coating for 3 and 7 days, respectively. Data are represented with respect to the Ti substrate. Statistical significance (P
IV. Discussion On the other hand, in the case of the HA/Ti02-layered coating, the initial burst release of Ca2+ was significantly reduced, and the release rate at prolonged periods was higher (-0.30 ppm/ day over 2-12 days) than that of the HA single coating (-0.15 ppmlday). Consequently, the HAP102 coating is considered to have higher chemical and mechanical stabilities than the HA single coating over the testing period used in this study. (4) In Vitro Cellular Responses The cellular responses to the HA/Ti02 coating were assessed using a human bone derived cell line. Figure 9 shows a typical electron micrograph of the cells cultured on the layered coating for 3 days. The cells were shown to grow and spread well, being in intimate contact with the surface. A similar cell growth morphology was observed on the HA coating (not shown here). The cell proliferation and the ALP level expressed by the cells were measured, after culturing for 3 and 7 days, respectively. The numbers of cells cultured for 3 days on the HA and HA/ TiOz coatings were higher than that on pure Ti, as shown in Fig. 10. However, no significant difference was observed between the HA and the HA/Ti02 coatings. The ALP activities of the cells cultured for 7 days on the HA and HA/Ti02 coatings were higher than that on the Ti substrate. However, in contrast to the cell proliferation, the ALP activity of the cells on the HA/ Ti02 coating was significantly higher than that on HA single coating (p<0.05).

Fig.9. Typical morphology of the osteoblastic cells grown on the HA/ Ti02 coating at I day after culturing. HA, hydroxyapatite; Ti02, titanium oxide.

Membranes, Films, and Coatings

The most remarkable change observed in the HA coating fabricated on top of the Ti02buffer layer by e-beam deposition was that of the grain morphology. As shown in the SEM image with high magnification, the grain sizes of the HA films deposited on the TiOz layer were brought down to the -tens of nanometers range after the crystallization process performed at temperatures ranging from 500" to 6Oo"C, which is in marked contrast to those of the HA films directly deposited on Ti (25&350 nm). Clearly, the intermediate TiOz layer influenced the evolution of the grain morphology of the HA film, as confirmed by the fact that its grain size was similar to that of the underlying Ti02 layer. It is inferred from this that the underlying Ti02 layer acts as a nucleation site for the apatite crystallization of the overcoated HA layer during the thermal treatment process. In this study, the HA layer was coated over the crystallized Ti02 layer. Therefore, it is reckoned that the presence of the crystallized Ti02 anatase phase favors the molecular ordering of the randomly distributed Ca and P ions, leading to the development of an apatite structure. This is comprehensible to some extent when considering the crystallography of the anatase Ti02 and HA, wherein the HA lattice matches well with a specific plane of the TiOz anatase structure, as has been previously suggested in the case of apatite nucleation onto anatase Ti02 under in vitro body solution conditions.13 In practice, the HA coating layer followed the grain morphology of the underlying Ti02 layer very well (see Fig. 3). Of particular note is the fact that the ultrafine grain morphology developed herein has never been achieved by directly coating HA on a Ti substrate. The ultrafine structure of the HA coating was observed to significantly enhance its mechanical and biological properties. The mechanical properties of the HA/Ti02-layered coating on Ti were assessed in terms of the adhesion strength of the coating layer and its stability under moist conditions. Because coatings deposited by e-beam evaporation are amorphous in the as-deposited state, a subsequent heat-treatment process is normally required to crystallize the structure. In the case of HA coatings, the crystallization procedure is essential to preserve their chemical stability, thereby preventing the abrupt dissolution of the coating and inducing favorable cellular responses. In practice, as-deposited calcium phosphate films without heat treatment have been shown to dissolve quickly (within a few hours or a day) and to elicit unfavorable cellular responses. During thermal treatment, however, concerns are often raised as to the generation of residual stress or even cracks, due to the thermal mismatch between the film and the substrate. As shown in Fig. 4, the adhesion strength of the HA single coating was observed to decrease significantly after the heat-treatment for crystallization, and the mechanical degradation became more severe as the heat treatment temperature increased. On the other 201

hand, in the case where the Ti02 buffer layer was used between HA and Ti, the adhesion strength of the coating layer did not change significantly following the crystallization process. Given that the thickness of the HA layer was similar in both cases (HA/Ti02 and HA), it can be inferred that this improvement in the adhesion strength of the coating layer was not caused by any variation in the thickness of the HA layer. Rather, it is believed that the intermediate Ti02 layer improves the bonding property of the HA film with respect to the substrate (herein Ti02-coated Ti).I4 Of special note, in the case of the HA single coating on Ti, is that the failure occurred mostly at the HA/Ti interface (Fig. 5(A)). However, the dominant failure mode observed in the HA/Ti02system was detachment at the HA/Ti02 interface, not at the TiOz/Ti interface (Fig. S(B)). Taken together, it is reasoned that the strengths of the interfaces are in the order: HA/Ti
ratio of 1.67. Therefore, as the grain growth of HA proceeds, the composition of the grain boundary phase is likely to be much different from that of HA. Actually, when a thin HA coating is immersed in a physiological saline solution, the dissolution of the HA is always initiated from the grain boundary region. Therefore, the small grains in the present HA/Ti02 coating are more likely to have a composition similar to that of the grain boundary phase than the large grains in the HA single coating. The cellular responses to the HA coating with the Ti02 buffer layer were significantly different from those to the HA coating directly deposited on Ti, particularly with regard to the ALP expression level, although the cell viabilities of the two types of coatings were similar. The significant differences in the grain morphology and dissolution behavior of the two types of coating would be expected to be reflected in their cellular responses. It has been reported that the initial cellular attachment on nanostructured substrates, including those made of ceramics and metals, is improved through a mechanism that is associated with the higher adsorption rate of the cell-adhesive proteins2 As such, the ultrafine-grained HA coating would be expected to have a more favorable interaction with biological molecules and provide superior cellular adhesion. However, the anticipated improvement in the cell adhesion on the ultrafine-grained HA coating was not observed in this study. Nevertheless, the level of osteoblastic phenotype ALP, which is one of the most important markers for osteoconductivity, was stimulated to a greater extent on the ultrafine-structured HA coating. In this work, the mechanism underlying this phenomenon was not able to be elucidated; however, it is reasoned that the morphological benefit and more stable release of the HA coating deposited on the Ti02 layer played a positive role in the expression of ALP by the cells. To elucidate the reasons for the effectiveness of the HA/Ti02 layer at the molecular and cellular level, further in vifro experiments are warranted.

V. Conclusions The Ti02-buffered HA coating on Ti produced by e-beam deposition was observed to have excellent properties, such as a nano-sized grain structure, good adhesion coating property, high mechanical and chemical stability in vifro, and a superior osteoblastic cellular response. The present findings suggest that the TiO2-buffered HA coating has good potential to be used as an alternative to the direct deposition of an HA film on Ti.

References ‘M. P. Casaletto, G . M. Ingo, S . Kaciulis, G. Mattogno. L. Pandolfi, and G . Scavia, “Surface Studies of In Vitro Biocompatihility of Titanium Oxide Coatin s,” Appl. Surf. Sci., 172, 167-77 (2001). ‘G. Giavaresi, L. Amhrosio, G . A. Battiston, U. Casellato, R. Gerbasi, M. Finia, N. N. Aldini. L. Martini, L. Rimondini, and R. Giardino, “Histomorphometric, Ultrastructural and Microhardness Evaluation of the Osseointegration of a Nanostructured Titanium Oxide Codling by Metal-Organic Chemical Vapour Deposition: An In Vivo Study,” Biomarerials, 25, 5583-91 (2004). ’ Y . M. Zhang, P. Bataillon-Linez,P. Huang. M. Zhao, Y . Han, M. Traisnel, K. W. Xu, and H. F. Hildehrand, “Surface Analyses of Micro-Arc Oxidized and Hydrothermally Treated Titanium and Effect on Osteohlast Behavior,” J. Biomed. Muter. Res.. 68A, 383-91 (2004). 4X. X. Wang, S. Hayakawa, K. Tsuru, and A. Osaka, “A Comparative Study of In Vitro Apatite Deposition on Heat-, H202-,and NaOH-TreatedTitanium Surfaces,” J. Biomed Mafer. Res., 54. 172-78 (2001). ’L. Can and R. Pilliar, “Calcium Phosphate Sol-Gel-Derived Thin Films on Porous-Surfaced Implants for Enhanced Osteoconductivity.Part 1: Synthesis and Characterization,”Biomaferiab,25, 5303-1 2 (2004). 6J. M. Choi, H. E. Kim, and I. S. Lee, “Ion-Beam-AssistedDeposition (IBAD) of HydroxyapatiteCoating Layer on Ti-Based Metal Substrate,”Biomaferials,21, 469-73 (2000). ’A. E. Porter, L. W. Hohhs, V. B. Rosen, and M. Spector, “The Ultrastructure of the Plasma-Sprayed Hydroxyapatite-BoneInterface Predisposing,” Biomaferials. 23, 725-33 (2002). ‘M. Sato, E. B. Slamovich, and T. J. Wehster, “Enhanced Osteohlast Adhesion on Hydrothermally Treated Hydroxyapatite/Titania/Poly(Lactide-co-Glycolide) SolGel Titanium,’’ Biomoferinls, 26, 1349-57 (2005).

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’P. A. Ramires, A. Wennerberg, C. B. Johansson, F. Cosentino, S. Tundo, and E. Milella, “Biological Behavior of SolGel Coated Dental Implants,” J. Muter. Sci. Mafer. Med., 14, 53945 (2003). ’“E. J. Lee, S. H. Lee, H. W. Kim, Y. M. Kong, and H. E. Kim, “Fluoridated Apatite Coatings on Titanium Obtained by Electron-Beam Deposition,” Biomarerink, 26, 3843-51 (2005). “H. W. Kim, H. E. Kim. V. Salih, and I . C. Knowles, “Hydroxyapatite and Titania SoCGel Composite Coatings on Titanium for Hard Tissue Implants; Mechanical and In Vitro Biological Performance,” J. Biomed. Muter. Res. Purl B: Appl. Biomater., 728, 1-8 (2005). 2H. W. Kim, Y. H. Koh, L. H. Li, S. Lee, and H. E. Kim, “Hydroxyapatite Coating on Titanium Substrate with Titania Buffer Layer Processed by SoLGel Method,” Biomaferials. 25, 2533-38 (2004). I’S. H. Oh, R. R. Finones, C. Daraio. L. H. Chen. and S. Jin, “Growth of Nano-Scale Hydroxyapatite Using Chemically Treated Titanium Oxide Nanotubes,” Biomarerials, 26, 4938-43 (2005). ‘‘Y. P. Lu, M. S. Li, S. T. Li. Z. G. Wang, and R. F. Zhu, “Plasma-Sprayed Hydroxyapatite+Titania Composite Bond Coat for Hydroxyapatite Coating on Titanium Substrate,” Biomarerials, 25, 4 3 9 H 3 (2004). ‘9.J. Webster, C. Ergun, R. H. Doremus, and W. A. Lanford, “Increased Osteoblast Adhesion on Titanium-Coated Hydroxylapatite That Forms CaTiO,,” J . Biomed. Mafer. Res.A, 61A,975-80 (2003).

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I‘M. Manso, M. Langlet, and J. M. Martinez-Duart, “Testing Sol-Gel CaTiO, Coatings for Biocompatible Applications,” Muter. Sci. Eng. C, 23 (3) 447-50 (2003). ”R. R. Kumar and M. Wang. “Functionally Graded Bioactive Coatings of Hydroxyapatite/Titanium Oxide Composite System,” Mafer. Lerr., 55, 133-7 (2002). “V. Nelea, C. Ristoscu, C. Chiritescu, C. Ghica, I. N. Mihailescu. H. Pelletier, P. Mille, and A. Cornet, “Pulsed Laser Deposition of Hydroxyapatite Thin Films on Ti-5Al-2.5Fe Substrates With and Without Buffer Layers,” Appl. Su$ Sci., 168. 127-31 (2000). ”B. T. Chou and E. Chang, “Plasma-Sprayed Zirconia Bond Coat as an Intermediate Layer for Hydroxyapatite Coating on Titanium Alloy Substrate,” J. Mafer. Sci. M a w . Med., 13, 589-95 (2002). 2oF. Suska, C. Gretzer, M. Esposito, P. Tengvall, and P. Thomsen, “Monocyte Viability on Titanium and Copper Coated Titanium.” Biomaferials, 26, 5942-50 (2005). 2’X.F. Xiao, R. F. Liu, and Y. 2.Zheng, “Characterization of Hydroxyapatite/ Titania Composite Coatings Codeposited by a Hydrothermal-Electrochemical Method on Titanium,’’ Surf. Coat. Techno/., 200, 4406-13 (2006). 22D.H. Kim, Y. M. Kong, S. H. Lee, 1. S. Lee, and H. E. Kim, “Composition and Crystallization of Hydroxyapatite Coating Layer Formed by Electron Beam Deposition,” J . Am. Ceram. SOC.,86 [l] 18G8 (2003).

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SOL-GEL ROUTES TO NANOSTRUCTURED PATTERNED FERROELECTRIC THIN FILMS WITH NOVEL ELECTRONIC AND OPTICAL FUNCTIONS Makoto Kuwabara****,',Yong Jun WU',~, Juan Li', and T. Koga' 1

Department of Applied Science for Electronics and Materials, Kyushu University, 6-1 Kasuga-kouen, Kasuga 8 16-8580, Japan 2Departmentof Materials Science and Engineering, Zhejiang University, Zheda road 38, Hangzhou 3 10027, China ABSTRACT We have successfully fabricated barium titanate (BaTi03: BTO) patterned thin films of < 1 pm in thickness on Pt/Ti/Si substrates by combining an electrophoretic deposition (EPD) method using a suspension of BTO nanoparticles and electron beam lithography, which was used to form resist molds on the substrates. A suspension of BTO nanoparticles (size=lO-15 nm) for EPD was prepared by dispersing BTO nanocrystals, which were synthesized by a sol-gel process using a high concentration alkoxides solution (1.1 mol/L), in 2-methoxyethanol added with acetylacetone BTO ceramic thin films obtained by firing the EPD green films at various temperatures between 600" and 800°C consisted of particles with diameters within 30 nm, and had a dielectric constant of about 100 and a dielectric loss < 1% and a very high dielectric strength of more than 500 kV/cm. The obtained nanostructured BTO ceramic thin films and sub-micrometer-sized ceramic pillars exhibited distinct piezoelectric butterfly curves and also clear ferroelectric hysteresis loops; this may prove that even BTO crystals of 30 nm in diameter can exhibit induced ferroelectric characteristics under high electric field. INTRODUCTION Deposition of nano-structured ceramic thin films on conducting substrates has attracted increasing attention due to novel electronic and optical functions that are provided exclusively by such materials, and moreover, from points of view of both materials science and processing technology. As one of the most important issues in electroceramics, preparation of barium titanate (BaTiO3 or BTO) thin films has been investigated extensively because of its promising electrical and optical properties. BaTi03 thin films are of significant interest for a variety of technological applications, such as multilayer ceramic capacitors (MLCC), sensors, actuators, thermistors, dynamic random access memory (DRAM) and electroluminescent elements.'-3 A variety of methods, such as sputtering, metal organic chemical vapor deposition (MOCVD), pulsed laser deposition (PLD), chemical solution deposition (CSD) have been investigated to prepare BTO

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thin film^.^-^ Compared with these methods, electrophoretic deposition (EPD) offers advantages of process simplicity, low cost, easily deposition of multilayer films of controlled thickness on complex shaped or patterned substrates.839It has been studied extensively for fabricating advanced coating, nanocomposites, laminated structures, and functional graded materials. Electrophoretic deposition of BTO thick films from micrometer or sub-micrometer-sized particles suspensions have been investigated by several groups'o-'2 and EPD of BTO thin films has been reported in our previous paper.13 Meanwhile, the continuing trend of miniaturization in the microelectronics industry requires integration of micropatterned BTO thin films into electroceramic packages. Various methods, including top-down and bottom-up approaches have been used to prepare micropatterned electroceramics. Within the top-down approach, focused ion beam patterning, electron beam direct writing, dry-etching, wet-etching, nanoimprint lithography and other lithography techniques have been widely investigated. The bottom-up approach can be classified into two main routes, viz., physical and chemical. i4;i'- However, because a relatively high sintering temperature is required to process BaTi03 ceramics by conventional methods, the development of new patterning techniques applicable to BTO is desired. In this study, a templating method using the sol-gel process, EPD and electron beam lithography was proposed to fabricate micropatterned BaTi03 thin films at a relatively low temperature. EXPERIMENTAL PROCEDURE Figure 1 shows the schematic representation of preparing BTO nano-particles and micropatterned BTO thin films. The BTO nano-particles were prepared by a high-concentration sol-gel High purity barium diethoxide (Ba(OC2H5)2, >99%) was dissolved in a mixed solvent of methanol (CH30H, >99%) and 2-methoxyethanol (CH30C2&0H7 EGMME, >99.8%) with a volume ratio of 3:2. After stirring the solution for 24 hours, titanium tetra-iso-propoxide (Ti(OiC3H7)4, >99%) was added in it to prepare a BTO precursor solution and then stirred for another 24 hours. The concentration of the precursor solution was 1.1 mol/l. A mixed solution of distilled water and EGMME with a volume ratio of 1:1 was used for hydrolysis. The molar ratio of H20 and Ba was 1O:l. All above processes were performed in a dry N2 atmosphere. The hydrolyzed solution, after stirred at -30 "C for 10 min, was sealed in an autoclave and aged in a 90°C oven for 1 hour to allow gel formation. After discarding the liquid generated in association with the gel contraction, the BTO gel was dispersed into EGMME using sonication for several hours. Then acetylacetone (CH3COCH2COCH3, Acac, >99%) was added and again sonicated to prepare BTO suspension. The volume ratio of EGMME and Acac was fixed as 9:l and the concentration of BTO nanoparticles in the suspension was controlled to be

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0.075 mol/l.

( . . , . p . " ' ? ] [ 7 ][ T I

1

CH30H + CH30C2H40H (3:2)

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Resist films on substrate

1

Exposing Hydrolysis (-30°C)

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Rinsing Post-baking Micropatterned substrate

ultrasonicating Electrophoretic deposition Polishing

i

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1 Sintering [

Micropatterned BTO ceramic films

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Fig. 1. Schematic representation of the method used for preparing micropatterned BaTi03 ceramic films.

For preparation of micropatterned BTO thin films, electron beam resist films were first formed on Pt/Ti/Si substrates by spin-coating at 3000 rpm for 20s, followed by pre-baking on a hot-plate at 18OoC for 3 minutes. The spin-coating/pre-baking cycle was repeated several times to obtain resist films with a thickness of about 800 nm. An EB lithography system (ELS-5700, Elionex, Japan) was then used to create latent micropatterns in the resist films. After developed and baked, resist molds with different micropatterns on Pt/Ti/Si substrates were fabricated. These micropatterned substrates were used as cathodes and the same size plates of stainless steel were used as anodes in the EPD process. A schematic diagram of the EPD system employed is shown in Fig. 2. The distance between the cathode and the anode was 2 cm. A Fig. 2. Schematic diagram of EPD system.

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DC voltage of 5 V was applied for EPD. As-deposited samples were dried in a 50°C oven in air. After removing BaTi03 nanoparticles layers superfluously deposited on the resist molds by mechanical polishing, the resist molds were dissolved in a remover to obtain nanoparticles patterns left on the substrates. The nanoparticles patterns were then calcined at 600°C for 30min to yield micropatterned BaTi03 ceramic thin films. The microstructures of the sol-gel-derived BTO particles were evaluated by high resolution transmission electron microscopy (HR-TEM, EM-O02BF, Topcon, H-9000NAR, Hitachi, Japan) and X-ray diffraction (XRD, M18XHF, Mac Science, Japan). The transmittance property of a BTO suspension in the wavelength of 300 to 900 nm was measured with a UVNIS spectrophotometer (U-550, Jasco, Japan). The surface morphology and thickness of both an EB resist mold and a deposited thin film were evaluated by field emission scanning electron microscopy (FE-SEM, S5000, Hitachi, Japan); FE-SEM images were taken on the surface and cross section of each sample. Scanning probe microscopy (SPM, Nano-R system, Pacific Nanotechnology, USA) was also used to characterize the surface of a micropatterned BaTi03 thin film. The strain hysteresis loop of a single BTO pillar was measured using SPM (contact mode) equipped with a ferroelectric properties measurement system (FCE system, Toyo Technical, Japan). All the measurements were carried out at room temperature. RESULTS AND DISCUSSION Figure 3 shows the transmittance of a BTO suspension prepared in this study (slightly yellowish) with a high value in the visible region and an absorption edge near 400 nm (about 3.1 eV), indicating that the BTO particles were well dispersed and no significant size effect on the absorption edge was observed for the BTO nanoparticles. The BTO suspension was very stable and no agglomeration or sedimentation occurred for more than two months. An appropriate addition of Acac was found to be effective to obtain a suspension of well-dispersed BTO nanoparticles. The reason is explained by an improved electrostatic stabilization of BTO particles caused by adsorption of protons, generated in the keto-enol equilibrium reaction of Acac, on the particles' surfaces. A representative TEM micrograph of BTO nanoparticles prepared in the present study is shown in Fig. 4, clearly indicating that the BTO particles were well crystallized and the

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01 300

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.

' . ' . ' . ' . 500 600 700 800 900 Wavelength (nm)

Fig. 3 . Transmittance spectrum of a BTO suspension between 300 and 900 nm.

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average particle size was about 10 nm. From a selected area electron diffraction (ED) pattern of the BTO nanoparticles, shown in the inset of Fig. 4,it has been found that the BTO particles crystallized in the perovskite cubic phase. This was also confirmed from an analysis of the XRD peak broadening. By using the Scherrer equation, the crystallites size was calculated to be 8.6 nm. This result is consistent with that of TEM observations. For EPD experiments, a BTO suspension with a BTO particles concentration of 0.075 moVl was used, and the BTO particles in the suspension were found to have an average particle size of -10 nm and a very narrow particle size distribution, measured with a laser particle size analyzer. This suggests that the Fig. 4. TEM image of BTO nanoparticles and primary BTO nanoparticles formed no significant an ED pattern ofthe particles (in the inset). agglomeration in the suspension. The small size of BTO nanocrystallites in the well dispersed suspension makes it possible to deposit BTO into micropatterned resist mold formed on Pt/Ti/Si substrates. SEM micrographs of BTO films deposited by EPD at different conditions are shown in Fig. 5. The BTO films deposited from the nanoparticles suspension had much flatter surfaces than those from the sub-micron particle suspension.".'2 A uniform microstructure and a very smooth surface were observed for the BTO films. The thickness of the deposited film was able to be easily controlled by adjusting the EPD conditions. BTO thin films with a desired thickness from -100 nm to about 1 pm have been fabricated in one step by adjusting the EPD voltage and duration time. Dielectric property measurements of BTO ceramic thin films prepared at temperatures between 600 and 800°C demonstrated that the films had a dielectric constant of about 100 and a dielectric loss <1%. The films were also found to have a very high dielectric strength >500 kV/cm. The cause that the BTO ceramic thin films exhibited such a relatively low dielectric constant is being investigated. Figure 6 a) and b) show SEM micrographs of a resist mold with a hexagonal array of air holes in it, which was formed on a Pt/Ti/Si substrate. The micropatterned resist mold was prepared by electron beam lithography, which is one of the most advantageous techniques for preparing high quality micropatterns with high resolution. The thickness of the resist film (the depth of air holes) was measured to be about 800 nm, and the edge of the air holes was confirmed to be very sharp. Figure 6 c) and d) show different views of a hexagonal array of BTO nanoparticles pillars as prepared by EPD. The BTO nanocrystallites are seen to have been compactly deposited and the

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Fig. 5 . FE-SEM micrographs of BTO films prepared under different EPD conditions: a) 5v/lmin; b) 15V/20min.

Fig. 6. FE-SEM micrographs of a), b) resist mold; c), d) as-prepared BTO micropatterns.

BTO pillars almost had the same diameter as that of air hole in the resist mold. The height of the pillars can be controlled in the polishing process. Other micropatterns of microlines or microgrids also have been successfully prepared. Figure 7

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Fig. 7. FE-SEM micrographs of a), c) resist mold; b), d) BTO micropatterns. a) shows an SEM micrograph of a resist mold with microlines on a Pt/Ti/Si substrate. The resist

microlines had a width of 400 nm with a spacing of 400 nm in between, and in Fig. 7 b) is shown an SEM micrograph of BTO microlines produced from the mold. In Fig. 7 c) and d) are shown an SEM micrograph of a resist mold with an array of square pillars with a width of 2.4 pm and a spacing of 600 nm in between and an SEM micrograph of a BTO microgrid produced from the resist mold, respectively. Figure 8 shows a 3D image of an as-deposited BaTi03 nanoparticles microgrid, evaluated with SPM over an area of 10 pmx 10 pm, indicating that the BTO microgrid had a height of about 200 nm and a robust structure with a sharp edge. Sintering of BTO

Membranes, Films, and Coatings

Fig. 8. SPM 3D image of the as-deposited BaTi03 micropatterns.

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Fig. 9. FE-SEM micrographs of sintered BTO micropatterns. nanoparticles pillars as deposited by EPD, whose SEM micrographs are shown in Fig. 6 c) and d), at 600°C for 30 min yielded robust BTO ceramic pillars without any appreciable disordering on the substrate, as shown in Fig. 9. To examine the ferroelectric properties of the BTO ceramic pillars,

c,

__

C

1

-30

-20

-10

10

20

.

1

30

measurements of strain-DC D C b b s to the t$ of SPM '$1 characteristics were carried out Fig. 10. Strain hysteresis loop of a sintered BaTi03 pillar. for some of the ceramic pillars. Figure 10 shows a quadratic strain hysteresis loop observed for a single BTO ceramic pillar, demonstrating that the BTO pillars sintered at a low temperature of 600°C exhibited induced ferroelectric characteristics under high electric field. CONCLUSION Well crystallized BTO nanoparticles in a pseudo-cubic perovskite phase with an average particle size of -10 nm and a narrow size distribution were synthesized by a high concentration sol-gel process. By using a mixed solvent of EGMME and Acac as a dispersion medium, transparent and stable BTO suspensions were successfully obtained. Crack-free BTO thin films of controlled thickness from -100 nm to about 1 pm were electrophoretically deposited on Pt/Ti/Si substrates. A new method of sol-gel process and nanoparticles EPD combined with an

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EB lithography technique were investigated to prepare BTO ceramic micropatterns at a relatively low temperature. Confirmation of good feature qualities and good adhesion between the micropattern and the substrate suggested that this method is promising as one for micropatterning electroceramic thin films on conductive substrates. Nanoparticles EPD-derived BTO ceramic pillars, sintered at a low temperature of 6OO0C, exhibited induced ferroelectric characteristics under high electric field. FOOTNOTES * Corresponding author. Email: [email protected] p . ** This work is partially supported by a Grant-in-Aid (No. 16206065) from the Japan Society for the Promotion of Science (JSPS). Y. J. Wu is grateful to the JSPS for a Fellowship. REFERENCES ‘Y. Yamashita, H. Yamamot, Y. Sakabe, “Dielectric Properties of BaTi03 Thin Films Derived from Clear Emulsion of Well-Dispersed Nanosized BaTi03 Particles,” Jpn. J. Appl. Phys., 43, 6521-24 (2004). 2K. Tanaka, K. Susuki, D. Fu, K. Nishizawa, T. Miki, and K. Kato, “Grain Size Effect on Dielectric and Piezoelectric Properties of Alkoxy-Derived BaTi03-Based Thin Films,” Jpn. J. Appl. Phy~.,43,6525-29 (2004). 3M. E. Marssi, F. L. Marrec, I. A. Lukyanchuk, M. G. Karkut, “Ferroelectric transition in an epitaxial barium titanate thin film: Raman spectroscopy and x-ray diffraction study,” J. Appl. Phys., 94,3307-12 (2003) 4H.Basantaka Sharma, H.N.K. Sarma, and A. Maningh, “Ferroelectric and Dielectric Properties of Sol-Gel Processed Barium Titanate Ceramics and Thin Films,” J. Muter: Sci., 34, 1385-90 (1999). 5 C.L. Jia, M. Siegert, and K. Urban, “The Structure of the Interface Between BaTiO3 Thin Films and MgO Substrates,” Actu. Muter:, 49,2783-89 (2001). 6Y. Sakabe, Y. Takeshima, and K. Tanaka, “Multilayer Ceramic Capacitors with Thin (Ba,Sr)Ti03 Layers by MOCVD,” J. Electrocerumics, 3, 115-21 (1999). I H. Kozuka, and A. Higuchi, “Single-layer Submicro-thick BaTi03 Coating from Poly(vinylpyrro1idone)-coating Sols: Gel-to-ceramic Film Conversion, Densification, and Dielectric Properties,” J. Muter: Res., 16,3 1 16-23 (200 1). ‘P. Sarkar, and P. S. Nicholson, “Electrophoretic Deposition (EPD): Mechanisms, Kinetics, and Application to Ceramics,” J. Am. Cerurn. Soc., 79, 1987-2002 (1996). 91. Zhitomirsky, and A. Petric, “Electrolytic and Electrophoretic Deposition of Ce02 Films,” Muter: Lett., 40,263-68 (1999)

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‘OS. Okamura, T. Tsukamoto, and N. Koura, “Fabrication of Ferroelectric BaTi03 Films by Electrocphoretic Deposition,” Jpn. J. Appl. Phys., 32,4 182-85 (1993). “R. Louh, and Y. Hsu, “Fabrication of Barium Titanate ferroelectric Layers by Electrophoretic Deposition Technique,” Mater. Chem. and Phys., 79,226-69 (2003). 12M. Nagai, K. Yamashita, T. Umegaki, and Y. Takuma, “Electrophoretic Deposition of Ferroelectric Barium Titanate Thick Films and Their Dielectric Properties,” J. Am. Ceram, SOC.,

76,253-55 (1993). I3J. Li, Y.J. Wu, H. Tanaka, T. Yamamoto and M. Kuwabara, “Preparation of Mono-dispersed Suspension of Barium Titanate Nanoparticles and Electrophoretic Deposition of Thin films”, J. Am. Ceram. SOC.,87, 1578-81 (2004). 14M. Alexe, C. Harnagea, and D. Hesse, “Non-conventional Micro- and Nanopatterning Techniques for Electroceramics,” J. Electrceramics, 12,69-88 (2004). I5J. Baborowski, “Microfabrication of Piezoelectric MEMS,” J. Electrceramics, 12, 33-5 1 (2004). I6H. Shimooka, and M. Kuwabara, “Preparation of Dense BaTi03 Ceramics from Sol-Gel-Derived Monolithic Gels,” J. Am. Ceram. SOC.,78,2849-52 (1995). 17H. Shimooka, and M. Kuwabara, “Crystallinity and Stoichiometry of Nano-structured Sol-Gel-Derived BaTi03 Monolithic Gels,” J. Am. Ceram. SOC.,79,2983-85 (1996).

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Preparation and Properties of Hydrothermally Stable y-Alumina-Based Composite Mesoporous Membranes Md. Hasan Zahir,+ Koji Sato, Hiroshi Mori, and Yuji Iwamoto* Japan Fine Ceramics Center, Hybrid Process Group, Nagoya 456-8587, Japan

Mikihiro Nomura and Shin-ichi Nakao Department of Chemical System Engineering, Faculty of Engineering, University of Tokyo, Tokyo 1 13-8656, Japan

We report here on our thermal and hydrothermal investigations comparing mesoporous 7-&03 membrane with single and double dopant membranes prepared by the sol-gel method. Improvements in the h drothermal stability of mesoporous y-A1203 by transition (Gal+) or rare-earth (La3+) cations are discussed, along with the effectiveness of double dopant (Ga3+La3’). The amounts of Ga203 oxide used varied between 0.0 and 30 mol% and those of La203 between 6 and 15 mol%. The thermal and hydrothermal (up to 75% steam) stability of nine types of membranes fabricated on an asymmetric porous aA1203 support by means of a dip-coating process was characterized by H2 gas permeation at 500°C. By conducting tests with wide variations in dopant concentrations, material characterizations, and gas permeance performance, we have been able to optimize the key parameters for hydrothermally stable systems. 1. Introduction EMBRANE reactors offer a new industrial paradigm, as they combine reaction and separation in one step.’ The application of high-temperature membrane reactors in the steam-reforming step, in particular, has the potential to achieve conversion efficiencies equivalent to those attained in conventional reactors at a significantly lower temperature of about 500°C.’,2A state-of-the-art microporous silica membrane is one candidate for developing novel membrane reactor^.^ Generally, a microporous silica-based membrane can be fabricated on a porous support, and an important structural feature is the intermediate layer, i.e., a mesoporous (2 nm< 0<SO nm) y-A1203membrane that is placed between a layer comprising a hydrogen perm-selective microporous silica-based membrane on the upper side and the surface of an u-AI2O3porous support on the inner The mesoporous intermediate y-A1203layer confers mechanical strength to the selective silica top layer. In the above-mentioned processes, high operating temperatures are necessary, but the reaction atmospheres usually contain considerable amounts of steam either because water is one of the reactants or because water is added to reduce coke formation. Thus, in many applications, the membranes must be sufficiently stable in environments with both increased temperature and high steam content. It is well known that transition aluminas, such as y-A1203, are metastable phases that gradually transform into u-A1203as T. Gur--contributing editor

Manuscript No. 21444. Received February I , 2006; approved May 2, 2006. This work has been supported by the New Energy and Industrial Technology Development Organization (NEDO) as a part of an R&D Project on “Highly Efficient Ceramic Membranes for High-Temperature Separation of Hydrogen” promoted by the Ministry of Economy, Trade, and Industry (METI), Japan. ‘Member, American Ceramic Society. ‘Author to whom correspondence should he addressed. e-mail: [email protected]

Membranes, Films, and Coatings

calcination temperatures increase. Generally speaking, the addition of suitable dopants to y-A1203confers an increased thermals as well as hydrothermal stability to the transitional aluminas.6 These dopants can significantly reduce membrane deterioration during high-temperature excursions. Recently, we prepared a new type of supported membrane composed of La203-doped Ga2O3-AI2O3, which can be used at above 500°C without hydrothermal deterioration.’ It has been reported that the GazO3-Al2O3catalyst with a solid solution species, prepared by the sol-gel method, showed quite high and stable catalytic activity at 500°C for NO reduction by hydrocarbon in the presence of steam.6 The high performance of the Ga2O3A1203 catalyst under a steam atmosphere probably originated from the presence of acid-base sites, a coke deposition removal ability, and the formation of a unique pore structure.8 This high tolerance for steam is quite interesting from a practical point of view. Several recent studies have addressed the stability of alumina membrane pore structures with respect to thermal and hydrothermal treatments, and its improvement with the addition of lanthan~m.’?~ In fact, lanthanum is present in all the washcoat compositions proposed for automotive depollution catalysts. Church et a/.’’ studied the effectiveness of rare earths (La, Ce, Pr, Nd, Sm, Yb) and alkaline earths (Ca, Sr, Ba) for stabilization of alumina carriers. The most effective dopants are larger ions with high ionic charges, i.e., La3+ and Ba2+, for rare and alkaline earths, respectively. These studies were carried out mainly over unsupported membranes, and only the thermal stability was studied based on the heat-treatment results. It is known, moreover, that a membrane’s performance depends on its permeability. In this connection, nothing is known about the gas permeance performance of La203-doped A1203 membranes under steam conditions at 500°C. Gallaher and Liu” reported that the pore growth of a commercial y-A1203 membrane was accelerated by the presence of steam after heat treatment at high temperatures. Chang et a/.” performed a comprehensive comparative study on the thermal and hydrothermal stability of alumina, titania, and zirconia membranes prepared by sol-gel techniques. It was found that the pore growth rates of alumina, titania, and zirconia membranes were enhanced in the presence of steam at elevated temperatures. Recently, Nijmeijer e f al.I3reported that y-A1203membranes could be separated from the u-A1203 supports without using an A1(H2P0& precursor solution during high-temperature steam treatment. Zahir et a ~Lin , and ~ Burggraaf: and Gallaher and Liu” did not complain about this type of separation problem, however, during hydrothermal treatment of supported y-A1203 membranes. More recently, Yeung et al.’ studied only the thermal stability improvements afforded by alkaline earth (Ba), rare earth (La and Ce), and I11 B (Y) dopants along with the effectiveness of multiple dopants. To the best of our knowledge, no report has appeared to date that can provide the gas permeation performance of y-A1203-basedcomposite membranes at 500°C under steam conditions. In this study, we have systematically 215

investigated the influence of Ga3+ alone as well as of Ga3+La3+ mixed species content over mesoporous y-A1203by H2 gas permeance as part of an ongoing study. In order to examine the stability of supported doped and nondoped y-A1203mesoporous membranes in steam at higher temperatures, several kinds of y-A1203-basedcomposite membranes were fabricated in this study by the sol-gel technique. The hydrothermal stability characteristics of the composite alumina membranes were examined by evaluating their gas permeation properties after exposure to steam (75%) at 500°C. The porestructure evolution during hydrothermal treatment was also studied directly on the membranes by nano-permporometer. 11. Materials and Experimental Procedures

(1) Synthesis Sols were prepared by reacting 0.5 mole of aluminium-tri-secbutoxide (Al(0-s-Bu)~= ATSB, 97% Aldrich, USA) with double-distilled water at 90°C. After the addition of ATSB, the mixture was maintained at 90°C for at least 1 h to evaporate off the butanol that formed. The mixture was subsequently cooled down to -60°C and peptized with 1M H N 0 3 at a pH of about 3. During the synthesis, the sol was stirred vigorously. The peptized mixture was refluxed at 90°C for 12 h, yielding a very stable 0.5M boehmite sol with a clear whitish-blue appearance. Doping of this sol was performed by mixing it with an aqueous solution of Ga(N03)2.8Hz0 (Kishida Chemical Co., Japan) and/or La(NO&. 6H20 (Wako Pure Chemical Co., Japan). The compositions [metal nitrate (mol%): boehmite (mol%)] of various precursor solutions for sol synthesis are shown in detail in Table I. The amount of Ga(N03)2.8 H 2 0 (for example) corresponding to the added 5-30 mol% Ga was calculated as (mol Ga/mol G a

+ mol ATSB) x 100%

(1)

The increase in sol viscosity upon addition of the nitrate solution was probably a result of a decrease in the pH due to the acidity of the nitrate solution. Doped boehmite sols were therefore prepared without the addition of HN03.

( 2 ) Membrane Preparation A dipcoating solution was obtained by diluting boehmite sol (y-AI00H) with a 3.5 wt% solution of polyvinyl alcohol (PVA, Kanto Chemical Co. Inc., Tokyo, Japan, MW = 72000) prepared by refluxing at 90°C for 3 h. It has been reported that the pore-size distributions of y-A1203membranes with the addition of PVA do not show a measurably altered pore structure.14 Unless stated otherwise, all y-AI2O3-basedcomposite membranes in this paper were prepared by adding PVA to the boehmite colloidal suspension. Before the dip-coating step, the graded alumina support was cleaned by rinsing in acetone. Finally, the dry homogeneous support was brought into contact with the boehmite sol. The Table I. Tested Samples Prepared by the Sol-Gel Method and their Composition in mol% Samples

Composition (mol%)

Abbreviated as

(100) (5:95) (l0:90) (20:SO) (30:70) (694) (6:30:64) (10:30:60) (15:30:55)

y-AI203 GA-5 GA-10 GA-20 GA-30 LA-6 LGA-6 LGA- 10 LGA- 15

All samples were calcined at 600°C for 3 h in air.

216

concentration of boehmite increased at the entrance of the pores, and gelation of the sol occurred. Owing to the capillary forces, water might be sucked into the pores of the support. After a certain period (on the order of seconds), the support was removed and the gel layer was dried and calcined. The dip-coated boehmite was converted into y-A1203by calcination. After dip coating, all the samples were calcined from ambient temperature to 600°C at a heating rate of I"C/min and then maintained at 600°C for 3 h, followed by gradual cooling (2OCjmin) to room temperature. In all cases, a layer of doped or non-doped alumina was deposited onto the graded support at a fixed slipcasting time of 10 s. The whole process of dipping, drying, and calcining was repeated once to repair any defects in the first yAl2O3layer. The porous support used in the present study was an a-alumina substrate (@ 6 mm, L 800 mm, mean pore diameter 1.1 pm, porosity 0.43%, and average pore diameter 80 nm: Noritake Co., Aichi, Japan).

( 3 ) Characterizations X-ray diffraction (XRD) patterns were identified using an X-ray diffractometer (Rigaku, RINT-2000, Japan) with CuKa radiation of 50 keV and 200 mA with a monochromator over a diffraction angle range of 20 from 10" to 80" at a scan rate of 2"/ min by analyzing the respective powdered samples prepared by calcining the original boehmite sols. The lattice parameter value was obtained with silicon used as an internal standard. The morphology and thickness of the membranes were examined by scanning electron microscopy (SEM; HITACHI S-4500, Japan) with an energy-dispersive X-ray (EDX) analyzer (HORIBA, EMAX-7000, Japan). For asymmetric membranes with very thin effective layers (< 5 pm), conventional N2 adsorption and/ or Hg porosimetry methods fail to provide accurate and reliable results. l5 The pore-size distribution measurements were therefore conducted directly on the membrane by a nano-permporometer (Seika Co., Tokyo, Japan). In order to understand the thermal and hydrothermal stabilities of doped and non-doped mesoporous y-A1203membranes, the H2 gas permeability was measured at 500°C in the apparatus shown in Fig. 1. The hydrothermal treatment was performed under a N2 to steam ratio of 1:3 at 500°C. The steam was condensed and collected from both the reject and permeate streams. Before the permeances were measured, the water flow was halted at the specified time intervals. The gas permeability performance was evaluated by the volumetric method at a constant pressure." More than five measurements were made and the average value was recorded. 111. Results and Discussion

( I ) Phase Structure The XRD pattern (Fig. 2(a)) of y-A1203 after dehydration at 600°C was consistent with that reported by other researchers,I6 and was attributed to the formation of less crystalline phases

Pressure Gauge

Mass Flow,

P.

0 I

&

/Mesoporous Ceramic Membrane Furnace

Flow Meter

Gas Cylinder

\

Stop Valve

Fig. 1. Schematic representation of a thermal and hydrothermal stability test unit with on-line gas permeation characterization.

Progress in Nanotechnology: Processing

I

I

I

I

I

I

0 = Present studv c

c

m

i

= JCPDS No. i9-0063

E .

0.800

c

{ 0.795

.-8

+: 4 0.790 0.785

0

5

10 15 20 25 GapOBcontent/rnol%

30

Fig.3. Change in the peak position assigned to y-A1203(440) and the lattice constant of Ga2O3-Al2O3solid solution as a function of mol% of Ga3+ ions.

I

10

20

30

I

40

50

2 thetaldegree

60

70

ao

Fig. 2. X-ray diffraction patterns of powder (a) y-AlzO3, (b) GA-5, (c) GA-10, (d) GA-20, and (e) GA-30.

due to particle-size broadening. The X-ray data (Figs. 2(bHe)) indicated slightly sharper XRD patterns with increasing Ga2O3 concentrations. Lines due to Ga2O3 were not detected, even on 30 mol% Ga203-doped A1203 (Fig. 2(e)), suggesting the presence of highly dispersed Ga2O3. Figure 2 also shows shifts of the XRD peaks at 28 = 45.5" and 66.5", which can be assigned to yA1203(400) and (440), toward lower angles, suggesting the formation of a uniform solid solution.677In order to examine these shifts in detail, the crystal structure of the Ga2O3-AI2O3,changes in the XRD peak assigned to y-A1203 (440), and its lattice constant with Ga2O3 content (mol%) were measured. The results are shown in Fig. 3. The lattice constant of y-A1203(440) alone was essentially identical to that of the JCPDS card No. 290063. A linear correlation was observed in the region of Cia203 content up to ca. 30 mol%. Haneda et aL6 also observed a linear line when it was plotted against the amount of added Ga2O3, which proved that the materials were only a solid solution. It should be noted that the diffraction peaks for GA-30 consist of single peaks without shoulder peaks or split peaks (Fig. 2(e)). Taking these results into account, we considered that composite oxides such as a solid solution of [Ga,A1~1-,,]203(x< I)] could be formed uniformly in a Ga2O3-AI2O3system. In this study, the Ga2O3 doping content was checked up to 30 mol% for two specific reasons. First, 30 mol% of Ga203was known to give the highest catalytic activity in the presence of steam,' and second, a phase transformation occurred when the GazO3 content increased to above 30 mol% and after calcination at above 800"C, as shown in the phase diagram of the Ga2O3-AI2O3system.17 The y-A1203itself represents a somewhat amorphous structure as revealed by XRD, and the addition of 6 mol% La203 produces a more or less amorphous structure, as shown in Fig. 4(a). No diffraction attributable to a La203-containing phase was detected. La203 is probably in an amorphous state, or else the particles are very well dispersed in the y-A1203phase, with the crystalline phases scarcely detected during scanning. As for the 2.65 wt% La2O3-AI2O3 sample, after calcination at 500°C under air, Beguin et a1." observed coverage of the alumina sheets by small particles exhibiting a high contrast. These particles probably consist of La203 that is not detected by XRD because of their small size. Membranes, Films, and Coatings

Figures 4(cHe) show the XRD patterns of powders with a composition of La:Ga:Al = X30:70-X (X= 5, 10, and 15 mol%) calcined at 600°C for 3 h. The XRD pattern of GA-30 was inserted again in Fig. 4(b) for purposes of comparison. A drastic change was observed after the addition of 6 mol% LazO3; the sharp peak of GA-30 was diminished and a quite amorphous powder was formed, which might have led to the production of ultrafine composite powders. Only the y-A1203 phase was detectable in the La203-doped Ga2O3-AI2O3 (referred to as LGA-6) system. The La203 signals were detected easily by EDX in the SEM. This implies that the La203 oxide on the surface of y-A1203is highly dispersed or low in crystallinity. When La3+ cations were added to the GA-30 with a composition of La:Ga:Al= 6:30:64, the La3+ cations may not have occupied either the tetrahedral or octahedral sites due to their larger ionic radius (0.104 nm) in comparison with that of Ga3+ ion (0.062 nm) and A13+ ion (0.051 nm). As the XRD data showed that no X-ray detectible phase was present other than that of y-A1203upon the addition of 6 mol% of La203 to GA30, the La203 oxide could be present either on the grain surface or as a thin layer over the GA-30 system. At a La203content of 10 mol%, the XRD pattern still represents the y-A1203phase structure. When the La203 content (X)increased up to 15

Fig.4. X-ray diffraction patterns of powder (a) LA-6, (b) GA-30, (c) LGA-6, (d) LGA-10, and (e) LGA-15.

217

10.0

I

2 9.0 -6 I

I

0

8.5

. 0 vX

m

C

8.0

a,

E

g

7.5 7 . 0 ' " 1 1 " 1 1 ' " 1 ' ' ' 1 1 1 1 ' ' 0 4 8 12 16

20 0 Timelh

4

8

12

16

20

Fig.5. (a) Effect of steam on the time course of H2 permeance for an a-AI203 porous support alone (O), mesoporous y-Alz03(V),GA-10 (A), GA-20 (+),and GA-30 (W) membranes; (b) porous support (o), LA-6 (W), LGA-10 (A), and LGA-15 (+), and LGA-6 (solid line 0) membranes at 500°C for 20 h. (The 0 h represents the Hz permeance at 500°C without the presence of steam.)

mol%, however, no y-A1203phase was detected. After calcination at 600"C, only a very low degree of crystalline Laz03could be detected at -15% or lower, but as the addition of La203 increased to 20% or higher, the crystalline La203 structure became readily apparent. This could imply a high degree of La203 dispersion on the GA-30 surface at submonolayer coverage, with a bulk-like, crystalline Laz03 structure failing to initiate any substantial formation until additions exceeding a monolayer are present.

( 2 ) Effect of Thermal Treatment y-A1203alone, as well as single oxide-doped AI2O3,i.e., GA-10, GA-20, and GA-30 samples, showed no change in H2 gas permeance during 20 h thermal treatment, suggesting that the permeance remained the same from 0 to 20 h. In Fig. 5(a), the 0 h represents H2 gas permeance at 500°C without the presence of steam. The thermal stability results were not plotted in detail in Fig. 5(a), because the permeance data might produce merely a straight line. It is important to note that the membrane thicknesses were examined by SEM before thermal and hydrothermal testing. The typical membrane thicknesses were observed to be in the range of S3.5 pm from a cross-sectional view of the

membrane^.^

(3) Effect of Hydrothermal Treatment The effect of hydrothermal treatment with a 75% steam stream on the membranes is shown in Figs. 5(a) and (b). After the first measurement at 0 h, the remainder of the permeance data was obtained under hydrothermal conditions. The H2 gas permeance of y-Alz03 drastically increased, i.e., from 7.44 to 8.79 x mol . (m2.s . Pa)-' within the first 2 h in steam. The permeance of the y-A1203membrane increased steeply with an increase in exposure time in the presence of steam. Gallaher and Liu" investigated the hydrothermal (up to 90% steam) stability of commercial y-A1203membranes by Hz gas permeation. They found that the H2 gas permeances increased rapidly, as with hydrothermal treatment. The results indicate that the y-A1203membrane is not a good candidate for use as an intermediate layer for fabrication of microporous gas separation membranes. The intermediate layer (y-A1203membrane) must be stable under hydrothermal conditions to minimize the defects of the microporous top layer. Like y-A1203, the Gaz03-doped membrane also showed a rapid initial increase in permeance, and the sub218

sequent permeance maintained a constant value. The H2 gas permeance increment sequence with respect to Ga2O3 doping was GA-30 > GA-20 > GA- 10, as shown in Fig. 5(a). In the case of the GA-30 membrane, however, the gas permeance showed considerable stability in comparison with that of y-Al203 under hydrothermal conditions. For a single dopant, 30 mol% containing Gaz03 in y-A1203 was found to be the most effective under hydrothermal conditions at 500"C, as shown in Fig. S(a). The effects of La203 doping on the hydrothermal stability of yA1203 as well as GA-30 membranes were then studied, and the results are summarized in Fig. 5(b). The H2 permeance of LA-6 membranes increased gradually, approaching that of the m-A1203 porous support within 20 h in steam. It has generally been assumed that the presence of Laz03 on the y-A1203surface might immobilize the surface ions and thereby reduce sinterin of the material and impede the phase transition to a-A1203?g9 It is interesting to note, however, that La203-doped A1203 membranes have higher heat resistance than A1203membranes in the absence of steam.' The situation is totally different under hydrothermal conditions. In the case of LGA-10, the permeance increased to a slightly higher degree than that of the LGA-6 sample. In the case of LGA-15, a very large difference was observed between the LGA-6 sample and LGA-15. The results revealed that increasing the amount of Laz03in the sol from 10% to 15% in the GA-30 system was not advantageous for fabrication of a hydrothermally stable intermediate layer. During hydrothermal treatment, two distinguishable phenomena were observed: a drastic increase in permeance for 2 h and very slow progress of subsequent increases upon introduction of a stream containing 75% steam. After the first 2 h, the permeance of LGA-6 membrane had increased only ca. 2% in the presence of steam. This can be considered a remarkable result, as both the thermal and hydrothermal stability of the yA1203membrane was insufficient at I500°C."

( 4 ) Meso-Pore Structure The N2 permeation and the calculated pore size distribution (PSD), obtained by permporometry, for a supported as-prepared mesoporous y-A1203membrane are shown in Fig. 6(a) ( 0 ) . The calculated PSDs show the distribution of the total number of pores within a certain range. In the case of y-A1203,it can be seen that the permeation at large Kelvin radii was almost Progress in Nanotechnology: Processing

I

.-nc L-

2 0.8 0

al 0 C

m

kl n

0.6 0.4 0.2

0.0 0

2

4

6

2 10 0 Kelvin diameter [nm]

8

4

6

8

10

Fig.6. Pore size distribution of (a) mesoporous y-Alz03,(b) GA-30, (c) LA-6, and (d) LGA-6 membrane before ( 0 ) and after steam treatment (0) for 20 h measured by a nano-permprometer.

zero, and that a strongly increased permeation was observed at around 4-4.5 nm. This implies that no pores or only a few pores were present with Kelvin radii larger than 4.5 nm. The calculated maxima was found at rk = 4.5 nm. The pure mesoporous y-A1203membrane showed an average pore size somewhat larger than those of GA-30, LA-6, and LGA-6 membranes. The pore size distributions of all doped as-prepared samples had a rather narrow distribution. Figures 6(bHd) ( 0 ) clearly show a sharp pore size distribution with an average Kelvin radius of 3.0, 3.0, and 3.2 nm for as-prepared GA-30, LA-6, and LGA-6 membranes, respectively. From the above permeability results, it was clear that permeability increases due to increases in pore size as steam exposure time increases. In the case of mesoporous y-A1203, the average pore size increased and the pore size distribution (Fig. 6(a), 0) broadened with prolonged heat exposure in the presence of steam. These results agree very well with those of Gallaher and Liu." A general trend was observed, i.e., the PSD shifted toward larger pore radii with prolonged steam exposure at 500°C for the y-A1203,GA-30, and LA-6 samples. Figure 6(b) (0)shows that the GA-30 membrane retained a somewhat smaller pore size (- 3.5 nm), however, after hydrothermal treatment. To the best of our knowledge, no permporometry measurement has yet been performed on doped y-A1203-based composite-supported membranes after hydrothermal treatment. Laz03-doped y-A1203,especially, might be an interesting system, as several studies have already reported on its high thermal stability using powder calcinations. In this study, we observed that the pore size of an LA-6 membrane shifted from 2.1 to 5.5 nm after hydrothermal treatment, as shown in Fig. 6(c) (0). It seems that the LA-6 sample was not effective for pore stabilization under hydrothermal conditions at 500°C. In the case of Membranes, Films, and Coatings

the LGA-6 sample, on the other hand, the pore size was unchanged after hydrothermal treatment. A sharp pore size distribution with an average Kelvin radius of 2.4 nm was found for the LGA membrane, as shown in Fig. 6(d) (0).

Fig.7. Scanning electron microscope images of top surface view of (a) y-AIzO, and (b) LGA-6 after hydrothermal treatment. 219

-

10.0

-n

9.5

- I

I

m

'?

N

-'E

"'O

9.0 8.5

. 0 T-

o) 0 C

8.0

E

7.5

m

n

G

7.0

0

10

20 30 Timelh

40

50

0

2

4

6

8

0

Kelvin diameterlnm

(a) Effect of steam on the time course of Hz permeance for an LGA-6 membrane on the a-AIz03porous support at 500°C for 50 h. (b) Pore size distribution of (a) mesoporous LGA-6 membrane before ( 0 ) and after hydrothermal treatment (0)for 50 h measured by a nano-permprometer. Fig. 8.

(5) Microstructure Very smooth morphology was observed only in the case of asprepared y-A1203,GA-30, and LGA-6 samples. It was apparent that a large number of Al2O3particles agglomerated on the surface after modification with La203. Because of these changes in the nanostructure of the membranes, doping with La203 alone was not an appropriate strategy for improving the hydrothermal stability of y-A1203. The sol showed a quite viscous appearance with the addition of Ga2O3 content up to 30 mol% to AI2O3. Even after the addition of 6 mol% La203to GA-30, the sol remained viscous, and thus appropriate for membrane fabrication on the a-AI2O3porous support. The surface of the steam-exposed LGA-6 membrane was smooth and crack-free, whereas morphology with spherical grains was found on the y-AI2O3 membrane, as shown in Figs. 7(a) and (b). The observed gas permeation was consistent with morphological changes in the case of the y-A1203sample after hydrothermal treatment. The rough morphology observed after hydrothermal treatment, in the case of LGA-10 and LGA-15 membranes, might be due to high viscosity and significant shrinkage during hydrothermal treatment for 20 h. ( 6 ) Durability of LGA-6 Membranes The H2 gas permeance results suggest that the most critical period was the first few hours of hydrothermal treatment. Our fabricated mesoporous LGA-6 membrane showed only a small change within the first 2 h. The membrane appeared to be relatively stable thereafter. The higher concentration of steam appears to have an increased mineralizing effect. It was not clear, however, whether a truly stable plateau had been reached under these conditions or whether the kinetics of the transformation had only slowed to the point where much longer exposure times might be necessary to observe significant changes in the membrane structure. To determine the durability of the LGA-6 membrane, the gas permeation performance was examined for more than 20 h in the presence of 75% steam. The LGA-6 membrane proved able to withstand a N2 to steam ratio of 1.3 at 500°C for more than 50 h with no sign of degradation, as shown in Fig. 8(a). The LGA-6 membrane was found to have high durability at 500°C in the presence of 75% steam. Figure 8(b) shows the pore size distribution of an LGA-6 sample after hydrothermal testing for more than 50 h. The pore radius did not shift after hydrothermal treatment. As expected, the pore size distribution shows good consistency with the gas permeance of the LGA-6 sample. As the membrane prepared in this study has a very narrow pore size distribution, with the average pore radius in the range of 3.W.5 nm, the contributions of ordinary diffusion and viscous Poiseuille flow to the gas fluxes were negligible 220

compared with those of Knudsen flow. Gas permeation through the y-AlzO3 membrane deviated from the Knudsen diffusion mechanism after hydrothermal treatment. The observed gas permeation was consistent with Knudsen diffusion in the case of the LGA-6 sample after hydrothermal treatment, as was expected based on its stable retention of a crack-free microstructure. We have demonstrated that the addition of combined 6 mol% La203 as well as 30 mol% Ga203 (referred as LGA-6) can improve the hydrothermal stability of y-AI2O3membranes. As stated above, the incorporation of well-dispersed La203into a solid solution of GA-30 might enhance the stabilization of pore growth under hydrothermal conditions. With respect to yA1203, certain stabilization conditions of the tests need to be recorded accurately: (i) the exact temperature of the test (thermodynamics and kinetics), (ii) the duration of the test (kinetics), and (iii) the type of atmosphere (dry or wet). For temperatures below I O O O T (5OO0C,for example), the conditions were not too severe, and some stabilization could be observed and reported. As it was obvious that an additive must remain on the surface in order to decrease the rate of surface diffusion, large ions were required to prevent dissolution into the bulk. Based on these considerations, lanthanides were a logical choice. Indeed, it was found that the addition of 6 mol% La203increases the thermostability of y-A1203.9.'9 GA-30 is widely recognized as a steam resistance catalyst,6 and we also observed considerable hydrothermal stability in the GA-30 membrane. The high steam tolerance was therefore considered to be due to the incorporation of well-dispersed La203into the solid solution of GA-30, which was thought to enhance the stabilization of pore growth under hydrothermal conditions.

IV.

Conclusions

The thermal and hydrothermal stability of mesoporous membranes was studied in terms of gas permeability, pore structure, and surface microstructure. The permeability and characterization results showed that the hydrothermal stability of the La203doped Ga2O3-AI2O3(referred to here as LGA-6) mesoporous membrane was greatly improved in comparison with mesoporous y-A1203, while the pore system remained highly ordered. The fabricated mesoporous membrane exhibits an extraordinarily high hydrothermal stability at 500°C for a very long service period. The ability of a fabricated mesoporous membrane to withstand steam is advantageous for many reaction systems, and offers possibilities for better exploiting the benefits of a membrane reactor configuration. The hydrothermal test indicates that the membrane prepared is viable for use as an intermediate layer for fabricating a microporous gas separation membrane. Progress in Nanotechnology: Processing

References ‘E. Kikuchi, “Membrane Reactor Application to Hydrogen Production,” Cafal. Today, 56.97-101 (2000). ’K. Jarosch and H. I. de Lasa, “Novel Riser Simulator for Methane Reforming Using High Temperature Membranes,” Chem. Eng. Sci., 54, 1451-60 (1999). ’8. N. Nair, T. Okubo, and Nakao, “Structure and Separation Properties of Silica Membrane,” Membrane, 25 (21 73-85 (2000). 4R. M. De Vos and H. Verweij, “High-Selectivity, High-Flux Silica Membranes for Gas Separation,” Science, 279, 1710-1 (1998). ’K. L. Yeung, J. M. Sebastian, and A. Varma, “Mesoporous Alumina Membranes: Synthesis, Characterization, Thermal Stability and Nonuniform Distribution of Catalyst,” J. Memhr. Sci., 131. 9-28 (1997). 6M. Haneda, Y. Kintaichi, T. Mizushima, N. Kakuta, and H. Hamada, “Structure of Ga2O3-AI2O3Prepared by SolGel Method and its Catalytic Performance for NO Reduction by Propene in the Presence of Oxygen,” Appl. Caral. B, 31, 8192 (2001). 7Md. H. Zahir, K. Sato, and Y. Iwamoto, “Development of Hydrothermally Stable SolGel Derived La20,-Doped Ga2O3-AI2O3 Composite Mesoporous Membrane,” J . Membr. Sci., 247, 95-101 (2005). *K. Shimizu, A. Satsuma, and T. Hattori, “Selective Catalytic Reduction of NO by Hydrocarbons on Ga2O3/AI203Catalysts,” Appl. Caral. B, 16, 319-26 (1998). 9Y. S. Lin and A. J. Burggraaf, “Preparation and Characterization of HighTemperature, Thermally Stable Alumina Composite Membrane,” J. Am. Ceram. Soc., 74 [I]219-24 (1991).

Membranes, Films, and Coatings

‘‘I. S. Church, N. W. Cant, and D. L. Trimm, “Stabilization of Aluminas by Rare Earth and Alkaline Earth ions,” Appl. Caral. A., 101. 105-16 (1993). “G. R. Gallaher and P. K. T. Liu, “Characterization ofCeramic Membranes I . Thermal and Hydrothermal Stabilities of Commercial 40 A Membranes,” J. Membr. Sci., 92, 2 9 4 4 (1994). ”C. H. Chang, R. Gopalan, and Y.S. Lin. “A Comparative Study on Thermal and Hydrothermal Stability of Alumina, Titania and Zirconia Membranes.” J. Membr.

Sci.,91, 2 7 4 5 (1994).

”A. Nijmeijer, H. Kruidhof, R. Bredesen, and H. Verweij, “Preparation and Properties of Hydrothemally Stable y-Alumina Membranes,” J. Am. Ceram. Soc., a,$’] 136-40 (2001). R. S. A. de Lange, J. H. A. Hekkink, K. Keizer, and A. J. Burggraaf, “Formation and Characterization of Supported Microporous Ceramic Membranes Prepared by SolGel Modification Techniques,’’ J. Memhr. Sci.,99, 57-75 (1995).

I5T. Tsuru, T. Hino, T. Yoshioka. and M. Asaeda, “Permporometry Characterization of Microporous Ceramic Membranes,” J. Membr. Sci., 186. 257-65 (2001). 16H. Schaper, E. B. M. Doesburg, and L. L. van Reijen, “The Influence of Lanthanum Oxide on the Thermal Stability of Gamma-Alumina Catalyst Sup ports,” Appl. Coral., 7,21 1-20 (1983). I7V. G. Hill, R. Roy, and E. F. Osborn, “The System AluminaGallia-Water,” J . Am. Ceram. S a c , 35, 136-42 (1952). “9. Beguin, E. Garbowski, and M. Primet, “Stabilization of Alumina by Addition of Lanthanum,” Appl. Cafal., 75, 119-32 (1991). I9Y. S. Lin, C. H. Chang, and R. Gopalan, “improvement of Thermal Stability of Porous Nanostructured Ceramic Membranes,” Ind. Eng. Chem. Rex, 33, 8% 70 (1994).

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Behavior of Silicon Oxycarbonitride Thin Films Derived from Poly(Urea)Methyl Vinyl Silazane Tsali J. Cross' and Rishi Raj University of Colorado at Boulder, Department of Mechanical Engineering, Boulder, Colorado 80309

Tsali J. Cross, Somuri V. Prasad, and David R. Tallant Sandia National Laboratories,* Materials and Process Sciences Center,Albuquerque, New Mexico 87185

A process for deposition of silicon oxycarbonitride films from poly(urea)methyl vinyl silazane (PUMVS) by spin coating precursor solutions onto a substrate, followed by polymerization, cross-linking and pyrolysis has been developed. The crosslinked polymer films (350 nm thick), deposited on variety substrates (e.g., silicon, sapphire, zirconia), were pyrolyzed in nitrogen or ammonia environments either in a hot isostatic press or in a tube furnace. Their microstructure was characterized using infrared and Raman spectroscopy. The tribological (friction and wear) behavior was evaluated in dry nitrogen and air with 50% relative humidity using a unidirectional linear wear tester in a ball-on-disk configuration. Wear surfaces, transfer films and wear debris were analyzed by scanning electron micrograph (SEM)/energy dispersive spectroscopy (EDS).

Introduction Silicon oxycarbonitride (SiCNO) ceramics are promising materials for microelectronics and optoelectronics, owing to high-temperature oxidation resistante, tunable band gap characteri~tics,~'~ adjustable transparency in the visible and IR and high-temperature thermal stability.6 Their refractory nature makes this class of materials attractive for tribo-

'

'Sandia is a multiprogram laboratory operared by Sandia Corporation, a Lockheed Martin Company, for the United States Department of Energy under contract DE-AC04-

~~AL~~OIIO. '[email protected]@sandia.gov 02006 The American Ceramic Society

Membranes, Films, and Coatings

logical applications in harsh environments. ',7-9 SiCNO films can be deposited by various techniques, including CVD," CVD bias-assisted hot filament," RF PECVD," electron cyclotron resonance PECVD,I3 pulsed laser d e p ~ s i t i o n , 'ion ~ beam sputtering,' RF sputtering," plasmajet CVD," reactive magnetron sputtering, and via silazane polymer-derived ceramic (PDC) routes. 1,19-23 Wear and friction (static or dynamic) limits the reliability of silicon microlectromechanical system (MEMS) devices, because of its high-wear and -friction coefficient. Surface modifications such as hard coatings, including Sic and W, solid lubricants, films deposited by atomic layer deposition, and adsorbed organic molecules have been explored for reducing friction and wear of

''

223

MEMS However, there is a limited selection of wear-resistant MEMS materials with low friction that can withstand high-temperature and corrosive environments. Deposition of SiCNO films by the polymeric, silazane-PDC routes, where liquid precursors are cast as thin films on a substrate and pyrolyzed in situ in the ceramic state, has additional benefits such as cost and ease of application. However, studies on the fundamental understanding of the mechanisms of friction and wear of the materials are limited. PDCs are refractory materials that have structure at the nano-scale and have high strength and hardPDCs are processed from liquid organometallic precursors by cross-linking the polymers into an infusible solid followed by controlled pyrolysis. PDCs remain inert and highly stable, both mechanically and chemically in harsh environ~nents.~~ Typical examples of PDCs include SiCNOs, and silicon oxycarbide (SiOC). The microstructure of PDCs depends on precursor synthesis, polymerization, degree of cross-linking, the temperature at which the cross-linked part is pyrolyzed, as well as the environment of the pyrolysis process. Polymer-derived silicon carbonitride typically contains phases of silicon atoms bonded to carbon and nitrogen in sp3 configurations, as well as phases of carbon in aromatic sp2 configuration^.^^ These ceramics retain their amorphous microstructure above 1400°C,32 where nucleation of S i c and Si3N4 phases is limited. It is speculated the carbon phases help to reduce coarsening of S i c and Si3N4 phases by physically and thermodynamically impeding their growth. The presence of carbonrich phases also presents an opportunity to tailor their microstructure and chemistry for improved friction and wear performance.

The objective of the current study was to synthesize SiCNO PDCs into thin films, characterize their microstructure and chemistry, and evaluate their tribological behavior in various environments.

Materials and Methods Synthesis and Processing of Thin Films

Thin films were processed with the liquid monomer poly(urea)methyl vinyl silazane (PUMVS, trade name CerasetTM,Kion Corp. Huntingdon Valley, PA). The chemical structure and a schematic representation of the polymerization steps through vinyl groups are shown in Fig. 1. The precursor for SiCNO films consisted of PUMVS, 30 wt% acetonitrile, and 5 wt% photoinitiator, 2,2 dimethoxy-2-phenyl-acetophenone (DMPA, Aldrich Inc., St. Louis, MO) that was polymerized with ultraviolet radiation. Acetonitrile served to improve precursor wettability and the dissolution of DMPA. Acetonitrile was selected as an appropriate solvent, based on its volatility with PUMVS and DMPA, and wetting behavior after spin coating. Tetrahydrofouran and heptane solvents also generated pinhole-free coatings after spin coating, but acetonitrile was selected because it is less toxic. The liquid monomer, PUMVS, was also used to spin coat and was thermally polymerized, cross-linked, and pyrolyzed in order to observe any effects that may result from the solvent and photoinitiator, DMPA. Films were deposited on commercially available wafers of (100) silicon (boron doped), (0001) sapphire (A1203), and partially yttria-stabilized zirconia (YZTP). The wafers were first cleaned ultrasonically

z -Si I

I

- CH =CH2

+tycHYH2 -Si

-

-Si

-

z- 1

H

(a)PUMVS

(b) Vinyl group polymerization

Fig. 1. (a) MoLecuLar structure of the monomer f i r poly(urea)rnethyl vinyl silazane (PUMVS), and (b) its polymerization through vinyl gfoups.

224

Progress in Nanotechnology: Processing

Fig, 2. Flow chart ofpolymer derived Silicon oycarbonitride thin film deposition process, (I) Thefilm is deposited by spin coating. (2) The coating polymerization was initiated either by ultraviolet or thermal energy. (3) The polymerized coating is pyrolyzed either in a hot isostatic press in nitrogen, or in a @wing tube fimace in nitrogen or ammonia gas.

in acetone, followed by methanol. They were then subjected to UV ozone cleaning for 15 min. Cleaning procedures, involving standard piranha solution (3: 1, H2SO*:H202) to first remove organics and a 48% hydrofluoric acid solution to remove native oxides, were found not to have an effect on film uniformity or its chemistry. Precursor solutions were dispensed using a syringe, and 0.2 pm syringe filter, with enough fluid to cover the entire 12 mm x 12 mm-diced wafers. The wafers were then spun at 4000 rpm for 30 s. The precursor films were polymerized either thermally at 350°C or by exposing to W radiation with a dose of 11.4 J/cm2 at 354 nm. They were cross-linked at 400°C under nitrogen in either a flowing tube furnace or in a hot isostatic press under 13.79 MPa (2000 psi) of isostatic pressure. One batch of cross-linked films were pyrolyzed in a hot isostatic press and flowing tube furnace in nitrogen, and another batch in a flowing tube furnace in ammonia. A flow chart of the steps in thin film synthesis is shown in Fig. 2. Sample inventories including the precursor, cross-linking, and pyrolysis

Membranes, Films, and Coatings

conditions are shown in Table I (nitrogen pyrolysis) and Table I1 (ammonia pyrolysis). Characterization

The structure and chemistry of the films on silicon were evaluated by Fourier transform infrared (FTIR), attenuated total reflectance (ATR), and microRaman spectroscopy. Raman spectra were obtained using the 458 nm line of an argon laser. The microscope accessory to the Raman system focused the laser beam on the samples, collected a 180" backscatter and transported the backscattered light to a triple spectrograph. A charge-coupled device detector detected photons dispersed by the spectrograph. We removed broad (fluorescence) background features from the resulting Raman spectra using a polynomial fitting technique. For FTIR, a spectrometer, fitted with a horizontal sample placement, W light exposure, and modified reading photoaccessory that collected data in transmission, was used. ATR spectra were obtained from an additional, separate

225

Table I.

Sample Inventory of Thin Films Cross-linked and Pyrolyzed in Nitrogen on Different Substrates Under Various Conditions Precursor, polymerization, cross-linking, and pyrolysis in nitrogen

Sample

Precursor

Cross-linking and pyrolysis Polymerization environment

FILM 1 PUMVS with 30 wt% Acetonitrile, 5wt% DMPA

W cured in open air or UV cured under nitrogen

Flowing tube furnace: nitrogen gas

FILM 2 PUMVS with 30 wt% Acetonitrile, 5wt% DMPA

W cured in

Hot isostatic press: nitrogen

FILM 3 Pure PUMVS

open air

gas

Thermally cured All processed in flowing tube in a glove box furnace under with 25 ppm nitrogen H 2 0 content

Cross-linking temperature control

Pyrolysis temperature control

FTIR results

HR1: 1200°C/h, HR2: 300"C/h, Dwelll: 90 rnin Dwell2: 60 or at 400°C 180 rnin at

Si bonds: Si-0, Si-C, Si-CH3 looooc C bonds: C-H HR1: 800"C/h, HR2: 40"C/h, Si bonds: Dwelll: 100 rnin Dwell2: 600 rnin Si-0, Si-C, at 400°C, at 1000°C Si-CH, C bonds: 14 MPa C-H HR1: 1200"C/h, HR2: 30O0C/h, Si bonds: Dwelll: 90 min Dwell2: 60 min Si-0, Si-C, Si-CH3 at 350°C at 825°C C bonds:

Raman results

Highly ordered graphene phases Highly ordered graphene phases

C-H FILM 4 PUMVS with on ZrOz 30wt% Acetonitrile, 5wt% DMPA

W cured in open air

Hot isostatic press: nitrogen gas

HR1: 800"C/h, Dwelll: 100 rnin at 400"C, 14 MPa

HR2: 40"C/h, Dwell2: 600 rnin at 1000°C HR3: 40"C/h, Dwell3: 600min at 1250°C

PUMVS, poly(urea)methyl vinyl silazane; DMPA, 2,2 dimethoxy-2-phenyl-acetophenone.

Table 11.

Sample Inventory of Thin Films Cross-linked and Pyrolyzed in Ammonia on Different Substrates Under Various Conditions Precursor, polymerization, cross-linking, and pyrolysis in ammonia

FILM 5 PUMVS with 30 wt% acetonitrile, 5wt% DMPA

W cured

FILM 6 PUMVS with

UV cured

30 wt%

acetonitrile, 5 w t % DMPA

226

Flowing in open air tube furnace: ammonia gas

Flowing in open air tube furnace: ammonia gas

HR2:

Si bonds: Ordered Si-0, graphene Si-N phases C bonds:

HR1 : 12OO"C/h, Dwelll: 90 min at 400°C

30O0C/h, Dwell2: 60 min at 1000°C

NO CCOSS-

HRl :

linking step

1200"C/h, Dwell2: 90 min at 1000°C

C-H Si bonds: No carbon Si-N detected C bonds:

C-H

Progress in Nanotechnology: Processing

horizontal sample mount with a diamond, singlebounce plate, and photoaccessory that was added on to the FTIR spectrometer. All infrared spectra from films were collected with FTIR in transmission, with the exception of FILM 4, where transmission through zirconia substrate was not possible and it was necessary to acquire FILM 4's spectrum with ATR. Film surfaces and cross-sections were examined in the SEM to assess the topography, substrate-film interfaces, and porosity in the films. Cross-sections of films were prepared by dicing the coated wafers after carefully scribing with a diamond-tipped pen at the back.

Tribological Testing Friction and wear tests were performed using a ballon-disk linear wear tester (Fig. 3). A Si3N4 ball (3.175 mm diameter) was used as the counterface. Normal load was applied by means of deadweights. Friction force was measured with a strain gage load cell. Measurements were made at three applied normal loads (i) 98 mN, which corresponded to "low" Hertzian contact stresses of 0.42, 0.58, and 0.62 GPa on silicon, zirconia, and sapphire substrates, respectively, (ii) an "intermediate" normal load of 294 m N corresponding to contact stresses of 0.61, 0.84, and 0.89 GPa on silicon, zirconia, and sapphire substrates, respectively, and (iii) a "high" normal load of 519mN, corresponding to contact stresses of 0.73, 1.01, and 1.08 GPa on silicon, zirconia, and sapphire substrates, respectively. The device was housed in an environmental chamber with provision to

Fig.3. A photograph of the linear wear tester employed for fiiction measurements.

Membranes, Films, and Coatings

monitor dew point and oxygen content. Mixing dry air with air passed through water bubblers provided the desired relative humidity in the test chamber. Measurements were made in a dry nitrogen environment (< 10 pprn 0 2 and < 100 pprn H 2 0 ) and in air with 50% relative humidity. Tests were run in unidirectional mode for 1000 cycles over a track distance of 1.6 mm at a sliding speed of 3.3 mm s-'. Results

Microstructure and Chemistry FTIR and ATR Spectroscopy: FTIR spectra of films before and after photoinitiated polymerization are shown in Fig. 4. The spectra were obtained on UVcured films in a nitrogen purged environment (Fig. 4b) as well as in air (Fig. 4c) in order to investigate the role of oxygen. Photoinitiated polymerization occurred through the vinyl groups as illustrated in Fig. l b where a reduction in the carbon-carbon double bonds (1600cm-') and increase in the methylene groups was evident. The bands near 2300cm-' (Fig. 4a) are because of nitrile groups within the solvent, acetonitrile. The solvent evaporates with the addition of heat generated from exposure to ultraviolet radiation, as shown by the disappearance of the nitrile bands in Figs. 4b and c. Homopolymerization of silazanes has been well characterized by Kroke et ~ 2 1 and . ~ ~others and are shown to polymerize through vinyl groups as well, along with hydrosilylation reactions at elevated temperatures. The presence of the strong Si-0 peak at 1100 cm-' for polymerizations performed in air shows that oxygen becomes chemically incorporated into the polymer network itself. FTIR spectra of the films after UV-initiated polymerization in an open air, and a nitrogen purged, low oxygen environment are shown in Fig. 5. The spectrum in Fig. 5c (FILM 3) is also representative of a 100% concentration PUMVS film, cross-linked at 400°C and pyrolyzed at 825°C in a nitrogen-purged glove box with 25 ppm H z O content, after etching the native oxide on the silicon substrate with 48% HF for 5 min. Because of the near absence of Si-N (900cm-') or N-H (3384cm-') bonds, it is believed that the thin films pyrolyzed between 700" and 1000°C in nitrogen are primarily composed of a material consisting of Si-0 (1098 cm-'), and Si-0-C bonds, as indicated by the peak between 800 and 850cm-' in Fig. 5c.

227

Fig. 4. Fourier transform infiared spectra of (a) poly(urea)methyl vinyl sihzune with 30 wt% acetonitrile and 5 wt% 2,2 dimethoxy-2phenyl-acetophenone, and (b) afer it was Wphotoinitiatedpolymerization in a nitrogen-purged, low oxygen environment, and (c) afer W photoinitiated polymerizution in air.

Films deposited on zirconia were further annealed at 1250"C, but their composition was not easily identifiable from its ATR spectrum, because of the heavily convoluted peak around 1100-1600 cm-' (Fig. 5d). Films deposited on silicon or sapphire evaporated almost entirely when annealed at 1250°C. FTIR spectra of films cross-linked and pyrolyzed in ammonia are shown in Fig. 6. Cross-linking of films at 400°C (Fig. 6a) before pyrolysis between 700" and 1000°C in ammonia, resulted in broadening of the base at 900cm-' of the large Si-0 peak at around llOOcm-' out to 900cm-', shown in Fig. 6b. The breadth of the underlying band around 3 100-3500 cm-', indicates the presence of 0-H groups. Thin films cross-linked and pyrolyzed in ammonia are therefore speculated to comprise of complex Si-0, Si-C, 0-H, and possibly Si-N bonds. The films that were directly pyrolyzed in ammonia, without cross-linking, showed the presence of N-H bonds at 3384 cm-', and a strong peak at 900 cm-' (Fig. 6c). The breadth of the under-

228

lying band around 3100-3500 cm-', indicates the presence of 0-H groups as well, but not conclusively. This suggests noncross-linked films are comprised of primarily SiN, bonds. Raman Spectroscopy: All films that were pyrolyzed in a flowing tube furnace in nitrogen, as well as HIPed films that were cross-linked under pressure, yield a Raman spectrum with only two peaks corresponding to the D and G bands, shown in Fig. 7. These Raman bands are characteristic of stacked sheets of six-membered, aromatic (graphene) carbon rings,34 similar to those in graphite. The D peak is typically located near 1350cm-', while the G peak is usually located near 1600 The shifts in these peaks, from 1350 and l6OOcm-', found in the spectra from films shown in Fig. 7, could be because of incorporation of silicon or nitrogen into the graphene Some Raman spectra, obtained from locations both on and off wear tracks, indicated reduction of graphene domain size,

Progress in Nanotechnology: Processing

Fig. 5. Typical Fourier transform infiared spectra o f j l m s at dzffient stages o f j l m processing. All films were processed in a nitrogen atmosphere and spectra are shown (a) after cross-linking at 400°C (6) after 600°C (c) (FILMS I , 2, 3) after 1000°C beat treatments (I and 3 b), and (d) after a 1250°C heat treatment (FILM 4).

presumably by crushing. No carbon Raman bands were found in the Raman spectrum for the ammonia-pyrolyzed film, that was not cross-linked before pyrolysis. We compared the Raman spectra of the PDC films in Fig. 7 to spectra obtained from samples of a phenolic composite pyrolyzed at temperatures from 375" to 2700°C. One of the Raman spectra of the phenolic composite samples, treated at 980°C is shown in Fig. 7. The Raman spectra of the pyrolyzed phenolic samples show the evolution of graphene-related Raman bands as graphite-like domains increase in size with increasing exposure temperatures. The spectrum of the film crosslinked (400°C) and pyrolyzed (1000°C) in ammonia is, except for enhancement in intensity and an upshift in frequency of its D peak, generally consistent with the spectrum of the 980°C phenolic composite. Enhancement in intensity and an upshift in frequency of the D peak is consistent with nitrogen incorporation into the graphene structure, based on Raman spectra of a pyrolyzed, nitrogen-containing polymer.36 The spectrum of the film cross-linked and pyrolyzed (1000°C) in nitro-

Membranes, Films, and Coatings

gen is intermediate in Raman band profile between spectra of phenolic composite samples treated at 980°C and 1200°C. However, its G peak is upshifted in frequency (at 1615 cm-' from about 1600 cm-' in spectra of pyrolyzed phenolic composites). Wang et aL3* attribute an upshift in G band frequency specifically to boron incorporation into the graphene matrix but more generally to graphene sheets adjacent to layers incorporating intercalant atoms. The Raman spectra in Fig. 7 of the ammonia and nitrogen pyrolyzed and cross-linked films are generally consistent with the formation of stacked graphene ring structures formed by thermal charring, as occurs in a phenolic resin matrix. The observed shifts in their D and G peak frequencies are because of incorporation of noncarbon atoms in the graphene matrix. Topography and Cross-Sectional SEM: A SEM of typical film surface, regardless of pyrolysis conditions, is shown in Fig. 8a. The topography was measured by optical profilometry. The average surface roughness (4)was

229

600

850

1100 1350 1600 1850 2100 2350 2600 2850 3100 3350 3600 3850

Wavenumber (cm-’)

Fig. 6 Typical Fourier transform inpared spectra o f j l m at diferent stages o f j l m processing. Alljlm were processed in an ammonia atmosphereand spectra are shown (a) after cross-linkingat 400°C in a flowing tubefitrnace in ammonia, (6) (FILM 5) after cross-linking and 1OOOOC pyrolysis, and (c) (FILM c;) after 100O”Cpyrolysisonly, without a cross-linking heat treatment.

6.7nm. Film thickness was measured from the SEM micrographs of the cross-sections; the average film thickness was 350 f40 nm. Films also appear to adhere well with all the substrates without any interfacial porosity; a typical cross-section of the film on silicon is shown in Figs. 8a and b. Films that were not cross-linked before pyrolysis were found to be porous (Fig. 8b) whereas cross-linking before pyrolysis resulted in dense films (Fig. 8c). This observation suggests preparation of a dense ceramic film requires a minimum degree of cross-linking before pyrolysis.

Tribology Fig. 7. Typical Raman spectra of silicon oxycarbonitridefilm cross-linked andpyrolyzed in nitrogen (FILMS I, 2) and ammonia (FILM 5). Also shown is a comparison to carbon referencefiom a phenolic char annealed at 980°C. Narrow bands are indicative of extended graphene structure.

230

Typical friction traces (friction coefficient as a hnction sliding cycles) of films processed in nitrogen environment (FILMS 1, 2) are given in Fig. 9. These measurements were made in dry nitrogen environment

Progress in Nanotechnology:Processing

Fig. 3. Friction coeficient versus number of sliding cycLes by a linear wear tester in a d?y nitrogen environment with contact stress 0.42 GPa tested on film deposited on silicon. Tbejiction traces are representative ofjilms cross-linked and pyrohzed in a nitrogen environment (FILM I), that were (a) polymerized in open air, and (b) polymerized and processed in a nitrogen-purged glove box

(FILM 3).

Fig 8. Scanning electron microscope micrograph of silicon OxycarbonitriAfilms on silicon substrates: (a) topography of a vpicaljlm, (6) cross-section of a Jilm that was cross-linked and pyrolyzed (FILM 5), and (r) cross-section of a pyrolyzed film that was not cross-linked (FILM G). A l l j l m were processed in ammonia.

at an applied load of 98 mN. FILM

3 was processed in

the glove box including the spinning and polymerization steps, whereas FILM 1 was spun and polymerized in air followed by cross-linking and pyrolysis in nitrogen. The film that was spun and polymerized in open lab air had an extremely short wear life (Fig. 9a). SEM micrographs of the wear scar clearly reveals generation

Membranes, Films, and Coatings

of large amount wear debris (Fig. lOa), partial exposure of substrate silicon, and transfer of the debris to the counterface ball (Fig. 1Oc). The EDS analysis on wear surface in Fig. 10b showed predominantly silicon and oxygen peaks. This behavior was representative of all the films that were spun and polymerized in lab air, irrespective of the subsequent processing steps in nitrogen. By contrast, films processed in the nitrogen glove box all the way beginning with spinning showed a steady state friction coefficient of 0.21 0.01 throughout the 1000cycle test. Films cross-linked and pyrolyzed in ammonia showed a progressive decrease in friction coefficient with sliding distance. As shown in Table 11, the films were spun and polymerized in lab air before cross-linking and pyrolysis in an ammonia atmosphere. Figure 11 shows a typical friction trace in dry nitrogen environment at a contact stress of 0.42 GPa on a silicon substrate. SEM micrograph of a wear scar on FILM 5 is shown in Fig. 12a. A higher magnification micrograph of the wear scar shows portions of the film inside the wear scar curling up. Figure 12c corresponds to the wear surface of the counterface ball, where in the presence of large number of rolled up debris can be seen. Friction behavior of films (pyrolyzed in nitrogen and ammonia atmospheres) in humid air is shown in

231

1.01

0.8

,

I

,

.

I

I

I

I

800

1000

i

0

200

400

600

Cycles

Fig. 11. Friction coeficient versus number of sliding y l e s by a Linear wear tester in a d y nitrogen environment with contact stress 0.42 GPa on a silicon substrate. Friction traces are representativeof

(FILM 5).

Deep grooves on the ball surface and ploughing of the silicon substrates are also evident in the SEM micrographs in Fig. 14. These results imply that the films, irrespective of the processing conditions, wore out almost instantly exposing the underlying substrate to silicon nitride ball counterface. Particulate wear debris at the sliding interface might have been responsible for the observed friction coefficients that are higher than those of the bare silicon. Discussion

Fig. 10. Scanning electron micrographs of a wear track at (a) low and (b) high magnification, and (c) the Si@*counte$ace ball aft.. a linear wear test performed in d y nitrogen on a j l m cross-linked and pyrolyzd in nitrogen (FILM 2), and deposited on silicon.

Fig. 13. It can be seen that the friction coefficients are high (- 0.8). It should also be noted from Table I11 that the friction coefficients are in fact higher than those measured on bare silicon substrates. Subsequent SEM analyses did not reveal wear debris roll-ups either on the wear surface or on the counterface ball (Figs. 14a and b).

232

A summary of the friction behavior for all films at various contact pressures and environments is shown in Table 111. There appears to be a strong correlation between film density, and pyrolysis atmosphere, and the friction and wear behavior. The degree of cross-linking before pyrolysis affected the rate of diffusing gas, and density of the films. Nitrogen pyrolysis produced a microstructure consisting of an amorphous Si-0, Si-C matrix surrounding graphene structures whose ordering is consistent with their thermal exposure, as revealed by FTIR (Fig. 5) and Raman spectroscopy (Fig. 7). Pyrolysis in ammonia produced a microstructure comprised of an amorphous matrix of Si-0, 0-H, Si-C, and possibly Si-N bonds surrounding graphene nano-clusters, as revealed by FTIR (Fig. 10) and Raman spectroscopy (Fig. 7). The

Progress in Nanotechnology: Processing

Fig. 13. Friction coe$cient versus number of sliding cycles by a linear wear tester in air with 50% relative humidity environment with contact stress 0.42 GPa. The friction traces are representative offilms deposited on silicon that were cross-linked andpyrolyzed in either a hot isostatic press or flowing tube&rnace in a nitrogen environment (FILMS I , 2), and cross-linked and pyrolyzed in a flowing tube&rnace in ammonia (FILM 5).

Si-CH3 E SiZ-NH

+ NH3 -+=Si-NH2 + CH4

+ H20

-+

= Si-NH2

= Si2-0 + NH3

Fig. 12. Scanning ekctron micrographs of a wear track at (a) low and (6) high magnification, and (c) the Sifl4counteface ball ajer a linear wear testpefomed in dry nitrogen on a film cross-linked andpyrolyzed in ammonia (FILM 5), and deposited on silicon.

retention of Si-N or Si-0 bonds during pyrolysis is anticipated from reactions with ammonia or with latent water, respectively as follows:

Membranes, Films, and Coatings

(a)

(b)

Several films were prepared to investigate the influence of the processing parameters, including cross-linking, pyrolysis environment, and the temperature cycle on the composition of the films, specifically nitrogen and oxygen contents. Subtle changes in their FTIR spectra indicated that the dehydrogenation reactions for the films are successful when the duration of crosslinking dwell times at 350°C are longer than 90min, and when the films are pyrolyzed at 700°C for longer than 15 min. However, their FTIR spectra of pyrolyzed films did not differ from those that were shown in either Fig. 5c or Fig. bb. Further, an observable change in the FTIR spectra was not shown even with a 95%Ar, 5%H2 mixture of gas that was employed during pyrolysis. The films with the lowest friction (in dry nitrogen environment) were obtained by pyrolyzing in an ammonia atmosphere, where, presumably more nitrogen may have been incorporated within the film during pyrolysis, based on the reaction in equation (a). Their

233

Table 111. Summary of Typical Friction Results Summary: friction and wear of films after 1000 cycles COF in dry nitrogen

Sample

COF low contact stress (GPa): Substrate 0.42 < ts < 0.62

Films pyrolyzed in nitro en Bare zro2 ~ 0 % 0.8f0.1 : Bare A1203 COF: 0.25 f0.02 Bare Si COF: 0.5 f 0 . 2

FILM FILM FILM FILM FILM

1

Si

2 2 3 4

A1203

A2O3 Si

Zr02

COF: COF: COF: COF: COF:

0.7 k 0.1 0.7k0.2 0.7 f0.1 0.21 kO.01 0.25 f 0.02

Films pyrolyzed in ammonia FILM 5 Si COF: 0.1 fO.01 COF: 0.22f0.02 FILM 5 A 1 2 0 3 FILM 6 Si COF: 0.7 f0.1

COF intermediate contact stress (GPa): 0.61 < ts < 0.89

COF: 0.9 f0.05

COF: 0.75 f0.05 COF: 0.65 f0.2

COF: 0.75 f0.05 COF: 0.25 k 0.1 20% RH in Air COF: 0.6 k0.1 COF: 0.7f0.2 COF: 0.7 f0.3 COF: 0.7 f0.1

COF: 0.25 f0.02

COF: 0.4 f0.02

COF: 0.67 f0.03

COF: 0.1 fO.01 COF: 0.45f0.05

COF: 0.5 f0.05 COF: 0.6f0.1

COF: 0.75 & 0.05 COF: 0.8 rfr 0.02 COF: 0.7 f0.1

low-friction behavior is a result of the formation of sheat-reducing roll-ups. However, these films do not have inherent resistance to wear, and the low friction is because of the formation of roll-ups. All the films showed high-friction and -wear behavior in humid environments, because of ploughing of the film from the formation of a harder third body phase of unknown composition, probably Si-0. It is speculated that nitrogen within films may improve its hardness and stiffness, while oxygen may deteriorate its mechanical properties, as would be expected for amorphous Si-N versus Si-0 based ceramics, in general. Residual stresses within the films may also play an important role, helping to facilitate roll-up formation. Further characterization of the mechanical properties was not pursued, because the films did not have the desired friction and wear response, particularly for microsystem applications. Although there were differences in the ordering of the carbon phases within nitrogen and ammonia pyrolyzed films, its effects on tribological behavior are not understood. However, it is clear that the graphene/ graphitic regimes did not dominate the tribological be-

234

COF high contact stress (GPa): 0.73 < ts < 1.08

COF in 50% RH in air COF low contact stress (GPa): 0.42 < a < 0.62

havior, because high friction was observed in humid environments. The lubricity of graphite in air is related to its lamellar, hexagonal sttucture. It is generally accepted that the ease of slip of the graphite basal planes increases from intercalation of adsorbed water and oxygen.37’38 Ongoing research on bulk polymer-derived SiCNO ceramics showed that the bulk materials exhibited friction coefficients of 0.2f0.25 in both dry and humid environments and they were extremely resistant to wear.39 This ongoing study revealed that pyrolyzing the material at temperatures greater than 1200°C was essential to achieve the desired tribological behavior. The high-temperature anneal has the effect on bulk SiCNO specimens in increasing its hardness and stiffness. Therefore, it was decided that higher temperatures were required to improve the tribological behavior of films. It is difficult to subject the films to high-temperature anneals because of the sublimation of the film and the reaction with the substrate. Despite these challenges, films were deposited on zirconia substrates and annealed at 1250°C. Unfortunately, these films were extremely thin, because of evaporation or sublimation during the

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0.8 -

0.0

0

-Annealed

200

film on zirconia

400 600 Cycles

800

1000

Fig. IS. Friction coefficient versus number of sliding ycles by a linear wear tester in a dry nitrogen environment with contact stress 0.42GPaforfilms deposited on zirconia substrates. Friction trace is typical f i r (FILM 4).

Conclusion

We have developed a process for deposition of SiC-

NO films from PUMVS. Films that were spun and po-

Fig. 14. Scanning electron micrographs of the (a) wear track and (6) Si&dcountprf.ce ball a j e r a linear wear test peformed in air with 50% relative humidity on a film cross-linkedandpyrolyzed in ammonia (FILM 5), and deposited on silicon.

anneal, which made it difficult to characterize them with

ATR infrared spectroscopy. EDS revealed that the film contained Zr, 0, and C, but not silicon or nitrogen. So, we are not sure if this particular film indeed belonged to the family of SiCNO PDCs. However, its friction coefficient remained low (0.25 0.02) and steady throughout the 1000-cycle test (Fig. 15). Furthermore, the wear scars were so shallow that they could not be identified with certainty even with a high-resolution SEM. “Low-temperature” pyrolysis techniques such as laser pyrolysis,’ and other may eventually be employed to apply thin film PDC’s onto low-melting substrates, and help alleviate any effects from thermal mismatch. However, before these engineering issues may be addressed, a better understanding of the relationship between the desirable chemistry, microstructure, and tribology of PDCs should be gained.

Membranes, Films, and Coatings

lymerized in air, and then pyrolyzed in nitrogen wore through immediately in dry nitrogen environment, which produced high friction (0.8). However, films that were spun, polymerized, and pyrolyzed (in nitrogen) in a low-HzO glove box, had reduced friction coefficients. Films pyrolyzed in ammonia showed an initial coefficient of friction of 0.2, which decreased progressively with sliding distance. The mechanism of low friction is because of the formation of wear debris roll-ups. All films wore through in humid environments without forming friction-reducing roll-ups. Cross-linking was the important processing parameter in affecting the density and composition of the films. Films deposited on zirconia substrates and annealed at 1250°C showed promising tribological behavior. References R. Riedel, H.-J. Kleebe, H. Schonfelder, and F. Aldinger, “A Covalent Micro/ Nano-Composite Resistant to High-Temperature Oxidation,” Nature, 374 526-528 (1995). A. Badzian, T. Badzian, R. Roy, and W. Drawl, “Silicon Carbonitride, a New Hard Material and Its Relation to the Confusion About ‘Harder Than Diamond’ C3N4,” Thin Sol. Films, 354 148-153 (1999). D. H. Zhang, Y. Gao, J. Wei, and Z. Q. Mo, “Influence of Silane Partial Pressure on the Properties of Amorphous SiCN Films Prepared by ECRCVD,” Thin Sol. Films, 377-378 607-610 (2001).

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F. J. Gomez, P. Prieto, E. Elizalde, and J. Piquerasa, “SiCN Alloys Deposited by Electron Cyclotron Resonance Plasma Chemical Vapor Deposition,” Appl. Phys. L e n , 69 773 (1996). 5. G. Soto, E. C. Samano, R. Machorro, and L. Cota, “Growth of S i c and SiCJ, Films by Pulsed Laser Ablation of S i c in Ar and N2 Environments,” /. Vac. Sci. Technol., A, Vac. Surf: Films, 16 1311-1315 (1998). 6. H. J. Seifert, J. Peng, H. L. Lukas, and F. Aldinger, “Phase Equilibria and Thermal Analysis of Si-C-N Ceramics,” J Alloys Camp., 320 251-261 (2001). 7. H.-J. Krauss and G. Mom, “Laser Pyrolysis of Polysilazane A New Technique for the Generation of Ceramic-like Coatings and Structures,” Kq Ens Muter., 206213 467-470 (2002). 8. P. Jedrzejowski, J. Cizek, A. Amassian, J. E. Hemberg-Saphieha, J. Vlcek, and L. Martinu, “Mechanical and Optical Properties of Hard SiCN Coatings Prepared by PECVD,” Thin Solid Films, 447-448 201-207 (2004). 9. E. Bertran, E. Martinez, G. Viera, J. Farjas, and P. Roura, “Mechanical Properties of Nanometric Structures of SilSiC, C/SiC and ClSiN Produced by PECVD,” Dia. Rel. Muter., 10 1115-1 120 (2001). 10. A. Bendeddouche, R. Berjoan, E. Beche, T. Merle-Mejean, S. Schamm, V. Serin, G. Taillades, A. Pradel, and R. Hillel, “Structural Characterization of Amorphous SiCxNyChemical Vapor Deposited Coatings,” /. App. Phys., 81 [91 6147-6154 (1997). 11. 2. Gong, E. G. Wang, G. C. Xu, and Y. Chen, “Influence of Deposition Condition and Hydrogen on Amorphous-to-Polycrystalline SiCN Films,” Thin SolidFilms, 348 [l-21 114121 (1999). 12. H. Sachdev and P. Scheid, “Formation of Silicon Carbide and Silicon Carbonitride by RF Plasma-CVD,” Dia. Rel. Muter., 10 1160-1 164 (2001). 13. D. H. Zhang, Y. Gao, J. Wei, and 2. Q. Mo, “Influence of Silane Partial Pressure on the Properties of Amorphous SiCN Films Prepared by ECRCVD,” Thin Solid Films, 377-378 607-610 (2000). 14. T. Tharigen, G. Lippold, V. Riede, M. Lorenz, K. J. Koivusaari, D. Lorenz, S. Mosch, P. Grau, R. Hesse, P. Streubel, and R. Szargan, “Hard Amorphous CSiJy Thin Films Deposited by RF Nitrogen Plasma Assisted Pulsed Laser Ablation of Mixed GraphitelSi3N4-Targers,” Thin Solid Films, 348 103-1 13 (1999). 15. 2. He, G. Carter, and J. S. Colligon, “Ion-Assisted Deposition of C N and Si-GN Films,’’ Thin SolidFilms, 283 90-96 (1996). 16. K. B. Sundaram and J. Alizadeh, “Deposition and Optical Studies of Silicon Carbide Nitride Thin Films,” Thin Solid Films, 370 151-154 (2000). 17. J. Wilden, A. Wank, M. Asmann, J. V. R. Heberlein, M. 1. Boulos, and F. Gitzhover, “Synthesis of Si-C-N Coatings by Thermal Plasmajet Chemical Vapour Deposition Applying Liquid Precursors,” Appl. Oqanomct. Chon., 15 841-857 (2001). 18. T. Berlind, N. Hellgren, M. P. Johansson, and L. Hultman, “Microstructure. Mechanical Properties, and Wetting Behavior of Si-C-N Thin Films Grown by Reactive Magnetron Sputtering,” Surf Coat. Technol., 141 145-1 55 (2001). 19. Y. W. Bae, H. Du, B. Gallois, K. E. Gonsalves, and B. 1. Wilkens, “Structure and Chemistry of Silicon Nitride and Silicon Carbonitride Thin Films Deposited from Ethylsilazane in Ammonia or Hydrogen,” Chon. Muter., 4 478483 (1992). 20. J. Bill and D. Heimann, “Polymer-Derived Ceramic Coatings on ClC-Sic Composites,”/. Eur. Ceram. SOL.,16 1115-1120 (1996).

4.

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21. M. R. Mucalo, N. B. Milestone, 1. C. Vickridge, and M. V. Swain, “Preparation of Ceramic Coatings from Pre-Ceramic Precursors: Part 1 S i c and “Si3N&i2N20’ Coatings on Alumina Substrates,” /. Muter. Sci, 29 44874499 (1994). 22. G. Motz and G. Ziegler, “Simple Processibiliry of Precursor-derived SiCN Coatings by Optimised Precursors,” Kq Ens Muter., 20621 3 475-478 (2002). 23. Y. D. Blum, R. M. Platz, and E. J. Crawford, “Glass Strengthening by Polymer-Derived Ceramic Coatings,” /. Am. Ceram. Soc., 73 [ l ] 170-172 (1990). 24. R. Maboudian, W. R. Ashurst, and C. Carraro, “Tribological Challenges in Micromechanical Systems,” Tribol. Lett., 12 95-100 (2002). 25. S. A. Henck, “Lubrication of Digital Micromirror Devices,” Tribal. Lett., 3 239-247 (1997). 26. T. M. Mayer, J. W. Elam, S. M. George, P. G. Kotula, and R. S. Goeke, “Atomic-Layer Deposition of Wear-Resistant Coatings for Microelectromechanical Devices,” Appl. Phys. Lett., 82 2883-2885 (2003). 27. L.-A. Liew, W. Zhang, L. An, S. Shah, R. Luo, Y. Liu, T. Cross, M. L. Dunn, V. Bright, J. W. Daily, and R. Raj, “Ceramic MEMS New Materials, Innovative Processing and Future Applications,” Am. Ceram. SOC.Bull., 80 [5] 2530 (2001). 28. R Riedel, G. Passing, H. Schonfelder, and R. J. Brook, “Synthesis of Dense Silicon-Based Ceramcis at Low Temperatures,” Nature, 355 [20] 7 1 4 7 1 7 (1992). 29. J. Bill, J. Seitz, G. Thurn, J. Durr, J. Canel, B. Z. Janos, A. Jalowiecki, D. Sauter, S. Schempp, H. P. Lamparter, J. Mayer, and F. Aldinger, “Structure Analysis and Properties of Si-C-N Ceramics Derived from Polysilazanes,” Phys. Stam Solidi A, [166] 269-296 (1998). 30. R. Riedel, “Advanced Ceramics from Inorganic Polymers”; pp. 1-50 in MateriuL Science and Technology: A Comprehensive Treatment, Vol. 178. Processing of Ceramics, Pan II. Ed., R. J. Brook. VCH, Wurzburg, Germany, 1996. 31. J. Bill, J. Schumacher, K. Miller, S. Schempp, J. Seia, J. Durr, H. P. Lamparter, J. Golczewski, J. Peng, H. J. Seifert, and F. Aldinger, “Investigations into the Structural Evolution of Amorphous SiCN Ceramics from Precursors,” 2 Metallkd, 94 [4] 335-351 (2000). 32. S. R. Shah and R. Raj, “Mechanical Properties of a Fully Dense Polymer Derived Ceramic made by a Novel Pressure Casting Process,” Acta Mater., 50 [16] 4093-4103 (2002). 33. E. Kroke, Y.-L. Li, C. Konetschny, E. Lecomte, and C. Fasel, “Silazane-Derived Ceramics and Related Materials,” Mater. Sci. Eng. R. A Rev. /., 26 97199 (2000). 34. Y. Wang, D. C. Alsmeyer, and R. L. McCreery, “Raman Spectroscopy of Carbon Materials: Structural basis of observed spectra,” Chem. Muter., 2 557-563 (1990). 35. A. C. Ferrari and J. Robertson, “Interpretation of Raman Spectra of Disordered and Amorphous Carbon,” Phys. Rev. B, 61 [20] 14095-14107 (2000). 36. D. Tallant Polyacrylonitrile, personnel communication. 37. M. Hokao, S. Himnaka, Y. Suda, and Y. Yamamoto, “Friction and Wear Properties of GraphitelGlassy Carbon Composites,” Wear, 237 54-62 (2000). 38. P. J. Bryant, P. L. Gutshall, and L. H. Taylor, “A Study of Mechanisms of Graphite Friction and Wear,” Wmr, 7 118-126 (1964). 39. T. Cross, S. V. Prasad, D. R. Tallant, and R. Raj. MANUSCRIPT IN PROGRESS.

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Synthesis and Tribology of Carbide-Derived Carbon Films A. Erdemir* and A. Kovalchenko Argonne National Laboratory, Energy Technology Division, Argonne, Illinois 60439

C. White, R zhu, A. Lee, and M. J. McNallan Department of Civil and Materials Engineering, University of Illinois, Chicago, Illinois 60607

B. Carroll and Y. Gogotsi Department of Materials Science and Engineering and A. J. Drexel Nanotechnology Institute, Drexel University, Philadelphia, Pennsylvania 19104

Carbide-derived carbon (CDC) films are produced at atmospheric pressure on the surfaces of carbide-based ceramic materials and coatings by a high-temperature chlorination process. These nanoporous carbon films contain carbon nanoonions and amorphous carbon, and may contain nanocrystalline diamond and graphite as well, depending on the synthesis conditions. The combination of such diverse carbon phases in one material or coating provides unique and potentially useful properties for a wide range of engineering applications. In this paper, we will present the results of a comprehensive study on the tribological behavior of these films. The friction coefficient of C D C in open air is comparable with that of graphite and is typically in the range of 0.15-0.25. However, the friction coefficients of C D C tend to decrease with decreasing humidity. In dry nitrogen, its friction coefficient is -0.1 or less. Such behavior is in contrast to that of crystalline graphite, which normally exhibits low friction at high humidity, but high friction at low humidity or in vacuum. The friction coefficient of C D C becomes increasingly lower under heavier loads; however, increasing sliding velocity does not seem to affect its frictional behavior significantly. Using a hydrogenation process that removes residual chlorine from the C D C film, the friction coefficients of CDC can be further lowered to values as low as 0.03. In an attempt to understand some of the underlying mechanisms, we carried out comprehensive chemical and structural studies of the sliding surfaces as well as bulk films and correlated these findings with the friction and wear behavior of C D C films.

Introduction Carbide-derived carbon ( C D C ) is a novel nanostructured material that holds promise for numerous ‘[email protected]

02006 The American Ceramic Society

Membranes, Films, and Coatings

industrial applications such as catalysis, hydrogen storage, molecular sieves, rolling, rotating, and sliding tribological systems.lP2It is produced o n the surfaces of metal carbides such as S i c a n d TIC by a high-tem erature chlorination process at atmospheric pressure?”Speci fically, the carbide-forming element is selectively etched out in a halogen-containing carrier gas at temperatures

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ranging from 600°C to 1200°C. The carbon atoms left behind rearrange themselves and form the layer of CDC. While pure chlorine can be used, the typical carrier gas composition for chlorination of the specimens described in this paper consists of Ar and small amounts (i.e., 3-10%) of Clz. Hydrogen may also be added to this mixture to control the reaction kinetics and, hence, the growth rates and structure of the CDC films. Because of the higher thermodynamic stability of metal chlorides in comparison with CCl, within the range of 600-10OO0C, chlorine reacts preferentially with the metallic part in carbides and leaves carbon behind as a thin film. Recent electron microscopic studies have revealed that CDC may contain not only amorphous carbon and turbostratic graphite but also nanocrystalline diamond and nested fullerene nanostructures such as nanotubes and nano-onion~.~ The process used in the production of CDC is unique mainly because of its ability to produce diamond-structured carbon at atmospheric pressure and temperatures below lOOO"C.* This process does not require the creation of a gas discharge plasma or other forms of high-energy activation that are needed for the synthesis of diamond films. Unlike diamond and diamond-like carbons (DLC), CDC can be grown in the thickness range from 1 to more than 100 pm at very high growth rates. Because of their high internal stresses and very low deposition rates, diamond and DLC films may not be grown very thick. In most instances, the DLC films are less than 5 p m thick, while diamond could be up to 2Opm in thickness. Because of the very low growth rates, it may take a very long time to deposit such carbon films on suitable substrates. The growth rate of CDC films is very high and takes only a few hours to produce films that are tens of micrometers thick. As C D C is derived from carbidebased substrates, there is no significant change in the original dimensions of the work pieces. At or near the interfaces between CDC and carbide-based substrates, a chemically graded interface provides a very smooth transition from a high hardness value near the CDC/SiC interface to a lower value at the external surface. This graded interface reduces the susceptibility of CDC to spallation and delamination in comparison with diamond or DLC films, which are vulnerable to spallation due to the build up of very high internal stresses at or near the interface. Because of its unique nanostructure, CDC possesses very attractive physical, mechanical, and tribological properties. In particular, recent systematic studies in

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our laboratory have confirmed that CDC can afford very low friction and wear to sliding tribological interfaces. Specifically, depending on the test conditions, this coating can provide friction coefficients of 0.03-0.3 and wear rates of 10-9-10-7 mm3/N m.5-8 Therefore, C D C holds promise for large-scale applications in rotating equipment industries where S i c is used as a higher end seal material. In recent years, CDC films were further optimized to yield much improved friction and wear properties. Specifically, the growth kinetics of these films were thoroughly examined in gas mixtures including Ar, C12, and H2 at temperatures ranging from 600°C to 1200°C. As a result of these systematic studies, the growth conditions that can lead to the formation of an ideal film were determined. Extensive mechanical and tribological studies of the optimized films have indicated a close relationship between the growth conditions and the performance of these coatings.'-* Specifically, the hardness of the CDC layers can be increased by reducing the temperature of the treatment and blending the chlorinating gas mixture with hydrogen. In this paper, we concentrate mostly on the mechanical and tribological properties of highly optimized CDC films on Sic. Furthermore, we explore the role of hydrogen in the friction and wear behavior of the CDC films. During our studies, we used Raman spectroscopy and scanning electron microscopy (SEM) to determine the microstructure and chemistry of the CDC films and correlated these findings with the friction and wear data obtained under a wide range of test conditions.

Experimental Details Synthesis of CDC Films The C D C films can be produced on the surfaces of all kinds of carbide-based bulk ceramics and their thin coatings. One of the most common carbide ceramics is S i c (HexoloyTM, Saint Gobain Ceramics, Niagara Falls, NY), which we have used extensively in our previous studies over the years. Figure l a shows a schematic representation of the chlorination process for which a short description is provided below, but more details are available in recent papers by Gogotsi and colleagues14 The interface between CDC film and Sic substrate is physically rugged and chemically graded, as shown in Fig. Ib. The S i c substrates used in our study were about 6 mm thick and prepared in disk and/or square shapes.

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Fig. 2. Photographs of (a) as-received and (6) carbide-derived carbon treated at 1000°C S i c disks showing a black carbon layer on the exposed suface. Fig. 1. (a) Schematic representation of the chlorination process used in the production of carbide-derived carbon (CDC) on S i c and (b) illustration of the graded inteface between CDC and Sic.

The nominal diameter of the disks was about 30mm, while squares were roughly 25.4 mm x 25.4 mm. One of the faces of these flat samples was highly polished to an RMS surface finish of better than 0.01 pm. Chlorination of disk and square S i c substrates was performed in gas mixtures consisting of various ratios of

Membranes, Films, and Coatings

C12:Hz and Ar at 1000-1100°C for 3-6h. The gases used in the tube furnace were supplied from cylinders of pure C12, pure Ar, and H2. They were passed through anhydrous CaSO4 and concentrated HZS04 to remove any traces of water vapor. The typical gas mixture used in our study included 9% C12-4.5% H2-86.5% Ar (by volume). The photographs in Fig. 2 show one of the S i c disks before and after C D C treatment. The disk changed its color from gray to black after chlorination. The coating thickness ranged from 20 to 1OOpm and

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average. Further details of the test method used in the tribological testing of CDC films can be found in ASTM G99, Section 3 (Metals Test Methods and Analytical Procedures, Volume 3.02, Wear and Erosion; Metal Corrosion).

Results Microstructure and Mechanical Properties Fig. 3. Scanning electron microscopy image showing the morphology of a carbide-derived carbon-coated su face at high magnifcation. The su face becomes ve’y rough ajer the chlorination treatment.

depended on the process parameters, such as temperature, gas flow rate, gas mixture, and duration. Figure 3 shows the physical condition of the as-produced CDC on S i c surface. As is clear, the surface of the CDC layer is rather rough after chlorination and removal of the outer powdery material. Such a rough surface might cause additional frictional forces during tribological testing; therefore, before subjecting CDC to tribological tests, we performed a very brief polishing on these surfaces. Some of the polished samples were hydrogenated by annealing at 800°C for 8 h in 5% H2/95% Ar. The treated specimens were analyzed by micro-Raman spectroscopy using an Ar ion laser (514.5 nm line) at x 500 magnification with a spot size of 2 pm. Post-treatment polishing of some samples to a mirror finish was performed manually on a polishing wheel using an alumina suspension.

Some of the chlorinated S i c samples were sectioned and mounted on an SEM stage for microscopic examination. Figure 4 shows the cross-sectional morphology of a typical CDC film produced in a Cl-Hz gas mixture at 1000°C. As is clear, the film on top has a distinct microstructure, and the interface is rugged, as shown in Fig. I b. The columnar morphology typical of most physical vapor-deposited films is not visible on CDC films, and except for some polishing marks, the overall microstructure is very dense and featureless. The polishing marks are most likely due to the fact that CDC is softer than S i c and hence can readily be scratched during the polishing of the harder Sic. The CDC film shown in Fig. 4 is about 50 pm thick and appears to be strongly bonded to its substrate. The microcracks shown in this photomicrograph were not there originally, but generated from a Vickers indentation close to the interface. Because of the graded interface and strong bonding between CDC and Sic, the crack propagated into the CDC and away from the interface. The cracks deflected toward the CDC instead of following the interface because CDC is under a tensile stress. This is supported by

Friction a n d Wear Tests The friction and wear properties of CDC films and as-received Sic samples were evaluated in a ball-on-disk machine under a wide range of test conditions. The contact load varied between 2 and ION, while the sliding velocity was in the range of 10-1 5 cm/s. Tests were performed in a number of environments, including dry nitrogen and open air (with 17-32% relative humidity). The total sliding distance varied between 120 and 180 m, depending on the wear track diameters. The test balls (9.55 mm in diameter) were made of sintered Si3N4 and had a surface finish of better than 0.01 pm centerline

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Fig. 4. Cross-sectional micrograph of a carbide-derived carbon film grown on S i c The microcrack shown was caused by a Vickers indentation near the interjace.

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0.8 nm. Therefore, the softness of these films may have, in part, been due to the porous nanostructure. The hardness of the transition layer between CDC and S i c was much higher, that is, 20-35 GPa. Films produced at 600-800°C were two to three times harder and more resilient than the films produced at or above 900°C.

Tribological Properties Fig. 5. Typical Raman spectrum of a carbide-derived carbonfilm used in our studies. This particular sample was chlorinated in 2:l Cl*:H,for 6 h at 1000°C.

numerous ripples and cracks in the outermost layer of CDC seen in Fig. 3. Detailed studies by Raman spectroscopy revealed two principal peaks for CDC films, as shown in Fig. 5. One of the peaks was centered at around 1340cm-', while the other was at 1590 cm-'. The position of these two peaks is close to that of D and G bands of graphite. The D band is downshifted, while the G band is upshifted compared with microscrystalline graphite. A much higher intensity of disorder-induced D band compared with G band, corresponding to in-plane vibrations of carbon atoms in graphite, suggests a nanocrystalline and disordered structure of CDC. Hence, based on the Raman results, we concluded that the CDC was highly graphitic. However, separate transmission electron microscopy studies at very high resolutions revealed a truly nanocomposite microstructure for CDC. Specifically, these studies have shown that the microstructure is composed of amorphous carbon, turbostratic graphite, carbon nano-onions, and nanocrystalline diamond phases3,* Diamond nanocrystals were found within the transition zone between S i c and CDC in some samples, while the graphitic phase was more dominant in the outer layer of CDC. In general, the structure and composition of the films depend on the synthesis conditions, such as temperature and environment. Detailed mechanical characterization studies by nanoindentation have revealed that the C D C films were rather soft, with typical hardness values ranging from 1 to 1.5 GPa. The elastic modulus of these films was also very low; that is, in the range of 8-12GPa. Such a low hardness and resilience may have been due to the highly graphitic nature of CDC, as revealed by Raman studies. It is important to remember that CDC films are porous at the nanoscale with a pore size of 0.7-

Membranes, Films, and Coatings

Tribological tests were performed in a pin-on-disk machine under a wide range of test conditions. Initially, we ran several tests with base (untreated) S i c in different environments and compared its friction coefficients with those of the as-grown CDC. Figure 6 shows the typical friction coefficients of S i c and CDC during sliding against the Si3N4 balls. The friction coefficient of base S i c is around 0.8 in open air and fluctuates rather heavily, but that of the C D C is around 0.25 and is very steady. These friction values are typical of SiCbased ceramics and pyrolytic graphite in open air. When tests were run in dry nitrogen, the friction coefficient of C D C became much lower, that is, less than 0.1. However, the friction coefficient of bulk S i c increased to 0.9. Figure 7 shows the wear tracks formed on the surface of a C D C film during tests in air and dry nitrogen. As is clear, the size of the wear track formed during testing in open air is much larger. As is apparent from Fig. 3, the physical roughness of as-grown CDC is rather high. The typical R, values of

Fig. 6 Friction coefficients of base S i c and carbide-derived carbon films in open air and d y nitrogen. Thefiiction coefficient of S i c in dty nitrogen was around 0.9.

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Fig. 7. Typical sizp of wear tracksformed on carbide-derived carbon during tests in open air and dry nitrogen.

such films were in the range of 0.3-0.4pm. By progressively polishing the surface of such films, we were able to lower the R, values to less than 0.1 pm. When we repeated some of the tribological tests on such polished CDC films, we obtained as high as 40% reductions in friction (depending on the test environment). Figure 8 shows the effect of mechanical polishing on the frictional behavior of CDC films in open air. It is well known that physically rough surfaces can cause high friction and severe wear losses in most sliding contacts. Specifically, if the sliding surfaces are very rough, a high level of mechanical interlocking can take place between surface asperities, and this condition leads to high frictional losses. A prime example is the inherently rough surface finish of microcrystalline diamond films, which can cause very high friction and severe wear losses during sliding contacts. In fact, recent systematic

Fig. 8. Effect ofpolishing on the friction coeficient of carbidederived carbon in laboratory air.

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studies have demonstrated the existence of an almost linear correlation between surface roughness and friction and wear coefficients of such diamond films. In general, the higher the surface roughness, the greater the friction and wear losses.9 Earlier studies by Erdemir and colleague^'^-'^ have determined that hydrogen plays an important role in the friction and wear behavior of diamond and DLC films. Specifically, DLC films grown in hydrogen-rich plasmas of a plasma-enhanced chemical vapor deposition system exhibited friction coefficients less than 0.01 in dry nitrogen. Based on these findings, we hypothesized that CDC films should also be able to provide very low friction coefficients after a hydrogenation treatment. Through a systematic study, we determined that a postprocess annealing of CDC in 5% H2+95% Ar at 800°C for 8 h was very effective. This hydrogen treatment decreased the IDIIG ratio of Raman spectra of CDC films. This change suggested that the CDC films had become more graphitic during hydrogen treatment, but a considerable disorder still existed. Hydrogen also removes residual chlorine trapped in pores of CDC. Figure 9 summarizes the results of friction tests on CDC films before and after the hydrogenation treatment. As is clear, the friction coefficients of as-grown CDC films are much higher in both air and dry nitrogen than those of the CDC films subjected to a post-hydrogenation treatment. The friction coefficients of the as-grown CDC in air and dry nitrogen are around 0.18 and 0.14, respectively, but after hydrogen treatment the value decreases to 0.07.

Fig. 9. Friction coe8cients of carbide-derived carbon films before and after the hydrogenation treatment.

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Fig. 10. Effect of contact load on thejiction coeficient of carbide-&rived carbon films.

In another series of tests, we investigated the effects of load on the friction of CDC films. Figure 10 shows the frictional behavior of hydrogen-treated C D C under 2 and 10 N loads. Specifically, it shows that under higher loads, the friction coefficient becomes even lower, that is, 0.03. Also note that the initial friction coefficient is lower and the time required to reach the steady-state friction value is shorter under the heavier load. The wear of sliding ball and disk surfaces was very difficult to quantify after wear tests. In all cases, a shallow wear track was formed on the CDC-covered disk side, and its size was strongly dependent on the test environment. The tracks formed in open air wete much wider than those formed in dry nitrogen and had smooth surfaces as shown in Fig. 7. The wear on the Si3N4 ball side was, in most cases, minimal (partially due to the short duration of the sliding tests) and hard to quantify. A typical 3D image of one of the wear scats formed on a Si3N4 ball is shown in Fig. 11 (a). A dark contrasting transfer layer was often observed on some of the sliding ball surfaces, especially after tests in dry nitrogen (see Fig. 11 (b)).

Discussion From the results of the tribological tests presented above, we can conclude that CDC has the ability to lower drastically the friction coefficients of Sic-based ceramics. In separate studies, we obtained similar results from TIC, WC, and other carbide-based ceramics. Such

Membranes, Films, and Coatings

Fig. 11. (a) Three-dimensional (30) and (6) optical images of Si@d ball sufaces afier a sliding test in d y nitrogen. 3 0 image suggests that not much wear had occurred, but the optical photograph reveals the presence of a transfir hyer on and around the contact spot. The width of wear scar shown is 300 pm.

a reduction in friction can be primarily attributed to the low-shear strengths of the phases that make up CDC. Such phases can easily shear during sliding contacts to provide low friction. The CDC films used in our tests were largely made of a graphitic phase. Graphite is a well-known solid lubricant and is able to reduce the friction of sliding tribological interfaces.l 4 At the same time, graphitic carbon in CDC forms a network struc-

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ture that does not contain large planar crystals of graphite that would break under load applied during friction tests. Among the other phases that make up CDC, we feel that nano-onions with their great ability to carry a load and roll over under sliding contact situations may have also played a significant role in lowering the frict i ~ n Nanocrystalline .~ diamond phases are also known to provide low friction when present on a sliding surface.15 Overall, the nanostructured nature of the CDC is a key to the impressive frictional properties shown in Fig. 6. As for the dramatic effects of the test environment on friction, studies to date have confirmed that a very strong relationship exists between the frictional behavior of most carbon films (diamond, DLC, or graphitic carbon films) and the species within the test environmerits. 10,15-19 Some of these studies”-”,” have shown that, if a film is hydrogen free or poor, then very high friction coefficients may result (especially during tests in inert or vacuum environments). It is now generally believed that the surface termination states of these carbon films are extremely important for their frictional behavior. If their sliding surfaces are fully terminated by hydrogen, then these films may provide fairly low friction coefficients. Conversely, if their surfaces are chemically active and contain dangling bonds, then the friction coefficients of these films may become very high. In addition to strong covalent bonding, other types of short- and long-range chemical and/or physical interactions (van der Waals forces, electrostatic and n-n* attractions, capillary forces, etc.) may exist between sliding carbon surfaces and contribute to overall friction.’ ‘-13 We believe that most of the frictional interactions mentioned above are relevant to the behavior of the CDC films tested in our studies. In dry and inert test environments, strong covalent and relatively weak n-n* interactions as well as van der Waals forces may act at the sliding interfaces and control the frictional behavior of CDC films. The n-n* interactions are relevant to the sliding interfaces of graphite or graphitic carbons, such as CDC. Also, the frictional behavior of the CDC films tested in open air could be largely controlled by the graphitic phase, and their friction coefficients (i.e., 0.2) measured in air are typical of graphite. However, when these films are tested in dry nitrogen, their friction coefficients become even lower; this behavior is not the case for pure graphite, which generally provides low friction in moist air. For such an observation, we provide the following explanation.

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As CDC results from selective etching of metal atoms in carbides by C12, some portion of C12 is retained within the growin CDC structure (either in the atomic or bonded form)!’ Also, the small amounts of hydrogen within the carrier gas can act in a similar fashion. Such possibilities may serve the same purpose as water molecules do in a graphitic structure by simply overcoming or reducing the detrimental effects of the “ S O called” n-n* interactions. Another halogen gas, fluorine, is often used in the making of fluorinated graphite, which provides very low friction and possesses friction properties that are relatively insensitive to large fluctuations in relative humidity.21 Also, fluorine doping of DLC films has substantial beneficial effects on the wear behavior of such films.22 We believe that a similar mechanism may have been responsible for the relatively lower friction coefficients of the CDC films in dry nitrogen. Energy-dispersive X-ray analysis of CDC has verified the presence of some Cl in the structure, but the detection of hydrogen was not possible by this method. Nanopores in CDC also allow accumulation of humidity due to capillary forces and termination of dangling bonds by functional groups throughout the coating. Chlorine transport through the growing CDC film has been identified as the rate-controlling step for growth of thick CDC layers on The activation energy for CDC growth is consistent with the chlorine being chemically adsorbed on the surfaces of the nanoscale pores in the CDC. Only a small fraction of the chlorine in the CDC can be removed by an 8-h heat treatment in N2 at 800°C, while a similar treatment in Ar-5% H2 removes 90% of the chlorine. In the presence of hydrogen, the thermodynamic driving force for formation of HCl contributes to the removal of chlorine from C D C , and excess hydrogen is available to replace chlorine atoms on adsorption sites on the CDC surfaces. As for the much lower friction coefficients (the lowest of all measured) of hydrogen-treated CDC films in dry nitrogen, we provide the following mechanistic explanation. Within the mostly amorphous and disordered graphitic microstructure of CDC films (especially near the top), many voids, defects, and active atomic and/or molecular sites may contain dangling or unsatisfied (T bonds. The existence of such bonds at sliding contact interfaces is undesirable (especially in inert and vacuum environments) as they can easily cause adhesive interactions during sliding motion. During the hightemperature hydrogenation, we believe that most of these bonds are effectively eliminated by the hydrogen

Progress in Nanotechnology: Processing

atoms. Such a situation will undoubtedly lead to a higher degree of chemical passivation of the carbon atoms at or near the CDC surface. When such surfaces are rubbed against other materials (Si3N4 in our case), they cause much lower adhesion (hence, friction). Some of the hydrogen may diffuse or migrate well into the structure and remain there as interstitials, while others may eliminate active 6 bonds within the film as well. Remember that C-H bonding is very strong (i.e., stronger than single C-C bonds), and that thermal desorption of hydrogen from carbon does not occur below 700°C.23 In short, elimination of the possibility of strong covalent bond interactions at sliding CDC interfaces is a major reason for the very low friction behavior of CDC films in dry nitrogen. This explanation is consistent with the mechanistic model proposed for the superlubricity of highly hydrogenated DLC

films.’ 1,1324

Conclusions

The results of our study have further confirmed that CDC films are self-lubricating and capable of providing very low friction coefficients to sliding surfaces, especially in dry nitrogen. In humid air, the friction coefficients are somewhat higher but still three to four times lower than the base Sic. High-temperature hydrogenation treatment has a dramatic effect on the frictional behavior of C D C films. It causes significant reductions in friction, especially during tests in dry nitrogen and under heavier loads. The as-grown C D C samples also yielded relatively lower friction in dry nitrogen. In open air, the friction coefficients of both the as-received and hydrogen-treated samples were similar (i.e., around 0.2). Based on these experimental observations, we proposed that the frictional behavior of CDC films in dry nitrogen was largely controlled by their subnanometer porosity and the presence of onion-like structures at sliding interfaces. Hydrogen treatment may have effectively reduced or eliminated dangling bonds, removed chlorine trapped in pores, and hence, provided the very low friction coefficients observed in dry nitrogen. The frictional behavior of these films in laboratory air was largely dominated by the adsorbed water molecules and/ or oxygen-containing hnctional groups covering the surface of CDC.

Membranes, Films, and Coatings

References Y. G. Gogotsi, J.-D. Jeon, and M. J. McNallan, “Carbon Coatings on Silicon Carbide by Reaction with Chlorine-Containing Gases,” J Muter. Chon., 7 [9] 1841-1848 (1997). 2. D. A. Ersoy, M. J. McNallan, and Y. Gogotsi, “Carbon Coatings Produced by High Temperature Chlorination of Silicon Carbide Ceramics,” Muter. Res. Inn., 5[2] 55-62 (2001). 3. Y. Gogotsi, S. Welz, D. A. Ersoy, and M. J. McNallan, “Nature, Conversion of Silicon Carbide to Crystalline Diamond-Structured Carbon at Ambient Pressure,” Nacure, 41 1[6835] 283-287 (2001). 4. S. Welz, Y. Gogotsi, and M. McNallan, “Neucleation, Growth, and Graphitization of Daimond Nanocrysrals During Chlorination of Carbides,” j.Appl. Phys., 93[7] 4207-4214 (2003). 5. D. A. Ersoy, M. J. McNallan, Y. Gogotsi, and A. Erdemir, “Tribological Properties of Carbon Coatings Produced by High Temperature Chlorination of Silicon Carbide,” Tribol. Trans., 43 809-81 5 (2000). 6. M. J. McNallan, D. Ersoy, R Zhu, A. Lee, C. White, S. Welz, Y. Gogotsi, A. Erdemir, and A. Kovalcbenko, “Nano-Structured Carbide Derived Carbon (CDC) Films and their Tribology,” j. Tringhua Sci. Tccbnol., 10 699-703 (2005). 7. 8. Carroll, Y. Gogotsi, A. Kovalchenko, A. Erdemir, and M. J. McNallan, “Effect of Humidity on the Tribological Properties of Carbide-Derived Carbon (CDC) Films on Silicon Carbide,” Tribol. Lett., 15 [l] 51-55 (2003). 8. A. Erdemir, A. Kovalchenko, M.J. McNallan, S. Welz, A. Lee, Y. Gogotsi, and 8. Carroll, “Effects of High-Temperature Hydrogenation Treatment on Sliding Friction and Wear Behavior of Carbide-Derived Carbon Films,” Surf Coat. Tccbnol., 188 588-593 (2004). 9. A. Erdemir, “Wear of Diamond and Diamond-Like Carbon Films,” J Eng. Tn’b~l.,216 387-400 (2002). 10. A. Erdemir and C. Donnet, “Tribology of Diamond, Diamond-Like Carbon and Related Films,” Modcrn Tribolo~Handbook. ed., 8. Bhushan. CRC Press, B o a Raton, FL, 871-899, 2000. 11. A. Erdemir, 0. L. Eryilmaz, and G. Fenske, “Synthesis of Diamond-Like Carbon Films with Superlow Friction and Wear Properties,” j. Vac. Sci. Tecbnol., A18[4] 1987-1992 (2000). 12. A. Erdemir, “Superlubricity and Wearless Sliding in Diamond-like Carbon Films,” M a w . Res. SOC.Symp. Proc., 697 391-403 (2002). 13. A. Erdemir, “The Role of Hydrogen in Tribological Properties of DiamondLike Carbon Films,” Surf Coat. Technof.,146-147 292-297 (2001). 14. A. Erdemir, “Solid Lubricants and Self-Lubricating Films,’’ Handbook of Modern Tribology, CRC Press, B o a Raton, FL, 787-818, 2001. 15. A. Erdemir, G. R. Fenske, A. R. Krauss, D. M. Gruen, T. McCauley, and R. T. Csencsits, “Tribological Properties of Nanocrystalline Diamond Films,” Surf Coat. Technol., 121 565-572 (1999). 16 S. V. Pepper, “Effect of Electronic Structure of the Diamond Surface on the Strength of the Diamond-Metal Interface,” J. Vuc. Sci. Technol., 20[3] 643646 (1982). 17. D. S. Kim, T. E. Fischer, and 8. Gallois, “The Effects of Oxygen and Humidity on Friction and Wear of Diamond-like Carbon Films,” Surf Coat. Tcchnol.. 49 537-542 (1991). 18. M. N. Gardos, Tribiological “Fundamentals of Polycrystalline Diamond Films,” Surf Coat. Techno[., 113[3] 183-200 (1999). 19. C. Donnet, “Recent Progress on the Tribology of Doped Diamond-like and Carbon Alloy Coatings: A Review,” Su$ Coat. Tecbnol., 101 180-186 (1998). 20. A. Lee, “Effects of Processing on Tribological Properties of Carbide Derived Carbon on Silicon Carbide,” Ph. D. Thesis in Materials Engineering, University of Illinois at Chicago, Chicago-Illinois, December, 2005. 21. R. L. Fusaro and H. E. Sliney, “Graphite Fluoride [(CF x)n]-A New Solid Lubricant,” ASLE Tram., 13 56-65 (1970). 22 C. Donnet, J. Fontaine, A. Grill, V. Patel, C. Jahnes, and M. Belin, “WearResistant Fluorinated Diamond-like Carbon Films,” Surf Coat. Technol.,9495 531-536 (1997). 23 C. Su and J. C. Lin, “Thermal Desorption of Hydrogen from the Diamond C(100) Surface,” Surf Sci., 406 149-166 (1998). 24 A. Erdemir, “Design Criteria for Superlubricity in Carbon Films and Related Microstructures,” Tribol. Inr., 37 [7] 577-583 (2004). 1.

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Nanotubes, Nanorods, and Nanowires

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DESIGN, FABRICATION AND ELECTRONIC STRUCTURE OF ORIENTED METAL OXIDE NANOROD-ARRAYS Lionel Vayssieres National Institute for Materials Science, International Center for Young Scientists Tsukuba, Ibaraki, Japan 305-0044

ABSTRACT Materials chemistry has emerged as one of the most consistent fabrication tool for the rational delivery of high purity functional nanomaterials, engineered from molecular to macroscopic scale at low cost and large scale. An overview of the major achievements and latest advances of a recently developed growth concept and low temperature aqueous synthesis method for the fabrication of purpose-built large bandgap metal oxide semiconductor materials and oriented nano-arrays are presented. Important fundamental insights of direct relevance for semiconductor technology, optoelectronics, photovoltaics, and solar hydrogen generation are revealed by in-depth investigations of the electronic structure of metal oxide nanostructures with new morphology and architecture carried out at synchrotron radiation facilities. INTRODUCTION The controlled fabrication of well-defined and well-ordered one-dimensional large bandgap semiconductor nanomaterials such as nanorods, nanowires and nanobelts as well as their large scale manufacturing at low cost remain is of crucial importance to unfold the exciting and promising future of nanodevices. In addition to the rational and economical manufacturing of nanostructures, a better fundamental knowledge of their electronic structure as well as their physical, chemical, electrical and interfacial and structural properties is necessary to hlly reveal and exploit their fascinating potentials. To achieve such ambitious targets, a new approach for the rational development of semiconductor nanostructures has been developed. A strategy based on the chemical and electrostatic lowering of the water-oxide interfacial energy of the systems' allowed a direct thin film growth by heteronucleation onto various substrates (amorphous, poly-, or single-crystalline) of large physical areas. Such approach has the capability to generate functional nanomaterials at large scale low cost and low temperature. For instance, crystalline metal oxide nanostructures consisting of oriented multi-dimensional arrays featuring building blocks of controlled morphologies, sizes, aspect ratios and orientations at nano-, meso-, and micro-scale are genuinely fabricated without template, surfactant, undercoating, or applied external field from the hydro1 sis-condensation of aqueous metal salts and complexes at mild temperatures, ca. below 100°C . In-depth investigations of the electronic structure of such novel nanostructures have been carried out by x-ray spectroscopies at synchrotron radiation facilities. Such studies include x-ray photoelectron3 spectroscopy as well as soft x-ray absorption and emission4 measurements, including polarization dependent' and energy dependent resonant inelastic x-ray scattering6 experiments. The electron structure mapping results reveal important fundamental knowledge of orbital character and symmetry, bandgap, and quantum confinement effects of direct relevance for semiconductor technologies, optoelectronics, gas sensors, photovoltaics and photocatalysis for solar hydrogen generation.

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FABRICATION METHOD Our strategy to generate large-area of advanced nano and micro-particulate thin films at low cost and large scale is a bottom-up aqueous chemical growth approach7 that is wellsupported by a classical thermodynamic model, monitoring the nucleation, growth, and aging processes via the experimental control of the interfacial free energy of the system'. RESULTS AND DISCUSSION Such a strategy has been well-illustrated on the size control of magnetite nanoparti~les~ over an order of magnitude, i.e. between 1 and 10nm. This concept and synthetic method allows the design and the creation of metal oxide nanomaterials with novel morphology, texture, and orientation which enables to probe, tune, and optimize their physical properties". Particulate thin films and 3-D arrays are obtained by direct growth onto various substrates from the condensation of aqueous precursors at low temperatures. Such an approach to material synthesis offers the ability to generate anisotropic nanoparticles and to control their orientation onto substrates. Such an approach has been successfully applied for the growth of advanced nano and microparticulate metal oxide materials such as three-dimensional arrays of ferric oxide nanorods' for photovoltaic12 and phot~catalytic'~ applications, nanocomposite arrays of trivalent iron and chromium corundum oxidesI4 as well as oriented nanorod-I5, microrod-I6 and mi~rotube-'~ arrays of zinc oxide for photovoltaic, optoelectronic devices and gas sensors". Ferroma netic 3-D array of metallic iron nanorods for magnetic devices", 2-D arrays of chromium oxideYo for non-linear optics and ma netoelectrics applications, 3-D oriented arrays of tin dioxide nanorods with square cross-section2 F for gas sensing and photocatalytic applications have also been fabricated by such a method. An example of the variety of metal oxide structures that can be obtained with such an approach is given in figure 1.

'

Figure 1 Field-emission scanning electron microscopy images of crystalline oriented arrays of Sn02 (a); ZnO (b); and a-Fe2O3 (c) grown onto various substrates by aqueous chemical growth

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Figure 2 (a) Polarization-dependent x-ray absorption spectra of ZnO (Wurtzite structure) 3-D arrays consisting of isotropic (spherical) and anisotropic (rod) morphology. The inset illustrates the XAS experimental geometry, where a-, b-axes define the sample surface plane, the c-axis is along the growth direction of the ZnO rods, E is the polarization of incoming photons and B indicates the angle of incidence with respect to the sample surface: 10' (E // c, lines) or 90' (E I c, dots); (b) Calculated polarization dependent 0 1s XAS of wurzite bulk ZnO, where energy is referred to valence-band maximum22;(c) Experimental and calculated x-ray absorption spectra of ZnO rods. The calculations of the XAS spectra include the DOS and the probability of transition to the Opx+y and Opz states5. Synchrotron radiation studies of the electronic structure of such materials have been carried out at beamline 7.01 at the Advanced Light Source (ALS), Lawrence Berkeley National Laboratory. Such beamline is equipped with a 99-pole, 5-cm period undulator and spherical grating monochromator. For instance, the determination of orbital symmetry and orbital character of the conduction band as well as the bandgap of 11-VI semiconductor ZnO (wurtzite structure) has been investigated. Strong polarization effects have been recorded on oriented

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nanorod arrays (fig. 2a) compared to spherical nanoparticle~~. The measured XAS peaks are confirmed by the calculated polarization dependent 0-xy and 0 - z 1sXAS spectra22of bulk ZnO (Fig. 2b, 2c). The XAS has relatively strong in-plane character at energies of around 11, 17, 25 and 29 eV, and out-of-plane character at energies of around 11, 14, 20, and 32 eV. The broad measured XAS between peak a1 and peak a2 is not seen in the calculated spectrum. It may partly be broadening and/or excitonic effects, but may also arise from native donor-like defects. Direct assessment of the bandgap by probing the occupied (by XES) and unoccupied levels (by XAS) of ZnO nanostructures and single crystals has also been carried out4. The XES spectra of bulk and nanostructured ZnO are displayed together with the corresponding XAS spectrum in Fig. 3. The 0 K emission spectrum reflects the 0 2p occupied states (valence band), and the 0 1s absorption spectrum reflects the 0 2p unoccupied states (conduction band). In the photon energy region of 530-539 eV, the x-ray absorption can be mainly assigned to the 0 2p hybridized with Zn 4s states. In the region of 539-550 eV the spectrum is mainly attributed to 0 2p hybridized with Zn 4p states. Above 550 eV, the contribution is mainly coming from 0 2 p Z n 4d mixed states. Stronger s-p-d hybridization was revealed in nanostructured ZnO since the contributions of features at 520 eV and 523 eV are enhanced. A well defined band gap can be observed between the valence-band maximum and conduction-band minimum. Our absorption-emission spectrum yields the fundamental band-gap energy of 3.3 eV, which is in agreement with the 3.4 eV found for bulk ZnO.

Figure 3 Oxygen x-ray absorption-emission spectrum reflected conduction band and valence band near the Fermi level of ZnO nanoparticles in comparison with bulk Zn04. The electronic structure of quantum rod arrays of hematite has also been investigated6. Their Fe 2p absorption spectrum is displayed in figure 4a. The spectral shape appears very similar to previous XAS measurements conducted on polycrystalline or single-crystal samples. The typical spectrum shows the spin-orbit interaction of the 2p core level that splits the L2 (2~112)and L3 (2~312)edges, and the p-d and d-d Coulomb and exchange interactions that cause multiplets within the edges. The ligand field splitting of 3d transition metals, being of the same order of magnitude as p-d and d-d interactions (1-2 eV), gives a 1.4-eV-energy splitting between the tZg (xy,yz, xz) and eg (x2-y2, 3zz-r2)orbitals. Figure 4b shows the Fe L-emission spectrum recorded with a higher photon-energy excitation (ca 750 eV). The spectral shape shows two peaks originating from the transitions of 3d orbitals to 2 ~ 1 1 2and 2~312core levels. A branching ratio (LdLa) of 0.8 is found for the a-Fe2O3 bundled nanorod arrays, which appears substantially larger than that of the single crystal. The RIXS spectrum recorded at the Fe L-edge of a-Fe203

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nanorods is shown in figure 4c. Several energy-loss features are clearly resolved. The low energy excitations, such as the strong d-d and charge-transfer excitations, are identified in the region from 1 to 5 eV. The 1-eV energy-loss features originate from multiple excitation transitions. The 4.1 and 6.4 eV excitation originates from charge transfer between oxygen 2p and iron 3d orbitals. The 2.5-eV excitation, which corresponds to the bandgap transition of hematite, appears significantly blue shifted compared to the reported 1.9-2.2 eV bandgap of single-crystal and polycrystalline samples as suggested above by the higher LB/La branching ratio observed in the L-emission spectrum (figure 3b) of the nanorods. Such direct observation of a substantial (0.30.6 eV) bandgap increase is successfully attributed to a 1-D (lateral dimension) quantum confinement effect in the designed bundled ultrafine nanorods of hematite (fig. 4d). Such a finding strongly suggests that these designed quantum rod-arrays would meet the bandgap and band edge requirements for solar hydrogen23generation without an applied bias. Further studies aiming at developing novel low-cost visible-active semiconductor photocatalysts are currently under investigation in our laboratories.

Figure 4 Fe 2p absorption (a), Fe L emission (b) and energy-dependent resonant inelastic x-ray scattering spectra (c) of a-Fe203 quantum rod-arrays (d). The inset (c) shows a schematic representation of the radiative de-excitation process for the two core excitatiom6

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CONCLUSION The ability to hierarchically design and fabricate, at low cost, large three-dimensional arrays of transition metal oxides consisting of anisotropic nanoparticles with controlled orientations and aspect ratios onto various substrates without the use of templates, surfactants, or applied external fields contributes not only to the development of smart and functional nanomaterials for optoelectronic, photoelectrochemical, gas sensors and magnetic devices, but also to a better fundamental understanding of orbital character and symmetry contributions, bandgap and quantum confinement effects of one-dimensional anisotropic nanostructures by indepth studies of their electronic structure by soft x-ray spectroscopies at synchrotron radiation facilities. REFERENCES 1

L. Vayssieres, “On the thermodynamic stability of metal oxide nanoparticles in aqueous solutions”, h t . J. Nanotechnol. 2005, 2(4), 41 1-439 2 L.Vayssieres, “On the design of advanced metal oxide nanomaterials”, Int. J. Nanofechnol. 2004, I(I-2), 1-41 3 A. Henningsson, A. Stashans, A. Sandell, H. Rensmo, S. Sodergren, L.Vayssieres, A. Hagfeldt, S. Lunell, H. Siegbahn, “Proton insertion in Polycrystalline W03 studied with electron spectroscopy and semi-empirical calculations”, in Advances in Quantum Chemistv, edited by E. J. Brandas and E. Brandas (Academic Press, 2004), Vol. 47, pp. 23-36 4 C. L. Dong, C. Persson, L. Vayssieres, A. Augustsson, T, Schmitt, M. Mattesini, R. Ahuja, C. L. Chang, J.-H. Guo “The electronic structure of nanostructured ZnO from x-ray absorption and emission spectroscopy and the local density approximation”, Phys. Rev. B 2004, 70(19), 195325 ’J.-H. Guo, L.Vayssieres, C. Persson, R. Ahuja, B. Johansson, J. Nordgren, “Polarizationdependent soft-x-ray absorption of highly oriented ZnO microrods” J. Phys. : Condens. Matter 2002,14(28), 6969-6974 6 L.Vayssieres, C. Sathe, S. M. Butorin, D. K. Shuh, J. Nordgren, J.-H. Guo, “One-dimensional quantum-confinement effect in a-Fe2O3 ultrafine nanorod arrays”, Adv. Mater. 2005, I 7(19), 2320-2323 7 L.Vayssieres, “Designing ordered nano-arrays from aqueous solutions”, Pure Appl. Chem. 2006, 78(9) 1745-1751 ‘L. Vayssieres, “PrCcipitation en milieu aqueux de nanoparticules d’oxydes: ModClisation de l’interface et contr8le de la croissance”, PhD. Thesis, UniversitC Pierre et Marie Curie, Paris, France (1995), pp. 1-145 ’L.Vayssieres, C. Chaneac, E. Tronc, J.P. Jolivet, “Size tailoring of magnetite particles formed by aqueous precipitation: An example of thermodynamic stability of nanometric oxide particles”, J. Colloid Interface Sci. 1998,205(2), 205-212 “L.Vayssieres, “Advanced semiconductor nanostructures”, C. R. Chimie 2006,9(5-6), 69 1-701 “L.Vayssieres, N. Beermann, S.-E. Lindquist, A. Hagfeldt, “Controlled aqueous chemical growth of oriented three-dimensional nanorod Arrays: Application to iron(II1) oxides”, Chem. Mater. 2001, I3(2), 233-235 12 N. Beermann, L.Vayssieres, S.-E. Lindquist, A. Hagfeldt, “Photoelectrochemical studies of oriented nanorod thin films of Hematite”, J. Electrochem. SOC.2000, I47(7), 2456-246 1

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T. Lindgren, H. Wang, N. Beermann, L.Vayssieres, A. Hagfeldt, S.-E. Lindquist, “Aqueous photoelectrochemistry of hematite nanorod-array”, Sol. Energy Muter. Sol. Cells 2002, 71(2), 23 1-243 14 L.Vayssieres, J.-H. Guo, J. Nordgren, “Aqueous chemical growth of aFe203-aCr203 nanocomposite thin films”, J. Nunosci. Nunotechnol. 200 1,1(4,, 385-388 15 L.Vayssieres, “Growth of arrayed nanorods and nanowires of ZnO from aqueous solutions”, Adv. Muter. 2003,15(5), 464-466 16 L.Vayssieres, K. Keis, S.-E.Lindquist, A. Hagfeldt, “Purpose-built anisotropic metal oxide material: 3D highly oriented microrod-array of ZnO”, J. Phys. Chem. B 2001, 105(17), 33503352 17 L.Vayssieres, K. Keis, A. Hagfeldt, S.-E. Lindquist, “Three-dimensional array of highly oriented crystalline ZnO microtubes”, Chem. Mater. 2001, 13(12), 4395-4398 18 J. X. Wang, X. W. Sun, Y. Yi, H. Huang, Y. C. Lee, 0. K. Tan, and L.Vayssieres, “Hydrothermally grown ZnO nanorod arrays for gas sensing applications”, Nanotechnology 2006,17(19), 4995-4998 ”L.Vayssieres, L. Rabenberg, A. Manthiram, “Aqueous chemical route to ferromagnetic 3D arrays of iron nanorods”, Nuno Lett. 2002,2(12), 1393-1395 20L.Vayssieres,A. Manthiram, “2-D mesoparticulate arrays of a-Cr203”, J. Phys. Chem. B 2003, 107(12), 2623-2625 21 L.Vayssieres, M. Graetzel, “Highly ordered Sn02 nanorod-arrays from controlled aqueous rowth”, Angew. Chem. Int. Ed. 2004, 43(28), 3666-3670 F2 C. Persson, C.L. Dong, L. Vayssieres, A. Augustsson, T, Schmitt, M. Mattesini, R. Ahuja, J. Nordgren, C.L. Chang, A. Ferreira da Silva, J.-H. Guo “x-ray absorption and emission spectroscopy of ZnO nanoparticles and highly oriented ZnO microrod arrays”, Microelectronics J. 2006,37(8), 686-689 23Solur Hydrogen and Nanotechnology, L.Vayssieres (Editor), Proceedings of SPIE-The International Societyfor Optical Engineering, SPIE Press Ltd., 2006, Vol. 6340, 336 pages

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ELECTROSPINNING OF ALUMINA NANOFIBERS Karin Lindqvist, Elis Carlstrom IVF Industrial Research and Development Swedish Ceramic Institute Argongatan 30 SE-43153 Molndal, Sweden Anna Nelvig, Bengt Hagstrom IFP Research Argongatan 30 SE-431 53 Molndal, Sweden ABSTRACT A straightforward manufacturing route of alumina nanofibers was established by the electrospinning of a commercial alumina sol mixed with polyethylene oxide. The molecular weight and concentration of the polymer were varied as well as the spin parameters in order to find suitable fabrication settings. The results showed that the polymer content should be at least 27 vol% in the green fibre in order to avoid brittle fibres and the molecular weight of PEO should be at least 400 000 g/mol. The shear and elongation rheological properties of the suspensions were measured. A contraction flow method for measuring the elongational viscosity was evaluated and found to be suitable for highly shear thinning systems with high viscosity. The fibre diameters of calcined samples were in the range 400-700 nm. The electrospinning should be performed in a dry environment facilitating better control of the fibre deposition. INTRODUCTION Electrospinning is a convenient and simple process to obtain ultra thin ceramic fibres. In its simplest form a ceramic precursor suspension is placed in a syringe and subjected to a high electric potential visa-vi a grounded collector electrode. A schematic drawing of the process is shown in figure 1. When a voltage is applied to the syringe the pending spherical drop at the needle tip is changed into a conical shape (Taylor cone) and, upon a further increase in applied voltage (producing an increased charge density in the suspension), the electrostatic repulsive forces will overcome the surface tension of the suspension and a fine charged jet is ejected from the Taylor cone towards the grounded collector. During its flight the jet is highly elongated and the suspension medium is evaporating leaving a fibrous non-woven mat on the collector. The “green” fibres so produced are calcined into ceramic fibres at elevated temperature. Powder dispersions, alkoxide solutions or colloidal sols can be mixed with polymers to create the ceramic precursor suspension. During the last few years a number of different ceramics have been spun into sub-micron fibres by the electrospinning technique[13. The rheological properties of the precursor suspension turn out to be important for the electrospinning process and for the formation of uniform fibres. Since the electrostatic forces available are comparatively weak the counter acting viscous forces have to be limited in order to create thin fibres. This will set an upper bound on viscosity. At lower viscosities surface tension will become a dominant factor and capillary break up of the liquid thread into droplets (Rayleigh instability) may produce electro spraying instead of electrospinning. To this end it is common practise to use a soluble polymer in the precursor suspension to regulate its viscoelastic

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properties. Several authors [2-41 have emphasised the importance of the elongational viscosity for the formation of uniform fibres. It was shown that the Rayleigh instability could be completely suppressed if a critical value of the elastic stress in the jet was exceeded. Yu [4] used an extensional rheometer measuring the filament break-up of the sample. Stading [5] developed a measurement system for shear thinning fluids that measured the elongational viscosity in contraction flow. This method gives the possibility to control the strain rate during the measurement, which is an advantage compared to measurements based on filament break-up. The parameters controlling the final diameter of electrospun fibres have been the subject of extensive research. Sigmund [ l ] modified the theory for prediction of the final jet radius to account for higher bulk electrical conductivity in ceramic-polymer suspensions. According to the theory the final green fibre diameter is dependent on the surface tension, the volume flow rate of the suspension, the volume current through the jet, the electrical conductivity and the electric field strength. A few papers report on the fabrication of electrospun aluminium oxide fibers. Azad [6] made transparent alumina fibers from aluminium pentadionate mixed with PVP in an acetone-ethanol solution. Larsen et al. [7] reported on the fabrication of alumina fibers from aged sol prepared from aluminium di-sec-butoxide ethylacetoacetat. The objective of this study was to establish a robust straightforward process for electrospinning of alumina fibers with controlled fibre morphology. There are commercially available ceramic sols commonly used as high refractory binders. These materials have high purity, small particle size, are delivered as stable colloidal water based dispersions and could therefore be suitable as raw material for electrospinning. No papers have to our knowledge been published on manufacturing of ceramic nanofibers from commercially available sols. Several water-soluble polymers could be candidates for facilitating the formation of ceramic fibers. Polyethylene oxide (PEO) is commonly used in electrospinning and was chosen in this study. The molecular weight and concentration of the polymer have been varied as well as the spin parameters in order to find suitable fabrication settings. The shear and elongation rheological properties of the suspensions have been measured. A contraction flow method for elongational viscosity was evaluated to see if this method was suitable for electrospinning suspensions. The fibre diameters were measured on calcined samples. EXPERIMENTS A commercially available colloidal alumina sol, Nyacol AL20 (Nyacol Nan0 Technologies Inc., USA) with a primary particle size of 50 nm was used as the ceramic precursor. Polyethylene oxide with molecular weights (Mw) of 100,000 (look), 400,000 (400k) and 900,000 (900k) (Sigma Aldrich, Sweden) were dissolved in the sol by magnetic stirring so that a suspension of polymer and sol was obtained. The conductivity (CDM210, Radiometer Analytical, France) was measured on the alumina sol as received and with the polymer dissolved. The suspensions were rheologically characterized by measuring the steady state shear viscosity and the phase shift in oscillatory shearing in a cup and bob (CC25) measurement system (StressTech, Rheologica, Sweden). The elongational viscosity was estimated through the measurement of the pressure drop over a specially designed contraction flow nozzle (Reologen i Lund, Sweden) that produces a constant elongational strain rate along its centre line, see figure 2. The measurement of elongational viscosity requires a rather strong shear thinning behaviour of the fluid in order to subtract the contribution from shear flow components to the measured pressure drop.

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Electrospinning was performed in the experimental set up schematically shown in figure 1. Initial experiments were performed in ambient atmosphere with the electrospinning parameters set to 20 kV (high voltage power supply ES5OP from Gamma High Voltage, Ormon Beach, FL, USA), a syringe-collector distance of 20 cm; a syringe of 0.8 mm and a flow rate of 0.056 ml/min. A grounded platinum foil was used as collector electrode. Electrospinning for fibre diameter measurements were performed in a Plexiglas box under flowing nitrogen gas with parameters set to obtain a single well defined Taylor cone at the tip of the syringe needle. The parameters for each system are shown in table 1. The fibres were calcined at 700°C in air. XRD was used to identify the alumina phases obtained after calcinations. The specific surface area was measured (Gemini 2300, Micromeretics) on green and calcined fibres. Samples for fibre diameter measurements were prepared by sedimentation of calcined and crushed fibres on a sample holder. The fibres were manually measured from SEM micrographs.

Figure 1. Schematic of the electrospinning Figure 2. Measuring system for contraction process flow measurement. Table 1. Parameters for electrospinning of fibers for diameter measurements. Sample Composition Voltage Collector distance Field strength (kV) (cm) (kV/cm) A 3wt%400k 10 15 0.67 B 6wt%400k 20 15 1.33 C 3wt%900k 10 10 1 D 3wt%900k 10 10 1

Flow rate (mumin) 0.05 0.05 0.005 0.01

RESULTS The results from the shear viscosity measurements and frequency sweep measurements are shown in figures 3 and 4.

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Figure 3. Shear viscosity vs. shear rate for polyethylene oxide dissolved in alumina sol.

Figure 4. Phase angle vs frequency for PEO dissolved in alumina sol Figure 3 shows that the suspension with 5 wt% PEO lOOk was Newtonian and had a low shear viscosity. No fibres were obtained with this sample but spherical particles from electrospraying, see figure 5. Severe dripping during handling also made this suspension unsuitable for electrospinning and no further analyses were made. The suspensions with PEO 400k and PEO 900k had higher shear viscosities and were shear thinning. Suspensions with higher concentrations of PEO 400k or 900k could not be dissolved properly in the sol due to very high viscosities. Oscillatory measurements in shear were performed to obtain the viscoelastic properties and the results are shown in figure 4. A phase shift of 90 degrees between strain and stress means that the material is completely viscous. The lower the phase shifts the more elastic

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is the suspension. Although the shear viscosity is moderate for the sample with 3wt% PEO 900k the elasticity is high at high frequencies indicating the importance of molecular weight for this property. Preliminary electrospinning experiments were performed in ambient environment without controlling humidity. The fibre formation process was erratic and fibres were deposited not only on the collector electrode but also on surrounding surfaces. The fibres were in fact “standing up” forming a 3-dimensional cotton like morphology in the applied electric field and fell back (more or less) on the collector when the voltage was turned off. Turning on the voltage again resulted in the “erection” of the fibres. It seems as the fibres were polarised in the field. The formation of 3-dimensional cotton like structure has so far not been recorded when pure polymer solutions are electrospun to fibres. However, Azad [8] obtained this type of behaviour when processing zirconia and ceria fibres in ambient atmosphere. Samples with 2wt% polymer in the alumina sol had a lot of broken fibres when studied in SEM, see figure 6. The sample was very brittle, which is believed to be due to the low polymer content (20 vol% in the green fibre). An addition of 3 wt% polymer resulted in more continuous fibres and no broken fibres were found in samples with 6wt % polymer added, see figure 7.

Figure 5. 5 wt% PEO 100k.Spherical particles from electrospraying, green sample.

Figure 6. 2wt% PEO 900k. Green broken Fig 7.6wt% PEO 400k green fibres. fibres due to low polymer content.

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10000 lu

5

W/o

400k -3 900k -+3 400k t .6

U

a

2

1000

b 0

<

.% 0

MW

100

0

v) .-

> -

lu

0 .+

10

lu

cn

K 0 rl

1 1

10

100

1000

Extension rate (s-1)

Figure 8. Elongational viscosity vs. extension rate for PEO alumina sol suspensions.

Figure 9. Results from XRD measurements on calcined fibres of 6wt%PEO 400k pressed into a thin disc shape. The elongational viscosity and fibre diameters were measured on the systems shown in table 1. Figure 8 shows the results from the measurements of elongational viscosity with the results corrected for the shear viscosity that has an influence on the registered force during the measurements. Only one extension rate could be accurately measured for the system with 3 wt% PEO 400k due to the low shear viscosity and the limited shear thinning of that system. There is a much larger difference between the behaviour of the systems in extension than in shear. For a Newtonian fluid the elongational viscosity is three times (Trouton's ratio) the shear viscosity. In

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this case, when the molecular weight of PEO is sufficiently high, the elongational viscosity is 50100 times larger than the shear viscosity and it can be assumed that the suspensions are strain hardening in elongation. Such behaviour will stabilise the filament and prevent capillary break UP.

Very little information about the actual elongational strain rates during electrospinning is available in the literature. Yu measured the jet thinning of aqueous PEO solutions from the Taylor cone up to the point where the whipping instability sets in. Based on the flow rate and the diameter progression along the spin line the strain rate can be measured in this region. Data from Yu indicate strain rates in order of magnitude around 10 s-'. Unfortunately, the strain rates in the whipping instability region are not known since it is difficult to assess the point along the fibre where stretching ceases. . The elongational rheometer used in this study cover a wide range of strain rates around 10 s'l so it seems reasonable to assume that it provides viscosity data in the relevant strain rate regime. Figure 9 shows that the main phase obtained when calcining to 700°C was y alumina. There is no amorphous phase in the sample but the broad peaks show that the grain size is very small, about 30 nm. The fibre diameters of calcined samples obtained from measurements on SEM micrographs are given in Table 2. A frequency diagram is shown in figure 10. All of the systems had uniform fibres without branches when the electrospinning parameters were set to give a single Taylor cone, which is required to get proper whipping of the jet. Comparing system C and D, having the same composition and electric field strength shows that an increased flow rate resulted in larger fibre diameter. Figure 11 shows a SEM micrograph of system D with 3 wt% PEO 900k.

Figure 10. Calcined fibres diameter frequency diagram

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Figure 11. 3 wt% PEO 900k calcined fibres electrospun in Plexiglas box under flowing nitrogen. Table 2. Results from measurements of conductivity, fibre diameter and specific surface area. Sample Composition Conductivity Average Standard Confidence Specific surface (mS/cm) fibre deviation interval area (m2/g) (95%) diameter (nm) (nm) Alumina sol 2.2 3wt% lOOk 2.0 A 3wt%400k 2.0 42 1 f 11 166.9 112 B 6wt%400k 1.8 557 f 47 165.6 330 33.3 (green) f 8 C 3wt%900k 1.9 526 97 D 3wt%900k 1.9 710 f 16 158.5 160 The value of the conductivity was independent of the molecular weight of the PEO and seems to be influenced mainly by the concentration of the sol in the sample. The specific surface area of the green sample of 6wt% PE0400k was only 33.3 m2/g while the calcined samples had much higher values. The large specific surface area of the calcined samples is mainly due to the porous crystal structure of the y - alumina. The surface of the alumina is “hidden” by the polymer before calcination of the sample.

DISCUSSION The choice of a commercial alumina sol as a source for the ceramic phase of electrospun nanofibres resulted in a straightforward preparation of the spinning suspension. The choice of polymer, molecular weight and concentration is crucial for the formation of uniform fibres. The initial electrospinning in this study showed that suspensions with PEOlOOk did not form continuous fibres but resulted in spherical particles due to the Rayleigh instability. Higher molecular weight or concentration of the polymer resulted in uniform fibres. However, too low polymer content gave very brittle fibres and green fibres broke during handling. It can therefore be concluded that a minimum amount of polymer, in this study about 27 vol%, is necessary to obtain continuous and usable fibres.

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The fixed initial spinning conditions often resulted in several simultaneous Taylor cones at the tip of the syringe needle. Spinning was also discontinuous in the sense that the cones appeared and disappeared in a stochastic fashion and drops were accumulated and released periodically at the needle tip. It was soon realised that in order to have one single active Taylor cone and a continuous spinning without dripping it was necessary to tune the spinning conditions for each suspension. To avoid multiple jets and dripping, as compared to the initial settings, it was necessary to decrease the flow rate from the syringe pump andor increase the field strength by increasing the voltage or decreasing the distance to the collector. The parameter combinations in Table 1 resulted in a single well defined Taylor cone and continuous spinning without dripping. There is a processing window for each suspension allowing a stable and well defined spinning. It was noted that the ambient humidity greatly affected the deposition of fibres. By using nitrogen blanket the relative humidity was brought close to zero and the fibres were then directed and deposited on the collector. Humidity may affect the surface resistivity of surrounding objects made from otherwise insulating materials like wood, paper and plastics via adsorption and it can be speculated that a low humidity will concentrate the electrical field lines from the charged syringe towards the grounded collector plate. However, the deposition was still 3-dimensional in the sense that a “cotton like” structure was built up from the collector surface, see figure 12.

Figure 12. Morphological feature of green alumina fibres electrospun under flowing nitrogen. This type of highly porous structure may be interesting for certain applications like high temperature filtration. This type of structure is not seen when pure polymer solutions are electrospun. In this case a 2-dimensional non-woven mat is formed on the collector electrode. This type of 2-dimensional structure was also recorded when alumina particles were dispersed in aqueous PEO solutions and electrospun. According to the theory [ 11 of predicting the green fibre diameter an increased flow rate and /or a decreased field will increase the fibre diameter. The bulk conductivity and the surface tension of the suspension will also have an influence on the fibre diameter. In the systems studied the conductivity is almost constant and the differences in surface tension will

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also be very small. The fibre diameters as influenced by the electrospinning parameters were not studied in a systematic manner. However, some observations regarding the possible effect of elongational viscosity can still be made. Comparing system A (3wt% PEO 400k) and B (6wt% PEO 400k) with the same flow rate but B having twice the field strength one would expect B to show a smaller diameter than A if the elongational viscosity had no effect. In fact, B shows thicker fibres. The higher elongational viscosity due to the higher polymer concentration in B might explain this. A similar argument is found by comparing A and C (3wt%PEO 900k). C having higher field strength and significantly lower flow rate than A, both should work in the direction of thinner fibres, still C shows thicker fibres. It should be noted that the fibre diameters were measured on calcined samples obtained by sedimentation on a sample holder. This sample preparation is not possible to use for green fibres that would dissolve during the sedimentation. However, the shrinkage of the calcined fibres will be the same for all suspensions containing 3wt%PEO regardless of the molecular weight. For the system containing 6wt% the shrinkage will be larger; that is the green fibres were coarser than measured on the calcined samples. This further stresses the influence of the elongational viscosity as stated above. The rheological measurements showed that there is a much larger difference in the extensional behaviour between the systems than in the shear behaviour. Contraction flow measurements are suitable for highly shear thinning systems with a rather high viscosity. This method would have been more suitable if the polymer suspension contained polymer of higher molecular weights or higher concentrations. Extensional rheology measurements based on filament break up is probably more suitable and also possible to use on solutions with polymers of low molecular weight and concentration. CONCLUSIONS We have shown that a straightforward manufacturing of alumina nanofibers is possible by the electrospinning of a commercial alumina sol mixed with polyethylene oxide. The polymer content should be at least 27 vol% and then have a molecular weight of at least 400 000. The concentration range should be 3-6 wt% of polymer in the sol. The electrospinning should be performed in a dry environment facilitating better control of the fibre deposition. The manufactured fibres had diameters in the range 400-700 nm and further optimisation can reduce these values. REFERENCES Sigmund, W., et al., Processing and Structure Relationships in Electrospinning of 1. Ceramic Fiber Systems. J. Am. Ceram. SOC., 89(2), 395-407.(2006). Daga, V., Helgeson, M., and Wagner, N., Electrospinning of Neat and Laponite-Filled 2. Aqueous Poly(ethy1ene oxide) Solutions. Journal of Polymer Science: Part B: Polymer Physics, 44, 1608- 16 17.(2006). Shenoy, S.L., et al., Role of chain entanglements on fiber formation during 3. electrospinning of polymer solutions: good solvent, non-specific polymer-polymer interaction limit., Polymer, 46, 3372-3384.(2005). Yu, J.H., Fridrikh, S.V., and Rutledge, G.C., The role of elasticity in the formation of 4. electrospun fibers. Polymer, 47,4789-4797.(2006). Stading, M. and Bohlin, L. Contraction flow for characterization of extensional properties 5. of semi-solid materials. in XZVth Znternational Congress on Rheology. 2004. Seoul, Korea: The Korean Society of Rheology.

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6. 7.

8.

Azad, A.-M., Fabrication of transparent alumina (A1203) nanofibers by electrospinning. Materials Science and Engineering: A, 435-436,468-473.(2006). Larsen, G., et al., A method for making Inorganic and Hybrid (Organic/Inorganic) Fibers and vesicles with Diameters in the Submicrometer and Micrometer range via Sol-gel Chemistry and Electrically Forced Liquid Jets. J. Am. Chem: Soc., 125( 11541155).(2003). Azad, A.-M., Matthews, T., and Swary, J., Processing and characterization of electrospun Y203-stabilized Zr02 (YSZ) and Gd203 -doped CeOa (GDC) nanofibers. Materials Science and Engineering: B, 123,252-258.(2005).

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ZnO Nanofiber and Nanoparticle Synthesized Through Electrospinning and Their Photocatalytic Activity Under Visible Light Haiqing Liu,'3$,5Jinxia Yang,$ Jianhe Liang,$ Yingxing Huang,$ and Chunyi TangS kollege of Chemistry and Materials Science, Fujian Normal University, Fuzhou 350007, China $Key Laboratory of Polymer Materials of Fujian Province, Fuzhou 350007, China

In this paper, we fabricate ZnO nanofibers and nanoparticles through electrospinning precursor solution zinc acetate(ZnAc)/ cellulose acetate(CA) in mixed-solvent N,N-dimethylformamidelacetone. Depending on the posttreatment of precursor ZnAc/CA composite nanofibers, both ZnO nanofibers and nanoparticles were synthesized after calcination of precursor nanofibers. The morphology and crystal structure of the ZnO nanofiber and nanoparticle were characterized by scanning electron microscopy, transmission electron microscopy, atomic force microscopy, and X-ray diffraction. It was found that the mean diameter of the ZnO nanofiber and nanoparticle was ca. 78 and 30 nm, respectively. The photodegradation of dye molecules such as Rhodamine B and acid fuchsin catalyzed by the ZnO nanofiber and nanoparticle was evaluated under the irradiation of visible light. Both morphological ZnO species showed strong photocatalytic activity. However, the ZnO nanofiber in the form of nanofibrous mats showed much higher efficiency than the nanoparticle although the latter has a smaller size than the former. The porous structure of ZnO nanofibrous mats is believed to improve the contacting surface areas between the catalyst and the dye molecules, while the aggregation of ZnO nanoparticle in the solution lowers the photocatalytic efficiency. I. Introduction

T

HE ZnO is an important

semiconductor material with a band gap of 3.37 eV and an excitation binding energy of 60 meV. Owing to its nontoxicity and strong oxidizing power to organic pollutants under visible light irradiation, ZnO has been applied as a photocatalyst in environmental cleaning.' However, it is well known that the photocatalytic activity of ZnO strongly correlates with the size of the ZnO specimen. That is, the larger the size of ZnO the higher the recombination rate of photoinduced electron-hole pairs at or near its surface, resulting in lower photocatalytic activity or vice versa. Owing to quantum confinement of nanostructured material, the separation rate of photo-induced charge carriers in semiconductor photocatalysts can be significantly increased by reducing its size to the nanometer range, and therefore enhance its quantum efficiency and photocatalytic activity. ZnO in the form of a nanoparticle exhibits strong photocatalytic oxidation to organic pollutants; but it is realized that the nanoparticle photocatalyst may repollute the treated water due to difficult recovery. Moreover, the nanoparticle aggre ation during aging may lower the photocatalytic efficiency. To overcome the drawbacks of ZnO nanoparticles, ZnO nanowires and nanoplatelets supported on an aluminum film have been synthesized. It is reported that they

B

exhibit a high surface-to-volume ratio and are very stable against aggregation, and they show higher efficiency than conventional nanoparticle in degrading organic pollutants."6 However, the yields of a nano-ZnO film from custom methods are relatively low and cannot meet the needs of photocatalytic application on a large scale. Therefore, new strategies have to be developed to prepare large quantities of ZnO films with high surface areas. It is well documented that a fibrous membrane composed of nanofibers exhibits large surface areas and porous which makes it a very promising morphology for photocatalyst. Among many methods such as hydrothermal and vapor- hase transport used for the preparation of ZnO nanofibers,bl1° electrospinning is the only method able to yield abundant quantities of continuous nanofibers for practical applications. The electrospinning technique has been applied to prepare a one-dimensional nanofiber with diameter down to tens of nanometers and length up to several centimeters, and the nanofibers are often collected as nanofibrous mats.' Although e-spinning has been used mainly for the fabrication of organic polymeric nanofibers, ceramic metal oxide nanofibers have also received considerable interest because of their potential applications in areas including photovoltaic devices, piezoelectric material, gas sensors, and solar cells.12 Metal oxide nanofibers are usually prepared by electrospinning a precursor metal salt solution with the help of a proper polymer, followed by calcination to decompose the polymer completely and turn metal salt into metal ~ x i d e . ' ~So . ' ~far, the only precursor system reported for the e-spun ZnO nanofiber is zinc acetate (ZnAc)/poly(vinylalcohol) in water. The diameter of a ZnO nanofiber as obtained is over 100 nm." Up to now, the photocatalytic activity of ZnO nanofiber has not been examined experimentally. In this work, we demonstrate the preparation of a ZnO nanofiber from a novel spinning solution system, i.e., N,N-dimethylformamide (DMF)/ acetone as a solvent, ZnAc as a precursor, and cellulose acetate (CA) as a fiber template. From this solution system, ZnO nanoparticles smaller than 40 nm in diameter are obtained by direct calcination of ZnAc/CA composite nanofibers, whereas ZnO nanofibers with diameter down to 23 nm are synthesized by calcination of Zn(OH)2/cellulosecomposite nanofibers, which is obtained from the hydrolysis of ZnAc/CA composite nanofibers in base solution. Additionally, for the first time, the photocatalytic activity of ZnO nanofibers is evaluated and compared with that of ZnO nanoparticles by measuring the photo-degradation of some dye molecules under visible light irradiation.

11. Experimental Procedure

J. Ninwontrihuting editor

(I)

Manuscript No. 23749. Received September 15,2M)7; approved December 9, 2007. The authors are grateful for the financial support by the Initiative Fund for the Returned Overseas Chinese Scholar administered by the State Education Ministry, and the Key Project of Natural Science Foundation of Fujian Province (Grant No. E0620001). 'Author to whom correspondence should be addressed. e-mail: haiqing.liu@gmdi~.com

Nanotubes, Nanorods, and Nanowires

Materials Zinc acetate dihydrate (Sinopharm Chemical Reagents Co., Shanghai, China) was recrystallized in distilled water. CA with a degree of substitution (DS) of 2.45 and an M , of 3.0 x lo4 was Obtained from Eastman. All Other chemicals were as received. 269

(2) Preparation of ZnO Nanofiber Transparent spinning solutions containing 6 or 10 wt% ZnAc were prepared by adding ZnAc to 20 wt% CA in a 2.1 (v/v) DMF/acetone solvent mixture, followed by magnetic stirring at ambient temperature. They were placed in a syringe with an 18-gauge stainless spinneret. The feeding rate was 5 pL/min, monitored by a syringe pump (TS2-60, Longer Precision Pump Co. Ltd., Baoding, China). An electrode was clamped on the spinneret and connected to a power supply (DW-P303-IAC, Tianjin Dongwen High Voltage Plant, Tianjin, China). A grounded counter electrode was connected to a collector aluminum foil, which was placed 10 cm away from the orifice. The electric field was maintained at 10 kV. The ZnAc/CA composite nanofibrous membranes on the grounded collector were dried in a vacuum oven at 90°C for 5 h. One part of the ZnAc/CA composite nanofibrous membrane was calcinated at 500°C in air for 5 h. The other part of the membrane was hydrolyzed in 0.1 N NaOH aqueous solution for 24 h at ambient temperature to transform the ZnAc/CA composite nanofibers into Zn(OH)2/cellulose nanofibers, and then washed with distilled water completely. The hydrolyzed membrane was vacuum dried at 50°C for 5 h. Subsequently, it was calcinated in the same way as the original ZnAc/CA composite nanofibrous membrane.

(3) Characterization The viscosity of the spinning solution was measured by a rotational viscometer (NDJ- 1, Shanghai Precision & Scientific Instrument Co. Ltd., Shanghai, China) at 30°C. Fourier transform infrared (IT-IR) spectra were collected on a Nicolet Avatar 360 (Nicolet, Pittsfield, MA) spectrometer in KBr form. The structure of the ZnO nanofiber was characterized by an X-ray diffractometer (Xpert MPD Pro, Philips, the Netherlands) with CuKu radiation. The diameter and morphology of the ZnO nanofiber and the nanoparticle were observed by scanning electron microscopy (SEM) (JSM-6300LV, JEOL), transmission electron microscopy (TEM) (JEM2010, JEOL), and atomic force microscopy (AFM) (NanoScope Ma).

( 4 ) Photocatalytic Activity of ZnO Nanofiber A lab-made photochemical reactor was set up according to reference.I6 A 500 W tungsten lamp was used as the visible light source. A light filter was placed between the lamp and the dye solution to only allow visible light (h>420 nm) to pass through. Five milligrams of ZnO nanofiber or nanoparticle was added to a cylindrical glass vessel containing 10 g of 45 ppm acid fuchsin, Amido black 10B, or 6 ppm Rhodamine B (RhB) aqueous solution, respectively. At a certain time interval, the degradation reaction was stopped and the solution was centrifuged. The dye concentration in the supernatant was measured on a UV-vis spectrometer (Lambda 850, PerkinElmer, Waltham, MA). Parallel degradation reactions under the same conditions were conducted for varied time intervals. Only one measurement was made for each batch. Two control experiments were also conducted to test: ( I ) the adsorption of dye on the ZnO nanofiber in darkness and (2) photodegradation of the dye when exposed to visible light without the presence of catalyst ZnO.

111.

Results and Discussion

( 1 ) Morphology and Structure of ZnO Nanofiber and Nanopwticle Figure 1 shows the SEM of ZnAc/CA composite nanofibers. Composite nanofibers generated from 6 wt% ZnAc solution are generally smooth and round, with fiber diameter in a narrow range of 115-185 nm (Fig. I(a)), while that from 10 wt% ZnAc solution are mixtures of beads and nanofibers with a large size variation (Fig. I(b)). The large nanofibers are 27&380 nm in diameter with a regular round morphology; whereas the size of the whisker-like nanofibers is <20 nm. Obviously, the diameter and size distribution of ZnAc/CA composite nanofibers increased with the concentration of ZnAc, while the fiber template CA concentration was kept constant at 20 wt%. We also found that the spinnability of the ZnAc/CA solution containing 10 wt% ZnAc became difficult. It has been reported that a higher polymer solution viscosity often results in a larger fiber diameter and size distribution in many electrospinning As listed in Table I, the viscosity of the spinning solution

Fig. 1. Scanning electron microscopy images of Zn(OAc)2/celluloseacetate (CA) composite nanofibers containing (a) 6 wt%, (b) 10 wt% Zn(OAc)2; (c) nanoparticles from calcination of (a) in air at 500°C for 5 h; (d) transmission electron microscopy image of (c); (e) atomic force microscopy of (c).

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Table I. Viscosity of Zn(OAc)2/CA Solutions Spinning solution

Concentration of Zn(0Ac)Z (wt%)

6

1 2

10

3

15

Concentration of CA (wt%)

( x 10) cp)

Viscosity

20 20 20

1.86 2.50 3.10

increased with the concentration of ZnAc. The molecular structure of cellulose and cellulose derivatives often contains residual carboxyl groups, which are introduced during the manufacture and posttreatment processes.19Therefore, the addition of ZnAc in CA solution would allow a chelating reaction between cation Zn2+ and carboxyl groug-COO- on CA to take place, resulting in a higher CA chain entanglement and an increase of solution viscosity. Surprisingly, the calcination of ZnAc/CA composite nanofibers at 500°C in air generated nanoparticles instead of nanofibers (Fig. I(c)). The size of grain-like ZnO nanoparticles is <40 nm as shown in the TEM image (Fig. I(d)). The AFM image also shows that ZnO nanoparticles are in the range of 2838 nm in diameter and 3CL130 nm in length in the longitudinal direction (Fig. l(e)). It is known that CA is a thermoplastic with a T, and a Td (decomposition temperature) of ca. 230" and 340"C,lX respectively. Therefore, the CA component in the ZnAcjCA composite nanofibers melts upon heating in an oven, resulting in transformation from fiber to molten CA. With a further increase in the oven temperature, CA is decomposed and ZnO nanoparticles are left behind. However, it is reported that when PVA is applied as a polymer fiber template to fabricate a metal oxide nanofiber, a metal nanofiber could be formed by direct calcination^.'^'^ It is believed that the T,,, and T, of PVA are very close, causing no apparent melting of PVA upon heating. In order to fabricate a ZnO nanofiber from a ZnAc/CA composite nanofiber, we firstly hydrolyzed the composite nanofiber to transform component CA into cellulose. In the hydrolysis process of ZnAc/CA, the acetyl group on the CA nanofiber is readily removed in a mild basic aqueous solution to regenerate cellulose nanofiber." Meanwhile, component ZnAc in the composite nanofiber can react with anion OH- to form precipitate Zn(OH)2. Compared with the morphology of the ZnAc/CA

composite nanofiber (Fig. I(a)), the surfaces of the Zn(OH)*/ cellulose nanofiber became coarse. Fibers were irregular and distorted, and intertwined at many loci. Some fibers merged together to form larger fibers (Fig. 2(a)). The merging and intertwining among fibers became drastic for composite nanofibers with a higher ZnAc content (Fig. 2(d)). Noticeable morphology change was hardly observed in the hydrolysis of single component CA nanofiber.18 Hence, the transformation of ZnAc into zinc hydroxide played an important role in the morphology change of the Zn(OH)2/cellulose nanofiber. The flocculent Zn(OH)2 precipitates merge and penetrate readily among the adjacent nanofibers, resulting in intertwining among Zn(OH)2/ cellulose nanofibers. Different from the calcination of the ZnAc/ CA nanofiber, free-standing ZnO mats (an image from the digital camera, Fig. 2(b)) composed of ultrafine nanofibers (Fig. 2(c)) were formed after calcination of the Zn(OH)2/cellulose nanofiber. Because of the thermal decomposition of cellulose component and dehydration of the Zn(OH), component, the fiber size decreased significantly to as small as 24 nm in diameter shown in its TEM image (inset, Fig. 2(c)). With an increase in the concentration of precursor ZnAc from 6% to 10% in the spinning solution, the mean diameter of the ZnO nanofiber increased from 72 nm to ca. 100 nm (Fig. 2(e)). It should be pointed out that the ZnO nanofiber morphology could be maintained even after several minutes of ultrasonication as shown in its TEM image (Fig. 2(c)), suggesting that it is hard to disintegrat into nanoparticles. This is really important in its application as a high-efficiency photocatalyst as will be discussed in the following section. The many bright spots appearing mostly at the end of the ZnO nanofibers (Figs. 2(c) and (e)) are not caused by impurities; instead, they resemble ZnO nanospheres formed by the fusing of nanofibers that are not completely deacetylated, similar to the nanoparticle formation in the calcination of ZnAc/ CA nanofibers as shown in Fig. 1. FT-IR spectra (not shown) confirm the chemical structure change that occurred during the course of hydrolysis and subsequent calcination of the ZnAc/CA composite nanofiber. After hydrolysis, the band at 1750 cm-' (vcx0) for ZnAc/CA disappeared in the spectrum of the Zn(OH)2/cellulosenanofiber. The FTIR spectrum of ZnO nanofibers showed its characteristic peak at 440 ern-'.'' The XRD pattern of ZnO nanofibers is shown in Fig. 3. It shows the characteristic peaks of a hexagonal

Fig. 2. Scanning electron microscopy (SEM) images of (a) Zn(OH)z/cellulosenanofibers obtained from hydrolysis of 6 wt% Zn(OAc)z/celluloseacetate (CA) in aq. NaOH solution, (b) ZnO nanofibrous mat after calcination of (a) in air at 500°C for 5h. This photo was taken with a digital camera. (c) SEM of ZnO nanofibers in (b); the inset shows transmission electron microscopy of (c); (d) Zn(OH)z/cellulose nanofibers obtained from hydrolysis of 10 wt% Zn(OAc)*/CA in an aq. NaOH solution, (e) ZnO nanofibers after calcination of (d).

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30

35

40

45 28 (degree)

50

55

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Fig. 3. X-ray diffraction pattern of ZnO nanofibers.

crystalline ZnO phase at 20 = 31.7" (loo), 34.4" (002), 36.2" (IOI), 47.5" (102), and 56.6" (1 lo), while no other peaks for impurities appear. This suggests that a pure ZnO nanofiber was obtained after the calcination of Zn(OH)2/cellulose composite nanofibers.

Fig.4.

( 2 ) Photocatalytic Activity of ZnO Nanojiher and Nanopavticle The photocatalytic oxidation of dyes acid fuchsin and amido black 10B by ZnO nanofiber under visible light was evaluated and shown in Fig. 4. After photocatalytic degradation for 30 min, the pink color of the acid fuchsin solution changed to very light pink. The UV spectra showed that nearly 95% acid fuchsin was degraded. In the same time interval, the dark blue color of amido black 10B turned into light blue due to the decomposition of almost 85% of amido black IOB. Control experiments showed that the photolysis of these two dyes is unnoticeable without the addition of a ZnO nanofiber under the radiation of visible light in 30 min. The ZnO nanofiber absorbs <5% dye in darkness. The results suggest that the ZnO nanofiber has strong photocatalytic degradation activity toward acid fuchsin and amido black 10B. We also conducted an experiment to compare the photocatalytic activity of the ZnO nanofiber and the nanoparticle. It was found that the degradation rate of acid fuchsin was 85% and 73% in 20 min for the ZnO nanofiber and nanoparticle, respectively, indicating that the photocatalytic activity of the ZnO nanoparticle is less than that of the ZnO nanofiber, although the former has a smaller size as mentioned above. RhB has often been used as a model dye molecule for hotocatalytic degradation by a transition metal oxide! The photocatalytic degradation of RhB by the ZnO nanofiber and nanoparticle is shown in Fig. 5. Catalyzed by the ZnO

UV-vis spectra (left) and appearance (right) of dye solutions after photo-degradationcatalyzed by ZnO nanofibers. (a) Acid fuchsin, (b) amido

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efficiency than easy-to-aggregate ZnO nanoparticles in degrading some dye molecules. ZnO nanofibrous mats degraded nearly 60% RhB in 2 h under visible light irradiation, whereas only 28% RhB was decomposed by the ZnO nanoparticle under the same conditions. ZnO nanofibrous mats are highly efficient photocatalysts and can be potentially used in the treatment of waste water.

P

References .3 .

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0

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Fig. 5. ZnO nanofiber- and nanoparticle-catalyzed photodegradation of Rhodamine B in an aqueous solution.

nanofiber, nearly 60% of the initial RhB was degraded after 2 h of visible light irradiation. Almost 90% RhB was consumed after the degradation time was extended to 19 h. In contrast, only about 28% and 50% RhB was decomposed on the ZnO nanoparticle in 2 and 19 h, respectively. Thus, the ZnO nanofiber showed much stronger photocatalytic activity toward RhB than the ZnO nanoparticle. Although the mean diameter of ZnO nanofibers is larger than the ZnO nanoparticle, the porous structure of ZnO nanofibrous mats (Figs. 2(b) and (c)) allows more ZnO surface areas to make contact with RhB molecules, whereas the aggregation of ZnO nanoparticle makes the available areas to RhB molecules less. This results in higher photocatalytic efficiency of ZnO nanofibrous mats. Many researches have also demonstrated that ZnO nanopalates and porous nanoparticle films possess much stronger photocatalytic activity than analogous nanoparticles.s.2’ From the almost linear relationship between (Ao-A)/Ao and I , a first-order degradation mechanism was indicated for both the ZnO nanofiber and the nanoparticle in the degradation period tested in this research. This is the same as the degradation mechanism for the other nano-ZnO morphologies in photocatalyzed degradation of organics.5s2’In the degradation period of 2-19 h, the degradation rate was 0.017 and 0.013 h-’ for the ZnO nanofiber and nanoparticle, respectively. The final products after photodegradation were not determined in this study, but it has been reported that COl and H 2 0 are obtained for the degradation of RhB b semiconductor metal oxides such as nano-Ti02 and Zn0?*lx

IV.

Conclusions

ZnO nanofibers and nanoparticles were synthesized by the electrospinning technique. The mean diameter of both nanostructures is < 100 nm. Both morphologies possess photocatalytic activity under visible light. Because of the porous structure of ZnO nanofibrous mats, it shows much higher photocatalytic

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’J. L. Yang, S. J. An, W. 1. Park, G. C. Yi, and W. Choi, “Photocatalysis Using ZnO Thin Films and Nanoneedles Grown by MetalLOrganic Chemical Vapor Deposition,” Adv. Mater.. 16, 16614 (2004). ’Y. J. Jang, C. Simer, and T. Ohm. ”Comparison of Zinc Oxide Nanoparticles and its Nano-Crystalline Particles on the Photocatalytic Degradation of Methylene Blue,” Maier. Res. BUN., 41, 67-77 (2006). ’Z. Y. Liu, D. D. Sun, P. Guo, and J. 0. Leckie, “An Efficient Bicomponent TiOz/SnOz Nanofiber Photocatalyst Fabricated by Electrospinning with a SideBy-Side Dual Spinneret Method,” Nano Lett., 7, 1081-5 (2007). 4B. Rohe. W. S. Veeman, and M. Tausch, “Synthesis and Photocatalytic Activity of Silane-Coated and UV-Modified Nanoscale Zinc Oxide,” Nanotechnology, 17, 277-82 (2006). ’C. H. Ye, Y. Bando, G. Z. Shen, and D. Golberg, “Thickness-Dependent Photocatalytic Performance of ZnO Nanoplatelets,” J. Phys. Chenr. B, 110. 15146-51 (2006). ~, 6F. Xu, Z. Y. Yuan, G. H. Du, T. Z. Ren, C. Bouvy, M. Halasa, and B. L. Su, “Simple Approach to Highly Oriented ZnO Nanowire Arrays: Large-Scale Growth, Photoluminescence and Photocatalytic Properties,” Nanotechnology, 17, 588-94 (2006). ’D. H. Reneker and I. Chun, “Nanometre Diameter Fibres of Polymer, Produced by Electrospinning,” Nanotechnology, 7, 21623 (1996). ‘C. Patel, S. Li, C. Wang, W. Zhang, and Y. Wei, “Electrospinning of Porous Silica Nanofibers Containing Silver Nanoparticles for Catalytic Applications,” Chem. Mater., 19, 1231-8 (2007). 9X. Xu, X. W. Sun, and B. J. Chen, “Field Emission from Gallium-Doped Zinc Oxide Nanofiber Array,” Appl. Phys. Leit.. 84. 154&2 (2004). ‘9. P. Fang, Q. Pang, X. G. Wen, J. R. Wang. and S. H. Yang, “Synthesis of Ultrathin ZnO Nanofibers Aligned on a Zinc Substrate,” SmaN, 2, 612-5 (2006). “C. Y. Tang, S. H. Ye, and H. Q. Liu, “Electrospinning of Poly(StyreneCo-Maleic Anhydride) (SMA) and Water Swelling Behavior of Crosslinked/ Hydrolyzed SMA Hydrogel Nanofibers,” Polymer, 48, 4482-9 I (2007). I’D. Li and Y. N. Xia, “Electrospinning of Nanofibers: Reinventing the Wheel?,” A h . Maier., 16, 1151-70 (2004). “W. Sigmund, J. Yuh. H. Park, V. Maneeratana, G. Pyrgiotakis. A. Daga, J. Taylor, and J. C. Nino, “Processing and Structure Relationships in Electrospinning of Ceramic Fiber Systems,” J . Am. Ceram. Soc., 89, 3 9 5 4 7 (2006). I4D. Li and Y. N. Xia, “Fabrication ofTitania Nanofibers by Electrospinning,” Nan0 Leu., 3, 5 5 5 6 0 (2003). I5H. Wu and W. Pan, “Preparation of Zinc Oxide Nanofibers by Electrospinning,” J . Am. Ceram. Soc., 89, 699-701 (2006). I6F. Peng, H. J. Wang, H. Yu, and S. H. Chen, “Preparation ofAluminum FoilSupported Nano-Sized ZnO Thin Films and its Photocatalytic Degradation to Phenol Under Visible Light Irradiation,” Muter. Res. BUN., 41, 2123-9 (2006). ”C. S. Ki, D. H. Baek. K. D. Gang, K. H. Lee, I. C. Um, and Y. H. Park, “Characterization of Gelatin Nanofiber Prepared from Gelatin-Formic Acid Solution,” Polymer, 46, 5094102 (2005). “H. Q . Liu and Y. L. Hsieh, “Ultrafine Fibrous Cellulose Membranes from Electrospinning of Cellulose Acetate,” J. Polym. Sci. PI. B-Polym. Phys., 40, 21 t$29 (2002). H. Q.Liu, S. K. Obendorf, M. 1. Leonard, T. J. Young, and M. J. Incorvia, “Adsorption of Aroma Chemicals on Cotton Fabric from Aqueous Systems,” J . Sur/act. Deterg., 8, 311-7 (2005). 20 H. Wu, R. Zhang, X. Liu, D. Lin, and W. Pan, ”Electrospinning of Fe, Co, and Ni Nanofibers: Synthesis, Assembly, and Magnetic Properties.” Chem. Mater., 19, 35061 I (2007). ”B. Pal and M. Sharon, “Enhanced Photocatalytic Activity of Highly Porous ZnO Thin Films Prepared by S o l C e l Process,” Maler. Chem. Phys., 76, 82-7 (2002). 0 ~~

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SYNTHESIS OF CARBON NANOTUBES AND SILICON CARBIDE NANOFIBERS AS COMPOSITE REINFORCING MATERIALS Hao Li Department of Mechanical and Aerospace Engineering University of Missouri at Columbia Columbia, MO, 6521 1 Abhishek Kothari, and Brian W. Sheldon Division of Engineering Brown University Providence, RI 02906 ABSTRACT The excellent mechanical properties of nanomaterials are driving research into the creation of strong and tough nanocomposite systems and new forms of nanomaterials. It is critical to select an appropriate reinforcing material with desired microstructures and properties to achieve better nanocomposite performance. The present study focused on synthesis and processing-microstructure relationships of multiwalled carbon nanotubes (CNTs) and Sic nanofibers with chemical vapor deposition (CVD). Various CNTs grown by CVD with anodic aluminum oxide (AAO) templates were examined with scanning and transmission electron microscope (SEM and TEM). It was demonstrated that the experimental conditions, especially catalysts and plasma, have significant impact on CNT growth rates and microstructures. Both catalyst and plasma can increase the deposition rate about one order of magnitude. In addition, catalysts promote the secondary growth of CNTs inside the primary CNTs and plasma may improve the stiffness of primary and secondary CNTs. The Sic nanofibers grown by CVD with catalysts were also investigated. Sic diameters match well with the diameters of precursor CNTs, indicating Sic nanofiber size is controlled by the catalysts originated from CNTs. Generally, CNTs fabricated with CVD-template method have disordered graphitic structures and thus have lower tensile strength, but the disordered structures may help for the load transfer between graphitic layers. In the present study, CVD-SIC nanofibers appear to be stiffer than multiwalled CNTs and may also serve as a good candidate for composite reinforcing materials. INTRODUCTION The high modulus and strength of nanomaterials, such as carbon nanotubes (CNTs) and carbon nanofibers (CNFs), has sparked tremendous interest in nanomaterial reinforced composites and development of novel materials for composite applications. Numerous efforts worldwide are addressing all aspects of this rapidly developing field, including synthesis, processing, characterization and integration.'-2 Reinforcing materials, as one of the three major components of the composite system, strongly affect composite properties, such as strength, toughness, conductivity, and thermal stability. The synthesis and selection of reinforcing materials is thus critical to composite design and manufacturing. CNTs and CNFs are the most widely investigated reinforcing materials for ceramic, metal, and polymer based nanocomposites, while other materials, such as clay, have also been intensively investigated as a reinforcing material.

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A good understanding of the composite system is required for the design and synthesis of reinforcing nanomaterials considering the objectives of the inclusion of nanomaterials vary for different systems. In a polymeric nanocomposite, the application of CNTs is expected to increase its strength for nanotube super strength and stiffness, while improving or at least keeping its toughness for nanotube flexibility. The extremely large interfacial area distinguishes nanocomposites from traditional microfiber reinforced composites. The impact of the interface over a few R , (radius of gyration, in the order of tens of nanometers) in polymer matrix is substantiated by extensive investigations on the behavior of thin polymer films, as well as empirical observations in traditional composites and filled rubber^.^ It is expected that the inclusion of nanomaterials will significantly affect the properties of the polymer matrix as well as the polymeric nanocomposites. In ceramic matrix nanocomposites, the motivation of adding nanomaterials is related to the toughness or the crack growth resistance of ceramics considering ceramics are already stiff and strong. Beyond mechanical property concern, there are other objectives of inclusion of nanomaterials in polymer or ceramic matrix, such as increase of thermal conductivity and electronic conductivity. A better understanding of the goal of the nanocomposite and the synthesis-microstructure-property relationships of nanomaterials is of great importance for the success of development of nanocomposites with desired properties. S i c microfibers are one of the most well studied non-oxide fibers for ceramic matrix composite, especially in harsh environments and at elevated temperature^.^ Sic nanofibers are expected to have superior thermal stability to CNTs in oxidizing environments. It was reported that Sic nanofibers with 20-30 nm size have tensile strength of over SOGPa, which is about half of the theoretical strength of CNT (IOOGPa).' It appears that S i c nanofibers may be a good candidate as reinforcing materials for various applications. The present study will focus on the synthesis and microstructure characterization of various CNTs and Sic nanofibers fabricated with CVD. It is hoped that the information collected here will help for the design of the composites and selection of reinforcing nanomaterials. EXPERIMENT Synthesis of carbon nanotubes Microwave plasma-enhanced chemical vapor deposition (MPCVD; AsTex 15009 was used to grow CNTs. The typical growth experiments were conducted with 100 sccm C2H2 and 400 sccm NH3 at about 10 torr. The deposition temperature was varied from 800 to 900 "C and deposition time varied from 5 minutes to 14 hours. Anodic aluminum oxide (AAO) templates (Whatman Inc.) were used for the growth of CNTs. The inner diameter of the pore of the AAO template is about 200 11111. Catalysts were also used for some experiments. In order to prepare catalyst nanoparticles, AAO templates were first immersed in 0.1 M or 1 M Ni(NO3)2 solution and then put in the CVD chamber and treated with 100 sccm H2 at 800 "C for about 2 hours. When plasma was employed, the input power was controlled at 1000 W. Synthesis of S i c nanofibers A hotwall CVD was used to grow S i c nanofibers. CNTs (Helix Material Solutions, Inc.) with diameter from 30-100 nm were used as the carbon source. A mixture of Si and Si02 powder with 2:1 molar ratio was put under the CNTs in a high-purity alumina crucible to generate volatile Si species, which reacted with the CNTs to form Sic nanofibers. The reaction temperature was controlled at approximately 1400 "C. The carrier gas composition was 5% H2 in Ar.

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Characterization The morphology of the CNTs and SIC nanofibers were examined with SEM (LEO 230) and more detailed information, such as atomic structures, was obtained by TEM (JEOL 20 10). All the CNTs fabricated in the present study were embedded in AAO templates. For all the TEM samples and some of the SEM samples, the AAO templates were etched off with KOH solutions at room temperature. No conducting coating was used for any SEM image.

Figure 1. SEM and TEM images of a CNT sample fabricated for 12 hours at 900°C. (a and b) and a CNT sample fabricated for 14 hours at 800 "C (c and d). (a) CNTs in templates, (b) TEM images, (c) and (d) Templates etched off. RESULTS AND DISCUSSION Catalyst Free CNTs (No Plasma) Fabrication of catalyst-free CNTs with such CVD-template method at lower temperatures were studied by Martin's group and Xu's group, but their primary research interests and efforts are not on their mechanical properties.6, The present study will lay emphasis on the correlation of microstructures and mechanical properties. Figure 1 depicts the SEM and TEM images of two samples grown at relatively high temperatures for 12 hours and 14 hours. Both samples turned black after deposition, but not much microstructural difference could be observed before and after deposition based on SEM images (Figure la). The deposition rate is very low, only -0.5nm per hour at 900 "C. CNT graphitic layers are not well ordered, but roughly parallel to the tube axis (Figure Ib). Figure Ic and Id shows the SEM images of CNT samples when AAO

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templates were etched off. Most of CNTs formed CNT bundles after matrix was removed. It is also observed that tubular structure of some CNTs collapsed. The SEM and TEM characterization results demonstrate that the CNTs fabricated with the CVD-template methods do not have perfect tubular structures. In a recent study of AAO matrix nanocomposite reinforced by CNTs fabricated by a similar method at 600 "C, the CNT strength is estimated to be about 15-25 GPa, significantly lower than the theoretical strength (-1OOGPa). The toughness of the AAO templates increase from 0.4 to 5 MPa m"'2 with this type of CNTs based on nanoindentation tests." Although the nanocomposite mechanical properties are not considered excellent, the significant toughness increase is very encouraging. In addition, toughening mechanisms, such as tube pull-out and crack bridging, were observed in such nanocomposites.8 The lower tensile strength estimated from the nanoindentation test seems to be attributed to disordered structure of CNT (Figure lb). A perfect multiwalled CNT may have higher tensile strength, but the very weak van der Waals forces between the tubes will significantly reduce the load transfer. Such phenomenon has been observed experimentally by Cumings and Zettl." In this case, only the outermost layer of multiwalled CNTs may strengthen the matrix. Recent work by Waters et al suggested that that disorder of this type can create more interactions between the graphitic layers in the multiwalled CNTS." While imperfect structures like this may reduce the axial modulus of the CNTs, it should also increase the interlayer shear strength. The effects of these changes on composite properties have not yet been studied, but this type of bonding may be beneficial for load transfer and could thus improve nanocomposite mechanical properties. This suggests a possible tradeoff between multiwalled CNT strength and inter-tube load transfer, such that there may be an optimal level of disorder for multiwalled CNT reinforced nanocomposites. Heat treatments were used to improve the degree the order of graphitic layer, so tailoring such order is possible experimentally. However, further investigation of these effects is required for a better understanding. Catalyst Assisted Growth of CNTs (No Plasma) Ni nanoparticles, reduced from Ni(N03)2, were uniformly deposited on the inner wall of the template pores. In the sample shown in Figure 2, 1 M Ni(NO3)2 was used for catalyst preparation. CNT wall thickness of this sample is in the range of 30-50 nm. It appears that the Ni catalysts significantly increased the deposition rates. The wall thickness is a little smaller when O.1M Ni(NO3)2 was used for catalyst preparation (images not shown here), but the thickness is still significantly larger than that of catalyst-free samples. More interestingly, secondary CNTs inside the bigger CNT were observed as shown in Figure 2c and illustrated in Figure 2d. The secondary CNTs are curled and tangled inside primary CNTs. Martin et al reported growth of CNT with metal catalyst nanoparticles at lower temperatures around 500-550 "C. At the lower temperature range, the growth of secondary CNTs was impeded and a hollow CNT could be fabricated without forming the secondary CNTS.~ The secondary CNTs was apparently promoted by a catalyst assisted vapor-liquid-solid (VLS) mechanism. A typical VLS process starts with the dissolution of gaseous reactants into nanosized liquid droplets on a catalyst metal, followed by nucleation and growth of nanotube or nanowire. The onedimensional growth is induced and dictated by the liquid droplets, whose size remains unchanged in most cases. 12, 13 We speculate that no liquid droplet formed at low temperatures, such as under 550 "C, so that no secondary CNTs could grow via the VLS mechanism. In this case, temperature acts as a switch that can turn on and off the secondary growth of CNTs. Figure 2b shows fracture surface of the same sample. No debonding between the primary CNTs and the

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AAO matrix is observed, suggestive of strong interfacial bonding. However the bonding between the primary CNTs and secondary CNTs is relatively weak. Future research will focus on detailed CNT microstructure characterization by TEM and nanocomposites mechanical testing by nanoindentation.

Figure 2. SEM images (a, b, and c) and schematic (d) of a CNT sample grown with Ni nanoparticles catalyst for 1 hour at 800 "C. (a) templates etched off and (b) templates fractured. Catalyst and Plasma Assisted Growth of CNTs When both catalysts and plasma was used, extremely high deposition rates were achieved. It is believed the high energy electrons in the plasma dissociate gases and generate atomic and ionized species and thus significantly increase the growth rates. CNT wall thickness reached about 30-50 nm in 5 minutes in the sample shown in Figure 3. CNTs (Figure 3a) did not fall or collapse after templates were etched off. It appears that the CNTs fabricated with plasma are stiffer than those fabricated without plasma. Secondary CNTs were also found in these types of CNTs, but they are less curled compared to those shown in Figure 2 and many of them popped out after AAO matrix was fractured. Both the primary and secondary CNTs are expected to have

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better structures as indicated by the above observations. It is believed that the high energy electrons contributed to better structures, such as high degree of ordering, of these CNTs. It will be of great value to further investigate the influence of plasma on the structures of CNTs at atomic and meso scales. The secondary CNT popout indicates they are under residual stress inside the primary CNTs and push primary CNTs against the AAO matrix. It is conceivable that the interfacial strength between the primary CNTs and matrix may be improved by the secondary CNTs. Further research will be required to investigate such effects.

Figure 3. SEM images of CNT grown with catalyst and plasma for 5 minutes at 700 "C. (a) templates etched off and (b) templates fractured.

Figure 4.S i c nanofibers grown for 7 hours at 1400 "C. Sic nanofibers Fabricated in a Hotwall CVD Figure 4 shows a typical SEM image of Sic nanofibers. The Sic nanofibers grown in a hotwall CVD are straight and appear to be much stiffer than all the CNTs in the present study.

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Many small beads could be easily identified at the end of S i c nanofibers, suggestive of vaporliquid-solid (VLS) growth mechanism. The metal catalysts originated from the original CNT precursor, which may contain 2% of metal nanoparticles or less. The diameters of these Sic nanofibers, ranging from 30-100 nm, match very well with those of CNTs. This is reasonable considering that the sizes of both are controlled by the same nanoparticles. Direct conversion of CNT to Sic is not observed based on SEM and TEM results. The lower end strength of the Sic nanofibers is about 24GPa (based on diameter of 100 nm) and nanofibers with smaller sizes are expected to have higher ~trength.’~ The average strength of those Sic nanofibers is certainly higher than the CVD catalyst-free multiwalled CNTs fabricated with the plasma off. The TEM characterization (not shown here) also reveals a 1-5 nm think amorphous coating on the surface of Sic nanofibers. Considering Sic microfibers were usually coated with a coating before added in matrix, this amorphous coating may also serve as an interphase to tailor the bonding strength between the nanofiber and the matrix of interest. CONCLUSIONS It is demonstrated that nanomaterials with different microstructures and different nanomaterials could be fabricated for the composite applications with CVD methods. Both the microstructure and the mechanical properties of CNTs could be tailored by modifying the experimental conditions. The deposition rates of catalyst-free CNTs without plasma are extremely low and it may not be practical to use such of CNTs as reinforcing materials despite the improvement of composite toughness. Catalysts and plasma can significantly speed up the deposition rates and plasma may also improve the stiffness of CNTs. Both catalyst and plasma may be desired for fabrication of CNT reinforce AAO nanocomposites. S i c nanofibers with diameters of 30-100 nm were also fabricated with a hotwall CVD. These nanofibers may serve as a good candidate for composite reinforcing materials considering their excellent mechanical properties and superior thermal and chemical resistance. ACKNOWLEDGEMENTS This work was supported by the National Science Foundation through a Nanoscale Interdisciplinary Research Team (NIRT) Grant, CMS-0304246. REFERENCES ‘W. A. Curtin and B. W. Sheldon, “CNT-reinforced Ceramics and Metals,” Materials Today, 7,44-49 (2004). *H. D. Wagner and R. A. Vaia, “Nanocomposites: Issues at the Interface,” Materials Today, 7,38-42 (2004). 3 F. Gao, “Clay/Polymer Composites: The Story,” Materials Today, 7,50-55 (2004). 4R.Naslain, “Design, Preparation and Properties of Non-Oxide CMCs for Application in Engines and Nuclear Reactors: An Overview,” Compos. Sci. Technol., 64, 155-170, (2004). 5 E. W. Wong, P. E. Sheehan, and C. M. Lieber, “Nanobeam Mechanics: Elasticity, Strength, and Toughness of Nanorods and Nanotubes,” Science, 227, 1971-75, (1997). 6 G. Che, B. B. Lakshmi, C. R. Martin, E. R. Fisher, and R. A. Ruoff, “Chemical Vapor Deposition Based Synthesis of Carbon Nanotubes and Nanofibers Using a Template Method,” Chem. Mater., 10,260-67, (1998).

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7J Li, C. Papadopoulos, J. Xu, M. Moskovits, “Large-Area Hexagonally Ordered Arrays of Carbon Nanotubes,” Appl. Phys. Lett., 75,367-69, (1999). 8 Z. Xia, L. Riester, W. A. Curtin, H. Li, B. W. Sheldon, J. Liang, B. Chang, and J. Xu, “Direct Observation of Toughening Mechanisms in Carbon Nanotube Ceramic Matrix Composites,” Actu Muter., 52,93 1-944, (2004). 9Z. Xia, W. A. Curtin, B. W. Sheldon, “Fracture toughness of highly ordered carbon nanotube/alumina nanocomposites, ” Transactions of the ASME, 126, 238-244, (2004). 10 J. Cumings, and A. Zettl., “Low-Friction Nanoscale Linear Bearing Realized From Multiwalled Carbon Nanotubes,” Science, 289, 602-604, (2000). 11 J. F. Waters, P. R. Guduru,a- M. Jouzi, J. M. Xu, T. Hanlon and S. Suresh, “Shell Buckling of Individual Multiwalleded Carbon Nanotubes Using Nanoindentation,” Appl. Phys. Lett., 87, 103109, (2005). 12 R. S. Wagner, W.C. Ellis, “Vapor-liquid-solid mechanism of single crystal growth,” Appl. Ph s. Lett. 4, 89-90 (1964). M. Law, J. Goldberger, and P. Yang “Semiconductor nanowires and nanotubes,” Annu. Rev. Muter. Sci. 34, 83-122 (2004). 14 W. Yang, H. Araki, C. Tang, S. Thaveethavorn, A. Kohyama, H. Suzuki, and T. Noda, “Single-Crystal Sic Nanowires with a Thin Carbon Coating for Stronger and Tougher Ceramic Composites,” Adv. Muter., 17, 1519-23, (2005).

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POLYMER FIBER ASSISTED PROCESSING OF CERAMIC OXIDE NAN0 AND SUBMICRON FIBERS Satyajit Shukla, Erik Brinley, Hyoung J. Cho, and Sudipta Seal Surface Engineering and Nanofabrication Facility (SNF) Advanced Materials Processing and Analysis Center (AMPAC) and Mechanical Materials Aerospace Engineering (MMAE) Department Engineering 38 1 4000 Central Florida Blvd. Orlando, Florida 328 16 Phone: 407-882-1 189 Fax: 407-882-1462 E-mail(s): [email protected], [email protected] ABSTRACT Nan0 and Submicron ceramic oxide fibers of tin oxide (Sn02) have been processed using the polymer fibers as templates. Highly porous fibrous mat of hydroxypropyl cellulose (HPC) polymer (molecular weight 80000 g/mol) has been obtained via electrospinning technique. In order to derive the ceramic oxide fibers using the polymer fibers as templates, the electrospinning characteristics of the HPC polymer has been established by varying the critical processing parameters such as polymer concentration, solvent type, tip-to-electrode distance, feeding speed, and applied voltage. Under selected processing conditions, the Sn-precursor is mixed with the polymer solution in alcohol, and then, the electrospun porous HPC polymer mat is converted to porous ceramic fiber network by using a suitable calcination treatment. Nan0 and submicron fibers of Sn02 have been deposited on the microelectromechanical system (MEMS) device, consisting oxidized silicon (Si/SiO2) wafer with gold (Au) electrode pattern. Such micro-device is highly suitable for the room temperature gas (typically hydrogen) sensing application. In conclusion, very effective polymer assisted processing of ceramic oxide nano and submicron fibers have been demonstrated for the low temperature gas sensing application. INTRODUCTION Electrospinning is a technique to produce the nano and the submicron sized polymer fibers.' A highly dense polymer solution is utilized for this purpose, which is placed in a syringe under the application of an electric field. If sufficiently high voltage is applied in between the syringe and the collector, then a polymer jet is ejected from the tip of the syringe. The polymer fibers are then deposited on the grounded substrate. Various electrospinning parameters, such as the molecular weight and the concentration of the polymer, the solvent type, the solution viscosity and the conductivity, the surface tension, the tip-to-collector distance, the applied voltage, the flow rate, and the surrounding atmosphere, significantly affect the characteristics of the electrospun polymer fibers.2 Although fibers of different polymers are formed via electrospinning technique, the electrospinning characteristics of the hydroxypropyl cellulose (HPC) polymer have not been investigated in the literature, which has been utilized earlier in the powder form, as a steric stabilizer, in the synthesis of nanocrystalline ceramic oxide powders.' In this investigation, we Nanotubes, Nanorods, and Nanowires

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demonstrate the use of electrospinning technique in forming the HPC polymer fibers and utilize them as templates in the synthesis of semiconductor tin oxide (Sn02) fibers. EXPERIMENTAL Anhydrous ethanol was purchased from Alfa Aesar (U.S.A.) and anhydrous 2-propanol, HPC polymer (molecular weight 80000 g/mol) and tin(I1) chloride (SnC12) were supplied by Sigma-Aldrich (U.S.A.). All the chemicals were used as received. The experimental set-up consists of a syringe, syringe pump, Cu-plate, Al-foil, and a high voltage power supply (0-30 kV). The syringe having a metallic tip holds the viscous polymer solution and is mounted over the syringe pump, which maintains the flow of the polymer solution to the tip of the syringe at a constant rate. The Cu-plate covered with Al-foil is placed in front of the syringe tip to collect the electrospun HPC nano and submicron fibers by adjusting the tip-to-collector distance as desired. The positive terminal of the high voltage power supply is connected to the metallic syringe-tip while the negative (ground) terminal is connected to the Cu-plate covered with A1 foil. Proper amount of the HPC polymer (15 wt%) was weighed and dissolved completely in the appropriate solvent under continuous stirring using the magnetic stirrer. The beaker containing the polymeric solution was covered with a paraffin tape during stirring. Proper amount of HPC solution is taken in the syringe, which is then placed on the syringe pump. After applying the desired electric potential in between the metallic syringe-tip and the Cu-plate, the HPC solution is continuously fed to the syringe-tip at the constant flow rate of 20 yllmin using the syringe pump. During the electrospinning process, the Al-foil gradually changed its color from grey to white indicating the deposition of the HPC fibers. The electrospinning of the HPC polymer dissolved in anhydrous ethanol and 2-propanol was conducted for the two different tipto-collector distances, 10 cm and 15 cm, by using 15 kV, 25 kV, and 30 kV applied voltages. To synthesize Sn02 fibers using the electrospun HPC fibers, 0.45 M of SnC12 was dissolved completely in 15 wt% HPC solution in ethanol. The tip-to-collector distance of 10 cm was selected with the applied potential difference of 15 kV. An oxidized silicon (Si/SiOz) wafer having the patterned gold (Au) electrodes was placed on the A1 foil to collect the HPC fibers with the Sn-precursor. After the electrospinning process, the Si/SiOz wafer was subjected to the thermal treatment at 700 "C using the heating rate of 30 "C/min and holding time of 1 h. The substrate was then cooled down to room temperature inside the furnace. The morphology and the average diameter of the electrospun HPC and Sn02 fibers were analyzed using the scanning electron microscopy (SEM) (JSM-6400F, JEOL, Tokyo, Japan). To avoid any surface charging during the SEM analysis, the fibers were sputter coated with 30 nm Au-Pd layer. The energy dispersive analysis of x-rays (EDX) was performed to study the bulk chemistry of the fibers. X-ray diffraction (XRD, Rigaku, Japan) analysis was performed using the Cu Ka radiation (wavelength=1.5418 A). RESULTS The SEM images of the HPC fibers obtained using the ethanol solution, with the tip-tocollector distance of 10 cm and the applied voltages of 15 kV, 25 kV, and 30 kV, are presented in Figs. l(a)-(c) respectively at low magnification. The HPC fibers are formed with beads under these processing conditions with typical bead-on-string morphology. The average bead size and

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Fig. 1. SEM micrographs, at low magnification, of HPC fibers obtained via electrospinning using anhydrous ethanol as a solvent and the tip-to-collector distance of 10 cm. In (a), (b), and (c), the applied voltage is 15 kV, 25 kV, and 30 kV respectively.

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their number density are observed to be a function of the applied voltage. With increasing applied voltage, the average bead size is noted to increase with the bead number density, Figs. l(a) and l(b). With further increase in the applied voltage, Fig. l(c), the bead number density increases; however, the average bead size is qualitatively noted to decrease. The SEM images of the HPC fibers obtained under high magnification are presented in Fig. 2. It is clearly noted that, the average HPC fiber diameter increases first with increasing applied voltage, Figs. 2(a) and 2(b), and then decreases, Fig. 2(c). The SEM image of the HPC fibers obtained using the anhydrous ethanol solution, with the tip-to-collector distance of 15 cm and the applied voltages of 15 kV, 25 kV, and 30 kV are presented in Figs. 3(a)-(c) respectively. It is qualitatively observed that, the average HPC fiber diameter, the average bead size, and their number density increase with increasing applied potential difference. However, the comparison of Figs. 1 and 3 shows that, for the tip-to-collector distance of 15 cm, no decrease in the average fiber size and average bead size is noted at the highest applied voltage of 30 kV. The SEM images of the HPC fibers obtained using the 2-propanol solution, with the tip-to-collector distance of 10 cm and the applied voltages of 15 kV, 25 kV, and 30 kV are presented in Figs. 4(a)-(c) respectively at low magnification. In contrast to the bead-formation tendency with the HPC solution in ethanol, Fig. 1, the HPC fibers are formed without any beads under all processing conditions when 2-propanol is used as a solvent. It is also qualitatively noted that, the average HPC fiber diameter increases first with increasing applied voltage, and then, it decreases. This variation is similar to the one that is observed in Fig. 2. Typical EDX analysis of the HPC fiber is presented in Fig. 5(a), where a strong C-peak is seen, which is originating from the HPC polymer (Note: Au and Pd peaks are due to sputtering process). The EDX spectra obtained from the calcination treatment of the HPC fibers with Sn2+ and C1- ions is presented in Fig. 5(b). Very strong C-peak, as observed in Fig. 5(a), has been reduced in intensity significantly after the calcinations treatment. The presence of Sn-peak is, however, noted with the absence of C1-peak. This possibly suggests the elimination of C1-peak from the bulk of the fibers after the calcinations treatment. The SEM image of the SnO2 fibers deposited on the Si-wafer with the Au-electrodes is shown in Fig. 6(a). Highly porous fibrous network of SnO;! fibers short circuiting the Au electrodes is observed in Fig. 6(a). The XRD spectra of Sn02 fibers is presented in Fig. 6(b), where (1 lo), (101), and (21 1) peaks of tetragonal crystal structure are identified by comparison with the JCPDS data file # 41-1445. DISCUSSION The HPC fibers are successfblly synthesized using the granular HPC powder via electrospinning technique. According to the proposed mechanism4i5of the fiber formation via electrospinning, the polymer solution held in the syringe forms a droplet with flat surface, which is held at the syringe tip by the surface tension force. When the high voltage is applied in between the collector and syringe tip, the droplet surface is pulled into the approximate shape of a section of a sphere by the electric forces and the surface tension. As the bulge forms, the charges produced by the applied electric field, move through the polymer solution to concentrate in the area that is protruding the most. The accumulation of the electric charges causes the droplet surface to protrude more, and since the charge per unit surface area is highest near the greatest protrusion, the surface is pulled into a conical shape, which is known as a Taylor cone.

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Fig. 2. SEM micrographs, at high magnification, of HPC fibers obtained via electrospinning using anhydrous ethanol as a solvent and the tip-to-collector distance of 10 cm. In (a), (b), and (c), the applied voltages are 15 kV, 25 kV, and 30 kV respectively.

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Fig. 3. SEM micrographs, at low magnification, of HPC fibers obtained via electrospinning using anhydrous ethanol as a solvent and the tip-to-collector distance of 15 cm. In (a), (b), and (c), the applied voltages are 15 kV, 25 kV, and 30 kV respectively.

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Fig. 4. SEM micrographs, at low magnification, of HPC fibers obtained via electrospinning using anhydrous 2-Propanol as a solvent and the tip-to-collector distance of 10 cm. In (a), (b), and (c), the applied voltages are 15 kV, 25 kV, and 30 kV respectively.

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C

Sn

h

0

1

2

3 4 5 Energy (keV)

6

7

8

Fig. 5. EDX analysis of the HPC fibers (a), and Sn02 fibers (b), which are obtained after calcining the HPC fibers, containing the Sn2+ and C1- ions, at 700 OC for 1 h.

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Fig. 6 . (a) SEM image of the SnO2 fibers over the Si/SiOz substrate having the patterned Au electrodes, (b) XRD spectrum of the SnO2 fibers.

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The charge per unit area at the tip of the cone increases as the radius near the tip of the cone decreases. As the applied potential is increased, a jet of charged polymer solution is pulled from the syringe tip and the electrospinning process begins. If the applied potential difference is increased further, multiple jets may be ejected from the Taylor cone.6 As the jet travels from the base to the collector, its diameter decreases and the length increases in a way that keeps constant mass of charged polymer jet per unit time passing any point on the axis. The electric charges, in the form of ions, tend to drift in response to the applied electric field, and they transfer the forces from the electric field to the polymer mass. As the electric charges move to the grounded terminal, they complete the electrical circuit, which provides the energy needed to accelerate the polymer. As the jet moves forward, the solvent gets evaporated depending on the surface tension of the polymer solution. This causes the jet diameter to decrease faster. As the jet diameter decreases with the simultaneous increase in its length, the electric charges are brought closer, and as a result, they start repelling each other. This causes the jet to expand in radial direction and stretch along its axis. Upon closer approach of the electric charges, the radial force become stronger than the cohesive strength of the charged polymer jet, and as a consequence, a single charged polymer jet splits into two or more jets, which is known as ‘~playing’.’~~ This jet division may take place several times in rapid succession producing large number small electrically charged fibers moving towards collector. The charge polymer fibers are, thus, deposited on the grounded substrate. Since the deposited polymer fibers are charged, they repel the polymer fibers, which are deposited subsequently. Hence, the deposited polymer fibers do not overlap and produce a highly porous non-woven polymer fibers mat, as observed in Figs. 1-4. The HPC fibers with beads are formed with a typical ‘bead-on-string’ morphology when anhydrous ethanol is used as a solvent instead of 2-propanol, which is similar to the observation reported earlier for other During the electrospinning, the beads are basically formed if the forces such as viscoelastic, charge-repulsion, and electrostatic attraction, which tend to elongate the charged polymer jet are overcome by the surface tension force, which tend to break up and spherodize the charged polymer jet. The formation of beads with anhydrous ethanol, hence, suggests that the surface tension force on the charged polymer jet is not completely overcome by the combination of all counter forces with ethanol as a solvent. This also indicates that, the HPC solution in 2-propnaol has higher solution viscosity and lower surface tension relative to those of the HPC solution in ethanol. Further, when ethanol is used as a solvent, it is noted that, the average HPC fiber diameter and the average bead size increase first with increasing applied voltage, and then decrease. The charge concentration within the polymeric solution increases with increasing applied voltage. As a result, more amount of charged polymer is possibly ejected from the syringe, which increases the average HPC fiber diameter. However, at higher applied voltage, the multiple jet formation and the splaying may take place, which not only reduce the average HPC fiber diameter but also the average bead size. Such effect is, however, only observed when the tip-to-collector distance is 10 cm and not for the working distance of 15 cm. It is likely that, the reduced electric field strength at higher working distance is not able to cause the multiple jet formation and the splaying to occur effectively, resulting in gradual increase in the average HPC fiber diameter and the average bead size with increasing applied potential difference. A detail discussion (qualitative and quantitative) on the effect of applied voltage and tip-to-collector distance on the avera e polymer (including HPC) fiber diameter and the average bead size could be found elsewhere.”

F

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In the present investigation, SnO2 fibers have been produced using the HPC fibers as templates. For this, Sn2+and C1- ions have been mixed with the HPC solution in ethanol and HPC/Sn2+/C1-composite fibers have been obtained. Using higher calcination treatment under the normal atmospheric conditions, the HPC polymer is decomposed with simultaneous removal of C1- ions, which results in the formation crystalline SnO2 fibers having the tetragonal crystal structure. Moreover, the porous structure of the Sn02 fibers short circuits the Au electrodes, which are patterned on the Si/SiOz-wafer; this in turn suggests the semiconducting nature of the SnO2 fibers. Such a design is highly suitable for the gas sensing application and work is in progress to measure the H2 sensitivity of the SnO2 fibers derived using the electrospinning process. CONCLUSIONS The HPC polymer fibers have been synthesized via electrospinning technique using two tip-to-collector distances and two different solvents with increasing applied voltage within the range of 10-30 kV. The HPC fibers with beads are obtained when ethanol is used as a solvent; however, the beads are eliminated using the HPC solution in 2-propanol. The average HPC fiber diameter and the average bead size are observed to be fknction of the applied voltage and the tip-to-collector distance, and they follow the trend which is well described by the known electrospinning mechanism. ACKNOWLEDGEMENTS The authors thank UCF, Florida Space Grant Consortium (FSGC), NASA-Glenn (NASA NAG 3275 1), KSC-NASA, National Science Foundation (NSF EEC-0 1367 10 and NSF CTS 0350572) (Seal) and (NSF CAREER, ECS-0348603 (Cho)) for hnding the sensor and the nano-technology research.

&-r%-RENB ID. Reneker and I. Chun, “Nanometre Diameter Fibers of Polymer Produced by Electrospinning”, Nunotechnology, 7 , 2 16-23 (1996). 2 Z. Huang, Y. Zhang, M. Kotaki, and S. Ramakrishna, “A Review on Polymer Nanofibers by Electrospinning and Their Applications in Nanocomposites”, Compos. Sci. Technol., 63,2223-53 (2003). 3S. Shukla, S. Seal, and R. Vanfleet, “Sol-Gel Synthesis and Phase Evolution Behavior of Sterically Stabilized Nanocrystalline Zirconia”, J. Sol-Gel Sci. Technol., 27, 1 19-36 (2003). 4 D. Reneker, A. Yarin, H. Fong, and S. Koombhongse, “Bending Instability of Electrically Charged Liquid Jets of Polymer Solutions in Electrospinning”, J. Appl. Phys., 87, 453 1-47 (2000). ’Y. Shin, M. Hohman, M. Brenner, and G. Rutledge, “Experimental Characterization of Electros inning: the Electrically Forced Jet and Instabilities”, Polymer, 42,09955-67 (2001). M. Demir, I. Yolgor, E. Yilgor, and B. Erman, “Electrospinning of Polyeurethene Fibers”, Polymer, 43,3303-09 (2002). ’5. Deitzel, J. Kleinmeyer, D. Harris, and N. Tan, “The Effect of Processing Variables on the Mo hology of Electrospun Nanofibers and Textiles”, Polymer, 42,26 1-72 (2001). H. Fong, I. Chun, and D. Reneker, “Beaded Nanofibers formed During Electrospinning”, Polymer, 40,4585-92 (1999).

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a?

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9S. Shukla, E. Brinley, and S. Seal, “Electrospinning of Hydroxypropyl Cellulose Polymer Fibers and Their Application in Synthesis of Nan0 and Submicron Tin Oxide Fibers”, Polymer 46, 12130-145 (2005). 10 K. Lee, H. Kim, H. Bang, Y. Jung, and S. Lee, “The Change of Bead Morphology Formed on Electrospun Polystyrene Fibers”, Polymer 44,4029-34 (2003).

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Growth of Quasi-Aligned AIN Nanofibers by Nitriding Combustion Synthesis Mohamed Radwant and Yoshinari Miyamoto*?** Department of Nano/Micro Structure Control, Smart Processing Research Center, Joining and Welding Research Institute, Osaka University, Osaka 567-0047, Japan

Quasi-aligned ALN nanofibers were formed by the nitriding combustion synthesis according to a unique micro-reactor model. A charge composed of aluminum and aluminum nitride diluent powders (40/60 mol%) with a mixture of yttria and ammonium chloride as additives (5 wt% each) was combusted a t low nitrogen gas pressures of 0.25 MPa. The FE-SEM images of as-synthesized AIN product showed the formation of ball-like grains (same shape and size as the original A1 reactant) that consisted of a thin surface nitride layer or crust cover quasi-aligned AIN nanofibers grown in the interior. The crosssectional view is sea anemone l i e . Formation of this novel morphology is believed to occur through a two-stage process. The first one occurs a t the preliminary stage of the combustion outside Al particles. After the ignition, the heat generated causes the sublimation and dissociation of ammonium chloride into various gaseous species. This effectively interrupts the combustion and slows down the increase of reaction temperature. In addition, yttria interacts with the native oxide layer present on the surface of Al particles and forms a stable AI-N-Y-0 crust. The second stage begins by the infiltration of various gaseous species such as HCtg), NHy,), and N2(,) through the crust into the molten A1 cores. The "crust-core" systems function as "micro-reactors" where both the nitridation and growth processes occur inside. The molten A1 cores are spontaneously halogenated to AICl3 vapors and the nitridation proceeds by the gas-gas reaction of AlCl3 and NH3/N2 vapors. The AIN nanofibers are then grown from the vapor phase quasi-aligned inside the micro-reactors by VLS and VS mechanisms.

I. Introduction

D

URING the past two decades, aluminum nitride (AIN) has attracted considerable interest in the electronics industry because it possesses an excellent combination of material properties, including high intrinsic thermal conductivity ( 320 W/mK), wide band gap (6.2 eV), high electrical resistivity (> 10" C2. cm), low dielectric constant (8.6), high mechanical strength, chemical stability, and low thermal expansion coefficient (4.2 x 10-6"C)-1 matches to both Si and G a N semiconductors. AIN components and substrates are used for various applications in power electronics (electrical engines), microelectronics (LSI circuits, sensor carriers), naval radio and defense systems, railway systems, aeronautical systems, and environmental Recently, the growth of one-dimensional (1-D) A1N nanostructures such as nanowhiskers; nanowire~,'-~nanofibers:

-

B. Derbydontributing editor

Manuscript No. 22034. Received July 12, 2006; approved January 25, 2007. Radwan is deeply grateful to the Japan Society for the Promotion of Science (JSPS) for the postdoctoral fellowship grant. 'Member, The American Ceramic Society. "Fellow, The American Ceramic Society. 'Author to whom correspondence should be addressed. e-mail: radwan@jwri. Osaka-u.ac.jp

Nanotubes, Nanorods, and Nanowires

nanobelts," nano-pillars," nanoribbons,12 n a n ~ t u b e s , ' ~ - ' ~ n a n o ~ o n e s , nanotips,I8 ~~.~~ and nanorods'' has become the focus of scientific research and proposed promising applications such as field emitters, flexible pulse-wave sensors, and nanomechanical resonators.16," Several methods have been developed for the preparation of 1-D A1N nanostructures, mostly by the nitridation at elevated temperatures of 1000"-1350°C in conventional horizontal-type tube furnaces. These methods can be categorized as follows: (i) A halide-assisted direct nitridation of starting charges contains halide compounds such as A1/A1X3>12 A1X3,14 or Al/NH4X7.' where X: CI, F, or I. The Al halide acts as a transport agent and supports the nitridation in the vapor phase and unidirectional growth by a vapor-solid (VS) mechanism. (ii) A catalyst-assisted direct reaction of metallic Al and NH3/N2. In general, a small amount of a metallic catalyst (0.5-1.0 mmol/l g Al) such as Ni,6 C O , ' ~or, ~A~U ' ~ , "is used. It forms eutectic quasi-molten droplets that act as active nucleation sites and promote the vapor-liquid-solid (VLS) growth mechanism of 1-D AIN nanostructures.6 (iii) A template-assisted carbothermic reduction and nitridation of A1203/AI mixture^.^ A1N nanowires have been synthesized in bulk from carbon nanotubes (CNTs), which act as a carbon source and template for confined carbothermic reduction and nitridation reactions. The synthesis mechanism was described by space-limited nucleation and the growth occurred from the vapor phase along the template channels. Oriented growth of 1-D AIN nanostructures has been reported by halide-assisted nitridation and deposition on catalyzed substrates (halide-CVD on catalyst-coated silicon or quartz wafers) at moderate temperatures of 700°-850"C,1'3'6317and the direct nitridation of Al in NH3/N2at 1100°C under confinement of a Ni-catalyzed alumina template.8 The combustion synthesis (CS) method (also called the selfpropagating high-temperature synthesis or SHS) is considered to be an economical approach for the production of pure AIN powder.2623 It involves the direct nitridation of metal aluminum, which is a typical exothermic reaction as follows: A1

+ iN2

-+ AIN(AHo

= -318 kJ/mol)

(1)

The heat evolved sustains the reaction so that no extra energy is needed except the small amount used for initiating the reaction and the reaction is completed very quickly within seconds. High conversions can be attained by careful selection of the combustion parameters such as nitrogen pressure, amount and type of diluents, etc. However, the literature of CS of AIN shows that the product morphology is complex and often consists of various grain morphologies such as agglomerated particles, whiskers, faceted particles, rods, pyramids, e t ~26. Although ~ ~ a uniform morphology is very important to engineer the properties of AIN-based materials or devices, a morphologycontrolled synthesis condition has not yet been realized. In this paper, we report novel observation of the growth of quasialigned AIN nanofibers inside the reacting Al particles during the combustion under a low nitrogen pressure. A new growth 295

model is proposed to explain the formation of this unique morphology. As far as we know, this mode of growth has not been reported in the literature before. 11.

Experimental Procedure

High-purity Al powder ( > 99.9%, Toyo Aluminum KK, Tokyo, Japan), AIN diluent powder (type H, >99.9%, Tokuyama K.K., Hino, Tokyo, Japan), NH4C1 ( > 99%, Nacalai Tesque, Inc., Kyoto, Japan), and Y203(99.9%, Shin-Etsu Chemical Co., Tokyo, Japan) were used as the starting materials. Figure 1 shows the morphological characteristics of Al and AIN powders; the average particle size is 23 and 0.5 pm, respectively. The reaction charge was composed of Al, AIN (AI/AlN=40/60 mol%), and promoting additives of NH4CI and YzO3 (5 wt% each). The reaction powders were mixed using mortar and pestle for 10 min and then sieved through a 212-pm sieve to disperse any large agglomerates. Fifty grams of the powder mixture was poured into a porous graphite container ($42 mm x 90 mm H) and packed by tapping to a relative density -60% of the theoretical density. The density was estimated by measuring the weight of the powder mixture in a fixed volume of the container. The graphite container was then placed in the combustion chamber. Two W-Re thermocouples (connected to a data acquisition system) were inserted into the center of the charge (one at the middle and another near the top surface) at a fixed distance of 30 mm and used to record the temperaturetime pattern of the combustion and determine the combustion speed by measuring the time lapsed for the wave passage between the two thermocouples. The chamber was evacuated and subsequently filled with high-purity N2 gas (99.999%) up to 0.25 MPa pressure. The combustion was initiated from the bottom by igniting a 2-g ignition pellet (AI+AIN, 1:l) placed at the bottom of the packed powder by passing an electric current (60 A, 20 V) for 10 s through a carbon ribbon under the pellet. The combustion reaction was completed in about 5 min and the chamber was then cooled to room temperature in -30 min. The reaction product was visually observed. The product phases were identified by Xray powder diffraction (XRD; JEOL, JDX-3530, Tokyo, Japan) using CuKa radiation. The morphology of as-synthesized powder was observed by field emission scanning electron microscopy (FE-SEM; ERA-8800, ELIONIX, Tokyo, Japan). Samples for SEM observation were coated with thin films of sputtered gold to reduce electrical charge-up. 111. Results and Discussion

The content of diluent (60 mol%) in the reaction mixture was chosen according to a previous study for the relationship between Al molar ratio and nitrogen pressure on the yield and properties of AIN p r o d u ~ t . ~It~ is , ~used ' to reduce the reaction temperature and prevent coagulation of melted aluminum

particles. The as-synthesized AIN cake was very fragile with only a white color. The microstructure of as-synthesized AIN powder was observed by FE-SEM using representative samples from three different locations in the product cake: top surface, side surface, and middle center. The grain morphologies of as-synthesized AIN particles are given in Fig. 2. The microstructure consists of two major types: aggregates of irregular particles (-0.5 pm, same as original AIN diluent) and ball-like grains (same size and shape as original Al particles) consist of thin crust ( 5150 nm) covers unique quasi-aligned AIN nanofibers grown in the interior. Their cross-sectional view is similar to oval disk of the sea anemone. Figure 3 shows the XRD pattern of the as-synthesized product. The diffraction lines are assigned to a hexagonal AIN structure similar to the bulk AIN powder reported (JCPDS-file 2 5 1 133). Residual metallic Al has not been detected. From the SEM observations, one can conclude that: ( I ) The AIN diluent did not participate in the formation of AIN particles and played only a passive role in controlling the combustion temperature and dispersing the Al particles. (2) There was no grain growth or sintering for both the formed and original AIN diluent particles. (3) The quasi-aligned AIN nanofibers were formed inside the reacting Al particles. As far as we know, this mode of growth, encapsulation in reacting particles, has never been reported before. The question is how were these balls of quasi-aligned AIN nanofibers formed? The formation of a shell-core system during the course of nitridation of Al metal is known in the direct nitridation (DN) method and reported as a "coreshell" mode1.29-34In this method, the nitridation takes place via three steps: nitridation at the surface of the particles with the formation of a crystalline nitride shell, breakaway or flowout of molten or vaporized Al core, and volume nitridation outside the shell with a remaining hole or an empty core. The final morphology of the AIN product is honeycomb like (a clear SEM photograph of this morphology can be seen in our recent paper by Radwan and Bahgat34).In the combustion synthesis method, both the surface nitridation and breakaway were also observed when moderate combustions were promoted by using small amounts of additives (low),,"I such as C and NH C13s38 and/or ignition under an appropriate nitrogen pressure.3LL The nitride skins were formed at the preheating stage of the combustion and then molten Al flowed out with the formation of an eggshell-type AIN morphology. Neither of the previous observations can account for the present SEM observations, which requires a new growth model without a breakaway. The typical temperaturetime history of the nitridation reaction (Fig. 4) sheds light on the behavior of the combustion. The temperature was measured in the middle center of the charge. The pattern shows a mild combustion with no explosive mode and has a relatively low rate of temperature increase. It took -53s to reach 600°C (below melting of Al), -73s to reach

Fig. 1. Scanning electron microscopy micrographs of starting powders: (a) A1 and (h) aluminum nitride diluent.

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Fig. 2. Field emission scanning electron microscopy images of the as-synthesized aluminum nitride product.

-

lOWC, 160 s to T', (1620"C), and lasted -27 s in the afterburning stage. We noticed that in the temperature-time histories of other combustion experiments without NH4C1 and Y203 additions, once the combustion wave begins, the temperature increases rapidly close to its maximum value. Owing to the

low combustion temperature, the grain growth and sintering of

AIN particles were avoided. The speed of the combustion reaction was determined by measuring the time lapsed for the wave passage between two thermocouples inserted into the center of

A1N

20

30

40

50

20 (7

60

70

80

Fig. 3. X-ray diffraction pattern of the as-synthesized aluminum nitride

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Fig. 4. Temperaturetimevariation at the combustion front.

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Fig.5. Schematic illustration of the growth model of quasi-aligned aluminum nitride nanofibers by the combustion synthesis method

the reaction bed at a (vertical) distance of 30 mm. The combustion had a slow speed (0.26 mmjs). In Fig. 5, we propose a new growth model to explain the formation of quasi-aligned AIN nanofibers by the CS based on a “micro-reactors’’ model. The formation of nanofibers occurred through two stages.

( I ) Formation of Micro-Reactors This stage started in the preliminary stage of combustion and occurred outside the Al particles through two steps. In the first step (below 600”C),NH4CI dissociated into HCI and NH3 vapors (NH4C1 sublimes at 350°C and dissociates at 520°C): NH~CI(S, V)

4

HCl(g)

+ NH,(g)

(2)

The sublimation and decomposition of NH4CI are endothermic and produce several gaseous species. These reactions absorb sufficient heat and disturb the direct nitridation of Al particles with N2 gas, which retards the wave propagation. It also provides enough time for another endothermic reaction between yttria and the surface alumina layer (at T 2 melting point of Al) with the creation of a thin crust on the surface of the Al particles:

+

+

YzO~(S) A1203(S) N2(g) - O ( s )stable thincrust

+

A1 - Y - N

(3)

The energy dispersive X-ray spectroscopy microanalysis showed the presence of yttrium in the composition of the thin crust (Fig. 6). The formation of these new crusts also suppresses the AI-N interaction and slows down the combustion propagation. Because of the low heat evolution and the presence of new protective crusts, there was neither breakaway nor explosion of molten cores.

Ammonium chloride plays a critical role because it produces hydrogen chloride, which can be considered to be a key intermediate product, acting as a “catalyst.” HCI(v) promotes the vaporization of molten aluminum cores into volatile aluminum chloride species and facilitates progress of nitridation through sequence of spontaneous chlorination and nitridation intermediate reactions. The semi-molten AI-Y-N-O crusts seemed to function as catalyzed self-substrates and provided active sites that promoted homogenous nucleation of AIN embryos on the inner surface of the crusts from the vapor phase by a VLS mechanism. The AIN nanofibers might condense from the vapor phase, after a critical (low) degree of supersaturation inside the “micro-reactors’’ is attained, and grow on the preceding embryos by a VS growth mechanism in an epitaxial way according to the classical crystal growth theory.43 No droplets could be observed at the tips of these nanofibers. This results in a unique oriented growth in the interior of the reacting particles normal to the inner crust surface. The first stage of the combustion, formation of micro-reactors, was the essential step for the formation of AIN nanofibers inside reactant particles. The postulation of nucleation of AIN by VLS in molten droplets or layer with a subsequent fiber growth through VS mechanism is in close agreement with the observations of Moya and colleagues.2494345 To the best of our knowledge, this mode of growth inside reactant Al particles has not been observed before and it was not expected in combustion reactions due to the high increase

( 2 ) Nitridation and Growth The crust-core systems function as unique “micro-reactors.’’ The various gaseous species present (HCl(g), NH3(g), N2(g)) diffuse through the crust (through pores or cracks) into the molten AIN cores. Both the nitridation and growth steps then occur inside the developed micro-reactors. The nitridation proceeds via spontaneous chlorination-nitridation sequences similar to that reported in our previous results of direct nitridation of an A1/NH4C1 mixture.42 Gaseous hydrogen chloride is very active and spontaneously reacts with molten Al cores to produce AICI, vapors, which are nitrided by a gas-gas reaction as follows: AI(I) 298

+ 3HCl(g)

+

AICI,(V)+#H2(g)

(4)

Fig.6. Energy dispersive X-ray spectroscopy pattern of the thin nitride crust. Progress in Nanotechnology: Processing

in nitridation temperature and fast speed of combustion. The current combustion condition was successful creating new micro-reactors during the combustion reaction, which promoted this mode of oriented growth inside.

IV. Conclusion Quasi-aligned AIN nanofibers were formed by the combustion synthesis according to a new micro-reactor model. This route might be a possible method for economical growth of AIN nanofibers, which are obtainable only by complicated nitridation reactions at elevated temperatures. AIN with this unique morphology can be used not only for AIN ceramics and composites but also in nanotechnology applications.

References ‘L. M. Sheppard, “Aluminum Nitride: A Versatile but Challenging Material,” Ceram. Bull., 69[11] 1801-12 (1990). *J. H. Harris, “Sintered Aluminum Nitride Ceramics for High-Power Electronic Applications,” JOM, 50 [6] 56-60 (1998). ’B. H. Mussler, “Advanced Materials and Powders-Aluminium Nitride (AIN),” Am. Ceram. Soc. Bull., 79 [6] 45-7 (2000). 4J. A. Haber, P. C. Gibbons, and W. E. Buhro, “Morphological Control of Nanocrystalline Aluminum Nitride: Aluminium Chloride-Assisted Nanowhisker Growth,” J. Am. Chem. Soc., 119, 5455-6 (1997). 5Y. Zhang, J. Liu, R. He. Q . Zhang. X. Zhang. and J. Zhu, “Synthesis of Aluminum Nitride Nanowires from Carbon Nanotubes,” Chem. Maler., 13,389p905 (2001). 6Q.Wu, Z. Hu. X. Wang, Y. Lu. K. Huo, S. Deng, N. Xu, B. Shen, R. Zhang, and Y. Chen. “Extended Vapor-Liquid-Solid Growth and Field Emission P r o p erties of Aluminium Nitride Nanowires,” J. Marer. Chem., 13, 2024-7 (2003). ’M. Radwan and M. Bahgat. “Novel Growth of Aluminium Nitride Nanowires,” J. Nanosci. Nanorechnol., 6 [2] 55861 (2006). ‘4. Wu, Z. Hu, X. Wang, Y. Hu, Y. Tian, and Y. Chen, “A Simple Route to Aligned AIN Nanowires,” Diamond Relat. Muter., 13, 3 8 4 1 (2004). 9H. Chen, Y. Cao, and X. Xiang, “Formation of AIN Nano-Fibers,” J. Cryst. Growrh, 224, 187-9 (2001). “Q. Wu, Z. Hu, X. Wang, Y. Chin, and Y. Lu, “Synthesis and Optical Characterization of Aluminum Nitride Nanobelts,” J. Phys. Chem. B, 107 [36] 97269 (2003). I’M. Yoshioka, N. Takahashi, and T. Nakamura, “Growth of the AIN NanoPillar Crystal Films by Means of a Halide Chemical Vapor Deposition under Atmospheric Pressure,’’ Marer. Chem. Phys., 86, 74-7 (2004). I2T. Xie, Y. Lin, G. Wu, X. Yuan, Z. Jiang. C. Ye, G. Meng, and L. Zhang. “AIN Serrated Nanoribbons Synthesized by Chloride Assisted Vapor-Solid Route,” Inorg. Chem. Commun., 7, 545-7 (2004). ”Q. Wu, Z. Hu, X. Wang, Y. Lu, X. Chen, H. Xu, and Y. Chen, “Synthesis and Characterization of Faceted Hexagonal Aluminum Nitride Nanotubes,” J. Am. Chem. Soc., 125, 101767 (2003). I4L.-W. Yin, Y. Bando, Y.-C. Zhu. D. Golberg. and M.-S. Li, “A Two-Stage Route to Coaxial Cubic-Aluminum-Nitride-Boron-Nitride Composite Nanotubes,” Adv. Mater., 16 [ I I] 929-33 (2004). ”V. N. Tondare. C. Bahsubramanian, S. V. Shende, D. S. Joag, V. P. Godbole, S. V. Bhoraskar, and M. Bhadbhade, “Field Emission from Open Ended Aluminum Nitride Nanotubes,” Appl. Phys. Lett., 80 [25] 4813-5 (2002). I6C. Liu, Z. Hu, Q . Wu, X. Wang, Y. Chen, H. Sang, J. Zhu, S. Deng, and N. Xu, “Vapor-Solid Growth and Characterization of Aluminum Nitride Nanocones,” J. Am. Chem. Soc., 127, 1318-22 (2005). ”C. Liu, Z. Hu, Q . Wu, X. Wang, Y. Chen, W. Lin, H. Sang, S. Deng, and N. Xu, “Synthesis and Field Emission Properties of Aluminum Nitride Nanocones,” Appl. Sur. Sci., 251, 220-4 (2005). ‘*S.-C. Shi, C.-F. Chen, S. Chattopadhyay, Z.-H. Lan, K.-H. Chen, and L.-C. Chen. “Growth of Single-Crystalline Wurtzite Aluminum Nitride Nanotips with a Self-Selective Apex Angle,” Adv. Funcr. Mater., 15 [5] 7 8 1 4 (2005).

Nanotubes, Nanorods, and Nanowires

I9S.-C. Shi. S. Chattopadhyay, C.-F. Chen. K.-H. Chen, and L.-C. Chen. “Structural Evolution of AIN Nano-Structures: Nanotips and Nanorods.” Chem. Phy.7. Lett., 418. 152-7 (2006). ”A. G. Merzhanov, “History and Recent Developments in SHS.” Ceram. Inr., 21, 371-9 (1995). *‘K. Tanihata and Y. Miyamoto, “Reaction Analysis on the Combustion Synthesis of Aluminum Nitride,” Inr. J. SHS, 7 121 209-17 (1998). ”V. V. Zakorzhevskii and I. P. Borovinskaya, “Regularities of Self-Propagating High-Temperature Synthesis of AIN at Low Nitrogen Pressures,” Inr. J. SHS., 7 [2] 199-208 (1998). 21 V. V. Zakorzhevskii, I. P. Borovinskaya, and N. V. Sachkova. “Combustion Synthesis of Aluminum Nitride.” Inorg. Mater., 38 [ I I] 1 1 3 1 4 (2002). 24J.S. Moya, J. E. Iglesias, J. Limpo, J. A. Escrioa. N. S. Makhonin. and M. A. Rodriguez, “Single Crystal AIN Fibers Obtained by Self-propagating HighTemperature Synthesis (SHS),” Acta Mater., 45 [8] 3089-94 (1997). 2SG.Jiang, H. Zhuang, J. Zhang, M. Ruan, W. Li, F. Wu, and 8 . Zhang. “Morphologies and Growth Mechanisms of Aluminum Nitride Whiskers by SHS Method-Part I.” J. Muter. Sci., 35, 57-62 (2000). 26 G. Jiang, H. Zhuang, J. Zhang, M. Ruan, W. Li, F. Wu, and B. Zhang, “Morphologies and Growth Mechanisms of Aluminum Nitride Whiskers by SHS Method-Part 2,” J. Mater. Sci., 35, 63-9 (2000). 27T. Sakurai, Y. Miyamoto, and 0. Yamada, “Combustion Synthesis of Fine and High-Purity AIN Powder and Its Reaction Control,” J. Soc. M a t . Sci. Jpn, 54 [6j 574-9 (2005) (in Japanese). 28T. Sakurai. 0. Yamada, and Y. Miyamoto. “Combustion Synthesis of Fine AIN Powder and Its Reaction Control,” Muter. Sci. Eng. A , 415. 40-4 (2006). 291. Kimura, K. lchiya, M. Ishii, N. Hotta, and T. Kitamura, “Synthesis of Fine AIN Powder by a Floating Nitridation Technique using an N2/NHJ Gas Mixture,” J. Mater. Sci. Lett., 8. 3 0 3 4 (1989). ’OH. Scholz and P. Greil, “Nitridation Reactions of Molten AI-(Mg, Si) Alloys,” J. Marer. Sci., 26, 669-77 (1991). ”K. Komeya, N. Matsukaze, and T. Meguro. “Synthesis of AIN by Direct Nitridation of Al Alloys,” J. Ceram. Soc. Jpn, 101 [I21 1319-23 (1993). ’*A.-J. Chang, S.-W. Rhee, and S. Baik, “Kinetics and Mechanisms for Nitridation of Floating Aluminum Powder,” J. Am. Ceram. Soc.. 78 [I] 3 3 4 0 (1995). ”T. Fujii, K. Yoshida, K. Suzuki. and S. Ito. “Direct Nitriding of Large Grains of Aluminum with 2 mm Size,” SolidStare Ionics, 141-142, 593-8 (2001). 14M. Radwan and M. Bahgat, “A Modified Direct Nitridation Method for Formation of Nano-AIN Whiskers,” J. Mater. Process. Techno/., 181, 99-105 9 7 ) . G. J. Jiang, H. R. Zhuang, W. L. Li, F. Y. Wu, B. L. Zhang, and X. R. Fu, “Mechanisms of the Combustion Synthesis of Aluminum Nitride in High Pressure Nitrogen Atmosphere (2);’ J. Marer. Synth. Process., 7 [I] 1-6 (1999). ‘k.-N. Lin and S.-L. Chung, “Combustion Synthesis of Aluminum Nitride Powder using Additives,” J. Marer. Re.7.. 16 [XI 2200-8 (2001). 37R.-C.Juang, C:J. Lee, and C . C . Chen, “Combustion Synthesis of Hexagonal Aluminum Nitride Powders under Low Nitrogen Pressure,” Mater. Sci. Eng. A. 357, 219-27 (2003). 38 C.-N. Lin, C.-Y. Hsieh, S.-L. Chung, J. Cheng, and D. K. Agrawal, “Combustion Synthesis of AIN Powder and Its Sintering Properties,” I n / . J. SHS, 13 [2] 93-106 (2004). I9S. M. Bradshaw and J. L. Spicer, “Combustion Synthesis of Aluminum Nitride Particles and Whiskers,” J. Am. Ceram. Soc., 82 [9] 2293-300 (1999). *J. Shin, D.-H. Ahn, M.-S. Shin, and Y.-S. Kim, “Self-Propagating HighTemperature Synthesis of Aluminum Nitride under Lower Nitrogen Pressures,’’ J . Am. Ceram. Soc., 83 [5] 1021-8 (2000). 4’G.J. Jiang, H. R. Zhuang, W. L. Li, F. W. Wu. B. L. Zhang, and X. R. Fu, “Mechanisms of the Combustion Synthesis of Aluminum Nitride in High Pressure Nitrogen Atmosphere (2),” J. Mater. Synth. Process., 9 [I] 49-56 (2001). 42M. Radwan, M. Bahgat, and A. A. El-Geassy, “Formation of Aluminium Nitride Whiskers by Direct Nitridation,” J. Eur. Ceram. Soc., 26 [I31 248S8 (2Ow. 4’W. B. Campbell, “Growth of Whiskers by Vapor-Phase Reactions”; pp. IS46 in Whisker Technology, Chapter 2, Edited by A. P. Levitt. John Wiley & Sons Inc., New York, 1970. “P. G. Caceres and H. K. Schmid, “Morphology and Crystallography of Aluminum Nitride Whiskers,” J. Am. Ceram. Soc., 77 [4] 977-83 (1994). 4SR. Fu, H. Zhou, L. Chen, and Y. Wu, “Morphologies and Growth Mechanisms of Aluminum Nitride Whiskers Synthesized by Carbothermal Reduction.” Muter. Sci. Eng. A , 266,4451 (1999). 0

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Synthesis and Optical Properties of Mullite Nanowires Han-Kyu Seong, Ungkil Kim, Myoung-Ha Kim, and Heon-Jin Choit Department of Materials Science and Engineering, Yonsei University, Seoul 120-749, Korea

Youngho Lee and Won-Seon Seo Reliability Assessment and Materials Evaluation Center, Korea Institute of Ceramic Engineering and Technology, Seoul 153-801, Korea

-

Single-crystalline mullite (3AI2O3 2Si02) nanowires have been synthesized on silicon substrates by forcing aluminum and chromium chloride powders to react under an ammonia gas flow. The diameter and length of the nanowires have uniform diameters of < 100 nm and several micrometers, respectively. High-resolution transmissionelectron microscope and selected-area electron diffraction analyses indicated that the nanowires were almost structural defect free and had a single-crystalline phase with a (0001 ) growth direction. The photoluminescence spectra showed that the mullite nanowires reached an emission peak at the center wavelength of 442 nm originating from the AI-0 bonds in preference to the A 1 4 and Si-0 bonds in the mullite bonding structure. Comparison of the photoluminescence between as-synthesized, oxygen-annealed, and plasma-etched nanowires indicates that the nanowires have few defects (i.e., oxygen vacancies).

1.

Introduction

D

shape, and size have been known to play a significant role in determining the properties of In this regard, one-dimensional nanostructures such as nanotubes and nanowires have been the focus of considerable research into their use to improve the physical and mechanical proper tie^.^^ Recently, carbon nanotubes (CNTs) have been reported as potential elements for reinforcement in ceramic matrix composites (CMCs) owing to their unique mechanical properties, such as excellent tensile strength (>200 GPa) and Young's modulus (1 TPa).6 Meanwhile, mullite whiskers have been investigated as reinforcements for high-temperature CMCS.~-" However, the whiskers are limited by whisker damages and matrix cracking during the course of processing and densification of composites. Such problems may be solved by using mullite nanowires that could have excellent mechanical properties due to a perfect single-crystalline nature free from defects. Their nano-size could be also useful for processing and densification of composites without reinforcement damage and matrix cracking. In this study, the synthesis and optical properties of single-crystalline mullite (3A1203 . 2Si02) nanowires using the chloride vapor transport (CVT) process are reported. IMENSIONALITY,

11. Experimental Procedure

The growth of the single-crystalline mullite nanowires was performed in a horizontal hot-walled CVT system (Fig. I). Si/Si02 substrates oxidized by the wet oxidation method were used. The solid Al (purity 99.5%) and CrCI2 (purity 99.9%) powders placed in quartz susceptors, respectively, were inserted into the center of a quartz tube at 2 in. intervals. The oxidized substrates, deposited with a 0.2-nm layer of Ni by sputtering, were placed in the quartz tube at a distance of 2 in. from the Al powders. The temperature of the furnace was increased at a heating rate of 5O"C/min from room temperature to the reaction temperature of 1000°C under a flow of ammonia gas at a rate of 20 cm3/min, maintained for 10 min under a constant flow of ammonia, and then cooled down to room temperature. The quartz tube was then degassed and purged with argon gas. The morphologies and crystal structures of the nanowires grown on the substrates were characterized using scanning electron microscopy (SEM) and X-ray powder diffraction. Further structure and stoichiometry analyses of the nanowires were performed using high-resolution transmission electron microscopy (HRTEM), selected area electron diffraction (SAED), and energy-dispersive X-ray spectroscopy (EDS). The photoluminescence (PL) spectra were measured at room temperature with an excitation wavelength of 325 nm (He-Cd C W laser). 111. Results and Discussion

Figure 2 shows the SEM images of the typical mullite nanowires synthesized on the substrates using a CVT process. The nanowires with a high aspect ratio were distributed over the entire area of the substrates (Fig. 2(a)) and had uniform diameters of < 100 nm and lengths of several micrometers (Fig. 2(b)). The overall crystalline nature of these nanowires was confirmed as a high-quality mullite structure using an X-ray diffraction pattern. The low-magnification TEM image in Fig. 3(a) shows that most nanowires have uniform diameters of < 100 nm and lengths of several micrometers. Although Ni was used as the catalyst and it

X. Miao--contributing editor

Manuscript No. 22766. Received February 3, 2007; approved February 28, 2007. This research supported in part by a grant from the Korea Research Foundation (MOEHRD, KRF-2005442- wO203) and the Second Stage of Brain Korea 21 Project in 2006. H. J. C. thanks the RAME Center, KlCET for the use of their facilities. 'To whom correspondence should be addressed. e-mail: [email protected]

Nanotubes, Nanorods, and Nanowires

Fig. 1. Schematic illustration showing the setup for the growth of mullite nanowires.

301

Fig.2. Scanning electron microscopy images of mullite nanowires grown on substrates by annealing at 1000°C for 10 min. (a) Low and (b) high magnification.

Fig. 3. (a) A low-magnification transmission electron spectroscopy (TEM) image of mullite nanowires. (b) Energy-dispersiveX-ray spectroscopy (EDS) spectrum of a selected individual nanowire in the image (a). (c) High-resolution TEM image of mullite nanowire. The inset shows the selected area electron diffraction pattern of the wire, recorded along the [110] zone axis.

302

Progress in Nanotechnology: Processing

clearly contributed to the growth of the nanowires, no Ni globules or elements in the nanowires were identified in the SEM observations and the EDS analysis. The failure to observe any Ni globules or elements suggests that the catalysts were etched out by the chloride vapors during the growth of nanowires. As shown in Fig. 3(b), the EDS analysis of a selected individual nanowire shows that it is composed of Al (80.7 wtX), Si (19.3 wt%), and 0 elements, which correspond to the stoichiometric composition of 3A1203.2Si02. The HRTEM observation (Fig. 3(c)) also shows that the mullite nanowires are smooth and almost structural defect free. The SAED analysis, recorded along the [ 1101 zone axis, further demonstrated that the nanowires are a single-crystalline form of the orthorhombic Pbum structure with a (0001) growth direction (inset of Fig. 3(c)). The HRTEM lattice image shows that the fringe spacing of the mullite nanowire corresponding to the interplanar distance of (0001) planes is 2.95 A. This lattice expansion on the [OOOI] plane compared with the bulk mullite may be ascribed to the quantum confinement effects and/or stresses generated on the curved surfaces of the nanowires, which is expected to be significant as the size approaches the nanometer scale.” In the present process, the major vapor species in the CVT reactor are Sic&, AIC13, NH3, H20, and HCI. Therefore, the overall reactions for the mullite nanowire growth could be as follows:

+

+

2SiC14(v) 6AICIj(v) 13HzO(v) = 3A1203 .2SiO~(s)+ 26HCl(v)

(1)

The SiCI4,AICI,, and H 2 0 vapors can be formed through the sub-reactions as follows:

+

+

~NH~(v) 2Al(s) 9cd&(s) = 2AICl,(v) IZHCI(v) 9Cr(s) 1 +2N2(v)

+

+

+

+

3sio2(s) 6CrC12(s) 4NH3(v) = 3SiC14(v) 6H2O(v) 6Cr(s) 1 +2N2(v)

+

+

Photoluminescence spectra of the mullite nanowires as-synthesized, annealed a t 0 2 for 2 h, and plasma etched f r o m CF,+O2 mixed gas at room temperature, respectively.

vacancies, the band gap of the silica becomes narrow dramatically.13 Similarly, an increase of the oxygen vacancies in the nanowires may induce a narrow band gap. As the mullite nanowires synthesized in this study are single crystalline with few defects, it could be used as novel reinforcements for hightemperature CMCs.

IV. Conclusions

(2)

(3)

As the reactants in Eq. ( 3 ) , SiCI4 and H20 vapors as the silicon and oxygen sources originate from the amorphous silica thin film on the oxidized substrate formed during wet oxidation. The CrC12 powder remains in the Cr powders on the quartz susceptor of the CVT reactor after the reaction and thus acts as a chloride source for the overall CVT process. Figure 4 shows a typical PL spectrum from the mullite nanowires in the as-synthesizedstate, annealed at 600°C in 0 2 for 2 h, and plasma etched by a CF4+7.8% O2 mixing gas at room temperature. The excitation wavelength was 325 nm at room temperature. All spectra show a strong PL peak at the center wavelength of 442 nm (2.8 eV in photon energy) that can be attributed to the quantum confinement effects and/or the AIL@ Si bonds of the mullite bonding structure. The intensity of the PL peak of annealed nanowires is similar to that of the as-synthesized one; however, it considerably increases for the plasmaetched nanowires. This result implies that the radiative centers leading to the measured PL are generated from the AIL0 bonds in preference to the AIM-Si bonds in the mullite bonding structure. Peng et a/.” observed that AI2O3 nanowires and nanobelts have various emission intensities in different thermal annealing atmospheres such as O2 and H2 owing to the optical transitions in oxygen-related defects, F+ (oxygen vacancy with one electron) center. In the mullite nanowires, however, the PL peak intensity and the position for the annealed nanowires were barely changed. This may be due to the fact that the mullite nanowires synthesized in this study had a perfect single-crystalline nature with few defects (i.e., oxygen vacancies). On the other hand, the increase of the PL intensity as well as red shifting of the peak position in the plasma-etched nanowires can be principally attributed to the increase of the oxygen vacancy density, resulting from the destruction of the Si-0 bonds in the nanowires. It has been reported that with the increase of the oxygen Nanotubes, Nanorods, and Nanowires

Fig. 4.

The mullite nanowires with high aspect ratios were synthesized on the substrates using CVT process. The nanowires had uniform diameters of < 100 nm and a length of a few micrometers. The TEM characterization and PL measurements from the nanowires with as-synthesized, annealed in oxygen, and plasma etched by a CF4+O2 mixing gas indicate the single-crystalline and defect-free nature of nanowires. The mullite nanowires could offer good opportunities for highly toughened ceramic, metal, and polymer composites. References ‘J. Ning, J. Zhang, Y. Pan, and J. Guo, “Fabrication and Mechanical Properties of Si02 Matrix Composites Reinforced by Carbon Nanotube,” Mazer. Sci. Eng., A357, 3 9 2 4 (2003). ’H. J. Choi. H. K. Seong, J. Chang, K. 1. Lee, Y. J. Park, J. J. Kim. S. K. Lee. R. He, T. Kuykendall. and P. Yang, “Single-Crystalline Diluted Magnetic Semiconductor GaN:Mn Nanowires,” A h . Muter., 17, 1 3 5 1 4 (2005). ’H. K. Seong, Y. Lee. J. Y. Kim, Y. K. Byeun, K. S. Hdn, J. G . Park. and H. J. Choi, “Single Crystalline AIGaN:Mn Nanotutes and Their Magnetism,’’ Adv. Mater., 18, 3019-23 (2006).

4A. Peigney, Ch. Laurent, E. Flahaut. and A. Rousset, “Carbon Nanotubes in Novel Ceramic Matrix Nanocomposites,” Ceram. IIII., 26, 6 7 7 4 3 (2000). ’M. B. Nardelli, B. 1. Yakohson, and J. Bernholc. “Brittle and Ductile Behavior in Carbon Nanotubes,” Phys. Rev. Leu., 81, 4 6 5 6 9 (1998). 6B. 1. Yakobson and P. Avouris, “Mechanical Properties of Carbon Nanotubes,” Appl. Phys., 80,287-327 (2001). ’T.I. Mah and K. S. Mazdiydsni, “Mechanical Properties of Mullite,” J. Am. Ceram. Sac., 66.699-703 (1983). ‘ S . Kanzaki, H. Tabata, T. Kumazawa. and S. Ohtd. “Sinteringand Mechanical Properties of Stoichiometric Mullite.” J. Am. Ceram. Suc., 68 [ I ] C - 6 c - 7 (1985). 9K. Okada and N. Otuska, “Synthesis of Mullite Whiskers and Their Application in Composites,” J. Am. Ceram. SOC.,74 [lo] 2412-8 (1991). ‘OH. J. Choi and J. G. Lee, “Synthesis of Mullite Whiskers.” J. Am. Ceram. Suc., 85 [2] 481-3 (2002).

“H. J. Choi, H. K. Seong, J. C. Lee, and Y. M. Sung, “Growth and Modulation of Silicon Carbide Nanowires,” J. Crystal Growth, 269, 472-8 (2004). ‘*X. S. Peng, L. D. Zhang, G. W. Meng, X. F. Wang, Y. W. Wang, C. 2.Wang, and G. S. Wu. “Photoluminescence and Infrared Properties of a-Al,O, Nanowires and Nanobelts,” J. Phys. Chem. B., 106. I 1163-7 (2002). ”H. R. Philipp, “Optical Properties of Non-Crystalline Si, SiO. SiO, and SiOz.” 0 J . Phys. Chem. SolidF, 32, 1 9 3 H 5 (1971).

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(Nao.8Ko.2)o.5Bio.5Ti03 Nanowires: Low-Temperature Sol-GelHydrothermal Synthesis and Densification Yu-Dong Hou,+ Lei Hou, Ting-Ting Zhang, Man-Kang Zhu, Hao Wang, and Hui Yan Department of Materials Science and Engineering, Beijing University of Technology, Beijing 100022, China

The sol-gel-hydrothermal processing of ( N ~ O . ~ K ~ . ~ ) ~ . ~ Bthe ~ . number ~ T ~ Oof~possible spontaneous polarization directions for (NKBT) nanowires as well as their densification behavior were the compositions near the MPB due to the coexistence of rhoinvestigated. The morphology and structure analyses indicated mbohedral and tetragonal phases. Therefore, NBT-based solid solutions that have an MPB structure and can be poled easily that the sol-gel-hydrothermal route led to the formation of phase-pure perovskite NKBT nanowires with diameters of 50have been studied r e c e n t l ~ . ~Among -~ the NBT-based binary lead-free piezoelectric systems, (Nao,sI&2)o.5Bio,5Ti03(NKBT) 80 nm and lengths of 1.5-2 pm, and the processing temperature systems have shown good piezoelectric performance and have was as low as l W C , but the conventional sol-gel route tended been the subject of the most extensive investigations because the to lead to the formation of NKBT agglomerated porous struccomposition was close to the MPB and the ceramics can be tured nanopowders, and the processing temperature was higher than 650°C. It is believed that the gel precursor and hydrotherpoled easily in a relatively low dc field ( 4 5 kvjmm). NKBT powders are traditionally prepared by a solid-state reaction.64 mal environment play an important role in the formation of the nanowires at a low temperature. Owing to the better packing In case of solid-state reactions, the starting materials are oxides o r carbonates of Bi, Na, K, and Ti. The particle size of these efficiency and therefore a good sinterability of the freestanding nanowhiskers, the pressed pellets made by NKBT nanowires starting materials is in the micrometer or submicrometer range. showed >98% theoretical density at 1100°C for 2 h. The solA perovskite phase-forming temperature of 800°C and above is gel-hydrothermal-derived ceramics have typical characteristics needed so that the components of the mixture have sufficient thermal energy to overcome the atomic/ionic diffusion barriers of relaxor ferroelectrics, and the piezoelectric properties were for the reaction. In most cases, the resulting powders are not better than the ceramics prepared by the conventional sol-gel nanocrystalline and undergo drastic agglomeration and yield an and solid-state reaction. inhomogeneous particle size as a result of the high-temperature treatment; therefore, they are unsuitable for enhancing the dielectric and piezoelectric properties of ceramics for high-perforI. Introduction mance uses. In comparison with a conventional solid-state reaction, the sol-gel process has some advantages, including exT present, lead zirconate titanate (PZT)-based ceramics are cellent compositional control, and homogeneity at the molecular the most widely used in electronic devices due to their high level due to the mixing of liquid precursors. However, sol-gel piezoelectric performance. However, the pollutant of toxic lead was regarded as a solid rather than solution process because solduring the fabrication and waste of products cause a crucial gel-derived precipitates are amorphous in nature and calcinaenvironmental problem. Therefore, there is an increasing intertions in air are inevitable for the formation of the crystalline est in developing lead-free piezoelectric ceramics to replace material. The sol-gel-hydrothermal processing represents an alPZT-based ceramics and to minimize lead pollution. It is well ternative to the calcinations for the crystallization of an objecknown that covalency between unoccupied states of the Pb6d in tive compound under mild temperatures. As a novel method to the perovskite structure, and O p states favor ferroelectric prepare oxide powders, the sol-gel-hydrothermal technique has ground states.' Compared with Pb2+, Bi3+ ions were in an the double advantage of both sol-gel and hydrothermal syntheisoelectronic state and also showed a long pair effect, which ses and has become attractive in the last decade due to its high encouraged studies of Nao.sBio.5Ti03(NBT) as an alternative degree of crystallinity, well-controlled morphology, high purity, to PZT ceramics. and narrow particle size distribution of the prepared powSodium bismuth titanate, NBT, which was found by Smolenders.'.'' In our previous work, K,,5Bi,,5Ti03 (KBT) nanowires skii et al.: is a kind of perovskite-type relaxor ferroelectric with with good sinterability were successfully synthesized by the sola Curie temperature Tc= 320°C. At room temperature, it has a gel-hydrothermal technique, and the dielectric properties of the rhombohedra1 structure (a = 0.389 nm, rx = 89.6"), and shows a derived KBT ceramics were superior to that prepared by all the relatively large remanent polarization (Pr= 38 pC/cm2). Howother methods reported Thus, the direct generever, it is difficult to pole NBT due to the high coercive field ation of NKBT nanowires with good sinterability at a low tem(Ec = 7.3 kvjmm), making it difficult to obtain the desirable piperature is of considerable interest. ezoelectric properties. In addition, unlike PZT ceramics, NBT In the present work, the sol-gel-hydrothermal process was has no morphotropic phase boundary (MPB), which plays a very presented as a new route to produce NKBT nanowires at a important role in PZT ceramics. The electromechanical propertemperature below 200"C, which is comparatively lower than ties show a maximum over a compositional range around that synthesized by the normal sol-gel route, which requires a the MPB of PZT, which can be attributed to an increase in temperature at least 650°C. The nanowires were characterized by X-ray diffraction (XRD), FT-IR, Raman, and transmission electron microscope (TEM) analysis. The densification behdvJ. Ninwontrihuting editor ior, the final microstructure of the sintered material, and the electric properties of the densified disks were investigated in detail. To the best of our knowledge, this is the first time that the Manuscript No. 22459. Received November 8, 2006; approved January 29, 2007. synthesis of NKBT nanowires has been reported, and the elecThis work was supported by the National Natural Science Foundation of China (Grant tric properties of sol-gel-hydrothermal-derived ceramics were No. 60601020) and the Natural Science Foundation of Beijing (Grant No. 4072006). 'Author to whom correspondence should be addressed. e-mail ydhou@bjut,edu.cn superior to that prepared by other methods.

A

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11. Experimental Procedure

The raw materials used in the present work were analyticalgrade bismuth nitrate pentahydrate (Bi(N03)2.5H20), sodium nitrate (NaN03), potassium nitrate (KNO,), tetrabutyl titanate (Ti(OC4H9)4)racetic acid (CH3COOH), and ethanol (CH,CH20H). Firstly, bismuth nitrate pentahydrate was dissolved in acetic acid; potassium nitrate and sodium nitrate were dissolved in C02-free distilled water, respectively. The mixture was then introduced into a prepared solution of a stoichiometric amount of tetrabutyl titanate in ethanol. After stirring vigorously for 2 h, a thin yellow homogeneous sol was formed. Then, the sol was heated at 80°C for 12 h to prepare a dry gel. The obtained gel precursor was added to a NaOH solution to form a suspension. The initial concentration of the NaOH solution was 10M. The as-prepared mixture was poured into a Teflon vessel, and then subjected to hydrothermal treatment at an appropriate temperature under auto-generated pressure for 48 h. After cooling, the product was filtered, washed with distilled water, and dried at ambient temperature. To compare the results with the conventional sol-gel process, a part of the gel was calcined at different temperatures from 300" to 800°C for 2 h in air. The crystal phase of the powders was determined using X-ray diffractometry (XRD; Model D8 Advance, Bruker AXS, Karlsruhe, Germany) in the 8-28 mode with graphitemonochromatized CuKu radiation (h = 0.154178 nm). The powder morphology was observed on a transmission electron microscope (TEM; Model JEM-2000 F, JEOL, Tokyo, Japan) and scanning electron microscope (SEM; Model S-3500N, Hitachi, Tokyo, Japan). Fourier transform infrared absorption spectra of the powders were obtained using an FT-IR apparatus (Model NEXUS670, Nicolet, Madison, WI). Raman scattering spectra of powders were recorded at room temperature from a Raman spectrometer (Model T64000, Jobin-Yvon, Paris, France) under backscattering geometry. Excitation was taken as the 488 nm line of an Art laser with a 50 mW output power. For the sintering experiments, the obtained powders were pressed into pellets with a diameter of 12 mm under an isostatic pressure of 150 MPa. Conventional sintering was performed at 50°C temperature intervals between 1000" and 1200°C for 2 h in a sealed alumina crucible. The bulk densities of the sintered pellets were measured by the Archimedes method. The microstructure of the sintered pellets was observed using SEM; (Model S-3500N, Hitachi, Tokyo, Japan) on the fracture side and free-top surfaces of the pellets. When the fracture surface did not reveal clear grain boundaries, the polished and thermally etched surface was observed. To measure the electrical properties, silver paste was coated on both sides of the sintered pellets and fired at 560°C for 30 min to form electrodes. The dielectric property and its dependence on temperature were measured using a precision LCR meter (Agilent 4284A, Agilent Technologies Inc., Palo Alto, CA) with an automated temperature controller. Before testing the piezoelectric properties, the specimens were poled in a silicone oil bath at 120°C by applying a dc field of 5 kV/mm for 30 min and aged for 24 h. The piezoelectric coefficient (&) was measured using a quasi-static piezoelectric 4, meter (Model ZJ3D, Institute of Acoustics, Chinese Academy of Sciences, Beijing, China). The electromechanical coupling factor (kp) and the mechanical quality factor (Q,,,)were estimated by the resonance and anti-resonance technique using an impedance analyzer (Agilent 4294A, Agilent Technologies Inc.).

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2 Theta (Degree) Fig. 1. X-ray

diffraction (XRD) patterns of powders hydrothermally treated at different temperatures: , perovskite. Inset: Fine scanning XRD patterns in 28=37"-48" for N%5Bio5TiO, (NBT), (Nao *KO2)o ,Bi, 5Ti03 NKBT, and KO5Bio 5Ti0, (KBT) powders hydrothermally treated at 160°C. to 48" for NBT, NKBT, and KBT powders hydrothermally treated at 160°C. It can clearly be seen that (1 11) and (200) peak positions of NKBT shift to high degrees compared with that of pure KBT, which can be attributed to the larger radius of K+ (1.33 A) than that of Na+ (1.02 A) in the A site of the AB03 s t r ~ c t u r eTo . ~ compare the results with the conventional sol-gel process, a part of the gel was calcined at different temperatures from 300" to 800°C for 2 h in air, and the XRD results are shown in Fig. 2. As can be seen in the figure, an amorphous phase was formed at a temperature below 400°C. When the temperature was increased to 450"C, some diffraction peaks corresponding to the Bi2Ti207pyrochlore phase appeared. However, for the powders heated at S O T , an obvious change was observed in the XRD patterns. The characteristic peaks of the NKBT perovskite phase appeared. After further increasing the temperature to 650°C and above, only a pure NKBT phase could be observed and there was no evidence of a second phase. The above results revealed that a temperature as high as 650°C is needed for the transformation from the pyrochlore to perovskite phase in a conventional sol-gel process, which is about 500°C higher than that required for the sol-gel-hydrothermal technique. Figures 3(a) and (b) show the images of NKBT samples synthesized by the sol-gel-hydrothermal and conventional sol-gel

111. Results and Discussion

Figure 1 shows the XRD patterns of powders hydrothermally treated at different temperatures. As can be seen from the XRD pattern for the sample synthesized at IWC, peaks corresponding to the perovskite phase had begun to appear, but the peaks were ill defined, which was indicative of the low crystallinity of this phase. Well-crystallized phases of NKBT were obtained for the samples hydrothermally treated at 160°C and above. The inset in Fig. I shows the fine scanning XRD patterns in 28 = 37"

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, 1

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2 Theta (Degree) Fig. 2. X-ray diffraction patterns of the dried gel heated at different

temperatures

m , perovskite, 0 , pyrochlore

Progress in Nanotechnology: Processing

2000

1500

1000

Wavenumber (cm-'>

500

Fig.4. Infrared spectra of the dried gel, and (Nao.&&sBio (NKBT) nanowires hydrothermally treated at 160°C.

Fig. 3. (a) Transmission electron spectroscopy (TEM) photograph of (Nan & 2)0 5BioSTi0, (NKBT) powders synthesized by the sol-gelhydrothermal method at 160°C. Inset: A typical NKBT nanowire and its SAED pattern. (b) Scanning electron microscope photograph of NKBT powders synthesized by the conventional sol-gel method at 650°C. Inset: High-magnification TEM image of powders.

method, respectively. From Fig. 3(a), it can clearly be observed that the sample synthesized at 160°C by the sol-gel-hydrotherma1 process demonstrates the morphology of nano-sized wirelike fibers, which are monodispersed and are not fused to one another. Furthermore, each nanowire is uniform in width along its entire length, with diameters of 5&80 nm and lengths of 1.5-2 pm. The inset in Fig. 3(a) shows the SAED pattern recorded from an isolated nanowire. The diffraction spots have been well indexed to the planes of NKBT, confirming the formation of a Nanotubes, Nanorods, and Nanowires

STiO3

single-crystal NKBT nanowire. This sol-gel-hydrothermal method for preparing NKBT nanowires was highly reproducible, and in all cases, similar morphologies were obtained. Compared with nanowires synthesized by the sol-gel-hydrothermal process, the powders synthesized at 650°C by the conventional sol-gel method consist of spherical particles with an average particle size of about I pm, as can be seen in Fig. 3(b). However, the high-magnification TEM micrograph in the inset of Fig. 3(b) highlights the particulate morphology of the crystallized NKBT particles. It can be seen that large aggregated particles consisting of many smaller primary particles of about 2 M nm are formed. This particulate morphology can be attributed to the porous nature of the as-prepared xerogel precursor, which minimized the interparticle contact by the open structure, thereby resulting in the formation of porous agglomerated powders in calcination. In order to further analyze the structure changes during the synthetic process, the FT-IR spectra of the gel and the NKBT nanowires hydrothermally treated at 160°C were obtained, respectively, and the results are shown in Fig. 4. The bands at approximately 1630 cm-l on the IR spectrum of the gel can be attributed to the bending mode of H-O-H. The peak at 1541 cm-' can be attributed to COO vibration. The sharp and intensive peak at 1390 cm-' is due to the presence of nitrate, while the weak peaks at around 1037, 920, and 830 cm-l can be attributed to CO vibration. The bands in the low-wavenumber region (45&650 cm-I) can be attributed to Ti-0 bond vibrations. After hydrothermal treatment at 160"C, nitrate and carbon peaks were reduced significantly, but a large band appeared around 600 cm-l, which can be attributed to the characteristic vibration of T i 4 octahedron and indicates the formation of the perovskite phase.13 Raman scattering is known to be an appropriate technique for the investigation of the short-range order and phase structure in perovskites. For comparison, pure KBT has been synthesized at 160°C by the sol-gel-hydrothermal method and the room temperature Raman scattering experiments of both KBT and NKBT have been investigated and the results are shown in Fig. 5. For pure KBT, there are four obvious peaks appearing, in an ascending order, at around 273, 330, 529, and 632 cm-' in the Raman shift ran e of 200-700 cm-', which is similar to already reported data.g4 For NKBT, it can be seen that the dominant peak at 273 cm-' roughly maintains its frequency position. This observation, together with the fact that a strong peak at essentially the same position has been observed in the Raman 307

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Fig. 7. X-ray diffraction patterns of the sol-gel-hydrothermal-derived (Nao.xKo.2)o.sBio.sTi03 ceramics sintered at different temperatures: , perovskite; 0 , pyrochlore.

700

Raman shift (cm-'1 Fig. 5. Raman spectra of G.5BiO.5Ti07(KBT) and (N+.~&.2)o,sBio.~TiO~ (NKBT) nanowires hydrothermally treated at 160°C.

spectra of BaTi03 and PbTi03, is an indication that this band must be dominated by T i 4 vibrations. However, compared with pure KBT, the Raman spectra of NKBT show quite a broad feature. The radii of Na+ and K+ areoquite different, r(Na+) = 1.02 A compared with r(K+) = 1.33 A. The decrease of the ionic radii, when going from K+ to Na+, leads naturally to a distortion of the structural framework. On the other hand, the incorporation of NBT and KBT leads to a cation disorder (Bi, K, Na) on the 12-fold coordinated site, which results in the broad feature of NKBT Raman spectra due to the overlapping of Raman modes. In the present work, NKBT nanowires, which were not obtained by other methods, were prepared by a simple sol-gelhydrothermal method. Owing to an extremely small size and anisotropy, the control of nucleation and growth of one-dimensional nanostructure materials is a big challenge. In our work, it is believed that the gel precursor and hydrothermal environment play a key role in the formation of the nanowires. It is known that the growth of nuclei in the calcinations of the gels was controlled by the short-range diffusion of ions in a limited space; therefore, it is difficult to control the morphology of the final products. In the hydrothermal process, the initial condition of

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Temperature ("C) Fig. 6. Relationship between the relative density and the sintering temperature of the (Nao.xKoz)o.sBio.sTi03ceramics prepared by the sol-gel and sol-gel-hydrothermal methods.

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the reactant exerts an impact on the crystal nucleation and growth, which is responsible for the morphology of the product. It is supposed that chained nanoclusters with a well-defined morphology are formed during the gelation of an NKBT sol by controlling the hydrolysis of the sol. These chained nanoclusters could serve as the nuclei of NKBT during hydrothermal treatment and gradually grow freely in an aqueous solution to form the nanowires. One-dimensional wire-like nanostructures have received considerable attention in recent years for providing a good system to study the dependence of properties, such as nonlinear optical phenomena, thermal transport, electrical, and mechanical properties on dimensionality and size r e d ~ c t i o n . ' ~It' ~is important to note that nanowires also show superior sinterability, which aided in obtaining a well-densified single-phase ceramic.' However, until now, work related to the ceramics derived from nanowires has rarely been reported. Thus, it is of great significance to investigate the sintering as well as electrical properties of the pressed pellets made by NKBT nanowires. Figure 6 gives the relative density as a function of sintering temperature for the NKBT ceramics prepared by sol-gel and sol-gel-hydrothermal methods, respectively. It can be seen that the pressed pellets made by NKBT nanowires showed >98% theoretical density at 1100°C for 2 h. However, the sol-gel-derived NKBT ceramics showed only 93% theoretical density at the same temperature. It is well-known that the relative density of a sintered ceramic is directly related to the green density of the pressed pellets, which in turn is highly de endent on the morphology of the precursor oxide powders.2032PAscan be seen in Fig. 3, the powder synthesized by the sol-gel-hydrothermal method yields free-standing nanowires of NKBT. The free-standing nanowires have a better packing efficiency and therefore a high green density of about 60% of the theoretical density. In comparison, the powder synthesized by the sol-gel method yielded a porous-structured nanopowder of NKBT, which does not have a high packing efficiency and is responsible for the low green density of about 50% of the theoretical density. Figure 7 shows the XRD patterns of the sol-gel-hydrothermal-derived NKBT ceramics sintered at different temperatures. At 1000"-1 IOWC, all peaks can be indexed to the pure NKBT perovskite phase. However, when the sintering temperature reached up to 115O"C, peaks corresponding to the Bi2Ti~07 pyrochlore phase were observed and peak intensities of Bi2Ti207 became stronger with a further increase in sintering temperature. The appearance of the Bi2TiZO7phase indicated that NKBT was decomposed due to the high volatilization of the potassium element at a high temperature, and the latter could be responsible

'

Progress in Nanotechnology: Processing

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Fig. 9. Temperature dependence of dielectric constant at various frequencies for the sol-gel-hydrothermal-derived (Nao.8Ko.2)o.sBio.STi03 (NKBT) ceramics sintered at 1100°C.

Fig.8. Scanning electron microscope (SEM) micrographs of the freetop surface of (Nao.8Ko.2)o.5Bio.5TiOz (NKBT) ceramics prepared by the sol-gel method at (a) 1100°Cand the sol-gel-hydrothermal method at (b) I 100°C and (c) 1200°C. Inset: SEM micrographs of the polished and thermally etched cross section of NKBT ceramics.

for the decrease of the relative density at a high temperature above 1150°C. Strong differences in the microstructure of the NKBT ceramics prepared by sol-gel and sol-gel-hydrothermal methods, respectively, are evidenced by the SEM photographs, as shown in Figs. 8(aHc). It can be seen from Fig. 8(a) that the visible intergranular porosity, irregular polyhedral shape of the grains, and inhomogeneous grain size distribution clearly evidence the lower degree of densification of the sol-gel-derived specimen sintered at 1100°C. Comparatively, the sol-gel-hydrothermalderived specimen sintered at 1100°C showed a dense and homogenous microstructure, as can be seen in Fig. 8(b). This agrees well with the relative density of 98% measured for the material. In addition, it should be noted that the sol-gel-hydrothermal-derived specimen sintered at 1200°C is less homogenous and the surface micrograph shows some evidences of melting grain, as can be seen in Fig. 8(c). A similar feature of the formation of amor hous phase has also been observed in a PZTbased ceramic.2PLeaching the obtained specimens with dilute nitric acid completely removes the amorphous phase. The origin Nanotubes, Nanorods, and Nanowires

of the amorphous phase can be attributed to vaporizing of the potassium component during the sintering process, and a further study has been carried out to investigate the detailed mechanism. Owing to the high density and homogenous microstructure, the sol-gel-hydrothermal-derived NKBT ceramics sintered at I100"C had been selected to investigate the dielectric and piezoelectric properties. Figure 9 shows the temperature dependence of dielectric constant at various frequencies for the sol-gelhydrothermal-derived NKBT ceramics sintered at 1100°C. A frequency dependence and a hump were observed between room temperature and 180°C ( T , , corresponding to a change from a ferroelectric phase to an anti-ferroelectric phase),23and a diffusion of the dielectric constant at temperatures higher than 320°C (T,, corresponding to a transition from an anti-ferroelectric phase to a paraelectric phase) was also found.8 As N a f, K+, and Bi3+ are randomly distributed in the 12-fold coordination crystallographic sites, the diffuse behavior of NKBT ceramics could reasonably be attributed to the compositional fluctuation and structural disorder in the arrangements of these cations at the A site. On the basis of the above results, it is obvious that for a temperature higher than T,, the variation of the dielectric constant does not follow the classical Curie-Weiss law. Whatever the measurement frequency, the dielectric constant varies according to the Uchino and Nomura function,24

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In (T-TJ Fig. 10. In (l/sr-l/&max) as a function of In (T-Tmax)at 10 kHz for the sol-gel-hydrothermal-derived (N~.8Ko.2)o.sBio.sTi03 (NKBT) ceramics sintered at 1 100°C.

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Table I. Comparison of the Dielectric and Piezoelectric Properties of NKBT Ceramics Prepared by Different Methods Method

Sol-gel-hydrothermal method Sol-gel method Solid-state reaction Solid-state reaction Solid-state reaction

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This work Sasaki et aL6 Jidng et 01.’ Li et

1030 1200 760

I/IZ-I/E,,,~~ = (T-TmaJ7/C,with y close to 1.51 as clearly shown in Fig. 10 for 10 kHz. In this relation, y acts as a diffusion coefficient and ranges between 1 (normal ferroelectric) and 2 (a complete diffuse phase transition (DPT) ferroelectric). This behavior would be significant of a relaxation process due to an Asite cation mixture. The dielectric and piezoelectric properties of the NKBT ceramics obtained by our sol-gel-hydrothermal methods are compared in Table I with those of ceramics prepared by other methods, including sol-gel and solid-state reactions. It can be found that excellent electrical properties (d33= 156 pC/N, k, = 0.35, Qm = 165, E, = 1220, and tan 6 = 0.022 at room temperature) are obtained for the sol-gel-hydrothermal-derived NKBT ceramics, which were superior to those of previously reported ceramics obtained by other m e t h o d ~ . 6 >Besides ~ . ~ ~ this, it should be noted that the sol-gel-hydrothermal-derived NKBT ceramics retained a high Td value of 180°C. It is known that Td is the depolarization temperature, which is an important factor for BNT-based piezoelectric ceramics in view of their practical use.26The enhanced electric properties in the sol-gel-hydrothermal-derived ceramics can be attributed to the good sinterability of nanowhiskers, which resulted in an increase in density and a more homogeneous microstructure.

IV. Conclusion Free-standing nanowires of NKBT have been synthesized successfully by using the sol-gel-hydrothermal method. In the solgel-hydrothermal process, the hydrothermal conditions create a gentle environment to promote the formation of crystalline NKBT nanowires at a very low processing temperature of 160”C, which reduced the synthesized temperature for the NKBT perovskite phase in the conventional sol-gel process by about 500°C. Owing to the good sinterability of nanowires prepared by the sol-gel-hydrothermal route, the NKBT ceramics with a relative density of 98% are easy to prepare at 1100°C for 2 h and the dielectric and piezoelectric properties were superior to that prepared by the conventional sol-gel route and solid-state reactions. The sol-gel-hydrothermal route, without the presence of catalysts and requiring no expensive equipment, will ensure higher purity in the products and greatly reduce the production cost, and thus offer a novel and simple synthetic route for onedimensional nanoscale materials and high-quality ceramics.

References ‘R. E. Cohen, “Origin of Ferroelectricity in Perovskite Oxides,” Nafure, 358, 136-8 (1992). ’G. A. Smolenskii, V. A. lsupov, A. 1. Agranovskaya, and N. N. Krainik, “New Ferroelectrics of Complex Composition,” Sov. Phys. Solid Stare, 2, 265 1-4 (1961). ‘H. Ishii. H. Nagata, and T. Takenaka, “Morphotropic Phase Boundary and Electrical Properties of Bisumuth Sodium Titanate-Potassium Niobate Solid-Solution Ceramics,” Jpn. J. Appl., 40 [Part I] 566C-3 (2001).

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4B. I. Chu, D. R. Chen, G. R. Li. and Q. R. Yin, “Electrical Properties of Na,, 2Bi112Ti03-BaTi03Ceramics.” J. Eur. Ceram. Soc., 22, 21 15-21 (2002). 5S. C. Zhao, G. R. Li, A. L. Ding, T. B. Wang, and Q. R. Yin, “Ferroelectric and Piezoelectric Properties of (Na, K)o 5BiosTi03 Lead Free Ceramics,” J. Phys. D, 39, 2277-81 (2006). ‘A. Sasaki. T. Chiba. Y. Mamiva. and E. Otsuki, “Dielectric and Piezoelectric Properties of (BiosNao.s)TiO~-(Bi~.S~.s)TiO, Systems,” Jpn. J . Appl. Phys. Part I , 38,5564-7 (1999). ’Y. M. Li, W. Chen, J. Zhou. Q. Xu,H. I. Sun, and M. S. Liao, “Dielectric and Ferroelectric Properties of Lead-Free Nao.5Bio5Ti03-K0.5Bi0.5Ti03Ferroelectric Ceramics,” Ceram. fnr., 31, 13942 (2005). *X.P. Jiang, L. Z. Li, M. Zeng, and H. L. W. Chan, “Dielectric Properties of Mn-Doped (Nao8Ko,~)0.5Bi0.sTiO~ Ceramics.” Marer. Lerf.. 60, 178690 (2006). 9Z. J. Li, B. Hou, Y. Xu, D. Wu, Y. H.Sun, W. Hu, and F. Deng, “Comparative Study of Sol-Gel-Hydrothermal and Sol-Gel Synthesis of Titania- Silica Composite Nanoparticles,” J. Solid Stare Chem., 178, I 3 9 W 5 (2005). ‘OH. Wang, L. Wang, J. B. Liu, B. Wang, and H. Yan, “Preparation of Phl_,LayTil-r,403(x = 0.28) Powders by a Sol-Gel-Hydrothermal Method,” Mater. Sci. Eng. B, 99, 495-8 (2003). “L. Hou, Y. D. Hou, X. M. Song, M. K. Zhu, H.Wang, and H. Yan, “SolGel-Hydrothermal Synthesis and Sintering of KO5Bio5TiOl Nanowires,” Marer. Res. Bull., 41, l 3 3 M (2006). ”Y. D. Hou, M. K. Zhu, L. Hou. J. B. Liu, J. L. Tang. H. Wang, and H. Yan, “Synthesis and Characterization of Lead-Free KOsBio5TiOl Ferroelectrics by Sol-Gel Technique.” J. Cryst. Growth, 273, 500-3 (2005). I3C.Y. Kim, T. Sekino. and K. Niihara. “Synthesis of Bismuth Sodium Titanate Nanosized Powders by Solution/SolCel Process,” J . Am. Ceram. Soc., 86. 1464-7 (2003). I4J. Kreisel, A. M. Glazer, G. Jones, P. A. Thomas, L. Abello, and G . Lucazeau, “An X-Ray Diffraction and Raman Spectroscopy Investigation of A-Site Substituted Perovskite Compounds: The (Nal-,K.Jo 5Bio.sTi03( 0 5 x 5 I) Solid Solution,” J . Phys.: Condens. Marrer. 12, 3267-80 (2000). ”H. J. Dai, E. W. Wong, Y. Z. Lu, S. S. Fan, and C. M. Lieber, “Synthesis and Characterization of Carbide Nanorods,” Narure. 375, 769-72 (1995). I6L. A. Bumm. J. J. Arnold, M. T. Cygan. T. D. Dunbar, T. P. Burgin. L. Jones 11, D. L. Allara, J. M. Tour, and P. S. Weiss, “Are Single Molecular Wires Conducting,” Science, 271, 1705-7 (1996). ”J. D. Holmes, K. P. Johnston, R. C. Doty, and B. A. Korgel, “Control of Thickness and Orientation of Solution-Grown Silicon Nanowires,” Science, 287, 1471-3 (2000). “J. J. Urban, J. E. Spanier, L. Ouydng. W. S. Yun, and H. Park, “Single Crystalline Barium Titanate Nanowires,” Adv. Mater., 15, 423-6 (2003). ”G. Feng. W. D. Nix, Y. Yoon, and C. J. Lee, “A Study of the Mechanical Properties of Nanowires Using Nanoindentation.” J . Appl. P h j x , 99,074304-1-10 (2006). 2oY, D. Hou, L. Hou. M. K. Zhu, and H. Yan, “Synthesis of (&,sBi,5)04. B% 6Ti03 Nanowires and Ceramics by Sol-Gel-Hydrothermal Method, Appl. Ph s Lerr.. 89, 243114-1-3 (2006). ‘A. Banerjee and S. Bose, “Free-Standing Lead Zirconate Titanate Nanoparticles: Low-Temperature Synthesis and Densification,” Chem. Marer., 16, 561&5 (2004). 22Y.D. Hou, M. K. Zhu, H. Wang, B. Wang, C. S. Tian. and H. Yan, “Effects of Atmospheric Powder on Microstructure and Piezoelectric Properties of PMZNPZT Quaternary Ceramics,” J. Eur. Ceram. Soc., 24. 3731-7 (2004). ”Y. Hiruma, H. Nagata, and T. Takenakd, “Phase Transition Temperatures and Piezoelectric Properties of (Bil~2Nal,2)Ti03-(Bili2Kli2)TiO~-BaTi03 LeadFree Piezoelectric Ceramics,” Jpn. J . Appl. Phys.. 45 [Part I] 7409-12 (2006). 24K.Uchino and S. Nomura, “Critical Exponents of the Dielectric Constants in Diffused-Phase-Transition Crystals,” Ferrwlecrr. Lerf. Sect., 44,55-61 (1982). ”Y. M. Li, W. Chen, Q. Xu, J. Zhou, and M. S. Liao, “Dielectric and Piezoelectric Properties of (Nao 8&,.2)o sBioSTi03Ceramics,” J. Inorg. Marer., 19, 81722 (2004). 26X, X. Wang, X. G. Tang, and H. L. W. Chan, “Electromechanical and FerLead-Free Piroelectric Properties of (Bil,2Nal,2)Ti03-(Bil~2Kl,2)Ti03-BaTi03 ezoelectric Ceramics.” Appl. Phys. Letr., 85, 91-3 (2004). n

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Synthesis and Characterization of Ce, --xCdx02-6Nanorods Jong So0 Leet and Sangtae Kim* Department of Chemical Engineering and Materials Science, University of California, Davis, California 956 16

Single crystalline Cel_,Gd,02-g nanorods with x = 0.05 and 0.1 as well as x = 0 were synthesized by a sol-gel process without using any template and/or applying high pressure. The nanorods grow preferentially along the [ l l o ] direction regardless of the dopant concentration. The diameter and the length of the synthesized nanorods range from 10 to 50 nm and from 200 to 700 nm, respectively. The length of the nanorods decreases with increasing dopant concentration. The role of the surfactant as a structuredirecting agent is critical to the growth of Ce, -,GdX02-g nanorods.

I. Introduction

C

ERIUM dioxide (Ce02) is a fluorite-structured nonstoichiomeric oxide with a greater structural tolerance to reduction. Its electrical property as well as oxygen storage capacity have been the center of attention for decades because of the potential for various applications such as exhaust catalysts, oxygen sensors, and, in particular, a solid electrolyte (SE) for solid oxide fuel cells (SOFCS).'-~In recent years, the oxygen ion conductivit of nano-structured Ce02 doped with acceptors such as GdY+ has been of particular interest as the nano-structured CeOz may serve as an SE for low-temperature S O F C S . This ~~ hypothesis is based on the assumption that the large surface/ interface area in the nanomaterials may serve as fast conduction pathways leading to high conductivity. Hence, the preparations as well as the oxygen ionic conductivity of Gd-doped Ce02 in the forms of nanocrystalline (crystallite size < 50 nm), thin films (two dimensional (2 D)), and ceramics (three dimensional (3 D)) have been extensively e ~ p l o r e d . ~Unlike other oxide conductors,' however, little attention has been paid to the synthesis of one dimensional (1 D) single crystalline nano-structured (e.g., nanorods) CeO2 and thus to their electrical properties, although they can serve as ideal model systems to study the size effects on their physical properties. In particular, preparation of 1 D Gddoped Ce02 has not been reported to date, while few results have described the synthesis of 1 D nominally pure Ce02 that are practically less relevant as far as the SOFC applications are concerned.*-I2 In this communication, we report a simple large-scale synthesis of single crystalline Cel-,Gd,02-6 (CGO) nanorods with x = 0.05 (CGO-5) and 0.1 (CGO-10) as well as x = 0 (CEO), based on sol-gel routes without using any template and/or applying high pressure typically needed for the synthesis of the 1 D Ce02 reported previously.8-12 J.-H. Lee--contnhuting editor

Manuscript No. 22076. Received August 12, 2006; approved October 16, 2006. J.S.L. acknowledges a partial support from the Korea Research Foundation (KRF2005-214-DOO305) for this work. 'Author to whom correspondence should be addressed. e-mail [email protected] 'Member, American Ceramic Society.

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11. Experimental Procedure

For the synthesis of 1-D CGO of interest, 30 mL of aqueous sulfuric acid solution (pH 2) containing 1.52 g of CeC13.7H20 (Aldrich, St. Louis, MO; 99.9%), together with appropriate amounts of Gd(N03).6H20 (Aldrich, 99.99%) and 5 g of cetyltrimethyl ammonium bromide (C19H42N. Br, CTAB, Aldrich, 99%), was prepared and stirred for l h at 60°C. Seventeen milliliters of aqueous ammonia (Aldrich, 28-30 wt%) was then added to the solution in a dropwise manner. The addition results in a change in the pH value of the solution from about 2 to about 11, leading to precipitation. The color of the solution changes from colorless to dark purple accordingly. After stirring for another hour, the glass beaker containing the mixture was sealed with an aluminum foil to prevent possible contaminations and was placed in a dry oven at 80°C for 36 h. The product was then separated from the solution using a centrifuge and washed multiple times with pure ethanol and deionized water. The collected product was dried at 60°C for 12 h in a dry oven. Finally, for crystallization, the product was calcinated at 550°C in N2 for 6 h and then in air for 2 h. The structure analysis of the as-calcinated product was performed using an X-ray diffraction (XRD) spectrometer (Scintag XDS-2000, Cupertino, CA) with CuKa radiation, a transmission electron microscope (TEM, Philips CM- 12, Hillsboro, OR), a high-resolution TEM (HRTEM, Tecnai F-20, Hillsboro, OR), and an energydispersive X-ray spectrometer (EDX, Gatan, Pleasanton, CA). 111.

Results and Discussion

Figure 1 shows the XRD patterns of the synthesized CGO-10 (a) and CGO-5 (b) samples. Also included in Fig. 1 is the synthesized CEO (c) sample for comparison. All the peaks revealed are indexed to confirm a fluorite structure. Neither secondary nor impurity phases are detected in the samples as can be seen in Fig. 1. Also shown in Fig. 1 is a peak (marked with a solid square) of a high-quality standard quartz (JCPDS, No. 33-1 161) used as a reference for the angular correction, which remains unshifted in all the samples. The XRD peaks of the CGO samples are found to be slightly shifted to the lower diffraction angle with increasing Gd content while those of the CEO sample remained nearly unshifted compared with those of the bulk Ce02 with a lattice constant of a = 0.541 nm (JCPDS, No. 43-1002). The shifts of the (220) and (311) peaks of the CGO-5 and CGO-10 are A(20) = 0.02, 0.04" and 0.03, 0.18", respectively, indicating that the lattice parameter of the CGO is increased with increasing Gd content relative to that of CEO. Such an increase in the lattice parameter can beattributed to the slightly larger radius of the dopan: Gd3+(1.05 A) compared with that of the host ion Ce4' (0.97 A) in the 1atti~e.I~ This result strongly supports the fact that well-crystallized CGO with the dopant uniformly distributed into the lattice has been synthesized. Figure 2 shows a TEM image (a) and its associated EDX spectrum (b), and an HRTEM image (c) of the CGO-5 sample. Figure 2(a) clearly demonstrates that the synthesized CGO-5 is in the form of nanorods with a diameter (d)and length (I) of 311

Fig. 3. Transmission electron microscope (TEM) (a) and its associated energy dispersive X-ray spectrometer spectrum (b), and high-resolution TEM image (c) of a 10 mol% Gd-doped Ce02 nanorods. Fig. 1. X-ray diffraction patterns of as-synthesized 10 mol% (a), 5 mol% (b) Gd-doped Ce02. and nominally pure C e 0 2 (c) samples.

about 1&50 nm and about 100-400 nm, respectively. A small amount of the nanopowder of CGO-5 (d- 10 nm) was found to coexist with the nanorods (see Fig. 2(a)). The EDX spectrum of Fig. 2(b) reveals the peaks associated only with Ce, Gd, and 0 (the Cu-related peaks come from the Cu grid). The relative atomic percent of Gd in the nanorod estimated based on the EDX data is about 6.8%, leading to the fact that the exact chemical composition of CGO-5 nanorods is Ceo.932Gdo.06802-s. The selected area diffraction (SEAD) patterns of one of the nanorods shown in the bottom inset of Fig. 2(a) indicate that the nanorod is a single crystalline phase with growth along the [110] direction. The HRTEM image (Fig. 2(c)) confirms that the CGO-5 nanorod is a structurally uniform single crystal with a regular periodicity of the lattice. The interplannar distance between the adjacent lattice fringes calculated from Fig. 2(c) is about 0.31 nm (see the inset), corresponding to that of the (1 11) lattice plane of the bulk Ce02. These results support the fact that the CGO-5 nanorod grows preferentially along the [I 101 direction. Figure 3 shows the TEM image (a), EDX spectrum (b), and HRTEM image (c) of the synthesized CGO-I0 sample. Similar to CGO-5, CGO-I0 was found to form nanorods coexisting with its nanopower (d< 10 nm, see in Fig. 3(a)). The volume fraction of the nanopowder is, however, found to be slightly higher than that in the CGO-5 sample. The CGO-10 nanorods are about lO(r200 nm long and are about 30-50 nm wide, indicating that those nanorods are shorter and thicker compared

with the CGO-5 nanorods shown in Fig. 2(a). Both the CGO-5 and the CGO-10 nanorods did not grow further even at a longer reaction time. This may imply that the length of the CGO nanorods and the dopant concentration are interrelated. Indeed, CEO nanorods reveal the smallest aspect ratio among the nanorods of interest in this study. The SEAD patterns of a selected single nanorod (see the bottom insets of Fig. 3(a)) of CGO-10, together with the EDX spectrum (Fig. 3(b)), indicate that the synthesized nanorod is a structurally uniform single-crystal CGO with growth along the [I 101 direction, consistent with the CGO-5 nanorods. This is also confirmed by Fig. 3(c) from which the interplannar distance between the adjacent lattice fringes is calculated to be about 0.19 nm, corresponding to that of the (220) lattice plane of the bulk CeO2. The exact chemical composition of the CGO-10 nanorods calculated from the EDX data was c % ssGdo 120z-s. For comparison, CEO nanorods (see the EDX spectrum shown in Fig. 4(b)) were also synthesized and their morphologies are shown in Fig. 4. Note that, unlike the CGO samples discussed above, the as-synthesized CEO sample consists of nearly 100% single crystalline nanorods that appear to be thinner (ca.l(r30 nm) and longer (ca. 200-700 nm) than the CGO nanorods. The CEO nanorods also grow along the [I 101 direction, which confirms the results reported previously.’ The interplannar distance between the adjacent lattice fringes calculated from Fig. 4(c) is about 0.27 nm, corresponding to that of the (200) lattice plane of the bulk Ce02. The cetyltrimethyl ammonium bromide (CTAB, see “Section II”), which serves as a structure-directing species and/or a reactant, has been used for the synthesis of various nano-

Fig. 2. Transmission electron microscope (TEM) (a) and its associated energy dispersive X-ray spectrometer spectrum (b). and high-resolution TEM image (c) of 5 mol% Gd-doped Ce02 nanorods.

Fig.4. Transmission electron microscope (TEM) (a) and its associated energy dispersive X-ray spectrometer spectrum (b), and high-resolution TEM image (c) of pure Ce02 nanorods.

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structure^.^^'^^'"'^ It was indeed found that only nanopowders with d<10 nm (see the inset in the upper right corner of Fig. 4(a)) are formed without using CTAB, indicating that the CTAB plays a critical role in the synthesis of CEO nanorods. A suggested growth mechanism of CGO nanorods is the following: cerium chloride is oxidized to form Ce(H20) (OH):-”)+ under basic medium at 60”C.x,12 The hydrated Ce4+ .hcorporates the ionic surfactant CTMA+ (C16H33(CH3)3N+) by exchange with the surface OH group when the pH value is higher than that of the isoelectric point (6.7S8) of the hydrous cerium This process leads to formation of colloidal ceria/surfactant mixtures, in which the ceria component is periodically arranged within the amorphous surfactant matrix.*’I2 Finally, the single crystalline Cel-,Gd,02-s nanorods are formed with removal of the surfactant species by calcinating the isolated colloidal ceria/ surfactant mixtures at 550°C. IV. Conclusions We have synthesized single crystalline 5 and 10 mol% Gd doped as well as nominally pure Ce02 nanorods using a simple sol-gel method. The growth of the nanorods is preferentially along the [I 101 direction regardless of the dopant concentration. The diameter and the length of the synthesized nanorods range from 10 to 50 nm and from 200 to 700 nm, respectively. The length of the nanorods decreases with increasing dopant concentration. This implies that the dopant may restrict the growth of the nanorods. The role of the surfactant (CTAB) as a structure-directing agent is critical to the growth of CGO nanorods. The single crystalline CGO nanorod will provide a unique opportunity as a model system to study the ion conduction property of the ceria electrolyte used in solid oxide fuel cells in a nanoscale regime at a fundamental level.

Nanotubes, Nanorods, and Nanowires

References ‘M. Mogensen, N. M. Sammes, and G . A. Tompsett, “Physical. Chemical and Electrochemical Properties of Pure and Doped Ceria,” Solid State lonics. 129.6394 (2000). ’B. C. H. Steele, “Ceramic Ion Conducting Membranes,” Curr. Opin. Solid Sfare Murrr. Sci.. 1 [5]684-91 (1996). ’N. Q. Minh, “Ceramic Fuel Cells,” J. Am. Cerum. Soc., 73 [3]56348 (1993). 4S. Kim and J. Maier, “On the Conductivity Mechanism of Nanocrystalline Ceria,” J. Electrochem. Soc., 149 [lo]573-83 (2002). ’A. Tschope, “Grain Size-Dependent Electrical Conductivity of Polycrystalline Cerium Oxide 11: Space Charge Model,” Solid Sfare lonics, 139 [34]26740 (2001). q. Suzuki, I. Kosacki, and H. U. Anderson, “Microstructure-Electrical Conductivity Relationship in Nanocrystalline Ceria Thin Films,” Solid Slate lonics, is; [4]I I 1-21 (2002). Z. R. Dai, Z. W. Pan. and Z. L. Wang, “Novel Nanostructures of Functional Oxides Synthesized by Thermal Evaporation,” Adv. Funct. Muter., 13 [I] 9-23

.----, ‘A. Vantomme, Z. Y. Yuan, G . Du, and B. L. Su, “Surfactant-Assisted Large-

tznnn

Scale Preparation of Crystalline Ce02 Nanorods,” Lungmuir, 21, 1132-5 (2005). %. Q.Han, L. Wu, and Y. Zhu, “Formation and Oxidation State of Ce02_, Nanotubes,” J. Am. Chem. Soc., 127, 12814-5 (2005). “B. Tang, L. Zhuo, J. Ge. G. Wang, Z. Shi, and J. Niu, “A Surfactant-Free Route to Single-Crystalline Ce02 Nanowires,” Chem. Commun., 3565-7 (2005). “K. Zhou, X. Wang, X. Sun, Q. Peng, and Y. Li. “Enhanced Catalytic Activity of Ceria Nanorods from Well-Defined Reactive Crystal Planes,” J . Caral., 229. 20612 (2005). 12D. Terribile, A. Trovarelli, J. Llorca, C. Leitenburg. and G . Dolcetti, “The Synthesis and Characterization of Mesoporous High-Surface Area Ceria Prepared Using a Hybrid Organic/lnorganic Route,” J. Carol., 178, 299-308 (1998). I3P. Li. 1. W. Chen, J. E. Penner-Hahn, and T. Y. Tien, “X-Ray Absorption Studies of Ceria With Trivalent Dopants,” J. Am. Ceram. Soc., 74 [S]95847 (200I). I4N. R. Jana, L. Gearheart, and C. J. Murphy, “Wet Chemical Synthesis of Silver Nanorods and Nanowires of Controllable Aspect Ratio,” Chem. Commun., 6174 (2001). ”N. R. Jana, L. Gearheart, and C. J. Murphy, “Wet Chemical Synthesis of High Aspect Ratio Cylindrical Gold Nanorods,” J. Phys. Chem. B, 105, 4065-7 (2001). I6M. Tornblom and U. Henriksson, “Effect of Solubilization of Aliphatic Hydrocarbons on Size and Shape of Rod like C16TABrMicelles Studied by 2H NMR Relaxation,” J . Phys. Chem. 8, 101, 6028 35 (1997). I7J. L. Blin, R. Flarnant, and B. L. Su, “Synthesis of Nanostructured Mesoporous Zirconia Using CTMABr-ZrOCl2. 8 H 2 0 Systems: A Kinetic Study of Synthesis Mechanism,” Int. J. Inorg. Marer., 3, 959-72 (2001). 0

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SYNTHESIS AND CHARACTERIZATION OF CUBIC SILICON CARBIDE (p-Sic) AND TRIGONAL SILICON NITRIDE (a-Si3N4) NAN0 WIRES Karine Saulig-Wenger, Mikhael Bechelany, David Cornu*, Samuel Bernard, Fernand Chassagneux and Philippe Miele Laboratoire des Multimatkriaux et Interfaces UMR 56 15 CNRS University Claude Bernard - Lyon I 43 Bd du 11 novembre 1918 F-69622 Villeurbanne Cedex, France David.Cornu@univ-lyon I .f? Thierry Epicier GEMPPM UMR 55 10 CNRS - INSA Lyon 20 Avenue Albert Einstein F-6962 1 Villeurbanne Cedex, France ABSTRACT By varying the final heating temperature in the range 1050°C - 1300"C,cubic silicon carbide ([MiC) and/or trigonal silicon nitride (a-Si3N4) nanowires (NWs) were prepared by direct thermal treatment under nitrogen, of commercial silicon powder and graphite. Long and highly curved p-Sic NWs were preferentially grown below 12OO0C, while straight and short a-Si3N4 NWs were formed above 1300°C.Between these two temperatures, a mixture of both nanowires was obtained. The structure and chemical composition of these nanostructures have been investigated by SEM, HRTEM, EDX and EELS. INTRODUCTION Numerous studies have been recently devoted to ceramic and metallic nanowires (NWs) due to their outstanding properties which can be tailored by varying their chemical composition and also their crystalline structure. The possible applications of these NWs run from nanoelectronics to composite materials. Among the series, silicon-based NWs, made of cubic silicon carbide @-Sic), silicon dioxide (Si02) or trigonal silicon nitride (a-Si3N4, are of particular interest due to their interesting mechanical' v2, electrical3y4and/or optical' properties. However, industrial applications clearly need a cheap growth method for the large-scale fabrication of NWs. Moreover, works are also devoted to coaxial nanocables (NCs) due to their possible uses as reinforcement agents for mechanical applications, the outer sheet of the NC acting as an interface between the nanowire and the matrix. Numerous routes have been reported for the fabrication of Sic NWs. The most promising ones for large scale production are (i) carbothermal reduction routes using carbon nanotubes (CNTs) as templates'* or a mixture of Si02 nanoparticles and active carbon' and (ii) methods based on VLS (Vapour Liquid Solid) growth mechanism e.g. assisted by catalytic metallic nan~particles'@'~. Only few techniques have been described for the preparation of silicon nitride nanowires. They usually

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315

correspond to those reported for Sic NWs but they are driven at higher temperature (1400'C to 165OOC). The three main routes are the following: (i) carbothermat reduction of sili~a'~-'', (ii) the use of carbon nanotubes as templates2' and (iii) a VLS growth technique2*.All those synthetic methods require however, either expensive starting materials such as CNTs or expensive equipment. We previously reported that the direct pyrolysis under nitrogen of commercial silicon powder in presence of graphite led to the formation of cubic silicon carbide (p-Sic) NWs22,23.These nanowires have diameters in the range 20 - 30 nm and micrometric lengths. We showed that when a mixture of argon and ox en is used instead of nitrogen, amorphous silica nanowires were preferentially obtained2 s2 . As an extent to these results, we report in the present paper the influence of the final heating temperature on the structure and chemical composition of the resulting siliconbased nanostructures.

18

EXPERIMENTAL All the experiments were conducted following the same experimental procedure. Silicon powder (Aldrich 99.999%, 60 mesh) was placed in an alumina boat containing a piece of graphite. This boat was then placed in the alumina tube of a horizontal tubular Furnace, the tube being previously degassed in vacuo before filling with nitrogen (electronic grade). Under the as flow (5 damin-'),the boat was heated up to the selected temperature (heating rate 200°C min-Ig), held for I hour then allowed to cool down to room temperature. Finally, powder was scraped from the alumina boat and analysed by scanning electron microscopy (SEM, Model N"S800, Hitachi), high-resolution transmission electron microscopy (HRTEM, Field Emission Gun microscope JEOL 20 1OF) and electron energy-loss spectroscopy (EELS, Digi-PEELS GATAN). RESULTS AND DISCUSSION In previous works, p-Sic nanowires have been obtained by heating silicon powder under In order to nitrogen at I 2OO0C, this temperature being held during I hour before cooling examine the influence of the final heating temperature on the yield, the crystallographic structure and the chemical composition of the resulting nanowires, five independent experiments have been conducted up to 1050"C, 1150"C, 12OO0C, 1250OC and 13OO0C, respectively. In all these experiments, the final temperature was held during 1h before cooling down. Each crude product was first analysed by SEM and representative images of each sample are shown in fig. 1.

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Fig. 1: SEN images of the crude products obtained at different final heating temperature: (a) 1O5O0C, (b) 1 15OoC,(c) 12OO0C,(d) 1250°C and (e) 1300°C.

The crude sample heated up to 1050°C contained only few nanowires (Fig. 1a), located on the surface of the residual silicon particfes. These NWs are straight with scattered diameters (from -10 to -300 nm) and short lengths (below -3 pm). When the experiment was conducted up to 11 5OoC, it resulted in the formation of larger amount of nanowires (Fig. Ib). Their diameters are in the same range but they are longer with lengths estimated above -8 pm. As illustrated by Fig. 1b, this result should be related with their highly curved shape. Nanowires exhibiting similar diameters and lengths were obtained at 1200°C (Fig. lc). in that case, the yield, estimated fiom the SEM

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images, was however significantly improved. In contrast, when the thermal treatment was driven up to 125OoC,the yield was not significantly improved (Fig. Id). At this temperature, two kinds of nanowires were observed: numerous long and highly curved nanowires similar to those obtained at 1200°C mixed with few straight nanowires (Fig. Id, white arrows). The latter are shorter and exhibit well-defined angles which can be interpreted as changes in direction of the N W axis (Fig. Id, white circles). The SEM image of the samples obtained after heating up to 1300°C shows a strong modification in the shape of the obtained NWs (Fig. le). No long and curved naxowires were detected but numerous straight N Ws were observed. Their diameters are compr'rable to those previously obtained but their lengths are shorter and below -8 pm. HRTEM investigations coupled with EELS analysis have been conducted in order to determine if there is a difference in structure and/or chemical composition within the two kinds of NWs observed by SEM. Figure 2a shows a HRTEM image of a typical long and curved nanowire obtained at 1200°C. On the corresponding EELS spectrum (Fig. 2b), two main features are observed at -1 00 eV and -284 eV corresponding to Si-L and C-K edges, respectively. As expected, further selected area electron diffraction (SAED) analysis showed that cubic silicon carbide (p-Sic) nanowires have been formed. According to HRTEM investigation, a high density of stackin faults was observed (fig. 2a) which can be related to high growth rate, as previously mentioned22*2 .

8

Fin. 2: HRTEM image of a p-Sic NW (a) with the corresponding EELS spectrum (b). In contrast, figure 3a shows a typical HRTEM image of a straight and short nanowire obtained at 1300°C. The chemical composition of the observed NW has been determined by EELS analysis performed on its core (Fig. 3b). Two main features are observed at -100 eV and -400 eV corresponding to Si-L and N-K edges, respectively. Features at -284 eV and -532 eV, corresponding to C-K edge and O X edges respectively, were not detected. A coating of -2.5 nm thickness was observed around the analysed NW. Its EELS analysis performed using a 2 nm probe revealed that it is composed of carbon. As illustrated by figure 3c, fast Fourier transformation (FFT) shows that the core of the NW is composed of the trigonal polymorph of silicon nitride (a-Si3N4). Complementary

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structural analysis showed that the carbon coating is amorphous. The obtained nanostructures can be considered as carbon sheathed silicon nitride nanocables (cI-S~~N~@C).

Fin. 3: (a) HRTEM image of a (r-Si3N4 NW coated with amorphous carbon with corresponding (b) EELS spectrum and (c) fast Fourier transformation (FFT) of the core of the NW. In previous works, we suggested that the silicon carbide nanowires were formed at 1200OC by a Vapour-Solid (VS) nucleation process, carbon being transported to silicon nanoparticles by nitrogen through the formation of CN-like derivatives22.23 . For the formation of a-Si3N4 NWs e.g. at higher temperature (>125O"C), we can suggest a similar VS growth mechanism with a preferential nucleation of silicon nitride onto the surface of the silicon nanoparticles due to the higher temperature reached during the experiment. The formation of an amorphous coating of carbon onto the NWs can be interpreted by considering that carbon transportation by nitrogen was still effective at that temperature. Comparative analysis of nanowires obtained at 1300°C revealed that main aSi3N4 NWs were free of carbon coating. Moreover, no SIC nanowires were detected in the sample obtained by heating up to 1300°C.This result is of primary importance for the determination of the

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growth mechanism of these nanowires because it indicates that their formation did not occurred when the temperature was increased, but either when the final temperature is reached or, less probably, during the cooling step. While only a-Si3N4 NWs were formed at 1300°C, a mixture of p-Sic and cx-Si3N4 NWs were obtained when the experiment was conducted up to 125OOC. This is evidenced by the bright field (Fig. 4a) and dark field (Fig. 4b) TEM images recorded on a whole mass of nanowires.

Fig. 4: bright field (a) and dark field (b) TEM images of a whole mass of p-Sic and a-Si3N4 NWs. The image in bright field reveals the presence of Si3N4 nanowires, which appear in bright in Fig. 4a. The dark field image 4b, made with a p-Sic reflection, reveals that the Si3N4 NWs were mixed with Sic NWs which appear in bright in this image. These observations clearly point out that between 1200°C and 1 300”C, a mixture of p-SiC NWs and a-Si3N4 NWs were obtained. CONCLUSION The direct thermal treatment of a silicon powder under nitrogen yielded the formation of cubic silicon carbide (p-Sic) in the range 1050°C - 1200°C. As we found, when the experiment was driven in the range 1200°C - 13OO0C,a mixture of P-SiC and trigonal silicon nitride (a-Si3N4) nanowires (Nws) were obtained. Moreover, when the experiment was conducted up to 13OO0C, only a-Si3N4 NWs were observed within the sample. The growth mechanism for these nanowires were presumed to be a VS process, starting with the nucleation of silicon carbide or silicon nitride nuclei onto the surface of the silicon nanoparticles. REFERENCES 1 E. W. Wong, P.E. Sheehan, C.M. Lieber, “Nanobeam mechanics: elasticity, strength, and toughness of nanorods and nanotubes”’,Science, 277,197 1- 1975 (1997). 2 Y. Zhang, N. Wang, R. He, Q. Zhang, J. Zhu, Y. Yan,“Reversible bending of Si3N4 nanowire“ J. Maier. Res., 15, 1048-1051. (2000)

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3K. W. Wong, X. T. Zhou, F. C. K. Au, H. L. Lai, C. S. Lee, S . T. Lee, “Field-emission characteristics of Sic nanowires prepared by chemical-vapor deposition”, Appl. fhys. Lett., 75, 29 18-2920. ( 1999) 4Z. Pan, H.-L. Lai, F. C. K. Au, X. Duan, W. Zhou, W. Shi, N. Wang, C.-S. Lee, N.-B. Wong, S.-. Lee, S . Xie, “Oriented silicon carbide nanowires: synthesis and field emission properties”, Adv. Muter., 12,1186-1 190. (2000) 5 D.P. Yu, L. Hang, Y.Ding, H.Z. Zhang, Z.G. Bai, J.J. Wang, Y.H. Zou, W. Qian, G.C. Xiong, S.Q. Feng, “Amorphous silica nanowires: intensive blue light emitters”, Appl. Phys. Lett., 73,30763078 $1998) W. Han, S . Fan, Q. Li, W. Liang, B. Gu, D. Yu, “Continuous synthesis and characterization of silicon carbide nanorods”, Chem. Phys. Lett., 265,374-378 (1997) 7 C. C. Tang, S . S . Fan, H. Y. Dang, J. H. Zhao, C.Zhang, P. Li, Q. Gu, “Growth of Sic nanorods prepared by carbon nanotubes-confined reaction”, J. Cryst. Growth, 2 10,595-599. (2000) ‘J-M. Nhut, R. Vieira, L. Pesant, J-P. Tessonnier, N. Keller, G. Ehret, C. Pham-Huu, M. J. Ledoux, “Synthesis and catalytic uses of carbon and silicon carbide nanostructures“, Catul. Today, 76,11-32. (2002) ’J. Pelous, M. Foret, R. Vacher, “Colloidal vs. polymeric aerogels: structure and vibrational modes“, J. Non-Cryst. Solids, 145,63-70. (1992) ‘OX. T. Zhou, H. L. Lai, H. Y. Peng, F. C. K. Au, L. S. Liao, N. Wang, I. Bello, C. S . Lee, S . T. Lee, “Thin p-Sic nanorods and their field emission properties”, Chem. Phys. Lett., 318,58-62. (20) “B. Q. Wei, J. W. Ward, R. Vajtai, P. M. Azjayan, R. Ma, G. Ramanath, “Simultineous growth of silicon carbide nanorods and carbon nanotubes by chemical vapor deposition“ Chem. Phys. Lett., 354,264-268. (2002) 12 X . T. Zhou, N. Wang, H. L. Lai, H. Y . Peng, I, Bello, N. B. Bello, N. B. Wong, C. S . Lee, S. T. Lee, ‘?-Sic nanorods synthesized by hot filament chemical vapor deposition” Appl. Phys. Lett., 74,3942-3944. (1 999) 13 H. L. Lai, N. B. Wong, X. T. Zhou, H. Y.Peng, F. C. K. Au, N. Wang, I. Bello, C. S . Lee, S. T. Lee, X. F. Duan, “Straight a-Sic nanorods synthesized by using C-Si-SiO2“, Appl. Phys. Lett., 76, 294-296. (2000) 14 Y. Zhang, N. L. Wang, R. He, X. Chen, J. Zhu, “Synthesis of S i c nanorods using floating catalyst”, Solid State Comm., 118,595-598. (2001) ”L. D. Zhang, G. W. Meng, F. Phillipp, “Synthesis and characterization of nanowires and nanocables”, Muter. Sci. and Eng. A, 286,34-38 (2000) 16X.C. Wu, W. H. Song, W. D. Huang, M. H. Pu, B. Zhao, Y. P. Sun,J. J. Du, “Simultaneous growth of a-Si3N4 and P-SiC nanorods“ Muter. Res. Bull., 36,847-852 (2001) “X. C. Wu,W. H. Song,B. Zhao, W. D. Huang, M. H. Pu, Y. P. Sun, J. J. Du, “Synthesis of coaxial nanowires of silicon nitride sheathed with silicon and silicon oxide“, Solid State Comm.,115, 683-686 (2000) 18 Y. H. Gao, Y. Bando, K. Kurashima, T. Sato, “Synthesis and microstructural analysis of Si3N4 nanorods”, Microsc, Microanal., 8,5- 10. (2002) 19 Y. H. Gao, Y . Bando, K. Kurashima, T. Sato, “Si3N4/SiC interface structure in SiCnanocrystal-embedded a-Si3N4 nanorods”, J. Appl. Phys., 91,15 15-1 5 19. (2002)

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2oW.Han,S. Fan, Q. Li, B. Gu, X. Zhang, D. Yu, “Synthesis of silicon nitride nanorods using carbon nanotube as a template”, Appl. Phys. Lett., 71,227 1-2273 (1997) *‘H. Young Kim, J. Park, H. Yang, “Synthesis of silicon nitride nanowires directly from the silicon substrates”, Chem. Phys. Lett., 372,269-274. (2003) 22K. Saulig-Wenger, D. Cornu, F. Chassagneux, G. Ferro, T. Epicier, P. Miele, “Direct synthesis of /j-SiC and h-BN coated P-SiC nanowires”, Solid State Comm., 124, 157-I61 (2002) 23K. Saulig-Wenger, D. Cornu, F, Chassagneux, G. Ferro, P. Miele, T,Epicier, “Synthesis and characterization of pC-Sic nanowires and h-BN sheathed pC-Sic nanocables”, Ceramic Transactions, (Ceramic Nanomaterials and Nanotechnology) 137,93-99. (2003) 24K. Saulig-Wenger, D. Cornu, F. Chassagneux, T. Epicier, P, Miele, “Direct synthesis of amorphous silicon dioxide nanowires and helical self-assembled nanostructures derived therefrom”, J. Mater Chem., 13,3058-3061. (2003) 25K. Saulig-Wenger, D. Cornu, P. Miele, F. Chassagneux, S. Parola, T. Epicier, “Synthesis of Si-based nanowires“, Ceramic Transactions, (Ceramic Nunomaterials and Nanotechnology 11) 148, 131-137. (2004)

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SYNTHESIS OF BORON NITRIDE NANOTUBES FOR ENGINEERING APPLICATIONS Janet Hurst David Hull NASA Glenn Research Center 2 1000 Brookpark Rd Cleveland, Ohio 44135 Daniel Gorican QSS Group 2 1000 Brookpark Rd Cleveland, Ohio 44 135 ABSTRACT Boron nitride nanotubes (BNNT) are of significant interest to the scientific and technical communities for many of the same reasons that carbon nanotubes (CNT) have attracted wide attention. Both materials have potentially unique and important properties for structural and electronic applications. However of even more consequence than their similarities may be the complementary differences between carbon and boron nitride nanotubes. While BNNT possess a very high modulus similar to CNT, they also possess superior chemical and thermal stability. Additionally, BNNT have more uniform electronic properties, with a uniform band gap of 5.5 eV while CNT vary from semiconductive to highly conductive behavior. Boron nitride nanotubes have been synthesized both in the literature and at NASA Glenn Research Center, by a variety of methods such as chemical vapor deposition, arc discharge and reactive milling. Consistent large scale production of a reliable product has proven to be difficult. Progress in the reproducible synthesis of 1-2 gram sized batches of boron nitride nanotubes will be discussed as well as potential uses for this unique material. INTRODUCTION In the last decade, significant attention from the scientific community has been focused on the area of nanotechnology and specifically upon nanotube synthesis. While carbon nanotubes have generated the bulk of interest to date, other compositions offer promise as well and may have advantages or complementary properties relative to carbon nanotubes for various applications. At NASA Glenn Research Center, where application interests are often focused on high temperature propulsion, both BN and Sic nanotube synthesis are currently under investigation for high temperature structural and electronic materials (1,2). The focus of the current effort is BNNT synthesis. While boron nitride nanotubes are known to be structurally similar to carbon nanotubes, inasmuch as both are formed from graphene sheets, much less is known about BNNT. In large part this is due to the difficulty in synthesizing this material rather than lack of interest. It has been found that BNNT have excellent mechanical properties with a measured Young’s modulus of 1.22 +/- 0.24 TPa (3). BNNT also have a constant band gap of about 5.5 eV (4). In contrast, Nanotubes, Nanorods, and Nanowires

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CNT vary from semi-conducting to conducting behavior depending on chirality and diameter of the product. Little progress has been demonstrated in the control of the chiral angle and hence the electronic properties of CNT. On the other hand, BNNT preferentially forms the zig-zag structure rather than the armchair or chiral structures due to the polar nature of the B-N bond (5). Recently, it has also been shown that BNNT systems have excellent piezoelectric properties, superior to those of piezoelectric polymers (6). Additionally, the expected oxidation resistance of BNNT relative to CNT suggests BNNT may be suitable for high temperature structural applications. This stability may be an important safety consideration for some applications, such as hydrogen storage, as carbon nanotubes readily bum in air. Many synthesis approaches have been tried with varying degrees of success. Among these approaches are pyrolysis over Co (7), CVD methods (8), arc discharge (9) laser ablation (lo), and reactive milling techniques (1 1,12). The reported approach developed at NASA Glenn Research Center produced BNNT of significant length and abundance. EXPERIMENTAL PROCEDURE BNNT were prepared by reacting amorphous boron powder in a flowing atmosphere of nitrogen with a small amount of NH3. Prior to heat treatment, fine iron catalyst particles were added in the range of up to several weight percent and briefly milled in polyethylene bottles with a hydrocarbon solvent and ceramic grinding media. Batch sizes of 2 grams are typically produced but the process should be easily scaleable to larger sizes. Milled material was applied to various high temperature substrates such as alumina, silicon carbide, platinum and molybdenum. Nanotubes were formed during heat treatments to temperatures ranging from 1100 C to 1400 C for brief times, 20 minutes to 2 hours. Nanotubes were imaged with a Hitachi S4700 field emission scanning electron microscope with a super thin window EDAX Genesis System energy dispersive spectrometer (EDS) or Phillips CM200 transmission electron microscope operated at 200 kV, with Gatan electron energy loss spectrometer (EELS). Thermogravimetric analysis (TGA) was done in air up to 1000 C. RESULTS BNNT synthesized by this method are shown in Figure 1. The low magnification photo in Figure l a shows an as-produced flake of BNNT. The flake was removed by tweezers from a 2.5 mm x 5 mm substrate or crucible. It is robust and easily handled in this form. Also, as the nanotubes are anchored within a growth media, there is no respiration hazard. BN nanotubes grow extensively both from the “top” and “bottom” of the growth media, where “top” refers to the side exposed to the atmosphere. Both top and bottom layers can be observed in Figure la as the edges were slightly rolled during handling. The nanotubes are quite long, 100 microns being common as shown in Figures l b and lc. On the bottom of the growth media, a layer of shorter nanotubes develops between the film and substrate or crucible. EDS also showed that the BNNT grow upon a film composed primarily of B, N, 0, Fe and some purities. Nanotube diameters can be very

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consistent throughout their length; however this is processing temperature dependent. Not surprisingly, variation in the heat treatment temperature resulted in somewhat different products. A lower processing temperature resulted in more fine, uniform nanotubes, as those shown in figure lb. Diameters of 20nm to 50 nm were typical. Higher temperatures and/or longer times resulted in large “nanotube” growth as shown in Figure 2a. Growth originated from fine nuclei but with excessive temperature, large structures quickly developed, up to a few microns in diameter. As shown in Figures 2a and 2b, secondary nucleation of small diameter BNNT also occurred on these larger structures. Higher processing temperatures, as well as excess catalyst concentration, resulted in interesting structures, such as nanohorns or nanoflowers, as seen in figure 2d. These structures also generally had considerable amounts of oxygen found by EDS, up to 6 wlo.

Transmission electron microscopy results are shown in Figures 3 and 4. TEM results showed the nanotubes to be nearly stoichiometric BN and a mixture of both straight walled nanotubes, as shown in Figure 3, and the “bamboo” structures, Figure 4. Predominately, the product from this method is multiwalled, commonly composed of 1530 lattice layers, although this can be affected by processing conditions. Diameters of the

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Figure 2. Examples of less typical structures synthesized at processing extremes. a) Adjacent extremes in size. b) open ended nanopods c) fine BNNT nucleated on larger tubes d) nanohorns.

Figure 3 TEM photos of typical straight walled BNNT

multiwalled BNNT were often in the 20-40 nm range, again determined by processing conditions and catalysts concentration. Figure 3 shows the atomic planes within the straight walled nanotubes exhibited lattice fringes at an angle of 12.5" with respect to the 326

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tube axis. This has also been noted elsewhere ( 5 , 13) and may be an indication of rhomohedral stacking order (12). In Figure 4 the typically highly faulted lattice walls of the bamboo structure are evident. These faulted short walls, with their open edge layers, have been suggested to be superior for hydrogen storage (14,15). The structure of the BN layers are analogous to stacked paper cups with potential hydrogen storage sites on the surface and between lattice planes, and also perhaps within the isolated voids. Predominately bamboo structures can be consistently produced by this processing method. Figure 5 shows a region of exclusively bamboo BNNT.

Figure 4 TEM photos of typical bamboo BNNT The temperature stability of as-produced BN nanotubes was investigated by thermal gravimetric analysis. The results were compared to commercial as-produced carbon nanotubes". Photos of the as-produced CNT and BNNT materials are shown in Figure 6 as well as those following heating in the TGA. The BNNT structure is clearly intact with the CNT decomposed, leaving behind the extensively oxidized iron catalyst. Figure 7 shows the TGA data in air, confirming that the carbon nanotubes have decomposed by 400 C. However, the BNNT are unaffected by the heat treatment with the exception of some slight weight gain from oxidation above 1000 C.

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Figure 6. As-processed BNNT and CNT samples were examined before and after exposure in air to 1000 C in a TGA. Photos on the left are as-processed samples. On the right are photos of material remaining following a TGA run to lOOOC

Figure 7.TGA results for BNNT and CNT samples showing the superior stability of BNNT relative to CNT

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CONCLUSION Boron nitride nanotubes were successfully and reproducibly grown by a NASA Glenn Research Center developed process. Currently 1-2 gram batches are being synthesized; however, the process is scaleable to much larger batch sizes. Sufficient amounts are now available so that boron nitride nanotubes are currently being incorporated into composites to provide high strength behavior at high temperature. It is also possible that this material may have many applications for sensors, electronics, piezoelectrics, among other applications of interest to NASA and the technical community. A preponderance of bamboo structured nanotubes could be achieved by careful additions of catalyst materials and control of processing conditions. These bamboo structures are of interest for hydrogen storage applications. The boron nitride nanotubes were found to be much more stable at high temperatures than carbon nanotubes. ACKNOWLEDGEMENTS This work was sponsored by the Alternate Fuel Foundation Technologies (AFFT) Subproject of the Low Emissions Alternative Power (LEAP) Project at the NASA Glenn Research Center. *Carbon Nanotechnologies, Inc., 16200 Park Row, Houston, TX REFERENCES

’. .J. Hurst NASA R&T 2002, NASA TM

- 2002-21 1333.

D. Larkin NASA R&T 2002, NASA TM - 2002-21 1333 N.G. Chopra, A. Zettl, “Measurement of the Elastic Modulus of a Multi-wall Boron Nitride Nanotube”, Solid State Commun. 105 (1998) 297. 4. X. Blaise, A. Rubio, S.G. Louie and M.L. Cohen, Europhys. Lett. 28 (1994) 335. 5 . L. Bourgeois, Y. Bando and T. Sato, “Tubes of Rhombohedra1 Boron Nitride”, J. Phys. D: Appl. Phys. 33 (2000)1902. S. M. Nakmanson, A. Cazolari, V. Meunier, J. Bernholc, and M. Buongiorno Nardelli, “Spontaneous Polarization and Piezoelectricity in Boron Nitride Nanotubes”, Phy. Rev. B 67,235406 (2003). 7. R. Sen, B.C. Satishkumar, A. Govindaraj, K.R. Harikumar, G. Raina, J.P. Zhang, A.K. Cheetham, C.N.R. Rao, “B-C-N, C-N and B-N Nanotubes Produced by the Pyrolysis of Precursor Molecules Over Co Catalysts”, Chem. Phys. Lett. 287 (1998) 67 1. *. P. Ma, Y. Bando, and T. Santo, “CVD Synthesis of Boron Nitride Nanotubes Without Metal Catalysts”, Chem. Phys. Lett. 337 (200 1) 6 1-64. 9. N.G. Chopra, R.J. Luyren, K. Cherry, V.H. Cresi, M.L. Cohen, S.G. Louie, A. Zettle, “Boron Nitride Nanotubes”, Science, New Series, 269 (1995) 966. 10. D.P. Yu, X.S. Sun, C.S. Lee, I. Bello, S. T. Lee, H. D. Gu, K. M. Leung, G.W. Zhou, Z. F. Dong, Z. Zhang, Appl. Phys. Lett 72 (1998) 1966. 2. 3.

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Y. Chen, L.T. Chadderton, J.S Williams, and J. Fitz Gerald, “Solid State Formation of Carbon and Boron Nitride Nanotubes”, Mater. Sci. Forum 343, 63 (2000). 12. Y. Chen, M. Conway, J.S Williams, J. Zou, “Large-Quantity Production of ‘I.

High-Yield Boron Nitride Nanotubes”, J. Mater. Res. Vol 17 No.8 (2002) 1896-1899. 13. R. Ma, Y. Bando, T. Sat0 and K. Kurashima, “ Growth, Morphology, and Structure of Boron Nitride Nanotubes”, Chem. Mater. (2001) 13 2965. 14. R. Chen, Y. Bando, H. Zhu, T. Sato, C. Xu, and D. Wu, “Hydrogen Uptake in Boron Nitride Nanotubes at Room Temperature”, J. Am. Chem. SOC.124,7672 (2002). 15. T.Oku, M. Kuno, I. Narita, “Hydrogen Storage in Boron Nitride Nanomaterials Studied by TGDTA and Cluster Calculations”, J. Phys. and Chem. of Solids, 65, 549 (2004).

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NOVEL PROCESS OF SUBMICRON-SCALE CERAMIC ROD ARRAY FORMATION ON METALLIC SUBSTRATE Kazuya Okamoto", Satoshi Hayakawa", Kanji Tsuru*,

**, Akiyoshi Osaka", **

"Biomaterials Laboratory, Graduate School of Natural Science and Technology, Okayama University, Tsushima, Okayama, 700-8530, Japan **Research Center for Biomedical Engineering, Okayama University, Tsushima, Okayama, 700-8530, Japan ABSTRACT Nano-scale rods of titania and tungsten oxide were fabricated on metallic titanium and tungsten substrates by heat-treatment in vacuo. The surface structure and morphology were characterized by Thin-film X-ray diffraction analysis (TF-XRD) and scanning electron microscopy. The TF-XRD analysis indicated that the thermal treatment in vacuo (6.7 ~10.'Pa) gave TiO, (rutile: PDF#21- 1276) and WO, (PDF#20-1324) nano-scale rod arrays. The mechanism of the nano-scale rod array formation was explained by the structure matching due to topotaxy and the crystal growth habit. INTRODUCTION One-dimensional nanostructures of polymers, metals and semiconductors in the forms of rods, tubes and others have attracted much attention for a broad range of potential applications, including catalysts, electronics, sensors, photonics, micromechanical devices, and biomedical devicesI4. Possible factors to optimize various chemical, biological and physical properties of the rod arrays are the chemical states of the metals, like valence and oxygen coordination, and the shape and size of the rods. Actually, Liu et al. pointed out that the rod shape was dependent on the growth mechanism596. They prepared submicrometer-scale rod arrays of titania on titanium (Ti) substrates with a glass-coating technique that involved selective dissolution of Recently, Gu et al. prepared chemically unstable ingredients of the glass into hot tungsten oxide nanowires with heating metallic tungsten (W) in an argon flow7, where the oxide nanowires growth resulted from the oxygen leakage of their system. Therefore, it is likely that the oxygen partial pressure in the heating atmosphere due to the oxygen leakage is one of the key factors in the rod formation process of these metallic oxides, or the partial pressure of oxygen during the heat-treatment controls the oxidation and the growth of rods. In this study, then, we heat-treated Ti and W substrates in an electric furnace in vacuo, and created new nanometer- to submicrometer scale rod array structures on the surfaces of metallic substrates. Nanotubes, Nanorods, and Nanowires

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EXPERIMENTAL PROCEDURE A sheet of commercially available pure Ti (a-Ti) and W (Nilaco, Japan) was cut to give pieces of Ti and W substrates with 1OxlOxO.l mm in size, which were subsequently cleaned in acetone. These substrates denoted as Ti-NT and W-NT. The Ti-NT samples were heated up to 700°C at a rate of 10"C/min in an electric furnace (MILA-3000, ULVAC-RIKO, Japan) in vacuo with vacuum pump ((6.7 xlO-*Pa); GLD-051, ULVAC KIKO, Japan), and kept there for 6 h, while the W-NT samples were heated up to 500°C at the same heating rate and kept there for 3 h. Then, the furnace was turned off, and the samples were let cooled down to room temperature in the furnace, before these samples were cleaned in acetone gently. These samples were denoted as Ti-VHT and W-VHT. For comparison, some of the Ti-NT and W-NT samples were heated in air under the same heating and cooling schedule. They were denoted as Ti-HT and W-HT. After they were coated with gold of 30 nm in thickness, their surface microstructure was observed under a scanning electron microscope (JSM-6300, JEOL, Japan) which was operated under 20 kV acceleration and 300 mA emission current. The crystalline phases of the samples were identified using thin-film X-ray diffraction (TF-XRD: CuKa) patterns taken by an X-ray diffract meter (RINT2500, Rigaku, Japan) operated at 40 kV-200 mA and at a scanning step of 0.02"/s. RESULTS Fig. 1 shows SEM photographs of Ti-HT and Ti-VHT. The titanium oxide covered the whole surface of Ti-HT and Ti-VHT. In Ti-VHT the random aggregation of submicron-scale rods with 1 pm in length was observed on the whole surface. TF-XRD patterns in Fig. 2 show a sharp and strong diffraction peak at 27.48" assigned to the (110) diffraction of rutile (PDF#21-1276). The peaks at 38" and 40" were assigned to a-Ti. Although the heat atmosphere was different, Ti-HT and Ti-VHT showed almost the same oxide

Fig. 1 SEM photographs of surface of the titanium substrates heated in air (Ti-HT) and in vacuo (Ti VHT).

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XRD patterns.

However, the diffraction peak width at 27.48' is narrow in Ti-VHT compared

with that in Ti-HT, indicating that the crystallite size of rutile is larger in Ti-VHT compared with that in Ti-HT.

101

Ti-NT 20

25

30

35

40

45

50

28idegree

Fig. 2 TF-XRD patterns of the titanium substrate.

See text for the sample codes.

Fig. 3 shows SEM photographs of W-HT and W-VHT, where both had nm-size crystallites but in different morphology. Aggregates of block-like crystallite of around 100 nm in size were formed in W-HT. In contrast, nm-size rod arrays were formed in W-VHT, where the size of each rod was tens of nanometers in diameter, and 100-300 nm in length. Fig. 4 shows TF-XRD patterns for W-NT, W-HT, and W-VHT. The presence of WO, was confirmed for the heat-treated samples. Note that W-HT gave a peak at 23.25', assigned to (001) of WO, (PDF#20-1324) whereas W-VHT gave one at 23.53' assigned to (020) of WO,.

Fig. 3

SEM photographs of surface of the tungsten substrates heated in air (W-HT) and in vacuo (W-VHT).

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."

h CI

C

1

d

2

W

5..

."

I

0

Y

v)

a

C

0

Y

C

A 0

Y

I

0

0

W-NT

c

10

I 20

I

30

I

40

50

A

I

60

70

80

2Bldegree

Fig. 4 TF-XRD patterns of the tungsten substrates. See text for the sample codes.

DISCUSSION The present Ti substrates gave stronger (002) diffraction than (101). This means that more a-Ti grains exposed their (002) plane to the heating atmosphere, and were oxidized. For Ti-VHT with the rutile nano-rod arrays in Fig. 2 the rutile (1 10) diffraction was by far strong accompanying very weak diffractions. We can interpret the result by considering that the topotaxy between the a-Ti (002) plane and rutile (1 10) plane. That is, a good correlation has been already pointed out between the Ti-Ti atomic distance of the a-Ti (002) plane (2.9506 A; 1 nm = lOA) and that of the rutile (110) plane (2.9587 A), and it favors the nucleation and subsequent growth of rutile layers with the (110) plane developed preferably"". Moreover, similar titania nano-rod arrays were formed on Ti due to heating at 700°C under the nitrogen atmosphere (not shown here). Therefore, those results strongly suggest that the oxygen leakage of the vacuum system during the heat-treatment controls the oxidation and the growth of rods by the vapor-solid growth process as reported by Funahashi et al." and that the decrease in the oxygen partial pressure in the heating atmosphere is favorable for obtaining the nano- and micro-scale rod array structure due to the topotaxial nucleation and growth of the rutile rods on the titanium substrate. Then, let us examine if a similar topotaxial nucleation and growth mechanism is applicable to the WO, rod formation on the tungsten substrate. The strength of the XRD peaks of W (211) and WO, (020) in Fig. 4 suggests the topotaxial correlation between those planes. Fig. 5 compares the atomic arrangements on the W (21 1) plane and WO, (020) plane. Contrary

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to the expectation, the topotaxial structure matching between those planes are not so good as that between the a-Ti (002) plane and rutile (1 10) plane. Thus, we can conclude that this inferior structure matching caused the random arrangements of the WO, nano-rod array. Indeed, Gu et al. reported that the growth direction of the nanowire of tungsten trioxide is along <010>, and this crystal growth habit is in good agreement with Fig. 4 where the (020) diffraction was the strongest.

Fig. 5 Comparison of the atomic arrangements based on the crystal structures of W (21 1) plane and WO, (020) plane.

CONCLUSION We fabricated nano- and submicron-scale rods of titania and tungsten oxide on metallic titanium and metallic tungsten surface by heat-treatment in vacuo. TF-XRD analysis indicated that thermal treatment in vacuo gave TiO, (rutile: PDF#21-1276) and WO, (PDF#20-1324) nanoand submicron-scale rod array. The mechanism of nano- and submicron-scale rod array formation was explained by the structure matching due to topotaxy and the crystal growth habit.

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ACKNOWLEDGMENT This study was supported by The Kazuchika Okura Memorial Foundation. REFERENCES ‘M. Gopal, W.J. Moberly Chan, L.C.De Jonghe, “Room temperature synthesis of crystalline metal oxides,” J. Muter: Sci., 32[22], 6001-6008. (1997). 2 U. Diebold, “The surface science of titanium dioxide,” Surf: Sci. Rept. 48, 53-229 (2003). 3Z.W. Pan, Z.R. Dai and Z.L. Wang, “Nanobelts of semiconducting oxides,” Science, 291 [5510], 1947-1949. (2001). 4P.Si, X. Bian, H. Li and Y. Liu, “Synthesis of ZnO nanowhiskers by a simple method,” Muter Lett., 57[24-251,4079-4082. (2003). 5Y.Liu, K. Tsuru, S. Hayakawa, A. Osaka, “Titanate Nano-rods Grown on Titanium

Substrates,” The 20th International Japan-Korea Seminar on Ceramics, 299-302, 2003.1 1.20-22.

Y. Liu, K. Tsuru, S. Hayakawa, A. Osaka, “In Vitro Bioactive Nano-Crystalline TiO, Layers Grown at Glass-Coatingmitanium Interfase,” J. Cerurn. SOC.Japan, 112[8], 452-457 (2004). ’G. Gu, B. Zheng, W. Q. Han, S. Roth, J. Liu, “Tungsten Oxide Nanowires on Tungsten Substrates,” Nuno. Lett., 8[2] 849-85 1 (2002). 8 Y. Liu, K. Okamoto, S. Hayakawa, K. Tsuru, A. Osaka, “Preparation of Titanate Nano-Rod Array on Titanium Substrates by Novel Microflux Method,” in Ceramic Nanomaterials and Nanotechnology 111, (Ceramic Transactions, Volume 159), American Ceramic Society, Westerville, Ohio, USA, pp. 193-200, (2004). 9 S. Hayakawa, Y. Liu, K. Okamoto, K. Tsuru, and A. Osaka, “Formation of Titania Submicron-Scale Rod Arrays on Titanium Substrate and In Vitro Biocompatibility,” in Nanoscale Materials Science in Biology and Medicine, edited by Cat0 T. Laurencin and Edward A. Botchwey (Mater. Res. SOC.Symp. Proc. 845, Warrendale, PA , 2005), AA6.9. 10 Y. Liu, K. Tsuru, S. Hayakawa and A. Osaka, “Topotaxial nucleation and growth of TiO, submicron-scale rod array on titanium substrates via sodium tetraborate glass coating,” J. Cerurn. SOC.Japan, 112[101,567-57 1, (2004). “R. Funahashi, I. Matsubara and M. Shikano, “Growth of Co-based Oxide Whiskers”, Chem. Mater., 13,4473, (2001). 6

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Tin Oxide Nanoparticle-Functionalized Multi-Walled Carbon Nanotubes by the Vapor Phase Method Wugang Fan, Lian Gao,*-+and Jing Sun State Key Lab of High Performance Ceramics and Superfine Microstructure, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050, China

The vapor phase method was used to prepare tin oxide (Sn02)functionalized multi-walled carbon nanotubes (MWNTs). SnOz nanoparticles were uniformly decorated on the outer surface of MWNTs. At the same time, parts of the MWNTs can be filled with S n 0 2 nanoparticles. This simple technique is of interest for a distinct reacting process and efficiency.

I. Introduction I N oxide (Sn02) is a typical n-type wide band gap semiconductor (Eg= 3.62 eV at 300 K), with potential applicability for gas sensor, transparent conducting electrodes, transistors, and solar cells.' Gas sensors made of carbon nanotubes have also been extensively studied. The work function of the multiwalled carbon nanotubes (MWNTs) is approximately equal to that of SnO2; the Schotty barrier between them is very low. Electrons travel through the Sn02 grains into the MWNTs, and then conduct in MWNTs with low resistance. Recently, Liang et aL2 have fabricated a low-resistance gas sensor to detect oxidizing and reducing gases from Sn02 layer-coated multi-walled carbon nanotubes. Wei et d3reported that a Sn02 gas sensor doped with carbon nanotubes exhibited a much higher sensitivity and can be operated at room temperature for detecting NO2 rather than the usually elevated temperature. The above work illustrated the potential of the composite as a superior gas sensor. Sn02-coated carbon nanotube composites were synthesized by various methods such as the chemical-solution route using SnC12 as a starting material: the one-step wet chemical m e t h ~ dand ,~ the microwave-polyol process.6 Here, we utilize the vapor phase method (VPT)7 for preparation of uniform Sn02 nanocrystalfunctionalized MWNTs. VPT is characterized by separating the reagents into solid and liquid phases. It has been used to synthesize silica nanobead-decorated MWNTs' and Ti02-coated MWNTs9 successfully. N,N-dimethyllformamide (DMF) was used as the solvent in the solid phase and SnCI4. 5H20 as the starting material in this study. The method proved to be simple and quite efficient.

T

acid (H2S04/HN03= 3 v/v) at 140°C for 3 h to oxidize and remove the amorphous carbon and catalyst. The acid treatment not only shortened MWNTs but also made the inert nanotube wall to active by grafting -COOH and -OH groups. In a typical experimental procedure, 0.2 g SnCI4.5H20 was added to 5 mL DMF solution. Then, acid-treated MWNT (20 mg) was sonicated in the above solution for 5 min as the solid phase. Later, 5 mL distilled water was added by a pipette as the liquid phase. Heat treatment on the sealed VPT instrument' (Fig. 1) was conducted at 100°C for 18 h in an oven and then the solid phase was washed with distilled water and ethanol several times. Thereafter, the precipitated solid collected by centrifugation was dried at 60°C in vacuum for 12 h.

( 2 ) Characterization of Sn02-Functionalized M WNTs The morphology and size of the composites were observed by a transmission electron microscope (TEM; Model JEM 2100F, JEOL, Tokyo, Japan). The phase composition was identified by power X-ray diffraction (XRD; Model D/max 2550 V, Rigaku Co., Tokyo, Japan), using CuKu ( h= 0.15406 nm). Energy-dispersive X-ray spectroscopy analysis (EDS, Model ISIS, Oxford Microanalysis Group, Link, High Wycombe, UK) was performed, using an accelerating voltage of 200 kV. Thermogravimetry-differential thermal analysis (TG-DTA) data were obtained using a thermal analysis meter (STA-449C, Netzsch, Bayern, Germany) in air atmosphere. The UV-visible absorption spectra of the samples were also obtained (Perkin Elmer, Lambda 950, Boston, MA). 111. Results and Discussion

Figure 2(a) reveals the XRD patterns of the Sn02 nanoparticlefunctionalized MWNTs. All the diffraction peaks could be indexed to the tetragonal Sn02 with lattice constants (JCPDS file

11. Experimental Procedure

( I ) Synthesis of the SnOrfunctionalized MWNTs Pristine MWNTs were prepared by decomposition of CH4, with diameters ranging from 15 to 40 nm and lengths ranging from 500 nm to 500 pm9 (Shenzhen Nanotech Port Co. Ltd. Shenzhen, China). We first refluxed MWNTs in mixed concentrated 9. Wei-ontributing

editor

Manuscript No. 21299. Received December 30, 2005; approved March 20, 2006. This work was supported by the National Key Project of Fundamental Research (Grant No: 2005CB6236-05). the National Natural Science Foundation of China (Grant No: 505721 14). 'Member, American Ceramic Society. 'Author to whom correspondence should be addressed. e-mail: [email protected]

Nanotubes, Nanorods, and Nanowires

Fig. 1. Vapor phase method instrument (a) container; (b) solid phase; (c) Teflon substrate; (d) holder; (e) liquid phase.

337

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a,

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Fig. 2. X-ray diffraction patterns of (a) tin-oxide functionalizedmultiwalled carbon nanotubcs (MWNTs), (b) acid-treated MWNTs.

No. 41-1445, a = 4.738, and c = 3.187 A). The width of the reflections is considerably broadened, indicating a small crystalline domain size. According to the Debye-Scherrer equation, d = 0.89h/P cos 8, the average crystalline size of the SnOz was roughly quantified as 2.6 nm. It is hard to elicit the characteristic peaks of MWNT from the XRD pattern (Fig. 2(b)) because of the good crystallization of Sn02 and their fine coating.

TEM micrographs of SnOz nanoparticle-functionalized MWNTs are shown in Fig. 3. The SnOz layer around the surface of MWNTs is overall and uniform. Weak characteristic SnOz diffraction circles were detected on this structure (inset in Fig. 3(b)). The coating layer was confirmed to be single-phase Sn02 by calculating the d spacing from the radius of rings. Figure 3(c) shows the high-resolution transmission electron microscopy (HRTEM) image of the composite. The SnOz nanoparticles present either 110 or 101 lattice fringes, which are in accordance with the cassiterite SnOz lattice fringe spacing given in the literature as 0.333 and 0.263 nm. Thus, this coating on the MWNTs is assured to be composed of nanosize Sn02 particles. In addition, parts of open-ended MWNTs were filled with discontinuous Sn02 nanoparticles (Fig. 3(d)). The filling occurs at the entire length in the hollow cavity, which can be attributed to the low surface tension of the DMF solution. The EDS was carried out to probe the composition of the functionalized nanoparticles, as shown in Fig. 4. The peaks of Sn and 0 were obviously observed; no peaks of other impurities occurred. Formation of the Sn02 coating on MWNTs’ surface can be explained by the following chemical reactions: SnC& ([)+2H20 (v) = SnOz (s)+4HCI (v); the scheme of the preparation process is illustrated in Fig. 5 . To avoid the rapid hydrolysis of SnCI4in H 2 0and for the good dispersion of MWNTs, we chose DMF as a reaction medium. This solvent is important as carbon nanotubes are practically insoluble in any solvent, and if the MWNTs

Transmission electron microscope images of multi-walled carbon nanotube (MWNT) composites. (a, b) Tin oxide (SnO2)-functionalized MWNTs and corresponding SAED; (c) High-resolution transmission electron microscopy of the MWNTs composite; (d) MWNTs filled with Sn02 nanoparticles. Fig. 3.

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Sn

I

Sn

1

1

2

3 Energy (keV)

200

4

Fig.4.

Energy-dispersive X-ray spectroscopy spectrum of a tin oxidefunctionalized multi-walled carbon nanotube. are aggregated, uniform functionalization will be difficult. Regarding the growth process, it is proposed that during the heat treatment, H 2 0 was evaporated from the bottom of the instrument and dissolved into a solid phase slowly. Here, the so-called “solid phase” means a mixture of MWNTs, DMF, and SnCI4. 5H20. This conce t was inferred from dry-gel for prep’ aration of zeolite in VPT.P,8 Tin(1V) ions in the solid phase were adsorbed on the sidewalls of acid-treated MWNTs, which possess a great deal of carboxyl group, through electrostatic. Then, during the convection of the solution, Sn02 nanoparticles tend to agglomerate and continuously crystallized on the surface of MWNTs. According to the above equation, the water amount in the liquid phase is excessive to meet the requirement of SnCI4 hydrolysis. Figure 6 shows the UV absorption spectrum of Sn02-functionalized MWNTs. The Sn02 nanoparticles prepared by this VPT have an Eg = 3.68 eV. The absorption band edge of a SnO2-functionalized carbon nanotube is estimated to be around 314 nm (3.95 eV) by extrapolating the steep slope in the curve to the long-wavelength side, showing an apparent blue shift. The composite showed a quantum size effect. TG-DTA was used to investigate the thermal properties and texture of the as-synthesized sample (Fig. 7). According to the plot, the composite began to lose its weight at the beginning of heating until 109.2”C because of evaporation of physisorbed water and HCI (ca. 8%). The exothermic peak at 274.7”C may be due to the decomposition of DMF bonded on MWNTs. The relatively pronounced weight loss spanning from 450” to 680°C could be attributed to burning of MWNTs (ca. 16%).

300

400 Wavelength

500

f 10

Fig. 6. UV-visible adsorption spectra of (a) tin oxide (SnO2)-functionalized multi-walled carbon nanotubes, (b) SnO2 nanoparticles prepared by the vapor phase method.

Fig. 7. Thermogravimetry-differential thermal analysis plot of tin ox-

ide-functionalized multi-walled carbon nanotubes. IV.

Conclusions

We have successfully prepared anatase Sn02-functionalized MWNTs by VPT. The method is simple and quite efficient. DMF had proved to be a good solvent in the process. This work presents an alternative route for the preparation of assembly of Sn02 nanocrystals and carbon nanotubes. UV absorption indicates that the composite exhibits an obvious blue shift relative to that of pure SnOz. Because Sn02 has been uniformly anchored on the MWNTs and parts of the MWNTs were filled with SnOz, this composite has potential application as a gas sensor used under ambient conditions. References ‘8.Cheng, J. M. Russell, W. Shi, L. Zhang, and E. T. Samulski, “Large-Scale. Solution-Phase Growth of Single-Crystalline Sn02 Nanorods.” J . Am. Chem. Soc., 126, 5972-3 (2004). ’Y.X. Liang, Y. J. Chen, and T. H. Wang, “Low-resistance Gas Sensor Fabricated from Multiwalled Carbon Nanotubes Coated with a Thin Tin Oxide Layer,” AppL Phys. ,!.err., 85, 666-8 (2004). ’B. Y. Wei, M. C. Hsu, P. G. Su, H. M. Lin, R. J. Wu, and H. J. Lai, “A Novel Sn02 Gas Sensor Doped With Carbon Nanotubes Operating at Room Temperature,” Sensors Acruarors B. 101, 81-9 (2004). 4W. Q . Han and A. Zettl, “Coating Single-Walled Carbon Nanotubes with Tin Oxide.” Nano ,!.err., 3, 681-3 (2003). ’L. P. Zhao and L. Gao, “Coating Multi-Walled Carbon Nanotubes with Thick La ers of Tin(lV) Oxide,” Carbon, 42, 1858-61 (2004). Y. Bai, Z. D. Xu, and Y.F. Zheng, “Microwave-Polyol Process for Functionalizing Carbon Nanotubes with Sn02 and Ce02 Coating,” Chem. Lerr., 35, 96-7 (2006). ’W. Y. Xu, J. X. Dong, 1. P. Li, J. Q . Li, and F. Wu, “A Novel Method for the Preparation of Zeolite ZSM-5,” J . Chem. Sor. Chem. Commun., 755-6 (1990). *W. G. Fan and L. Gao, “Silica Nanobeads-Decorated Multi-Walled Carbon Nanotubes by Vapor-Phase Method,” Chem. Lcrr., 34, 954-5 (2005). %. G. Fan, L. Gao, and J. Sun, “Anatase Ti02-Coated Multi-Wall Carbon Nanotubes with Vapor Phase Method,” J. Am. Ceram. Soc., 89, 731-3 (2006).0

6.

Fig. 5. Scheme for preparing tin oxide nanoparticle-functionalized multi-walled carbon nanotubes by the vapor phase method.

Nanotubes, Nanorods, and Nanowires

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Electrospinning: A Simple and Versatile Technique for Producing Ceramic Nanofibers and Nanotubes Dan Li,t Jesse T. McCann, and Younan Xia' Department of Chemistry, University of Washington, Seattle, Washington 98195-1700

Manuel Marquez7 INEST Group, Research Center, Philip Morns USA Inc., Richmond, Virginia 23234

Electrospinning is a remarkably simple method for generating nanofibers of polymers. When combined with conventional solgel processing, it provides a versatile technique for producing ceramic nanofibers with either a solid, porous, or hollow structure. This article presents a brief overview of recent progress in preparation of ceramic nanofibers by electrospinning, with a focus on an introduction to experimental procedures and analysis of several technical issues that are vital for a successful electrospinning experiment. We also highlight the unique capabilities of this technique in processing ceramic materials into nanostructures, and illustrate some potential applications of these nanostructures.

I. Introduction

P

ROCESSING of ceramic materials with effective control at the nanometer scale has been actively pursued as a problem of materials science and engineering in recent years. In particular, the preparation of one-dimensional (ID) ceramic nanostructures such as fibers, wires, rods, belts, tubes, spirals, and rings has received great interest owing to their potential applications in many technologically im ortant areas such as electronics, photonics, and mechanics.'-' A large number of chemical and/or physical methods, mostly based on the bottom-up and templatedirected routes, have been demonstrated for producing I D ceramic nanostructures with various com ositions by controlling the nucleation and growth processes.I-'Top-down approaches such as photolithography, soft lithography, and electrospinning have also been exploited to fabricate 1D ceramic nanostructures.&1° Among various top-down approaches, electrospinning appears to be the most straightforward and versatile technique for generating 1D nanostructures."" Like many conventional spinning processes that have found widespread use in the fiber industry, electrospinning also involves the ejection of a viscous solution o r melt jet from an orifice. Subsequent drawing and solidification of the jet lead to the formation of uniform, thin fibers. Unlike the conventional spinning processes, the driving forces for electrospinning are based on electrostatic interactions, instead of mechanical stretching or pulling. The most remarkable feature associated with electrospinning is that this techG. Messing-contributing editor

Manuscript No. 21349. Received January 8, 2006; approved January 31, 2006. Presented at the 91h International Ceramic Processing Science Symposium, Coral Springs, FL, Jan. 8-1 I , 2006. 'Author to whom correspondence should be addressed. e-mail: [email protected].

edU

'INEST Group, Research Center, Philip Morns USA Inc.. Richmond. Virginia. 'NCTCN Center, Physical and Chemical Properties Division, NIST, Gaithersburg, Maryland.

Nanotubes, Nanorods, and Nanowires

nique is capable of spinning fibers with diameters down to tens of nanometers. On account of this unique feature as well as the simplicity and versatility of this technique, electrospinning has gained considerable attention in the past several years. A large number of nanofibers consisting of different materials have been successfully fabricated using this technique by many research groups. The research activities and trends in this field have recently been summarized in several review articles."" Electrospinning was originally invented 70 years ago and has been used primarily for producing ultrathin polymer Recent efforts by several research groups have made it possible to generate ceramic nanofibers and nanotubes directly using this technique. This article presents a brief overview of recent progress in this research area, with a focus on the work carried out in our group. After briefly introducing the basic setup and principles of electrospinning, this article will concentrate on the general experimental procedures for electrospinning ceramic nanofibers and nanotubes, and will discuss a number of technical difficulties or problems that are frequently encountered in electrospinning experiments. Finally, we will highlight several unique features associated with the technique and illustrate the potential of electrospinning in future research related to ceramic processing and application. It is expected that this article can help the readers quickly become acquainted with the basic principles and particularly the experimental procedure for preparing 1D ceramic nanostructures. 11. How Electrospinning Works

In the generic design, an electrospinning setup consists of three major components (Fig. 1): a high-voltage power supply, a spinneret, and an electrically conductive collector. For most experiments, an ordinary hypodermic metallic needle and a piece of aluminum foil work well as the spinneret and collector, respectively. The solution for spinning is loaded in a plastic syringe, which is connected to the metallic needle. In order to control the quality of electrospun fibers, the syringe is often connected to a syringe pump so that a constant and adjustable feeding rate of the solution can be maintained. For some experiments, especially for electrospinning of ceramic nanofibers, the setup needs to be placed in a closed box so the humidity and/or atmosphere can be varied in a controllable fashion. Depending on the end use of resultant fibers, the collector can be constructed from various materials and in various configurations. The complex electrohydrodynamics involved in electrospinning have been extensively studied by a number of groups from a theoretical point of view."-" Here, we only briefly introduce the basic working principle of an electrospinning process. Like electrospray, electrospinning is a technique based on electrostatic interactions. When a high voltage is applied to the spinneret, the pendant droplet of the solution at the orifice becomes highly 341

Fig. 1.

Schematic of a typical setup for electrospinning.

electrified and is deformed into a conical shape as a result of the electrostatic repulsions between the surface charges, as well as the electrostatic attractions from the oppositely charged collector. Once the voltage has surpassed a threshold value, the electrostatic forces can overcome the surface tension and force the ejection of a liquid jet from the orifice. The electrified jet then experiences a rapid bending and whipping process in which the jet is continuously stretched and elongated by electrostatic repulsive forces, leading to the formation of a long, thin, and uniform thread. Unlike the conventional spinning processes, the drawing and thinning of fibers in electrospinning is accomplished through electrostatic repulsions of charges on the liquid jet itself, making it a contactless stretching process. As the solvent rapidly evaporates during the spinning process, the diameter of the thread is constantly reduced, leading to the formation of an ultrathin fiber. In the end, the charged fiber is deposited on the grounded collector placed under the spinneret as a result of electrostatic attraction. Depending on the rheological properties of the solution and other spinning parameters, the diameter of the resulting fiber can be varied from tens of nanometers to a few micrometers. Although the operation of electrospinning seems to be very simple, it is not always feasible to apply this technique to the fabrication of high-quality fibers. One of the major problems that one has to deal with is the spherical or spindle-like beads that may appear in a spinning process, especially when very thin fibers are produced by, for example, decreasing the concentra-

tion of a polymer solution. As a matter of fact, elimination of beads is the most difficult task in an electrospinning experiment. Some of the parameters that influence the morphology of electrospun fibers include viscosity, surface tension, and the density of net charges of the The surface tension always tends to convert the liquid jet into one or many spherical droplets in order to minimize the surface energy. In contrast, the electrostatic repulsions between surface charges on the jet tend to increase the surface area, and thus favor the formation of a thin jet rather than beads. Viscoelastic force also resists rapid changes in shape. It is the interplay of these three major forces that determines the final morphology of the product. One must keep in mind that the relative intensities of these forces may change during the spinning process as a result of the elongation of the jet and the evaporation of solvents. Particularly, as the jet attenuates, surface tension may become much more pronounced than the other forces. As a result, surface tension may dominate the process, leading to the formation of beads. Furthermore, variation in the properties of a solution during the spinning process needs to be fully considered when electrospinning is applied to preparation of ceramic nanofibers because chemical reactions such as hydrolysis and condensation are interwoven with the jetting and spinning processes. 111. Preparation of Ceramic Nanofibers by Electrospinning

( I ) General Procedure Electrospinning has been largely limited to the fabrication of nanofibers from organic polymers as it is relatively convenient to prepare a polymer solution or melt with a propriate rheological properties required for electrospinning.“” Ceramics are usually considered to be not directly spinnable although it is possible (at least, in principle) to electrospin ceramic nanofibers from their melts at extremely high temperatures. Like conventional spinning processes for generating ceramic preparation of ceramic fibers by electrospinning has to rely on the use of spinnable precursor^.^'-^^ Figure 2 shows a typical procedure for preparing ceramic nanofibers through electrospinning, with preparation of Ti02 nanofibers serving as an example. The procedure consists of three major steps: (i) Preparation of an inorganic sol or a solution containing a matrix polymer together with an alkoxide, salt, or polymer precursor. (ii) Electrospinning of the solution to generate composite nanofibers consisting of the matrix polymer and precursor. The spinning experiments are usually performed in a well-controlled environment at room temperature. (iii) Calcination, sintering, or chemical conversion of the precursor into the desired ceramic at an elevated temperature, with concomitant removal of all organic components from the precursor fibers.

Fig.2. Schematic illustration of a typical procedure for preparing ceramic nanofibers by electrospinning. The SEM image in the middle shows Ti02 nanofibers electrospun from an ethanol solution that contained titanium tetraisopropoxide (0.1 g/mL) and poly(viny1 pyrrolidone) (PVP, 0.03 g/mL). The relative humidity was 30%. The as-spun fibers were made of a composite of amorphous titania and PVP. After calcination at 500°C in air, PVP was removed and amorphous titania was converted into anatase (see the SEM image on the right) (Li and Xia2’). Reprinted with permission from D. Li and Y. Xia, “Fabrication of Titania Nanofibers by Electrospinning,” Nuno. Lett., 3, 551-60. 2003, American Chemical Society.

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(2)

Different

Types of Precursor Solutions

Preparation of a precursor solution with appropriate rheological properties is the key step for successful electrospinning. There are several types of precursor solutions that have been tested for preparing oxide-based nanofibers. Each type has its merits and shortcomings. Larsen and several other groups have demonstrated that inorganic sols with controllable viscosities prepared by hydrolysis and condensation of metal alkoxides could be directly used for electrospinning.21-26Uniform fibers made of and Si02 have been successTi02/Si02, A1203, PbZr,Til -x03, fully prepared using this approach. Although electrospinning of inorganic sols is relatively straightforward and capable of generating ceramic fibers made of various amorphous oxides, the resultant fibers are usually more than several hundred nanometers in diameter. The ability to control the size and uniformity of the nanofibers is quite limited because of the difficulty in precisely and reproducibly controlling the rheological properties of a sol. For example, an inorganic sol is a thermostatically unstable system. The viscosity changes with aging time, which makes it difficult to spin uniform fibers continuously. In addition, it may take a long period of time (up to several days or weeks) to prepare a spinnable sol. We have recently developed a new approach of continuously electrospinning ceramic nanofibers with controllable diamet e r ~ . ~In ’ this method, the precursor solution contains a nonor slightly hydrolyzed metal alkoxide-for example, titanium tetraisopropoxide or Ti(OzPr)4-for Ti02 nanofibers. To make the solution spinnable, a polymer soluble in both alcohol and water-typically, poly(viny1 pyrrolidone) or PVP-is added to regulate the rheological properties of the solution. It is very convenient to prepare such a solution by co-dissolving the precursor and polymer in alcohol. More importantly, the solution is sufficiently stable to enable continuous production of nanofibers from hours to days. In addition to the ease and simplicity associated with this polymer-assisted spinning process, it is also feasible to control the diameter of the resulting fibers. As shown in Fig. 3, the diameter of Ti02 nanofibers could be varied from tens of nanometers to several micrometers with relatively narrow size distributions by simply altering the concentration of PVP and/ or the ratio of the alkoxide precursor to PVP. By adjusting the composition of the precursor solution as well as the spinning parameters, we have demonstrated that Ti02 nanofibers thinner than 20 nm could be routinely prepared. Because the spinnability of the solution is mainly determined by the polymer (PVP in this case), one can easily extend this method to process many different types of oxides into nanofibers by simply changing the sol-gel precursor. It is worth noting that some metal alkoxides or salts may be incompatible with PVP in a solvent. Upon mixing, they tend to precipitate out of solution. This problem can be solved by changing the alkoxides or adding suitable additives to the solution. For example, tin isopropoxide, a precursor to Sn02, is not soluble in the PVP/ propanol solution. However, when small amounts of acetylacetone and acetic acid are added, a clear solution can be prepared for continuous electrospinning.28We have successfully used this method to prepare a variety of oxide-based nanofibers including, for example, Si02, V2O5, Sn02, ITO, Ge02, NiFe204, LiCo02, BaTi03, and mixed oxide^.*^.^^ Other groups have also adopted this strategy to produce oxide nanofibers for various applications.3G34Some of these studies have also shown that ordinary salts, instead of more expensive alkoxides, could serve as precursors to a number of ceramic^.^"^^ In addition to PVP, poly(viny1 alcohol) (PVA), poly(viny1 acetate), or poly(ethy1ene oxide) (PEO) have all been used as the matrix polymer to host an inorganic precursor.3s54 A wealth of metal oxides that include AI2O3,CuO, NiO, Ti02-Si02, Ge02 V205, MnzO3, Mn304, Zr02, ZnO, Co304, Nb2O5, Moo3, MgTiO3, PdO, Ce02, and LiMn204have been fabricated as fibrous structures. The fibers prepared using these polymers and with metal salts as precursors are often hundreds of nanometers Nanotubes, Nanorods, and Nanowires

400

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,

3501 A

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TiO,/PVP fibers TiO, fibers

0.03

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/ 30 -

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-m-Ti02/PVP fibers -0-TiO, fibers

0.04 0.08 0.12 Ti(OiPr), concentration (g/mL)

0.16

Fig. 3. Plots demonstrating the ability to control the diameter of electrospun nanofibers by simply adjusting the composition of spinning solutions (Li and XiaZ7).Reprinted with permission from D. Li and Y . Xia, “Fabrication of Titania Nanofibers by Electrospinning,” Nuno. Lett., 3, 55560. 2003, American Chemical Society.

in diameter or have relatively broad size distributions. It remains a challenge to prepare uniform fibers with diameters thinner than 50 nm when PVP is replaced by these polymers.

(3) Control Over the Spinning Process There is not much difference in basic operation between the electrospinning of a ceramic precursor solution and an organic polymer solution. However, as reactions such as hydrolysis, condensation, and gelation of ceramic precursors are often involved in the spinning process, special care must be taken to ensure that the process runs continuously and evenly. The problems that one might encounter in an electrospinning experiment include the blockage of the spinneret and non-uniformity of the resulting fibers. When such problems arise, the most likely cause is that the hydrolysis and gelation rates of the ceramic precursors are not well controlled. If gelation occurs too rapidly in the spinning jet, it is imaginable that the jet will become less stretchable, leading to the formation of thicker fibers or blockage of the spinneret. There are two approaches that could help solve these problems. One is to change the precursor or to add a catalyst to the spinning solution to adjust the hydrolysis and gelation rates. The other approach is to control the surrounding around the jet. An atmosphere with less humidity and/or saturated with the solvent vapor can substantially reduce the hydrolysis and gelation rates and thus facilitate continuous spinning. 343

( 4 ) Control Over the Ceramic Nanofbers Non-woven mats or membranes made of ceramic nanofibers offer two types of pores: the relatively large pores formed through the interconnection of fibers and pores much smaller in size within individual fibers. This novel class of porous materials is useful for applications such as catalysis, adsorption, separation, and tissue engineering. Balkus and c o - w o r k e r ~ ~ ~ ~ ’ have recently demonstrated that a number of oxides could be processed into mesoporous nanofibers by electrospinning inorganic sols to which structure-directing agents were incorporated. Our preliminary results indicate that porous ceramic nanofibers can also be prepared from normal ceramic precursor solutions by controlling the spinning conditions. For example, we found that humidity could affect the structures of electrospun nanofibers. Spinning of Ti(OiPr)4/PVP solutions in an atmosphere with a low relative humidity ( ~ 4 0 % yielded ) fibers with smooth surfaces (Fig. 2). If the relative humidity was greater than 70%, the resulting fibers were very rough and highly porous (Fig. 4(A)). The mechanism is similar to the one involved in the formation of pores in certain types of electrospun polymer nan~fibers.~’ It has been suggested that rapid evaporation of a highly volatile solvent during spinning could significantly cool down the jet and cause condensation of moisture on the jet as well as phase separation within the jet, leading to the formation of pores on the surface of the fibers.’* Pores could also be produced within the fibers using new electrospinning procedures. We have demonstrated that porous titania nanofibers could be prepared by co-spinning the Ti(OiPr)4/PVP solution with a polystyrene (PS) solution using a coaxial ~ p i n n e r e tDuring . ~ ~ the spinning process, the PS and PVP chains were brought together to form a more or less homogeneous mixture in certain regions of the liquid jet. As the solvents evaporate quickly in the spinning process, the two polymer phases separated to generate nanoscale domains of PS within the Ti02/PVP matrix. When both polymer components were removed by calcination, the resultant fibers became highly porous in the PS-rich regions (Fig. 4(B)). In addition to porosity, the morphology of electrospun ceramic fibers could also be varied by controlling the spinning conditions. For example, the morphology of BaTi03 fibers could be manipulated by changing the electrospinning parameters. Cylindrical polycrystalline fibers could be generated by electrospinning an isopropanol solution containing PVP and the sol-gel precursor (barium titanium ethylhexano-isopropoxide). By increasing the concentration of the Sol-gel precursor and feeding rate while electrospinning in a low-humidity environment, flat ribbons with a mesoporous structure could be fabricated (Fig. 4(C)). Larsen et af. 6o have also recently demonstrated that submicron-thick silica ribbons could be prepared by electrospinning. It is worth pointing out that although most work on elecFig.4. SEM images showing electrospun ceramic nanofibers with vartrospinning of ceramic nanofibers has focused on oxide ious morphologies obtained by adjusting the spinning parameters. (A) materials, it is also possible to electrospin non-oxide ceramic Titania nanofibers with very rough surfaces prepared by electrospinning nanofibers if suitable precursors and calcination atmosphere the same solution presented in Fig. 2 at a relative humidity of 70%. (B) are applied. For example, Larsen et al. have fabricated S i c Porous Ti02nanofibers prepared by co-electrospinning a TiOz precursor solution with a polystyrene solution, followed by calcination in air at fibers by electrospinning a sol solution containing Novolac 500°C in air (Li and Xia54. (C) Barium titanate (BaTiO,) nanobelts resin and tetraethylorthosilicate, followed by conversion of prepared by electrospinning an isopropanol solution containing both these fibers into S i c through high-temperature pyrolysis. Snedbarium titanium ethylhexano-isopropoxide(0.05 g/mL) and poly(viny1 don and co-workers have recently fabricated boroncarbide/ pyrrolidone) (0.08 g/mL). Reprinted with permission from D. Li and carbon nanofibers by electrospinning of a single-source Y . Xia, “Direct Fabrication of Composite and Ceramic Hollow Nanopolymeric ceramic precursor, followed by pyrolytic ceramic fibers by Electrospinning,” Nuno. Lett., 4, 933-8. 2004, American conversion. Chemical Society. IV.

Preparation of Ceramic Nanotubes by Electrospinning

Inorganic tubular nanostructures are of fundamental and technological importance in a number of areas. In the literature, methods for fabricating such nanostructures are dominated by the bottom-up and template-assisted synthetic approaches. Recently, our g r o ~ p ~and~ the , ~Larsen ~ - ~group@ ~ have independently demonstrated that electrospinning could be directly

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utilized to fabricate continuous ceramic nanotubes, providing a simple, top-down approach for preparing extremely long ceramic nanotubes.

(1) Setup and General Procedure The basic idea for preparing ceramic nanotubes using electrospinning is to prepare nanofibers with a core/sheath structure by co-spinning two liquids, followed by selective removal of the Progress in Nanotechnology: Processing

Fig. 5. Scheme showing a typical procedure for direct preparation of ceramic nanotubes by electrospinning.The spinneret (left panel) was fabricated from two coaxial capillaries, through which heavy mineral oil and an ethanol solution containing poly(viny1 pyrrolidone) (PVP)and Ti(0iPrk were cospun to form continuous nanofibers with a core/sheath structure (middle panel). To show the core/sheath structure clearly, the sample shown in the middle image has been extracted with octane to remove the oil core. The walls of these tubes were made of a compositecontaining PVP and amorphous TiOz. Pure anatase nanotubes were obtained by calcining the as-obtained sample in air at 500°C. Both TEM (right panel) and SEM (inset) clearly indicate a uniform wall thickness for these nanotubes (Li and Xia54. Reprinted with permission from D. Li and Y. Xia, “Direct Fabrication of Composite and Ceramic Hollow Nanofibers by Electrospinning,” Nano. Lett., 4, 933-8. 2004, American Chemical Society.

core material. Figure 5 illustrates the typical pr~cedure.’~ The spinneret contains two coaxial capillaries, which can be easily constructed by inserting a polymer-coated silica capillary into a stainless-steel needle. In a typical experiment, two viscous liquids are simultaneously fed through the inner and outer capillaries, respectively. When a right combination of liquids and operation conditions is used, a coaxial jet can be formed under the application of a suitable voltage. Subsequent stretching and whipping of this jet by electrostatic forces results in the generation of coaxial nanofibers. As shown in Fig. 5 , such core/sheath nanofibers (TiOz-PVP/oil) can be readily prepared by co-spinning heavy mineral oil (its viscosity is 66.5-70.4 centistokes at 40°C according to the supplier’s data) and an ethanol solution containing PVP and Ti(OiF’r)4. Selective removal of the oil phase by solvent extraction gives tubular fibers with their walls consisting of Ti02/PVP. Titdnia nanotubes can be obtained by simultaneously removing PVP and oil through calcination of the as-spun fibers at an elevated temperature in air. By controlling the spinning parameters such as electric field strength, the concentration of sheath liquid, and the feeding rates for both liquids, the wall thickness and inner diameter of the nanotubes can be varied in the range from tens of nanometers to several hundred nanometers.

(2) Solutionsfor Coaxial Electvospinning Although the concept of electrospinning core/sheath nanofibers and nanotubes seems to be simple and straightforward, formulation of the solutions for coaxial electrospinning is rather complicated. One has to take many experimental parameters into account for a successful experiment, particularly when a new composition or morphology is desired. It is expected that the following lessons that we have learned from our experiments could be used by others to design new coaxial spinning experiments. Firstly, both the core and sheath solutions must be sufficiently viscous. During a co-spinning process, the elongation of the sheath is mainly driven by electrostatic repulsions between the surface charges. Rapid stretching of the sheath causes a strong viscous stress, which is then passed onto the core liquid. The shear force stretches the oil phase and elongates it along with the sheath through mechanisms such as viscous dragging and/or contact friction. If the rheological properties are properly controlled, the core and sheath liquids can be stretched and elongated at the same magnitude throughout the spinning process, resulting in the formation of core/sheath nanofibers with continuous, uniform inner/outer diameters. During the elongation process, the final morphology of the inner liquid is also determined by the competition between surface/interface tension and shear force. It was found that if an oil of lower viscosity, for example, light mineral oil (according to the supplier’s data, its Nanotubes, Nanorods, and Nanowires

viscosity is 14.2(r17.00 centistokes at 40”C), was used, the resultant fibers might exhibit beads in addition to hollow structures. Secondly, immiscibility of core and sheath liquids (including solvents) seems to be necessary for the formation of continuous and uniform hollow fibers.59Preferably, the liquid source for the core is not volatile. If the two co-spinning solutions or their solvents are miscible, the two solutions will be more or less mixed up during the bending and stretching of the jet. In particular, rapid diffusion and evaporation of solvents during spinning could significantly cause the mixing of the two phases, making it difficult to achieve well-defined core/sheath structures. Thirdly, the mechanical strength of the sheath materials is important for preparing robust nanotubes. The presence of a ceramic sol-gel precursor in the sheath solution or the use of inorganic sols as the sheath source is thought to be necessary for the formation of stable, coaxial jets and thus robust tubular structures. Although several groups have recently claimed that polymer nanofibers with a core/sheath structure could be prepared by spinning two pol mer solutions with miscible solvents using a coaxial ~ p i n n e r e t : ~so~far there have been no reports regarding the feasibility of directly preparing polymer nanotubes by selectively removing the core polymer from as-spun core/ sheath fibers. It is likely that organic polymer nanotubes are not robust enough to maintain the tubular morphology during the removal of the core material. Another explanation is that the two polymers partially mix during the spinning process, making it difficult to obtain well-defined nanotubes. We have also failed to obtain polymer nanotubes by co-spinning two polymer solutions@or co-spinning PVP solutions with immiscible oil. In fact, when a PVP solution was co-spun with heavy mineral oil, it was even difficult to obtain stable coaxial jets. Addition of Ti(OiPr)4 to the PVP solution greatly improved the stability of coaxial jets to obtain core/sheath nanofibers with uniform diameters. It is believed that the hydrolysis and condensation of Ti(01Pr)~increase the viscosity of the PVP solution, and thus allow the PVP sheath to transmit the viscous stress to the interface between core and sheath more effectively during the spinning process. The in situ formation of a ceramic network also strengthens the sheath in the electrified jet to prevent the oil from leaking out during the rapid stretching of coaxial jets. More importantly, the inorganic gel in the walls significantly enhances the stability of the resulting nanotubes. These experiments clearly demonstrate the important role of sol-gel reactions in preparing stable nanotubes. V.

Enhancement of Sol-Gel Processing by Electrospinning

Processing of ceramic materials via the sol-gel route has been widely used in conventional ceramic production processes. In

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Fig. 6. (A) SEM image of a uniaxially aligned array of Sb-doped Sn02 nanofibers prepared by electrospinning.(B) SEM image showing a fiberbased device fabricated by simply transferring a single Sb-doped Sn02 nanofiber onto a pair of gold electrodes. The inset in (B) shows a typical I-V curve measured from the device (Li er aL2*). Reprinted with permission from D. Li, Y. Wang, and Y. Xia, “Electrospinning of Polymeric and Ceramic Nanofibers as Uniaxially Aligned Arrays,” Nuno. Left., 3, 1167-71. 2003, American Chemical Society.

the last few decades, sol-gel processing has been extensively developed for low-temperature preparation of many ceramic materials into coatings, thin films, fibers, and monoliths for various applications such as catalysis, solar cells, rechargeable batteries, sensing, electronics, and p h o t o n i ~ s . ’ ~Combined ~” with electrospinning, the sol-gel processing technique is now capable of producing ceramic nanofibers and nanotubes with controllable compositions, sizes, and shapes and in relatively large quantities. This section briefly highlights some of the most interesting features associated with this technique. The combination of electrospinning with conventional sol-gel processing not only makes it possible to generate nanoscale fibers, it also makes it feasible to manipulate the as-spun fibers and assemble them into ordered structures. Electrospun nanofibers are highly charged in the spinning process. This feature enables electrostatic manipulation of the nanofibers. We have demonstrated that electrospun nanofibers tend to be uniaxially aligned over the insulating area between two grounded parallel electrodes as a result of electrostatic interaction between fibers and collectors?8 The simplest demonstration of this method is that uniaxially aligned arrays of nanofibers can be readily formed over the void gap of a metallic tweezer that is placed on the grounded electrode. The aligned fibers can be readily transferred onto other substrates for various applications. Figure 6(A) shows the SEM image of an array of Sb-doped SnOz nanofibers fabricated using this simple approach. It is worth noting that by controlling the collection time, single fibers could be collected across the gap, manually transferred onto a solid substrate, and then incorporated into the fabrication of single nanofiber-based devices. Figure 6(B) shows a single Sb-doped Sn02 nanofiber

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transferred on a pair of gold electrodes separated by 20 pm. The inset shows a nonlinear current-voltage curve measured from this nanofiber.28 We have further demonstrated that the configuration of fibers deposited on insulating areas can be controlled by the shape and size of patterned electrodes. By using specially designed electrodes as the collector, multiple layers of well-designed nanofibers with different compositions could be conveniently assembled to generate highly ordered, hierarchically porous architectures for the fabrication of ceramic devices with attractive f~nctionalities.~~”” The marriage of electrospinning with sol-gel processing can also provide solutions to some technical problems that are difficult to deal with using conventional methods. In conventional sol-gel processing, it is difficult to prepare a homogeneous and stable sol or gel reproducibly because hydrolysis and condensation of metal alkoxides or salts are generally very sensitive to local reaction conditions. If two or more precursors with different hydrolysis rates are involved, it becomes even more challenging to control the homogeneity of the sol. Phase separation often occurs in these systems, destabilizing the sol. These problems can be avoided using electrospinning. In electrospinning, the sol-gel reactions (hydrolysis, condensation, and gelation) are confined to an extremely small space (the spinning jet and the final nanofiber). Furthermore, because of the rapid bending and elongation of the jet, solvent molecules in the jet can quickly evaporate, and each component is rapidly solidified and fixed in the fibers. All of these features can prevent or greatly minimize the phase separation of different compositions in the final product. Reduction in phase separation has been observed in electrospun nanofibers composed of polymer blends!’ This unique feature is instrumental for the preparation of complex oxide ceramics, ceramic/polymer nanocomposites, as well as uniformly doped ceramics. For example, we have found that the nanocrystallites in electrospun NiFe204 nanofibers were more uniform in dimension than those powders prepared using a conventional sol-gel process (Figs. 7A and B) because NiO and Fe203 in the as-obtained nanofibers were thoroughly and evenly mixed.29 Another remarkable feature that distinguishes the electrospinning approach from other techniques for preparing ceramic nanofibers and nanotubes is that it is simple and convenient to incorporate or encapsulate other materials into the nanofibers and nanotubes. Note that doping or incorporation of foreign materials into a solid matrix material is often required for many applications of ceramics such as catalysis and chemical sensors. In an electrospinning process, functional materials (molecules or nanoparticles) can be easily incorporated into nanofibers by adding these materials or their precursors to the spinning solutions. If the material is not dispersible in the solution, it could be introduced through a co-spinning route from a separate solution. As for the encapsulation and functionalization of nanotubes, functional molecules (e.g. fluorescent dyes and silanes) and nanoparticles (e.g. metallic or magnetic nanoparticles) can be readily incorporated into the tubes by simply dispersing the corresponding molecules, precursors, or nanoparticles into the core phase prior to spinning.6,6’Incorporation of functional components into the tubes could lead to many new applications. For example, adding different silanes to the oil allows for convenient derivatization of the inner surfaces with different functional groups, which could significantly enhance the application of these nanotubes in nanofluidics, absorption, and separation?’

VI.

Concluding Remarks

Electrospinning is a simple and versatile method for fabricating ceramic nanofibers and nanotubes. The above discussion on the unique features of this technique clearly demonstrates that a combination of electrospinning with conventional sol-gel processing not only provides a powerful tool for producing inorganic fibers and tubes with diameters down to the nanometer scale but also offers new and simple approaches to manipulate Progress in Nanotechnology: Processing

Fig.8. SEM images showing that electrospun titania nanofibers could be decorated with gold nanostructures through photocatalytic reduction of HAuC14 in the presence of poly(viny1 alcohol) (PVA). The morphology of deposited gold nanostructures could be controlled by varying the M (Li et al.’?. concentration of PVA: (A) 0.1 M; and (B) 5.0 x Reprinted from Chemical Physics Leffers,394, D. Li, J.T. McCann, M. Gratt, and Y. Xia, “Photocatalytic Deposition of Gold Nanoparticles on Electrospun Nanofibers of Titania,” 387-91, Copyright 2004, with permission from Elsevier.

Fig.7. (A, B) TEM images of NiFe204nanofibers (A) and powders (B) prepared by electrospinning and using the conventional sol-gel process, respectively. (C) Magnetic properties of these two samples at 5 K (Li ef Reprinted with permission from D. Li, T. Herricks, and Y. Xia, “Magnetic Nanofibers of Nickel Ferrite Prepared by Electrospinning,” Appl. Phys. Left., 83, 4 5 8 6 8 . Copyright 2003, American Institute of Physics.

the fibers and to further create complex, nanostructured architectures and systems. As it is now possible to fabricate a variety of ceramic nanofibers and nanotubes using electrospinning, future research in this area should be shifted to the exploration of interesting properties and applications associated with these 1D nanostructures and their hierarchical assemblies. First of all, electrospun ceramic nanofibers are promising for applications in those areas where high porosity and surface areas are desired, for example, as components for structural reinforcement and thermal insulation; as active units in sensing; as membranes for purification and separation; as supports in catalysis; and as electrodes in Nanotubes, Nanorods, and Nanowires

energy conversion or storage devices that include batteries, dyesensitized solar cells, supercapacitors, and actuators. A number of publications have recently demonstrated some applications in these a ~ p e c t s ? l - ~ ~ It is also attractive to explore ceramic nanofibers in unconventional applications, for example, as templates for synthesizing new types of nanostructures; as functional coatings for controlling the wettability of surfaces; as nanoscale channels in micro- and nanofluidics; and as ultra-small waveguides in optic devices. For instance, we have used titania nanofibers to generate metal nanostructures with controllable shapes on the surface of the fibers via photocatalysis (Fig. 8).76It has also been demonstrated that superhydrophobic surfaces could be easily obtained by coating the target material with electrospun silica nanofibers, followed by functionalization with a fluorine-derivatized ~ilane.~’ Electrospun nanofibers are extremely long. Such long nanofibers and nanotubes made of ceramics also provide an ideal model system to investigate some fundamentally important problems, especially some phenomena at the nanometer scale in ceramic science. For example, many properties of polycrystalline ceramic materials (e.g. gas sensing or superconductivity) are strongly related to the grain size and grain boundary as well as the domain size of each phase if the material is composed of multiple phases. Most ceramic nanofibers prepared by electrospinning are polycrystalline. The grain or domain size can be comparable with the diameter of the fibers. It would be enlight347

ening to study how each grain or phase forms and evolves within the confined, nanoscale, ID nanostructures during sintering and how the electronic, optic, and magnetic properties of the resultant nanofibers differ from the corresponding bulk materials or other forms of nanostructures. Figure 7(C) demonstrates that the difference in shape of magnetic nickel ferrite could be significant enough to result in different magnetic proper tie^.^^ As both electrospinning and sol-gel processing techniques are being constantly improved, it is expected that progress in understanding the properties and applications of ceramic nanofibers and nanotubes will be accelerated in the coming years.

Acknowledgments This work has been supported in part by an AFOSR-MURI grant on smart skin materials awarded to the UW, and a research fellowship from the David and Lucile Packard Foundation. Y. X. is a Camille Dreyfus Teacher Scholar (2002) and an Alfred P. Sloan Research Fellow (2000). This work was performed in part at the Nanotech User Facility (NTUF) at the UW, a member of the National Nanotechnology Infrastructure Network (NNIN) funded by the NSF.

References ‘Y. Xia, P. Yang, Y. Sun, Y. Wu, B. Mayers, B. Gates, Y. Yin, F. Kim, and H. Yan, “One-Dimensional Nanostructures: Synthesis, Characterization, and Applications,” Adv. Muter., 15, 353~89 (2003). ’M. Law, J. Goldberger, and P. Yang, “Semiconductor Nanowires and Nanotubes,” Annu. Rev. Mater. Res., 34, 83-122 (2004). ‘Z. L. Wang, “Functional Oxide Nanobelts: Materials, Properties and Potential Applications in Nanosystems and Biotechnology,” Annu. Rev. Phys. Chem., 55, 159-96 (2004). 4Y. Yin, B. Gates, and Y. Xia, “A Soft Lithographic Approach to the Fabrication of Nanostructures of Single Crystalline Silicon with Well-Defined Dimensions and Shapes,” Adv. Marer., 12, I 4 2 6 9 (2000). ’Y. Sun, D. Y. Khang, F. Hua, K. Hurley, R. G . Nuzzo, and J. A. Rogers, “Photolithographic Route to the Fabrication of Micro/Nanowires of Ill-V semiconductors,” Adv. Funct. Muter., 15, 3 M (2005). ‘D. Li and Y. Xia, “Electrospinning of Nanofibers: Reinventing the Wheel,” Adv. Murer., 16. 1151-70 (2004). ‘D. H. Reneker and I. Chun, “Nanometre Siameter Fibres of Polymer, Prcduced by Electrospinning,” Nanotechnology, 7, 216-23 (1996). 8A. Frenot and I. S . Chronakis, “Polymer Nanofibers Assembled by Electrospinning,” Curr. Opin. Colloid Interf. Sci., 8, 64-75 (2003). 9Z.-M. Huang, Y.-Z. Zhang, M. Kotaki, and S. Ramakrishna, “A Review on Polymer Nanofibers by Electrospinning and their Applications in Nanocomposites,” Compos. Sci. Technol., 63, 2223-53 (2003). ‘9 Subbiah, . G. S. Bhat, R. W. Tock, S . Pararneswaran. and S. S. Ramkumar, “Electrospinning of Nanofibers,” J. Appl. Polym. Sci., 96, 55749 (2005). I’D. H. Reneker, A. L. Yarin. H. Fong, and S. Koombhongse, “Bending Instability of Electrically Charged Liquid Jets of Polymer Solutions in Electrospinnin ” J. Appl. Phys., 87, 4531-7 (2000). ”A. L. Yarin, S. Koombhongse, and D. H. Reneker, “Taylor Cone and Jetting from Liquid Droplets in Electrospinning of Nanofibers,” J. Appl. Phys., 90,4836 46 (2001). ‘’Y. M. Shin, M. M. Hohman, M. P. Brenner, and G. C. Rutledge, “Experimental Characterization of Electrospinning: The Electrically Forced Jet and Instabilities,” Polymer, 42, 995S67 (2001). I4M. M. Hohman, M. Shin, G. C. Rutledge, and M. P. Brenner, “Electrospinning and Electrically Forced Jets. 1. Stability Theory,” Phys. Fluids, 13, 2201-20 (2001). ISM. M. Hohman, M. Shin, G. C. Rutledge, and M. P. Brenner, “Electrospinning and Electrically Forced Jets. 11. Applications,” Phys. Fluids, 13, 2221 (2001). “S. V. Fridrikh, J. H. Yu, M. P. Brenner, and G. C. Rutledge, “Controlling the Fiber Diameter During Electrospinning,” Phys. Rev. Lett., 90, 144502 (2003). ‘’5. J. Feng. “The Stretching of an Electrified Non-Newtonian Jet: A Model for Electrospinning,” Phys. Fluids, 14, 3912-26 (2002). “H. Fong. 1. Chun, and D. H. Reneker. “Beaded Nanofibers Formed During Electrospinning.” Polymer, 40,4585-92 (1999). I9S. Sakka, “Fibers from the SolGel Process”; pp. 1 4 M l in Sol-Gel Technologyfor Thin Films. Fibers, Performs, Electronics, and Specialty Shapes, Edited by L. C. Klein. Noyes Publications, 1988. C. Lacourse, “Continuous Filament Fibers by the S o K e l Process”; pp. 18499 in Sol-Gel Technology for Thin Films. Fibers, Performs, Electronics, and Specialty Shapes, Edited by L. C. Klein. Noyes Publications, 1988. 2’G.Larsen, R. Velarde-Ortiz. K. Minchow, A. Barrero, and 1. G . Loscertales,“A Method for Making Inorganic and Hybrid (Organic/Inorganic) Fibers and Vesicles with Diameters in the Submicrometer and Micrometer Range Via SolCel Chemistry and Electrically Forced Liquid Jets.” J. Am. Chem. Soc., 125. I1545 (2003). 22W. Kataphinan, R. Teye-Mensah, E. A. Evans, R. D. Ramsier. D. H. Reneker, and D. J. Smith, “High-Temperature Fiber Matrices: Electrospinning and Rare-Earth Modification,” J. Vac. Sci. Technol. A. 21, 15748 (2003). ”S.-S. Choi. S. G . Lee, S. S. Im, and S. H. Kim, “Silica Nanofibers from Electrospinning/SolGel Process,” J. Mater. Sci. Lett., 22, 891-3 (2003).

’w.

348

24 Y. Wang, R. Furlan, I. Ramos, and J. I. Sdntigano-Aviks, “Synthesis and Characterization of Micro/Nanoscopic Pb(ZrOs2Ti0an)O, Fibers by Electrospinnin ,” Appl. Phys. A , 78. 1043-7 (2004). ”Y. Wang and J. J. Santigano-Aviles, “Synthesis of Lead Zirconate Titanate Nanofibres and the Fourier-Transform Infrared characterization of Their Metallo-Organic Decomposition Process,” Nanorechnology. 15, 32-6 (2004). 26 G. Zhang, W. Kataphinan. R. Teye-Mensah, P. Katta, L. Khatri, E. A. Evans. G . G. Chase, R. D. Ramsier, and D. H. Reneker, “Electrospun Nanofibers for Potential Space-Based Applications,” Mater. Sci. Eng. B, 116, 353-8 (2005). 27D. Li and Y. Xia, “Fabrication of Titania Nanofibers by Electrospinning,” Nano Lett., 3, 55540 (2003). 2RDD.Li. Y. Wang, and Y. Xia, “Electrospinning of Polymeric and Ceramic Nanofibers as Uniaxially Aligned Arrays,” Nono Lett., 3, 1167-71 (2003). 29D. Li, T. Herricks, and Y. Xia, “Magnetic Nanofibers of Nickel Ferrite Prepared by Electrospinning,” Appl. Phys. Lett., 83. 4586-8 (2003). 30 V. Tomer, R. Teye-Mensah, J. C. Tokash, N. Stojilovic, W. Kataphinan, E. A. Evans. G. G. Chase, R. D. Ramsier. D. J. Smith, and D. H. Reneker. “Selective Emitters for Thermophotovoltaics: Erbia-Modified Electrospun Titania Nanofibers,” Sol. Eng. Muter. Sol. Celk. 85. 477-8 (2005). ”J. Yuh, J. C. Nino. and W. A. Sigmund, “Synthesis of Barium Titanate (BaTi03) Nanofibers Via Electrospinning.” M a w . Lett., 59, 3645-7 (2005). 32M. Azad, T. Matthews, and J. Swary, “Processing and Characterization of Electrospun YzO’-Stabilized ZrO2 (YSZ) and Gd203-DopedCeO2 (GDC) Nanofibers,” Mater. Sci. Eng. B. 123, 2 5 2 4 (2005). 33M. A n d , “Fabrication of Yttria-Stabilized Zirconia Nanofibers by Electrospinning,” Muter. Lett., 60,67-72 (2006). ”Z. W. Fu, J. Ma, and Q. Z. Qin, “Nanostructured LiCo02 and LiMn204 Fibers Fabricated by a High Frequency Electrospinning,” Solid Stare Ionicc, 176. 1635-40 (2005). ’’H. Dai, J. Gong, H. Kim, and D. Lee, “A Novel Method for Preparing Ultrafine Alumina-Borate Oxide Fibres Via an Electrospinning Technique,’’ Nanotechnology, 13, 674-7 (2002). %B. Ding, H. Kim, C. Kim, M. Khil, and S. Park, “Morphology and Crystalline Phase Study of Electrospun Ti02-Si02 Nanofibres,” Nanotechnology, 14. 532-7 (2003). 37C.Shao, H.-Y. Kim. J. Gong, B. Ding, D.-R. Lee, and S.-J. Park. “Fiber Mats of Poly(Viny1 Alcohol)/Silica Composite Via Electrospinning.” Mater. Lett.. 57. 157: (2003). P. Viswanathamurthi. N. Bhattarai. H. Y. Kim, D. R. Lee, S . R. Kim, and M. A. Morris, “Preparation and Morphology of Niobium Oxide Fibres by Electrospinning.” Chem. Phys. Lett., 374, 79-84 (2003). ”P. Viswanathamurthi, N. Bhattarai, H. Y. Kim, and D. R. Lee, “Vanadium Pentoxide Nanofibers by Electrospinning,” Scr. Mater., 49, 5774 I (2003). 40N. Dbarmaraj, H. C. Park, B. M. Lee, P. Viswanathamurthi, H. Y. Kim, and D. R. Lee, “Preparation and Morphology of Magnesium Titanate Nanofibres Via Electrospinning,” Inorg. Chem. Commun., 7, 43 1-3 (2004). 4’P. Viswanathamurthi, N. Bhattarai, H. Y. Kim. D. 1. Cha. and D. R. Lee, “Preparation and Morphology of Palladium Oxide Fibers Via Electrospinning,” Muter. Lett.. 58, 3368-72 (2004). 42P.Viswanathamurthi, N. Bhattarai, H. Y. Kim, M. S. Khil, D. R. Lee, and E.K. Suh, “ G d 2 Fibers: Preparation, Morphology and Photoluminescence Prope r t a ” J. Chem. Phys., 121, 441-5 (2004). P. Viswanathamurthi, N. Bhattarai, H. Y. Kim. and D. R. Lee, “The Photoluminescence Properties of Zinc Oxide Nanofibres Prepared by Electrospinnin ,” Nanotechnology, 15, 320-3 (2004). ’N. Dharmaraj, H. C. Park, C. K. Kim. H. Y. Kim, and D. R. Lee, “Nickel Titanate Nanofibers by Electrospinning,” Muter. Chem. Phys., 87, 5-9 (2004). 45H. Guan, C. Shao, B. Chen, J. Gong, and X. Yang. “A Novel Method for Making CuO Superfine Fibres Via an Electrospinning Technique,” Inorg. Chem. Commun., 6 , 1409-1 I (2003). &H. Guan, C. Shao, B. Chen, J. Gong, and X. Yang. “Preparation and Characterization of NiO Nanofibres Via an Electrospinning Technique,’’ Inorg. Chem. Commun., 6. 1302-3 (2003). 47 H. Guan, C. Shao, S. Wen, B. Chen, J. Gong, and X. Yang, “A Novel Method for Preparing Co304Nanofibers by Using Electrospun PVA/Cobalt Acetate Composite Fibers as Precursor,” Muter. Chem. Phys.. 82, 10024 (2003). 48 X. Yang, C. Shao. H. Guan, X. Li, and J. Gong, “Preparation and Characterization of ZnO Nanofibers by Using Electrospun PVA/Zinc Acetate Composite Fiber as Precursor,” Inorg. Chem. Commun., 7, 176-8 (2004). 49C. Shao, H. Guan, Y. Liu, J. Gong, N. Yu, and X. Yang, “A Novel Method for Making Zr02 Nanofibres Via an Electrospinning Technique,” J. Cryst. Growth, 267, 3 8 0 4 (2004). M C. Shao. X. Yang, H. Guan, Y. Liu, and J. Gong, “Electrospun Nanofibers of NiO/ZnO Composite,” Inorg. Chem. Commun., 7, 625-7 (2004). ”C. Shao, H. Guan, Y. Liu, X. Li, and X. Yang, “Preparation of Mn203 and Mn304 Nanofibers Via an Electrospinning Technique,” J. Solid State Chem.. 177, 2628-31 (2004). 52 X. Yang. C. Shao. H. Guan, X. Li, and H. Gong, “Preparation and Characterization of ZnO Nanofibers by Using Electrospun PVA/Zinc Acetate Composite Fiber as Precursor,” Inorg. Chem. Commun., 7. 176-8 (2004). ”X. Yang, C. Shao, Y. Liu, R. Mu. and H. Guan. “Nanofibers of CeO, Via an Electrospinning Technique,” Thin Solid Fib, 478, 228-3 1 (2005). ”N. Yu, C. Shao. Y. Liu, H. Guan, and X. Yang, “Nanofibers of LiMn204by Electrospinning,” J. Colloid Inrert Sci., 285, 1 6 3 4 (2005). ”S. Madhugiri, W. Zhou, J. P. Ferraris, and K. J. Jr. Balkus, “Electrospun Mesoprous Molecular Sieve Fibers,” Micropor. Mesopor. Muter., 63, 75-84 (2003).

Progress in Nanotechnology: Processing

5hM.Sudha. B. Sun, P. G. Smirniotis, J. P. Ferraris, and K. J. Balkus, ”Electrospun Mesoporous Titanium Dioxide Fibers,” Micropor. Mesopor. Mufer.,69, 77-83 (2004).

57M.Macias, A. Chacko, J. P. Ferraris. and K. J. Balkus, “Electrospun Mesoporous Metal Oxide Fibers.” Micropor. Mesopor. Murer.. 86, 1-13 (2005). 5H S. Megelski, J . S. Stephens, D. B. Chase, and J. F. Raholt, “Micro- and Nanostructured Surface Morphology on Electrospun Polymer Fibers,’’ Macrornokcdes, 35, 8 4 5 M (2002). 59D. Li and Y. Xia, “Direct Fabrication of Composite and Ceramic Hollow Nanofibers by Electrospinning.” Nuno Lerr.. 4, 933-8 (2004). 60 G. Larsen. S. Noriega, R. Spretz, and R. Velarde-Ortiz, “Electrohydrodynamics and Hierarchical Structure Control: Submicron-Thick Silica Ribbons with an Ordered Hexagonal Mesoporous Structure,” J . Muter. Chem., 14, 2372-3

“H. Yu, S. V. Fridrikh, and G . C. Rutledge, “Production of Suhmicrometer Diameter Fibers by Two-Fluid Electrospinning.” Adv. Mufer., 16, 1562-6 (2004). 67Y.Z. Zhang, Z. M. Huang, X. J. Xu, C. T. Lim, and S. Ramakrishna, “Preparation of Core-Shell Structured PCL-r-Gelatin Bi-Component Nanofibers by Coaxial Electrospinning,” Chem. Muter., 16. 34069 (2004). 68D. Li, A. Babel, S. A. Jenekhe. and Y . Xia, “Nanofibers of Conjugated Polymers Prepared by Electrospinning with a Two-CapilkdrySpinneret,” Adv. Mater., 16. 20624 (2004). 69D. Li, Y. Wang, and Y. Xia, “Electrospinning Nanofibers as Uniaxially Aligned Arrays and Layer-By-Layer Stacked Films,” Adv. Murer.. 16. 3614 (2004).

“D. T. Welna. J. D. Bender, X. L. Wei, L. G . Sneddon, and H. R. Allcock, “Preparation of Boron-Carhide/Carbon Nanofibers from a Poly(norborneny1decahorane) Single-Source Precursor Via Electrostatic Spinning,” Adv. Mater., 17,

70 D. Li, G . Ouyang, J. T. McCann, and Y. Xia, “Collecting Electrospun Nanofibers with Patterned Electrodes,” Nuno Letf.,5, 9 1 3 4 (2005). ”Y. Gu, D. Chen, and M. Jiao, “Synthesis and Electrochemical Properties of Nanostructured LiCo02 Fibers as Cathode Materials for Lithium-Ion Batteries,” J . Phys. Chem. B,, 109, 179014 (2005). 12 M. Y. Song, D. K. Kim, K. J. Ihn. S. M. Jo, and D. Y. Kim, “Electrospun Ti02 Electrodes for Dye-Sensitized Solar Cells,” Nanotechnology. 15, 1861-5

S5$ :2

(2004).

(2004).

(2005).

D. Li, J. T. McCann, and Y. Xia. “Use of Electrospinning to Directly Fabricate Hollow Nanofibers with Functionalized Inner and Outer Surfaces,” SmaN, 1, 8(2005). 63J. T. McCann, D. Li, and Y. Xia. “Electrospinning of Nanofiben with Core-Sheath, Hollow. or Porous Structures,” J. Mater. Chem., 15, 735-8 (2005).

64G. Loscertales. A. Barrero, M. Marquez, R. Spretz, R. Velarde-Ortiz. and G. Larsen. “Electrically Forced Coaxial Nanojets for One-Step Hollow Nanofiber Design,” J. Am. Chem. Soc., 126. 53767 (2004). 65 2. Sun, E. Zussman, A. L. Yarin, J. H. Wendorff, and A. Greiner, “Compound Core-Shell Polymer Nanofibers by Co-Electrospinning,” Adv. Marer., 15, 1929-32 (2003).

Nanotubes, Nanorods, and Nanowires

73M.Y. Song, Y. R. Ahn, S. M. Jo, D. Y. Kim. and J. P. Ahn, “Ti02 SingleCrystalline Nanorod Electrode for Quasi-Solid-State Dye-Sensitized Solar Cells,” Apjl. Pkys. Leff.,87, 113113 (2005). 4P. 1. Gouma. “Nanostructured Polymorphic Oxides for Advanced Chemosensors,” Rev. Adv. Mater. S c i , 5, 147-54 (2003). 15M.Sawicka, A. K. Prasad, and P. 1. Gouma, “Metal Oxide Nanowires for Use in Chemical Sensing Applications,” Sensor Left., 3, 3 1-5 (2005). 16D. Li, J. T. McCann, M. Gratt, and Y. Xia, “Photocatalytic Deposition of Gold Nanoparticles on Electrospun Nanofibers of Titania,” Chem. Phys. Left., 3!24+ 387-91 (2004). H.-M. Shang, Y.Wang, K. Takahashi, G. Z. Cao, D. Li, and Y. Xia, “Nanostructured Superhydrophobic Surfaces,” J . Mater. Sci., 40,3587-91 (2005). 0

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Nanocomposites and Nanostructures

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Chemical Precipitation Synthesis and Optical Properties of ZnO/Si02 Nanocomposites Huaming Yang,+ Yu Xiao, Kun Liu, and Qiming Feng Department of Inorganic Materials, School of Resources Processing and Bioengineering, Central South University, Changsha 410083, China

SiOz nanowires, made from natural chrysotile, were used to synthesize zinc oxide (ZnO)/SiOz composites by chemical precipitation. The as-prepared samples were characterized by X-ray diffraction, scanning electron microscopy, transmission electron microscopy (TEM), energydispersive X-ray spectroscopy, and Fourier transform infrared. Their optical properties were studied by a ultraviolet-vis spectrophotometer and a fluorescence spectrophotometer.Structural analysis revealed that the crystal size of ZnO crystallites is <20 nm, and the leached SiOz nanowires were amorphous. TEM analysis showed that the size of ZnO particles in SZ8 was mainly in the range from 15 to 20 nm and dispersed uniformly on the surface of SiOz nanowires. The photoluminescence spectra showed that ZnO/SiOz composites have stronger emitting intensity at the blue-green band than pure ZnO synthesized under the same reaction conditions. Therefore, the composites will be of great interest in the application of luminescence material. The as-prepared ZnO/SiOz composites can be used as photocatalysts for waste water treatment because they separate much more easily away from solution.

ZnO coated on other materials has been reported. Sun et a/.’’ have reported the attachment of well-crystallized ZnO nanoparticles onto 4-aminobenzene monolayer-grafted multi-walled carbon nanotubes (MWNTs) using water-in-oil microemulsions and have hoped to find applications in field-emission display devices. Zhu et a1.I6have reported ZnO nanoparticles on beaded MWNTs for ultrafast nonlinear optical switching. However, MWNTs are expensive and not easy to synthesize; therefore it is meaningful to find a substitute for MWNTs. Si et a/.17have reported a 2-3 nm ZnO shell coated on Fez03 particles through a hydrolysis process of Zn2+, with the result the as-prepared Fe203/ZnO-based sensors showed a higher response to various combustible gases. In this work, we put forward an easy method to load ZnO nanoparticles onto Si02 nanowires. The as-prepared ZnO/Si02 composites have a stronger emitting intensity than pure ZnO, which is of great interest in the application of luminescence material. In addition, the asprepared ZnO/Si02 composites can be used as a photocatalyst for waste water treatment because they separate much more easily away from suspension.

I. Introduction

11. Experimental Procedure

oxide (ZnO) has attracted a lot of attention in the past because it is a very important electronic and photonic material, with a wide direct band gap of 3.37 eV and a high binding excition energy of 60 meV. A great many morphologies of ZnO have been reported, such as nanowires, nanobelts, tetrapods, hexagonal nanoprisms, nanotubes, and nano~heets.’-~It is known that the size and morphology of particles can affect their properties. Some researchers have already performed some work on the optical properties of ZnO with different nanostructures.’ It is now necessary to find a simple synthesis route through which the morphology of ZnO is controllable. So far, apart from nanoparticles, one-dimensional ZnO nanowires are the most attractive morphology because they are found to be suitable for nanoscale applications.’ However, the category of two-dimensional ZnO nanowires is also significant.” It can be obtained by electrodeposition, solvothermal route or hydrothermal method, chemical bath deposition, and so Among them, chemical precipitation is a low cost and simple method; however, surfactants are always applied because they are effective in controlling the morphology of ZnO. In this paper, we introduce a simple method to synthesize ZnO nanosheets through chemical precipitation without surfactants but with SiO2 nanowires.

( I ) Synthesis of ZnO and ZnO/SiO1 Composites Si02 nanowires were prepared according to the literature.18 NaOH and ZnAcz were of analytical purity. Deionized water was used throughout this work. Two hundred milliliters 0.07M ZnAc2 was put into a beaker and stirred by a PTFE-coated magneton at room temperature. A 0.3M NaOH solution was dropped into the ZnAcz solution until the pH value was 8. The white precipitate was filtered and washed, and then dried at 110°C for 12 h. The as-prepared white powder was ground and heated at 400°C for 2 h. The end product was named Z8. For the preparation of ZnO/Si02 composites, 0.104 g Si02 nanowires were dispersed into 236 mL of deionized water by ultrasonic for 10 min, and then mixed quickly with 200 mL 0.07M ZnAc2. The mixture was stirred by a PTFE-coated magneton for 12 h at room temperature. Then, the 0.3M NaOH solution was dropped into the ZnAc2 solution until the pH value was 8. The white precipitate was filtered and washed, and then dried at 110°C for 12 h. After that, the white powder was ground and heated at 400°C for 2 h. The end product was named SZS8. Furthermore, SZ8 was prepared in the same way as SZS8 but without the 12 h prestirring process.

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Manuscript No. 23944. Received November 4, 2007; approved January 13, 2008. The work wits financially supported by the National Natural Science Foundation of China (50574102,50774095), the Program for New Century Excellent Talents in University (NCET-05-0695). and the Dissertation Innovation Fund of Central South University (134375212). ‘Author to whom correspondence should be addressed. e-mail: [email protected]. edu.cn

Nanocomposites and Nanostructures

( 2 ) Characterization Powder X-ray diffraction (XRD) patterns were obtained using a fully automatic XRD apparatus (Rigaku D/max 2550 VB+, Rigaku International Corp., Tokyo, Japan) with an acceleration voltage of 40 kV and an emission current of 300 mA. A scanning electron microscope (SEM, JEOL JSM-6360, JEOL Corp., Tokyo, Japan) was used for the morphology analysis of the samples. Transmission electron microscopy (TEM) analysis was performed on a JEM 3010 (JEOL Corp.), equipped with energy-dispersiveX-ray spectroscopy (EDX). Fourier transform 353

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Fig. 1. X-ray diffraction patterns of the as-prepared samples. (a) SZX, (b) SZSX, (c) Si02 nanowires.

infrared (FT-IR) spectroscopy of the samples was carried out using a Nicolet NEXUS 670 spectrophotometer with KBr discs (Nicolet Instrument Corp., Madison, WI). Ultraviolet (UV)-vis absorption spectra of the solid samples were measured using a Shimadzu UV-2450 spectrophotometer (Shimadzu Corp., Kyoto, Japan). Photoluminescence (PL) spectra of the solid samples were measured at room temperature using a Hitachi F4500 with a 290 nm filter (Hitachi Ltd., Tokyo, Japan). The excitation wavelength was 254 nm.

111. Results and Discussion

(I)

Structure and Morphology

Figure 1 illustrates the XRD patterns of SZ8, SZS8, and Si02 nanowires. In Fig. I(c), the broad peak, ranging from 15" to 30" in the 20 angle confirms the amorphous nature of the asprepared Si02 nanowires." In Fig. I@) and (b), all the diffraction peaks of SZ8 and SZS8 can be indexed to wurtzite-structured ZnO (PDF card no. 80-0074). Furthermore, no characteristic peak of amorphous SiOz was detected because the content of Si02 nanowires was much less than that of ZnO in both SZ8 and SZSS. The average size of each plane can be determined by the Sherrer equation (D= 0.9h/(@os0)).For SZ8, the values of (IOO), (002), and (101) are 22, 21, and 21 nm, respectively. Meanwhile, for SZS8, the values of (IOO), (002), and (101) are 18, 15, and 16 nm, respectively. Therefore, it is concluded that the long prestirring process (before adding NaOH solution) affected the crystal size of ZnO in ZnO/SiO2 composites. The morphology of SZSS has been studied by SEM. As shown in Fig. 2(a), only a few nanoparticles were found on the surface of curly Si02 nanowires; however, some new sheets appeared. In the l i t e r at ~ r e,the '~ authors considered that the intensity ratio of (100) reflection and (002) reflection relates to the shape of ZnO. According to the values of (IOO), (002), and (IOI), ZnO in SZS8 could not be in the nanosheet shape. The area in the black pane is magnified, as shown in Fig. 2(b). From Fig. 2(b), it can be clearly seen that the sheets are between 0.5 and 1 pm in diameter, while the thickness is <50 nm. As is known to all, Zn2+ ions hydrolyzed in water lead to the formation of Zn(OH)2 colloidal particles in a weak acidic solution. These Zn(OH)2 colloidal particles were positively charged in the weak acidic solution, and therefore they were easily adsorbed on the surface of the negatively charged SiOz nanowires during the stirring process." Meanwhile, after a long stirring process, the concentration of Zn2+ ions at the Si02/solution interface was higher (the concentration of Zn2+ ions in the solution is lower, which has been proved in our former experiment), perhaps forming a layered distribution of Zn2+ ions. We hypothesize that when NaOH solution was added, Zn2+ ions at the SO 2/ solution interface reacted with OH- ions to form Zn(OH)2 colloidal particles first, forming small-sized nanosheets. Then, with the increase of OH- ions, Zn2+ ions in the bulk solution

354

Fig. 2.

Scanning electron microscopy images of SZSX.

were transformed into Zn(OH), colloidal particles, which were adsorbed onto the brim of the nanosheets by electrostatic force. Finally, the nanosheets, as shown in Fig. 2(b), were formed. The morphology of SZ8 was also studied by SEM. As shown in Fig. 3, a large quantity of nanoparticles is dispersed uniformly on the surface of the Si02 nanowires compared with the SEM image of pure Si02 nanowires." Confined by the magnification, the machine is unable to see the loading situation very clearly. Therefore, the sample of SZ8 was further characterized by TEM in order to obtain clearer images. Figure 4(a) shows that nanoparticles ranging from 15 to 20 nm in diameter were loaded onto SiOz nanowires. Figure 4(b) is

Fig. 3. Scanning electron microscopy image of SZ8. Progress in Nanotechnology: Processing

Fig. 4. (a) Transmission electron microscopy image of SZX; (b) Energy-dispersive X-ray spectroscopy spectrum taken from the black pane area of (a); (c) Transmission electron microscopy image of SZ8; (d) High-resolution-transmissionelectron microscopy image enlarging the selected area in (c); (e) (HRTEM) image enlarging the nearby area of the selected area in (c).

the EDX spectrum taken from the selected area of Fig. 4(a), which only shows Zn, Si, and 0 elements (peaks of copper and carbon are not indexed in the figure). The EDX result confirms that the SiOz nanowires obtained from chrysotile are pure, and the acid-leaching method is effective. Fig. 4(c) is also a TEM image of SZ8, and the selected area was magnified as shown in Fig. 4(d). From Fig. 4(d), it is clearly seen that the Si02 nanowire is amorphous, which is consistent with the result of XRD. The nanoparticles show clear lattice fringes, and the measured lattice spacing is 0.248 nm, which corresponds to the (101) lattice plane of hexagonal ZnO crystal. The edges of the ZnO nanoparticles were further magnified, which confirms that the ZnO nanoparticles are well crystallized. From Fig. 4(a), it can be clearly seen that some ZnO nanoparticles are attached to the ZnO nanoparticles that are attached to the S O 2 nanowires. Because these two kinds of ZnO nanoparticles are of the same size, it is deduced that positively charged Zn(0H)z nanoparticles formed in the liquid first and then adsorbed onto the Si02 nanowires by electrostatic force. After the drying process and thermal treatment, the ZnO nanoparticles were attached to the Si02nanowires steadily by the Si0-Zn bond. The surface of silica nanowires was magnified as shown in Fig. 4(e). Some small ZnO crystallites, about 5 nm in diameter, were found as marked in white loops, two of which have been magnified. We think that these crystallites were not adsorbed but grew on the silica nanowires. It is reported that Zn2+can be adsorbed by SiOz through the following process2":

When the NaOH solution was added, the adsorbed Zn2+ hydrolyzed. With the continuous addition of NaOH solution, the Nanocomposites and Nanostructures

Zn(OH)2 particles grew. The region where ZnO nanoparticles grew is expected to be highly defective, making it difficult for ZnO nanoparticles to grow epitaxially.'6 Because the two ZnO/ SiOz samples prepared are fully alike, the volumes of the same quantity of SZS8 and 2 8 were compared, as shown in Fig. 5. The increased volume of SZS8 compared with 28 is due to the addition of Si02 nanowires.

( 2 ) FT-IR Spectroscopic Analysis FT-IR spectra of the samples are shown in Fig. 6. In Fig. 6(a), the FT-IR spectrum of ZnO displays two characteristic peaks at 499 and 430 cm-'. The FT-IR spectrum of Si02 nanowires shows three characteristic peaks at 804, 945, and 1091 cm-' which is attributed to the ring structure of Si04 tetrahedra:' terminal Si-OH deformation band, and Si-GSi symmetric stretching mode,22 respectively. Compared with the spectrum of Si02 nanowires, the s ctra ,with of SZS8 and SZ8 are different between 700 and 1350 cm- p". decreasing peak intensity at 1091 cm-I, increasing peak intensity

Fig.5. Digital photo of SZSX (left bottle) and 28 (right bottle).

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Fig.6. (a) Fourier transform infrared (FT-IR) transmittance spectra of 28, SZS8, SZS8, and Si02 nanowires and FT-IR absorbance-fittingspectra of (b) S O z nanowires, (c) SZS8, and (d) SZ8 at low range.

at 945 cm-I, and the loss of a peak at 804 cm-I. In the literature,22the authors considered that the silylating agents grafting onto the silica surface resulted in the loss of the Si-OH deformation mode originally present at 945 cm-'. Actually, the peak of the Si-0H deformation mode became just a shoulder peak. In order to confirm the reaction between the surface of SiOz nanowires and ZnO, the FT-IR spectra of Si02,SZS8, and SZ8, ranging from 700 to 1350 cm-', were deconvoluted. The fast Fourier transform-filtering algorithm is used to create a smooth spectrum, and the second derivative procedure is used to find the hidden bands. The process of fitting work had been carried out until a high fitting degree (coefficient is over 0.99) was reached. As shown in Fig. 6(b), the absorbing band of Si02between 1000 and 1350 cm-l can be deconvoluted into two peaks. The peak at 1184 cm-l may be due to asymmetric vibrational Si-0- stretching.23As shown in Figs. 6(c) and (d), an obvious shift of the peak at 800 cm-l to a higher frequency appears in both SZS8 and SZ8, with almost the same value. This may be caused by thermal treatment, which may lead to the rearrangement of the ring structure of the Si04 tetrahedra.24 A comparison in the intensities of the deconvoluted peaks is shown in Table I. It is

Table I.

Results of the FT-IR Deconvolution

FT-IR, Fourier transform infrared

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found that the intensity value of the two higher frequency peaks decreased by almost the same amount for both SZ8 and SZS8 because less S O 2quantity is contained in the samples for FT-IR characterization. Meanwhile, this also confirms that corresponding bonds exist in the internal Si02 and, therefore, will not be affected by the reaction with ZnO. Compared with the intensity ratio of the two higher frequency peaks, it is found that the intensity ratio of the two lower frequency peaks increases relatively. In the l i t e r a t ~ r ethe , ~ ~authors considered the peak at 950 cm-l as symmetric vibrational Si-0-Ti stretching, and the intensity was related to the dispersion of Ti02 over the silica surface. Therefore, we deduce that the increase in peak intensity at 936 cm-' was brought on by the reaction of Si-0H and ZnO particles.

(3) UV-vis and P L Analysis Comparison of the UV-vis absorption spectra samples is shown in Fig. 7. The value of a direct band gap of the semiconductor determines its optical property and application. Much literature has shown that the method calculates the direct band ap of ZnO. The absorption onset can be obtained by plotting In ([*/Io) versus energy h ~ . ~Because ' the absorption edge lines were close in the same figure, they were separated in order to draw the tangents. As shown in Fig. 7, the band gap of ZnO/Si02 composites is bigger than pure ZnO, while that of SZS8 is bigger than that of SZ8. This is because the crystal size of SZS8 is smaller than that of SZ8, as confirmed by XRD results, and we know that the band gap value of smaller crystals is higher. We further carried out room temperature PL measurement to examine the quality of the products. Figure 8 shows the PL spectra of all the samples at room temperature. The PL spectrum of ZnO crystals always exhibits a UV emission peak at 380 nm, which is attributed to the recombination of free excitons. However, it is not observed in all the PL spectra of the samples,

8

Progress in Nanotechnology: Processing

prestirring process before adding NaOH solution, some ZnO existed in the form of nanosheets with a crytal size of about 15 nm. The as-prepared ZnO/Si02 composites have bigger band gap energy and stronger emitting intensity at the bluegreen band than pure ZnO synthesized under the same reaction conditions. Therefore, ZnO/Si02 composites would play an important role in the PL material. The as-prepared ZnO/Si02 composites can also potentially be used in the treatment of polluted water because they are innocuous and easily separated away from water.

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Fig. 7. Ultraviolet-vis spectra of the samples prepared under pH = 8.

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Fig.8. Photoluminescence emitting spectra of the samples prepared under pH = 8.

ID. Banerjee, S. H. Jo, and Z. F. Ren, “Enhanced Field Emission of ZnO Nanowires,” Adv. Muter., 16, 2028-32 (2004). 2E. C. Greyson, Y. Babayan, and T. W. Odom, “Directed Growth of Ordered Arrays of Small-Diameter ZnO Nanowires,” Adv. Mater., 16, 1348-52 (2qp4).

W. 2. Wang, B. Q. Zeng, J. Yang, B. Poudel, J. Y. Huang, M. J. Naughton, and 2. F. Ren, “Aligned Ultralong ZnO Nanobelts and Their Enhanced Field Emission,” Adv. Muter., 18, 3275 8 (2006). 4A. B. DjuriSiC, W. C. H. Choy. V. A. L. Roy, Y. H. Leung, C. Y. Kwong. K. W. Cheah, T. K. G. Rao, W. K. Chan, H. F. Lui, and C. Surya, “Photoluminescence and Electron Paramagnetic Resonance of ZnO Tetrapod Structures,” Adv. Funct. Mater., 14, 85(2004). ’X. M. Sui, Y. C. Liu, C. L. Shao, Y. X. Liu. and C. S. Xu, “Structural and Photoluminescent Properties of ZnO Hexagonal Nanoprisms Synthesized by Microemulsion with Polyvinyl Pyrrolidone Served as Surfactant and Passivant,” Chem. Phys. Lert.. 424, 3 W 4 (2006). 6Y. Sun, G . M. Fuge, N. A. Fox, D. J. Riley, and M. N. R. Ashfold, “Synthesis of Aligned Arrays of Ultrathin ZnO Nanotubes on a Si Wafer Coated with a Thin ZnO Film,” Adv. Muter., 17, 2477-81 (2005). ’A. Eftekhari, F. Molaei, and H. Arami, “Flower-Like Bundles of ZnO Nancsheets as an Intermediate between Hollow Nanosphere and Nanoparticles,” Marer. Sci. Eng. A , 437, 44650 (2006). *A. B. DjuriSiC and Y. H. Leung, “Optical Properties of ZnO Nanostructures,” Sma//, 2, 94461 (2006). 9C. So&,A. Zhang, B. Xiang, S. A. Dayeh. D. P. R. Aplin, J. Park, X. Y. Bao, Y. H. Lo, and D. Wang, “ZnO Nanowire UV Photodetectors with High Internal Gain,” Nono Lett., 7, 1003-9 (2007). ‘OS. J. Chen, Y. C. Liu. C. L. Shao. R. Mu, Y. M. Lu, J. Y. Zhang, D. Z. Shen, and X. W. Fan, “Structural and Optical Properties of Uniform ZnO Nanosheets,” Adv. Mater., 17, 586-90(2005). “L. F. Xu, Y. Guo, Q . Liao. J. P. Zhang, and D. S. Xu, “Morphological Control of ZnO Nanostructures by Electrcdeposition,” J. Phys. Chem. B, 109, 1351922 2005)

which agrees with some l i t e r a t ~ r e . ~ ~Mainly, .~’ four emission peaks are observed in SZ8 and SZS8: a violet emission peak at around 419 nm, two blue emission peaks at around 450-475 nm, and a green emission at around 536 nm. The blue emission may result from the recombination of oxygen vacancies with oxygen interstitials or other defects, while the green emission may result from the recombination of a photogenerated hole with the sin 1 ionized oxygen vacancies at the surface of ZnO particles.’ !Y’ The PL spectrum of pure ZnO shows only three emission peaks, which are almost at the same position as the corresponding peaks of ZnO/Si02 samples. By comparing the intensity of the peak at 419 nm with the peak at 473 nm or the peak at 536 nm, it is found that the ratio of SZ8’s peak is smaller than SZS8’s peak. From the morphologies of ZnO existing in SZ8 and SZS8, we deduce that the specific surface area of ZnO in SZ8 is higher. Because oxygen vacancies are mainly on the surface of ZnO particles, the higher specific surface area leads to a higher concentration of oxygen vacancies and consequently a stronger intensity at the blue and green emission peaks.29330 Hence, we deduce that the increase of emission intensity is attributed to the increase of the specific area of ZnO in ZnO/ Si02 composites.

IV.

Conclusions

ZnO/Si02 composites were synthesized by chemical precipitation. Without a prestirring process, most of the ZnO nanoparticles, ranging from 15 to 20 nm in diameter, dispersed uniformly on the surface of the Si02 nanowires. With a Nanocomposites and Nanostructures

“S. Kar, A. Dev. and S. Chaudhuri, “Simple Solvothermal Route to Synthesize ZnO Nanosheets, Nanonails. and Well-Aligned Nanorod Arrays,” J. Phys. Chem. B, 110, 17848-53 (2006). ”D. Scarano, S. Bertarione, F. Cesano, J. G. Vitillo. and A. Zecchina, “Plate-Like Zinc Oxide Microcrystals: Synthesis and Characterization of a Material Active Toward Hydrogen Adsorption.” Cata/. Today, 116. 433-8 (2006).

I4V.Pachauri, C. Subramaniam, and T. Pradeep, “Novel ZnO Nanostructures Over Gold and Silver Nanoparticle Assemblies,” Chem.Phys. Left., 423, 2 4 M (2006).

”J. Sun, L. Gao, and M. Iwasa, “Noncovalent Attachment of Oxide Nanoparticles onto Carbon Nanotubes using Water-in-Oil Microemulsions,” Chem. Commun., 832-3 (2004). I6Y. W. Zhu, H. 1. Elim. Y. L. Foo, T. Yu, Y. J. Liu, W. Ji, J. Y. Lee, 2. X. Shen. A. T. S. Wee, J. T. L. Thong, and C. H. Sow, “Multiwalled Carbon Nanotubes Beaded with ZnO Nanoparticles for Ultrafast Nonlinear Optical Switching,” Adv. Marer.. 18, 587-92 (2006). ”S. F. Si, C. H. Li, X. Wang, Q . Peng, and Y. D. Li, “Fe20,/Zn0 Core-Shell Nanorods for Gas Sensors,” Sens. Actuators B. 119, 52-6 (2006). ‘*K. Liu, Q. M. Feng, Y. X. Yang, G. F. Zhang, L. M. Ou, and Y. P. Lu, “Preparation and Characterization of Amorphous Silica Nanowires from Natural Chrysotile.” J. Nan-Cryst. Solids, 353, I S3&9 (2007). I9Zs. Csogor, M. Nacken, M. Sameti, C.-M. Lehr, and H. Schmidt, “Modified Silica Particles for Gene Delivery,’’ Mater. Sci. Eng. C, 23, 93-7 (2003).

’%. Janusz, J. Patkowski, and S. Chibowski, “Competitive Adsorption of Ca2’ and Zn(I1) Ions at Monodispersed SiO2/Electrolyte Solution Interface.” J. Co/. fnt. Sci., 266,25948 (2003). ’IT. M. H. Costa, M. R. Gallas, E. V. Benvenutti, and J. A. H. da Jornada. “Infrared and Thermogravimetric Study of High Pressure Consolidation in Alkoxide Silica Gel Powders,” J. Non-Crysf.So/ids, 220, 195-201 (1997). 22M. G. Fonseca. A. S. Oliveira, and C. Airoldi. “Silylating Agents Grafted onto Silica Derived from Leached Chrysotile,” J. Colloid Interface. Sci., 240, 533-8 (2001).

23M. Cozzolino, M. Di Serio, R. Tesser, and E. Santacesaria, “Grafting of Titanium Alkoxides on High-Surface Si02 Support: An Advanced Technique for the Preparation of Nanostructured Ti02/Si02 Catalysts,” Appl. C a r d A, 325, 256-62 (2007).

’%. Diaz, R. Perez-Hernindez, A. Gomez-Cortes, M. Benaissa, R. Mariscal, and J. L. G. Fierro, “CuO-Si02 SoCGel Catalysts: Characterization and Catalytic Properties for NO Reduction,” J. Coral., 187. 1-14 (1999). 357

25S.Monticone, R. Tufeu. and A. V. Kanaev, “Complex Nature of the UV and Visible Fluorescence of Colloidal ZnO Nanoparticles,” J. Phys. Chem. B, 102, 28:6M2 (1998). F. H. Zhao, W. J . Lin, M. M. Wu, N. S. Xu. X. F. Yang, Z. R. Tian, and Q. Su. “Hexagonal and Prismatic Nanowalled ZnO Microboxes,” Inorg. Chem., 45, 3256-60 (2006). “J. P. Cheng, X. B. Zhang, X. Y. Tao, H. M. Lu, 2. Q. Luo, and F. Liu, “FineTuning the Synthesis of ZnO Nanostructures by an Alcohol Thermal Process,” J. Phys. Chem. B, 110, 10348-53 (2006).

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28 Z. G. Chen, A. Z. Ni, F. Li, H. T. Cong, H.M. Cheng, and G . Q. Lu, “Synthesis and Photoluminescence of Tetrapod ZnO Nanostructures,” Chem. Phys. Leu., 434, 301-5 (2007). 29H. M. Xiong, D. P. Liu, Y. Y . Xia, and J. S. Chen, “Polyether-Grafted ZnO Nanoparticles with Tunable and Stable Photoluminescence at Room Temperature,” Chem. Murer., 17, 3 0 6 2 4 (2005). ‘OF. Xu, K. Yu, Q. Li, and 2. Q. Zhu, “Two-Dimensional Growth and Field Emission Properties of ZnO Microtowers.” J. Phys. Chem. C, 111, 4099-104 (2007). 0

Progress in Nanotechnology: Processing

Low-Temperature Processing of Dense HydroxyapatiteZirconia Composites Yougojoti Nayak, Raghunath Rana, Swadesh Pratihar, and Santanu Bhattacharyya* National Institute of Technology, Department of Ceramic Engineering, Rourkela-7G9008, India

Hydroxyapatite (HA)-YTZP (2, 5, 7.5, and 10 wt% Zr02) composite powders prepared from inorganic precursors were characterized by FTIR, DSC/TG, XRD, and TEM. The calcined powders had HA and dc-Zr02, which undergo structural changes between 650°C and 1050°C. TEM of calcined powder showed larger HA particles (100nm) and smaller ZrOz particles ( 5 5 0 nm). HA and HA-2wt% ZrO2-sintered samples had 98% density and it was (90-95%) for HA-5, 7.5, and 10 wt% Zr02. The bending strength of HA-2wt% ZrO2 composites was 72 MPa. The grain sizes of HA showed a refinement with Zr02 addition.

Introduction Hydroxyapatite (HA) ceramics have widespread applicability in the biomedical field partially for repair or replacement of bone tissues. Owing to the low strength and toughness of HA, its use is limited mainly to the nonload-bearing components, for example, small unloaded or low loaded implants, coating on metal implants, and as a bioactive filler powder or as a second phase in polymer-based composites. 1-3 For load-bearing applications, the research on the use of HA as a bioactive filler in ceramic and metallic matrix composites did not succeed due to the deleterious reactions taking place between HA and the matrix phase.*-' However, the use of HA in load-bearing parts can be explored, provided the strength and toughness of HA can be increased.' Among many oxide reinforcements, ZrO2 is an important one, which has been used to increase the strength and toughness of Al203, mullite, spinel, and many other '[email protected] 0 2008 The American Ceramic Society

Nanocompositesand Nanostructures

ceramic materials. High sintered density and ultra fine particles (in the nanorange), along with the use of stabilizers, will ensure Z r 0 2 mostly in the tetragonal phase, leading to improved mechanical properties of the composites. A number of approaches have been attempted to prepare HA-Zr02 composites containing a higher fraction of Z r 0 2 (20-40 vol%).'-' In these composites, the major problems encountered were the densification of the composites and the decomposition of HA to TCP.' Actually, the decomposition of HA to T C P is accelerated at higher sintering temperatures (> 1300°C). With Z r 0 2 addition, the decomposition temperature is further lowered (in the range 1000- 1050"C).'0 Efforts to achieve densification at a lower sintering temperature involved hot pressing/hot isostatic pressing, spark plasma sintering,' and use of additives (calcium fluoride (CaF,)),' etc. All the above composites in the HAZrO2 system involved a high volume fraction of Z r 0 2 addition. On the other hand, the available literature on the processing, densification behavior, and mechanical properties of HA-Zr02 composites containing a low volume fraction of ZrO2 ( I 20 ~ 0 1 % )is comparatively

'

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limited.8,'2,'3 Kim et aL8 studied the effect of CaF2 addition on the reaction sintering and mechanical properties of HA-Zr02 composites containing 20 vol% Z r 0 2 . Addition of CaF2 promoted densification and retarded the HA to T C P decomposition. The densification temperature also decreased with CaF2 addition. The highest flexural strength was 125 MPa and fracture toughness was 2 MPa/m'/2, which were higher than pure HA. Pyda et all2 studied the effect of chemical composition and morphology of zirconia particles on the properties of HA-Zr02 particulate composites containing 11.9 vol% Ca or Y-stabilized ZrOz. The composites were hot pressed in the temperature range 11501250°C in an argon atmosphere. It was observed that isometric ZrO2 crystallites incorporation to the HA matrix increased both the strength (18OMPa) and the toughness (1.1 MPa/m'/2), which were higher than that of pure HA. Ahn et studied the effect of nano zirconia reinforcement on the strength of hotpressed HA-Zr02 composites containing 1.5-8 wt% Z r 0 2 . They observed the highest strength and hardness at 3 wt% Z r 0 2 , which subsequently decreased at higher Z r 0 2 loading. They also reported that a lower volume fraction of Z r 0 2 addition will help to retain both HA and t-Zr02 phases, which may be due to the combined effect of matrix constraint and uniform dispersion of fine zirconia particles. Thus, it is expected that HAZrO2 composites containing a lower volume fraction of t-Zr02 may yield a transformation-toughening effect that can lead to improved mechanical properties of HAzirconia composites.7,8,12-17 The aim of the present investigation is to observe the effect of low volume percent TZP addition on the density, strength, and microstructure of HA-Zr02 composites. An attempt has also been made to analyze the observed results with respect to phase analysis and microstructure of the samples for different HA-Zr02 composites. The exact effect of Z r 0 2 addition (both at low and high Z r 0 2 loading) cannot be understood unless a systematic investigation in this system is carried out.

Experimental The HA-zirconia composite preparation consisted of two parts: (a) Preparation of nano-TZP (2.5 mol% Y2O3) powder using the precipitation method.

360

(b) Preparation of an HA-zirconia composites (containing 2, 5, 7.5, and 10wtYo TZP) by the reverse strike precipitation method, denoted as HZ2, HZ5, HZ7, and HZ1 0, respectively. The precursors used for TZP (2.5 mol% Y2O3) were ZrOC12 . 8 H 2 0 (Loba Chemicals, India) and Y2O3 (Indian Rare Earths Limited, India). The concentration of ZrOCl2 solution was 0.75 mol/L in which the requisite amount of Y2O3 powder was dissolved. Precipitation was carried out using an N H 4 0 H (1:l) (Oster Chemicals, India) solution at p H 10. The precipitates were washed with water to remove chloride ions, followed by drying and calcination at 850°C/2 h to obtain TZP powder. The outline of the powder preparation steps is shown in Fig. la. The HA-Zr02 composite powders were prepared from an equimolar aqueous solution (1.63 mol/L) of both Ca(NO&. 2 H 2 0 (Oster Chemicals, India) and (NH4)2HP04 (Nice Chemicals, India). These two solutions were mixed together, along with the addition of a few drops on H N 0 3 to obtain a colorless solution. In the usual method of HA-Zr02 composite powder preparation, N H 4 0 H is added to the above-mixed solution. However, in the present investigation, the mixed solution was added dropwise to a beaker containing calcined TZP powder dispersed in N H 4 0 H , while the beaker was being vigorously stirred. This allowed better and uniform reaction of Ca(N03)2and (NH,J2HP04 with NH40H and for the present investigation, this precipitation process has been termed as the "reverse strike precipitation" method. The outline of the precipitation process is shown in Fig. 1b. The as-precipitated amorphous powders having varying proportions of TZP were subjected to hot water and propanol washing to prevent particle agglomeration on drying of these powders. It has been reported that the presence of hydroxyl and other ions (e.g., CI-) affects the crystallization and densification behavior of many ceramic powders. l 3 The washed powders were oven dried for 24h. The FTIR spectra of the as-dried precipitated powder of HA and HA-TZP were observed in a Perkin Elmer FTIR spectrophotometer (Spectrum RX- 1). The spectrum was taken in the KBr pellet mode in the wave number range 4400-400 cm-'. A small part of the dried powder was subjected to thermal analysis in the DSC/TG mode (Netzsch 402 C) till 1200°C at a heating rate of 1O"C/ min in air. The dried powders were calcined in different temperature ranges from 650°C to 850°C/2h in air. The phase analysis of the calcined powders and sintered

Progress in Nanotechnology: Processing

length = 20 mm) as well as in diametral compression of the cylindrical disk (12.5 mm diameter) at a crosshead speed of 0.2 mm/min. The microstructures of the sintered and polished pellets were studied by SEM USM6480LV). In all the above characterizations of the HATZP composite, HA was used as the control sample.

I

Addiumof*xccrr Ni,oH l o imnasothopH (9-10) I

I

Results and Discussion

I +

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1.63 M of

I

1.63 M of

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2.5 mow Y-TZP powder dispasedin (1:l) MI3solution kept in

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FTIR The FTIR spectra of pure HA (Fig. 2) show all the peaks corresponding to (OH)- and stretching, and no extra peak was observed. The broad peak between 3750 and 2900 cm-' corresponds to the hydrogen-bonded 0-H group stretch of HA and water. Thus, the processed HA was stochiometric. The three strong peaks at around 1060, 570, and 606cm-' could be attributed to the P-0 vibration modes. The band at 570 and 606cm-' corresponds to the bending mode of and the strong band at 1060 and 963 cm-' is due to the vibration of the group. The absence of broadening of the peak at 1060cm-' further confirms that no decomposition of HA to P-TCP occurs at 850°C. The decrease in broadness of the peak between 3750 and 2900cm-' with an increase in calcination temperature indicates the removal of absorbed H20.

Fig. I . (a) Flow diagram f i r the synthesis of 2.5Y-TZP (6) flow diagram f i r the synthesis of hydroyapatite-(2, 5, 7.5, and 10 wt%).

pellets was studied by XRD (Philips PW1830 Holland, the Netherlands). The relative amount of the different phases present in the sintered pellet was calculated from the relative XRD peak intensities of the different phases. The particle morphology was studied by TEM (Philips CM 200). The densification behavior was studied in a dilatometer (Netzsch DL 402 C). The powders were uniaxially compressed at 280 MPa using 3 wt% PVA as the binder. The density of the sintered samples was measured by the Archimedes principle using kerosene as the immersion liquid. The strength of the sintered composites was measured using three-point bending (span

Nanocomposites and Nanostructures

I

I

4000

I

3500

I

3000

I

2500

I

2000

IV

1500

I

1000

500

Wave number (cm-')

Fig. 2. FTIR spectra ofhydroyapatite as a firnction of calcination temperature.

361

DSC/TG The DSClTG curve of the precipitated powders of HA is shown in Fig. 3a. The curve shows a total weight loss of 5.78%, which rakes place in two stages. It has a broad but sharp endothermic peak in the temperature range 3O-10O0C, with a peak at 70°C. This endothermic peak is associated with a weight loss of 3.55%, which could be related to the loss of adsorbed water from HA. Following this peak, there is a broad and diffused exothermic peak in the temperature range 250400°C. This exothermic peak also shows a gradual weight loss of about 2.23%. X-ray diffraction of the precipitated powder heated at 400°C shows peaks corresponding to HA only. Thus, this diffuse exothermic

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Fig. 3. DSC/TG curve of (a) hydroxyapatite powder, (b) HZ2 powder.

362

peak is due to the crystallization of HA from the precipitated amorphous powder. Figure 3b shows the DSC/TG curve of the precipitated HZ2 powder. The nature of the curve and the total weight loss (4.74%) is comparable with that of pure HA (5.78%). This curve also shows one broad endothermic peak in the temperature range 30-1 OO'C, with a peak at about 70"C, along with a weight loss of 3.21% corresponding to the loss of adsorbed water. The weight loss is lower than pure HA as the amounr of HA is less by 2%. Similarly, the second phase of weight loss from 100°C to 400°C is also lower in HZ2 in comparison with pure HA. However, the exothermic peak of HZ2 is at a lower temperature and it is sharper as compared with pure HA. This implies that the addition of TZP reduces the crystallization temperature and also improves the crystallization behavior of HA. However, no further change in either the crystallization behavior or the crystallization temperature could be observed in HA with a still higher number of TZP (viz. HZ5, HZ7, and HZ10) samples. Further, the DSC/TG curves of both HA and HATZP powders show a broad exothermic peak in the temperature range 800-1000°C. This exothermic peak is not associated with any weight loss. Therefore, the peak should correspond to phase changedordering in HA and/or ZrOz. In order to validate the DSC curve, the positions of the first three peaks of HA (or triplets) (d = 2.81, 2.72, and 2.78A0, respectively) were monitored as a function of calcination temperature. The results are shown in (Fig. 4b), where the left-hand pattern (Pattern A) shows the complete XRD profile of HA and HA-TZP composites, and the right-hand pattern (Pattern B) shows the "zoomed-in" pattern of the same in the 28 range 30-35". It is clear from the patterns that at 650"C, the triplets are not well separated and only two peaks (d = 2.8 1, 2.72A") are distinguishable. With an increase in the calcination temperature to 850"C, besides the above two peaks a third peak (as split peak) appears at d = 2.78. Finally, at 1050"C, all the three peaks could be clearly distinguished. It could also be seen from the zoomed-in pattern that the peak position shifts toward a lower d value and the peak intensity and sharpness increase. Thus, we can infer that at 65OoC, H A is poorly crystallized (disorder state) and it becomes perfectly crystalline at 1050°C. This phase change from poorly crystallized HA to perfectly crystallized HA is indicated as up a broad exothermic peak in the temperature 800-1000°C. A similar observation was made

Progress in Nanotechnology: Processing

HA is stable up to 850°C. The t-ZrO2/c-ZrO2 peaks could be detected from H Z 2 in Zr02-containing samples (HZ2 to HZ10). The intensity of Z r 0 2 peak increases at a higher Z r 0 2 content. No other phases like TCP, CaZr03, or m-ZrO2 could be detected. As discussed in the DSC/TG section, with an increase in the calcination temperature, the "d" value of HA peaks shifted to a lower value (Fig. 4b). From the XRD pattern, it is clear that in the calcined powder, only HA and t-Zr02 are present. No other phases like TCP, CaZr03, or m-Zr02 could be detected. The intensity of t-Zr02 increases with an increase in wt% of zirconia addition.

Particle Morphology of Calcined Powder Figures 5a, b shows the TEM photograph of calcined HA and HZ2 powder, respectively. The small and spherical particles indicate Z r 0 2 , while the lighter ones indicate HA. It could be seen that HA particles are both spherical and cube shaped in the size range 50-100 nm. Mostly, the Z r 0 2 particles are seen on the surface of HA particles. Some clusters of Z r 0 2 are also seen in the top part of the Fig. 5b.

Densifcation Behavior and Phase Analjsis of Sintered Compacts Fig.4. (a) XRD analysis of bydroxyapatite (HA), HA-ZrO, compositepowder calcined at 850aC/2b, (b) expanded region of XRD pattern to show the peak sbifing of HA.

from the XRD pattern of calcined HA-TZP powders. Moreover, the presence of Z r 0 2 also contributes to an exothermic peak because t-Zr02 prepared from an amorphous powder undergoes the following phase change: amorphous-cubic-tetragonal. It has been observed in our earlier study18 that the cubic-tetragonal phase change in Z r 0 2 takes place around 850°C. This could also be observed as an exothermic peak in DSC. Thus, in HA-TZP powder, we could observe two distinct peaks: one due to HA phase change and the other due to Z r 0 2 phase change.

Phase Analysis of Calcined Powder Figure 4a shows the XRD pattern of calcined HA and HA-TZP (HZ2, HZ5, HZ7, and HZ10) composite powders calcined at 850°C. The pattern shows that

Nanocomposites and Nanostructures

Figure 6 shows the nonisothermal sintering behavior of HA and HZ2, HZ5, HZ7, and HZlO composites. However, a significantly lower density is observed at 7.5 and 10wtYo TZP additions. The reduction in density at 7.5 and 1Owt% TZP occurred due to the higher sintering temperature of TZP and hence could not be densified at 1250°C. However, for HZ2, it appears that some other mechanism is playing a role in improving the densification. The relative sintered density at 1250°C is given in Table 11. The XRD pattern (Fig. 7) shows phases in the sintered samples of HA and HZ2, HZ5, HZ7, and HZ10. The XRD pattern shows that beyond 2wt% TZP additions, the sintered composites also contain P-TCP and CaZrO3. In HZ2, although the shrinkage starts at nearly the same temperature as that of HA, HZ2 has higher shrinkage in comparison with HA. This implies that the sintered density of HA increases on addition of 2 wt% TZP.

Microstructural Development on Sin tering The HZ2, HZ5, HZ7, and HZlO samples were sintered in air in the temperature range 1150-1250°C

363

Fig. 5. TEM of (a) hydroxyapatite (HA) powder and (6) HZ2 powder calcined at 850°C/2h.

for 2 h. The microstructure development of different HA-TZP composites as a function of varying TZP contents was studied for samples sintered at the highest temperature. Figures 8a-c shows the microstructure of different HA-TZP composites: (a) HZ2, (b) HZ7, and (c) HZ1O. All these samples were sintered at 1250°C for 4 h. For comparison, the microstructure of pure HA is also shown in Fig. 8d, which shows a mixed grain size for HA. The larger grains of HA are in 4-6pm, while

the smaller grains are in the range 1-2pm. Figure 8a shows that H Z 2 samples have a dense microstructure with an HA grain size in the range 2 - 4 p m , although some larger grains of 5-6pm could also be seen. The microstructure also contained void spaces and bright small grains (< 1 pm) distributed in the matrix. It appears that many of these void spaces were created due to the pullout of ZrOz grains during specimen preparation. The microstructure of HZ7 (Fig. 8b) clearly shows

Fig. 6. Nonisothermal densiJcation curve of bydroxyapatite (HA), HA-ZrO, composites.

Fig. 7. XRD analyJis of hydroxyapatite (HA), HA-Zr02 composites sintered at 1250" C/4h.

364

Progress in Nanotechnology: Processing

Fig. 8. SEM of hydroxyapatite (HA), HA-ZrO2 composites sintered at 125O"C/4h (a) HZ2, (b) HZ7, (c) HZlO, and (d) HA.

the effect of TZP addition on the grain size distribution of HA. In this sample, although some large grains (- 5 pm) are present, many small grains (1-2 pm) could also be seen. The ZrOz pullout during polishing is evident and some ZrOz particles are also present as clusters. The microstructure is further refined in HZlO (Fig. 8c). Although this microstructure had a mixture of a large grain size ( - 5 p m ) and very small grain size (0.5 pm), the distribution became narrow. The grain size distribution is shown in Table 11. The bright clusters of ZrOz grains can be seen near the multiple grain junctions. The microstructure (Figs. 8a-c) thus shows that addition of TZP to HA refines the grain size and produces an increased population of finer grains. However, the distribution of TZP is not uniform and it produces grain clusters. The TZP cluster size increases with increasing weight fraction of TZP. The effect of TZP addition on the grain size distribution is shown in Table I.

Table I.

Mechanical Properties Table I1 lists the diametral compression as well as the bending strength of sintered HA and HA-TZP composites. The table shows that following the density trend, the strength of HZ5, HZ7, and HZlO decreases with increasing TZP content. However, the strength increases significantly for H Z 2 composites. This is an interesting observation and needs more careful and detailed study. As already stated, this study only deals with the effect of TZP addition on the density, strength and microstructure. The exact effect of low TZP addition requires detailed TEM analysis, which will be taken up in a future study. Although the microstructure indicates grain refinement with TZP addition, the addition of TZP also increases the porosity due to incomplete densification due to the higher sintering temperature of TZP. Thus, the mechanical properties decrease with increasing TZP content.

Grain Size Distribution of Hydroxyapatite (HA), HA-Zr02 Composites Sintered at 125OoC/4h ~

HA Grain size (pm) Average Mean

Coarse

HZ2 Fine

Coarse

HZ7 Fine

Coarse

HZlO Fine

Coarse

Fine

3.63f1.15 1.13k0.3 3.64f1.76 0.93f0.36 2.7011.09 0.79f0.39 1.57f0.45 0.62k0.24 2.48 k 1.49 2.73 f 1.94 1.50+ 1.17 0.99 f0.57

Nanocomposites and Nanostructures

365

Table 11.

Density and Mechanical Properties of Sintered Compact of Hydroxyapatite (HA) and HA-Zr02 Composite

Relative Diametral Three-point sintered compression bending Composition density (Yo) strength (MPa) strength (MPa)

HA HZ2 HZ5 HZ7 HZlO

98.0 99.6 96.3 92.4 90.0

13.0f 1.03 34.5 k 2.763 11.3f1.04 17.0f 1.362 13.0 L- 0.907

35f 1.2 72 3.6 31 f 1.55 40f2 36L- 1.75

Phases in sintered HA sample (volo/o) HA

t-ZrOz

CZ

TCP

100 96.50 64.9 71.2 79

3.50 11.3 12.7 14.50

-

-

-

17.6 11 6.5

6.2 5.1 Tr.

Conclusions

5.

M. Knepper, 8. Milthrop, and S. Moricca, “Interdiffusion in Short-Fibre Reinforced Hydroxyaparite Ceramics,” J Mater. Sci. Mater. Med., 9 [ 101

HA and its composite with zirconia containing 2, 5, 7.5, and 10wtYo TZP were prepared by the reverse strike precipitation method. A DSC/TG study revealed that addition of TZP to HA reduced the decomposition

6.

Y.-M. Kong, S. Kim, and H.-E. Kim, “Reinforcement of Hydroxyapatite Bioceramic by Addition of Z r 0 2 Coated with A203,”J. Am. Ceram. Soc., 82

7.

2.Shen, E. Adolfsson, M. Nygren, L. Gao, H . Kawaoka, and K. Niihara,

589-596 (1998).

weight loss and promoted low-temperature crystallization of HA. TEM analysis of calcined powders showed spherical particles of TZP dispersed in cuboidal/spherical HA grains. The sintered density of HA increases with 2wt% TZP addition and it reduces with further addition of TZP. XRD revealed that in the sintered samples, only HA and t-Zt02 were present at 2wt% TZP addition. For a higher TZP content (5, 7.5, and 10wt%), both CaZr03 as well as a-TCP could also be seen, along with HA and t-Zr02. The bending strength of HA (35 MPa) increases to 70 MPa at 2 wt% T Z P addition, followed by a decrease at a higher TZP content. SEM microstructures show a grain refinement of HA with TZP addition. The decrease in strength at higher TZP addition could be related to the improper densification of the composite at a higher TZP loading.

References L. L. Hench, “Bioceramics,”J. Am. Ceram. Soc., 81 [7] 1705-1728 (1998). L. L. Hench and J. Wilson, Science, 226 1111 630-636 (1984). G. De With, H. J. A. Van Dijk, N. Hattu, and K. Prijs, “Preparation, Microstructure and Mechanical Properties of Dense Polycrystalline Hydroxyapatire,”J Mater. Sci., 16 [6] 1592-1598 (1981). K. Ioku, S. Somiya, and M. Yoshimura, “Hydroxyapatite Ceramics with Tetragonal Zirconia Particles Dispersion Prepared by HIP Postsinrering,” J. Ceram. Soc. Jpn. Int. Ed., 99 [3] 196203 (1991).

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1111 2963-2968 (1999).

“Dense Hydroxyapatite-Zirconia Ceramic Composites with High Strength for Biological Application,” Adu. Mater., 13 [3] 214-216 (2001). 8. H.-W. Kim, Y.-H. Koh, B.-H. Yoon, and H.-E. Kim, “Reaction Sintering and Mechanical Properties of Hydroxyapatite-Zirconia Composites with Calcium Fluoride Addition,” J. Am. Ceram. Soc., 85 [6] 1634-1636 (2002).

9. J. Li, L. Hermansson, and R. Soremark, “Synthesis and Sinrering of Hydroxyapatite-Zirconia Composite,” J. Marer. Sci.: M a w . M e d , 4 [ 11 50-54 (1993). 10. R. Ramachandra Rao and T. S. Kannan, “Synthesis and Sintering of Hydroxyapatite-Zirconia Composites,” M a w . Sci. Eng. C, 20 [ 1-21 187-193 (2002). 11. J. Li, H. Liao, and L. Hermansson, “Sintering of Partially Stabilized Zirconia

and Partially Stabilized Zirconia-Hydroxyapatite Composites by Hot Isostatic Pressing and Pressureless Sintering,” Eiomaterialr, 17 1181 17871790 (1996). 12. W . Pyda, A. Slosaryk, Z. Paszkiewicz, A. Rapaa-Kmita, M. Haberko, and A. Pyda, “Effect of Chemical Composition and Morphology of Zirconia

13.

14

15

16. 17.

18.

Particles on Properties of HAP-Zirconia Particulate Composites,” Key Eng Mater., 2 0 6 2 1 3 1567-1570 (2002). E. S. Ahn, N. J. Gleason, and J. Y. Ying, “The Effect of Zirconia Reinforcing Agents o n the Microstructure and Mechanical Properties of Hydroxyaparite-Based Composites,” J. Am. Ceram. Soc., 88 [ 121 3374-3379 (2005). S. H. Kim, H. C. Lee, H. G. Bang, and S. Y. Park, “Effect of MgFz Additive on the Mechanical Properties in HydroxyapatitelZirconia Composites,’’ Mater, Sci. Forum, 510-511 4 7 8 4 8 1 (2006). F. Mezahi, A. Harabi, S. Zouai, S. Achour, and D. B. Assolant, “Effect of Stabilized Zr02, A203and T i 0 2 on Sintering of Hydroxyapatite,” Mater. Sci. Forum. 492-493 241-246 (2005). A. R. Kmita, A. Slosaryk, and Z. Paszkiewicz, “Mechanical Properties of HAp-Zr02 Composites,”J. Eur. Ceram. Soc., 26 1481-1488 (2006). C.-Y. Chiu, H. C. Hsu, and W. H. Tuan, “Effect of Zirconia Addition on the Microstructural Evolution of Porous Hydroxyapatire,” Ceram. Int., 33 715-718 (2007). R. P. Rana, S. K. Pratihar, and S. Bhattachalyya, “Effect of Powder Treatment on the Crystallization Behaviour and Phase Evolution of Alz03-High Z r 0 2 Nanocomposites,” J Mater. Sci., 41 7025-7032 (2006).

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SYNTHESIS AND CHARACTERIZATION OF CHALCOGENIDE NANOCOMPOSITES J. Martin and G. S. Nolas Department of Physics, University of South Florida, Tampa, Florida 33620 ABSTRACT Lead Telluride (PbTe) nanocomposites were prepared by densifying 100 - 150 nm PbTe nanocrystal powders synthesized from a low temperature alkaline aqueous solution reaction in a high yield (> 2 grams per batch). Densification using Spark Plasma Sintering (SPS) successfully integrated nanostructure dispersions within a bulk matrix. We report the synthesis and low temperature transport measurements on novel chalcogenide nanocomposites, including resistivity, Seebeck coefficient, and thermal conductivity. INTRODUCTION Thermoelectric effects couple thermal and electric currents, allowing for the solid-state inter-conversion of heat and electricity. The thermoelectric figure of merit, ZT=POT/K,defines the effectiveness of different thermoelectric materials, where S is the Seebeck coefficient, o is the electrical conductivity, T is the absolute temperature, and K is the total thermal conductivity. Recent progress in a number of higher efficiency thermoelectric materials (room temperature ZT > 1) can be attributed to the nanoscale enhancement of thermoelectric properties. These new materials demonstrate increased Seebeck coefficient and decreased thermal conductivity due to the phenomenological properties of nanometer length scales, introducing both quantum confinement and interfacial scattering effects. One consequence of nanostructure is the increase of interfaces. These interfaces serve to scatter phonons more effectively than electrons and reduce the thermal conductivity. Additionally, the presence of interfacial energy barriers filters the carrier energy traversing the interface, restricting those energies that limit the mean carrier energy.' This increases the Seebeck coefficient, as its value depends on the mean carrier energies relative to those at the Fermi level. Experimental evidence in expitaxial films of n- and p-type lead chalcogenides verifies the enhancement potential of this mechanism.* The primary research into nanostructured enhancement of thermoelectric properties remains limited to thin films, heterostructures, and nanowires. For example, p-type BizTe3/Sb2Te3 10 angstrod50 angstrom supperlattice structures demonstrated a room temperature ZT = 2.4. The thermal conductivity in this system was reduced by a factor of 2 compared to other Bi2Te3 alloys.J Harman and co-workers reported a room temperature ZT = 1.6 in PbTe/PbTeSe quantum-dot superlattices (QDSLS)~that contain PbSe nanodots imbedded in a PbTe matrix. Kong and co-workers reported an enhancement in Si/Ge supper lattice^.^ Recently, Heremans and co-workers reported an increased Seebeck coefficient for PbTe with the inclusion of Pb precipitate nanostructures.6 One method to incorporate nanoscale dimensions into bulk materials is by ball-milling.7,s This procedure rapidly grinds powders to sub-micron dimensions. Ball-milled Si-Ge nanocomposites demonstrated an increased Seebeck coefficient and a reduced thermal conductivity.' Although the electrical conductivity also increased, the overall thermoelectric performance of the material was enhanced. Ball-milled PbTe materials have also demonstrated thermoelectric enhancement.l o However, the syntheses discussed leave suspect a direct correlation between nanostructure and thermoelectric enhancement, due to unaccounted lattice

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strain affects. This stresses the need to directly synthesize nanocrystals and incorporate them into bulk structures. SYNTHESIS AND CHARACTERIZATION Lead telluride nanocrystals were synthesized from the low temperature reaction of a selenium or tellurium alkaline aqueous solution and a lead acetate trihydrate solution.' Several experiments were performed to optimize product yield by varying both the alkaline and precursor lead acetate trihydrate concentrations. To prepare PbTe nanocrystals, the precursor solutions were prepared separately at 90" C by dissolving elemental Te in a 20 M KOH aqueous solution and by dissolving Pb(CH3C00)2-3H20 in distilled water. After 30 minutes, the lead acetate trihydrate solution was dripped into the rapidly stirring deep purple alkaline solution to immediately form PbTe nanocrystals. After 5 minutes, the reaction mixture was removed from the heat source and dilute HN03 was added to flocculate the nanocrystals. The solution was removed from the grayish-black precipitates that were subsequently washed 4 times with the dilute nitric acid solution, removing lead hydroxide impurities. Excess lead acetate trihydrate in the reaction favors the formation of easily removable impurities. The precipitates were then XRD analysis washed 4 times with distilled water and dried overnight in a fume hood. confirmed nanocrystal phase purity. This procedure reproducibly synthesizes 100 - 150 nm spherical PbTe nanocrystals, confirmed by TEM (Figure I), with a high yield of over 2 grams per batch. High yields of nanocrystals are required to synthesize practical nanocomposite materials. PbTe nanocrystals were subjected to Spark Plasma Sintering (SPS) to achieve 95 YO theoretical density. In the SPS procedure, a pulsed DC current conducts through both the graphite die and the sample under high pressure to minimize nanostructure deterioration. Prior to SPS, five batches of PbTe nanopowders were mixed together in a glass vial for each sample. The final densities are 7.67 g/cm3 (sample PbTel) and 7.75 g/cm3 (sample PbTe2). Densifying solely the nanocrystals results in a uniform dispersion of non-conglomerated nanostructure within a bulk matrix.

'

-

-

Figure 1. TEM image of PbTe nanocrystals after precipitation.

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Figure 2. SEM image indicating the nanostructure of PbTel .

SEM (Hitachi S-800) images of a PbTel fracture surface indicate the preservation of nanostructure following the SPS procedure with grain diameters ranging from 100 nm to over one micron (Figure 2). Figure 3 shows the standard X-ray diffraction scans for the two PbTe nanocomposites post SPS. All samples exhibit peaks characteristic of PbTe. The successive spectra are normalized and shifted in intensity for clarity. While the PbTe nanopowder spectra indicates phase purity, the sintered nanocomposites exhibit a PbTe03 impurity phase. Densified nanocomposites of PbTe were cut into 1 x 2 x 5 mm parallelopipeds for transport property measurements. Four-probe resistivity, Seebeck coefficient, and steady-state thermal conductivity were measured on the same sample from 300 K - 12 K. Table 1 summarizes the room temperature data. Both samples are p-type with large room temperature Seebeck coefficients of 330 pV/K (Figure 4). The temperature dependence and the magnitude of the Seebeck share consistency throughout the measured temperature range. Transport measurements also demonstrate a correlation between resistivity and porosity (Figure 5). An increase of only 1.5 % theoretical density between PbTel and PbTe2 decreased the room temperature resistivity by a factor of 2. Since porosity has no affect on the Seebeck coefficient'2, increasingly denser nanocomposites should facilitate optimal thermoelectric performance. The thermal conductivity values also correlate to the difference in sample density. As shown in Table I, PbTel demonstrates a slightly lower thermal conductivity than PbTe2 due to the increased porosity.13 The room temperature ZT for PbTe2 is 0.1 without optimization or doping.

-

-

Table 1. Percent theoretical density, resistivity, Seebeck coefficient, and total thermal conductivitv at 300 K for the PbTel and PbTe2 nanocomDosites. Sample Density (%) S (pV/K) K (W-m-lK-') p (mOhm-cm) PbTe 1 94 328 2.2 24.9 PbTe2 95 324 2.5 12.6

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Figure 3. XRD spectra for the two PbTe nanocomposites post SPS procedure. Arrows indicate PbTe03 impurity.

Figure 4. Temperature dependence of Seebeck coefficient for the two PbTe nanocomposites. Closed circles identify PbTel and open triangles identify PbTe2.

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Figure 5. Temperature dependence of resistivity for the two PbTe nanocomposites. Closed circles identify PbTel and open triangles identify PbTe2. CONCLUSIONS Dense PbTe nanocomposites were prepared by densifying PbTe nanocrystal powders synthesized from high yield (> 2 grams per batch) alkaline aqueous solution reactions. These nanocomposites dimensionally integrate nanostructure within a bulk matrix without conglomeration. Transport measurements indicate a strong correlation between resistivity and porosity, suggesting increasingly denser nanocomposites might hrther optimize the transport properties for thermoelectric properties. To our knowledge, these samples represent the first preparation of dense nanocomposites from solution-phase synthesized nanocrystals, including transport measurements. The authors acknowledge support by GM and DOE under corporate agreement DE-FC26-04NT42278. REFERENCES

B. Y. Moizhes and V. A. Nemchinsky, In Proceeding for the 1 lthInternational Conference on Thermoelectrics, Institute of Electrical and Electronics Engineers, Inc., 1992. Y. I. Ravich, In CRC Handbook of Thermoelectrics, edited by D. M. Rowe, pages 67-73, CRC Press New York, 1995. R. Venkatasubramanian, E. Siivola, T. Colpitts, B. O’Quinn, Nature 413, 597 (2001). 4 T. C. Harman, P. J. Taylor, M. P. Walsh, and B. E. LaForge, Science 297,2229 (2002). 5 T. Kong, S. B. Cronin, M. S. Dresselhaus, Applied Physics Letters 77, 1490 (2000). 6 J.P. Heremans, C.M. Thrush and D.T. Morelli, J. Appl. Phys. 98,063703 (2005). 7 K. Kishimoto and T. Koyanagi, J. Appl. Phys. 92,2544 (2002). 1

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P.C. Znai, W.Y. Zhao, Y. Li, X.F. Tang, Z.J. Zhang and M. Nino, Appl. Phys. Lett. 89, 0521 1 1 (2006). 9 M. S. Dresselhaus, G. Chen, M. Y. Tang, R. G. Yang, H. Lee, D. Z. Wang, Z. F. Ren, J. P. Fleurial and P. Gogna, Proc. Mater. Res. SOC.886 3 (2006). l o J. P. Heremans, C. M. Thrush, and D. T. Morelli, Thermopower Enhancement in PbTe nanostructures, Physical Review B 70, 115334 (2004). 11 W. Zhang, L. Zhang, Y. Cheng, Z. Hui, X. Zhang, Y. Xie, and Y. Qian, Materials Research Bulletin 35,2009 (2000). 12 L. Yang, J. S. Wu and L. T. Zhang, J. Alloys and Compounds 364,83 (2004). l 3 I.Sumirat, Y. Ando, S. Shimamura, J. Porous Mater. 13,439 (2006).

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SELF ASSEMBLED FUNCTIONAL NANOSTRUCTURES AND DEVICES Cengiz S. Ozkan Department of Mechanical Engineering Bourns Hall, A305 University of California, Riverside Riverside, CA 92521 ABSTRACT This paper reports the self assembly of functional nanostructured materials including multi-walled Carbon Nanotube-Quantum Dot (CNT-QD) heterojunctions using the Ethylene Carbodiimide Coupling procedure (EDC). Thiol stabilized ZnS capped CdSe quantum dots containing amine terminal groups (QD-NH2) were conjugated with acid treated Multi-Walled Carbon Nanotubes (MWCNT) ranging from 400 nm to 4pm in length. SEM, TEM, EDS and FTIR were used to characterize the conjugation process. The versatile electrical properties of carbon nanotubes (CNT) make them promising candidates for nano electronic devices, especially tunneling diodes and transistors"*. In previous CNT based nano-electronic devices, control over electrical properties of the devices has been limited. In addition, researchers have relied on overlapping CNTs2 for forming junctions, which introduces local bending that could adversely affect the electrical properties of the tubes3. Alternatively, covalent modifications of carbon nanotubes with metal colloids4 and semiconducting quantum dots ( Q D s ) ~have , ~ been reported for the synthesis of heterojunctions for electronics applications. These approaches were associated with carbon nanotube sidewall functionalization. Here, the synthesis of heterojunctions with controlled conjugation of water stabilized, amine terminating, ZnS coated CdSe QDs (QD-NH2) to acid treated ends of multi walled CNTs (MWCNT) are reported. Scheme 1 illustrates the steps involved in the synthesis of the heterojunctions. Carboxyl groups were introduced to the ends of as-grown MWCNTs (Nanostructured & Amorphous materials, Inc., Los Alamos) by refluxing at 130 "C in concentrated nitric acid for 24 hours. The tips of MWCNTs, which have the highest defect sites, get oxidized first. Oxidized MWCNTs are shorter and have terminating carboxylic groups that impart a hydrophilic nature and facilitate further functionalization. ZnS capped CdSe QDs (Evident Technologies, Inc., New York) were used in the functionalization of the MWCNTs. The use of ZnS coating over the CdSe core improves the quantum yield by passivating7 surface dangling bonds (carrier trap sites) and also, enabling them for use in biosystems'. Scheme 1. Conjugation of ZnS capped CdSe QDs to MWCNTs. Nanocomposites and Nanostructures

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To prepare water stabilized QDs (QD-NH2), ZnS capped CdSe nanocrystals were suspended in chloroform by sonication for 30 minutes. Equal volumes of 1.0 M 2-aminoethane thiol hydrochloride (AET) were added to this solution. This resulted in a two-phase mixture with the aqueous AET forming an immiscible layer above the organic chloroform-QD suspension. The mixture was stirred vigorously on a magnetic plate for 4 hours. When ZnS capped CdSe QDs were reacted with AET, the mercapto group in the thiol group in AET bonded to the Zn atoms and the amine groups rendered the QDs hydrophilic, in addition to facilitating further functionalization possibilities. The aqueous phase containing QD-NH2 was extracted in phosphate buffer saline (PBS, pH=8.5). Successive washing in PBS produced a good suspension of water stabilized QD-NH2 which were used for the synthesis of MWCNT-QD heterostructures via the two-step coupling procedure using 1-ethy1-3-(3-dimethyl-aminopropyl)carbodiimide HC1 (EDC, Pierce Chemicals, Inc.) in the presence of N-hydroxysuccinimide (sulfo-NHS, Pierce Chemicals, Inc.) The EDC reaction was carried out in PBS for 8 hrs at 50°C under continuous mixing. Characterizations of the heterostructures were conducted using scanning electron microscopy (SEM), transmission electron microscopy (TEM), Fourier transform infrared spectroscopy (FTIR) and energy dispersive spectroscopy (EDS). Figure 1(A) is an SEM image of QDs conjugated at the end of a MWCNT. No sidewall functionalizations were observed. MWCNTs produce multiple carboxylic groups upon oxidation at their ends and these results in the conjugation of multiple QDs at the ends. Figure 1. (A) SEM image of water soluble QDs before conjugation. (B) SEM image of QDs conjugated at the ends of a MWCNT (4pm) (C) QDs at the ends of a MWCNT (500nm). Previous research6 indicated sidewall QD conjugations for MWCNTs exceeding 200 nm in length. Our conjugation technique is specific that even for MWCNTs as long as 4 pm (Figure 1 (B)), QDs are observed only at the ends. Figure 1(C) is a SEM image indicating QDs conjugated at the ends of a 500 nm long MWCNT. In all the images, we have observed the absence of side wall functionalization. Further TEM (Figure 2 (A) & (B)) and EDS ( Figure 2 (C)) confirms the presence of QDs at the ends of the MWCNTs. The TEM image in Figure 2(A) shows a MWCNT with QDs at its end. Figure 2 (B) is the image of the same QD cluster at a higher magnification. Figure 2(C) shows the EDS data for the QD cluster in Figure 2 (B) using a 1.50 nm probe size. The spectra shows clear peaks corresponding to the elements of Cd, Se, Zn, and S. Notice that, Cu signal is generated from the TEM grid. The potential argument for the clusters to be Fe nano particles (from the MWCNT growth process) instead of QDs can be excluded after observing absence of the characteristic Fe peaks in EDS data. The MWCNT-QD conjugates were also characterized with FTIR, using an AgCl cell in a Bruker Equinox 55 FTIR spectrometer. Figure 3 is the FTIR spectra of oxidized MWCNTs (blue curve) and the MWCNT-QD conjugates (red curve). With plain oxidized MWCNTs, absorption peaks are observed at 1644 cm-1, 1704 cm-1 and 3403 cm-1 (peaks designated A, B, and C), which are characteristic of carboxylic and phenolic groups on acid

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treated MWCNTs. For the MWCNT-QD conjugates, new absorption peaks appear at 1653 cm-1, 2977 cm-1 and 3314 cm-1 (D, E, and F), which indicate the C=O stretch in amides, C-H stretch and N-H stretch in amides. C-H and N-H peaks are bigger than the amide C=O peak due to the presence of free QDs in the sample. Nevertheless, a slight blue shift of carboxylic C=O stretch to amide C=O stretch and the appearance of C-H and NH peaks indicate the formation of covalent MWCNT-QD conjugations, via the amide bond formation.

In conclusion, a method to synthesize heterojunctions of individual MWCNTs with QDs, preferentially at the MWCNT ends has been reported. This was confirmed by SEM, TEM and EDS analysis. Due to the mild and Figure 2. (A) SEM image of water soluble QDs well-controlled oxidation of the MWCNTs, before conjugation. (B) SEM image of QDs conjugation occurs only at the ends for conjugated at the ends of a MWCNT (4pm) nanotubes ranging from 400 nm to 4 pm in (C) QDs at the ends of a MWCNT (500nm). length. The controlled conjugation process preserves the electronic properties of the MWCNTs and enables the nanoassembly of heterojunctions. They can be used as building blocks for various nanoscale electronic or optoelectronic devices and three dimensional hierarchical assemblies of multilayered systems.

Figure 3. FTIR spectra of Oxidized MWCNTS (blue) and MWCNT-QD conjugates (red). Absorption Peaks are observed at 1644 cm-', 1704 cm-' and 3403 cm-I in the FTIR spectra for oxidized tubes. New peaks develop at 1653 cm-I, 2977 cm-l and 3314 cm-l in the FTIR spectra of MWCNT-QD conjugates, indicating formations of MWCNT-QD conjugates via amide bond formation.

NanocomDosites and Nanostructures

REFERENCES I (a) Tans, S. J.; Verschueren, A. R. M.; Dekker, C. Nature 1998, 393, 49. (b) Li, J.; Papadopoulos, C.; Xu, J. Nature 1999, 402, 253. (c) Yao, Z.; Postma, H. W. C.; Balents, L.; Dekker, C. Nature 1999, 402(6759), 273. (d) Ahlskog, M.;Tarkiainen, R.; Roschier, L.; Hakonen, P. Appl. Phys. Lett. 2000, 77, 4037. (e) Zhou, C. W.; Kong, J.; Yenilmez, E.; Dai, H. J. Science 2000, 290, 1552. ( f ) Ahlskog, M.; Hakonen, p.; Paalanen, M.; Roschier, L.; Tarkiainen, R. J. Low Temp. Phys. 2001, 124, 335. (g) Rosenblatt, S.; Yaish, Y.; Park, J.;

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Gore, J.; Sazonova, V.; McEuen, P. L. Nan0 Lett. 2002,2,869. 2 Fuhrer, M. S.; Nygard, J.; Shih, L.; Forero, M.; Yoon, Y. G.; Mazzoni, M. S. C.; Choi, H. J.; Ihm, J.; Louie, S. G.; Zettl, A.; McEuen, P. L. Science 2000,288,494. 3Dai, H. J. Surface Sci. 2002, 500,218. 4(a) Liu, J.; Rinzler, A. G.; Dai, H.; Hafber, J. H.; Bradley, R. K.; Boul, P. J.; Lu, A.; Iverson, T.; Shelimov, K.; Huffman, C. B.; Rodriguez-Macias, F.; Shona, Y.-S.; Lee, T. R.; Colbart, D. T.; Smalley, R. E., Science 1998, 280, 1253. (b) Azamian, B. R.; Coleman, K.S.; Davis, J. J.; Hanson, N.; Green, M.L.H. Chem. Commun. 2000.4,366. 'Banerjee, S.; Wong S. S. Nan0 Lett. 2002,2, 195. 6Haremza, J. M.; Hahn, M. A.; Krauss, T. D. Nan0 Lett. 2002,2, 1253 7Hines, M. A.; Gnyotsionnest, P. J. Phys. Chem. 1996, 100,468. k h a n , W. C. W.; Nie S. M. Science 1998,281,2016.

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CARBON NANOTUBE (CNT) AND CARBON FIBER REINFORCED HIGH TOUGHNESS REACTION BONDED COMPOSITES P. G. Karandikar, G. Evans, and M. K. Aghajanian M Cubed Technologies, Inc. 1 Tralee Industrial Park Newark, DE 19711

ABSTRACT Reaction bonded Sic and B4C offer low density (light weight), high hardness, high stiffness, high thermal conductivity, low CTE, excellent ballistic resistance, complex shape producibility and high volume producibility - properties needed in many armor, thermal management, semi conductor capital equipment, and aerospace mirrors and structures markets. However, similar to other ceramics, their toughness and strength are lower compared to incumbent metallic materials. Here we demonstrate the use of carbon fiber and carbon nanotube (CNT) reinforcements for increasing toughness of reaction bonded ceramics. The reaction bonding process depends on the reaction between carbon and molten silicon to achieve infiltration of particulate or fibrous preforms (e.g. Sic, B&, etc.). Thus, if unprotected, carbon fibers and CNTs will convert to Sic in this process. In this work, innovative processing was conducted to successfully incorporate both carbon fibers and carbon nanotubes in reaction bonded materials. Cf/SiC composites were obtained with quasi-isotropic low CTE (< 1 ppm/K) and high fracture toughness (6-10 MPa m112). Fracture toughness of reaction bonded Sic was increased from 4 to 7 MPa m112(a 73% increase) using CNTs. INTRODUCTION Reaction bonded Sic and B4C materials’-’ offer high specific stiffness (stiffnesddensity - E/p) and thermal stability (thermal conductivity / coefficient of thermal expansion - WCTE), and are used for applications such as lithography equipment, optical structures, mirrors, thermal management and aerospace components. Properties of several candidate materials that were historically considered for these applications are compared with those of M Cubed materials in Table 1. In some aerospace applications however, high toughness and strength are also required. Beryllium (Be) offers the best combination of specific stiffness, high strength and high toughness (8-10 MPa m ‘I2), and is therefore extensively used in aerospace structural components. However, Be has many limitations such as high cost, low thermal stability and above all, health hazards associated with Be dust. Reaction bonded Sic and B4C can compete with Be in terms of specific stiffness and have better thermal stability than Be. However, they suffer from the limitation of low (4MPa m ‘ I 2 ) toughness.

’

THE REACTION BONDING PROCESS Reaction bonded Sic was first developed in the 1940’~’-~. Other terms for the process include ‘reaction sintered’ and ‘self bonded”. Figure 1 shows a schematic of this process. In this process, the preform containing the reinforcement and a carbon precursor or binder is “carbonized” in an inert atmosphere above 600°C to convert the precursor to carbon. Next, the preform is placed in contact with Si metal or alloys of Si in an inert or vacuum atmosphere and

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heated to above the melting point of the alloy. Due to the spontaneous wetting and reaction between carbon and molten Si, the preform is infiltrated completely. The carbon in the preform reacts with the Si forming Sic and in the process bonds the reinforcement together. Some residual Si remains. Table 1.

Summary of properties of selected materials for aerospace applications.

p = Density, a = CTE (-50 to 100°C), k = Thermal Conductivity, E = stiffness, Krc = Fracture Toughness

Figure 1. A schematic of the reaction bonding process

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A major advantage of this process is that the volume of the reaction-formed Sic is 2.3 times larger than the volume of the reacted carbon. Thus, by infiltrating Si into preforms that contain high carbon contents, ceramic bodies rich in Sic can be produced. In addition, like water, silicon expands on freezing. As a result, unlike other materials made by casting-type processes, RB composites do not show any shrinkage porosity. The reaction bonding process has several advantages relative to traditional ceramic processes (e.g., hot pressing, sintering). First and foremost, volume change during processing is very low (generally well less than l%), which provides very good dimensional tolerance control and eliminates the need for final machining. In addition, the process requires relatively low process temperatures and no applied pressure, which reduces capital, tooling and operating costs. Moreover, fine high surface area powders capable of being densified are not required, which reduces raw material cost. In the earlier work5 on reaction bonded B4C it was shown that an unwanted reaction between the Si and B4C phases would often occur during the infiltration process, which resulted in cracking of the parts. To prevent excessive reaction from occurring, coarse particles (e.g., 300 microns) were used. This limited the surface area for reaction, but also limited the strength of the resultant material. M Cubed has fbrther optimized the reaction bonding process6-", to produce relatively fine grained Sic and B4C ceramics with favorable mechanical and ballistic properties as well as good machinability. In the specific case of reaction bonded B4C, the unwanted reaction between Si and B& was suppressed by alloying the Si infiltrant with boron (B). The binder formulation was optimized to yield preforms with high strength to allow intricate green machining using CNC machines. Preform bonding technology was developed to allow fabrication of complex shapes, ribbed structures, box structures, cooling channels etc. The shrinkage from the preform stage to final infiltrated stage was reduced to less than 0.5%, allowing net-shape component fabrication with minimal finish machining (that too only for high precision components). This process technology allows fabrication of intricate large structures (500 lb, 1.5m x 0.75 m x 0.3 m), as well as mass production of smaller components (10s of thousand per month). In spite of these refinements, one limitation of these materials remains, namely their lower fracture toughness than Be. As a result, their applications have been limited to stiffness controlled applications and they are not viable for applications needing high strength and toughness. Therefore, the present work pursued two approaches for enhancing the fracture toughness of reaction bonded materials: ( 1) Carbon fiber reinforcement, (2) carbon nanotube (CNT) reinforcement. CARBON FIBER REINFORCED RB SIC

Fiber reinforcement (e.g. Sic, A1203,and carbon) has been shown to increase toughness of ceramics'*. Out of all the fibers, carbon fibers are the most cost competitive due to their wider-scale usage in polymeric composites (e.g. graphite-epoxy). Other fibers (e. g. Sic, A1203) have proven to be cost prohibitive and their use is limited to very high temperature applications (>lOOO"C) where carbon fibers cannot be used due to their oxidation problem. Continuous carbon fiber based composites are typically anisotropic and laminate designs are used for making most of the components. The laminates are typically quasi-isotropic (i.e. the properties are uniform in the plane of the laminate). As described earlier, the reaction bonding process depends on the good wetting and chemical reaction between carbon and molten silicon. This wetting and reaction lead to spontaneous infiltration of preforms containing carbon by molten silicon. For processing carbon fiber reinforced Sic composites by reaction bonding, the key challenge is prevention of attack on

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the carbon fibers by molten silicon and their conversion to Sic. In the p a ~ t ' ~this - * problem ~ was overcome using carbon, Sic or BN coatings deposited on the fibers in a separate process step by techniques such as chemical vapor deposition (CVD) or chemical vapor infiltration (CVI). Such processes are time consuming and expensive. M Cubed has d e ~ e l o p e d in-situ '~ formed coatings that protect carbon fibers from attack by molten silicon. Figure 2a shows a micrograph of a woven Cf/SiC composite made by reaction bonding and using the in-situ formed protective coating. Energy dispersive analysis of X-rays (EDAX) was conducted on the fibers in this microstructure. Figure 2b shows an EDAX pattern from one such fiber. In this pattern, predominantly carbon is seen with very minor amount of Si. Thus, the fiber was not converted to Si. Typically, the EDAX pattern of Sic predominantly shows Si because in the presence of silicon, carbon X-rays get absorbed and do not reach the detector. Thus, the fact that the EDAX pattern from the fiber shows predominantly carbon is very significant in proving that the fiber was not converted to Sic. A dense carbon coating formed on the fiber surface during the initial processing steps is postulated to form a dense, protective Sic coating on the fiber surface during silicon infiltration, which prevents further attack by molten silicon and complete conversion to silicon carbideI4'l6.

Figure 2. (a) Microstructure of CdSiC composites (b) EDAX analysis from fibers shows predominantly carbon indicating that the fibers are protected from reacting with molten silicon.

To-date several different Cf/SiC composites have been made with different fibers, fiber architectures and in-situ formed coatings. The focus of the present development effort was obtaining composites with quasi isotropic low CTE and high toughness for some specific aerospace applications. Typically 2 x 2 x 0.125 and 6 x 6 x 0.125 inch panels were fabricated. The Cf/SiC composite plates were subjected to extensive microstructural, physical, thermal and mechanical characterization. The densities of the test samples were determined using the Archimedes principle per ASTM C373. For microstructural observations, specimens were sectioned and polished and observed by optical and scanning electron microscopes (SEM). Specimen elastic moduli were measured using strain gages in a tensile test. The flexural strengths were determined using a four-point bend testing apparatus per ASTM C 1161. Fracture toughness measurements were made on selected specimens by the Chevron notch method

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(ASTM C1421). Strength and toughness measurements were made on at least 5 specimens of each type. CTE was measured by laser interferometry. Conventional dilatometry techniques are not sufficiently accurate for measuring CTEs of low expansion materials. Thermal conductivities were measured by the axial rod method. The measured properties of these composites are summarized in Table 2. All three composites on which toughness measurements were made show toughness in excess of 6 MPa m”. The data also shows that very low CTE (< 1 ppm/K) Cf/SiC composites were successfully produced. It also shows that the CTE of the composite can be controlled by changing the fiber type (pitch based, PAN based, high modulus, high strength, extent of graphitization etc.) and architecture (unidirectional, cross ply, quasi-isotropic, woven etc.). The ultimate bend strength (UBS) and the modulus (E) of each composite are controlled by the fiber architecture, interface type (fiber pull out) and the extent of fiber-silicon reaction.

Table 2.

Properties of C$SiC composites made by reaction bonding for low CTE applications

IUni-directional

E

2.22

312

9.7

--

-1.02

y: 122 --

*The interface designations refer to various proprietary in-situ coatings. Several components have been fabricated from these composites. Figure 3 shows a photo of a lightweight mirror made out of Cf/SiC. Work continues to further enhance the toughness and strength of these composites.

Figure 3. Photos of the back and the front of a lightweight Cf/SiC mirror.

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ADVANTAGES OF CNT REINFORCEMENT AND PROCESSING CHALLENGES

Carbon nanotubes (CNTs) were first observed by Iijima in 1991. Since then, a significant effort has been spent in trying to make and characterize n a n ~ t u b e s ' ~The - ~ ~properties . of two typical carbon fibers are compared with the properties of carbon nanotubes in Table 3. Clearly, the properties of CNTs are significantly better than those of carbon fibers. Therefore, CNTs offer a very high reinforcing potential.

Table 3.

Comparison of properties of carbon nanotubes with the properties of carbon fibers.

I

Property

Diameter (micrometer) Density (g/cc) Elastic Modulus (GPa) Ultimate Tensile Strength (GPa) Thermal Conductivity (w/mK) CTE (ppm/K) Electrical Resistivity (micro-ohm-m)

Carbon Fibers Carbon Nanotubes T300 P120 1) (pircn oased) 7 10 0.05 1.76 2.17 -2.0 23 1 827 1000-1400 3.75 2.41 7-10 8 640 >2000 -0.6 -1.45 - 1 (isotropic) 18 2.2
CNTs are made by a variety of techniques such as chemical vapor deposition (CVD), laser evaporation, carbon-arc etc. CNTs are available in either single walled or multi-walled form. Figure 4 shows SEM of CVD-grown multi-walled CNTs. It can be seen from Figure 4 that the CVD-grown multi-walled CNTs (the most cost-effective form) are inter-twined like spaghetti. Thus, effective means are required to de-agglomerate the CNTs. Processing of CNT composites presents many challenges: (1) Difficulty in uniformly dispersing , (2) Difficulty in aligning in certain directions, (3) Achieving a good CNT-matrix bond, (4) In the case of ceramic matrices - preventing reaction between the CNTs and matrix due to high processing temperature/pressure, and ( 5 ) In the case of reaction bonding, an additional challenge is the attack and conversion to S i c by molten silicon.

Figure 4. SEM of CVD grown multi walled carbon nanotubes (MWNTs).

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FABRICATION OF RB CNT/SiC As described earlier, M Cubed has developed a technique to protect carbon fibers from attack by molten silicon during reaction b ~ n d i n g ’ ~This . technology was extended to process CNT/SiC composites by reaction bonding22. In addition, process modifications were made to overcome some of the processing difficulties listed in the previous section (e.g. dispersing CNTs). This technology was applied to make several 2 x 2 x 0.25 inch sample plates for characterization. The CNT/SiC composite plates were subjected to extensive microstructural, physical, thermal and mechanical characterization. The densities of the test samples were determined using the Archimedes principle per ASTM C373. For microstructural observations, specimens were sectioned and polished and observed by optical and scanning electron microscopes (SEM). Specimen elastic moduli were measured by an ultrasonic measurement technique. The time of flight of a shear wave and a longitudinal wave through the sample were measured to determine the Poisson’s ratio of the material, as well as the elastic modulus. The flexural strengths were determined using a four-point bend testing apparatus per ASTM C 1161. Fracture toughness measurements were made on selected specimens by the Chevron notch method (ASTM C1421). Strength and toughness measurements were made on at least 5 samples of each type. CTE was measured between 0 and 100°C using dilatometry. Thermal conductivity (K) was calculated from measured thermal diffusivity (D) and specific heat (Cp) data (K = Cp . p . D; where p is the density of the material). Thermal diffbsivity was measured using a laser flash method (ASTM E1461-92). Specific heats were also measured by the laser flash technique. Electrical resistivity measurements were carried out using the four point probe. Figure 5 compares the microstructure of a CNT/SiC composite with the microstructure of a standard reaction bonded Sic. Identical starting Sic powders were used for making both materials. The Sic content in the CNT/SiC is a little bit higher due to higher amount of reaction formed Sic. The property data for the two materials are compared in Table 4.The table shows that the fracture toughness of the standard reaction bonded Sic was increased by 73% using CNT reinforcement, while maintaining or improving all other properties. Also, a 78% reduction was obtained in electrical resistivity by using CNTs. This is another indication that CNTs still exist in the final composite and are not converted to Sic. As listed in Table 3, CNTs have a very low electrical resistivity. On the other hand, S i c has high electrical resistivity. Figures 6 and 7 show high resolution SEM micrographs of the fracture surface of two CNT/SiC test specimens. Here individual CNTs can be seen pulled out of Si/SiC. These observations unequivocally demonstrated that the CNTs were successfully protected from complete reaction with molten silicon. Similar to the case of carbon fibers, it is postulated that CNTs are protected by an in-situ formed Sic coating. Based on these microstructural observations, the higher toughness in CNT/SiC composite compared to the baseline Sic can be attributed to one or more of the following mechanisms: crack deflection by CNTs, crack bridging by CNTs, and CNT pullout. A CNT/SiC demonstration component was successfblly fabricated. A photo of this component is shown in Figure 8. The high fracture toughness obtained in the CNT/SiC composite while maintaining most of the other properties of the baseline RE3 Sic, will allow this material to start competing with Be for many aerospace applications. Work is underway to further enhance the strength and fracture toughness of this material.

Nanocomposites and Nanostructures

Figure 5. Comparison of microstructures of the RB Sic and RB CNT/SiC. Table 4. Comparison of properties of RB CNT/SiC and RB Sic.

I

Property

Stiffness (GPa) Bend Strength (MPa) Toughness (MPa m Electrical Resistivity (Ohm-cm) CTE 25-1OOC (ppm/K) Thermal Conductivitv (w/mK)

384

359 288 4 1.097 2.9 188

374 285 6.9 0.245 2.72 185

4 (1) 73 (78) (6) (1.5)

Progress in Nanotechnology: Processing

Figure 6. SEM of the fracture surface of RE3 CNT/SiC example sample 1. The higher magnification picture clearly shows individual CNTs pulled out.

Figure 7. SEM of the fracture surface of RE3 CNT/SiC example sample 2. The higher magnification picture clearly shows individual CNTs pulled out.

CONCLUSIONS Reaction bonded composites offer very high mechanical and thermal stability. However, like other ceramics, their toughness is lower than metallic materials such as A1 and Be. Carbon fibers and CNTs can enhance toughness of reaction bonded ceramics. However, both these reinforcements can be attacked by molten silicon during the reaction bonding process. In this work, CdSiC composites were successfilly fabricated by reaction bonding by using innovative in-situ formed protective coatings. C$SiC composites with toughness of 6 to 9 MPa m”* and quasi-isotropic low CTE (<1 ppm/K) were produced. Similarly, CNTs were successfilly dispersed and protected during RB SIC composite fabrication. Fracture toughness of RE3 S i c was increased by 73% using CNT reinforcement.

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Figure 8. A photo of a CNT/SiC demonstration component.

ACKNOWLEDGEMENTS This work was supported by the following US DoD SBIR contracts: F29601-03-C-0012, FA9453-04-M-0255, and HQ0006-06-C-7443. REFERENCES 1 K.M. Taylor, “Cold Molded Dense Silicon Carbide Articles and Methods of Making the Same,” U.S. Pat. No. 3 205 043, Sept. 7, 1965. ’P.P. Popper, “Production of Dense Bodies of Silicon Carbide,” U.S. Pat. No. 3 275 722, (1966). 3C.W. Forrest, “Manufacture of Dense Bodies of Silicon Carbide,” U.S. Patent No. 3 495 939, (1970). 4R. Morrell, Handbook of Properties of Technical and Engineering Ceramics, HMSO Publications, London, U.K., 1985. %.M. Taylor and R.J. Palicka, “Dense Carbide Composite for Armor and Abrasives,” U.S. Pat. No. 3 765 300, (1973). 6 M. Aghajanian, B. Morgan, J. Singh, J. Mears and B. Wolffe, “A new family of reaction bonded ceramics for armor applications,” in Ceramic Armor Materials by Design, Ceramic Transactions, Vol. 134., J. W. McCauley et a1 editors, (2002) 527-540. 7 P. Karandikar, M. Aghajanian, and B. Morgan, “Complex, net-shape ceramic composite components for structural, lithography, mirror and armor applications,” Ceramic Engineering and Science Proceedings, Vol. 24 [4], (2003) 561-566. 8 M. Waggoner, B. Rossing, M. Richmond, M. Aghajanian, and A. McCormick, “Silicon carbide composites and methods for making same,” US Patent No. 6,503,572 (2003) ’P. G. Karandikar and M. K. Aghajanian, “Ultrasonic NDE of reaction bonded ceramics,” Ceramic Transactions 175, Advances in CMC-XI, (2006) 49-6 1. 10 M. Waggoner, B. Rossing, M. Richmond, M. Aghajanian, and A. McCormick, “Silicon carbide composite bodies and methods for making same,” US Patent No. 6,919,127 (2005).

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”P. G. Karandikar, M.K. Aghajanian, D. Agrawal, and J. Chang, “Microwave Assisted (MASS) processing of metal-ceramic and reaction bonded composites”, Ceramic Engineering and Science Proceedings Vol. 27 [2] (2007) 435-446. 12 P. G. Karandikar and Tsu-Wei Chou, “Chapter 10: Structural Properties of Ceramic Matrix Composites,” in Handbook of Continuous Fiber-Reinforced Ceramic Matrix Composites, R. Lehman et al. (Editors), American Ceramic Society (1995) 355-429. 13 M. P. Borom, W. B. Hillig, R. N. Singh, W. A. Morrison, and L. V. Interrante, “Fiber containing composites,” US Patent 5,O 15,540, May 1991. I4P. G. Karandikar, J.R. Singh, and C. A. Anderson, “Low-expansion metal ceramic composite bodies and methods for making same,” US Patent 7,169,465 January 2007. 15 G. Dietrich, T. Haug, A. Kienzle, C. Schwartz, H. Stover, K. Weisskopf, and R. Gadow, “Fiber reinforced composite ceramics and method of producing the same,” US Patent 6,26 1,981, July 2001. 16G.Dietrich, G. Gross, T. Haug, and K. Rebstock, “Brake unit consisting of a brake disk and a brake lining,” US Patent 6,079,525, June 2000. I7P. G. Collins, and P. Avouris, “Nanotubes for electronics,” Scientific American, 283 (6) 2000, pg. 62-69. 18 M. Treacy, T. W. Ebbesen, and T. M. Gibson, “Exceptionally high Young’s modulus observed for individual carbon nanotubes,” Nature 381 (1996) pg. 680-687. 19 E. W. Wong, P. E. Sheehan, C. M. Lieber, “Nanobeam mechanics: elasticity, strength and toughness of nanorods and nanotubes,” Science 277 (1997) pg. 1971-1975. 20R. S. Ruoff and D.C. Lorents, “Mechanical and thermal properties of carbon nanotubes,” Carbon 33 (7) 1995 pg. 925-930. 21M.Dresselhaus, G. Dresselhaus, P. Elkund, and R. Saito, “Carbon Nanotubes,” Physics World January 1998 (http://physicsweb.org/article/world/ 11/ 1/9/ 1) 22E.T.Thostenson, P. G. Karandikar, and Tsu-Wei Chou, “Fabrication and characterization of reaction bonded silicon carbide/carbon nanotube composites,” J. Physics D: Applied Physics, Vol. 38, Issue 21, (2005) 3962-3965.

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SYNTHESIS, CHARACTERIZATION AND MEASUREMENTS OF ELECTRICAL PROPERTIES OF ALUMINA-TITANIA NANOCOMPOSITES Vikas Somani and Samar J. Kalita Department of Mechanical, Materials and Aerospace Engineering University of Central Florida Orlando, FL 328 16-2450

ABSTRACT In this study, we have synthesized and characterized nanocrystalline alumina-titania (Al2O3Ti02) composites via a simple sol-gel process, using aluminum propoxide and titanium propoxide as precursor chemicals. Propanol and 2-methoxy ethanol were used as solvent and stabilizer, respectively. The as-formed gel was heat treated at 400"C, to obtain amorphous powder. The amorphous powder was subsequently calcined at 700°C and 900°C. Phase evolution, phase composition, crystal structure and crystallite-size of the synthesized powder were determined using X-ray diffraction (XRD) technique. Crystallite-size was further confirmed by high-resolution transmission electron microscopy (HR-TEM). HR-TEM results of powder calcined at 700°C showed agglomerates of powder particles, with particle-size in 15 - 20 nm range. The synthesized powder was uniaxially pressed using a steel mold and then sintered at elevated temperature (10001500°C) for densification study and electrical property measurements. XRD technique was used to study phase composition of the sintered pellets. Dielectric constant and dissipation factor of the sintered pellets were measured. The effects of sintering time, temperature and phases present on the electrical properties of the sintered pellets, were studied. INTRODUCTION A1203-Ti02 systems are promising candidates for use in electronic devices as dielectrics. Si02, the most widely used dielectric till date, is reaching its physical limits due to miniaturization. It has been shown that Si02 films with thickness less than 7 nm lead to significant increase in leakage current. This drives the need to find new dielectrics. A number of reports in the last few years are based on potential dielectric materials like Hf02, A1203, ZrO2, Ta205, TiO2, BaTi03 and SiN.' In particular, Hf02 and A1203 with dielectric constants (K) of 25 and 9, respectively, are believed to be materials of choice in near-future applications. Recently, there has also been an upsurge in research of compound dielectric materials, for example ZrSi04 and HfSi04.2

-

-

Alumina (A1203)has excellent insulating properties (K 9), a high band gap 9.9 eV, and is resistant to most chemicals. While, titania (Ti02) possesses a much higher K 80-120 depending on its crystal ~tructure.~ The disadvantage with A1203is its relatively low dielectric constant, which

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389

limits its utility and life-span to only a few years. Ti02, on the other hand, is unstable on silicon resulting in high leakage current. Aluminum titanate (A12Ti05,AT), a compound of A1203and Ti02 is known to have excellent thermal shock resistance, low thermal expansion coefficient (0.2 X and high melting point.4' AT is usually synthesized by solid state reaction between to 1.0 X equimolar mixture of A1203 and Ti02 powders, at temperatures higher than 1400°C.63 The endothermic reaction can be written as: A1203+ Ti02

-+

A12Ti05

AT dissociates to corundum and rutile in the temperature range of 75O"C-128O0C, and to overcome this demerit, stabilizers like Fe2O3, MgO or Si02 are added. It has been established that nanocrystalline powder of various ceramic materials show enhanced sintering behavior, better mechanical and electrical properties.' Sol-gel processing technique has been successfully exploited to synthesize nanopowder and thin films. However, only a few studies have been reported on sol-gel synthesis of aluminum titanate, as well as aluminatitania nanocomposites. We hypothesized that nanocomposite systems of A1203 and Ti02 would help us unite the merits of both A1203 and Ti02 and possibly develop new material systems for future dielectrics. With this hypothesis, we synthesized of A1203-Ti02 nanocomposites by sol-gel technique; and systematically studied their sintering and densification behavior along with their dielectric properties. EXPERIMENTAL PROCEDURE Powder Synthesis A1203-Ti02 nanocomposite powder was synthesized using aluminum isopropoxide (reagent grade, Fisher Scientific, USA) and titanium (IV) isopropoxide (purity 98+%, Fisher Scientific, USA) as precursor chemicals and propanol (1 N, Alfa Aesar, Ward Hill, MA) as the solvent. Hydrochloric acid (HC1) was used to control the pH of the solution. In order to prepare the sol 20 ml of propanol was heated in a beaker to 50°C. Subsequently, stoichiometric amount of aluminum isopropoxide was added into hot propanol and refluxed at 120"C, for 2 h. The refluxed solution was allowed to cool in air. In another beaker, stoichiometric amount of titanium (IV) isopropoxide was dissolved in 5 ml propanol. The as-formed aluminum isopropoxide solution was then added dropwise into the solution of titanium (IV) isopropoxide in propanol using a burette. The resultant solution was continuously stirred, during the entire process. Hydrochloric acid was added during this process to maintain the pH of the solution between 3.5 and 7.0. As-produced A1-Ti sols at different pH were refluxed at 120"C, for 2 h. The resulting solutions were filtered and then dried in a muffle furnace to obtain amorphous powder. The as-formed amorphous powders were calcined at 700°C for 2 h. Based on XRD phase analysis of these powders, new sol was prepared with pH 5.5 only. Amorphous powder obtained from this sol was calcined separately in batches at different temperatures, viz., 400"C, 700°C and 900°C in a muffle furnace under atmospheric conditions, for crystallization. A flowchart of the powder synthesis process is shown in figure 1.

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Powder Characterization

1

Aluminum Isopropoxide + X-ray diffraction (XRD) propanol technique was used to study the phase evolution and phase transformation during synthesis as a function of solution pH and Reflux 12OoC,2 h calcination temperature. During + propanol initial phase of this work, I agglomerated powder obtained from solutions having different pH I I (3.5, 5.5 and 7.0), calcined at Add Ti isopropoxide solution 700°C for 2 h, were ground into a dropwise into Aluminum solution fine powder, by mortar and pestle and analyzed for phases. The results showed evolution of crystalline phases in powder synthesized using solution with pH of 5.5. Hence, we focused on studying the characteristics of powder synthesized, using solution with pH of 5.5 only. Powder Heat the sol at 4OO0C, 2 h agglomerates obtained after calcination at 400"C, 700°C and 900°C for 2 h, in separate batches, Calcination at different were ground into a fine powder temperature using mortar and pestle, and analyzed for phase composition. A Figure 1. Flowchart showing steps involved in the Rigaku diffractometer (Model DMAX-B, Rigaku Co., Tokyo, synthesis of alumina-titania composite nano-powder. Japan) equipped with Ni filtered CuKa radiation (h =O. 1542 nm) at 40 kV and 40 mA settings was employed for the purpose. Based on XRD analysis, powder calcined at 900°C for 2 h was examined for morphology, grain size and lattice fringes using a HR-TEM (Model Tecnai - Philips F30, FEI Co., Hillsboro, OR).

1

1

-

~

Sintering and Densification Study As-synthesized alumina-titania composite nanopowder calcined at 700°C was compacted in a steel mold, using a uniaxial single action manual hydraulic press (Carvar Press Inc, USA). These green structures were sintered at elevated temperatures (1000, 1100, 1200, 1300, 1400 and 15OO0C), to investigate their sintering behavior, densification characteristics and phase transformation as a result of sintering. Some of the sintered structures were also characterized for electrical properties. During the preparation of the green specimens, a dry P.T.F.E film (made with Dupont Krytox) was

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sprayed on the inside surface of the steel mold and the punch, to lubricate and reduce friction. Uniaxially compacted green pellets with average dimensions of 12.5 mm in diameter and 3.0 mm in thickness, were sintered in a high temperature programmable muffle furnace in ambient atmosphere, and at different temperatures for 2 h. A heat treatment cycle was developed to allow better densification and to avoid cracking. These pellets were initially heated to 400°C at a heating rate of 8 "C/min and homogenized at this temperature, for 30 min. In the second step, temperature was increased to the desired final level; at a slow heating rate of 4"C/min. Pellets were subsequently cooled to room temperature, at 10"C/min. Three specimens were heat-treated at each sintering temperature and the values here reported, are the average of the three samples. Geometric bulk density (p,) of each pellet was evaluated from the measurements of the mass of specimen and its volume (determined by dimensional measurements) using equation 1. Geometric bulk density (p,)

= Mass

/ Volume

(1) Linear and volumetric shrinkages associated with sintering at different temperatures were also calculated from geometric measurements of green and sintered specimens. Phase Analysis of the Sintered Specimens

XRD technique was used to study the effect of sintering temperature on phase purity and phase transformation of the alumina-titania nano-composites. Ceramic structures sintered at different temperatures were ground into a fine powder, using mortar and pestle. The fine powder thus obtained, was used for XRD analysis in a Rigaku diffractometer, as stated under Powder Characterization. Electrical Characterization Alumina-titania nano-composite structures sintered at IOOO"C, 1200°C and 1400°C for 2 h, were used to measure electrical properties. Silver paste (electrode) was applied on both sides of these sintered pellets and the resulting structures were dried at 400°C for 30 min, in a muffle furnace, in air. The dielectric constant and dissipation factor of these sintered pellets were measured at 1 MHz, using a precision LCR meter from Agilent technology (Model 8284A) and 16451 B dielectric test fixtures. The effects of phase composition, density and porosity of the sintered structures on the electrical properties were investigated. RESULTS AND DISCUSSIONS Figure 1 is a flowchart, showing the steps in the synthesis of alumina-titania composite nanopowder. During the synthesis process, it was interesting to observe the changes of color in the amorphous and the calcined powders. Amorphous powder obtained from sols, maintained at all three pH viz., 3.5, 5.5 and 7 , were initially light brown in color. After calcination, the color of the powder from the sol of pH 5.5 changed into white, while no color change was observed in case of powders from the other two sols (PH 3.5 and 7.0).

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22

27

32

2-T heta (Degrees)

37

Figure 2. XRD patterns of powder synthesized from solution maintained at different pH. Calcination was done at 700°C for 2 h. The phases present were 8-A1203and anatase Ti02 with an unknown peak, which are marked as 1,2 and U, respectively. The powder synthesized from solution maintained at three different pH values, viz., 3.5, 5.5 and 7.0 were analyzed, using XRD (figure 2). No distinct peaks were observed in the XRD pattern of powder synthesized from sol of pH 3.5. This XRD pattern confirms that the synthesized powder was amorphous in nature. XRD pattern of powder synthesized from sol of pH 7.0 shows a few low intensity broad peaks, confirming that it is not an optimum pH to obtain crystalline powder. On the other hand, XRD pattern of powder synthesized from sol of pH 5.5, showed clear broad peaks of 8A1203 and anatase Ti02. From this XRD pattern, it is clear that the powder is crystalline and was of a very fine size. It is well known that the reaction rate of molecular precursors used in the sol-gel process, is highly dependent on the processing variables like pH, concentration, solvent and temperature. In this process, we fixed the solvent and concentration but altered the pH and temperature, to find the optimum synthesis condition. The formation of a-A1203 from y- A1203is known to occur, by a series of polymorphic transformation on heating.’

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Here, it can be observed that @-A1203is stable at relatively low temperatures. This is due to the presence of Ti02 and structural changes, forced as an effect of the nanosize of the particles. XRD results showed that variations in pH do play a role in the synthesis. These XRD patterns also revealed that sol having a pH of 5.5 was good for the synthesis of alumina-titania composite powder. Therefore, in the next stage of our research, we studied the phase evolution as a function of calcination temperature, in powders synthesized using sol with pH of 5.5. The powders were calcined at 400"C, 700°C and 900"C, for 2 h. XRD pattern of the powder calcined at 900°C is shown in Figure 3. Figure 4a shows the TEM micrograph of the powder calcined at 700"C, for 2 h. Although, the calcined powder was ground manually using a mortar and pestle they are in the form of agglomerates. These agglomerates were found to be 0.3 to 0.6 pm in size. The agglomerates in turn, are formed of very fine particles, of about 15 - 20 nm in size. A HR-TEM micrograph of the particle is shown in figure 4b. The lattice fringes and ordered arrangement of atoms is apparent which, confirms the crystalline nature of the particles. Uniaxially compacted alumina-titania nanocomposite pellets were sintered at 1OOO"C, 12OO0C, 1300"C, 1400°C and 1500"C, for 2 h. Figure 5a is the plot of variation in sintered density, with increase in the sintering temperature. Volume shrinkage and linear shrinkage were calculated from the difference in the geometric measurements between green and sintered pellets and the calculated results were plotted as a function of sintering temperature, as shown in figures 5b and 5c, respectively. It is clear from these plots that geometric sintered density increased sharply from 1.39 g/cm3 at 1000°C to 2.15 g/cm3 at 1200°C. The corresponding volume shrinkage in this range almost doubled from 24% to 46%. The average sintered densities of the pellets, sintered at 1200°C

t

23

1

I

I

24 25 26

I

I

I

I

I

I

I

27 28 29 30 31 32 33 2-T-

(eesee)

I

34

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I

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Figure 3. XRD pattern of powder calcined at 900°C.

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Figure 4. (a) TEM micrograph of A1203-Ti02 composite nanopowder agglomerate calcined at 700°C for 2 h. (b) HR-TEM micrograph of the nanopowder showing the lattice fringes.

Figure 5a. Plot showing variation in sintered density of A1203- Ti02 composite pellets sintered at lOOO"C, 1200"C, 1300"C, 1400°C and 1500°C in furnace atmosphere.

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Figure 5b. Plot showing variation volume shrinkage of A1203- Ti02 composite pellets sintered at various temperatures in furnace atmosphere.

Figure 5c. Plot showing variation in linear shrinkage of A1203-Ti02 composite pellets sintered at various temperatures in furnace atmosphere.

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and 1300°C were found to be almost constant; though, the volume shrinkage increased from 45% to 56%. Thereafter, the sintered density increased to 2.6 g/cm3 in pellets sintered at 1400°C and the corresponding volume shrinkage at this temperature was found to be 62%. The highest sintered density of 2.99 g/cm3 is achieved at 1500"C, corresponding to volume shrinkage of almost 65%. The highest sintered density achieved is 80.7% of theoretical density of A12Ti05. As expected, an increase in sintering temperature led to an increase in sintered density. The linear shrinkage of the pellets followed a trend almost similar to the volume shrinkage. The linear shrinkage increased from 11% to 26% as temperature increased from 1000°C to 1400°C. It is to be noted that the molar volume of a-Al203, rutile-Ti02 and AlzTiOs are 25.56, 28.82 and 49.20 cm3/mol, respectively, which result in a net volume expansion, on transformation to A12Ti05. It is clear from the XRD patterns that the phase transformation to Al2TiOs phase occurs in the range of 1000°C to 1400°C. The linear shrinkage in this temperature range shows this effect, as it decreases from 25% at 1200°C to 22% at 1300"C, and then again increases to 32% at 1500°C.

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Figure 6a. XRD patterns of the nanocomposite sintered at lOOO"C, 1100°C and 1200°C for 2 h. A and T represent a-A1203and rutile-Ti02, respectively.

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The pellets used for sintering and densification study were ground into a fine powder, to study the effects of sintering temperature on phase transformation using XRD technique. XRD analysis was also performed to determine the crystallite size. Figure 6a shows the XRD patterns of samples sintered at lOOO"C, 1100°C and 1200°C for 2 h. The peaks in the XRD patterns are distinct and sharp indicating complete crystallization and all the phases have been identified. The phases present in the sintered structures were a-Al2O3 and rutile-TiO2. XRD plots of powder heat treated at 1300"C, 1400°C and 1500°C for 2 h, is shown in Figure 6b. Aluminum titanate phase begins to appear at 1300°C and as temperature increase its percentage increases. Initially at 1300"C, percentage of aluminum titanate phase was 15%, which sharply increased to 65% at 1400°C and at 15OO"C, pure aluminum titanate phase was obtained. The molar volume of a-&03, rutile-Ti02 and A12Ti05 are 25.56, 28.82 and 49.20 cm3/mol resulting in volume expansion on transformation to A12Ti05 at elevated temperature^.^ This volume expansion leads to development of compressive stresses and associated strain, resulting in peak shift to higher angles, as observed in figure 6b between 1300°C and 1400°C.

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2-Theta (Degrees) Figure 6b. XRD patterns of the A1203-Ti02 nanocomposite sintered at 1300"C, 1400°C and 1500°C for 2 h. 1 , 2 and 3 represent a- &03, rutile-Ti02 and A12Ti05 phases, respectively.

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The crystallite size of the powder calcined at various temperatures for 2 h, was calculated using the Scherrer's formula (2) Bcvs = (k h) / (L cos6) where h is the wavelength of the X-ray used (1.5406 A), 6 is the Bragg angle, k is a constant and L is the full width at half maximum. The crystallite size at 700°C of anatase-Ti02 calculated using (101) peak as the reference, was 18 nm and of 6-A1203calculated using (102) peak as reference was 26 nm. With the increase in sintering temperature, phase transformation occurred in both a-A1203 and rutile-Ti02 resulting in particle size of 24 nm at 1200"C, for both phases. Further increase in the sintering temperature resulted in grain coarsening of the two phases and evolution of aluminum titanate phase. In the specimens sintered at 1300"C, the grain size of a-A1203,rutile-Ti02 and AT (calculated using (110) peak as reference) phases were found to be 48 nm, 45 nm and 54 nm, respectively. The results clearly indicate that nano-features were retained in the composite form even after pressure-less sintering for 2 h, at 1300°C.

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Figure 7. Dielectric constant of sintered pellets as a function of sintering temperature. Dielectric constants of the sintered pellet structures have peak value at 1200°C (Figure 7) with the highest value of 10.32, in pellets sintered for 2 h. Dielectric constants of bulk materials depend on phase composition, defects, density and grain size." The phases present at both 1000°C and 1200°C are crystalline a-Al2O3 and rutile-Ti02 but the pellets have a marked difference of density, with the density at 1200°C higher. Even though the density increases at 1400"C, the dielectric constant shows a sharp decrease. This is clearly because of the increased volume fraction of the A12Ti05 phase, It must be pointed out that pellets sintered for 2 h have a composite composition of a-Al203, rutile-Ti02 and A12Ti05. This seems to be a result of higher density and the affect of volume expansion, due to the phase transformation taking place. The dissipation factors of the pellets sintered at 1200°C and 1400"C, were found to be 0.52 and 0.02, respectively. The formation of A12TiOS phase has a favorable effect on the dissipation factor and it decreases sharply. Thus, though the formation of A12Ti05 phase decreases the dielectric constant, it improves the

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dissipation factor. Thus, a proper phase composition can be identified to maintain a balance between the two, according to the requirements. CONCLUSIONS A1203-Ti02 nanocomposite powders were successfully synthesized via a simple sol-gel process. The effect of pH on the synthesis route has been investigated. Best results were obtained for powders synthesized from sol maintained at a pH of 5.5. A12Ti05 phase began to appear at 1300°C and its phase percentage increased with the increasing sintering temperature, resultin in pure A12Ti05 at 1500°C. The densification study showed the highest density of 2.99 g/cm in pellets sintered at 1500”C, for 2 h. Sintered density increased with the increase in sintering temperature from 1000°C to 1500°C. Results also showed that the presence of A12Ti05 decreases the dielectric constant but significantly improved the dissipation factor. Highest dielectric constant of 10.32 was achieved in pellets sintered at 1200”C, for 2 h. Crystallite-size of A1203, Ti02 and A12Ti05 in structures sintered at 1300°C was calculated to be 48 nm, 45 nm and 54 nm, respectively, using the Schemer’s formula.

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REFERENCES ‘P. Singer,“High K dielectrics: No easy solution”, Semiconductor International, 38 (2003)

R. W. Murto, M. I. Gardner, G. A. Brown, P. M. Zeitzoff, H. R. Huff, “Challenges in gate stack engineering”, Solid state Technology, 43-48 (2003)

2

3

M. Wallace, “Challenges for the characterization and integration of high-k dielectrics”, App. Surface Sc., 231-232,543-55 l(2004)

4

L. Stanch, J. R. Groza, L. Stoica , C. Plapcianu, “Influence of powder precursors on reaction sintering of A12Ti05”, Scripta Materialia, 50, 1259-1 262 (2004)

5R. G. Duan, G. D. Zhan, J. D. Kuntz ,B. H. Kear, A. K. Mukherjee, “Processing and microstructure of high-pressure consolidated ceramic nanocomposites”, Scripta Materialia, 51,1135-1 139 (2004) 6

D. M. Ibrahim, A. A. Mostafa, T. Khalil, “Preparation of tialite (aluminium titanate) via the urea formaldehyde polymeric route”, Ceramics International, 25,697-704 ( 1999)

7T. S. Liu, D.S. Perara, “Long-term thermal stability and mechanical properties of Aluminium I Mat. Sc., 33,995-1001 (1998) titanate at 1000-1200T”, . 8

C. Suryanarayana and C. C. Koch, Hyperfine Interactions, 130,5 (2000)

9J. G. Li and X. Sun, “Synthesis and sintering behavior of a nanocrystalline a-alumina powder”, Acta muter, 48, 3 103-3112 (2000) ‘OC. Hube, C. Elissalde, V. Hornebec, S. Mornet, M. Treguer-Delapierre ,F. Weill, M. Maglione, “Nano-ferroelectric based core-shell particles: towards tuning of dielectric properties,” Ceramics International, 30, 1241-1245 (2004)

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A New Ternary Nanolaminate Carbide: Ti3SnC2 Sylvain Dubois,' Thierry Cabioc'h, Patrick Chartier, Vitronique Gauthier, and Michel Jaouen Laboratoire de Metallurgie Physique, UMR 6630 du CNRS, Universitk de Poitiers, 86962 Futuroscope Chasseneuil Cedex, France

We report on the synthesis of Ti,+,SnC, (MAX phases) by hot isostatic pressing starting with titanium, tin, and carbon powders. In addition to the already known TizSnC compound (211 MAX phase), a new 312 MAX phase, Ti3SnC2, is formed. Its lattice parameters, deduced from Rietveld analysis of X-ray diffraction patterns, are found to be a=0.31366 nm and c = 1.8650 nm.

T

HE ternary M,, IAX, carbides and compounds, referred to as 211, 312, and 413 MAX for n = 1, 2, and 3, respectively, have attracted great attention recently.6 Here, M denotes an early transition metal, A is an A-group element (from IIIA to VIA) and X is either C or N. In these compounds, near-close-packed transition metal carbide and/or nitride layers are interleaved with layers of pure A-group element; MAX phases are thus considered as nanolaminated-layered materials. In the M2AX or 211's phases, every third layer is an A layer,' in the M3AX2 or 312's, every fourth layer: and in the M4AX3or 413's, every fifth layer.7s8The metallic nature of the bonding and the nano-layered nature of the structure give rise to a unique combination of metallic and ceramic properties. For example, these carbides and nitrides exhibit high strength and stiffness at high temperatures, good resistance to thermal shock and oxidation, high-thermal and electrical conductivities as well as high damage tolerance; they are also readily machinable.' About 50 such phases are known, most of them belonging to the 21 1 family (the so-called H-r Hagg phases). Indeed, there are only three 312 (Ti3SiC2,' Ti3GeC; and Ti3AIC2'@) and three 413 (Ti4A1N3,' Ti4SiC3," and Ta A1C3'23'3)phases known to date. If chemical vapor deposition' 1(CVD) and, more recently, physical vapor deposition' (PVD and mainly magnetron sputtering) methods are used to synthesize thin films of MAX phases, the hot isostatic pressing (HIP) route is mainly used because probably it is the most efficient for bulk materials processing. In this paper, results obtained by revisiting the HIP synthesis of Ti2SnC, already performed in 1963 by Jeitschko et d.,' are discussed. This work was performed using powder mixtures of titanium (150-250 pm, 99.5%), tin (2-20 pm, 99O/0), and carbon (10-40 pm, 99.8%) to yield a final composition with the appropriate stoichiometry 50:25:25 (Ti:Sn:C at.%). The powders were thoroughly ball mixed in a turbula for 1 h in order to ensure good mixing and to break up agglomerates. A steel container was then filled with the powder reactant mixture. The extremity of the container was welded under high vacuum using electron beam. Then, the green bodies were placed in the hot isostatic press and subjected to the following temperature and pressure cycle: (1) the samples were heated at 5"C/min to 850°C under an atmospheric pressure of

(2) the samples were soaked at 850°C for 1 h while the HIP was pressurized with Ar to 120 MPa, (3) once pressurized, the heating was resumed, with a heating rate of 25"C/min, up to 1315"C, (4) samples were held at 1315°C and 120 MPa for 4 h before cooling to room temperature. After HIPing, the sample was machined to remove the encapsulating steel container, sliced using a diamond wheel and roughly polished using S i c paper for microscopic observations. The latter were performed using a scanning electron microscope (SEM, JEOL 5600 LV) equipped with an energy dispersive X-ray spectrometer (EDXS, Oxford Isis 300) for chemical analysis. X-ray diffraction (XRD) experiments were carried out on a Siemens D-5000 diffractometer using the CuKu radiation (Bragg-Brentano geometry) for phase identification. A typical SEM micrograph of the HIPed sample is displayed in Fig. 1. It shows that the material is not fully dense because of the presence of numerous pits and holes. Furthermore, a marked contrast can be observed between the core of the sample and its external part, which looks darker. This external ring extends between 200 pm and about 500 pm depending on the observed area. At higher magnification (Fig. 2), one observes black precipitates and elongated plates (light gray) embedded in a more uniform matrix (dark gray). From EDXS analyses, black precipitates are assigned to T i c while light gray elongated plates are enriched in Fe (mainly Fe3Sn2 from EDXS analyses). The composition of the dark gray areas varies from the core of the sample to its periphery: the ratio Ti/Sn is close to 2 in the first region and close to 3 in the second one. The related XRD pattern is shown in Fig. 3. The most striking point is to observe well-defined diffraction peaks at 28 = 9.6" and 19.12' (diamonds in Fig. 3) that look like replicates of the (0002) and (0004) reflections of the TizSnC phase. From EDXS analyses, one can assume that there exists an unknown 312 MAX phase contributing to the measured XRD pattern. To assess this assumption, XRD data were refined using the MAUD softwarei7assuming the presence of a

G Hug--contnhuting editor

Manuscript No. 22780. Received February 7, 2007; approved March 18, 2007. 'Author to whom correspondence should he addressed. e-mail: sylvain.dubois@ univ-poitiers.fr

Nanocomposites and Nanostructures

Fig. 1. Back-scattered scanning electron microscope micrograph of the cross section of the cylindrical HIPed sample. The dashed line is a guide for the eye (see text).

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Fig. 2. Back-scattered scanning electron microscope micrograph of the core of the sample.

mixture of Ti2SnC, Tic, and of a hypothetic Ti3SnC2 MAX phase (space groupe P63/rnrnc). Despite the true conditions not being fulfilled to perform a Rietveld analysis (the sample is not milled to obtain a powder because we intend to perform transmission electron microscopy (TEM) observations), the calculated pattern is in rather good agreement with the experiment (Fig. 3). For the probed volume, the analysis gives a composition of about 92% of Ti2SnC, 5% of Ti3SnC2, and 3% of TIC. The lattice parameters of the 21 1 phase are a = 0.31665(1) nm and c = 1.37123(6) nm, values slightly different from those reported by Jeitschko and Nowotny3 (a = 0.3186 nm, c = 1.363 (a = 0.316265(9$ nm) and more recently by Vincent et nm, c = 1.36789(5) nm) and Barsoum and colleaguesIb3’ (a=0.31635 nm and c = 1.3675 nm). Regarding the new 312 MAX phase, Ti3SnC2, its lattice parameters are a = 0.3136(1) nm and c = 1.866(1) nm. It should be pointed out here that impurities have been excluded from the Rietveld fit, the missing diffraction peaks can be attributed to secondary phases (mainly Fe3Sn2 and Fe3Sn). To have a better determination of the 312 MAX phase lattice parameters, an XRD experiment was performed illuminating the external part of the sample. Indeed, SEM observations and EDXS analyses have shown that such a region is nearly free of the 21 1 phase. The related XRD pattern as well as its Rietveld fit are plotted in Fig. 4. As the illuminated sample surface is not perfectly flat, the background is noisier than in Fig. 3, but the overall agreement between the experimental and calculated patterns is quite good. Here too, impurities that are at a quite low level are excluded from the fit. If one excludes these minority phases, the probed volume was found to be composed of Ti3SnCz (78%) and TIC (22%) only. The 312 lattice parameters are found to be a = 0.31366(2) nm and c = 1.8650(2) nm, values that are, within the error bars, in agreement with those quoted above. In the Wycoff notations, there exists four inequivalent sites in the hexagonal P63/rnrnc space

Fig. 4. Observed/calculated/difference X-ray diffraction Rietveld plots for Ti3SnC2,TIC (A)and impurities (Fe3Sn2, Fe2Sn) (*).

group for 312 MAX phases: MI in 2a, MIIin 4f, A in 2b, and X in 4f. For the atoms in 4f positions, the ideal values of the two internal dimensionless parameters z/c are 1jS = 0.125 and 1/16 = 0.0625 for MI1and X, respectively. From our analysis, Ti3SnC2 departs from the ideal stacking because we obtain zjc = 0.1204(3) for TiIl and zjc = 0.072(1) for C atoms, a behavior already observed in MAX phases.20’21 Finally, the unexpected presence of Fe in our sample can be discussed. Its presence likely originates from the steel container used during the HIPing process. Nevertheless, it is surprising that this quite large amount of Fe (10 at.% from EDXS) does not forbid the MAX phases growth,22 a point that deserves further study. In conclusion, we discovered a new member of the 3 12 MAX phase family: Ti3SnC2. Its latticc parameters have been determined from Rietveld refinement of the XRD data. Further analyses using TEM techniques are currently in progress along with HIP experiments to try to obtain Ti3SnC2 single-phase material.

References ‘W. Jeitschko. H. Nowotny, and F. Benesovsky. “Kohlen-Stoffhaltige Ternare Verbindungen H-Phase,” Monatsh. Chem., 94. 672 (1963). ’W. Jeitschko, H. Nowotny, and F. Benesovsky, “Die H-Phasen Ti2TIC, Ti2PbC. NbJnC, Nb2SnC und TazGaC.” Monarsh. Chem., 95,431 (1964). ’W. Jeitschko and H. Nowotny, “Die Kristallstructur von Ti3SiC2-Ein Neuer Komplxcarbid-Typ.” Monarsh. Chem.. 98. 329 (1967). 4H. Wolfsgruber, H. Nowotny. and F. Benesovsky, “Die Kristallstuktur Von Ti3GeC.” Monarsh. Chem., 98. 2403 (1967). 5V. H. Nowotny, “Strukturchemie Einiger Verbindungen der Ubergangsmetalle Mit Den Elementen C, Si, Ge, Sn,” Prog. Solid State Chem., 5, 27-70 (1971). ‘M.W. Barsoum and T. El-Raghy, “Synthes~sand Characterization of a Remarkable Ceramic: Ti3SiC2,” J . Amer. Ceram. Sue., 79 [q 1953-6 (1996). ’M. W. Barsoum. L. Farber, I. Levin, A. Procopio, T. El-Rdghy, and A. Berner, “HRTEM of Ti4AIN3;or Ti,A12N2 Revisited,” . I Am. . Ceram. Soc., 82 [9] 2545-7 (1999). ‘C. J. Rawn, M. W. Barsouni, T. El-Rdghy, A. Procopio, C. M. Hoffmann, and Layered M,+]AX, Nitride,” Muter. C. Hubbard. “Structure of Ti4AIN,-,-A Res. Bulletin, 35. 1785 (2000). 9M. W. Barsoum, “The M N + I A X NPhases: A New Class of Solids ThermoI 4 1201 dynamically Stable Nanolaminates.” Prog. Solid Stare Chem., 28 [ (2000). ‘OM. A. Pietzka and J. C. Schuster. “Summary of Constitutional Data on the Aluminum-Carbon-Titanium System.” J . Phase Equilibria. 15 [4] 392400 ( 1994). “J.-P. Palmquist, S. Li, P. 0. Persson, J. Emmerlich, 0. Wilhelmsson, H. Hogberg, M. 1. Katsnelson, B. Johansson, R. Ahuja, 0. Eriksson. L. Hultman, and U. Jansson, “M,, + ,AX. Phases in the Ti-SiK System Studied by Thin-Film Synthesis and ab initio Calculations,” Phvs. Rev. B, 70, 165401 (2004). ”I.-P. Palmquist. T. El-Raghy, J. Howing, 0. Wilhemsson, and M. Sundberg, “Crystal Structure and TEM Characterization of Ta4AIC,” 30rh Inrernarional ConJcrence on Advanced Ceramics and Composites, Cocoa Beach, January 2006. Abstract #ICACC-SI-184, 2006. ”B. Manoun, S. K . Saxena, T. El-Raghy, and M. W. Barsoum. “High-pressure X-Ray Diffraction Study of Ta4AIC3,” Appl. Phys. Lett.. 88, 201902 (2006). I4J. J. Nickl, K. K. Schweitzer, and P. Luxenberg. “Gasphasenabscheidung Im Systeme Ti-C-Si,” J. Less Common Metals, 26, 283 (1972).

A.

Fig. 3. Observed/calculated/difference X-ray diffraction Rietveld plots for TizSnC (no symbol), Ti3SnC2 (+), T i c (A),and impurities (Fe3Sn2, FezSn) (*). 402

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”J.-P. Palmquist, U. Jansson, T. Seppanen, P. 0.A. Persson, J. Birch, L. Hultman, and P. Isberg. “Magnetron Sputtered Epitaxial Single-Phase Ti3SiC2 Thin Films,” Appl. Phys. Letf.,81 [5] 835-7 (2002). 16M. W. Barsoum, G. Yaroschuk, and S. Tyagi, “Fabrication and Chardcterization of MISnC (M =Ti, Zr, Hf and Nh),” Scripta Materialia, 37 [lo] 1583-91 (1997). ”L. Lutterotti, S. Matthies, and H.-R. Wenk, “MAUD: A Friendly Java Program for Material Analysis Using Diffraction.” IUCr: Newslett. CPD. 21 [I41 (1999). “H. Vincent, C. Vincent, B. F. Mentzen, S. Pastor, and J. Bouix, “Chemical Interaction Between Carbon and Titanium Dissolved in Liquid Tin,” Marer. S d . Eng.. A256, 83 (1998).

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I9T. El Raghy, S. Chakrabrty, and M. W. Barsoum. “Synthesis and Characterization of Hf2PbC ZrzPbC and M2SnC(M =Ti. Hf. Nb or Zr),” J. Eur. Ceram. Sac., 20, 2619-25 (2000). ’‘(3. Hug, M. Jaouen, and M. W. Barsoum, “X-Ray Absorption Spectroscopy, EELS, and Full-Potential Augmented Plane Wave Study of the Electronic Structure of Ti2AIC, Ti2AIN, Nb2AIC, and (Tio sNbo.s)2AIC,”Phys. Rev. B, 71.024105 (2005). ”G. Hug, “Electronic Structures of and Composition Gaps Among the Ternary Carbides Ti2MC,” Ph~vs.Rev. B, 74, 184113 (2006). 22N.Tzenov, M. W. Barsoum. and T. El-Raghy, “Influence of Small Amounts of Fe and V on the Synthesis and Stability of TilSiC2.” J . Eur. Ceram. Soc., 20, 8 0 1 4 (2000). 0

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Fabrication of a Nano-Si3N4/Nano-C Composite by High-Energy Ball Milling and Spark Plasma Sintering Xin Xu,+.$Toshiyuki Nishimura, Naoto Hirosaki,* Rong-Jun Xie,* and Hidehiko Tanaka* Nano Ceramics Center, National Institute for Materials Science, Tsukuba, Ibaraki 305-0044, Japan

A nano-Si3N4/5 wt% nano-C composite was successfully fabricated for the first time via high-energy ball milling, followed by spark plasma sintering. The milling promoted the amorphization of the starting powders; most of the carbon particles were transformed into nano-size and embedded in the amorphous phase. This, combined with a low sintering temperature and a rapid densification rate, prevented the reactions between carbon and the other starting powders, leading to a uniform nanolnanocomposite microstructure. Nano-sized carbon grains with an average diameter of about 10 nm were homogeneously dispersed in nano-sized (about 70 nm) Si3N4grain boundaries. The hardness of the obtained nano-ceramics is comparable with that of conventional Si3N4, whereas Young’s modulus is significantly decreased.

I. Introduction

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ILICONNITRIDE-based ceramics have been widely studied as candidates for technical and engineering applications due to their excellent mechanical roperties and good corrosion and thermal shock resistance.13‘They are suitable for a number of industrial applications, such as under high pressure, high temperature, and corrosive environments. T o meet the requirements for different severe environments, a second phase is sometimes added. For example, high refractory second phases (Sic, TIC, etc.) were added to improve high-temperature mechanical properties.’ An electrically conducting second phase (TIN) was added to im rove ele~tricity.~ h-BN was added to improve machinablit y.? Si3N4 ceramics have found many tribological applications, which sometimes need to be under non-lubricated sliding contact. However, the sliding contact of Si3N4-Si3N4 self-mated tribopairs under dry conditions often has a high friction coefficient and a high wear rate.6 Many researchers have reported self-lubricating Si3N4 com osites including carbon, boron nitride, and m01ybdenurn.~-? Among these materials, carbon possesses one of the highest solid lubricating functions. Hyuga et a/.” reported that the friction coefficient of the Si3N4/carbon fiber composite was around 30% of that of a monolithic Si3N4 ceramic. However, the fracture strength decreased significantly from 950 to 550 MPa even with only 1 wt% carbon fiber addition,” because large carbon fibers acted as fractural origins. Thus, it is very important to decrease the grain size of added carbon. On the other hand, nano-composite materials containing two distinct phases of material in which one or both phases are B. Derhy--contrihuting editor

Manuscript No. 22235. Received September I I , 2006; approved January 12, 2007. This work was financially supported by the Ministry of Education, Culture, Sports. Science, and Technology of Japan under Scientific Research No. 17360345(Section (B)) and from Chinese Academy of Sciences under the program for Recruiting Outstanding Overseas Chinese (Hundred Talents Program). ‘Member, American Ceramic Society. ‘Author to whom correspondence should be addressed. e-mail xuxin@!ustc,edu.cn ‘Department of Materials Science & Engineering, University of Science and Technology of China, Hefei, Anhui 230026, China

Nanocomposites and Nanostructures

present in nanometer dimensions have gained considerable interest recently. This interest arises because the properties of these materials often differ from composites in which the components have a large dimension. As Niihara” proposed the concept of “nano-composite ceramics,” many researchers have shown that nano-composites could exhibit improved mechanical properties. Recently, soft and weak materials have also been used as dispersoids. It was reported that the mechanical properties of Si3N4 ceramics were improved by nano-sized soft and weak ceramic dispersoids such as h-BN. For example, the 5 vol% BN-Si3N4 nano-composites showed improved strength despite a decrease in Young’s modulus due to the soft h-BN a d d i t i ~ nWang . ~ et a/.13found that the highly shear-deformable graphite in the A1203 composites helps redistribute the stress field under indentation, conferring the composites with contactdamage resistance. Thus, the incorporation of nano-carbon particles is also expected to increase some mechanical properties, such as fracture strength, contact-damage resistance, etc. The purpose of this study is to find a new process to disperse nano-sized carbon particles homogeneously in a nano-sized Si3N4 matrix. As fine carbon powder will easily react with Si3N4, SO2, and oxide additives at a high temperature, the addition of carbon to a Si3N4 powder was mostly a method to incorporate S i c grains into Si3N4 It is impossible to prepare Si3N4/Cnano-composite through conventional methods. To prevent reactions between carbon and the other starting powders, the ceramics should be sintered at a lower temperature within a very short time. We have reported the preparation of nano-sized Si3N4ceramics through high-energy ball milling, followed by spark plasma sintering. The ceramics could be fully densified at a lower temperature within 1-2 min.16 In this paper, a Si3N4/C nano/nano-composite will be prepared for the first time through high-energy ball milling, followed by spark plasma sintering. We will discuss in detail the synthesis process and microstructural features of the composite. 11. Experimental Procedure

(1) Materials Starting powders were 90.73 wt% submicrometer P-Si3N4powder (NP500grade, Denki Kagaku Kogyo Co., Tokyo, Japan) with an average particle size of 0.5 pm, 7.85 wt% Y2O3 (99.9% pure, Shin-etsu Chemical Co., Tokyo, Japan), and 1.42 wt% A1203 (99.9% pure, Sumitomo Chemical Co., Tokyo, Japan). The amount of carbon (Activated graphite powder, Wako Pure Chemical Industries, Osaka, Japan) with an average particle size of 5 pm was 5 wt% based on the weight of the other starting powders. The starting powders were mixed in ethanol using silicon nitride balls for 4 h. After drying, the as-received powder mixture was then highenergy ball milled using silicon nitride balls 5 mm in diameter and silicon nitride pots 350 mL in volume. The ball-to-powder weight ratio was 20: I, the milling speed was 475 rpm, and the milling time was 6 h. Powder mixtures were compacted in a carbon die (15 mm in inner diameter and 30 mm in outer diameter) and sintered using SPS (Sumitomo Coal Mining Co. Ltd., Tokyo, Japan) under a compressive stress of 30 MPa. The temperature was measured using an optical pyr405

ometer through a 5.5-mm depth hole on the outer surface of the graphite die. Heating was at 300"C/min and cooling at 600"C/ min. The linear shrinkage of the specimen was obtained by directly measuring the movement of the crosshead. All processing steps were conducted in a N2 atmosphere to avoid oxidation.

(2) Characterization X-ray diffractometry (XRD, Philips PW 1700, Eindhoven, the Netherlands) was used to identify phases of the starting and milled powders and the sintered samples using CUKEradiation at a scanning rate of 2"C/min. The observation of the particle morphology of the powders was performed on a field emission high-resolution transmission electron microscopy (HRTEM; JEM-3000F, JEOL, Tokyo, Japan). Thermogravimetric analyses (TG) of as-received and milled powders were carried out in flowing air from room temperature (RT) to 600°C at a heating rate of 2"C/min on a Rigaku TG 8120 thermal analyzer (Tokyo, Japan). The bulk densities of sintered samples were measured based on the Archimedes principle. The fracture surfaces of sintered samples were observed by scanning electron microscopy (SEM, Hitachi S-5100, Tokyo, Japan). The TEM specimen was prepared from the sintered sample using standard procedures for ceramics. The specimen was then observed on a field emission HRTEM (JEM-3000F, JEOL). The grain size was estimated from the SEM and TEM images. Raman scattering of sintered samples was measured at RT by a laser Raman spectrophotometer (JASCO NR-1800, Tokyo, Japan). The hardness test was conducted with a Vicker hardness tester at a load of 10 kg. Young's modulus of the specimen was determined using ultrasonic equipment (EXPLORER-9000, Matec Instruments, Hookinton, MA) with a fundamental frequency of 15 MHz. 111.

Results and Discussion

Figure 1 shows the typical XRD patterns of as-received and milled powders. Carbon peaks cannot be detected in both powders, because the starting carbon powder is mostly amorphous. The sharp diffraction peaks suggest that the starting powders contain coarse polycrystalline grains of Si3N4,Y203,and A1203. The high background and the weakened peaks after milling indicate that the starting powders have mostly been transformed into an amorphous phase. Residual P-Si3N4 peaks are greatly broadened, suggesting significant refinement in the particle size of crystalline P-Si3N4. To reveal more information, the internal structure of milled powders has been investigated by TEM technique. Figure 2 shows the TEM image and the corresponding select area diffraction (SAD) pattern for milled powders. The existence of a 2000

t

Fig. 2. Transmissionelectron microscopy image and the corresponding select area diffraction pattern for milled powders.

diffused halo ring and several sharp spots in the SAD pattern indicate that the milled powders were composed of an amorphous structure with many nano-crystallines embedded within. The TEM image indicates many particles with a diameter of lO(t200 nm. It could be concluded that the particles are mainly composed of an amorphous phase with many nano-Si3N4particles embedded in them, as shown e lse ~he re .' ~~' ' Typical thermal analysis curves for as-received and milled powders heated in flowing air are shown in Fig. 3. A large weight loss ( > 4 wt%) can be found for the as-received powders, which could be attributed to the oxidation of carbon, whereas for milled powders, nearly no weight loss could be detected, indicating that most of the carbon particles have been embedded into the amorphous phase during the milling process. The amorphous phase separates free carbon from oxygen gas in air, and prevents their reaction. The shrinkage curves of as-received and milled powders are plotted in Fig. 4. After being sintered at 1600°C for 5 min, the as-received powders reach only 85% pth, whereas the milled powders could be completely densified with a very fast densification rate. Milled powders contain a lar e amount of liquid phase and much finer P-Si3N4 particles.'' This promotes particle rearrangement at a lower temperature with a higher shrinkage rate, achieving full densification in a short time. The as-received powders could only be fully densified at a higher temperature of 1780°C with the same holding time of 5 min. The XRD pattern of milled powders after being sintered at 1600°C for 5 min is shown in Fig. 5(a). Comparing this pattern

As-received

Milled

As-received

I

10

20

30

40

20

50

60

70

Fig. 1. X-ray diffractometry patterns of as-received and milled powders.

406

200

I

400 Temperature ("C)

I

600

Fig. 3. Thermal analysis curves of as-received and milled powders in flowing air.

Progress in Nanotechnology: Processing

-

s -ge

1500

1.0

Milled

h c .m

v

a,

g

c

As-received

Si,N, / C composite

D hand

G hand

I

- 0.8 7a,3 a,

> .c

-ma,

900

- 0.6

600

0

200

Time (s)

oc

400

40t

Shrinkage curves of milled and as-received powders.

with that shown in Fig. I, we can say that the densification step leads to a rapid crystallization of the amorphous phase. The only crystalline phase present in the composite is P-Si3N4;SIC or other reaction phases could not be detected. For the ceramics from as-received powders sintered at 1780°C for 5 rnin (Figs. 5(b) and (c)), the phase composition is not uniform throughout the specimen. S i c was formed near the surface of the specimen, while only P-Si3N4 could be detected in the center region. Figure 6 shows the Raman spectra of monoclinic silicon nitride nano-sized ceramics16 and the composite we obtained in this research. The peaks at 1584 and 1350 cm-l for composite are the well-known graphite G band and disorder-induced D band," confirming the existence of free carbon in the form of a partial amorphous graphite. The sintered body from milled powders is crushed for TG analysis using the same method as that for starting powders; a weight loss of about 2.8% was detected, which could be attributed to the oxidation of free carbon. Because the ceramics cannot be completely crushed and perhaps some carbon grains are dispersed within Si3N4grains, the weight loss is lower than the added carbon content. Figure 7 shows the sintered microstructures from as-received powders (1780°C, 5 min) and milled powders (16OO0C, 5 rnin). The as-received powders lead to a nano-SiC/micro-Si3N4composite. Matrix Si3N4 grains are homogeneous and most of them are nearly the same as the starting Si3N4powder in size, indicating that SPS sintering is suitable for suppressing grain growth. New-forming BC grains with an average grain diameter of 90 nm agglomerate and disperse within the Si3N4 grain boundaries. A nano/nano-composite microstructure has been obtained from milled powders. Based on the microstructure of nano-sized Si3N4ceramics prepared using the same method," it could be concluded that the small grains are carbon, which were

0 500

loo0

1500

2000

Wavenumher (cm-') Fig. 6. Raman spectra of monoclinic silicon nitride ceramics and the composite we obtained.

homogeneously dispersed in nano-sized silicon nitride grain boundaries. The homogeneous dispersion of carbon in the matrix is important for retaining the mechanical properties, because the aggregration of soft carbon particles will lead to a decrease

N Si3N4

N

20

c

c

1200

Fig.4.

80

30

40

20

50

Fig. 5. X-ray diffractometry pattern of ceramics sintered at (a) 1600°C for 5 min from milled powders; (b) 1780°C for 5 rnin from as-received powders (inner region); and (c) 1780°C for 5 rnin from as-received powders (outer region).

Nanocomposites and Nanostructures

Fig.7. Fracture surface of composite (a) sintered at 1780°C for 5 rnin from as-received powders; (b) sintered at 1600°C for 5 rnin from milled powders.

407

Thermodynamic analysis of the possible reactions in a Si-YAI-O-N-C system has been reported in many reference^.'^.'^.'^-^^ Many possible reactions between the added carbon powder and the other starting powders may occur during sintering: C(s)

Fig. 8. Transmission electron microscopy images of the nano/nanocomposite sintered at 1600°C for 5 min from milled powders. (a) Lowmagnification; (b) high-resolution transmission electron microscopy image indicating the Si3N4and carbon grains.

of the mechanical properties. Both Si3N4and carbon have narrow grain size distributions. The average grain diameters are 10 and 70 nm for carbon grains and silicon nitride grains, respectively. Monolithic nano-sized silicon nitride ceramics have been prepared using the same methods. The average grain diameter of Si3N4 grains is 68 nm.I6 Although a longer milling time has been used, the grain diameter of Si3N4 grains in the obtained nano/nanocomposite is similar to that in monolithic silicon nitride nano-ceramics, because the carbon decreases the milling efficiency by acting as a lubrication agent. Figure 8 shows the TEM images of a nano/nano-composite obtained in this research. Figure 8(a) shows a typical nano/ nano-composite structure, in which nano-sized carbon particles are homogeneously dispersed at grain boundaries. As shown in the high-resolution image of Fig. 8(b), the fringe spacing of 0.34 nm is typical of graphite, confirming that the nano-grains are carbon. The layer-structure nano-C grains are composed of a disordered layer, which existed in the grain boundaries of Si3N4 grains. The structure of carbon shows that it is partially amorphous, which is in good agreement with the G band and the D band in the Raman spectra. 408

+ Si02(s)++SiO(g)+ CO(g)

(1)

All reversible reactions proceed to the right only if the gaseous reaction products (CO, SiO, A120, Y 2 0 , N2) could be removed simultaneously. It is very difficult for the gaseous reaction products to escape freely by solid-state diffusion. Hence, the reactions can be suppressed by a high densification rate. In this research, the sintered density of milled powders increased to more than 90% pth within 1 min, changing the ceramics from open to closed porosity. However, only this cannot entirely prevent all reactions from occurring. As-received powders could also be nearly densified within 1 min at a higher sintering temperature of 1780”C, whereas the Si3N4/Ccomposite still could not be obtained due to the high reactivity of carbon particles at a high temperature. The Sic is formed through reaction (4). There is no carbon or Sic in the inner region and the amount of newly forming S i c phase is low, indicating that reactions (1)-(3) might occur during sintering, and some of the free carbon is transformed into volatile gas and escapes from the specimen. Hence, the achievement of a Si3N4/C nano/nano-composite could also be related to two other reasons: (I) a lower densification temperature, which leads to small driving forces for the reactions proceeding to the right14 and (2) incorporation of most of the carbon particles into the amorphous phase during milling, which prevents all the possible gaseous reaction products from escaping in the beginning. In conclusion, the high densification rate, incorporation of carbon particles into the amorphous phase, and a lower sintering temperature prevent all reactions from occurring, making it possible to prepare a dense Si3N4/Cnano/ nano-composite. The nano/nano-composite obtained has a hardness of 14.7 GPa, a density of 3.15 g/cm3, and Young’s modulus of 210 GPa. Compared with conventional Si3N4 ceramics,2 it has a lower density, lower Young’s modulus, and similar hardness. The low Young’s modulus will lead to improved thermal shock fracture resistance, and the existence of nano-carbon grains is expected to improve contact-damage resistance, and tribological properties. Detailed study of the effect of carbon addition on the mechanical, contact-damage resistance, and tribological properties of Si3N4/Cnano/nano-composite will be carried out soon. IV.

Conclusion

We have demonstrated that a nano-Si3N4/5wt% nano-C composite can be fabricated by a novel method. Nano-sized carbon particles with an average grain diameter of 10 nm were homogeneously dispersed at the nano-sized (- 70 nm) Si3N4 grain boundaries. The nano/nano-composite has a lower density, lower Young’s modulus, and similar hardness compared with conventional Si3N4ceramics.

References ‘F. L. Riley, “Silicon Nitride and Related Materials,” J . Am. C e r m . Soc., 83 [2] 24545 (2000).

*G. Petzow and M. Herrmann, “Silicon Nitride Ceramics.” Strucr. Bun
47-167 (2002).

’M. Sternitzke. “Structural Ceramic Nanocomposites,” J . Eur. Ceram. Soc., 17

[9] 1061-82 (1997).

4C.C Liu and J.L Huang, “Microelectrode Discharge Machining of TiN/Si3N4 Composites,” Br. Ceram. Trans., 99 [4] 149-52 (2000).

Progress in Nanotechnology:Processing

’T. Kusunose, T. Sekino. Y. H. Choa, and K. Niihara, “Fabrication and Microstructure of Silicon Nitride/Boron Nitride Nanocomposites,” J. Am. Cerum. Soc., 85 [I I] 2678-88 (2002). ‘P. Anderson and K. Holmberg, “Limitations on the Use of Ceramics in Unlubricated Sliding Applications Due to Transfer Layer Formation.” Wear, 175 [ 1-21 1-8 (1994). 7A. Gangopadhyay and S. Jahanmir, “Friction and Wear Characteristics of Silicon Nitride-Graphite and AlumindGraphite Composites,” Triho. Trans., 34 [2] 25745 (1991). ‘A. Skopp and M. Woydt, “Ceramic-Ceramic Composite-Materials with Im roved Friction and Wear Properties,” Tribo. fnt., 25 [I] 61-70 (1992). Iizuka and H. Kita, “Tribological Behavior of Mo5Si3 Particle Reinforced Si3N4Matrix Composites,” Wear. 258 ( S 6 ] 877-89 (ZOOS). ‘OH. Hyuga, K. Hirao, M. 1. Jones, and Y. Yamauchi, “Processing and Tribological Properties of Si,Nd/Carbn Short Fiber Composites,” J. Am. Ceram. Soc., 861p] 1081-7 (2003). H. Hyuga, M. 1. Jones, K. Hirao, and Y. Yamauchi, ‘‘Influence of Carbon Fibre Content on the Processing and Tribological Properties of Silicon Nitride/ Carbon Fibre Composites,” J. Eur. Cerum. SOC.,24 [5] 877-85 (2004). 12K. Niihara, “New Design Concept of Structural Ceramics-Nanocomposites,” J. Ceram. SUC.Jpn.. 99 [lo] 974-82 (1991). ”X. Wang. N. P. Padture, and H. Tanaka, “Contact-Damage-Resistant Ceramic/Single-Wall Carbon Nanotubes and Ceramic/Graphite Composites,” N a f . Mafer.,3 [8] 5 3 9 4 (2004).

4.

Nanocomposites and Nanostructures

I4M. Herrmann. C. Schuber. A. Rendtel, and H. Hubner. “Silicon Nitride/ Silicon Carbide Nanocomposite Materials: 1. Fabrication and Mechanical Properties at Room Temperature,” J. Am. Ceram. Sor., 81 [5] 1095-108 (1998). I5J. F. Yang. G. J. Zhang, N. Kondo. and T. Ohji. “Synthesis and Properties of Porous Si3N4/SiCNanocomposites by Carbothermal Reaction between Si3N4and Carbon,” Acra Muter.. 50 [I91 483140 (2002). I6X. Xu. T. Nishimura, N. Hirosaki, R.4. Xie, Y. C. Zhu, and Y. Yamamoto. “New Strategies to Prepare Nano-Sized Silicon Nitride Ceramics,” J. Am. Ceram. Soc.. 88 [4] 934-7 (2005). I7X. Xu, T. Nishimura, N. Hirosaki, R.-J. Xie. Y. Yamamoto, and H. Tdnaka, “Fabrication of p-Sialon Nano-Ceramics by High-Energy Mechanical Milling and Spark Plasma Sintering,” Nmorerknology, 16 [9] 1569-73 (2005). ‘‘L. G. Cancado, M. A. Pimenta. B. R. A. Neves. M. S. S. Dantas. and A. Jorio, “Influence of the Atomic Structure on the Raman Spectra of Graphite Edges.” Pkvs. Rev. L e f t , 93 [24] 247401 (2004). I9P. Greil, G . Petzow, and H. Tanaka, “Sintering and HlPing of Silicon Nitride Silicon Carbide Composite Materials,” Ceram. Int., 13, I s 2 5 (1987). ’OR. Wusirika, “Problems Associated with the Melting of Oxynitride Glasses,” J. Am. Ceram. Soc., 67, C232-3 (1984). ”K. Watari, M. Kawamoto, and K. Ishizaki. “Carbon Behavior in Sintered Silicon Nitride Grain Boundaries.” Mafer. Sci. Eng.. AlW. 89-95 (1989). 22T.Grande, H. Sommerset, E. Hagen, K. Wiik, and M. A. Einarsrud, “Effect of Weight Loss on Liquid-Phase-Sintered Silicon Carbide,” J. Am. Ceram. Soc., 80 [4] 1047-52 (1997). 0

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Conversion of Bamboo to Biomorphic Composites Containing Silica and Silicon Carbide Nanowires Teresa L. Y. Cheung and Dickon H. L. Ngt Department of Physics, The Chinese University of Hong Kong, Shatin, New Territories, Hong Kong, China

Luk Bamboo had been converted into biomorphic composites containing high-purity cristobalite-SiOz and bsilicon carbide (Sic) nanowires after sintering at 1200" and 1400°C, respectively. The fabrication process was simple, in which no catalyst was needed and the sintering temperature was low. The procedure included pyrolysis of biotemplates, infiltration of a Si-containing reactant, and sintering. Both the S O t and Sic nanowires were grown by a two-stage growth mechanism, in which the impurities from raw bamboo acted as catalysts. We successfully demonstrated that inexpensive Luk bamboo could provide a breakthrough, cost-effective, and eco-friendly route for self-assembling onedimensional nano-structures in highly porous biomorphic materials.

I. Introduction

N

UMEROUS

kinds of naturally existing biological materials possess unique and sophisticated structures, which have been developed and optimized by nature through ages. By mimicking the hierarchically built structural morphologies of biological materials, smart materials with fascinating designs can be achieved. Bones, nacre, sea urchin teeth, and other tough hybrid materials have been studied for their specific structures, compositions, and properties.' Particularly, attention has been focused on their biomineralization processes, i.e., the self-assembling nucleation and growth of mineralogical phases guided by organic templates to form unique micrometer-sized or nanometer-sized architectures. Although the details of these metamorphoses remain unclear, many attempts have been made to convert biological materials into biomorphic products for practical industrial applications. For example, tobacco mosaic virus was used as a template to produce nickel and cobalt nanowires2 Another example is the conversion of wood into biomor hic alumina: silicon carbide or titanium carbide103'Pcomposites.In particular, sub-wavelen th silica nanowires have attracted increasing attention12-" due to their luminescence and wave-guiding properties for compact-sized photonic devices. With high-strength and high-temperature stability, Sic also provides a wide range of applications in refractory, reinforcements, medical implantations: and electronic

biomorphic composites. Previous works only showed that biomorphic material acted as a template for composite synthesis. Our work demonstrates that Si02 and Sic nanowires could be self-assembled on the highly porous biomorphic template.

11. Experimental Procedure

( 1 ) Materials Luk Bamboo (Bumbusu Textilis var. Fusca) was collected from the hillside of the Chinese University of Hong Kong campus. The 98% tetraethylorthosilicate (TEOS) solution (Batch No. 06426BC, Sigma-Aldrich, St. Louis, MO) was used for infiltration, and it provided the Si source for the Si02 and Sic nanowires in the composites. (2) Fabrication Procedures Freshly cut 5-mm thick bamboo slices were rinsed with water before use. The pyrolysis process was performed at 1100°C for 10 h in an Ar atmosphere to remove water and certain organic groups from the fresh bamboo slices so that charcoal biotemplates were produced. A low heating rate of l"C/min was used in order to minimize thermal stress. Pyrolyzed biotemplates were then soaked in the TEOS solution for 5-7 days to allow infiltration of the TEOS reactant. Thereafter, the infiltrated biotemplates were sintered in an Ar atmosphere at 1200"-1500"C for different durations ranging from 1 to 10 h.

( 3 ) Characterization Scanning electron microscopy (SEM; Model 1450VP, LEO, Cambridge, U.K.) and transmission electron microscopy (TEM; CM120, Philips, Hillsboro, OR), both equipped with an energy-dispersive X-ray spectrometer (EDS), were used to characterize the morphology and chemical composition of the samples. X-ray diffractometry (XRD) was performed on the sintered products by a Huberadiffractometer (Rimsting, Germany) with CuKcr (1.5405981 A) radiation, and the X-ray generator (RU-300, Rigaku, Spring, TX) was operated at 40 kV and 160 mA.

device^.'^-'^

Some common fabrication approaches used to transform wood into Si02 and S i c biomorphic composites are gas-phase liquid-phase infiltrati~n',~ and sol-gel methods.'.' The gas-phase and liquid-phase infiltration methods often require repeated infiltrations, and show low conversion efficiency? Sol-gel methods excel due to their uniformity and relatively low reaction temperature, but highly accurate and specific conditions are usually needed.',' In this work, we present a simple and convenient route to transform bamboo into nano-structured R. Riedel--contributing editor

Manuscnpt No. 21942. Received June 21.2006; approved September 14,2006. 'Author lo whom correspondence should be addressed. e-mail [email protected]

Nanocompositesand Nanostructures

111. Results and Discussion

( 1 ) Bamboo Biotemplates After pyrolysis, there was significant volume shrinkage in the bamboo slices. The weight loss in the samples was estimated to be about 85 wt%. However, as examined by SEM, the tubular structure of bamboo was still retained as shown in Fig. l(a). Figure l(b) shows the lateral structure, and some tiny pores were observed on the channel walls as shown in the inset. EDS was conducted, and the EDS spectrum revealed that the biotemplate mainly contained carbon with < 10 at.% impurities, including Ca, Si, and 0 (Fig. l(c)). The XRD diffraction result obtained from the biotemplate is presented in pattern (a) in Fig. 2. The two broad peaks located at 22" and 44" corresponded to C(ooo2) and C(ooo4)of the carbon biotemplate.63' 411

Fig. 1. SEM images showing the (a) cross-sectional and (b) lateral views of the raw bamboo biotemplate after pyrolysis and (c) the corresponding SEMEDS spectrum. SEM, scanning electron microscopy; EDS, energy-dispersiveX-ray spectrometer.

( 2 ) Silica Nanowires After infiltration of TEOS and sintering at 1200°C for 10 h, a thin white layer was found on the surface of the biotemplate. As shown in the SEM image in Fig. 3(a), a large number of nanowires with an average diameter of -200 nm and a length of -20 pm were found on the walls of the tubular structure. The EDS revealed that these nanowires contained Si and 0. As can be seen from the XRD pattern (b) in Fig. 2, the peaks at 21.6" and 35.6" indicated that this sample contained the p-Si02 cristobalite. Some of the nanowires were extracted from this sample and examined by TEM (Fig. 3(b)). The EDS spectrum is shown

0 p-Sic SiO, (cristobalite)

20

25

30

35

40

45

50

55

60

65

70

75

2 0 (degree)

Fig. 2. X-ray diffractometry patterns of (a) raw charcoal biotemplate, and the infiltrated biotemplates sintered at (b) 1200"C, (c) 1300"C, (d) 1400°C for 10 h, and (e) 1500°C for 1 h. 412

in Fig. 3(c), which shows that the body of the SiOz nanowire contained Si and 0. Selected area electron diffraction (SAED) was conducted on one of these nanowires. The clear and single SAED pattern shown in the inset of Fig. 3(b) confirmed that the wire was single-crystal SO2. It was also observed that many of these Si02 nanowires exhibited round or ellipsoidal droplet tips. EDS analysis on the tip indicated the presence of Si, 0,and < 10 at.% of impurities including Mg and Ca (Fig. 3(d)). Using a similar procedure, another infiltrated sample was sintered at 1300°C for 10 h. As shown in Fig. 4, a combination of Si02 clusters and nanowires was obtained in this sintered sample. Its XRD pattern was similar to that of the sample sintered at 1200°C for 10 h (pattern (c) in Fig. 2).

( 3 ) S i c Nanowires Some infiltrated samples were sintered at higher temperatures. Figure 5(a) shows the sample sintered at 1400°C for 10 h, in which long nanowires with a diameter of 20&500 nm were observed. For the nanowire sample sintered at 1500°C for 1 h shown in Fig. 6(a), the diameter was smaller (- 50 nm). The XRD pattern (e) in Fig. 2 containing peaks at 35.7", 60.0", and 7 1.8" indicated that these nanowires contained cubic-structured O-SiC. Apart from these three peaks, one extra shoulder peak (labeled S) located at 33.6" in the XRD pattern was a characteristic peak indicating typical stacking faults on the (1 11) planes in cubic p-SiC.'8,'9 The TEM image in Fig. 5(b) shows the morphology of a typical S i c nanowire in the biotemplate sintered at 1400°C.The SAED pattern (inset in Fig. 5(b)) taken o n the body of this nanowire demonstrated that it was a singlecrystal O-SiC. At the same time, we observed that most of the S i c nanowires had smooth round tips (Fig. 5(b)). I n the EDS spectra shown in Fig. 5(c), mainly Si and C peaks were found in the body of the wire. On comparing the intensities of the EDS spectra, the Si to C ratio was relatively higher on the catalystfree smooth round tip than on the body. However, some of these Progress in Nanotechnology: Processing

Fig.3. (a) SEM image of the sample sintered at 1200°C for 10 h. (b) TEM image of the top portion of a nanowire extracted from the sample shown in (a). (c) and (d) are the TEM-EDS spectra taken at the body and the droplet tip of the silica nanowire shown in (b), respectively. SEM, scanning electron microscopy; EDS, energy-dispersive X-ray spectrometer; TEM, transmission electron microscopy.

nanowires were found to have a droplet tip as shown in Fig. 5(d). At the catalyst (droplet) tip shown in Fig. 5(d), elemental Fe, Si, and Ni were detected by EDS (Fig. 5(e)). Figure 6(b) shows a high-resolution TEM image of a Sic nanowire in the sample sintered at 1500°C for 1 h. The SAED pattern (inset in Fig. 6(b)) exhibited an extra set of diffraction spots alon the (1 10) zone axis that indicated twinning in this nanowire." A magnified image of region c (circled region) in this wire is shown in Fig. 6(c). Based on the lattice spacing measurement, it was evident that the growth direction of these nanowires was [ 11I] with a d-spacing of 0.25 nm. In the same image, a twin boundary (marked T) and several stacking faults (marked F) are also shown.

( 4 ) Growth Mechanism To explore the formation mechanism of the nanowires, we first attempted to confirm the origin of the reactants, i.e., the Si sources for SiO2 and Sic. Bamboo itself contained some inherent SiOz, but the amount was not sufficient to form the abundant SiOz or Sic nanowires as observed. A control experiment was performed to demonstrate the requirement of an external Si supply. Biotemplates, without being subjected to infiltration of TEOS, were sintered separately at 1200" and 1400°C for 10 h, and no Si-containing nano-structures were observed in these samples. Most of the nanowires were found at the sample surface and the walls of the large channel. Attempts were made to break the samples to investigate the interior region, but much fewer nanowires were found. EDS analyses suggested that the bulk was mainly made up of C, indicating that the large channels allowed much easier infiltration of TEOS solution. TEOS was physically infiltrated into the biotemplates as shown in Fig. 7(a). The solution was mixed with the impurities on the surface of the biotemplates. As TEOS vaporized at 160"C, the gaseous molecules slowly decomposed into SiOz and other components by thermal decomposition. The reaction was given by2':

-

Fig. 4. Scanning electron microscopy image of the sample sintered at 1300°C for 10 h.

Nanocomposites and Nanostructures

where g, 1, or s indicated that the component was in the form of gas, liquid, or solid, respectively. In the present setup, it was expected that the nanowires were formed by a two-stage growth mechanism, which was analogous to the Sic whiskers produced from rice hulls2' and coconut shells.22In the first stage, illustrated in Fig. 7(b), impurity droplets were formed and acted as catalytic nucleation sites. In the second stage, the nanowire growth followed the vapor-liquid-solid (VLS) mechanism. 413

Fig.5. (a) Scanning electron microscopy image of the sample sintered at 1400°C for 10 h. (b) Transmission electron microscopy (TEM) image of the silicon carbide (Sic) nanowire with a Si-rich tip and (c) the corresponding energy-dispersiveX-ray spectrometer (EDS) spectra taken on the tip and on the body. (d) TEM image of the S i c nanowire with a catalytic droplet tip and (e) the corresponding EDS spectrum of the catalytic tip.

Schematic diagrams of the growth of the Si02 and S ic nanowires are shown in Figs. 7(b) and (c), respectively. ( A ) Growth of SiOl Nanowires: In the first stage, at an elevated temperature, the impurities mixed with Si02 and formed discrete liquid droplets on the surface of the tubular structures in the biotemplates (Fig. 7(b)). As Mg and Ca impurities were detected at the tips of the Si02 nanowires, we proposed that they had played a major role in the growth process. Such a claim was supported by the observation of droplet-like particles on the surface of the infiltrated biotemplate that was sintered at 700°C for only 1 h, as shown in Fig. 8(a). According to a recent report,23the following reactions may have occurred: Si,O (8) + SL-1 (8) 2Si0 (g) + Si (s) 414

+ SiO (8)

+ Si02 (s)

(x

> 1)

(2)

(3)

As illustrated in Fig. 7(b)(i), the impurity droplets may have been rapidly surrounded by the SiO vapor species produced according to Eq. (2). SiO acted as a gaseous intermediate and diffused into the impurities droplets, which then acted as nucleation sites for the wire growth during sintering. Based on the fact that the tip of the S O 2 nanowires usually contained droplet tips (Figs. 8(b) and (c)), we proposed that the growth was analogous to the VLS process, in which Si02 precipitated out upon saturation in the impurity droplets according to Eq. (3). As the amount of solidified SiOz increased, it pushed up the impurity droplet and formed the basis for the SiOz nanowire (see Fig. 7(b)(ii)). ( B ) Growth of S i c Nanowires: The formation mechanism of Sic nanowires was similar to that of the SiOz nanowires as described earlier, only that it occurred at a higher temperature. Based on the fact that Fe-rich seeds were observed at the tips of some of the Sic nanowires, we propose that FeSi2 Progress in Nanotechnology: Processing

Fig. 6. (a) Scanning electron microscopy and (b) high-resolution transmission electron microscopy (HRTEM) images of the sample sintered at 1500°C for 1 h. (c) A magnified region from (b).

droplets had participated in the growth process. As the eutectic point of FeSi2 was 1207°C,24it acted as the liquid-phase catalytic growth center. However, the majority of the Sic nanowires possess catalyst-free tips, and a relatively high Si to C ratio (see Fig. 5(c)) was detected at the tips. In such a case, we propose that silicon-rich droplets could also act as the catalytic growth centers (see Fig. 7(~)(i)).'~-~' When the sintering temperature was increased to 1400°C and held for 10 h, the initially formed Si02 nanowires were melted, and the molten silica reacted with the

biotemplate. Sic were produced by the silica carbotherma1 reduction; accordingly:

+

Si02 (s, I ) C (s) + SiO ( g ) SiO ( g )

+ CO ( g )

+ 2C (s) + Sic (s) + CO (8)

SiO2 (s, I )

+ 3C (s)

+

+

Sic (s) 2CO (8)

(4) (5)

(6)

It was described by Eq. (4) that at hi her temperature, molten SiOz was reduced by C to form SiO?. The gaseous products SiO diffused into the catalytic growth centers (Fig. 7(c)). As the reaction proceeded, the silica layer became thinner, making the carbon from the biotemplate come in contact with the droplets more effectively, and ready for more intensive reactions. According to Eq. (S), the dissolved SiO gas in the droplet reacted with C to form solid Sic. Summing up Eqs. (4H6) were the overall reaction. The solidified Sic accumulated and pushed up the catalytic droplet, and formed the basis for a growing Sic nanowire (Fig. 7(c)(ii)).

8

IV. Conclusions

Fig.7. Schematic diagrams showing the growth of the sio2 and silicon carbide (Sic) nanowires. (a) Tetraethylorthosilicate (TEOS) was physically infiltrated into the biotemplates, and then slowly decomposed into SOz. (b) Si02 or (c) Sic nanowires were grown by the twostage growth at different temperatures.

Nanocomposites and Nanostructures

We had successfully demonstrated that inexpensive Luk Bamboo provided the impurities as a catalyst for the growth of Si02 (0-cristobalite) and p-Sic nanowires at different temperatures. Conventional fabrications of nano-structured materials had always used high-purity chemicals and catalysts under accurately controlled experimental conditions, which significantly increased the cost of production. In our work, no additives were required, and the cost of starting materials was relatively lower. This facile method suggested a breakthrough route for PreParing nano-structures in a carbon matrix in porous biological structures. 415

Scanning electron microscopy images o f the samples sintered at (a) 700”C, (b) 1300°C. a n d (c) 1400°C f o r 1 h . T h e o p e n squares in (a) indicate the impurity droplets, a n d the a r r o w s in (b) and (c) indicate the S b M g - C a droplet tips on the SiOz nanowires.

Fig.8.

Acknowledgments This work was supported by the United College-Lee Hymn Foundation, Endowment Fund Research Grant Scheme (Project code: CAI 1066), and the RGC Earmarked Research Grant (Project code: 2150421/4233-04E).

”L. M. Tong, J. Y. Lou, Z. Z. Ye, G. T. Svdcha, and E. Mazur, “SelfModulated Taper Drawing of Silica Nanowires,” Nanotechnology, 16 [9] 1445-8

(2005).

l3Y. C. Lin and W. T. Lin, “Growth of Si02 Nanowires Without a Catalyst Via Carbothermal Reduction of CuO Powders,” Nanorechnology, 16 [9] 1648-54 (2005).

References ‘A. H. Heuer. D. J. Fink, V. J. Laraia, J. L. Arias, P. D. Calvert, K. Kendall, G. L. Messing, J. Blackwell, P. C. Rieke, D. H. Thompson, A. P. Wheeler, A. Veis, and A. 1. Caplan, “Innovative Materials Processing Strategies: A Biomimetic Approach,” Science, 255 [5048] 1098-1 105 (1992). ’M. Knez, A. M. Bittner, F. Boes, C. Wege, H. Jeske, E. Maih. and K. Kern, “Biotemplate Synthesis of 3-nm Nickel and Cobalt Nanowires,” Nano Lett., 3 [8] 107%82 (2003).

’C. Z. Li and J. H. He, “Easy Replication of Pueraria Lobata Toward Hierarchically Ordered Porous y-A1203,” Longmuir, 22 [6] 2827-31 (2006). 4P. Gonzalez, J. Serra. S. Liste, S. Chiussi, B. Leon, M. Perez-Amor, J. Martinez-Fernandez, A. R. Arellano-Lopez, and F. M. Varela-Feria. “New Biomorphic SIC Ceramics Coated With Bioactive Glass for Biomedical Applications,” Biomaterials, 24 [26] 4827-32 (2003). 5H. Sieber, C. Hoffmann. A. Kaindl, and P. Greil, “Biomorphic Cellular Ceramics,” Adv. Eng. Mater., 2 [3] 105-9 (2000). ‘E. Vogli, J. Mukerji, C. Hoffman, R. Kladny, H. Sieber, and P. Greil, “Conversion of Oak to Cellular Silicon Carbide Ceramic by Gas-Phase Reaction With Silicon Monoxide.” J . Am. Ceram. Sac., 84 [6] 1236-u) (2001). ’A. Munoz, J. M. Fernandez, and M. Singh, “High Temperature Compressive Mechanical Behaviour of Joined Biomorphic Silicon Carbide Ceramics.” J. Eur. Ceram. Sac., 22. 2727-33 (2002). *Y. S. Shin, C. M. Wang, and G . J. Exarhos, “Synthesis of SIC Ceramics by the Carhothermal Reduction of Mineralized Wood with Silica.” Adv. Muter., 17 [I] 73-7 (2005).

’J. M. Qian, J. P. Wang, and 2. H. Jin, “Preparation of Biomorphic S i c Ceramic by Carhothermal Reduction of Oak Wood Charcoal,” Mater. Sci. Eng. A , 371 [2] 229-35 (2004).

‘“B. H. Sun, T. X. Fan, and D. Zhang, “Porous TIC Ceramics Derived from Wood Template,” J. Porous Mater.. 9 [4] 275-7 (2002). “T. Ota, M. Imaeda, H. Takase, M. Kobayashi, N. Kinoshitd, T. Hirashita, H. Miyazaki, and Y. Hikichi, “Porous Titania Ceramic Prepared by Mimicking Silicified Wood,” J . Am. Ceram. Soc., 83 [6] 1521-3 (2000).

416

I4C. L. Shao, H. Y. Kim. J. Gong, and D. Lee, “A Novel Method for Making Silica Nanofibres by Using Electrospun Fibres of Polyvinylalcohol/Silica Composite as Precursor,” Nanotechnology. 13 [5] 635-7 (2002). ”G. W. Ho, A. S. W. Wong, D. J. Kang, and M. E. Welland, “Three-Dimensional Crystalline SIC Nanowire Flowers.” Nanotechnology. 15 [S] 996-9 (2004). “H. K. Seong, H. J. Choi, S. K. Lee, J. 1. Lee, and D. J. Choi, “Optical and Electrical Transport Properties in Silicon Carbide Nanowires,” Appl. Phys. Lett., 85 7 1256-8(2W). .&!I S . Wu, S. Z. Deng, N. S. Xu, J. Chen, and J. Zhou, “Needle-Shaped Silicon Carbide Nanowires: Synthesis and Field Electron Emission Properties,” Appl. P h g Lert., 80 [20] 3829-31 (2002). K. Koumoto, S. Takeda. and C. H. Pai, “High-Resolution Electron Microscopy Observations of Stacking Faults in p-Sic,’’ J . Am. Ceram. Soc., 72 [lo] 1985-7 (1989).

I9H. I. Dai, E. W. Wong. Y. Z. Lu, S. S. Fan. and C. M. Lieber, “Synthesis and Characterization of Carbide Nanorods,” Nature, 375 [6534] 769-72 (1995). 20 M. M. IslamRaja, C. Chang, J. P. McVittie, M. A. Cappelli, and K. C. Saraswat, “Two Precursor Model for Low-Pressure Chemical Vapor Deposition of Silicon Dioxide from Tetraethylorthosilicate,” J . Vuc. Sci. Techno/. B, I 1 [3] 720-6 (1993).

”H. J. Choi and J. G . Lee, ‘Stacking Faults in Silicon Carbide Whiskers,” Ceram. hit.. 26 [ I ] 7-12 (2000). 22 A. Selvam, N. G. Nair, and P. Singh, “Synthesis and Characterization of SIC Whiskers from Coconut Shells,” J . Mater. Sci. Lert., 17, 57-60 (1998). 23 Y. N. Xia, P. D. Yang, Y. G . Sun, Y. Y. Wu, B. Mayers, B. Gates, Y. Yin, F. Kim. and H. Q. Yan, “One-Dimensional Nanostructures: Synthesis, Characterization, and Applications,” Adv. Muter., 15 [ S ] 353-89 (2003). 24D.P. Yu, Q. L. Hang, Y. Ding, H. Z. Zhang, 2. G. Bdi,J. J. Wang. Y. H. Zou, W. Qian, G. C. Xiong, and S. Q. Feng, “Amorphous Silica Nanowires: Intensive Blue Light Emitters,” Appl. Phvs. Lert., 73 [21] 30768 (1998). 25T. Seeger. P. Kohler-Redlich, and M. Ruhle, “Synthesis of Nanometer-Sized S i c Whiskers in the Arc-Discharge,” Adv. Muter.. 12 [4] 279-82 (2000). 26T. Belmonte, L. Bonnetain. and J. L. Ginous, “Synthesis of Silicon Carbide Whiskers Using the Vapour-Liquid-Solid Mechanism in a Silicon-Rich Droplet,” J . Marer. Sci., 31, 2367-71 (1996). 0

Progress in Nanotechnology: Processing

Novel Processing to Produce PolymerKeramic Nanocomposites by Atomic Layer Deposition Xinhua Liang, Luis F. Hakim,* Guo-Dong Zhan,* Jarod A. McCormick, Steven M. George, and Alan W. Weimer*>+ Department of Chemical and Biological Engineering, University of Colorado, Boulder, Colorado 80309

Joseph A. Spencer I1 and Karen J. Buechler ALD NanoSolutions, Broomfield, Colorado 80020

John Blackson and Charles J. Wood Dow Chemical Company, Midland, Michigan 48667

John R. Dorgan Department of Chemical Engineering, Colorado School of Mines, Golden, Colorado 80401

An innovative process to uniformly incorporate dispersed nanoscale ceramic inclusions within a polymer matrix was demonstrated. Micron-sized high density polyethylene particles were coated with ultrathin alumina films by atomic layer deposition in a fluidized bed reactor at 77°C. The deposition of alumina on the polymer particle surface was confirmed by Fourier transform infrared spectroscopy and X-ray photoelectron spectroscopy. Conformal coatings of alumina were confirmed by transmission electron microscopy and focused ion beam cross-sectional scanning electron microscopy. The results of inductively coupled plasma atomic emission spectroscopy suggested that there was a nucleation period. The results of scanning electron microscopy, particle size distribution, and surface area of the uncoated and nanocoated particles showed that there was no aggregation of particles during the coating process. The coated polymer particles were extruded by a heated extruder at controlled temperatures. The successful dispersion of the crushed alumina shells in the polymer matrix following extrusion was confirmed using cross-sectional transmission electron microscopy. The dispersion of alumina flakes can be controlled by varying the polymer particle size.

I. Introduction

P

OLYMERIC materials

are widely used in packaging applications. Biomedical uses of plastic materials have been widespread and the combination of ceramics and certain polymers is the choice for medical The automobile industry also has embraced plastics to improve efficiency and improve manufacturing methods. However, the strength and some other properties, such as thermal stability, permeability to gases and organic solvents, and flame retardance of the pure polymer R. Riedel--contributing editor

Manuscript No. 21954. Received June 28, 2006 approved August 21,2M)6. Supported by the National Science Foundation, under Grant 0400292, US. Department of Energy, under STTR Grant DE-FG02-03ER86157, and the GAANN Program in Functional Materials, US. Department of Education. Support from Department ofEnergy does not constitute an endorsement by Department of Energy of the views expressed in the article. 'Member. American Ceramic Society. 'Author to whom correspondence should be addressed. e-mail: alan.weimer@ colorado.edu

Nanocornposites and Nanostructures

are often not enough for end use.4 Confinement of polymer and mineral pedigrees is one of the effective ways to improve material performance.' Work has primarily been done with nanoscopic montmorillonite clay."' Two widely adopted approaches to forming polymer/inorganic nanocomposites are high shear mixin of the preformed polymer with the ceramics (compoundingf-lo and in situ polymerization of monomer that has been premixed with the ceramics. Both approaches are feasible at the bench scale, but ceramics are not homogeneously dispersed in the polymer matrix at a nanoscopic level' and there are voids between ceramics and polymer. Commercialization will also require a low-cost continuous process. Previous studies show that the combined effects of aspect ratio and dispersion of clay particles ultimately control the mechanical properties of the nanocomposite, with dispersion playing a major role.576Relatively few compounding studies have appeared in the literature as a route toward polymer/ceramic nanocomposites. There is therefore a need to chemically bond ceramics and polymer and disperse ceramics homogeneously throughout the polymer matrix. A novel process to promote intimate mixing is to coat polymer particles with ultrathin, uniform ceramic films. The coated polymer particles can be extruded into pellets or films. During the high shear/high stress extrusion process, the shell on the polymer particle surface will crack and the shell remnants will be dispersed homogeneously throughout the polymer matrix. By means of this novel technique, the mechanical and barrier properties of this kind of polymer/ceramic nanocomposite may be further improved and new advanced features may appear. There are many problems in depositing inorganic films on polymer surfaces by conventional methods. Chemical vapor deposition (CVD) and plasma-enhanced CVD (PE-CVD) processing have been reported for polymer surface coating. However, typical CVD processes generally operate at temperatures ( - 3OO0-50O0C) much higher than the softening and melting temperatures of the polymers (- 125"-250°C). CVD techniques are not able to effectively control the use of precursor gases or to inherently control the location and the thickness of the ceramic film. In addition, both CVD and PE-CVD will leave defects and pinholes in the deposited inorganic films.12-I s Atomic layer deposition (ALD) provides unparalleled advantages over other techniques to deposit inorganic films on polymer surfaces.

'-''

417

ALD is a surface controlled layer-by-layer process, which deposits low impurity content, pin hole-free, conformal, and ultrathin flexible films.l”19 The film thickness is inherently controlled by self-limiting sequential surface chemical reactions, so precursors are used efficiently. ALD has been successfully demonstrated using a fluidized bed reactor (FBR).2&23A FBR has the main advantages of excellent gas/particle contact and thermal efficiency, and its control is easy due to stable operating conditions. AI2O3is non-flammable and has a melting point of 2050°C. The chemical and thermal stability of Al2O3 allows its application as a good diffusion barrierF4 From a toxicological viewpoint, A1203is non-toxic, but the montmorillonite clay can lead to toxic byproducts as the product ages, which may mean that many clay-based nanocomposites will never be suitable for food packaging applications. Therefore, AI2O3 is a good alternative to montmorillonite clay. High density polyethylene (HDPE) is a widely used polymer and a good candidate for experimentation. Polyethylene and A1203 are also biocompatible. Combining these two materials could make a stronger polymer with many potential applications. For example, along with the typical arthroplasty applications for polyethylene, successful biocompatibility has recently been observed for an A1203/polyethylene blood pump.3 The main objective of this research is to develop a new costeffective efficient process to fabricate uniform polymer/ceramic nanocomposites. In this paper, the successful deposition of ultra thin AI2O3films on micron-sized HDPE particles by ALD at the temperature of 77°C is reported, and the successful dispersion of A1203flakes in the polymer matrix following the extrusion process is demonstrated.

11. Experimental Procedure

AI2O3films have been deposited on several substrates, using repeated exposures of trimeth laluminum (TMA) and H 2 0 in an . ~ ’ ~ ALD ~ is derived from the ABAB.. . ...~ e q u e n c e . ~ ’ - ~ ~A1203 following binary CVD reaction: 2A1(CH3)3+ 3H20 --t A1203 + 6CH4

(1)

This binary reaction can be divided into two half-reactions:

+ Al(CH3), [AIOAI(CH3),]*+ CH4 (2) (B) AI(CH,)* + H2O AIOH* + CH4 (3) where * indicate the surface species.”” In each half-reaction, a (A) AIOH*

4

+

gas-phase precursor reacts with a surface functional group and forms CH4 as a by-product. The surface reaction continues until all the available surface functional groups have reacted. The experimental ALD-FBR is shown in Fig. 1. The reactor itself was composed of a 3.5 cm inside diameter stainless steel tube with a 10 pm pore size porous metal disc as the gas distributor. A 316 L porous metal filter element (1.9 cm ID x 15.24 cm long; 0.5 pm pore size) was used at the inside top of the reactor column to prevent particles from leaving the system. The reactor was encased by a clamshell-typefurnace and bolted to a platform that rested on four large springs. The reactor was maintained at low pressure by a vacuum pump (Model 2063, Alcatel, Paris, France), and the dosing header could also be pumped down directly using a smaller separate pump (Model 2008A, Alcatel). A vibration system (Model

Fig. 1. Schematic diagram of atomic layer deposition-fluidized bed reactor (ALD-FBR): ( I ) pressure transducers, (2) metal filter, (3) reaction column, (4)distributor plate, (5) vibro-motors, (6) spring supports, (7) pneumatic valves, (8) reactant containers, (9) mass flow controller.

41a

Progress in Nanotechnology: Processing

CD36210, Martin Engineering, Marine City, MI) was utilized to overcome some of the interparticle forces and improve the quality of fluidization. High purity N2 gas was used as the purge gas to remove the unreacted precursor and any CH4 formed during the reaction. The purge gas flow was fed in through the distributor of the reactor and its flow rate was controlled by a MKS" mass flow controller (Model 1179, MKS, Boulder, CO). Piezoelectric transducers (Model 902, MKS) were located below the distributor plate and at the outlet of the reactor column to measure the pressure drop across the bed of the particles. All valves used to provide the transient dosing were automatically controlled through Labview'" from National Instruments (Austin,TX). Pressure measurements were recorded to monitor the progress of each dosing cycle. Two different sizes of HDPE particles (Lyondell Chemical, Houston, TX) were used. One had an average size of 16 pm, and the other had an average size of 60 pm. 'The density of both primary particles was 952 kg/m3. The peak melting point was 134°C. For a typical run, about 20 g of HDPE particles were loaded into the reactor. The feeding lines were kept at about 70°C to avoid excessive absorption of H 2 0 on the internal walls of the system that could promote CVD reactions. The minimum pressure inside the reactor was about 10 Pa and the minimum fluidization superficial gas velocity was determined by measuring the pressure drop across the bed versus the N2 superficial gas velocity. Precursors, TMA (Sigma Aldrich, St. Louis, MO) and deionized H20, were fed separately through the distributor of the reactor using the driving force of their vapor pressures. The flow rate of TMA and H2O was adjusted using needle valves to ensure that a precursor pressure was high enough for particle fluidization. The reaction temperature was 77"C, which was lower than the softening/melting point of the HDPE particles. Before the reaction, the particles were dried at 77°C under a continuous N2 flow for 3 h. During each coating cycle, the precursors were fed for enough time so that saturation of all active sites occurred for every dose. A typical coating cycle occurred with the following sequence: dose TMA, purge N2, evacuate; dose H20, purge Nz, evacuate. In this manner, there is no overlap between the two reactants, and no CVD reactions occur. A Fourier transform infrared (FTIR) spectrometer (Model 750 Magna-IR, Thermo Nicolet, Waltham, MA) was used to analyze the composition of the HDPE particles before and after coating. The particles were milled with FTIR grade potassium bromide (Sigma Aldrich) to form a very fine powder. This powder was then compressed into a thin pellet using a hydraulic press and polished stainless steel die. An X-ray photoelectron spectroscopy (XPS) system (Model PHI 5600, Physical Electronics, Chanhassen, MN) with a high-energy resolution analyzer was used for this study. Aluminum concentration on HDPE particles was analyzed by inductively coupled plasma atomic emission spectroscopy (ICP-AES; Model ARL 3410+, Thermo Electron, Waltham, MA). Analysis by ICP-AES was achieved by placing the coated HDPE particles in a strong base solution (NaOH) to dissolve the A1203 films from the HDPE particles. The HDPE particle itself will not dissolve at normal laboratory conditions. The conformality of the A1203coatings on the HDPE particles was evaluated by transmission electron microscope (TEM; Model CM 10, Philips, Eindhoven, the Netherlands). The morphology of the HDPE particles before and after coating was investigated by scanning electron microscope (SEM; Model JSM-6400, JEOL, Tokyo, Japan). The size distGbution of HDPE particles was performed using an Aerosizer particle size analyzer (Model 3225, TSI, Shoreview, MN). Surface area analysis was, performed using a physisorption analyzer (Model Autosorb -1, Quantachrome, Boynton Beach, FL). The coated particles were extruded by a bench-sized, heated extruder (Bonnot, Uniontown, OH) at controlled temperatures. 16 and 60 pm HDPE particles were extruded at 135" and 175"C, respectively. To extrude a ribbon of polymer, the ribbon die was attached downstream of the heated barrel of the extruder. HDPE/A1203 nanocomposite films comprising various levels

Nanocomposites and Nanostructures

of concentration and morphologies of nanoscale ceramic flakes were formed. The structure information of the nanocomposite films was confirmed by cross-sectional TEM. 111. Results and Discussion

(1) Testfor Composition of A 1 2 0 3 Films on HDPE Particles The composition of HDPE particles (16 pm) before and after coating was characterized by ex sifu FTIR spectroscopy. As shown in Fig. 2, the FTIR spectrum of the reference alumina sample shows the A1203bulk vibrational mode at the frequency of IlOCL500 cm-l and the vibration of the OH group at the frequency of 370&3000 ern-'.'' No above-mentioned A1203 and OH group features are observed for uncoated HDPE particles. An A1203vibrational mode and a broad OH group feature appear for coated particles after 25 and 50 cycles. This is a direct confirmation of the composition of the A1203films on the polymer surface. For HDPE samples, the features at 3000-2800, 1460, and 720 cm-' are attributed to C-H stretching, deformation, and rocking modes of CH2 groups.25 XPS measurements were also performed on uncoated and A1203coated HDPE particles (I6 pm) after 50 cycles. The analysis was performed using an aluminum source, pass energy of 187.85 eV, and an energy step of 0.2 eV. In Fig. 3, the spectrum for the uncoated HDPE particles shows a photoelectron peak at 284.7 eV (C, Is). In contrast, the carbon spectrum for coated HDPE particles reveals much weaker photoelectron intensity at 284.7 eV. This reduction of carbon signal is expected if the A1203film conformally covers the entire polymer particle. The carbon XPS signal cannot be completely attenuated as some of it corresponds to surface carbon. Photoelectrons from the A1203-coated HDPE particles are observed at 118.7 eV (Al, 2s), 73.9 eV (Al, 2p) and 530.7 eV (0,1s). It is clearly evident that there is only a single peak centered at 73.9 eV, which corresponds to A 1 4 bonds of A1203. The absence of a shoulder region around 72.5 eV, which corresponds to AI-AI bonds, clear1 confirms that the aluminum metal is not present in our So, the XPS results corroborate the FTIR results and films!', verify the composition of deposited A1203 films on the HDPE particles.

(2) Uniformity of A1203 Films on HDPE Particles In order to study the uniformity of A1203films on HDPE particles, TEM analysis was performed at 100 kV on the coated particles (16 pm) after 50 cycles. The TEM image in Fig. 4 shows that an A1203film was successfully coated on the particle surface. The contrast between the film and the particle substrate is given by the difference in density between A1203 and HDPE. The thickness of the A1203 films is about 23+4 nm, which

4.0

I

3.0

2.0

4000

3500

3000

25M) 2000 1500 Wavenumbers(cm- 1)

1000

e,

2

e

2

U!

0.0 500

Fig. 2. Fourier transform infrared spectra of uncoated high density polyethylene particles, AlzO,-coated HDPE particles, and reference A1203 powders. 419

Fig. 3. X-ray photoelectron spectroscopy spectra of (a) uncoated and (b) AI2O3-coatedhigh density polyethylene particles (16 pm) after 50 cycles.

represents a growth rate of about 0.5 nm per coating cycle at this experimental condition. The A1203 films appear to be very uniform and smooth. Focused ion beam (FIB) cross-sectional SEM imaging allows precise observation at the edge interface of the polymer and A1203film. The FIB cross-sectional SEM image of HDPE particles (60 l m ) after 100 cycles is shown in Fig. 5. A1203islands began to grow below the polymer surface and the film merged into a linear layer as it grew. Approximately 35k7 nm thick A1203films were coated on the polymer surface. This thickness represents a growth rate of about 0.4 nm per coating cycle at this experimental condition. The SEM image also shows that the A1203films appear to be very uniform and smooth. The A1203film growth rate was much higher than the 0.1 10.13 nm per cycle of an ALD process reported in the literature.” Recent FTIR measurements of A1203 ALD on low density polyethylene (LDPE) indicated the presence of hydrogen-bonded H 2 0 molecules on the AI2O3 surface.25This higher growth rate may be explained by the presence of hydrogen-bonded H2 0 .This H 2 0can react with TMA to deposit additional A1203 by CVD.25Another reason is the increase in the surface coverage of reactants at lower Though the reaction kinetics is slower at lower temperatures, the growth rate is determined by the higher surface Also, it is important to mention that the growth rate of films may vary with the size and geometry of the substrate?2 For particles with a high ratio of curvature, more active sites on the surface are exposed to the gas phase reactants. The different initial surfaces may partly explain the discrepancy between A1203 growth rates on the HDPE particles and on some other substrates.

Fig. 5. Focused ion beam cross-sectional scanning electron micrograph of AI2O3-coatedhigh density polyethylene (HDPE) particle (60pm)after 100 cycles.

( 3 ) Nucleation and Linear Growth of A1203 Films After Nucleation The concentration of aluminum on HDPE particles was analyzed by ICP-AES. ICP-AES provides the concentration in parts per million (ppm) by mass of aluminum in relation to the HDPE particles. The ICP-AES aluminum concentration versus number of coating cycles is shown in Fig. 6. The average diameter of the particles was 16 pm. The lower growth rate of A1203 before 25 cycles shows that there is a delay before film growth starts, which verifies that a nucleation period is needed for the deposition of A1203 on an HDPE particle surface.32From this plot, the nucleation period is 10 cycles at this experimental condition. The AI2O3ALD is conventionally thought to begin with native hydroxyl groups on the surface. HDPE, however, is one kind of saturated hydrocarbon, which lacks typical chemical functional groups such as hydroxyl species that are necessary to initiate the growth of an inorganic film. So, the fundamental concept of Al2O3ALD cannot take place on the HDPE particle surface. The nucleation of A1203 ALD on HDPE requires a mechanism that does not involve the direct reaction between TMA and HDPE. Consequently, an alternative mechanism is needed to explain the A1203ALD on HDPE. HDPE has a porous surface, which is due to the interstitial space between individual molecules as HDPE does not have the regular lattice-type structure found in metals. Both HDPE and TMA are nonpolar, so it is expected that TMA has a reasonable solubility in the HDPE particles, and TMA can adsorb onto the surface of the polymer and subse uently diffuse into the nearsurface regions of the During the ALD reaction,

0

20

40

60

80

Coating cycles

Fig. 4. Transmission electron micrograph of A1203-coatedhigh density polyethylene particle (16 pm) after 50 cycles.

420

Fig. 6. Aluminum concentration on high density polyethylene particles versus coating cycles.

Progress in Nanotechnology: Processing

TMA will be first exposed to the HDPE particles and diffuse into the bulk of the polymer matrix; therefore, the incoming H20 will react efficiently with TMA molecules at or near the surface of the polymer particles and A1203 clusters will be formed. The pores on the particle surface will become smaller and will gradually close with progressive coating cycles. After several coating cycles, the A1203clusters will eventually merge to create a continuous adhesion layer on the polymer particle surface. This phenomenon can be observed in Fig. 5. AI2O3clusters with hydroxyl groups will provide a “foothold” for the deposition of A1203 films on the polymer. As shown in Fig. 6, the concentration of aluminum is almost directly proportional to the number of coating cycles after 25 cycles, which indicates a constant growth rate and a linear dependence between the film thickness and number of growth cycles after a nucleation period. The model of the predicted growth mechanism is illustrated in Fig. 7.

( 4 ) Efiect of Coating on Particle Size Distribution and Surface Area Fine particles will aggregate during fluidization because of interparticle forces, such as Van der Waals forces.2o SEM was used to analyze the morphology of the HDPE particles (16 pm) before and after coating. SEM analysis was performed at 15 kV. Figure 8 shows that no aggregates were coated; rather, particles were coated individually. This is also confirmed by the results of particle size distribution (PSD) of HDPE particles (I6 pm) before and after coating. The PSD curves for uncoated particles and A1203-coatedparticles after 50 cycles are shown in Fig. 9. As shown in the plot, the size of particles remains fairly unchanged after the coating process, meaning that no aggregates were being coated. If aggregates of particles were coated and glued together, the size distribution of particles after coating would drastically shift to the right. In addition to PSD analysis, Brunauer-Emmett-Teller (BET) measurements indicated that the surface area of A1203 coated HDPE particles (16 pm) after 50 cycles was 0.72k0.02 m2/g, which was very close to that of uncoated HDPE particles (0.70+0.03 m2/g). A drastic decrease in surface area, which did not occur, would have indicated necking of particles.20This result also indicated that the individual particles were coated as opposed to necking multiple particles together in the FBR. ( 5 ) Structure Information of Nanocomposite Films The 16 pm A1203-coatedHDPE particles after 75 cycles and the 60 pm A1203-coatedHDPE particles after 100 cycles were extruded to crush A1203shell coatings. Remnants of the crushed shell coatings were then dispersed throughout the polymer. The extruded nanocomposite films were cut using a microtome to achieve a thickness of approximately 100 nm for TEM analysis. The cross-sectional TEM images of the nanocomposite films are shown in Fig. 10. These two images show a scattering of nano-

Fig. 8. Scanning electron microscopy of (a) uncoated and (b) A1203coated high density polyethylene particles (16 pm) after 75 cycles.

sized inclusions of A1203throughout the samples. The brightest spots are areas in the films where the rough microtoming penetrated the films. In Fig. 10(a), the smaller image on the top left corner represents one of the A1203flakes at higher magnification, which indicates that A1203 flakes were formed of much smaller A1203 particles. The desired loading percent of A1203 can be controlled by adjusting starting polymer particle size. In the case of HDPE particles with the size of 16 pm, as shown in Fig. 10(b), more A1203flakes were dispersed in the matrix, and the A1203 flakes were dispersed more homogeneously. Hence, the dispersion of A1203flakes can be controlled by varying the polymer particle size.

IV. Conclusions This work represents the first successful attempt to fabricate polymer/ceramic nanocomposites by extruding nanocoated 1.2

x

0

20

40

Size (pin) Fig. 7. Proposed Al2O3growth mechanism.

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Fig. 9. Particle size distribution of uncoated and A1203-coated high density polyethylene particles (16 pm) after 50 cycles.

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References ‘M. A. Araujo. C. M. Vaz, A. M. Cunha, and M. Mota, “In-Vitro Degradation Behaviour of Starch/EVOH Biomaterials,” Polym. Deyrud. Stuhil., 73 [2] 23744 (2001).

’H. Warashina. S. Sdkano. S. Kitamura, K.4. Yaniauchi. J . Yamaguchi, N. Ishiguro. and Y. Hasegawa, “Biological Reaction to Alumina. Zirconia, Titanium. and Polyethylene Particles Implanted onto Murine Calvaria.” Biornuteriuh, 24 [21] 365S61 (2003). ’Y. Takami, T. Nakazawd, K. Makinouchi, J. Glueck. and Y. Nose, “Biocompatibility of Alumina Ceramic and Polyethylene as Materials for Pivot Bearings of a Centrifugal Blood Pump.” J. Uiomed Murrr. Res.. 36 [3] 381-6 (1997).

4J. H. Yeo. C. H. Lee, C.-S. Park, K.-J. Lee. J.-D. Nam, and S. W. Kim. “Rheological, Morphological, Mechanical, and Barrier Properties of PP/EVOH Blends,” Adv. Pulym. Techno!., 20 [3] 191-201 (2001). ’P. Maiti, K. Yamada, M. Okamoto, K. Ueda, and K. Okamoto. “New Polylactide/Layered Silicate Nanocomposites: Role of Organoclays,” Chem. Muter, 14 [I‘,1 465461 (2002). A. Usuki. Y. Kojima. M. Kawdsumi, A. Okada, Y. Fukushima, T, Kurauchi, and 0. Kamigaito, “Synthesis of Nylon 6-Clay Hybrid,” J . Murer. Res., 8 [5] 1179-84 (1993).

’Y. Kojima, A. Usuki, M. Kawasumi, A. Okada, T. Kurauchi. 0. Kamigdito, and K. Kaji, “Fine-Structure of Nylon-&Clay Hybrid,” J . Polym. Sci. PI. BPolym. Phys., 32 [4] 625-30 (1994). *P. B. Messersmith and E. P. Giannelis, “Synthesis and Barrier Properties of Poly(epsilon-Capro1actone)-Layered Silicate Nanocomposites,” J . Polym. Sci. Pol, Chem., 33 [7] 1047-57 (1995). 9J. W. Gilman, L. Jackson, A. 9. Morgan, R. Harris, E. Manias, E. P. Giannelis. M. Wuthenow, D. Hilton, and S. H. Phillips, “Flammdbility Properties of Polymer-Layered-Silicate Nanocomposites. Polypropylene and Polystyrene Nanocomposites,” Chem. Muter, 12 [7] 186673 (2000). “E. L. Cussler, S. E. Hughes, W. J. Ward, and R. Ark. “Barrier Membranes,” J. Membr. Sci., 38 [2] 161-74 (1988). “F. Breme, J. Buttstaedt, and G. Emig, “Coating of Polymers with TitaniumBased Layers by a Novel Plasma-Assisted Chemical Vapor Deposition Process,” Thin Solid Films,377, 755-9 (2000). ”A. P. Roberts. B. M. Henry, A. P. Sutton, C. R. M. Grovenor, G. A. D. Briggs, T. Miyamoto, M. Kano, Y. Tsukdhara, and M. Yanaka. ”Gas Permeation in Silicon-Oxide/Polymer (SiO,/PET) Barrier Films: Role of the Oxide Lattice, Nano-Defects, and Macro-Defects,” J . Membr. Sci., 208 [I-21 75-88 (2002).

Fig. 10. Cross-sectional transmission electron micrograph o f (a) high density polyethylene/AI2O3nanocomposite extruded from 60 pm A1203coated H D P E particles after 100 cycles and (b) HDPE/A1203 nanocomposite extruded from 16 pm A1203-coated H D P E particles after 75 cycles.

polymer particles. Micron-sized HDPE particles were coated with an ultrathin A1203film in a fluidized bed reactor by atomic layer deposition at a large scale. The FTIR and XPS revealed that A1203films were deposited on the polymer particle surface. TEM and FIB cross-sectional SEM revealed ultrathin and conformal A1203coatings. A nucleation mechanism for A1203atomic layer deposition on the polymer surface was confirmed. The results of ICP-AES suggested a nucleation period of 10 coating cycles, after which, a linear dependence between the film thickness and number of growth cycles was verified. The results of SEM, particle size distribution, and surface area of the uncoated and nanocoated particles showed that the particles were not coated as agglomerates during the coating process, rather as individual particles. A1203-coatedHDPE particles were successfully extruded into HDPE/AI2O3nanocomposite films by a heated extruder at controlled temperatures. Cross-sectional TEM indicated that nanoscale A1203 flakes were successfully dispersed in the polymer matrix. The dispersion of Alz03 flakes can be controlled by varying the polymer particle size. The process, firmly depositing nearly perfect nanometer thick ceramic films on polymer particle surfaces by atomic layer deposition and then extruding the coated polymer particles into final products, will provide unparalleled opportunities to produce quality nanocomposites with improved mechanical properties and reduced permeability in a continuous high throughput process at low cost.

Acknowledgement The authors thank Lyondell Chemical for providing the HDPE particles.

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”A. S. D. Sobrinho, G. Czeremuszkin, M. Latrkhe. G. Dennler. and M. R. Wertheimer, “A Study of Defects in Ultra-Thin Transparent Coatings on Polymers,” SurJ Cour. Tcchnol., 119, 1204-10 (1999). I4A. S. 0 . Sobrinho. G. Czeremuszkin, M. Latreche, and M. R. Wertheimer, “Defect-Permeation Correlation for Ultrathin Transparent Barrier Coatings on Pol mers,” J . Vuc. Sci. Technol., A, 18 [ I ] 149-57 (2000). “G. L. Graff, R. E. Williford. and P. E. Burrows. “Mechanisms of Vapor Permeation Through Multilayer Barrier Films: Lag Time Versus Equilibrium Permeation,” J. Appl. Phys., 96 [4] 184CL9 (2004). I6T. Suntola, “Atomic Layer Epitaxy.” Thin Solid Films, 216 [I] 8 4 9 (192). S. M. George, A. W. Ott, and J. W. Klaus, “Surface Chemistry for Atomic La er Growth.” J . Phys. Chem., 100 [31] 13121-31 (1996).

‘A. C. Dillon. A. W. Ott, J. D. Way, and S. M. George, “Surface Chemistry of A1203Deposition Using AI(CH3)3 and H 2 0 in a Binary Reaction Sequence,” Surf. Sci., 322 [ 1-31 2 3 M 2 (1995). I9A. W. Ott, J. W. Klaus, J. M. Johnson, and S. M. George, “Al2O3Thin Film Growth on Si( 100) Using Binary Reaction Sequence Chemistry,” Thin Solid Films, 292 [I-2] 13544 (1997).

wJ. R. Wank, S. M. George, and A. W. Weimer, “Nanocoating Individual Cohesive Boron Nitride Particles in a Fluidized Bed by ALD,” Powder Technol., 142 [I] 59-69 (2004). ”5. R. Wank, S. M. George. and A. W. Weimer. “Coating Fine Nickel Particles with A1203 Utilizing an Atomic Layer Deposition-Fluidized Bed Reactor (ALDFBR).” J. Am. Cerum. Soc., 87 [4] 762-5 (2004). 22 L. F. Hakim, S. M. George. and A. W. Weimer. “Conformal Nanocoating of Zirconia Nanoparticles by Atomic Layer Deposition in a Fluidized Bed Reactor,” Nanotechnology, 16 [7] S375-81 (2005). 23 L. F. Hakim, J. Blackson. S. M. George, and A. W. Weimer, “Nanocoating Individual Silica Nanoparticles by Atomic Layer Deposition in a Fluidized Bed Reactor,” Chem. Vupor Depos., 1 1 [lo] 420-5 (2005). 24G.D. Wilk, R. M. Wallace, and J. M. Anthony, “High-kappa Gate Dielectrics: Current Status and Materials Properties Considerations.” J. Appl. Phys., 89 [lo] 5243-75 (2001). *’J. D. Ferguson. A. W. Weimer, and S. M. George, “Atomic Layer Deposition of A120, Films on Polyethylene Particles.“ Chem. Muter.. 16 [26] 5602-9 (2004).

z6N. Kawakami. Y. Yokota, T. Tachibana. K. Hayashi. and K. Kobdshi, “Atomic Layer Deposition of A1203 Thin Films on Diamond,” Dium. Relur. Muter., 14 [II-I21 201S8 (2005). 27 A. Paranjpe, S. Gopinath, T. Omstead, and R. Bubber. “Atomic Layer Deposition of AIO, for Thin Film Head Gap Applications,” J . Electrochem. Soc., 148 [9] G465-71 (2001).

28 Y.-C. lung. H. Miura, K. Ohtani. and M. Ishida, “High-Quality Silicon/Insulator Hetermpitaxial Struclures Formed by Molecular Beam Epitdxy Using A1203 and Si.” J. Crysr. Crowrh. 1% [I] 88-96 (1999). 29 W. S. Yang, Y. K. Kim, S.-Y. Yang, J. H. Choi, H. S. Park, S. I. Lee, and J.-B. Yoo, “Effect of Si02 Intermediate Layer on AI2O3/SiO2/nt-Poly

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Si Interface Deposited Using Atomic Layer Deposition (ALD) for Deep Submicron Device Applications.” Surf Coal. Techno/., 131 [ 1-31 79-83 c2p’. M. D. Groner, F. H. Fabreguette, J. W. Elam, and S. M. George, ”LowTemperature A1203 Atomic Layer Deposition,” Chem. Marer., 16 [4] 6 3 9 4 (2004).

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”J. W. Klaus, 0. Sneh, and S. M. George, “Growth of Si02 at Room Temperature with the Use of Catalyzed Sequential Half-Reactions,” Science. 278 [53AS] 1934-6 (1997). C. A. Wilson, R. K. Grubbs, and S. M. George, “Nucleation and Growth During A1203 Atomic Layer Deposition on Polymers,” Chem. Muter., 17 [23] 562S34 (2005). 0

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NTRA-TYPE NANOCOMPOSITES CERAMIC MATERIALS

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STRENGTHENED

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TOUGHENED

Seong-min Choi, Sawao Honda, Shinobu Hashimoto, and Hideo Awaji Nagoya Institute of Technology Gokiso-cho, Showa-ku, Nagoya 466 8555, Japan Corresponding author : [email protected] ABSTRACT We proposed a new concept of strengthened and toughened mechanisms in ceramic matrix nanocomposites. It was clarified that intra-type nanocomposites has a useful structure to improve fracture strength and fracture toughness of ceramic materials. In this work, we used aalumina with average grain size of 260 pm and Moo3 with purity 99.9 % as starting materials. We mixed alumina powder and neutralized Moo3 solution with ethanol after dissolving Moo3 in 28 % ammonia solution. We fabricated alumina-molybdenum nanocomposites using a PECS method. The results showed that the maximum values of fracture strength and fracture toughness of the specimens sintered at 1450°C were 860 MPa and 5.22 MPa.m'/*, respectively. The sintered specimen carried out heat-treatment at 800"C, 900"C, and 1000°C for 0 to 10 min in Ar atmosphere. The fracture toughness increased to 5.67 MPa.m'/2 after annealing at 800°C for 10 min because dislocations diffused in the matrix. INTRODUCTION It is desirable for ceramics in high-temperature structural applications to have improved mechanical performance. Nanocomposites have been showing a possibility for increasing mechanical properties. Niihara and his group proposed several nano-structures, such as intra-type, inter-type, intrdinter-type, and nanohano-type, and reported many papers about nanocomposites'". Although, many research groups studied to improve the properties of ceramic rnaterial~~-~, the strengthening and toughening mechanisms of nanocomposites were obscure. Recently, we noticed that an intra-type nano-structure is improved mechanical properties and proposed a new concept of strengthening and toughening mechanisms in nanocomposites. In order to clarify both the strengthening and toughening mechanisms, our group reported several results in ceramic-metal and ceramic-ceramic nanocomposites*-". In this paper, we fabricated alumina-molybdenum(Mo) nanocomposites using a PECS (Pulse Electric Current Sintering) method and discussed the relationship between dislocations movement and improvement of mechanical properties.

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THEORY Toughening mechanism Figure 1 shows a schematic illustration of the toughening mechanism of nanocomposites. Dispersed dislocations within the matrix grains after annealing for alumina/molybdenum nanocomposites are described in this figure. Sub-grain boundaries are generated around the nano-size molybdenum particles and these dislocations become sessile dislocations at room temperature, shown in Fig. l(A). In this situation, when a tip of a propagating large crack reaches this area, these sessile dislocations in the matrix grains will operate as nano-crack nuclei in the vicinity of the propagating crack tip, shown in Fig. l(B). The high stress-state in the FPZ is then released by nano-crack nucleation, and the nano-cracks expand the FPZ size and consequently enhance the fracture toughness of the materials.

Fig. 1 Schematic description of the toughening mechanism in nanocomposites.

Strengthening mechanism Sintered alumina has residual stresses in the grains and grain boundaries because of anisotropic thermal expansion, Young’s modulus along the crystalline axes, and the crystallographic mis-orientation across the grain boundaries. In the sintered polycrystalline alumina, it is, therefore, conceivable that the large crack along a grain boundary created by synergetic effects of both residual stresses and processing defects will become equivalent to the grain size of the material and that the weakest crack generated along the boundary in the specimen dominates the strength of the specimen. Also, the fracture toughness of grain boundaries is usually lower than that of the matrix grains. Hence, polycrystalline alumina ceramics exhibit mainly intergranular fracture mode, as schematically shown in Fig. 2(A).

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Figure 3(A) shows the SEM (Scanning Electron Microscopy) observation of the fracture surface of monolithic alumina. Nanocomposites, however, will yield dislocations around the particles, and the dislocation creation releases residual stresses in the matrix grains. Consequently, the defect size along the grain boundaries reduces in nanocomposites, as shown in Fig. 2(B). Also, the dislocations in ceramics are difficult to move in ceramics at room temperature, serve as origins of small stress concentrations, and create nano-cracks around the propagating crack tip. These nano-cracks slightly reduce the strength of the alumina matrix and the reduction of both the residual stress along the grain boundaries and the strength in the matrix is attributable to a change in the fracture mode from intergranular fracture in monolithic alumina to transgranular fracture in nanocomposites. Also, the fracture surface of the transgranular mode of nanocomposites is not a simple planar cleavage plane. Several steps are frequently observed on the surface and this phenomenon is likely to be evidence of nano-cracking in the wake of the FPZ. Figure 3(B) shows the SEM micrograph of the fracture surface of alumina/copper(Cu) nanocomposites fabricated by us. Reduction of both the defect size along the grain boundaries and the tensile residual stresses in the matrix grains due to dislocation activities results in improvement of the strength of nanocomposites.

Fig. 2 Schematic description of the strengthening mechanism in nanocomposites.

Fig. 3 SEM micrographs of fracture surfaces of monolithic alumina (A) and alumindcopper nanocomposites (B).

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EXPERIMENTAL PROCEDURE Materials Preparation We used a-A1203 (TM-D, Taimei Chem. Japan ; a mean grain size 260 nm) and Moo3 (Mituwa Chem. Japan; purity 99.9%) as starting materials to fabricate Al2O3/Mo nanocomposites. We dissolved Moo3 powder in 28 % ammonia solution and neutralized this solution with ethanol. We mixed a-A1203 and neutralized solution. In this time, we controlled Mo content from 1 to 5 ~01%.The following Mo contents are apparent value. The mixtures are dried in evaporator at 70°C for 30 min and calcined at 550°C for 3 h. The calcined mixtures are reduced under high purity H2 gas atmosphere at 900°C for 2 h. We sintered the reduced mixtures to a cylinder shape in a graphite die (inner size : 20 mm) at 1250"C, 1350"C, 1450"C, and 1550°C with heating rate of 5O"C/min for 5 min under 30 MPa in vacuum atmosphere using a PECS method. The sintered specimens were cut into the sample dimension 2 x 2 ~ 1 0mm3 and polished with 3 pm diamond paste. Heat treatment and Characterization The sintered specimen was heat treated at 8OO"X to 1000°C for 0 to 10 min in Ar atmosphere. The sintered and annealed specimens were measured the physical and mechanical properties. Bulk densities of the specimens after sintering were measured by the Archimedes' method. The fracture strength was obtained by a three-point bending test with a cross head speed of 0.5 m d m i n and a span size of 8 mm for four specimens. The fracture toughness was measured using a single edged V-notch beam (SEVNB) method" with a notch depth of 1 mm for specimens with similar size of the fracture strength specimen. The SEVNB method uses a sharp V-shaped notch to estimate the intrinsic fracture toughness of the material. After a three-point bending test, the fracture surfaces were observed by SEM. RESULTS AND DISCUSSION Figure 4 shows XRD results of the a-A1203/Mo samples after (a) calcinations, (b) reduction, and (c) sintering. The calcined samples show a-Al203 and Moo3 peak. But, reduced and sintered samples show a-Al2O3 and Mo metal peaks. From this result, it is known that molybdenum oxide was transformed to molybdenum metal due to the reduction in H2 atmosphere and that we could

Fig. 4 XRD patterns of (a) calcined, (b) reduced, and (c) sintered specimens

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obtain a-Al2O3lMo nanocomposites.

Fig. 5 Relative density of monolithic A1203 and a-A1203/Monanocomposites

Fig. 6 Vickers’ hardness of monolithic A1203 and a-A1203/Mo nanocomposites

Figure 5 shows the relative densities measured by Archimedes’ method. In the case of monolithic alumina, we obtained dense alumina even at 1250°C. But, we could not obtain dense nanocomposites at 1250°C. From this result, it is known that nanocomposites should sintered above 1350 to fabricate dense aluminahlo nanocomposites. Figure 6 shows the Vickers’ hardness of monolithic alumina and nanocomposites. The Vickers’ hardness of monolithic alumina at 1250°C is slightly higher than that of nanocomposites. The Vickers’ hardness of monolithic alumina decreases with increasing sintering temperature because the grain size Fig. 7 Fracture strength of monolithic increased from 0.5 pm at 1250°C to 10 pm at A1203 and a-Al2O3lMo nanocomposites 1550°C. But, in the case of nanocomposites, the Vickers’ hardness shows higher than that of monolithic alumina above 1350°C because Mo

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particles dispersed in alumina matrix decreased residual stresses of alumina in grains. The maximum value of Vickers' hardness is 20 GPa in nanocomposites. Figure 7 shows the fracture strength of monolithic alumina and alumina/Mo nanocomposites measured by the three-point bending test. In the case of monolithic alumina, the fracture strength decreased with increased sintering temperature due to grain growth of alumina, and the value, 680 MPa at 1250°C is higher than that of nanocomposites. In contrast, the fracture strength of nanocomposites shows higher than that of monolithic alumina above 1350°C. The fracture strength of 1 vol% Mo nanocomposites increased to 860 MPa at 1350°C. But, this value decreased with increasing sintering temperature. In the case of 5 vol% Mo nanocomposites, maximum value is 860 MPa sintered at 1450°C and decreased at 1550°C. From these results, it is clear that the fracture strength of nanocomposites depends on the amount of Mo dispersed in the matrix.

Fig. 8 SEM microphotographs of fracture surface of 5 vol% Mo nanocomposites (a) 1250"C, (b) 135OoC,(c) 1450"C, and (d) 1550°C.

We observed microstructures of the fracture surface of 5 vol% Mo nanocomposites as shown in Fig. 8. The specimens sintered at 1250°C and 1350°C show inter-granular fracture mode. But, the fracture mode of nanocomposites above 1450°C was transformed from inter-granular mode to intra-granular mode. Although, the sample sintered at 1550°C shows intra-granular fracture mode, the fracture strength decreases comparing with the fracture strength of the specimen sintered at 1450°C because of grain growth. Figure 9 shows the fracture toughness measured by SEVNB method. The tendency of

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fracture toughness with sintering temperature is similar to the fracture strength. The maximum value of the fracture toughness is 5.22 MPa*m'/2in 5 vol% Mo nanocomposites sintered at 1450°C. From this result, it is know that the optimum sintering temperature to generate dislocation in sintered specimen is 1450°C because the specimen shows the maximum values of the fracture strength and fracture toughness.

Fig. 9 Fracture toughness of monolithic A1203and

Fig. 10 Fracture toughness of a-A1203/ 5 vol%

a-A1203/Monanocomposites

Mo nanocomposites sintered at 1450 "C after annealing

Figure 10 shows the fracture toughness of the specimen sintered at 1450°C after heat treatment at 800"C, 900"C, and 1000°C for 0 to 10 min in Ar atmosphere. We carried out heat treatment after sintering in order to diffuse dislocations in the matrix. From this result, the fracture toughness of the specimen annealed at 900"C, and 1000°C markedly decreased compared with that of the specimen before annealing. It is conceivable that the dislocations removed from the grains of alumina matrix. On the contrary, the fracture toughness of the specimen annealed at 800°C was higher than that of the specimen before annealing because the dislocations diffused in the matrix. The maximum value of the fracture toughness is 5.67 MPa*m'/*in the specimen annealed at 800°C for 10 min.

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CONCLUSIONS In this work, we used a-alumina and Moo3 as starting materials. We mixed alumina powder and neutralized Moo3 solution with ethanol after dissolving Moo3 in 28 % ammonia solution. We fabricated Al2O3/Mo nanocomposites using a PECS method. The maximum values of the fracture strength and fracture toughness were 860 MPa and 5.22 MPa.m”2 sintered at 1450”C, respectively. The specimen sintered at 1450°C was heat-treated at 800”C, 900”C, and 1000°C for 0 to 10 min. The fracture toughness increased to 5.67 MPa.m”2 in the specimen annealed at 800°C for 10 min. From this result, it is clarified that appropriate heat treatment after sintering is important to improve the fracture toughness. REFERENCE 1 K. Niihara, “New design concept of structural ceramics” J. Ceram. Soc. Japan, 99, 974-82 (1991). 2M. Nawa, K. Yamazaki, T. Sekino, and K. Niihara, “Microstructure and mechanical behaviour of 3Y-TZPMo nanocomposites processing a novel interpenetrated intragranu I ar microstructure” J. Muter: Sci. 3 1(1 l), 2849-58 (1996). 3 Y. Hayashi, Y. H. Choa, J. P. Singh, and K. Niihara, “Mechanical and electrical properties of ZnO/Ag nanocomposites” Ceramic Transactions, 96,209- 18 (1999). M. Sternitzke, “Review : Structural Ceramic Nanocomposites” J. Eur: Ceram. Soc., 17, 1061-1082 (1997). 5 G. Pezzotti, T. Nishida and M. Sakai, “Physical limitations of the inherent toughness and strength in ceramic-ceramic and ceramic-metal nanocomposites” J. Ceram. Soc. Jpn., 103 , 90 1-909 (1995). 6L.C. Steams, J. Zhao and M.P. Harmer, “Processing and microstructure development in A1203-SiC ‘nanocomposites”’ J. Eur: Ceram. Soc. 10,473-477 (1992). I W. L. E. Wong, S. Karthik, and M. Gupta, “Development of hybrid Mg/A1203 composites with improved properties using microwave assisted rapid sintering route” J. Muter. Sci. 40( 13), 3395-402 (2005). 8 H. Awaji, S-M. Choi, and E. Yagi, “Mechanisms of toughening and strengthening in ceramic-based nanocomposites” Mech. of Muter:, 34,4 11-22 (2002). 9 T. Matsunaga, U. Leela-adisorn, Y. Kobayasi, S-M. Choi, and H. Awaji, “Fabrication of alumina-based toughened nanocomposites” J. Ceram. Soc. Japan, 113, 123-25 (2005). 10 S-M. Choi and H. Awaji, “Nanocomposites-a new material design concept” Sci. and Tech. ofddv. Muter:, 6,2-10 (2005). “H. Awaji and Y. Sakaida, “V-notch technique for single-edge notched beam and chevron notch method”J. Am. Ceram. SOC.,73,3522-23 (1990).

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PREPARATION AND PROPERTIES OF MULLITE - BASED IRON MULTI - FUNCTIONAL NANOCOMPOSITES Hao WANG, Weimin WANG, Zhengyi FU, State Key Lab of Advanced Technology for Materials Synthesis and Processing, 122 Luoshi Road, Wuhan University of Technology, Wuhan, Hubei Province, P. R. China, 430070 Tohru SEKINO and Koichi NIIHARA Institute of Scientific and Industrial Research, Mihogaoka 8-1, Osaka University, Ibaraki, Osaka, Japan, 567-0047

ABSTRACT Mullite-based nanocomposite powders with embedded iron nanoparticles were synthesized by reduction of sol-gel prepared A15.4Fe0.6Si2013 solid solution in hydrogen. Structural characterization studied by TEM revealed that the intra-granular iron nanoparticles along with inter-type iron grains were obtained in mullite matrix. The magnetic properties of nanocomposite powders after reduction at 1200°C were measured by SQUID. It suggests that the intergranular and intragranular a-iron nanoparticles had the ferromagnetic and superparamagnetic behavior at room temperature, respectively. Dense Mullite-based iron composites were prepared from the nanocomposite powders reduced at different temperatures by hot pressing. The mechanical and magnetic properties were strongly affected by the reduction temperature. Flexural strength of mullite based iron nanocomposite reduced at 1200°C and successively hot pressed at 1650°C is one times higher than monothetic mullite. Furthermore, dense composites also behave ferromagnetic property at room temperature. INTRODUCTION Mullite (3A1203.2Si02) is widely researched for high-temperature structural, electronic, and optical applications ['I. Compared with its outstanding high temperature strength and creep resistance, mullite exhibits low fracture toughness and strength at room temperature compared with other engineering ceramics [21. To overcome this problem, some studies have been carried out by incorporating molybdenum particles in mullite matrix due to the similar thermal expansion coefficient values between them [3341. However, research on the influence of other metal inclusions on the properties of mullite composite is inadequate. Based on the novel concept of the ceramic-based nanocomposites proposed by Niihara [51, techniques were developed to disperse Ni, Co, Mo, W, etc., nanoparticles into alumina or zirconia by reduction of metal compounds in ceramics matrix [677381. The mechanical properties of those nanocomposites are strongly enhanced, and other fimctions such as ferromagnetism can be

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provided by the metal nanoparticles, which shows the possible application of ceramic nanocomposites with multifunction. In contrast, nanocomposite powders with dispersion of metal nanoparticles can be obtained by the reduction of homogenous metal oxides solid solution [9,101, which also shows the possibility to gain metal particles dispersed composites after sintering. IV 0 Mullite is a non-stoichiometric compound with the chemical formula of A1V'2[A12+2xSi2-2x] (0 I x I1) [I1]. Different amounts of transition metal cations mullite can be incorporated in mullite lattice by substitution of A13+or Si4+at different positions. In the case of Fe-doped mullite, 12 wt% Fez03 can be dissolved in mullite matrix under 1200°C by replacing the A13+ at octahedral site [I2]. In the present study, mullite-based iron nanocomposite powders were prepared from the reduction of A15.4Fe0.6Si2013solid solution at different temperatures. The dense composites were gained from the consolidation of nanocomposite powders by hot pressing at 1650°C. The formation, microstructure and magnetic properties of nanocomposite powders as well as the microstructure, mechanical and magnetic properties of the dense composites were investigated. EXPERIMENTAL PROCEDURE Sample preparation To synthesize A15.fie0.6Si2013 solid solution, aluminium nitrate (Al(N03)3.9H20; High Purity Chemicals, 99.9%) and iron nitrate (Fe(N03)3-9HzO;High Purity Chemicals, 99.9%) were used as the precursors for alumina and iron oxide respectively, while tetraethyl orthosilicate (TEOS, Si(OC2H5)4; Wako, 95%) as the precursor for silica. The synthesis process was given in a previous study in detail [I3]. Briefly, solid solution of A15.4Fe0.6Si2013was prepared from the diphasic gel by sol-gel method and calcinations. To obtain the iron dispersed mullite nanocomposite powders, solid solution powders were reduced by flowing hydrogen gas from 1100°C to 1400°C with an interval of 100°C for 1 h. The dense mullite-iron nanocomposite samples were prepared by reduction and hot pressing. Solid solution powders were dry ball milled for 24 h by alumina balls and screened through a No. 200 mesh sieve, then put into graphite dies. After hydrogen reduction from 1100°C to 1400°C with an interval of 100°C for 1 h with a heating rate of 20"C/min, the nanocomposite powders were successively consolidated at 1650°C under an applied load of 30 MPa in argon for 1 h by uniaxial hot pressing. The sintered bodies were cut and polished to the size of 3 x4 x40 mm for rectangular bars. Characterization The crystalline phases of composite powders before and after reduction were recorded between 10" to 90" at 4" 2Wmin by X-ray diffraction analysis (XRD; RU-200B, Rigaku Co. Ltd., Japan) using Cu Ka radiation. The powder XRD patterns for calculation of lattice parameters were measured using 0.5" 2Wmin with pure silicon as an internal standard. The microstructure of nanocomposite powders was observed by transmission electron microscope (TEM; H-8 1OOT,

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Hitachi Co. Ltd., Japan). To investigate the microstructure configuration of dense composites, polished surfaces thermally etched at 1450°C for 1 h in argon atmosphere were observed using scanning electron microscope (FE-SEM; S-5000, Hitachi Co. Ltd., Japan). Vickers indentation, with a load of 49N for 15s, was used to determine the fracture toughness and hardness of composites. More than five independent measurements were performed for each experimental point. Young's modulus was tested using the resonance vibration technique. Flexural strength was measured by the three-point bending test with a span of 30 mm at a crossed speed of 0.5 m d m i n at room temperature. At least five bars were tested for one experimental point. Magnetization measurements were conducted by means of a Quantum Design MPMS superconducting quantum interference device (SQUID) magnetometer. RESULTS AND DISCUSSION Nanocomposite Powders Figure 1 shows the XRD profiles of samples before and after reduced at 1100-1400°C by hydrogen flow. Besides a small quantity of silica in some samples, the mullite and a-Fe are the only phases detected in all the samples. Because of the different stability of metal cations under reduction atmosphere in the solid solution, only Fe3+ can be reduced from the lattice of solid solution. The separation of iron phase from the lattice of mullite-iron oxide solid solution results in the decrease of lattice parameters and volume of unit cell. The lattice parameters and volume of unit cell for the samples reduced at different temperature are summarized in Table 1. The volume of unit cell sharply shrinks with the raise of reduction temperature till 1200°C and then keeps in stabilization. It reveals that the Fe3+cation is completely transformed to iron phase when the reduction temperature exceeds 1200°C.

Figure 1. XRD profiles of samples before and after reduction in hydrogen flow at various temperatures.

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Table 1. Summary of lattice parameters for the samples reduced at different temperatures in hydrogen flow. Volume of unit Reduction a (4 b (4 c (4 cell (A3) temperature ("C) Before reduction

7.5757h0.0035

7.7177k0.0044

2.9021h0.0019

169.678

1100 1200 1300 1400

7.5437k0.0011 7.5439h0.0013 7.5413h0.00 14 7.5419h0.00 12

7.6926h0.0013 7.6898k0.0012 7.6909*0.00 13 7.6909+0.0013

2.8847*0.0004 2.8837h0.0004 2.8842*0.0005 2.8840*0.0005

167.402 167.287 167.283 167.284

The TEM images of reduced powders are shown in Figure 2. Two kinds of a-Fe nanoparticles with different particle size were observed. Some a-Fe grains with a size larger than 200nm are formed by the aggregation and growth of the iron nanoparticles located odnear the surface of the solid solution grain. a-Fe nanoparticles embedded in the mullite grains were also observed with a particle size less than 10nm. The formation of intra-granular nanoparticles is due to the precipitation of iron nanoparticles inside parent grains by changing the chemical state of iron through reduction. Mullite grain acts as a matrix to prevent the nanoparticles from further aggregation and growth.

Figure 2. The TEM images of reduced powders The magnetic hysteresis loops for mullite-iron nanocomposite powders reduced at 1200°C measured at temperature of 10 K (a) and 300 K (b) is shown in Figure 3. Saturated magnetizations (Ms), remanence ratio (Mr/Ms) and coercivity (Hc) of nanocomposite powders are 15.87 emdg, 1.9%, 380e at 300 K and 16.86 emdg, 3.0%, 450e at 10 K, respectively. It is known that for a superparamagnetic particle system, both the coercive field and remnant magnetization increase with decreasing temperature below the superparamagnetic-ferromagnetic transition temperature. And this increase is due to the lower thermal activation energy of spins at low temperature. In our results, the variation of coercivity and remanence ratio with temperature implies the existence of superparamagnetic iron nanoparticles in mullite-iron nanocomposite

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powders. As known, the magnetic behavior of metal nanoparticles is tightly related with their particle size. In the case of iron nanoparticles, the critical size of superparamagnetism is 14 nm [I4]. From the observation of TEM, most of the intragranular iron nanoparticles are less than the critical size, which show the superparamagnetic behavior at room temperature. On the other hand, intergranular iron grains, with the size of hundreds of nanometers, are in the state of ferromagnetism.

Figure 3. The hysteresis loop for A15.4Fe0.6Si2013 powders reduced at 1200°C. Insert shows enlargement of the plot near the origin ((a) at 10 K, (b) at 300 K). Microstructural Characteristics and of Properties Sintered Bodies The thermally etched surfaces of hot pressed samples and monolithic mullite are presented in Figure 4. Samples show the similar microstructure of some large acicular and rectangular grains dispersed in a matrix of fine equiaxed grains. Klug et. al. [14] reported that the single-phase mullite boundaries changed with temperature above -1 600°C and joined at -1 890°C at a composition 77 wt% A1203. Therefore, the single-phase mullite prepared from diphasic gel precursors will enter the mullite + liquid region of the phase diagram when heated between -1650-1 850°C. After the congregation and growth of embedded iron nanoparticles and iron grains through diffusion at high temperature sintering, the hybrid composite of mullite and iron was formed with iron particles locating in the triple join junctions (Figure 5).

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Fig. 4.SEM photograph for thermally etched surfaces of hot pressed samples. (a: mullite; b: sample reduced at 1200 "C; c: sample reduced at 1300"C)

Fig. 5. Location of iron grain in dense composite The mechanical properties of the studied specimens are given in Table 2. Flexural strength primarily rises to 400 MPa in the sample reduced at 1200°C and then sharply declines with the increase of reduction temperature, while the fracture toughness shows the maximum value in the sample reduced at 1300°C. Due to the incompletely transformation of Fe3+to iron in the sample reduced at llOO°C, the tsf is smaller than that of the sample reduced at 1200°C. When the

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reduction temperature exceeds 1200”C, all of the Fe3+ in solid solution is completely reduced. The iron phase is in the same content in composites. However, the grain size of iron phase is obviously related with reduction temperature. The iron grains grow large as the raise of reduction temperature. That is, the larger the iron grain size, the larger the critical flaw of composites. Thus the flexural strength declines as the raise of reduction temperature higher than 1200°C. The same result was reported in the research of mullite-Mo composites [31. By the contribution of iron inclusion in composites, the toughness of composites slightly increases compared with monothetic mullite. The decline of Young’s modulus and Vickers hardness are obviously resulted from the effect of soft metal phase. Table 2. Mechanical and magnetic properties of Mullite-iron composites Flexural Indentation Vickers Young’s strength toughness hardness modulus IMP4 [MPa*m”2] P a l [MPaI Mullite 207 f 22 1.86 f 0.06 12.30 226 Reduced at 1100°C 295 f 11 2.05 f 0.11 10.15 212 Reduced at 1200°C 405 f 54 2.40 f 0.13 10.82 213 Reduced at 1300°C 255 f 18 2.03 f 0.1 I 11.40 203 Reduced at 1400°C 2.13 f 0.22 204 f 14 11.34 21 1 The room temperature magnetic behavior of composite reduced at 1200°C is presented in Figure 6. Composite exhibits the typical ferromagnetism by the formation of hysteresis loop. Figure 7 shows the saturated magnetization and coercive force versus reduction temperature. The saturated magnetization of composites increased with reduction temperature till 1200°C. As the result of aggregation and growth of iron grain during sintering, the coercive force is around 200e.

Figure 6. The hysteresis loop for A15.4Fe0.6Si2013powders reduced at 1200°C and hot pressed at 1650°C. Insert shows enlargement of the plot near the origin.

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1i00

1200

Reduction temperature VC)

Figure 7. The saturated magnetizations and coercive forces of dense samples reduced at different temperature. CONCLUSION Dense mullite-based iron nanocomposites have been prepared through the reduction of A15.4Fe0.6Si2013 powders and successive hot pressing. Nanocomposite powders after reduction have an intra-type structure with the iron nanoparticles smaller than 10 nm embedded in mullite grains along with large iron grains (2 200 nm). The mechanical and magnetic properties of dense composites are strongly affected by the reduction temperature. Flexural strength of mullite-based iron nanocomposite reduced at 120OOC and successively hot pressed at 165OOC is one times higher than monothetic mullite. Furthermore, dense composites also behave ferromagnetic property in room temperature. ACKNOWLEDGEMENT H. Wang thanks the National Natural Science of China (No. 50102003) in part for financial support. REFERENCES ‘I. A. Aksay, D. M. Dabbs and M. Sarikaya, “Mullite Processing, Structure, and Properies”, J. Am. Ceram. Soc., 74,2343- 58 (1991). 2T. -I. Mah and K. S. Mazdiyasni, “Mechanical Properties of Mullite”, J. Am. Ceram. SOC.,66,699-703 (1983). 3J. F. Bartolome$ M. Diaz, J. Requena, J. S. Moya and A. P. Tomsia, “Mullite/ Molybdenum Ceramic- Metal Composites”, Acta. Mater., 47, 389 1-99 (1999). 4 J.‘ F. Bartolomei, M. Diaz and J. S. Moya, “Influence of the Metal Particle Size on the Crack Growth Resistance in Mullite-Molybdenum Composites”, J. Am. Ceram. Soc., 85, 2778-

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84 (2002). ’K. Niihara, “ New Design Concept for Structural Ceramics: Ceramic Nanocomposites”, J. Cerum. SOC.Jpn., 99,974-82 (199 1). 6 T. Sekino, T. Nakajima, S. Ueda, K. Niihara, “Reduction and Sintering of a Nickel-dispersed-alumina Composite and Its Properties”, J. Am. Cerum. Soc., 80, 1139-48 (1997). 7T. Sekino, K. Niihara, “Fabrication and Mechanical Properties of Fine-tungsten-dispersed Alumina-based Composites”, J. Muter: Sci., 32, 3943-49 (1997). *H. Kondo, T. Sekino, N. Tanaka, T. Nakayama, T. Kusunose, K. Niihara, “Mechanical and Magnetic Properties of Novel Yttria-stabilized Tetragonal ZirconiaN Nanocomposite Prepared by the Modified Internal Reduction Method”, J. Am. Cerum. Soc., 88, 1468- 73 (2005). 9V. Carles, Ch. Laurent, M. Brieu, A. Rousset, “Synthesis and characterization of Fe/Co/Ni alloys-MgO nanocomposite powders”, J. Muter. Chem., 9, 1003-9 (1999). “0. QuCnard, E. D. Grave, Ch. Laurent, A. Rousset, “ Synthesis, Charaterization and Thermal Behavior of Fe0.65Co0.35-MgA1204 and Fe0.65Ni0.35-MgA1204Nanocomposite Powders”, J. Muter. Chem., 7,2457- 67 (1997). l 1 W. E. Cameron, “Mullite: a Substituted Alumina”, Am. Min., 45, 157-61 (1977). 12 M. Ocafia, A. Caballero, T. Gonzalez-Carrefio, C.J. Serna, “Preparation by Pyrolysis of Aerosols and Structural Characterization of Fe-doped Mullite Powders”, Muter. Rex Bull., 35, 775-88 (2000). I3H. Wang, T. Sekino and K. Niihara, “Magnetic Mullite-iron Composite Nanoparticles Prepared by Solid Solution Reduction”, Chem. Lett., 34,298-99 (2005). l 4 F. J. Klug, S. Prochazka and R. H. Doremus, “Alumina-silica phase diagram in the mullite region.”, J. Am. Cerum. Soc., 70, 750- 59(1987). I5J. F. Loffler, J. P. Meier, B. Doudin, J. Ansermet and W. Wagner, “Random and exchange anisotropy in consolidated nanostructured Fe and Ni: Role of grain size and trace oxides on the magnetic properties”, Phys. Rev. B, 57,29 15-24 (1998).

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ELECTROSPINNING OF CERAMIC NANOFIBERS AND NANOFIBER COMPOSITES Junhan Yuh, Hyun Park and Wolfgang M. Sigmund*

Department of Materials Science & Engineering, University of Florida, 225 Rhines Hall, P.O. Box 116400, Gainesville, FL 3261 1, USA ABSTRACT Lanthanum cuprate nanofibers and barium titanate nanofibers were synthesized by electrospinning utilizing sol-gel precursors. The effects of electrospinning parameters (ceramic precursor concentration and electric voltage) on fiber diameter before/after calcinations were investigated in detail for lanthanum cuprate nanofibers. A series of experiments designed by response surface methodology shows as dried fiber diameter increases calcined fiber diameter increases. As voltage decreases and ceramic precursor concentration increases, dried fiber diameter increases for most combinations of voltage and precursor concentration. But there are also conditions available where the opposite is true. This paper explains how contradicting reports on fiber diameter dependence in the literature can be unified in a larger picture. Furthermore, the microstructure evolution in the nanofiber was studied. The morphology, microstructure and crystal structure were investigated by SEM, TEM and XRD. INTRODUCTION Over the last few decades, research and development of one dimensional nano materials such as nanotubes and nanofibers have been of great interest due to their unique structure and properties, i.e. high aspect ratio, large specific surface area and chemical / mechanical stabilities that can be utilized for diverse applications in the field of nano devices. Nanomaterials can be used as building blocks in nanotechnology.[[ 1, 211 Currently, several ceramic nanowires have been developed by various techniques, e.g. template growth, laser ablation and chemical vapor deposition (CVD)[2-51. But a major disadvantage of these methods is the need of multiple processing steps to produce nanowires. Etching for the template removal and purification processes are good examples of such extra processing steps required as they are both time and cost consuming processes. [ 1,6] A few years ago, research on electrospinning of ceramic fibers started since this process provides direct access to nanostructured ceramic nanofibers and requires less steps compared to other methods. In electrospinning of ceramic nanofibers, a high electric field is applied to a metallic syringe needle containing sol-gel precursors mixed with dissolved polymer. Under the influence of the electric field, a droplet of the polymer solution at the needle tip is deformed into a conical shape. If the electrostatic repulsive force overcomes the surface tension, a thin charged jet is ejected. The jet moves towards a grounded plate. Due to the presence of polymer entanglements, the jet remains stable and does not transform into spherical droplets. As the solvent evaporates, this jet is stretched to many times its original length to produce continuous, ultrathin fibers of polymer. Recently, binary ceramic nanofibers such as titania, vanadium oxide, niobium oxide, silica and copper oxide have been produced via electrospinning, [7-141 and a few ternary ceramic nanofibers were fabricated.[ 15, 161 Electrospun nanofibers come naturally with a high surface to volume ratio and are interesting for applications in catalysis and as electrodes in solid oxide fuel cells (SOFCs) with lanthanum cuprate being one of the parent materials for its *To whom correspondence should be addressed: Tel: 1-352-846-3338, Fax: 1-352-846-3355, e-mail: [email protected] Nanocornpositesand Nanostructures

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cathode.[ 171 SOFCS offer the advantage of high energy conversion efficiency combined with low pollution for power generation. Lanthanum cuprate that is doped with divalent ions such as strontium, forms an oxygen deficient perovskite structure.[ 181 This yields a perovskite oxide that is electrically conductive, which made it a promising candidate for the electrode material for SOFCs [19] and gas sensors.[20] We will compare the results found for electrospun lanthanum cuprates with another electrospun perovskite, i.e. barium titanate. This ferroelectric material is widely used in various electronic applications and with the current miniaturization trend; development of nanoscale materials with ferroelectricity is crucial. Since electrospinning is a novel technique for ceramics only few reports have been published to now. Therefore, we present a systematic study of the ceramic precursor concentration on the fiber diameter. We also report on the electrospinning synthesis of two ternary perovskite oxide nanofibers, lanthanum cuprate and barium titanate. Results will be presented in empirical models that describe effects of variables (ceramic precursor concentration, electric voltage) on fiber diameters. Parameter optimization and derivation of empirical models was based on response surface methodology.[2 1,221 EXPERIMENTAL The chemicals used for the electrospinning solution were lanthanum nitrate hydrate and copper nitrate, and polyacrylamide (M,=lO,OOO, 50 wt.% solution in water) for lanthanum cuprate. Barium titanate precursors were chosen as barium acetate, titanium isopropoxide and poly(viny1 pyrrolidone) (PVP, M,= 1,300,000). All chemicals were used as received. Electrospinning solutions were prepared by mixing lanthanum nitrate hydrate, copper nitrate and polyacrylamide with stirring for 24 hours. Ceramic precursor concentrations of lanthanum cuprate were determined by experimental design (Table 1). The solutions were taken into 5 ml syringes and pumped out of the syringe using a syringe pump (BSP 99, Braintree Scientific Inc., Braintree, USA). The distance between the metal needle and collector (grounded aluminum foil) was 5cm. A voltage according to the experimental design (Table 1) was applied to the metal needle and fibers were collected on an aluminum foil. Fibers were dried for 12 hours at 1OO"C, followed by calcination for 10 hours at 600°C. For the preparation of barium titanate via sol-gel processing, barium acetate and titanium isopropoxide were dissolved in acetic acid at a 1:l molar ratio. After dissolution of barium acetate and titanium isopropoxide, PVP as solution in ethanol (PVP 0.2g and ethanol 3mL) was added and everything stirred at room temperature for 2 h. This solution was loaded into the syringe. Electrospinning was conducted by applying an electric field of 2,143 V to the precursor mixture. As-synthesized non-woven mats of nanofibers were dried in a vacuum oven at 80°C for an hour. For crystallization an annealing process with holding temperatures of 65OoC and 75OoC for 60 minutes was applied. Morphologies and dimensions of fibers were investigated using field emission scanning electron microscopy (JSM-6335F, JEOL, Peabody, USA) with an acceleration voltage of 15kV. X-ray diffraction (APD 3720, Philips, New York, USA) was used to characterize the phases. Transmission electron microscopy operating at 200kV was utilized for microstructural analysis. Solutions were characterized for their surface tension by the Wilhelmy plate method and solution viscosities were measured with a rheometer (MCR 300, Paar Physica USA, Ashland, USA). Measurements were done at room temperature. The solution was pre-sheared at 800s-' for one minute, then, kept stationary for one minute to equilibrate and then, the shear rate was

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ramped from O.OOls-' to 1100s-'. The time duration between two measurements was 15 seconds. Electrical conductivities were measured with a conductivity meter (AB30, Fisher Sci. Fairlawn, USA). A series of experiments was designed using two variables (ceramic precursor concentration, electric voltage). Experimental conditions were determined according to Doptimal design. This design selected design points that minimize the variance of the model regression coefficients. To derive quadratic equations, at least 3 levels of experimental conditions for each factor were investigated. Designed experimental conditions are given in Table 1.

12

0.055

10

RESULTS 1. Lanthanum cuprate 1.1 Response surface methodology for dried fiber diameter Fig. 1 shows scanning electron microscopy (SEM) images of dried fibers and heat-treated fibers. They are typical for electrospun ceramic nanofibers. As synthesized dried fibers are nanocomposites of the sol-gel ceramic and the polymer. On heating, the polymer burns out and the microstructure of the nanofibers develops. The fiber ceramic microstructure can be controlled by the annealing process, but not the fiber diameter. For fiber diameter quantification at least 50 fiber diameters were measured on several SEM pictures by randomly picking fibers and measuring their diameter vertically to their length.

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Fig 1: SEM pictures of (a) dry as-synthesized lanthanum cuprate/polyacrylic acid composite nanofibers, (b) lanthanum cuprate nanofibers after annealing (600°C for 1h) Measurement results for each experimental condition are summarized in Table 2. Validities F of the model and regression coefficients are summarized in Table 3 and Table 4, respectively.

where Ddried fiber is the dried fiber diameter in nm, V is the coded variable for voltage with

v = (Uapp[ie&l7 . 5 ) /

7.5

where Uapplied is the applied potential in kV. C is the coded variable for the ceramic precursor concentration defined as C=Conc.-0.055/0.045 with Conc. using molar units. It follows from equation (1) that the ceramic precursor concentration is more important to determine the dried fiber diameter than Uappjied. Furthermore it clearly states that as uapplied decreases and ceramic precursor concentration increases, fiber diameter increases. Additionally, there is an interaction between voltage and ceramic precursor concentration, which is more effective than voltage alone. This can be seen in Fig. 2 where a contour plot of the dried fiber diameter as a function of voltage and ceramic precursor concentration is presented. Two main areas can be seen. In one area where 0.01 M < Conc. < 0.03 M Ddriedfiber increases with increasing Uapplied up to 2 1 kV .[111

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Table 2: Coded Variables C and V for conc. and voltage, respectively, with the resulting fiber diameters for dried and calcined lanthanum cuprate nanofibers.

Fig. 2: Contour plot of dried fiber diameter as a hnction of ceramic precursor concentration and voltage for lanthanum cuprate nanofibers.

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1.2 Response surface methodology for calcined fiber diameter For calcined fibers the same methodology was applied as for dried fibers. Measurement results of each condition are also summarized in Table 2. A SEM image of calcined lanthanum cuprate fibers is shown in Fig. l(b). An analysis was done to confirm the model. The regression coefficients and the validity F for each model are summarized in Table 3 and Table 4, respectively. Table 3 : Validity F of the models for lanthanum cuprate nanofibers

Dried fiber model Calcined fiber model

73.21

0.948

0.934

12.63

0.76

0.699

Table 4:Regression coefficients of each model for lanthanum cuprate nanofibers

Variable

Dried fiber model

Calcined fiber model

Uapplied

-32.29

-19.32

Conc.

78.3

16.63

V'

-22.52

9.29

52.94

16.3

-35.13

-20.9 1

cx v

For calcined fibers the following equation (3) is found: Dcalcined

(nm) = 70.95- 19.32xV

+ 16.63K + 9.29xV' + 16.30K'-

20.91 XVXc

(3)

where Dcalcined is the calcined fiber diameter in nm. C is the coded variable for the ceramic precursor concentration defined as C=Conc.-0.055/0.045 with Conc. using molar units, and V is the coded variable for voltage as mentioned earlier. Equation (3) has similarities to equation (1) and some differences. It is similar in that the ceramic precursor concentration is more important than voltage to determine the calcined fiber diameter, and as voltage decreases and ceramic precursor concentration increases, fiber diameter increases. There is also an interaction between voltage and ceramic precursor concentration, which is more important. The difference, however, is found in that there is only a very small

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region where calcined fiber diameter increases with increasing voltage (25 kV and 0.01 M). The major difference is a new saddle point. It is displayed in the contour plot of the calcined fiber diameter as a function of voltage and ceramic precursor concentration in Fig, 3. This saddle point can be used to minimize the calcined fiber diameter.

Fig. 3: Contour plot of calcined fiber diameter as a function of ceramic precursor concentration and voltage for lanthanum cuprate nanofibers 2. Barium titanate The second electrospun ceramic nanofiber material that is presented is a ferroelectric. Here, we focus on the microstructure evolution in the nanofiber with heat treatment. The X-ray diffraction (XRD) patterns of electrospun barium titanate nanofibers with heat treatment at different temperature are shown in Fig. 4(a). The XRD pattern of as-synthesized and dried at 80°C sample shows their amorphous nature. After heat treatment at 65OoC the onset of perovskite BaTi03 is evident, although there are some non-perovskite peaks (-27O, 20). Non-perovskite peaks were due to the complex structural chemistry of barium titanate and at this temperature, 65OoC, perovskite and non perovskite peaks are coexist. XRD pattern obtained from fibers heat treated at 75OoC shows well-defined, high intensity perovskite barium titanate peaks with no detectable non-perovskite phases. The increment of the intensity of the peaks corresponding to temperature increment indicates that amount of the tetragonal barium titanate phase was increased.

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Fig. 4: (a) XRD patterns of barium titante nanofibers heat treated at various conditions: 80"C, 650°C and 750°C (b) SEM image of barium titanate nanofibers after heat treatment at 750°C for 1 h. For comparison reasons to the study on fiber diameter of lanthanum cuprates, SEM images of electrospun barium titanate after heat treatment at 750°C were done and quantitatively analyzed (see Fig. 4(b)). Fiber diameters were measured for 50 fibers of both dry as-synthesized and annealed (750°C) nanofibers. Average fiber diameter before heat treatment was found to be 273 nm and after heat treatment at 750°C the average diameter was reduced to 116 nm. This is equivalent to 60% of diameter reduction upon annealing. The reduction can mainly be attributed to polymer burn out during the heat treatment. The shrinkage in fiber diameter does not directly correspond to the above presented equations (1) and (3) since a different polymer has been used with higher molecular weight thus influencing surface tension and viscosity. Furthermore, the different precursor concentrations as well as the ionic constituents for barium titanate also alter the bulk nanofiber conductivity. This demonstrates that the models for one ceramic system cannot easily be transferred to another ceramic material. During annealing the material crystallizes and for BaTi03 equiaxed grains of 25-50 nm diameter become visible. They were verified to be perovskite barium titanate by transmission electron microscopy (TEM) (see figure 5 ) and selected area diffraction, which is published elsewhere.[29] The upper right inset is a high resolution image of an individual grain, which shows the lattice fringes. This confirms the crystallinity of the individual grains.

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Fig. 5: TEM image of barium titanate nanofiber annealed at 750°C for 1 h. Right upper inset: high resolution image of barium titanate nanofiber showing lattice fringes with scale bar of 2 nm.

DISCUSSION

1. Process optimization For applications in nanotechnology the knowledge for nanofiber minimization is crucial. For the investigated system based on lanthanum cuprate a minimum for calcined fibers was found. When the highest voltage (25kV) is paired with a medium ceramic precursor concentration (0.07M) the thinnest calcined nanofibers can be fabricated yielding dried fibers with Ddried = 152 nm, and calcined fibers with diameters of 6 1 nm. 2. Effects of voltage Up to this paper, the influence of electrospinning voltage on fiber diameter had not been clearly established.[23] A previous study showed that fiber diameter decreased with increasing voltage [24,25] while others showed the opposite trend[26]. The contour map presented in Fig. 2 shows and explains both trends simultaneously. It is obvious that both observations are true. At a low voltage, the fiber diameter increases with increasing voltage. The relationship reverses at 17 kV, now the fiber diameter decreases with increasing voltage. 3. Effects of ceramic precursor concentration As seen in the presented contour a plot, there is an impact of the ceramic precursor concentration on the final nanofiber diameter. These differences can be explained by a recently introduced mode1.[27] Table 5 lists the main parameters that control fiber diameter from a fundamental physics point of view. It is found that as the ceramic precursor concentration increases conductivity increases due to the increase in ionic strength. Most groups reported that increasing conductivity decreases the fiber diameter. They reported that salts such as sodium chloride and lithium chloride were used to increase the conductivity. These salts ionize when they are added to electrospinning solution and increase the conductivity through the bulk of the

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fiber. In our systems the conductivity of the ceramic precursor solutions is typically 3 orders of magnitude higher than for pure polymer solutions. [ 11, 23, 281 Thus, fiber diameter increases with increasing ceramic precursor concentration for most conditions. Table 5: Viscosity, conductivity and surface tension data as a function of ceramic precursor concentration for lanthanum cuprate.

Conc. (M)

I

0.01 0.055

Viscosity (Pa s)

I

0.1

0.14 0.13 0.12

Surface tension (mN/m)

Conductivity (mS/cm)

I

11.69 12.09 18.96

I

~

76.24 77.2

I

78.96

SUMMARY Lanthanum cuprate nanofibers and barium titanate nanofibers were synthesized by electrospinning. A series of experiments designed by response surface methodology shows that the ceramic precursor concentration is a more important factor than voltage to determine the response (fiber diameter). As voltage decreases and ceramic precursor concentration increases, the dried fiber diameter is found to increase. Calcined fiber diameters show the same trend. Using the derived equations, the processes were optimized to produce dried fiber diameters of 152 nm, and calcined fiber with diameters of 6 1 nm at the highest voltage (25 kV) and medium ceramic precursor concentration (0.07 M). BaTi03 fibers were found to form perovskite 150°C above the standard reported temperature for sol-gel derived BaTi03. Electrospun barium titanate was still amorphous at 600°C which might be caused by the strong electric field. Since the crystallographic study of nanofibers confirmed tetragonal perovskite barium titanate after heat treatment at 75OoC, ferroelectric behavior can be expected. To confirm ferroelectricity of electrospun barium titanate nanofibers, scanning probe microscopy (SPM) and electrostatic force microscopy (EFM) investigations are ongoing.

ACKNOWLEDGEMENTS The authors acknowledge funding from the Department of Energy under grant no. DEFG26-03NT4 1614.

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REFERENCES 1. 2. 3. 4. 5. 6. 7. 8.

9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19.

Y. Dzenis, “Spinning continuous fibers for nanotechnology,” Science, 304(5679), 19171919 (2004). C.N.R. Rao, A. Govindaraj, G. Gundiah, and S.R.C. Vivekchand, “Nanotubes and nanowires,” Chem. Eng. Science, 59(22-23) 4665-467 1 (2004). S. Iijima, “Helical Microtubules of Graphitic Carbon,” Nature, 354(6348), 56-58 (199 1). A.M. Morales, and C.M. Lieber, “A laser ablation method for the synthesis of crystalline semiconductor nanowires,” Science, 279(5348), 208-21 1 (1998). C.R. Martin, “Nanomaterials - a Membrane-Based Synthetic Approach,” Science, 266(5193), 1961-1966 (1994). D. Li, Y.L. Wang, and Y.N. Xia, “Electrospinning of polymeric and ceramic nanofibers as uniaxially aligned arrays,” Nano Lett., 3(8), 1167-1171 (2003). K.H. Lee,H.Y. Kim, M.S. Khil, Y .M. Ra, and D.R. Lee, “Characterization of nanostructured poly(epsi1on-caprolactone) nonwoven mats via electrospinning,” Polymer, 44(4), 1287-1294 (2003). S.A. Theron, A.L. Yarin, E. Zussman, and E. Kroll, “Multiple jets in electrospinning: experiment and modeling,” Polymer, 46(9), 2889-2899 (2005). A.L. Yarin, and E. Zussman, “Upward needleless electrospinning of multiple nanofibers,” Polymer, 45(9), 2977-2980 (2004). A. Frenot, and I.S. Chronakis, “Polymer nanofibers assembled by electrospinning,” Curr. Opin. CoZZ. Inter. Sci.,8( l), 64-75 (2003). Z.M. Huang, Y.Z. Zhang, M. Kotaki, and S. Ramakrishna, “A review on polymer nanofibers by electrospinning and their applications in nanocomposites,” Comp. Sci. Tech., 63( 19,2223-2253 (2003). J.M. Deitzel, J. Kleinmeyer, D. Harris, and N.C.B. Tan, “The effect of processing variables on the morphology of electrospun nanofibers and textiles,” Polymer, 42( 1) 26 1272 (2001). M.M. Demir, I. Yilgor, E. Yilgor and B. Erman, “Electrospinning of polyurethane fibers,” Polymer, 43(1 l), 3303-3309 (2002). X.H. Zong, K. Kim, D.F. Fang, S.F. Ran, B.S. Hsiao, and B. Chu, “Structure and process relationship of electrospun bioabsorbable nanofiber membranes,” Polymer, 43( 16), 44034412 (2002). H.Y. Guan, C.L Shao, Y.C. Liu, N. Yu, and X.H. Yang, “Fabrication of NiC0204 nanofibers by electrospinning,” Solid State Commun., 131(2), 107-109 (2004). D. Li, T. Herricks, and Y.N. Xia, “Magnetic nanofibers of nickel ferrite prepared by electrospinning,” Appl. Phys. Lett., 83(22), 4586-4588 (2003). N. Dharmaraj, H.C. Park, D.M. Li, P. Viswanathamurthi, H.Y. Kim, and D.R. Lee, “Preparation and morphology of magnesium titanate nanofibres via electrospinning,” Inorg. Chem. Comm., 7(3), 43 1-433 (2004). H.C. Yu, and K.Z. Fung, “La~-,SrCuO~.~-delta as new cathode materials for intermediate temperature solid oxide he1 cells,” Mat. Res. Bull., 38(2), 23 1-239 (2003). H.C. Yu, and K.Z. Fung, “Electrode properties of Lal-,SrxCuO~.s-deltaas new cathode materials for intermediate-temperature SOFCs.,” J. Power Sources, 133(2), 162-168 (2004).

Nanocornposites and Nanostructures

453

20. 21. 22. 23. 24. 25.

26. 27. 28. 29.

454

X.H. Zhou, Q.X. Cao, Y. Hu, J.X. Gao, and Y.L. Xu, “Sensing behavior and mechanism of La2Cu04-SnO2 gas sensors,” Sen. Actuators B., 77( 1-2), 443-446 (2001). D.C. Montgomery, “Design and analysis of experiments,” 5th ed., New York: John Wiley. xii, 684 (2001). R.H. Myers, and D.C. Montgomery, “Response surface methodology : process and product in optimization using designed experiments,” Wiley series in probability and statistics. Applied probability and statistics, New York: Wiley. xiv, 700 (1995). D. Li, and Y.N. Xia, “Electrospinning of nanofibers: Reinventing the wheel?,” Adv. Materials, 16(14), 1151-1170 (2004). Y.M. Shin, M.M. Hohman, M.P. Brenner and G.C. Rutledge, “Experimental characterization of electrospinning: the electrically forced jet and instabilities,” Polymer, 42(25), 9955-9967 (2001). J.S. Lee, K.H. Choi, H.D. Ghim, S.S. Kim, D.H. Chun, H.Y. Kim and W.S. Lyoo, “Role of molecular weight of atactic poly(viny1 alcohol) (PVA) in the structure and properties of PVA nanofabric prepared by electrospinning,” J. Appl. Poly. Sci., 93(4), 1638-1646 (2004). S.L. Zhao, X.H. Wu, L.G. Wang, and Y. Huang, “Electrospinning of ethyl-cyanoethyl cellulose/tetrahydrofuran solutions,” J. Appl. Poly. Sci.,2004. 91( 1): p. 242-246. H. Park, “Fabrication of lanthanum copper oxide nanofibers by electrospinning,” Ph. D dissertation, University of Florida (2005) W.K. Son, J.H. Youk, T.S. Lee, and W.H. Park, “The effects of solution properties and polyelectrolyte on electrospinning of ultrafine poly(ethy1ene oxide) fibers,” Polymer, 45(9), 2959-2966 (2004). J.H. Yuh, J. Nino, W. Sigmund, “Synthesis of barium titanate (BaTi03) nanofibers via electrospinning, Materials Letters, in press, 2005

Progress in Nanotechnology:Processing

Microstructure and Properties of Spark Plasma-Sintered ZrOz-ZrBz Nanoceramic Composites Bikramjit Basu’ and T. Venkateswaran Laboratory for Advanced Ceramics, Department of Materials and Metallurgical Engineering, Indian Institute of Technology, Kanpu, India

Doh-Yeon Kim School of Materials Science and Engineering and Centre for Microstructure Science of Materials, Seoul National University, Seoul 151-744, Korea

In a recent work,’ we have reported the optimization of the spark plasma sintering (SPS) parameters to obtain dense nanostructured 3Y-TZP ceramics. Following this, the present work attempts to answer some specific issues: (a) whether ZrOz-based composites with ZrBz reinforcements can be densified under the optimal SPS conditions for T Z P matrix densification (b) whether improved hardness can be obtained in the composites, when 30 vol% ZrBz is incorporated and (c) whether the toughness can be tailored by varying the ZrO2-matrix stabilization as well as retaining finer ZrOz grains. In the present contribution, the SPS experiments are carried out at 1200°C for 5 min under vacuum at a heating rate of 600 K/min. The SPS processing route enables retaining of the finer t-ZrO, grains (100-300 nm) and the ZrOIZrBz composite developed exhibits optimum hardness up to 14 GPa. Careful analysis of the indentation data provides a range of toughness values in the composites (up to 11 MPa.m’12), based on Y z 0 3 stabilization in the ZrOz matrix. The influence of varying yttria content, t-ZrOz transformability, and microstructure on the properties obtained is discussed. In addition to active contribution from the transformation-toughening mechanism, crack deflection by hard second phase brings about appreciable increment in the toughness of the nanocomposites.

A

I. Introduction MONG

the important structural ceramics, stabilized zirconia ceramics, in particular yttria-stabilized tetragonal zirconia ceramics (Y-TZP), are widely investigated, because of their excellent combination of fracture toughness and strength. Since the pioneering work of Garvie et al., research on Zirconia ceramics has been triggered rapidly during the last three decades.’-’ The high toughness of the tetragonal Zirconia monoliths originates from the stress-induced transformation of the tetragonal ( t ) phase to the monoclinic (m)phase in the stress field around the propagating cracks, a concept widely known as “Transformation t o ~ g h e n i n g . ” Various ~,~ dopants like yttria, ceria, magnesia, calcia, etc., play a major role in the stabilization of the hi h temperature tetragonal phase in the sintered microstructure. 87,8 In spite of having better fracture toughness, the moderate hardness ( FZ 1&12 GPa) of the TZP monolith, in comparison with other structural ceramics (Si3N4 or SiAlON- 16-22 GPa; A1203-upto 19 GPa), limits their use in several structural a p plicati~ns.~ ‘ I Various attempts”-I2 have been made to exploit D. B. Marshali-contrihuting editor

Manuscript No. 20623. Received June 1. 2005; approved February 25,2006. ‘Author to whom Correspondence should be addressed. e-mail [email protected]

Nanocornposites and Nanostructures

the toughness of tetragonal Z r 0 2 and to improve hardness by the composite approach, in which hard reinforcements are incorporated into the Z r 0 2 matrix. Among various hard phases, a transition boride like ZrB2 is used in the present work, because of its excellent hardness and wear properties. In view of its importance as a structural material, toughness tailoring and enhancement of the hardness are considered as important issues in developing Zr02-based composites. ZrB2, a potential high-temperature material, is characterized by a high meting point, high hardness, and resistance to molten metals, and good electrical and thermal properties.’%’’ Despite having such advantages, the difficulty in densification of boride ceramics arises from a high melting point, highly directional covalent bonds, and a low self-diffusion coefficient. Basu et al.” developed zirconia-toughened 30 vol% ZrB2 composites (- 9597% dense) by hot pressing in vacuum at 1450°C for 1 h. Optimum combination of mechanical properties, i.e., fracture toughness of 9 MPa .m‘” and a hardness of 13 GPa, was obtained with the optimized hot-pressed Zr02-ZrB2 composite. While the addition of ZrB2 can potentially improve the hardness of the composites, achieving full density in Zr02-ZrB2 composites requires adoption of a hi h processing temperature of 1400”-1500”C in hot pressin$‘ or Sinter-HIPing,I6 and the total processing time in either route is usually 3-4 h or longer. To this end, spark plasma sintering (SPS) has the potential to achieve maximum densification at a lower temperature in a shorter time. In our recent research, we have been able to demonstrate that fully dense nanostructured Y-TZP monoliths can be obtained after SPS processing at 1200°C for 5 min.’ Following this, the present work is taken up to incorporate 30 vol% ZrB2 particulate Y-TZP matrix in order to develop a hard and tough TZP nanoceramic composite using the SPS route. This contribution also analyzes the microstructure-mechanical property relationship of the newly developed materials.

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-

-

11. Experimental Procedures

( I ) Starting Powders In the present investigation, commercially available co-precipitated nanosized-zirconia (average primary particle size 27 nm) starting powders (Tosoh grade: 3 mol% Y-Zr02, TZ-3Y) were mixed with varying amounts of yttria-free monoclinic powders (Tosoh grade: 0 mol% Y-ZrO2, TZ-OY) to obtain “mixed grades” having an overall yttria content of 2, 2.25, and 2.5 mol%, respectively. Additionally, co-precipitated 2 mol% YZrO2 powders (Tosoh TZ-2Y grade, Tosoh Corporation, Minato-Ku, Tokyo, Japan) were also used in composite production. Details of the starting powder nomenclature are given in Table I. ZrB2 starting powder is synthesized in the laboratory by a Selfpropagating High-temperature Synthesis process (SHS). In the SHS process, commercial Zr02 powders (99% purity, particle 455

Table I. Nomenclature Assigned for the Nanoceramic Composites Developed Sample designation

TM2B TM2.25B TM2.5B T3B T2B

Mol% yttria in ZrOz matrix

Comments

2 mol% yttria 2.25 mol% yttria 2.5 mol% yttria 3 mol% yttria 2 mol% yttria

Mixed grade Mixed grade Mixed grade Co-precipitated grade Co-precipitated grade

B stands for ZrB2

size < 5 pm), boron (99% purity, particle size <40 pm), and Mg (99% purity, particle size < 100 pm) were taken in a stoichiometric amount and mixed thoroughly in a planetary ball mill (PM40&Retsch, Haan, Germany). The mixture was ignited from the top by an electric arc in argon atmosphere and the velocity of the combustion wave has been calculated to be as high as 25 cm/s. The reaction was carried out in a closed chamber, flushed, and filled with IOLAR-1 argon. After the completion of the SHS reaction, the reacted product consisted of ZrB2, MgO, and unreacted Mg. The reaction by-product, is., MgO, was removed by dissolving in dilute HCI and washing with distilled water, and subsequently metallic Mg was leached out. No further milling was required to yield fine ZrB2 powders. The as-synthesized ZrB2 powders were observed to be spherical, with a uniform size distribution (average size -0.30.6 pm). X-Ray diffractometer (XRD) analysis (not shown) of the powders obtained showed the predominant presence of ZrB2. More details of the synthesizing route of ZrBz powders and sintering of ultra fine ZrB2 powders can be found elsewhere." The "mixed-grade" yttria-doped zirconia powders as well as TZ-3Y and TZ-2Y powders are further mixed with ZrB2 powder in the required proportion (70:30 volume ratio) on a planetary mill using WC balls on WC vials in acetone milling media. The milling was carried out for 24 h and the ratio of milling balls to powder was 4:l.

( 2 ) SPS Processing Appropriate amounts of dried premixed powders were placed in a graphite die having 10 mm internal diameter, and proper dressing with a graphite sheet was performed inside the die to avoid contamination. Lower and upper graphite punches were inserted into the central hole of the graphite die. This whole die assembly was transferred carefully into the SPS chamber (SPS, Eltech Co., Ulsan, Korea) and placed in between the graphite electrodes. These electrodes transfer pressure and current to the die assembly. A vacuum of 70 mtorr and a pressure of 40 MPa was maintained throughout the experiment. A DC current of 11.5 kA and a DC voltage of 5-10 V was applied during the SPS experiments. The pulse frequency during spark sintering varies around 30-40 KHz. The sintering of Zr02-ZrBz nanocomposite was carried out at 1200°C at a heating rate of 600 K/min and a holding time of 5 min. The selection of the processing parameter was based on our earlier experience in obtaining fully dense Y-TZP nanoceramics in the SPS route.' At the end of the experiment, the current flow was stopped and the pressure was released gradually. The temperature during the entire SPS processing was monitored by focusing an optical pyrometer on the graphite die. The graphite die used was of high-quality commercial graphite, and a temperature calibration revealed that the temperature difference between the graphite die and the powder sample was limited to a maximum of 100°C in the temperature range of 1200°-1300"C. The final thickness of the spark plasma-sintered composite was around 2-3 mm. The removal of the graphite sheet around the sample was ensured before further characterization.

(3) Chavactevization The density of all the SPS-processed materials was measured according to the Archimedes principle in distilled water. All the 456

samples were smoothly polished using a diamond paste. Ihe phase identification of starting powders as well as polished and fractured samples was performed using an XRD (Isodebyeflex 2002 and 1001, Rich Seifert & Co., Ahrensburg, Germany). XRD data were collected using a diffractometer (Intel CPS-120, Santa Clara, CA), equipped with a copper-rotating anode. The XRD patterns were collected at a scan rate of O.OS"/min with a fixed counting time of 3 s. The volume fraction of the m-Zr02 (V,) was calculated by measuring intensities of (1 11) and ( 1 1 1) reflections of the monoclinic phase as well as the (1 1 1) peak of the tetragonal phase and following the formula of Toraya e / ~ 1 . In ' ~order to estimate quantitatively the ability of /-Zr02 to transform to m-Zr02 in the crack tip stress field, the transformability, defined as the difference in the m-Zr02 phase content (YO)calculated from the XRD patterns obtained from fractured and polished surfaces, is computed. Further, the microstructural characterization of all polished and fracture surfaces was carried out using scanning electron microscopy (FE-SEM, JSM-6330F, Philips, Eindhoven, The Netherlands). The elastic modulus ( E ) was measured using the pulse-echo resonance frequency method (Tektronix TDS 200, Panametrics Model 5800). The Vickers hardness measurements were performed on smoothly polished surfaces at a load of 10 kg for a dwelling time of 15 s, using a Universal Hardness tester. The reported hardness values are the average of five indentation measurements. As far as the toughness measurement is concerned, it should be noted that the tou hness of brittle materials is dependent on which are widely classified into long the test crack and short crack methods. Long crack methods include the single edge-notched beam (SENB) and single edge V-notched beam (SEVNB) techniques.20 Short crack techniques involve measurement of the crack lengths (radial/median) around hardness indentations, from which the toughness data can be approximated using various reported model^.'^-^^ Note that the absolute toughness values of brittle materials cannot be measured by the indentation technique, for which one has to adopt the long crack fracture toughness measurement techniques, e.g., S E W , SEPVB, Cheveron notch beam (CNB), etc. However, in order to compare the toughness property of the newly developed composites, the indentation technique has been used in the present work. Additionally, it is now widely recognized that careful use can yield reproducible results for indentation toughness measurements. In the present work, the fracture toughness of the SPSed composites was calculated by measuring the crack lengths onginating from the corners of the Vickers indents. At least five indentations were taken on smoothly polished surfaces to measure hardness and toughness. It should be mentioned here that the indentation method is now routinely used for convenience to compute the fracture toughness of small and relatively brittle specimens, which are otherwise hard to machine into standard test samples (for example, SENB, SEVNB, etc.). A critical summary of the toughness values of various Y-TZP monoliths, presented in a recent review article: also indicates that the indentation method is, by far, the most widely used method in measuring the toughness of transformation-toughened ceramics.

technique^,"^^

III. Results and Discussion ( I ) Densijication SPS processing of the Zr02-ZrB2 composites was carried out at 1200°C for a holding period of 5 min in vacuum at a heating rate of 600 K/min, and the density data are provided in Table 11. The developed composites exhibit maximum densification (- 95%99% pth) under the selected SPS conditions. The theoretical density is calculated following the rule of mixture, assuming a theoretical density of 6.10 and 6.09 g/cm3 for Zr02 and ZrB2, respectively. The mixed-grade 2 mol% yttria-stabilized grade (TM2B) exhibited a slightly lower densification of around 95% Progress in Nanotechnology: Processing

Table 11. Overall Properties, i.e., Relative Density (RD), Elastic Modulus (4,Hardness (Ifvla), and Fracture Toughness ( 4 3 ,of the ZrOrZrBz nanocomposites, SPSed at 1200°C Material designation

TM2B TM2.25B TM2.5B T3B T2B

Relative Density

(YOpth)

Elastic Modulus (E, GPd)

Vickers Hardness ,hf ( GPa)

95.4 97.8 98.4 98.4 98.5

250 260 257 266 26 1

12.3 +O. 1 13.1 k0.2 13.7 t 0 . 1 13.9t0.3 13.9k0.2

The standard deviation in the hardness data measured is also shown

pth when compared with the co-precipitated powder-based composites. The sintering of nanocrystalline 3Y-Zr02 powders to full density via the SPS route was effectively carried out at 1200°C for a holding period of 5 min at a heating rate of 600 K/ min in vacuum.' Recent work in our research group indicated that the full density in a similar composite could be obtained after sinter-HIP at 1400°C for 1 h under a 110 MPa argon pressure.I6 Also, Basu et d."reported -97% pthin Zr02-ZrB2 (70/ 30) composites after hot pressing (HP) at 1450°C for 1 h in vacuum. Therefore, it can be stated that the full density in ZrO2ZrB2 (70/30) composites can be obtained by the SPS route at 2O0"-25O0C lower temperatures with reduced overall processing time of 30 min, which is much less than that by the HP or HIP route. The reduction of the total processing time is a useful advantage of the SPS process. The fact that the Zr02 composite containing 30 vol% of difficult-to-sinter powder particulates, i.e., ZrB2, is sintered to near-theoretical density under optimal SPS conditions for ZrOz matrix densification further reconfirms such an advantage. From the SPS point of view, the passage of current depends on the electrical conductivity property of the powder compact. In case of a powder compact containing only ZrO2 powders, the major part of current passes through the graphite die, with a smaller fraction passing through insulating Zr02 particles. In a recent work of Vanderbiest and co-workers, finite-element (FE) calculations26of the temperature field during the field-assisted sintering technique (FAST) have revealed that the temperature distribution in case of Zr02 is more homogeneous, while that of conducting ceramic particle compacts like TiN is non-homogeneous in the presence of more thermal gradients. In the present case, because of the presence of a larger amount of conducting ZrB2 phase, a comparatively larger fraction of the total current will pass through the powder compact and therefore, spark discharges in the void spaces become more frequent in the composite powder compact. We believe that this will enhance the mass transport, leading to increased neck growth and hence, enhanced densification is obtained for the composites at a much lower temperature by the SPS route. It can be mentioned here that the driving force in the case of conductive ceramics (ZrB2)is joules heating effect.27As has been well established, the oxide layers on the non-oxide particle surfaces, like borides, are formed as a result of selective oxidation of base transition metal. However, the surface cleaning of ZrBz particles, during plasma sintering might possibly have occurred and this could enhance the densificdtion of the composite material at a much lower temperature. Moreover, due to the use of finer starting powders of Zr02 (average particle size -27 nm) and ZrB2 (average particle size -0.3-0.6 pm), the increase in sintering driving force, i.e., the decrease in surface-free energy, triggers densification, ensuring grain boundary diffusion as the dominant mass transport path. From the above discussion, the enhancement of the driving force for sintering because of the electric field-assisted faster heating, pressure application, and surface activation, combined with grain boundary diffusion as the dominant mass transport mechanism, enable rapid densification of the developed Zr02ZrBz nanocomposite in the SPS process. Nanocornposites and Nanostructures

( 2 ) Microstructure XRD phase analyses of the starting powders and polished surfaces of the SPS-processed materials are shown in Fig. 1. The XRD spectra obtained clearly indicate the predominant phase for the Zr02 matrix as t-Zr02. However, there is also an indication for the presence of m-Zr02 in the TM2B composite. Otherwise, the addition of ZrB2 is found to have no observable influence on t-Zr02 retention. Note the apparently lower relative density of the TM2B composites as compared with other composites (see Table 11), which is due to the fact that the presence of rn-ZrO2 is not taken into account while calculating the theoretical density. Therefore, the amount of porosity would be, in fact, less in the TM2B composite as otherwise could be expected from the density data of Table 11. Selected SEM fractographs of both mixed and co-precipitated ZrO2 powder-based composites are shown in Fig. 2. Based on SEM images, the microstructure is characterized by the presence of coarser tabular/elongated ZrB2 particles (- 2-3 pm) and equiaxed nano-Zr02 particles (- 1W300 nm). A schematic illustration of the phase assemblage in the microstructure is presented in Fig. 2(d). A comparison of Figs. 2(aHc) indicates that the ZrOz grain sizes in T3B composite are relatively finer ( 1 0 s 150 nm) than T2B and TM2.25B composites. The average Zr02 grain size in the T2B composite appears to be around 200-300 nm, while the presence of few coarser ZrOz grains of sizes 300

25

I

2

35 40 28 (degrees)

30

A L d

TM2.5BA

50

55

after SPS at 1200°C

t-ZrO, m-ZrO, ZrB,

TM2.25B

45

JL 1

), J,

n 20 (b)

25

30

35 40 28 (degrees)

45

50

55

Fig. 1. XRD spectra obtained with starting powder (a) and SPS-synthesized nanocomposites, sintered at 1200°C at a heating rate of 600 K/ rnin and holding period of 5 min in vacuum (b).

457

Fig. 2. SEM fractographs of spark plasma-sintered ZrOrZrBz nanoceramic composites, sintered at 1200°C for a holding period of 5 min in vacuum: T3B grade (a), T2B grade (b), and TM2B grade (c). The presence of finer Zr02 grains (10&300 nm) and coarser ZrB2 particulates (2-3 pm) can be distinguished. A model ceramic nanocomposite microstructure with nanosized matrix particles reinforced with microsized reinforcement particulates, as observed in the newly developed materials, is shown in (d).

nm or more can be critically noted in the TM2B composite. As far as ZrB2 size is considered, a closer observation of Figs. 2 ( a t (c) indicates that considerable grain growth occurs for boride particles from 500 to 600 nm (starting particle size) to 2-3 pm (SPSed material). Qualitatively, the grain growth appears to be more in the case of conductive particles than insulating ZrOz in our SPSed composites. The fractographs further reveal the transgranular fracture of ZrB2 particles, and predominantly intergranular fracture of n a n o - Z Q grains. The ZrB2 particulates are well bonded in the nano-ZrO2 matrix in all the composites. The fact that a finer microstructure (Zr02- 10&300 nm and ZrB2 2-3 pm) is obtained after SPS processing indicates that the spark plasma processing route is effective in retaining of a fine-grained microstructure in a dense nanocomposite. The correlation between microstructure and properties will be discussed in the subsequent sections.

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( 3 ) Mechanical Property The mechanical property data, presented in Fig. 3, Tables I1 and 111, provide some interesting facts regarding the mechanical behavior of the Zirconia nanoceramic composites developed. The E-modulus data, as presented in Table 11, reveal that the addition of ZrB2 increases modulus up to 266 GPa. The measured E-modulus varies between 250 and 266 GPa. Also, the lower elastic modulus of 250 GPa, as measured with the TM2B composite, can be attributed to the presence of m-Zr02 and/or porosity than that in other composites. As far as the hardness is considered (see Table 11), SPS-processed nanoceramic composites exhibit moderate hardness varying in the range of 12-14 GPa. Higher hardness is measured with the composites, synthesized using co-precipitated Z r 0 2 powders (Table 11). Importantly, it should be pointed out here that no observable hardness improvement is recorded, despite incorporating 30

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458

vol% of harder ZrB2 particulates into a nanosized zirconia matrix. Such an observation can be explained based on similar observations made in other ceramic systems, including monolithic ZrB2.2S-30In an earlier study, Green also measured a substantially lower hardness of 10-12 GPa in AI2O3composites, containing more than 10 vol% Zr02.28Such an observation was attributed to the presence of microcracks, which were either

I

I

13

I1

1:

10

1

, . 2.0

0

y

P

71 6

E

,

. , .

I

2.2

2.4

2.6

.

z

2.8

mol% Y203 in ZrO, matrix

.

t

16

3.0

Fig.3. Plot of hardness and toughness versus yttria stabilization Zr02 matrix for spark plasma-sintered (120O0C, 5 min) nanocomposites. Different symbols represent the mechanical property measured with various composites based on the ZrO2 matrix processed from either co-precipitated or mixed-grade Zr02 powders: , H,,, of the mixed grades; 0, H , , , of the co-precipitated grades; A, 6,. of the co-precipitated grades (using the Palmqvist formulae); and A, KI, of the mixed grades (using the Plamqvist formulae). Progress in Nanotechnology: Processing

formed because of postfabrication cooling or developed because of indentation stresses. In a recent work, Chamberlain et measured a lower hardness of 14.5 GPa in a pressureless-sintered (2150"C, 9 h, helium atmosphere) monolithic ZrB2. In contrast, a much higher hardness of 23 GPa was measured in hotpressed monolithic ZrB2 (1900"C, 45 min, vacuum). Such a significant difference in hardness was attributed to the differences in the ZrB2 grain sizes, which are expectedly much coarser (average size 9.1 pm) in the pressureless-sintered material. A similar observation was also made by Lee and S ~ ey er,~'who measured a decrease in hardness from 18.3 to 16.5 GPa, as a result of an increase in the grain size from 2.2 to 3.1 pm, respectively, in case of pressureless-sintered (90% theoretical density) monolithic B4C. From the mechanistic point of view, it is quite likely that the frequency with which the movement of dislocations is impeded will decrease with an increase in grain size and this would lower the stress required for deformation.29s30 Therefore, at a given indent load, more deformation zone around the indented region can be expected in coarser-grained material. In the present case, the absence of any observable hardness increment in SPSed composites can be attributed to both the microcracking effect and the coarser particle size of ZrB2. The microcracking effect is quite probable, because of the presence of residual stress. In the Zr02-ZrB2 materials developed, the thermal residual stress is developed due to two factors: the first one is the anisotropy in the coefficient of thermal expansion "C-' (CTE) of the tetra onal zirconia matrix (a, = 11.4 x and a,= 7.1 x lo-' "C-' )31 as well as that of ZrB articulates (a, = 6.7 x lop6 "C-' and a, = 5.6 x lop6 "C-l )T !'he other factor is due to the CTE mismatch between Zr02 matrix and ZrB2 particles. Because of the tensile residual stress in ZrOz, the microcracking potentially takes place either during post-fabrication cooling or in the indented region. Besides possible contribution from microcracking, the coarser ZrB2 particles can potentially, like in other ceramic system^,^"^^ contribute to lowering the hardness of Zr02-ZrBz nanoceramic composites. In order to compute the toughness data, SEM analysis of the Vickers indents (taken at 10 kg) was carried out and the crack lengths around the indentations were measured from the SEM images. Typical SEM images of the indented composites material surfaces, as shown in Fig. 4, clearly indicate the presence of radial cracks, which predominantly emanate from the indent comers. All the relevant indentation data, i.e., the indent diagonal length ( 2 4 and total crack length (2c),used in hardness and toughness measurements are provided in Table Ill. The indentation data reveal that l/a varies in the range of 0.34.9. In estimating the toughness values, we have considered various indentation fracture mechanics models, which can be broadly classified into two categories: (a) The toughness can be measured using Anstis's modellg for brittle materials, exhibiting radial-median cracks ( / / a> 2.5):

-

where E is the elastic modulus, H the Vickers's hardness, P the indent load, and c the half crack length: (b) The toughness of brittle materials, exhibiting Palmqvist type cracks (0.25 < //a < 2.5) can be computed by a different equation2?

where 2a is the average indent diagonal length, 2c is the crack length, and 1 = c-a. In both the above equations, the dimensionless quantity, q, is a constant for a given indenter geometry, provided the volume is conserved within the "plastic zone" (adjacent to indentation). In case of transformation-toughened materials, the crack driving force-induced t-Zr02 transformation additionally modifies the Nanocomposites and Nanostructures

Fig.4. SEM topography images of the Vickers indents and indentation-induced radial crack pattern in the SPS-processed zirconia nanoceramic composites: T38, T2B, TM2B.

value of the dimensionless quantity (q). It is worthwhile to note that Kaliszewski et a/." have critically analyzed the transformation effect around the indented zone and reported that an appropriate value of q = 0.019 (instead of q = 0.016 as widely used for non-transforming ceramics) can be used when the transformation zone size as well as the volume fraction of tZrO2 (transformed within the transformation zone) are both experimentally obtained for Zr02 materials. As such measure459

Table 111. Indentation Data, i.e., Average Indent Diagonal Length (2a) and Total Crack Length ( 2 4 , as well as the Fracture Toughness Values for the SPSed ZrOrZrB, Nanocomoosites Material designation

TM2B TM2.25B TN2.5B T3B T2B

Indent diagonal Crack length Indentation toughness ( 2 4 (pm) (2c) (pm) //u (MPa. ml'*)

122 1I8 1I5 I15 1I4

158 185

197 210 166

0.3 0.6 0.1 0.9 0.4

11.4&1.0 8.7k0.9 7.8*0.3 6.9L0.7 10.0&1.2

The crack length parameter (I) is defined as I = c-u. The standard deviation in the measured toughness data is also shown.

ments have not been performed, Anstis's formulation has not been used in the present work. It should now be realized that the choice of dimensionless constant (q) from various models can considerably influence the value of Klc, calculated using Eq. (I) or (2). In an earlier work, Niihara ef ~ 1 experimentally . ~ ~ obtained the value of q to be 0.0089 and 0.0122 for l/a ratios varying in the range of 0.25-2.5 and 1-2.5, respectively. In the present study, we have chosen Eq. (2) with q = 0.0089, because the l/a ratio lies in the range of 0.34.9 (see Table 111). Besides the above-mentioned reason, the use of Anstits's formula would be inappropriate as the l/a ratio (measured in our work) is much less than 2.5, a value characterizing the radial-median crack pattern. A similar analysis in using the toughness equations has been performed in a recent work, reporting the mechanical properties of the SPS-processed zirconia-containing composite~.~' The fracture toughness data are presented in Table 111 and plotted in Fig. 3. The toughness data can he interpreted in terms of the difference in Zr02 matrix stabilization. For the co-precipitated powder-based Zr02 matrix, a decrease in overall yttria content from 3 to 2 mol% significantly increases the fracture toughness from 6.9 (T3B composite) to 10.0 MPa . m'" (T2B composite). It can be noted that nanocrystalline T3 monoliths, SPSed at 1200°C for 5 min, is reported to have a fracture toughness of -4.5 MPa.mii2.' This indicates that ZrB2 addition increases the fracture toughness of the composites. For the mixedgrade composites (TM based), the decrease in overall Y content from 3 to 2 mol% increases the tou hness from 6.9 MPa . mli2 (T3B composite) to 11.4 M P a . m ' j (TM2B composite). The toughness is also observed to increase systematically as the Y content is lowered from 3 to 2.5 mol% and finally to 2 mol%. Because of a greater increase in toughness by decreasing the Y content from 2.5 to 2 mol% rather than by decreasing in Y content from 3 to 2.5 mol%, an additional experiment was conducted to develop a composite with a ZrOz matrix having 2.25 mol% Y stabilization. The mechanical property measurement on the TM2.25B composite indicated the moderate toughness improvement as revealed by KlCdata: 8.7 MPa . mli2 for the TM2.25B grade as compared with 7.8 MPa.rnii2 for the TM2.5B grade. Broadly, the toughness data indicate that the toughness of the Zr02-ZrB2 composites can be tailored by carefully optimizing the Y stabilization of the ZrO2 matrix, when all are SPSed under identical sintering conditions. However, a higher toughness is measured with the mixed-grade TM2B composite, when compared with the co-precipitated powder-based T2B composite. This certainly indicates that although the hardness of TM2B (12.3 GPa) is inferior compared with that of T2B (13.9 GPa), the toughness property is better in the mixed-grade TM2B composite than T2B material.

( 4 ) Toughening Mechanisms XRD results, as mentioned in Section III(2), indicated that tZr02 is predominantly retained in all the SPSed composites (polished surfaces). Only in the TM2B composite, a trace amount of m-Zr02 content was also recorded. Furthermore, 460

the m-Zr02 content was also measured on the fracture surfaces in order to evaluate t-Zr02 transformation during fracture crack propagation. The transformability, the difference between the fracture and the polished surface m-Zr02 content (%), varies around 50%-60% for all the composites (see Fig. 5). A critical comparison of the toughness data (Table 111) and the transformability data (Fig. 5 ) indicated that the toughness shows an almost linear correlation with transformability. In particular, the high-toughness TM2B composite also exhibits the highest transformability. However, the observation that both TM2.5B and TM2.25B composites experience a slightly higher transformability, but lower toughness, than the T2B composite should be attributed to the use of two different ZrO2 starting powders used in the composite production. In the TM-grade composites, the mixed powders, i.e., a mixture of 0 mol% yttria and 3 mol% yttria-containing powders, are used, while T2B is processed from co-precipitated ZrO2 starting powders. Such a difference in starting powders is reported to result in a difference in the toughness of sintered Y-TZP monoliths (all hot pressed under identical conditions), as reported el~ewhere.~ Nevertheless, the toughness increase in both the TM-grade composites as well as co-precipitated powder-based composites (T2B and T3B) commensurate well with the increase in measured t-Zr02 transformability. This is expected as the toughness of the transformation toughened zirconia is directly pro ortional to the volume fraction of transformable t-Zr02.3,! Also, a large amount of monoclinic Zr02, relative to that on the polished surface, with a maximum of 60% in the case of the TM2B composite, has been measured on fracture surfaces (Fig. 5). Such an observation should be attributed to the fact that f-Zr02 is largely transformable in the SPSed composites and hence transformation toughening is considered as a toughening mechanism. The increased transformability, an indication of a higher transformable t-Zr02 content, clearly implicates a higher transformation-toughening contribution in the SPSed Zr02 nanocomposites. Regarding the microstructure-toughness relationship, the influence of grain size on toughness needs to be considered. It is noted in our work that finer t-Zr02 grain sizes (10&300 nm) are retained and such finer grains are also able to transform in the crack tip stress field, leading to a higher transformation toughness in the nanocomposites. To this end, it can be recalled that such a finer zirconia microstructure (in case of Y-TZP monoliths) should be extremely stable and should not transform at

c

Etn

60

A

50-

V

A

A

V

c

-l

40-

s?

30-

e'?I " 201 1.75

*

I

2.00

I

'

I

1

I

2.25 2.50 2.75 mol% yttria in ZrO, matrix

8

I

3.00

'

3.25

Fig. 5. Percentage transformability of the Zr02 matrix against the mol% of yttria content in that matrix at sintering temperatures of 1200°C. The f-Zr02 transformability is defined as: [% m-ZrOz on the fractured surface-% m-Zr02 on the polished surface]. The transformability data of various composite grades processed from the use of Zr02 powders are indicated by different symbols: V,co-precipitated grades; and A, mixed grades.

Progress in Nanotechnology: Processing

predicted that in addition to deflection by boride particles, the crack tip stress field energy is utilized in debonding the ZrBz particles and thus provides less driving force for further crack propagation.

IV. Conclusions

Fig.6. Crack deflection by the hard second-phase ZrB2 particles (a) and crack wake debonding of the coarser ZrBz particles at the interface (b) in the T2B nanocomposite, SPSed at 1200°C for 5 min (heating rate:

600 Kjmin).

the crack tip, according to the earlier literature reports3“ Therefore, the attainment of higher fracture toughness (up to 10 MPa . m1I2)with the newly developed composites implicates that other factors, besides grain size, should also be considered. As far as the transformation toughening is concerned, another important factor, apart from the yttria content and grain size, is the residual stress. Because of the CTE mismatch, as mentioned earlier, Zr02 grains are subjected to residual tension in a sintered microstructure. Therefore, the critical stress needed to induce the transformation of ZrOz grains in the crack tip would be reduced accordingly, as also predicted from the established l i t e r a t ~ r e . ~This .~’~ increases ~ the driving force for stress-induced t-Zr02 phase transformation. This should also contribute to the enhanced transformability of f-ZrOz. Coming to the other factors influencing toughness, the contribution of additional toughening mechanisms besides transformation toughening needs to be considered. The crack pathmicrostructure interaction in the SPSed nanocomposites is illustrated in Fig. 6. Figure 6(a) clearly reveals the increase in crack path tortuosity (more like a sinusoidal crack path) because of crack deflection by hard and coarser ZrBz particles. A closer observation of Fig. 6(b) also reveals the crack wake debonding of the ZrB2, as evident from the crack propagation around the ZrOz/ZrBZinterface. From the above observations, it can be Nanocornposites and Nanostructures

(a) Near theoretical density of the yttria-stabilized tetragonal zirconia nanocomposites, reinforced with 30 vol% of ZrBz by the SPS route, is obtained at 1200°Cwith a holding time of 5 min and a heating rate of 600 K/min under vacuum. The sintering temperature, which is 200”-250”C less, as well as the total processing time (- 20-25 min) are considerably less in comparison with established densification processes like hot pressing or sinter-HIPing. (b) The composite materials developed exhibit a finer microstructure, with ZrB2 exhibiting an average grain size of 2-3 pm and ZrOz in the range of 100-300 nm. Also, XRD results indicate that r-Zr02 is predominantly retained in all the SPSed materials. No noticeable hardness improvement, despite incorporating 30 vol% ZrB2 particulates, is recorded in spark plasma-sintered ZrOz composites, and the hardness of the developed composites remains moderate, varying in the range of 12-14 GPd. (c) Careful use of indentation data provides an estimate of toughness properties, and the computed toughness values of SPSed Zr02-ZrBz composites vary in the range of 6 1 1 MPa .m’”, depending on the YzO3 stabilization level of the Zr02 matrix. An important observation has been that the toughness can be tailored by varying the Y203 content, followed by mixing 3Y-Zr02 with undoped ZrOz in the starting powder mixture. Considerable amount of f-Zr02 is observed to transform to m-ZrO2 on the fracture surface and the t-ZrOz transformability broadly increases with reduction in overall yttria stabilizer content of Zr02 matrix. Such experimental observations suggest that the transformation toughening is an active toughening mechanism in the SPSed Zr02 nanocomposite. A critical observation is that a high transformation-toughening contribution is obtained with a ZrOl matrix having finer tetragonal grains (1W300 nm). The experimental observations also suggest that both crack deflection by ZrB2 particulates and crack wake debonding at ZrOz/ZrB2interface contribute toward attaining the toughness obtained.

Acknowledgments T. Venkateswaran thanks the members of the Sintering and Microstructure laboratory and the Thin film & Microstructure laboratory (TFML) for their cooperation during SPS experiments. The authors also thank Dr. Suman K Mishra, National Metallurgical Laboratory, Jamshedpur, India, for supplying finer ZrB, powders.

References ‘B. Basu. J. H. Lee, and D. Y. Kim, “Development of Nanocrystalline Wear Resistant Y-TZP Ceramics,” J. Am. Ceram. Soc., 87 [9] 17714 (2004). 2R. C. Garvie, R. H. Hannink. and R. T. Pascoe. “Ceramic Steel?,” Narure, 258, 703-8 (1975). ’B. Basu. J. Vleugels, and 0. Van Der Biest, “Toughness Tailoring of YttriaDoped Zirconia Ceramics,” Mater. Sci. Eng., A 380, 2 15-2 I (2004). 4B. Basu, J. Vleugels, and 0. Van Der Biest, “Microstructure-Toughness-Wear Relationship of Tetragonal Zirconia Ceramics.” J. Eur. Ceram. Soc., 24 [7] 203 I40 (2004). 5A. G. Evans, “Perspective on the Development of High-Toughness Ceramics,” J. Am. Ceram. Soc., 73 [2] 187-206 (1990). 6A. G. Evans and R. M. Cannon, “Toughening of Brittle Solids by Martensitic Transformations,” Acra Marer., 34 [5] 761-800 (1986). ’P. E. Royes Morel and I.W. Chen, “Transformation Plasticity of Ce203-Stahilized Tetragonal Zirconia Polycrystals I,Stress Assistance and Autocatalysis,” J. Am. Ceram. Soc., 71, 343 (1988). ‘B. Basu, “Toughening of Y-Stabilized Tetragonal Zirconia Ceramics.” Inr. M a w . Rev., 50 141 239-56 (2005). ’J. B. Watchman, Mechanical Properties of Ceramics: Chapter 11, pp. 3 9 1 4 8 . John Wiley & Sons Inc., U.S.A. ‘OR. G. Munro. “Evaluated Material Properties for a Sintered u-Alumina.” J. Am. Ceram. Soc., 80 [8] 1919-28 (1997).

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“B. Basu, J. Vleugels, and 0 . Van Det Biest, ”Development of Zr02-ZrB2 Composites,” J. Alloys Comp.. 334 [ 1-21 2 0 0 4 (2002). ”B. Basu, J. Vleugels, and 0. Van Der Biest, “Transformation Behaviour of Yttria Stabilized Tetragonal Zirconia Polycrystals-TiB2 Composites.” J. Muter. Res.. 16 [7] 2158-69 (2001). ‘ k . Mortz, “Zirconium Diboride.” Am. Ceram. SOC.Bull., 73 [6] 141-2 (1994). I4R. A. Cutler, “Engineering Properties of Borides”; p. 787 in Engineered Materiols Handbook, Ceramics and Gla.$ses,Vol. 4 , ASM Handbook. Edited by S. J. Schneider. The Materials Information Society, U.S.A., 1991. I5C. F. Powell, p. 349 in Borides in High Temperature Materials and Technology, Edited by 1. E. Campbell, and E. M. Sherwood. John Wiley & Sons Inc., New York, 1967. I6S. Das Bakshi, “Processing and Properties of TZP-ZrB2 Composites by Sinter-HIPing”; MTech Dissertation, Indian Institute of Technology. Kanpur, November 2004.. ”S. K. Mishra, S. Das, S. K. Das, and P. Ramachandra Rao, “Sintering Studies on Ultra Fine ZrBZ Powder Produced by Self-Propagating High-Temperature Synthesis Process.” J. Muter. Res., 15 [II] 2499-504 (2000). ‘*H. Toraya, M. Yoshimura, and Somiya, “Calibration Curve for Quantification Analysis of the Monoclinic-Tetragonal Zr02 System by X-Ray Diffraction.” J . Am. Cerum. Soc., 67,C119-21 (1984). I9G. R. Anstis, P. Chantikul, B. R. Lawn, and D. B. Marshall, “A Critical Evaluation of Indentation Techniques for Measuring Fracture Toughness,” J. Am. Cerum. Soc., 64, 553-7 (1981). ’“B. R. Lawn, Frucrure of Brittle Solids, 2nd edition, Cambridge University Press, Cambridge, U.K., 1992. ”M. S. Kaliszewski, G. Behrens, A. H. Heuer, M. C. Shaw, D. B. Marshall, G. W. Dransmann, and R. W. Steinbrech, “Indentation Studies on Yz03-Stabilized Zr02: 1. Development of Indentation-Induced Cracks,” J. Am. Cerum. Soc., 77, 1185-93 (1994). 2zK. Niihara, R. Morena, and D. P. H. Hasselman, “Evaluation of 4, of Brittle Solids by the Indentation Method with Low Crack-To-Indent Ratios,” J . Muter. Sci. L p f t . , I, I S 6 (1982).

462

*ID. K. Shetty, 1. G. Wright, P. N. Mincer. and A. H. Clauer. “Indentation Fracture of WC-Co Cermets,” J . Muter. Sci.. 20, 1873-82 (1985). z4P. Chantikul, G. R. Anstis, B. R. Lawn, and D. B. Marshall, “A Critical Evaluation of Indentation Techniques for Measuring Fracture Toughness,” J. Am. Ceram. Soc., 64 191 53943 (1981). 2sS. Palmqvist, “Occurrence of Crack Formation During Vickers Indentation as a Measure of the Toughness of Hard Materials.” Arch Eisenhuetteniws., 33. 629-333 (1962). 26 K. Vanmeensel, A. Laptev. J. Hennicke, J. Vleugels, and 0. Van Der Biest. “Modeling of the Temperature Distribution During Field Assisted Sintering,” Acta Muter., 53,4379-88 (2005). ”5. R. Groza and A. Zavalianges, “Sintering Activation by External Electric field,” Muter. Sci. Eng., A287, 171-7 (2000). 28D. J. Green, “Critical Microstructures for Microcracking in Al2O3-ZrO2 Composites,” J. Am. Cerum. Soc.. 65 [I21 6 1 0 4 (1982). 2yL. Chamberlain, W. C. Fahrenholtz, and G. E. Hilmans. “Pressureless Sintering of Zirconium Diboride,” J. Am. Ceram. Soc., 89 121 4 5 M (2006). ’OH. Lee and R. F. Speyer, “Hardness and Fracture Toughness of Pressureless-Sintered Boron Carbide (B4C),” J. Am. Ceram. Soc., 85 [5] 1291-3 (2002). ”S. Schmauder and H. Schubert, “Significance of Internal Stress on the Martensite Transformation in Yttria-Stabilised T e t r a g o d Zirconia Polycrystals Durin Degradation,” J. Am. Cerum. Soc., 69 171 53440 (1986). f2R. N. L Okamoto. M. Kusakaari, K. Tanakd, H. Inui, and S. Otani, “Temperature Dependence of Thermal Expansion and Elastic Constants of Single Crystals of ZrB2 and the Suitability of ZrB2 as a Substrate for GaN Film,” J . Appl. Ph s ,93[I] 88-93 (2003). ‘R. Kumar, K. H. Prakasb, P. Cheng, and K. A. Khor, “Microstructure and Mechanical Properties of Spark Plasma sintered Zirconia-Hydroxyapatite NanoComposite Powders.” Acta Muter., 53 [8] 2327-35 (2005). 34B. Budiansky, J. W. Hutchinson, and I. C. Lambropoulos, “Continuum theory of Dilatant Transformation Toughening in Ceramics,” Int. J. Solids Strucr., 19,337 (1983). 0

Progress in Nanotechnology: Processing

Homogeneous Zr02-A1203 Composite Prepared by Nano-Zr02 Particle Multilayer-Coated A1203 Particles Yu Jia, Yuji Hotta,*" Kimiyasu Sato, and Koji Watari* National Institute of Advanced Industrial Science and Technology (AIST), Nagoya 463-8560, Japan

ZrO2-AI2O3 nanocomposite particles were synthesized by coating nano-Zr02 particles on the surface of A1203 particles via the layer-by-layer (LBL) method. Polyacrylic acid (PAA) adsorption successfully modified the A1203 surface charge. Multilayer coating was successfully implemented, which was characterized by 6 potential, particle size. X-ray diffraction patterns showed that the content of Z r 0 2 in the final powders could be well controlled by the LBL method. The powders coated with three layers of nano-Zr02 particles, which contained about 12 wt% Zr02, were compacted by dry press and cold isostatically pressed methods. After sintering the compact at 1450°C for 2 h under atmosphere, a sintered body with a low pore microstructure was obtained. Scanning electron microscopy micrographs of the sintered body indicated that ZrO2 was well dispersed in the A1203 matrix. I. Introduction

I

is well-known that ZrO2-AI2O3composites and nanocomposites have improved mechanical properties compared with A1203and Zr02,e.g., fracture toughness and slow crack growth propagation threshold.'-' On the other hand, monodispersed nanoparticles are advantageous in the fabrication of ceramics with improved properties: mechanical, electrical, thermal, ionic conducting, catalytic, and optical."' The homogeneous microstructure of the composites has a considerable effect on properties such as strength, wear resistance, and fatigue.699-' Therefore, the homogenous microstructure is quite important to improve the composites properties. Although attainment of a very fine and homogenous microstructure using classical powder-mixing processes has been proved to be almost impossible,'2-'4 a homogeneous microstructure with no apparent aggregates was obtained via colloidal processes. s23 In view of this point, sol-gel processes, using an alcoholic solution of zirconium alkoxides, have been widely However, characteristics of the particles prepared by the sol-gel process were sensitive to the experimental conditions. On the other hand, several aqueous-based methods such as the forced hydrolysis method or coprecipitation method, using an aqueous solution of zirconium salts such as the zirconium oxychloride (ZrOC12.8H20), have been investigated to produce nanocomposites particle^.^'-^^ In these processes, adjustment of pH and addition of electrolytes were generally carried out to prepare the composite particles. However, some aggregation of zircon hydroxide precipitation remained in the as-synthesized composite powders which would result in inhomogeneous final composite ceramics. Otherwise, the layer-by-layer (LBL) coating method T

'

J. Halloran+ontributing

editoi

Manuscript No. 20797. Received July 29, Z W , approved October 3, 2005.

This work has been suDmrted by a colbdbordtion of "Research and Develoument of the Low Environmental Burden Process.' of NGK INSULATORS, LTD. and AiST. 'Member, American Ceramic Society. 'Author to whom correspondence should be addressed. e-mail: [email protected]

Nanocomposites and Nanostructures

was well established for the preparation of nanocomposite part i c l e ~ 26f , ~ ~Using the LBL method, it is possible to prepare coated core particles of different shapes and sizes with uniform layers of diverse composite as well as controllable thickness. In the present research, A1203 powders coated with nanoZr02 particles were prepared via the LBL method. The influence of polyacrylic acid (PAA) adsorption and nano-Zr02 particle coating on the composite particles prepared was investigated to produce ZrO2-AI2O3composites.

11. Experimental Procedure Ten grams a-AI2O3 (TMDAR, average particle diameter: 150 nm, Taimei Chemicals, Nagano, Japan) was dispersed in a 100 mL, 0.03M NaCl solution by ultrasonicating with a homogenizer at 40 W of power for 2 min. The pH of the suspension was adjusted to 4, and aged for 3 h. At the same time, 100 mL PAA (MW 5000, Wako Pure Chemical Industry, Osaka, Japan) solution (0.6 wt% in 0.03M NaCl solution) was adjusted to pH 4, and aged for 3 h. Then, the PAA solution was mixed with the A1203suspension for 15 h. The mixed suspension was separated and washed with distilled water by centrifuging twice to remove the excess PAA. PAA-adsorbed AI2O3particles were dispersed in 100 mL of a 0.03M NaCl solution, and the pH was adjusted to 5 for shell deposition. In shell deposition, 100 mL ZrOC12.8H20 (Sigma-Aldrich Chemical Company, St. Louis, MO) with an 8 mol% YCI3.6H2O (Wako Pure Chemical Industry) solution was mixed with 50 mL diglycol (Wako Pure Chemical Industry) and aged for 8 h, and then the pH was adjusted to 5 by NaOH. After the mixture was held for 16 h, nanoZr02 particles were obtained. The suspension was added to the PAA-adsorbed A1203 suspension at a rate of 5 mL/min, and stirred for 12 h. Then, the suspension was separated and washed with distilled water by centrifuging three times to remove the residual nanoparticles. After repeating the steps of the PAA adsorption and shell coating, nano-Zr02-coated A1203 powders were prepared. The powders were dried and uniaxially pressed at 10 MPa in a steel die to make pellets. These pellets were cold isostatically pressed (CIPed) under 150 MPa for 20 min. Finally, the compacts were sintered at 1350", 1450", 1500", and 1550°C for 2 h. The morphology of the synthesized nano-Zr02 particles was observed using a transmission electron microscope (TEM, JEM2010, JEOL, Tokyo, Japan). The 6 potential of nano-Zr02 particles, A1203,and composite particles was characterized using a laser electrophoresis analyzer (Model 502, Nihon Rufuto Co. Ltd, Tokyo, Japan). Particle size distribution of the powders was analyzed with a laser-scattering particle size distribution analyzer (LA-920 Horiba, Kyoto, Japan). Scanning electron microscopy (SEM, JSM5600N, JEOL) was used for observing the microstructures of the sintered bodies. X-ray diffraction (XRD) patterns of the samples were obtained with a Rigaku RINT2550 diffractometer (Tokyo, Japan) using monochromatic CuKu radiation operated at 40 kV and 200 mA. As a reference, 'JCPDS file number 17-0923.

463

80.0

I

I

-80.0 I

PH

Fig. 2. 5 potential of as-received Alz03particles ( ), Polyacrylic acidadsorbed particles ( O ) , and nano-ZrOz particles (A). 60.0 I

I

-20.0 -40.0 -60.0 -80.0

Fig. 1. Transmission electron microscope micrograph. (a) As-synthesized particles prepared from O.05M ZrOClz with diglycol for 72 h. (b) High-resolution transmission electron microscope image with the corresponding power spectrum.

XRD patterns of A1203mixed with 5, 10, 15, and 20 wt% of yttria-stabilized zirconia powders (Aldrich Chemical Company) were identified. The open porosity of sintered bodies was measured by the Archimedes method.

1 '

I

PH

Fig. 3. 5 potential of Polyacrylic acid-adsorbed powders (filled symbols) and nano-ZrOz-coated particles (open symbols): one layer (squares), two layer (triangles), and three layer (diamonds), together with the one of nano-ZrOz particles (filled circle).

is 9.2, while the IEP,H of the as-synthesized nano-Zr02 particles is 6.2. The two kinds of particles were reversely charged at a pH of 6.2-9.2, but the 5 potential was not so high for both the particles at the same pH values. Therefore, the coating process that

Ill. Results and Discussion A TEM micrograph of the as-synthesized nano-Zr02 particles is shown in Fig. I . It can be observed from Fig. 1 (a) that the particles have a monodispersed size of ca. 30 nm. Figure 1 (b) shows a high-resolution transmission electron microscope image with the corresponding power spectrum (the square of the Fourier transform of the TEM image). It was determined that the particles are crystalline in the tetragonal structure without the presence of defaults. The distance between lattice planes that m. was calculated from the power spectrum is about 3 x The faced particle is parallel to the (1 11) plane when crystal lattice unit cell parameters a = 0.512 nm and c = 0.525 nm were selected.t Figure 2 shows the 5 potential of as-synthesized ZrOz particles, as-received A1203,and PAA-adsorbed particles. It can be observed that the isoelectric point (IEP,H) of as-received A1203 'ICPDS file number 17-0923.

464

Fig. 4. Particle size distribution of as-received A1203 (no symbols) and composite powders: one layer coated (squares), two layer coated (triangles), and three layer coated (diamonds).

Progress in Nanotechnology: Processing

resulted from electrostatic attraction was difficult. In view of this, PAA was utilized to modify the AI2O3surface charge. After PAA adsorption, the IEP,H of A1203shifted from 9.2 to 2.7. In the range pH of 2.7-6.2, the surface of nano-Zr02 particles was positively charged while the PAA-adsorbed A1203 surface was negatively charged. Figure 3 shows the 5 potential of PAA-adsorbed and shell-coated composite powders. When the PAAadsorbed AI2O3particles were coated with nano-Zr02 particles, the IEP,! of the composite particles changed from 2.7 to 5.8, which is similar to the one of as-synthesized Zr02 particles. At a pH of 5 or so, PAA-adsorbed particles were negatively charged while nano-Zr02 particles were positively charged. When the two suspensions were mixed, nano-ZrOz particles were deposited

on the A1203 surface because of the electrostatic attraction. Whenever one layer shell coated composite particle was adsorbed with PAA again, the IEP,H of the particles changed to 2.5, which made the subsequent coating of Zr02 particles on the composite powders possible. Figure 4 shows the particle size distribution of the composite particles. The dSOof the as-received AIZO3particles was 154 nm while those of one layer-, two layer-, and three layer-coated composite particles were 236, 307, and 376 nm, respectively. It was found that after coating with one layer of the shell, the particle size increased by about 70 nm, which means that coating thickness was about 35 nm. The coating thickness agreed with the diameter of the synthesized nano-Zr02 particle (Fig. I). The results of the particle size distribution and 6-potential measurements of composite powders indicated that nano-Zr02 particles were adsorbed on the AIZO3surface. Figure 5 shows SEM micrographs of sintered samples formed by three-layer-coated composite powders. From the micrograph of the sintered body, it can be observed that Zr02 particles (bright parts) were dispersed in the A1203matrix (dark parts) in an intergranular manner, situated at grain boundaries or triple points and intragranular, as indeed some ZrO2 particles were found within A1203 grains. When the sintering temperatures were increased to 1500" and 1550"C, no apparent pores were observed. Moreover, when the sintering temperature increased, it was observed that the bright part of Zr02 phase at the triple points of grain boundaries increased, indicating that Zr02 p a ticles were concentrated at triple points of grain boundaries. Besides, a very fine dispersion of Zr02 within the A1203matrix could be observed. It is well-known that the homogeneous dispersion of the second phase in the matrix is advantageous as it improves the composite properties such as mechanical strength.g-' Figure 6 shows the X-ray intensity ratio of reference samples. As a reference for estimation of the composition of the composites, XRD patterns of A1203 mixed with 5, 10, 15, and 20 wt% of Zr02 powders were identified at first. Although A1203 (104) had the highest X-ray intensity, in this area Zr02 ( I 10)also had a high X-ray intensity. Therefore, the intensity ratio of Z r 0 2 (1 I 1) to A1203( I 16) was used as a reference. By comparing the X-ray intensity ratio of Zr02 ( I 1 1) to AI2O3(1 16) of the as-synthesized powders with the curve shown in Fig. 6, it is estimated that three-layer-shell coated composites powders contained 12 wt% of Zr02 while one-layer-shell-coated one contained about 4 wt% Zr02. Therefore, it was indicated that the content of the Z r 0 2 in the composite could be well controlled by coating with different layers. Assuming that alumina has a density of 4.00 g/ cm3 and yttria-stabilized zirconia has a density of 6.08 g/cm3,the

'

2.5

0

Fig. 5. Scanning electron microscopy micrographs of the sintered composites formed by three-layer-nano-Zr02-coated A1203 powders: (a) 1450"C, (b) 1500"C, and (c) 1550°C. Bright parts, ZrOz phase; dark parts; AI2O3phase.

NanocomDosites and Nanostructures

0

5

10

15

20

25

Zr02 contentlwtq, Fig. 6. X-ray intensity ratio of Zr02 ( I I l):A1203(1 16) of the reference

samples. 465

70

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%x

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O 20

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Fig. 7. Open porosity as a function of different sintering temperatures: pure A1203 (triangles), 1-layer-coated composites (diamonds), 3-layercoated composites (squares).

B

theoretical density of the 3-layer-coated powders is 4.17 /cm3, while I-layer-coated powders had a density of 4.05 g/cm . Figure 7 shows the open porosity of composites sintered at different temperatures. Composites that formed by shell-coated powders sintered at 1450°C had quite low open pores (2%). Generally, hot press sintering has been applied in order to produce pore-free ZrO2-Al203 composites. However, it is possible that ZrO2-AI2O3composite with low open-pore structures were produced under an atmosphere of normal pressure using nanoZ r 0 2 particle-coated composite powders.

IV. Conclusions The surface coverage of nano-Zr02 on the surface of A1203 particles was characterized by 5 potential, particle size distribution, and SEM. The content of Zr 0 2 in the coated particles was well controlled by the LBL method. Moreover, we demonstrated that a ZrOrAI2O3 composite with a low open-pore structure can be produced by nano-Zr02coated A1203powders. The structures consisted of dispersion of Z r 0 2 in the composite with a low open pore when the body of the coated powders was sintered at 1450°Cin air. One important result is that the present research offered a processing method of composite material dispersed Zr 0 2 in an AI2O3 matrix using nano-ZrO2-coated A1203particles.

References IN. Claussen, “Fracture Toughness of Alumina with an Unstabilized Z r 0 2 Dispersed Phase,’’ J . Am. Ceram. Soc., 59, 49-51 (1976). *M. Ruhle, A. H. Heuer, and N. Claussen, “Transformation and Microcrack Toughening as Complementary Processes in Zirconia-Toughened Alumina,” J. Am. Ceram. Soc., 69, 195-7 (1986).

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‘D. Casellas, I. Rafols, L. Llanes, and M. Anglada. “Fracture Toughness of Zirconia-Alumina Composites,” J . Ref. Met. Hard Mofer.. 17. 11-20 (1999). 4A. H. De Aza, A. H. Chevalier. G. Fantozzi, M. Schehl. and R. Torrecikas, “Crack Growth Resistance of Alumina, Zirconia, and Zirconia Toughened Alumina Ceramics for Joint Prostheses,” Biomurerials, 23, 93745 (2002). ’R. H. J . Hannink, P. M. Kelly, and B. C. Muddle, “Transformation Toughening in Zirconia-Containing Ceramics,” J . Am. Ceram. Soc., 83, 461-87 (2YO). F. F. Lange, “Powder Processing Science and Technology for Increased Reliability,” J . Am. Ceram. Soc., 72, 3-15 (1989). ’S. Badhuri and S. B. Badhuri, “Enhanced Low Temperature Toughness of AI20,/ZrO2 Nano/Nanocomposites.” Nanosfrucf. Mafer., 8. 75543 (1997). ‘V. Sergo, G. Pezzotti, 0. Sbaizero, and T. Nishida. “Grain Size Influence on Residual Stresses in Alumina/Zirconia Composites,” A d a Mater., 46, 1701-10 (1998). ’W. H. Rhodes, “Aggelomerate and Particle Size Effect on Sintering YttriaStabilized Zirconia,” J . Am. Ceram. Soc., 64, 19-22 (1981). ‘OH. Ferkel and R. J . Hellmig, “Effect of Nanopowders Deagglomeration on the Densities of Nanocrysta\line Ceramic Green Body and their Sintering Behavior.” Nanustruct. Muter.. 11, 617-22 (1999). “T. S. Suzuki, Y. Sakka. and K. Hiraga, “Colloidal Processing for Fine Particles of A1203-15vol% Zr02 System,” J. Jpn. Soc. Pouder Mrtall., 44,3 5 M l (1997). ”T. S. Suzuki, Y. Sakka, K. Nakano, and K. Hiraga, “Effect of Ultrosonication on the Microstructure and Tensile Elongation of Zirconia-Dispersed Alumina Ceramics Prepared by Colloidal Processing.” J . Am. Ceram. Soc., 84, 21324 (2001). ”W.H. Tuan, R. Z. Chen, T. C. Wang, C. H. Cheng. and P. S. Kou. “Mechanical Properties of AI2O3/ZrO2Composites,” J . Eur. Ceram. Sac., 22, 2827-33 (2002). I4S. Deville, J. Chevalier, G. Fantozzi, J. F. Bartolone, J . Requend, J. S. Moya. R. Torrecillas, and L. A. Diaz , “Low-Temperature Ageing of Zirconid-Toughe n d Alumina Ceramics and its Implication in Biomedical Implants,” J. Eur. Ceram. Soc., 23, 2975-82 (2003). ISM. Schehl, L. A. Diaz, and R. Torrecillas, “Alumina Nanocomposites from Powder-Alkoxide Mixtures,” Acta Mafer., 50, I12S-39 (2002). I6D. Enache, M. R. Auberger, K. Esterle, and R. Revel, “Preparation of AI2O3ZrOl Mixed Supports, their Characteristics and Hydrothermal Stability,” Colloid! Surfaces A., 220, 223-33 (2003). ”B. Fegley, P. White, and H. K. Bowen. “Preparation of Zirconia Alumina Powders by Zirconium Alkoxide Hydrolysis,” Commun. Am. Ceram. Soc., 68, Cqg-2 (1985). P. Durin, .I.Tartaj, J. F. Fernadez. M. Villegas, and C. Moure, “Crystallization and Sintering Behaviour of Nanocrystalline Y-TZP Powders Obtained by Seeding-Assisted Chemical Coprecipitation,” Ceram. Inf., 25, 125-35 (1999). ‘9.Ogihara, N. Mizutani, and M. Kato, “Processing of Monodispersed Zr02 Powders,” Ceram. In!., 13, 3 M (1987). 20 S. Bhattacharyya, S . K. Pratihar, R. K. Sinha, R. C. Behera, and R. I. Ganguly, “Preparation of Alumina-High Zirconia Microcomposite by Combined GelPrecipitation,” Mater. Left., 53, 425-3 I (2002). ”S. Kikkawa, A. Kijima, K. Hirota, and 0. Yamamoto, “Crystal Structure of Zirconia Prepared with Alumina by Coprecipitaiton,” J . Am. Ceram. Sac., 85, 721-3 (2002). 22A. Mondal and S. Ram, “Al”Stabilized c-Zr02 Nanoparticles at Low Temperature by Forced Hydrolysis of Dispersed Metal Cations in Water.” Solid Stare. h i e s , 160. 169-81 (2003). *’P. G. Rao, M. Iwasa, T. Tanaka. 1. Kondoh. and T. Inoue. “Preparation and Mechanical Properties of AI2O3-I5wt% ZtOz Composites.” Scripta Mater., 48, 43741 (2003). 24C.M. Wang, “Microstructure Homogenous Improvement in Si3N4by a Powder Coating Method,” J . Mater. Sci.. 31, 4709-18 (1996). ”M. Ohmon and E. Matijevic, “Preparation and Properties of Uniform Coated Inorganic Colloidal Particles: 8. Silicon on Iron.” J. Colloid. Interface Sci.. 160. 192-288 (1993). *‘G. Desai, Z. Xu. and J. A. Lewis. “Synthesis and Properties of CaAI2O4-coated A1203 Microcomposites Powders,” J . Am. Ceram. Soc., 78, 2881-8 (1995).

Progress in Nanotechnology:Processing

Preparation of a Highly Conductive AI2O3/TiN lnterlayer Nanocomposite through Selective Matrix Grain Growth Xihai Jin and Lian Gao*>' Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050, China

An electroconductive Tm/A1203 nanocomposite was prepared by a selective matrix grain growth method, using a powder mixture of submicrosized a-AI2O3, nanosized y-A1203, and TiN nanoparticles synthesized through an in situ nitridation process. During sintering, a self-concentration of TiN nanoparticles at the matrix grain boundary occurred, as a result of the selective growth of large a-Al2O3 matrix grains. Under suitable sintering conditions, a typical interlayer nanostructure with a continuous nanosized TIN interlayer was formed along the A1203 matrix grain boundary, and the electroconducting behavior of the material was significantly improved. Twelve volume percent TiN/ A1203 nanocomposite with such an interlayer nanostructure showed an unprecedentedly low resistivity of 8 x lop3 R - c m , which was more than two orders lower than the TiN/AI2O3 nanocomposite without such an interlayer nanostructure.

I. Introduction

I

recent years, multifunctional ceramic nanocomposites have attracted increasing attention."' In these materials, functional nanodispersoids are incorporated into conventional structural ceramics, giving the nanocomposites high strength and new functionalities, such as magnetic, dielectric, electroconductive and machinable properties, etc. Among various kinds of multifunctional nanocomposites, electroconductive and machinable nanocomposites are especially attractive, because of the great potential of fabricating complex-shaped ceramic component through electrical discharge machining o r lathing. In a machinable nanocomposite, the weak interface between machinable nanodispersoid and matrix grain accounts for the high machinability.' -9 As to electroconductive nanocomposites, the interconnecting network of electroconducting nanodispersoids serves as the current path and contributes to the conductivity."-15 If the functional nanodispersoids are homogeneously dispersed at the matrix grain boundary by forming a continuous interlayer, i.e., forming an "interlayer nanostructure," the electroconductivity or machinability can be significantly improved while keeping the nanodispersoid content low. In addition, the formation of interlayer nanostructure also shows great promise for toughness improvement in a metalkeramic nanocomposite, because of the strong shielding effect of the ductile metal interlayer on crack propagation. Under ideal conditions, even low-temperature ductility can be expected in the metalkeramic interlayer nanocomposite, which is facilitated by the grain-boundary sliding and dislocation movement in the highly ductile nano-metal inerlayer. N

J. Smialek-contrihuting editor

Manuscript No. 20888. Received August 17, 2005; approved October 10,2005. This work was supported by the National Natural Science Foundation of China under Grant number 50402006 and Shanghai Nanotechnology Promotion Center under Grant num,ber 0452nm086. Member. American Ceramic Society. 'Author to whom correspondence should he addressed. e-mail: [email protected]

Nanocornposites and Nanostructures

Except for the multifunctional nanocomposite mentioned above, the formation of an interlayer nanostructure shows great advantage for grain-boundary stress design in conventional structural nanocomposites as well, by utilizing the thermal expansion coefficient difference between the nanodispersoid and the matrix phase. Numerous cases have shown that a compressive thermal stress at the matrix grain boundaries can lead to significant enhancements in the reliability and mechanical properties of ceramic composite^.'^-^' Yet, unfortunately, no successful fabrication of interlayer nanocomposite has been reported in the literature, despite the fact that many great potentials of interlayer nanocomposite might have been recognized. The reason for this mainly lies in the strong tendency of a nanodispersoid to be expelled to matrix triple-point junctions when direct contact between matrix grains forms during sintering. This leads to considerable difficulty in obtaining a continuous nanodispersoid interlayer along the matrix grain boundary, especially at a low nanodispersoid content. In the present paper, a novel method for the fabrication of interlayer nanocomposite was invented, which used the selective matrix grain growth (SMGG) as a means to manipulate the nanodispersoid distribution in the matrix. The A1203/TiN interlayer nanocomposite prepared by this method showed a low electrical resistivity of 8 x f2.cm at a low TIN content of 12 vol%, which was more than two orders lower than the A1203/ TIN nanocomposite without such an interlayer nanostructure.

11. Experimental Procedures

A1203 and TIN mixture powder was prepared by high-temperature nitridation of AI2O3 and Ti02 nanocomposite powder synthesized through heterogeneous precipitation, using a-A1203, AI(N03)3, and Ti(0Buk as raw materials. First, a submicrosized a-A1203 (0.3 pm) powder was ball milled in distilled water at pH 10 for 24 h. Then, designated amounts of A1(N03)3 and T ~ ( O B Umixing )~ ethanol solution and 4 N ammoniated water were added drop by drop into the a-Al2O3 suspension under vigorous stirring. During the whole process, the pH value of the reaction system was maintained at pH 9-10. The precipitate mixture obtained was washed with distilled water for four times, dewatered with absolute ethanol alcohol for two times, and then dried and heated at 100" and 300°C, successively. Nitridation was carried out in a tube furnace (10 cm in diameter) at 900°C for 5 h, using flowing ammonia gas (1.5 Ljmin) as the nitridating agent. Sintering was conducted in vacuum at a pressure of 100 MPa and a sintering temperature of 1200-1500°C for 3-5 min, using a Dr Sinter 1020 SPS apparatus (Sumitomo Coal Mining Co., Ltd., Tokyo, Japan). Electrical resistivity was measured by the fourprobe method. Transmission electron microscopy (TEM, JEOL, JEM-200CX, Tokyo, Japan) was used for powder morphology observation, and x-ray diffraction (XRD, D/max 2550V, Rigaku, Tokyo, Japan) was used for crystalline phase identification. The microstructure of the sintered sample was characterized by scanning electron microscopy (SEM, Instrument JSM 6700F, JEOL, Tokyo, Japan) on mirror-polished surfaces.

467

Fig. 1. Schematic illustration of the microstructure evolution of an interlayer nanocomposite prepared by the selective matrix grain growth method (a) nanocomposite powder compact; (b) nanocomposite fully densified but still with considerable matrix nanoparticles unconsumed by matrix template grains; (c) nanocomposite showing a typical interlayer nanostructure after nearly complete consumption of the matrix nanoparticles; (d) oversintered nanocomposite with a damaged interlayer nanostructure.

111. Results and Discussion The fundamental theory underlying the concept of SMGG can be understood as follows: when a ceramic powder with a bimodal particle size distribution is sintered at high temperature, the large particles will preferentially grow large by consuming the small ones of the same chemical composition due to thermal dynamic reasons.18-20In other words, selective grain growth of large particles occurs in the material during sintering. In a nanocomposite system, such a selective growth of large matrix grains can be used for interlayer nanocomposite preparation. Figure 1 shows the schematic illustration of the preparation of an interlayer nanocomposite by the SMGG method. To facilitate SMGG, a starting powder with a bimodal matrix particle size distribution is used, containing submicro- and nanosized matrix particles, and second-phase “nanodispersoids” (Fig. I(a)). During sintering, the submicrosized matrix particles, which are denoted as “templates,” preferentially grow large by consuming the surrounding matrix nanoparticles; meanwhile, the nanodispersoids are expelled to the inter-template region as the template boundaries approach each other. This results in a self-concentration of the nanodispersoid at the inter-template boundary. When the matrix nanoparticle content in the starting powder is sufficiently high, an intermediate sintering state as shown in Fig. I(b) can be expected, due to the high sinterability of the nanosized particles. Figure I(b) shows that the material has been fully densified but still with a considerable amount of matrix nanoparticles unconsumed by the template grains. Later heating induces further grain growth of the templates through the SMGG process and the continuous concentration of nanodis-

Fig.2. X-ray diffraction patterns of (a) AI20,/TiN powder and (b) AI2O3/TiNsample sintered at 12WC, where , 0, and 0 represent uAl203, TIN, and A12Ti05, respectively. 468

persoids at the inter-template boundary. At this stage, the unconsumed matrix nanoparticles could serve as pinning agents and prevent the nanodispersoids from growing large and moving toward the triple-point junctions of template grains. Under proper sintering conditions, a nanocomposite with a microstructure as shown in Fig. I(c) can be obtained. In this nanocomposite, the nanodispersoids and a few residual matrix nanoparticles are homogeneously distributed in the intertemplate boundaries, forming a continuous nano-interlayer. Here, an important point should be reiterated. That is, the interlayer nanostructure formed through the SMGG process is not very stable and can only be obtained with careful control of the sintering program. If the template grains continuously grow large through inter-template mass transfer that occurs after complete consumption of the matrix nanoparticles by template grains, many second-phase nanodispersoids will be expelled to the template triple-point junctions (Fig. I(d)), resulting in damage of the interlayer nanostructure. To evaluate the feasibility of the method, an electroconductive AI2O3/TiN interlayer nanocomposite was prepared according to the SMGG approach described above, using a powder mixture of 68 vol% submicrosized a-Al203, 12 vol% TIN, and 20 vol% nanosized A1203 that assumably derived from A1(N03)3. Figure 2(a) shows the XRD pattern of the powder mixture. Only reflections characterized to a-A1203 were observed. This indicates that the nanosized A1203 might have transformed to a-A1203 during the nitridating process, when the seeding effect of submicrosized a-AI203 and the long heating time were considered. TIN was in an amorphous state, as indi-

Fig. 3. Transmission electron microscopy image of the A1203/TiN

powder mixture. Progress in Nanotechnology: Processing

Scanning electron microscopy images of the polished surfaces of A1203/TiNnanocomposite sintered at (a) 12WC, (b) 1250°C,and (c) 14Oo"C, respectively. The bright phase was TIN, and the dark-gray phase was AI2O3. Fig. 4.

cated by the broad humps appearing at 20 angles corresponding to TIN peak positions. TEM observation found a bimodal particle size distribution in the powder (Fig. 3), with submicrosized a-AI2O3(-0.3 pm) and nanosized (1&20 nm) a-A1203and TiN particles. The large particle size difference between submicroand nanosized A1203should facilitate selective grain growth of the submicrosized A1203by consuming the nanosized ones during sintering. Density measurement indicated that the powder exhibited a high sinterability and could be sintered to nearly full density above 1200°C by the spark plasma sintering (SPS) technique. Figure 2(b) shows the XRD pattern of the sample sintered at 1200°C. Crystalline a-A1203and TIN were the major phases in the sample, as deduced from their strong XRD reflections. In addition, because of the incomplete nitridation of Ti02 during the powder preparation process or a probable oxidation of nanosized TiN when handling the powder in air, a small amount of AI2TiO5 was also detected in the sample as a result of the chemical reaction between A1203and Ti02 during sintering. Figures 4(a) and (b) show the SEM images of the polished surface of a TiN/AI2O3 nanocomposite sintered at 1200" and 1250"C, respectively. Two phases could be clearly observed on the polished surfaces, according to their difference in contrast. The bright one was TIN, and the dark-gray one was A1203.The TiN nanodispersoids were homogeneously distributed in the A1203 matrix and formed into a continuous interlayer along many A1203grain boundaries. This suggested that a nearly perfect interlayer nanostructure had formed in these samples. Figure 4 (c) shows the SEM images of the polished surfaces of TiN/AI2O3 nanocomposites sintered at 1400°C. Although the samples sintered at 1200" and 1250°C showed a quite similar Nanocomposites and Nanostructures

microstructure, a drastic change in microstructure occurred as the sintering temperature increased to 1400°C. In comparison with the sample sintered at 1200" and 1250"C, TIN dispersoids in the later sample were much larger and more widely distributed in the A1203matrix. A few interconnecting TiN networks could be observed in them. This indicated that a large proportion of TIN nanodispersoids were expelled to the matrix grain triple junctions, when the overgrowth of A1203 matrix grains occurred after the complete consumption of fine A1203particles by the large A1203 templates at a high sintering temperature. And the interlayer structure had broken down. In general, the microstructure evolution process of the TiN/AI2O3 nanocomposite was in good agreement with the SMGG mechanism illustrated in Fig. l . Besides, all the samples showed a pore-free microstructure independent of the sintering temperature. This again implied that these samples had been fully densified. Figure 5 shows the electroconducting behavior of the TiN/ A1203 nanocomposite sintered at different temperatures. The sample sintered at 1200°C showed a low resistivity of 8 x R . cm; then, the resistivity increased slowly with sintering temand 1.5 x R.c m at 1250 and perature to 9.6 x 1300"C, respectively. A further increase in sintering temperature led to a rapid increase in resistivity. For the sample sintered at 1500"C, a moderate resistivity of 4 R . cm was observed, which was more than two orders higher than that of the sample sintered below 1300°C. The considerable change in electroresistivity with sintering temperature can be understood in terms of a strong correlation between the microstructure and electroconductivity. For the sample sintered at low temperature (< 1300"C), a more perfect interlayer nanostructure was formed as shown in Figs. 5(a) and (b), and the interspace between TIN 469

Acknowledgements We acknowledge the helpful discussion with Dr. Zhao Zhe of the Stockholm University with SPS.

References

10-3

1200

1250

1350 I400 Temperature/ “C

1300

1450

1500

Fig. 5. Electroconducting behavior of 12 vol% TiN/A1203nanocomposite sintered at different temperatures.

dispersoids was small. As a consequence, the TIN dispersoids had a high probability of forming an interconnecting current path, which led to an enhanced electroconductivity as expected. With increasing the sintering temperature and A1203over grain growth, the interlayer nanostructure broke down. The resulting increase in the interspace between TIN dispersoids hindered the formation of an interconnecting current path in the sample and caused the degradation of electroconducting property.

IV. Conclusions In summary, SMGG is a promising method for the fabrication of an interlayer nanocomposite. Under suitable sintering conditions, the TiN/A1203 nanocomposite prepared by the SMGG method showed a typical interlayer nanostructure, and highly interconnecting TIN networks that served as current paths were formed in the A1203 grain boundary. For a TiN/AI2O3 interlayer nanocomposite containing 12 vol% TIN, an unprecedentedly low resistivity of 8 x lo-’ Q . cm was observed, which was more than two orders lower than the TiN/AI2O3nanocomposite without such an interlayer nanostructure. What is more important is that this method is not just limited to fabrication of an electroconductive interlayer nanocomposite, and has a broad meaning for the fabrication of other nanocomposites as well.

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‘W.-P. Tai and T. Watanabe, “Preparation and Mechanical Properties of A1203 Reinforced by Submicrometer Co Particles,” J. Mafer. Sci.,33, 5795-810 (1998). 2T. Sekino, T. Nakajima, S. Ueda, and K. Niihara, ”Reduction and Sintering of a Nickel-Dispersed-Alumina Composite and Its Properties,” J. Am. Cprum. Soc., 80; 1139-48 (1997). S.-T. Oh, M. Sando, and K. Niihara, “Processing and Properties of Ni-Co Alloy Dispersed A1203 Nanocomposites.” Scripfa Mafer., 39, 1413-8 (1998). 4S.-T. Oh, M. Sando, and K. Niihara, “Preparation and Properties of Alumina/ Nickel-Cobalt Alloy Nanocomposites.,” J. Am. Ceram. Soc., 81, 3013-5 (1998). ’J.-G. Li, L. a d o , and J. K. Guo, “Mechanical Properties and Electrical Conductivity of TiN-A120, Nanocomposites,” J. Eur. Ceram. Soc., 23, 69-74 (2003). 6L. Gao, J.-G. Li. T. Kusunose, and K. Niihara. “Preparation and Properties of TiN-Si3N4 Composites,” J. Eur. Ceram. Soc., 24, 381-6 (2004). ’T. Kusunose, T. Sekino, Y.-H. Choa, and K. Niihara. “Fabrication’and Microstructure of Silicon NitrideiBoron Nitride Nanocomposites,” J. Am. Cerum. Soc., 85, 2678-88 (2002). ‘K. Niihara T. Kusunose, S . Kohsaka, T. Sekino, and Y.-H. Choa. “MultiFunctional Ceramic Composites through Nanocomposite Technology,” Key Eng. Mafer., 161-163, 527-35 (1999). 9.Kusunose, T. Sekino, Y.-H. Choa. and K. Niihara. “Machinability of Silicon Nitride/Boron Nitride Nanocomposites,” J. Am. Cerum. Soc.. 85. 2689-95 (2002). ‘OX. H. Jin, L. Gao, J. G. Li, and S. Zheng, “Influence of Microstructure Evolution on the Electroconducting Behavior of lntragranular TiNjZTM Nanocomposites,” J. Am. Ceram. Soc., 87, 162-5 (2004). ”A. Bellosi, S . Guicciardi. and A. Tampieri, “Development and Characterization of Electroconductive Si,N4-TiN Composites,” J. Eur. Ceram. Soc., 9.8>93 (1992). 12A.Bellosi and G. D. Portu, “Preparation and Properties of Electroconductive A1203-BasedComposites,” J . Eur. Ceram. Soc., 10, 307-15 (1992). ”K. Tsugeki. T. Kato, Y. Koyanagi, K. Kusakabe. and S. Morooka. ”Electroconductivity of Sintered Bodies of a-AI2Oi-TiN Composite Prepared by CVD Reaction in a Fluidized Bed.” J. M a w . Sci., 28, 3168-72 (1993). I4F. Lux, “Review: Models Proposed to Explain the Electrical Conductivity of Maxtures made of Conductive and Insulating Materials,” J. Muter. Sci.. 28, 285301 (1993). I’M. Tajika, H. Matsubara, and W. Rafaniello, “Microstructures and Properties in Aluminum Nitride-Titanium Nitride Composite Ceramics,” Ma/er. Left., 41, 139-44 (1999). 16X. H. Jin and L. Gao, “Effects of Powder Preparation Method on the Microstructure and Mechanical Performance of ZTA/LaAll, O 1Composites.” ~ J . Eur. Ceram. Soc., 24, 653-9 (2004). I7K. Niihara, “New Design Concept of Structural Ceramics: Ceramic Nanocomposite,” J . Ceram. Soc. Jpn., 99, 974-82 (199 I). I’M. M. Seabdugh, I. H. Kerscht, and G. L. Messing, “Texture Development by Templated Grain Growth in Liquid-PhawSintered u-Alumina,” J. Am. Ceram. Soc., 80, 1181-8 (1997). I9C. Duran, S. Trolier-McKinstry. and G. L. Messing, “Fabrication and Electrical Properties of Textured Sro s3Baod7Nb206Ceramics by Templated Grain Growth,” J. Am. Cerum. Soc., 83, 2203-13 (2000). 20 Y. M. Kan. P. L. Wang, Y. X. Li. Y. B. Cheng, and D. S. Yan, “Fabrication of Textured Bismuth Titanate by Templated Grain Growth using Aqueous Tape Casting,” J . Eur. Ceram. Soc., 23. 2163-9 (2003).

Progress in Nanotechnology:Processing

Preparation and Microstructure of Multi-Wall Carbon NanotubesToughened Al2O3 Composite Jinpeng Fan, Daqing Zhao,+ Minsheng Wu, Zening Xu, and Jun Song Department of Mechanical Engineering, Tsinghua University, Beijing 100084, China Key Laboratory for Advanced Manufacturing by Materials Processing Technology, Tsinghua University, Beijing 100084, China

With multi-wall carbon nanotubes (MWNTs) as reinforcement, a 12 vol% MWNTs/alumina (A1203) ceramic composite was obtained b hot pressing. A fracture toughness of 5.55k0.26 MPa-ml/{ 1.8 times that of pure A1203 ceramics, was achieved. Experimental results showed that the enveloping of carbon nanotubes (CNTs) with sodium dodecyl sulfate (SDS) is effective in changing the hydrophobicity of CNTs to hydrophilicity and improving the dispersion of CNTs in aqueous solution. Enveloped with SDS, CNTs can be homogeneously mixed with A1203 at a microscopic level by heterocoagulation. This mixing method can obviously improve the chemical compatibility between CNTs and A1203, which is important for enhancement of interfacial strength between them.

S

1.

Introduction

INCE

their first discovery by Iijima,' carbon nanotubes (CNTs) have been the focus of novel material research. Exp e r i m e n t ~ ~have , ~ shown that multi-wall carbon nanotubes (MWNTs) have exceptional mechanical properties. Yu et aL2 measured Young's modulus of the outermost layer of MWNTs as varying from 270 to 950 GPa and the tensile strength of this layer as ranging from 11 to 63 GPa. Thus, CNTs have considerable application potential in composite materials. Most of the investigations on CNT-based composites have focused on enhancing the mechanical properties of a polymer matrix composite.46 Only a few of them are related to a ceramic matrix composite, and successful cases are even fewer. Zhan et a/.' synthesized a nanocrystalline single-walled carbon nanotubes/alumina (A1203)composite by the spark-plasma sintering process, and the fracture toughness estimated by the indentation microfracture method was found to be three times higher than pure nanocrystalline AI2O3. MWNT-reinforced A1203 composites have been the subject of several previous investigations. Peigney et synthesized CNT/Fe/AI2O3composite powders in situ, and composite powders were hot pressed to obtain CNT/Fe/AI2O3composites.I0.' The CNTs were dispersed very well in the matrix, but the fracture toughness and flexure strength were both significant1 lower than those of carbon-free Fe/AI2O3 composites. An et a/.72studied the influence of CNT content on the tribological properties of CNT/AI2O3 composites. It was found that when the C N T content is up to 4 wt%, the composite microhardness reaches the maximum, about 30% more than that of pure A1203. Siege1 et a/.I3used MWNTs to toughen A1203matrix, achieving a 24%

'

R. Kerdns-contributing editor

Manuscript No. 20564. Received May 15, 2005; approved September 15, 2005. 'Author to whom correspondence should be addressed. e-mail: zhaodqatsinghua. edu.cn

Nanocomposites and Nanostructures

increase in fracture toughness. Recently, we successfully obtained MWNT/AI2O3composites by hot pressing. We measured the fracture toughness of the samples directly through the single edge-notched beam method. The result showed that the fracture toughness of the composites reached 5.55k0.26 M P a . m1'2 with an 80% increase. Two crucial problems exist in CNTs/Al2O3composite material research. The first one is the dispersion of CNTs in a matrix. The CNTs twist into an aggregate easily because of Van der Waals forces and their fibriform structure. It is difficult to separate and disperse these rope-like aggregates in solution. The second problem is the poor chemical compatibility between CNTs and a ceramic matrix, which leads to a weak interfacial strength. In order to strengthen the matrix, the interfacial strength must be maintained at a certain level that is neither too strong nor too weak. A strong interfacial strength does not allow CNTs to be pulled out from the matrix and fails to strengthen the matrix. But if the interfacial strength is too weak, the CNTs can be easily pulled out from the matrix and will also fail to strengthen the matrix. In this communication, the solutions to the problems and their influence on the microstructure and mechanical properties of composites are discussed.

11. Experimental Procedure

( I ) Sample Preparation Three disk samples of 50 mm diameter were obtained using the following procedures: (i) Eighty grams u-AI2O3 powders (HTAL-01, Nanjing High Technology Nan0 Co. Ltd, China; particle size -200 nm, purity >99.0%) were hot pressed in argon atmosphere at 1500°C at a uniaxial pressure of 20 MPa for 1 h. The sample obtained was marked as O#. (ii) Eighty grams A1203 and 12 vol% (MWNTs versus A1203) MWNTs (Fluidization Laboratory of Tsinghua University) were dispersed in 700 mL de-ionized water with ultrasonic agitation for 30 min. The two suspensions obtained were mixed and dispersed with ultrasonic agitation for 30 min. The mixed suspension was filtrated with water to slurries and then ball milled for 5 h. The slurry was dried in an oven, and the MWNT/ A1203composite powders were obtained. Finally, the composite powders were hot pressed under the same conditions as in step (i). The composite sample obtained was marked as l#. (iii) Eighty grams A1203powders were dispersed in 800 mL 0.5 wt% polyethylene glycol (PEG2000) solution with ultrasonic agitation for 30 min. Twelve volume percent MWNTs (MWNTs versus AI2O3)were dispersed in 700 mL 1 wt% sodium dodecyl sulfate (SDS) solution with ultrasonic agitation for 30 min. The above two suspensions were mixed and dispersed with ultrasonic agitation for 30 min. The mixture was filtered, the slurry obtained was ball milled for 5 h, and then dried in an oven. Finally, the MWNT/AI2O3 composite powders were hot pressed under

471

the same conditions as in step (i), and the new sample was obtained and marked as 2#.

( 2 ) Measurement and Characterization The flexure strength was measured by the three-point bending method, in which the size of the test specimen was 36 mm x 4 mm x 3 mm. The bending span and the load speed for the mmjs, respectively. The strength test was 30 mm and 8.3 x fracture toughness was measured by the single edge-notched beam method, in which the test specimen was 4 mm (wide) x 6 mm (thick) x 30 mm (long). A notch with a depth of about 2.5 mm was cut in the center part of the test specimen. A bending span of 24 mm and a load speed of 8.3 x mmjs was applied for the toughness test. For each sample four specimens were fabricated and tested, and the mean values and standard deviations were obtained. The densities of all samples were obtained using the Archimedes principle. The theoretical densities of the samples were estimated according to the volume fraction of MWNTs and A1203.The relative density is the ratio of the measured density to the theoretical density. The microstructures of all samples after rupture were observed with SEM (JSM-6301F, JEOL Co., Ltd, Tokyo, Japan). 111. Results and Discussion

( 1 ) Mechanical Properties The mechanical properties of all samples are presented in Fig. 1. Compared with 0#, the fracture toughness of I# increased slightly by 6%, but the flexure strength decreased sharply by 32%. The fracture toughness of 2# reached 5.55k0.26 MPa. m1I2, increasing by 80% in comparison with O#; the flexure strength was 314.4k22.4 MPa with a slight decrease by 4%, owing to a few small aggregates of CNTs in the matrix. Dispersion of CNTs in A1203 Matrix The agglomeration of raw MWNTs was considerable, and the aggregate size was generally -20 pm, and the agglomeration of CNTs in composite I# was also considerable (Fig. 2). Another macroscopic proof for poor dispersion of CNTs in I # was the density of the composite. CNT aggregates will reduce the relative density of the composite because of their porous rope-like structure. As shown in Table I, the relative density of the sample containing CNTs (I#) is 91.7%, whereas that of the CNT-free sample (O#) is 97.2%. The dramatic decrease of relative density in I# suggests that many CNT aggregates remain in the matrix. SDS is a useful dispersion agent for CNTs.14 To prove its dispersion effect, the solubility of CNTs in SDS solution and deionized water was measured. 0.5 g raw MWNTs was dispersed in 50 mL de-ionized water with and without 1 wt% SDS by ultrasonic agitation for 30 min, respectively. After centrifugal sedimentation at 4000 rpm for 30 min, the large aggregates in solution settled at the bottom, while dissolved CNTs remained

(2)

Severe agglomeration of MWNTs in fracture surface of I#: (a) aggregate pull out from matrix and (b) residual part of aggregate when the other part was pulled out.

Fig.2.

in the supernatant. The sediment and the supernatant were separated, and the latter was dried in an oven at 140°C, and the dry CNTs were weighed to calculate the solubility. The supernatant of the CNTs solution without SDS became transparent after centrifugal separation, which means no CNTs had dissolved in de-ionized water, while the supernatant of the CNTs solution with SDS remained dark and the solubility reached 1.1 mg/mL. So, with the addition of SDS, the dispersion of CNTs in the matrix had improved noticeably. Figure 3 shows that the largest aggregate size of CNTs in the composite had reduced drastically to about 2 pm. The relative density of 2# had also increased to 96.4% because of an improvement in dispersion. The improvement in dispersion was attributed to the following factors: ( I ) the dispersion capability of water had improved because the surface tension of water decreases when SDS is added; (2) when CNTs are enveloped by SDS molecules, the increased steric hindrance between CNTs can counteract the Van der Waals force; and (3) the surface of CNTs enveloped by SDS molecules can trap more anions in solution, which will reduce CNT aggregation because of electrostatic repulsion.

Table I.

Densities and Relative Densities of the Samples

~~

Sample

Fig. 1. The flexure strength and fracture toughness of samples.

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O# 1# 2#

Density (sicm')

Theoretic density (gicm')

Relative density (%)

3.88 3.36 3.51

3.97 3.68 3.68

97.1 91.3 95.4

Progress in Nanotechnology: Processing

Fig.3. Well-dispersed MWNTs in composite 2# and some MWNTs that join two Al2O3grains (marked by arrow).

The agglomeration of CNTs considerably influenced the mechanical properties of composite I#. When the load reaches a certain value, the aggregates in the composite break and separate into two parts. One part is pulled out from the aggregate (marked by an arrow in Fig. 2(a)), while the other remains in the matrix (area A in Fig. 2(b)). Because of its poor load-carrying ability, the function of this kind of aggregate in the matrix is similar to pores with the same size as aggregates. Consequently, the existence of aggregates may decrease the flexure strength drastically. However, when aggregates are broken and pulled out, they have to overcome the interaction force between CNTs. This process may consume some fracture energy and may slightly improve the strength of the composite. For composite 2#, CNTs were well dispersed in the matrix. Some CNTs that were embedded into the matrix can be clearly observed from Fig. 3 (marked by an arrow). When the composite ruptures, these embedded CNTs may strengthen the matrix through bridging of CNTs or a pull out mechanism.

(3) Interface Between CNTs and Matrix To enhance the interfacial strength between CNTs and the matrix properly, the chemical compatibility between them must be improved. An aqueous medium was selected to disperse CNTs in the matrix. It is well known that A1203is hydrophilic, while CNTs are hydrophobic. First of all, when CNTs are enveloped

Fig. 5.

The pull out of MWNTs on the fracture surface of 2#.

by SDS molecules, the sulfonic group of SDS makes CNTs hydrophilic. In the following process, heterocoagulation was used to mix CNTs and the A1203powder, which can effectively improve their chemical compatibility. A1203particles trap cations in aqueous solution, while CNTs trap anions when enveloped by SDS molecules. Zeta potential measurements showed that the zeta potential value of CNTs dispersed in the SDS solution was -54.3 mV, while that of A1203 in the PEG solution was 28.4 mV. When CNTs and A1203 solution are mixed, the CNTs become coated with small A1203 particles by electrostatic attraction. Simultaneously, Al2O3 particles of larger sizes become coated with individual CNTs or small CNT ropes. Therefore, CNTs and AI2O3particles can be mixed homogeneously at the microscopic level. Composite 2# showed a different microstructure (Fig. 3-Fig. 5 ) compared with composite I#: area A in Fig. 2(b) represents a typical pristine porous aggregate of CNTs, whereas the pores between tubes in Fig. 4 are filled with A1203, and there are no obvious interspaces between CNTs and the matrix, and so the interfacial strength is enhanced. This kind of microstructure might be formed during a sintering process when the CNTs and matrix are combined by sintering of adherent A1203particles. The interface strength between CNTs and A1203matrix was enhanced because of their improved chemical compatibility. So, when the composite ruptured, load could be effectively transferred from the matrix to CNTs through the interface between them. Thus, CNTs can bridge the cracks, contributing to the improvement in fracture toughness. In Fig. 5, some CNTs pulled out from the matrix can be clearly observed.

IV. Conclusions By adding 12 vol% MWNTs as reinforcement, a MWNTs/ AI2O3 composite was obtained by hot-pressed sintering. The fracture toughness increased to 5.55k0.26 MPa. m”2, 1.8 times that of pure A1203,and the flexure strength varied slightly. The SDS is effective in appropriately changing the surface property of CNTs from hydrophobicity to hydrophilicity, which is important to improve the dispersion of CNTs in aqueous solution. In addition, the chemical compatibility between CNTs and A1203can be obviously improved by the heterocoagulation mixing method, which contributes toward obtaining a well-bonding interface between CNTs and A1203. Acknowledgments Fig.4. Well-combined interfaces between MWNTs and matrix of 2#: (A) carbon nanotubes/AlzO3 compound body and (B) A1203 grain.

Nanocomposites and Nanostructures

The authors would like to thank Drs Qingyu Shi and Hongwei Hao for their helpful discussions. The authors also thank the Fluidization Laboratory of Tsinghua University for kindly providing CNTs.

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References ‘S. lijima, “Helical Microtubules of Graphitic Carbon.” Nature (London), 354, 5 6 8 (1991). *M. F. Yu, 0. Lourie, M. J. Dyer, K. Moloni, T. F. Kelly, and R. S . Ruoff, “Strength and Breaking Mechanism of Multiwalled Carbon Nanotubes Under Tensile Load,” Science (Washington, D C ) , 287, 63740 (2000). ‘M. R. Falvo, G . J. Clary, R. M. Taylor, V. Chi, F. P. Jr. Brooks, S. Washburn. and R. Superfine, “Bending and Buckling of Carbon Nanotubes Under Large Strain,” Nature (London), 389,5 8 2 4 (1997). 4L. S. Schadler, S . C. Giannaris, and P. M. Ajayan, “Load Transfer in Carbon Nanotube Epoxy Composites,” App/. Phys. Leu., 73 [26] 38424 (1998). ’H. D. Wagner, 0. Lourie, Y. Feldman, and R. Tenne, “Stress-Induced Fragmentation of Multiwall Carbon Nanotubes in a Polymer Matrix,” Appl. Phys. Lett.. 72 [2] 188-90 (1998). 60.Lourie and H. D. Wanger, “Evidence of Stress Transfer and Formation of Fracture Clusters in Carbon Nanotube-Based Composites,” Compos. Sci. Technol., 59 [6] 975-7 (1999). ’G. D. Zhan, J. D. Kuntz, J. Wan, and A. K. Mukherjee, “Single-Wall Carbon Nanotubes as Attractive Toughening Agents in Alumina-Based Nanocomposites,” Nut. Murer., 2 [I]3 8 4 2 (2003).

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‘A. Peigney. Ch. Laurent, F. Dobigeon, and A. Rousset, “Carbon Nanotubes Grown In-Situ by a Novel Catalytic Method,” J . Muter. Res., 12 [3] 613-5 (1997). ’A. Peigney, Ch. Laurent. 0. Dumortier, and A. Rousset, “Carbon NanotubesFe-Alumina Nanowmposites. Part 1: Influence of the Fe Content on the Synthesis of Powders,” J. Eur. Cerum. Sac.. 18 [I41 1995-2004 (1998). “Ch. Laurent, A. Peigney, 0. Dumortier, and A. Rousset. “Carbon Nanotubes-F‘e-Alumina Nanocomposites. Part II: Microstructure and Mechanical Properties of the Hot-Pressed Composites,” J . Eur. Cerum. Soc.. 18 [I41 200513 (1998). “A. Peigney. Ch. Laurent, E. Flahaut. and A. Rousset, “Carbon Nanotubes in Novel Ceramic Matrix Nanocomposites,” Cerum. Int., 26 [6] 677-83 (2000). 12J. W. An, D. H. You, and D. S. Lim, ”Tribological Properties of Hot-Pressed Alumina-CNT Composites,” Wear, 255, 6 7 7 4 1 (2003). I3R. W. Siegel, S. K. Chang. B. J. Ash. J. Stone. P. M. Ajayan, R. W. Doremus, and L. S. Schadler, “Mechanical Behavior of Polymer and Ceramic Matrix Nanocomposites,” Scripru Muter., 44, 20614 (2001). 14B. Vigolo, A. Pe’nicaud. C. Coulon, C. Sauder, R. Pailler, C. Journet, and P. Bernier, “Macroscopic Fibers and Ribbons of Oriented Carbon Nanotubes,” Science (Washington, D C ) , 290, 13314 (2000).

Progress in Nanotechnology:Processing

Three-Dimensional Assemblies of Zirconia Nanocrystals Via Shape-Preserving Reactive Conversion of Diatom Microshells Samuel Shian,* Ye Cai, Michael R. Weatherspoon, Shawn M. Allan, and Kenneth H. Sandhage**>+ School of Materials Science and Engineering, Georgia Institute of Technology, Atlanta, Georgia 30332

The synthesis of three-dimensional (3-D) assemblies of zirconia nanocrystals via the shape-preserving reactive conversion of biologically reproducible, silica-based microtemplates (diatom microshells) is demonstrated for the first time. Silica diatom microshells were first converted into magnesia replicas via an oxidation-reduction displacement reaction with magnesium gas. The magnesia replicas were then converted into zirconia replicas via a metathetic displacement reaction with zirconium tetrachloride gas. By utilizing both types of shape-preserving displacement reactions, bioclastic silica microtemplates may be converted into functional nanocrystalline assemblies with a variety of complex, but well-controlled 3-D shapes and chemistries for catalytic/chemical, biological, electrical, thermal, mechanical, and other applications.

I.

T

Introduction

attractive chemical, electrical, thermal, optical, and mechanical properties of zirconia-based ceramics have led to the use of these materials in a wide range of applications, including gas sensors and gas separators, fuel cells, catalysts and catalyst supports, thermal protection coatings, refractories, optical filters, cutting tools and milling media, biomedical implants, and high-temperature electrodes and heating elements.'-' Microscale zirconia-based structures with intricate three-dimensional (3-D) shapes and fine (submicrometer) features would be very attractive for a number of these and other application^.^^'" Ceramic microfabrication processes based on the buildup of 2-D patterned layers (e.g., micro-stereolithography, ink jet printing, co-firing of stacked ceramic tapes) have been used to generate 3-D structures with feature sizes on the order of 100 pm or larger.19-23However, the direct assembly of complex-shaped 3-D zirconia microstructures with finer features in large quantities at low cost remains a significant challenge. Numerous examples of 3-D assembly of mineralized structures can be found in Certain micro-organisms are adept at synthesizing oxide microstructures with complex shapes and with fine (submicrometer) features. A spectacular range of intricate 3-D structures are generated by a uatic micro-algae known as diatoms (order B n c i N a r ~ ~ ~ h L . Diatoms c e ~ e ~ are ~~~ single-celled eukaryotic micro-organisms that form rigid cell walls (frustules) comprised of amorphous ~ i l i c a . ~Each ' - ~ ~diatom species assembles a frustule with a particular 3-D shape that is decorated with specific patterns of fine features (e.g., pores, ridges, protuberances with characteristic dimensions in the range HE

'

''

M. Firestone-contributing editor

Manuscript No. 20753. Received July 7. 2005; approved August 16, 2005. Supported by the Air Force Office of Scientific Research (Hugh De Long, Joan Fuller, Program Managers). 'Member, American Ceramic Society. "Fellow, American Ceramic Society. 'Author to whom correspondence should be addressed e-mail sandhage@mu. gatech.edu

Nanocomposites and Nanostructures

of tens to hundreds of nanometer^).^'-^^ A given frustule morphology is replicated with a high degree of fidelity upon diatom reproduction. As a result, sustained culturing of a particular diatom species can yield enormous numbers of cells with similarly shaped frustules (e.g., more than one trillion copies of a arent cell could be generated after 40 reproduction c y ~ l e s ) . ~ ~ * ~ ~ S u c h massively parallel and genetically precise (species-specific) assembly of complex 3-D microstructures under ambient (marine o r freshwater) conditions greatly exceeds the capabilities of current synthetic (man-made) microfabrication processes. The estimated 104-103 extant diatom species provide a diverse selection of frustule morphologies for use in various applicat i o n ~ . ~ ' , The ~ ~ , range ~ ' of potential applications would be significantly enhanced if such frustules could be converted into other compositions without a loss of the frustule shape. Recent work has demonstrated that zirconia frustule replicas may be generated by coating silica frustules with a thin, sol-gel-derived film and then removing the underlying silica through selective d i s ~ o l u t i o n While . ~ ~ such a sol-gel approach may be used to produce frustule replicas with a variety of chemistries, alkoxidebased precursors tend to be relatively expensive. The purpose of this paper is to demonstrate that 3-D silica-based diatom frustules may be converted into nanocrystalline zirconia replicas through a series of shape-preserving gas/solid displacement reactions with lower-cost reactants.

11. Experimental Procedure

Experiments aimed at the direct conversion of silica-based diatom frustules into zirconium-bearing replicas were conducted by reacting the frustules with zirconium tetrafluoride vapor. Cylindrical Aulucoseira diatom frustules (obtained in the form of diatomaceous earth) and ZrF4 powder (99Oh pure, Advanced Research Chemicals Inc., Catoosa, OK) were placed in two separate nickel crucibles. The molar ZrF4:Si02 ratio within the crucibles was varied from 0.36:l to 3.6:l in different experiments. The crucibles were loaded inside a nickel alloy tube (5.1 cm diameter, 20.3 cm long, Inconel Alloy 600, C G Metals, Duluth, GA) that was then welded shut within an argon atmosphere glovebox. The sealed ampoules were then heated at 5'C/ min to a fixed temperature in the range of 250"-800"C and held at this temperature for 2 h. After cooling to room temperature, the sealed tubes were cut open and the reacted specimens were removed. A series of displacement reactions was also examined for converting silica-based frustules into zirconia replicas. An oxidation-reduction reaction with magnesium gas was first used to convert the silica frustules into magnesia-based replicas. A metathetic reaction with zirconium tetrachloride gas was then used to convert the magnesia replicas into zirconia. Magnesia-based replicas of Aulacoseira frustules were generated by reaction with Mg(g) at 900°C for 1.5 h in a manner similar to that described The MgO-based replicas were then exposed to a 0.49M NaOH solution at 60°C for 3 h to selectively dissolve residual silicon from the magnesia. MgO diatom

475

replicas and anhydrous ZrC14 powder (99.9% pure, Advanced Research Chemicals Inc.) were loaded into two separate nickel crucibles. The molar ratio of ZrC14:Mg0 placed within the crucibles was chosen to be 0.52: 1,0.86: 1, or 1.2:1. The two crucibles were then placed inside a nickel alloy tube (5.1 cm diameter, 20.3 cm long, Inconel Alloy 600, C G Metals) that was welded shut within an argon atmosphere glovebox. The sealed ampoules were then heated at 5"C/min to 650°C and held at this temperature for 2 h. After cooling to room temperature, the sealed tubes were cut open and the reacted diatom frustules were withdrawn. The reacted frustules were then immersed and stirred (100 rpm) for I h in de-ionized water heated to 90°C. The frustules were then removed from the water using vacuumassisted filtration. 111.

Results and Discussion

Initial experiments were conducted to directly convert cylindrical Aulucoseira diatom frustules, of the type shown in Fig. I(a), into zirconium oxyfluoride replicas via a metathetic reaction with ZrF4 gas (in a manner analogous to the conversion of diatom frustules into titanium oxyfl~oride~'). For reaction temperatures at or below 350°C, no reaction was detected within 2 h. At higher temperatures, and for all starting molar ZrF4:Si02 ratios examined, the silica frustules were found to disintegrate. A partially disintegrated frustule is shown in Fig. I(b). Relatively large, faceted crystallites of the type shown in Fig. I(c) were observed to form away from the surfaces of these disintegrating frustules. Energy-dispersive X-ray (EDX) analyses (Fig. I(d)) indicated that these crystallites contained Zr, F, and 0, which was consistent with a zirconium oxyfluoride compound.42 The frustule disintegration, and concurrent formation of Zr-O-F-bearing crystals away from the frustule surfaces, suggested that the silica had reacted to form an oxygen-bearing vapor species (e.g., Si20F6(g),or SiOF2(g)) that migrated away from the frustule surface and then reacted further with ZrF4(g) to form the Zr-&F-bearing crystals. A similar type of reactive frustule disintegration, and faceted T i - GF crystal formation away from the frustules, was observed upon exposure of diatom silica frustules to TiF4(g) at temperatures in the range of 500"700°C and at TiF4:Si02 molar ratios in excess of 4.9:1?' Although the exact nature of such reactive frustule evaporation in the presence of ZrF4(g) was not evaluated, the occurrence of such disintegration under all of the reaction conditions examined indicated that a different reaction path was needed in order to generate zirconia frustule replicas. An alternative approach was adopted that avoided the direct exposure of silica frustules to a zirconium halide reactant. A series of displacement reactions was used to convert the silica frustules first into magnesia-based replicas (via reaction with magnesium gas) and then into zirconia replicas (via reaction of the magnesia with zirconium tetrachloride gas). Prior work has shown that diatom frustules can be converted into magnesiabased replicas via the following net oxidation-reduction displacement reaction3'"? SiO2 + 2Mg(g) -+ 2MgO

+ {Si}

(1)

where {Si} refers to silicon dissolved within a Mg-Si liquid (generated by the continued reaction of elemental silicon with excess magnesium vapor) that tended to migrate away from the MgO-based frustule replica.3934Magnesia replicas generated by reaction (I) may then be exposed to zirconium tetrachloride gas to form zirconia by the following metathetic displacement reaction: 2Mg0

+ ZrC&(g)

+

ZrOz

+ 2MgC12

(2)

This net reaction is thermodynamically favored. The standard Gibbs-free ener y change per mole of reaction (2) at 650°C is -75.0 k.J/mol?fM The equilibrium partial pressure of ZrC& 476

Fig. 1. Secondary electron images of: (a) a starting, cylinder-shaped Aulucoseira diatom frustule; (b) a partially disintegrated frustule after exposure to ZrF&) for 2 h at 600°C (molar ZrF4:Si02 reactant ratio = 0.36: 1); and (c) faceted Zr-O-F-bearing crystals formed near the partially disintegrated frustule shown in (b). (d) Energy-dispersiveX-ray analysis of a crystal obtained at the position X in (c).

vapor required for reaction (2) to proceed spontaneously at 650°C (assuming unit activities for pure condensed phase reacTorr, which is much tants and products) is only 4.3 x smaller than the pressure of ZrC4 vapor generated from condensed ZrC4 at this temperature (the sublimation temperature of ZrCb is only 336"C).43.4 Upon cooling to room temperature, the MgClz product of this reaction may be selectively removed from the Zr02 product by dissolution in water. Indeed, zirconium chloride was preferred over zirconium fluoride as a gaseous reactant because the solubility of magnesium chloride in water is much higher than for magnesium flu0ride.4~ A secondary electron image of a magnesia replica of an Aulacoseira frustule is shown in Fig. 2(a). The shape of the starting Si02-based frustule (Fig. l(a)) was preserved after conversion into MgO. Secondary electron images obtained after reaction of the magnesia with zirconium chloride gas at 650°C for 2 h are shown in Figs. 2(b) and (c). EDX analysis obtained at the position X in Fig. 2(b) is shown in Fig. 2(d). The reacted frustules were embedded in a matrix phase that was Mg and CI rich, which was consistent with a solidified MgC12-rich liquid (note: while the melting point of MgC12 is 714"C, a eutectic liquid can form in the MgC12-ZrC4 system at only 426"C46). The Mg-CI-rich matrix phase could be selectively removed from the reacted frustules by dissolution in water. Secondary electron images of isolated frustules exposed after this dissolution treatment are shown in Figs. 3(aHc). EDX analysis of such an isolated frustule is shown in Fig. 3(d). The absence of a magnesium peak (at 1.25 keV) in the EDX pattern indicated that the frustules had been fully converted into zirconia within 2 h at 650°C. The cylindrical frustule shape, and the protruding rim and open hole at the end of the frustule, were preserved after conversion into zirconia. X-ray diffraction (XRD) analyses obtained from the zirconiaconverted frustules are shown in Fig. 4. Diffraction peaks for the monoclinic and tetragonal zirconia polymorphs were detected, whereas diffraction peaks for magnesium oxide were absent Progress in Nanotechnology: Processing

Fig.2. Secondary electron images oE (a) a MgO-based frustule replica after exposure to Mg(g) at 900°C for 1.5 h and then immersion in a 0.49M NaOH solution at 60°C for 3 h; (b), (c) the products of reaction between MgO diatom replicas and ZrCl,(g) at 650°C for 2 h (molar ZrC1,:MgO reactant ratio = 0.861); and (d) energy-dispersive X-ray analysis obtained at the position X in (b).

(Powder Diffraction File, Card No. 45-946 for MgO, No. 371484 for ZrOz (monoclinic), No. 50-1089 for ZrOz (tetragonal), International Center on Diffraction Data, Newtown Square, PA). As the molar ZrC14:Mg0 reactant ratio was increased from 0.52:l (near the stoichiometry of reaction (2)) to 1.2:1 (a significant excess of ZrCI4), the relative amount of monoclinic zirconia generated upon reaction at 650°C for 2 h increased at the expense of tetragonal zirconia. Transmission electron micrographs of cross-sections of zirconia-converted frustules generated with varied molar ZrC1,:MgO reactant ratios are shown in Figs. 5(a)4c). Corresponding selected area electron diffraction (SAED) patterns obtained from these cross-sections are presented in Figs. 5(dHf). Consistent with XRD analyses, the SAED analyses indicated an increase in the amount of monoclinic zirconia relative to tetragonal zirconia with an increase in the starting molar ZrC4:MgO reactant ratio. The crystallite size distribution was also observed to change with an increase in this reactant ratio. The zirconia crystallites formed with a ZrCh:MgO molar ratio of 0.52:l (Fig. 5(a)) were relatively small (-40 nm) and uniform in size. However, for a ZrCl4:MgO molar ratio of 0.86:l (Fig. 5(b)), the zirconia crystallites exhibited a bimodal size distribution; that is, a population of considerably larger crystallites (several hundreds of nanometers in size) was detected along with the finer crystals. This bimodal crystallite size distribution was consistent with XRD analyses of the average tetragonal and monoclinic zirconia grain sizes. The widths of major, non-overlapped diffraction peaks for monoclinic zirconia (the (- 11 1) and (1 11) peaks near 28.2" and 3 IS", respectively, in Fig. 4) and tetragonal zirconia (the (011) peak near 30.3" in Fig. 4) were measured at half of the maximum intensity values and these values were inserted into the Scherrer equation4' to calculate the average crystallite sizes for the tetragonal and monoclinic zirconia phases. For specimens with a ZrC4:MgO reactant ratio of 0.86: I , the average crystallite sizes for tetragonal and monoclinic zirconia were 48 and 380 nm, respectively. A similar bimodal crystallite size distribution was detected for a molar ZrC4:MgO Nanocornpositesand Nanostructures

Fig.3. (a, b) Secondary electron (SE) images of nanocrystalline ZrOz replicas of Aulacoseira diatom frustules generated by the reaction of MgO frustule replicas with ZrC14(g) at 650°C for 2 h (molar ZrC4:Mg0 reactant ratio = 0.52:l) and then immersion in de-ionized water at 90°C for 1 h. (c) Higher magnification SE image of the region indicated by the rectangle shown in (b). (d) Energy-dispersiveX-ray analysis of a nanocrystalline ZrQ2 replica.

reactant ratio of 1.2:l (Fig. 5(c)), although a higher concentration of the larger crystals was detected. The enhanced formation of relatively large zirconia crystallites, and of monoclinic zirconia, at 650°C with an increase in the amount of the ZrC14 reactant was consistent with prior reports on the effect of chlorine on the tetragonal-to-monoclinic zirconia transformation at et a/?' reported temperatures below I O O O " C . ~Karapetrova *,~~ that zirconia specimens with relatively high initial chlorine contents (up to several percent, synthesized from oxychloride precursors) exhibited rapid grain growth between 600" and 8Oo"C, and transformed more rapidly to monoclinic zirconia over this temperature range, than specimens containing little initial chlorine (on the order of 200 pm, synthesized from nitrate precursors)?' Pasquevich et al.4 9found that the growth rate of isolated crystallites, and the tetragonal-to-monoclinic transformation,

Fig.4. X-ray diffraction (XRD) analyses of the Zr02 frustule replicas generated by the reaction of MgO-based replicas with ZrC4(g) at 650°C for 2 hand then immersion in de-ionized water at 90°C for 1 h. The three XRD patterns are associated with different molar ZrC1,:MgO reactant ratios.

477

ceramic compounds, with well-controlled 3-D morphologies for a variety of catalytic/chemical, biological, electrical, thermal, mechanical, or other applications.

IV. Conclusions The conversion of silica-based diatom frustules into nanocrystalline zirconia replicas, via the use of gas/sohd displacement reactions, has been demonstrated for the first time. Initial experiments involving the reaction of the silica frustules with zirconium tetrafluoride gas resulted in disintegration of the microshells by a reactive evaporation process. Such reactive disintegration was avoided by converting the silica frustules into magnesia replicas prior to reaction with a zirconium halide gas (thereby avoiding the direct exposure of the silica frustules to a zirconium halide gas reactant). Magnesia-based replicas were generated via an oxidation-reduction displacement reaction between cylinder-shaped Aulucoseiru diatom frustules and magnesium gas at 900°C. A metathetic displacement reaction between the magnesia replicas and zirconium tetrachloride gas at 650°C then yielded zirconia replicas embedded in a magnesium chloride-rich matrix. Selective dissolution of the magnesium chloride in water yielded freestanding, nanocrystalline assemblies of zirconia that retained the starting cylindrical Aulucoseiru frustule shape. This serial displacement reaction process may be used to generate microscale zirconia assemblies with a wide variety of other controlled 3-D shapes available among the tens of thousands of species of extant diatoms and other silica-forming microorganisms. By utilizing self-replicating microshells as templates, large numbers of nano-crystalline zirconia microassemblies with similar, specific 3-D shapes may be produced for use as controlled-shape powders, reinforcement phases, or microdevice components for a host of applications.

Fig. 5. Transmission electron images of ion-milled cross-sections of ZrOz frustule replicas generated by the reaction of MgO replicas with ZrCls(g) at 650°C for 2 h, using molar 2rCb:MgO reactant ratios of: (a) 0.52: 1, (b) 0.86: I, and (c) I .2: I , and then immersion in deionized water at 90°C for 1 h. (dHf) Corresponding selected area electron diffraction patterns obtained from the specimens shown in (aHc), respectively (f and m refer to the tetragonal and monoclinic polymorphs, respectively, of zirconia).

occurred at faster rates in Ar-CI2 gas mixtures (with a chlorine partial pressure of 300 Torr) than in pure argon at 940°C!9 The present work demonstrates, for the first time, that a series of gas/solid displacement reactions may be used to exchange the silica in biologically reproducible, 3-D microtemplates with a first oxide and then a second oxide, without a loss of the original microtemplate shape. This general approach is not limited to the synthesis of cylinder-shaped assemblies of zirconia nanocrystals. Naturally occurring biosilica microstructures with a wide variety of 3-D shapes are generated by diatoms, radiolaria, sponges, plants, mollusks, echinoderms, arthropods, and other silica-forming organisms.2636 Future advances in genetic engineering may yield an even greater variety of biosilica morphologies. Indeed, recent advances in the sequencing of a diatom genome, and in the development of transformation-based approaches for introducing and expressing genes in diatoms, are the first steps toward the genetic engineering of shape-tailored diatom micros hell^.^^^^' Large-scale culturing (of the type currently used in aquaculture operations) of naturally occurring or genetically modified diatoms (or other organisms) could be used to produce enormous numbers of biosilica microtemplates with targeted, similar 3-D shapes.35 Shape-preserving oxidationreduction and/or metathetic displacement reactions of the type discussed in this work could then be used to convert such biologically self-assembled silica microtemplates into functional microparticles, comprised of nanocrystalline zirconia or other 478

References ‘I. Birkhy and R. Stevens, “Applications of Zirconia Ceramics,” Key Eng. Muter., 122-124, 527-51 (1996). ’J. W. Schwank and M. DiBattista, “Oxygen Sensors: Materials and Applications,” M R S Bull., 24 [6] 4 4 8 (1999). ‘D. J. M. Burkhard. B. Hanson, and G. C. Ulner, “Zr02 Oxygen Sensors: An Evaluation of Behavior at Temperatures as Low as 300”C,” Solid State lonics, 47, 169-75 (1991). 4N. Q. Minh. “Ceramic Fuel Cells,” J. Am. Ceram. Soc., 76 [3] 563-88 (1993). 5H. Armendariz, B. Coq. D. Tichit, R. Dutartre. and F. Figueras, “Influences of Some Synthesis Parameters and Activation Procedures on the One-Step SoCGel Synthesis of Sulfated-Zirconia Catalysts, Followed by TG-DSC and Mass Spectrometry,” J. Catal., 173 [2] 345-54 (1998). 6M. Peters, C. Leyens, U. Schulz. and W. A. Kaysser, “EB-PVD Thermal Barrier Coatings for Aero-Engines and Gas Turbines,” Adv. Eng. Muter., 3 [4] 193204 (2001). ‘A. M. Anthony, “High Temperature Refractory Applications of Zirconia”; pp. 437-54 in Advances in Ceramics. Vol. 3. Science Tecknolugy OJ Zirconiu, Edited by A. H. Heuer and L. W. Hobhs. American Ceramic Society, Westerville, OH, 1981. *Q.Zhang X. Li, J. Shen. G. Wu, J. Wang. and L. Chen. “Zr02 Thin Films and ZrOz/Si02 Optical Reflection Filters Deposited by Sol-Gel Method,” Mater. Lett., 45 161 3 1 1 4 (2000). 9R. C. Garvie, “Structural Applications of Zirconia-Bearing Materials”; pp. 465-79 in Advances in Ceramics. Vol. 12, Science Tecknolugy ofzirconin, Edited by C. M. Riihle and A. H. Heuer. American Ceramic Society, Westerville. OH, 1984. ‘OK. Demhinski. L. Dupont, and A. M. Anthony, “Heating Elements of Composite Zirconia for High Temperature Furnace ( T z 1400°C) in Oxidizing Atmosphere,” Sci. Ceranz., 9, 375-81 (1977). “S. A. Catledge, M. Cook, Y. K. Vohra, E. M. Santos. M. D. McClenny, and K. D. Moore, “Surface Crystalline Phases and Nanoindentation Hardness of Explanted Zirconia Femoral Heads,” J. Muter. Sci. Muter. Med., 14 [lo] 863-7 (2003). ”A. V. Shevchenko, E. V. Dudnik, A. K. Ruban, V. P. Red’ko, V. M. Verexhaka, and L. M. Lopato, “Nanocrystalline Powders Based on ZrOz for Biomedical Applications and Power Engineering,” Powder Metall. Met. Cerum., 41. 55843 (2002). ”0. Vasylkiv, Y. Sakka. and V. V. Skorokhod, “Low-Temperature Processing and Mechanical Properties of Zirconia and Zirconia.-Alumina Nanoceramics,” J. Am. Ceram. Soc., 86, 299-304 (2003). I4M. Boaro, A. Trovarelli. J.-H. Hwang, and T. 0. Mason, ”Electrical and Oxygen Storage/Release Properties of Nanocrystalline Ceria-Zirconia Solid Solutions,” Solid Sfare Ionics. 147, 8>95 (2002). ”M. Mamak, N. Coomhs. and G. Ozin. “Self-Assembling Solid Oxide Fuel Cell Materials: Mesoporous Yttna-Zirconia and Metal-Yttria-Zirconia Solid Solutions,” J . Am. Chem. Soc.. 122, 8932-9 (2000).

Progress in Nanotechnology: Processing

“G. Pang, S. Chen, Y. Zhu, 0. Palchik, Y. Koltypin. A. Zaban. and A. Gedanken, “Preparation and Characterization of Monodispersed YSZ Nanocrystals,” J. Phys. Chem. B, 105. 4647-52 (2001). I7E. L. Crepaldi. G. J. de, A. A. Soler-Mia, A. Bouchara, D. Grosso, D. Durand. and C. Sanchez, “Controlled Formation of Highly Ordered Cubic and Hexagonal Mesoporous Nanocrystalline Yttria-Zirconia and Ceria-Zirconia Thin Films Exhibiting High Thermal Stability,” Angew. Chem. In!. Ed.. 42 [3] 347-51 (2003). ‘“G. Soyez, J. A. Eastman. L. J. Thompson, G.-R. Bai, P. M. Baldo, A. W. McCormick. R. J. DiMelfi, A. A. Elmustafa, M. F. Tambwe, and D. S. Stone, “Grain-Size-Dependent Thermal Conductivity of Nanocrystalline YttriaStabilized Zirconia Films Grown by Metal-Organic Chemical Vapor Deposition,” Appl. Phys. Lett., 77, 1155-7 (2000). ’X. Zhang, X. N. Jiang, and C. Sun, “Micro-Stereolithography of Polymeric and Ceramic Microstructures.” Sens. Actuators A. Phys., 77, 149-56 (1999). 20M.I. Cima, M. Oliveira, H.-R. Wang, E. M. Sachs, and R. Holman, “SlurryBased 3DP and Fine Ceramic Components”; pp. 21623 in Solid Freeform Fabrication Symposium Proceedings, Edited by J. J. Beaman. H. L. Marcus, D. L. Bourell, 1. W. Barlow, and T. Crawford. University of Texas, Austin, TX, 2001. ”8. A. Vojak, S.-J. Park, C. 1. Wagner, J. G . Eden, R. Koripella, J. Burdon, F. Zenhausern, and D. L. Wilcox, “Multistage, Monolithic Ceramic Microdischarge Device Having an Active Length of 0.27 mm,” Appl. Phys. Lett., 78, 1340-2 (2001). *’X. Zhao, J. R. G. Evans. M. J. Edirisinghe, and J.-H. Song, “Direct Ink-Jet Printing of Vertical Walls,” J . Am. Ceram. Soc., 85, 21 13-5 (2002). ”J. E. Smay. G. M. Gratson. R. F. Shepherd, J. 111 Cesarano, and J. A. Lewis, “Directed Colloidal Assembly of 3D Periodic Structures,” Adv. Mater., 14, 1279-83 (2002). “H. A. Lowenstam and S. Weiner, “Mineralization by Organisms and the Evolution of Biomineralization,” p. 191 in Biomineraliiolion and Biological Metal Accumulation, Edited by P. Westbroek and E. W. D. de long. Reidel Publishing Co, Hingham, MA, 1983. ”C. C. Perry, S. Mann. R. J. P. Williams, F. Watt, G. W. Grime, and J. Takacs, “A Scanning Proton Microprobe Study of Macrohairs from the Lemma of the Grass Phuluris Canariensb L,” Proc. R. Sac. 8. Biol. Sci.,222 [I2291 43945 (1984). %. Mann and G. A. Ozin, “Synthesis of Inorganic Materials with Complex Form,” Nature, 382, 3 13-8 ( I 996). ”5. Aizenberg, A A. Tkachenko, S. Weiner, L. Addadi, and G. Hendler, “Calcitic Microlenses as Part of the Photoreceptor System in Brittlestars,” Nature, 412 [6849] 819-22 (2001). 28 J. C. Weaver, L. 1. Pietrasanta, N. Hedin, B. F. Chmelka, P. K. Hansma, and D. E. Morse, “Nanostructural Features of Demosponge Biosilica.” J. Struct. Biol., 144, 271-81 (2003). ’9.C. Sundar, A. D. Yablon, J. L. Grazul, M. Ilan, and J. Aizenberg, “Fiberoptical Features of a Glass Sponge,” Nature, 424 [6951] 899-900 (2003). ’OE. Bauerlein, “Biomineralization of Unicellular Organisms: An Unusual Membrane Biochemistry for the Production of Inorganic Nano- and Microstructures,” Angew. Chem. Int. Ed., 42, 614-41 (2003). ”F. E. Round, R. M. Crawford, and D. G. Mann, The Diatoms: Biology & Morphology ofthe Genera. Cambridge University Press, Cambridge, U.K., 1990. ’%. A. Crawford, M. 3. Higgins, P. Mulvaney, and R. Wetherbee, “Nanostructure of the Diatom Frustule as Revealed by Atomic Force and Scanning Electron Microscopy,” J. Phycol., 37, 543-54 (2001). ”M. Hildebrand and R. Wetherbee, “Components and Control of Silicification in Diatoms”; pp. 11-57 in Progress in Molecular and Subcellular Biology, Vol. 33, Edited by W. E. G. Muller. Springer-Verlag. Berlin, 2003. ’‘V. Martin-Jezequel, M. Hildebrand, and M. A. Brzezinski, “Silicon Metabolism in Diatoms: Implications for Growth,” J . Phycol., 36, 82140 (2000).

Nanocomposites and Nanostructures

’9.Lebeau and J. M. Robert, “Diatom Cultivation and Biotechnologically Relevant Products. Part I: Cultivation at Various Scales,” Appl. Microbiol. Biotechnol., 60,612-23 (2003). 36D. Werner, “Introduction with a Note on Taxonomy”; pp. 1-17 in The Biology oj’Diat0m.v. Edited by D. Werner. University of California Press, Berkeley, CA, 1977. 37 J. Parkinson and R. Gordon. “Beyond Micromachining: The Potential of Diatoms,” Trends Biotechnol., 17, 190-6 (1999). ”J. Zhao, C. S. Caddis, Y. Cai, and K. H. Sandhage, “Free-Standing Microscale Structures of Nanocrystalline Zirconia Nanocrystals with Biologically Re licable 3-D Shapes,” J. Mater. Res.. 20 (21 282-7 (2005). ‘K. H. Sandhage, M. B. Dickerson, P. M. Huseman, M. A. Caranna, J. D. Clifton, T. A. Bull, T. J. Heibel. W. R. Overton. and M. E. Schoenwaelder, “Novel, Bioclastic Route to Self-Assembled, 3-D, Chemically Tailored Meso/ Nanostructures: Shape-Preserving Reactive Conversion of Biosilica (Diatom) Microshells,” Adv. Mater., 14, 429-33 (2002). 40 K. H. Sandhage, M. B. Dickerson, P. M. Huseman, F. M. Zalar. M. R. Rondon, and E. C. Sandhage, “A Novel Hybrid Route to Chemically-Tailored, ThreeDimensional Oxide Nanostructures: The BaSlC (Bioclastic and Shape-Preserving Inorganic Conversion) Process,” Ceram. Eng. Sci. Proc., 23 [4] 6 5 3 4 (2002). 4’R. R. Unocic, F. M. Zalar, P. M. Sarosi, Y. Cai, and K. H. Sandhage. “Anatase Assemblies from Algae: Coupling Biological Self-Assembly of 3 - 0 Nanoparticle Structures with Synthetic Reaction Chemistry,” Chem. Commun..(71 7 9 5 4 (2004). 4zP. Joubert and B. Gaudreau, “Zirconium(1V) Oxyfluorides,” Rev. Chim. Miner., 12 [4] 289-302 (1975). 431. Barin, Thermochemical Data ofPure Substances, Vol. 11, 3rd edition, VCH Verlagsgesellschaft mbH, Weinheim, Germany, 1995. “K. V. Gourishankar, M. Karaminezhad Ranjbar, and G. R. St.Pierre, “Revision of the Enthalpies and Gibbs Energies of Formation of Calcium Oxide and Ma nesium Oxide,” J. Phase Equilibria. 14 [5] 601-1 1 (1993). 4BR. C. Weast, Handbook ofchembtry and Physics, 55th edition, pp. B-1054. CRC Press, Cleveland, OH, 1974. 4%’. A. Tverskov and 1. S. Morozov, “System ZrC14-MgC12-KCI”; p. 340 in Phase Diagramsfor Ceramists 1969 Supplement, Vol. 11. Edited by E. M. Levin. C. R. Robbins, and H. F. McMurdie. The American Ceramic Society, Westerville. OH, 1985. 47B. D. Cullity, Elements of X-Ruy DifJraction, 2nd edition, p. 101. AddisonWesley Publishing Co, Reading, MA, 1978. 48E. Karapetrova. R. Platzer. J . A. Gardner, E. Tome, I. A. Sommers, and W. E. Evenson, “Oxygen Vacancies in Pure Tetragonal Zirconia Powders: Dependence on the Presence of Chlorine During Processing,” J. Am. Ceram. Sac., 84 [I] 65-70 (2001). 49D M. Pasquevich, F. Lovey, and A. Caneiro, “Structural and Microstructural Changes in Zirconia in Dilute Chlorine Atmosphere.” J. Am. Ceram. Sac., 72 [9] 1664-7 (1989). 50 E. V. Armbrust, J. A. Berges, C. Bowler, B. R. Green, D. Martinez, N. H. Putnam, S. Zhou, A. E. Allen, K. E. Apt, M. Bechner, M. A. Brzednski, B. K. Chaal, A. Chiovitti, A. K. Davis, M. S. Demarest, J. C. Detter, T. Glavina, D. Goodstein, M. Z. Hadi, U. Hellsten. M. Hildebrand. B. D. Jenkins, J. Jurka, V. V. Kapitonov, N. Kroger, W. W. Y. Lau, T. W. Lane, F. W. Larimer, J. C. Lippmeier, S. Lucas. M. Medina, A. Montsant. M. Obornik, M. S. Parker, B. Palenik, G. J. Pazour, P. M. Richardson, T. A. Rynearson, M. A. Saito, D. C. Schwartz, K. Thamatrakoln, K. Valentin, A. Vardi. F. P. Wilkerson, and D. S. Rokhsar, “The Genome of the Diatom Thalassiosira Pseudonana: Ecology, Evolution, and Metabolism,” Science, 306,79-86 (2004). ”L. A. Zaslavskaia, J. C. Lippmeier, P. G. Kroth, A. R. Grossman, and K. E. Apt, “Transformation of the Diatom Phneodacty/um Tricornutum (Bacillarrophyceae) with a Variety of Selectable Marker and Reporter Genes,” J . Phycol., 36,379-86 (2000).

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