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Copyright © 2009. Nova Science Publishers, Incorporated. All rights reserved. Nanopowders and Nanocoatings: Production, Properties and Applications : Production, Properties and Applications, Nova Science Publishers,

Copyright © 2009. Nova Science Publishers, Incorporated. All rights reserved. Nanopowders and Nanocoatings: Production, Properties and Applications : Production, Properties and Applications, Nova Science Publishers,

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Copyright © 2009. Nova Science Publishers, Incorporated. All rights reserved.

NANOPOWDERS AND NANOCOATINGS: PRODUCTION, PROPERTIES AND APPLICATIONS

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Nanopowders and Nanocoatings: Production, Properties and Applications : Production, Properties and Applications, Nova Science Publishers,

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NANOPOWDERS AND NANOCOATINGS: PRODUCTION, PROPERTIES AND APPLICATIONS

V. F. COTLER Copyright © 2009. Nova Science Publishers, Incorporated. All rights reserved.

EDITOR

Nova Science Publishers, Inc. New York

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Copyright © 2010 by Nova Science Publishers, Inc. All rights reserved. No part of this book may be reproduced, stored in a retrieval system or transmitted in any form or by any means: electronic, electrostatic, magnetic, tape, mechanical photocopying, recording or otherwise without the written permission of the Publisher. For permission to use material from this book please contact us: Telephone 631-231-7269; Fax 631-231-8175 Web Site: http://www.novapublishers.com NOTICE TO THE READER The Publisher has taken reasonable care in the preparation of this book, but makes no expressed or implied warranty of any kind and assumes no responsibility for any errors or omissions. No liability is assumed for incidental or consequential damages in connection with or arising out of information contained in this book. The Publisher shall not be liable for any special, consequential, or exemplary damages resulting, in whole or in part, from the readers‘ use of, or reliance upon, this material.

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Independent verification should be sought for any data, advice or recommendations contained in this book. In addition, no responsibility is assumed by the publisher for any injury and/or damage to persons or property arising from any methods, products, instructions, ideas or otherwise contained in this publication. This publication is designed to provide accurate and authoritative information with regard to the subject matter covered herein. It is sold with the clear understanding that the Publisher is not engaged in rendering legal or any other professional services. If legal or any other expert assistance is required, the services of a competent person should be sought. FROM A DECLARATION OF PARTICIPANTS JOINTLY ADOPTED BY A COMMITTEE OF THE AMERICAN BAR ASSOCIATION AND A COMMITTEE OF PUBLISHERS. LIBRARY OF CONGRESS CATALOGING-IN-PUBLICATION DATA Nanopowders and nanocoatings : production, properties, and applications / [edited by] V. F. Cotler. p. cm. Includes bibliographical references and index. ISBN:  (eBook)

1. Nanoparticles--Industrial applications. 2. Powders--Industrial applications. 3. Perovskite. 4. Coatings. I. Cotler, V. F. TA418.9.N35N3435 2009 620'.43--dc22 2009035164



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CONTENTS Preface

xi

Research and Review Chapters Chapter 1

Chapter 2

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Chapter 3

Chapter 4

Chapter 5

Chapter 6

Chapter 7

Perovskite Nanopowders: Synthesis, Characterization, Properties and Applications Xinhua Zhu

1

Electrochemical Nanocoatings on Titanium for Biomaterial Applications Kyo-Han Kim and R. Narayanan

69

Preparation, Characterization and Potential Application of Magnetic Materials as Sorbents for the Removal of Contaminants Chiung-Fen Chang, Ching-Yuan Chang, Wolfgang Höll and Matthias Franzreb

97

Effects of Ag/In Additives and Crystallization Kinetics on the Resistive Characteristics of Amorphous SbTe Chalcogenide Films Chung-Wei Yang, Chien-Chih Chou and Truan-Sheng Lui

123

Chemical Vapor Synthesis (CVS) of Inorganic Nanopowders H. Y. Sohn and Taegong Ryu

147

Nano- and Microstructural Silicon Powders in the Synthesis and Storage of Hydrogen А. А. Kovalevskii, А. S. Strogova, V. А. Labunov and А. А. Shevchenok Semiconductor Ceramic Materials Produced from AIIBVI Nanopowders N.N. Kolesnikov, E.B. Borisenko, V. V. Kveder, D.N. Borisenko, A.V. Timonina and B.A. Gnesin

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179

195

x

Contents

Expert Commentary The Electrochemical Synthesis of the Tungsten Carbide Nanopowders and Carbon Nanotubes Kh.B. Kushkhov and Kh. M. Berbekov Kabardino

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Index

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209 211

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PREFACE In nanotechnology, a particle is defined as a small object that behaves as a whole unit in terms of its transport and properties. It is further classified according to size: In terms of diameter, fine particles cover a range between 100 and 2500 nanometers, while ultrafine particles, on the other hand, are sized between 1 and 100 nanometers. Similarly to ultrafine particles, nanoparticles are sized between 1 and 100 nanometers, though the size limitation can be restricted to two dimensions. Nanoparticles may or may not exhibit size-related properties that differ significantly from those observed in fine particles or bulk materials. Nanopowders are agglomerates of ultrafine particles, nanoparticles, or nanoclusters. Adding nanoparticles to the polymer matrix of a coating does not make it a nanocoating. A nanocoating is synthesized using molecular engineering techniques to create a nanostructured polymer/coating. A coating with nanoparticles added to the polymer matrix will only incrementally improve the physical properties. Nanoparticles or nano-dirt can be used as an "additive" in a coating to reinforce the polymer matrix, reduce UV degradation of the substrate, improve chemical resistance and change the coatings electrochemical properties. The technique of using nanoparticles to maximize the physical surface properties of a coating are limiting. Adding nanoparticles to a polymer matrix depend on the molecular structure of the polymer backbone, size and amount of nanoparticles added, particle dispersion throughout the coating, structure and functionality of the nanoparticles. This new and important book gathers the latest research from around the globe in the study of these dynamic fields. Chapter 1 - Perovskite materials display a wide spectrum of attractive properties, such as ferroelectricity, piezoelectricity, dielectricity, ferromagnetism, magnetoresistance, and multiferroics, which make them attractive for applications in ferroelectric random access memories, multilayer ceramic capacitors, transducers, sensors and actuators, magnetic random access memories, and the potential new types of multiple-state memories and spintronic devices controlled by electric and magnetic fields. Following a similar trend to the miniaturization as the conventional CMOS (complementary metal oxide semiconductor) devices, the down-sized electronic devices based on perovskite electronic ceramic materials have also been developed. Advances toward nanoscale electronics have increased interest in the field of perovskite nanopowders. Perovskite nanopowders are versatile matrices for generating transition- and rare-earth metal oxides that exhibit a broad spectrum of properties and functions related to the following characteristics: (a) nearly innumerable combinations of metal cations can be accommodated within perovskite structural systems, (b) by reduction/reoxidation processes, nonstoichiometry (i.e., controlled amounts of ordered

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xii

V. F. Cotler

oxygen vacancies) can be introduced into the structure. In turn, high oxygen ion mobility or modified electronic and magnetic features can be implemented, and (c) the design of composite structural systems containing perovskite building units (perovskite slabs of different thicknesses) allows fine-tuning electronic and magnetic properties. Conventionally, perovskite nanopowders are prepared by solid-state reactions between the corresponding oxides or oxides and carbonates at temperatures above 1000oC. However, the resulting microstructures of perovskite nanopowders obtained from this method, are not suitable for the miniaturization of electronic devices, due to their significant particle agglomeration, poor chemical homogeneity, and coarse large particle sizes. To resolve the problems and to produce homogeneous and stoichiometric perovskite nanopowders, recently wet chemical methods have been developed. In this chapter, an overview of the state of art in perovskite nanopowders is presented, which covers their synthesis, characterization, properties and applications. The underlying chemical and physical aspects of the synthesis process of perovskite nanopowders are discussed, focusing on understanding the reaction mechanism and phase transformation from amorphous to the crystalline perovskite phase that occur during the synthesis process. Following a review of the synthesized methods, the microscopic and spectroscopic characterizations of perovskite nanopowders are summarized. And then a wide range of properties and applications of perovskite nanopowders are also addressed. Finally, a perspective on the future outlook of perovskite nanopowders is provided. Chapter 2 - Titanium and its alloys are the materials of choice for most dental and orthopedic applications. Advantages of these materials include biocompatibility, good resistance to corrosion and excellent mechanical properties. However, bone response and implant success depend on the chemical and physical properties of the surface. Integration of the titanium implants with bone tissue can be improved and accelerated by the presence of hydroxyapatite coating or oxide tubular layers on the implant surface. Adhesion of cells such as osteoblasts to the implant surface is an important prerequisite to subsequent cell functions. Nanometer-sized hydroxyapatite grains improve bioactivity and improve osteoblast functions and are better than the micron-sized hydroxyapatite. Hydroxyapatite is produced on titanium by electrochemical deposition from electrolytes containing calcium and phosphorus precursors. This process uses titanium cathode. Advantages of the cathodic process include processing at ambient temperatures, dimensional conformity, flexibility of grain size, no post-treatment and obtaining thin coatings with less residual stress. By modifying the applied voltages, nano-dimensional hydroxyapatite has been obtained as a thin coating on titanium. Recently ultrasonic agitation has been employed to obtain a thin coating of hydroxyapatite. This coating contains nano-sized apatite that shows a promising osteoblast cell activity. Nanotubes of TiO2 have attracted increasing scientific and technological attention due to the increased exploitation of specific functional properties of TiO2 in various applications. Compared with the flat TiO2 layers, the nanotubuar layes of TiO2 can serve as suitable substrate for hydroxyapatite growth in biomedical applications and can be prepared by various techniques such as sol–gel, electrophoretic deposition and anodic oxidation. Anodization is preferred to the sol–gel and electrophoretic deposition as it provides strongly adherent TiO2 layer that the other two approaches generally do not produce. TiO2 nanotubes have been formed by anodic oxidation in fluoride-based acid electrolytes and these have thicknesses of up to a maximum of 500nm. Use of neutral NaF based electrolytes can produce high aspect ratio self-organized TiO2 nanotubes with thicknesses

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Preface

xiii

higher than 2μm. Both these electrolytes however produce corrugated or rippled tube walls. Use of fluoride-containing glycerol electrolyte is shown to produce smooth tubes of very high aspect ratio. Ultrasonic agitation of the electrolyte has also been employed to produce good quality TiO2 nanotubes. This article describes the cathodic deposition of nano-hydroxyapatite first and the features of anodic nanotubular TiO2 later. The processes are detailed and the properties of the coatings including the cell behavior are indicated. Chapter 3 - Particles (e.g., adsorbents and catalysts) of smaller size are favorable to achieve fast mass transfer rates in interactions on surfaces because they can provide larger specific areas of external surface and shorter lengths of internal diffusion paths. Furthermore, ultrafine particles are of special interest because of their novel behavior and high potential as components of superior composite materials. However, the simple separation of ultrafine particles from liquid phases is still a difficult task. To solve the problem of liquid-solid separation resulting from ultrafine particles, magnetic separation technology can be applied. The development of such a combined innovative method using magnetic particles to remove contaminants and separating or recovering the used magnetic particles can be viewed as a great challenge for the scientists and engineers in the environmental and related fields. This paper summarizes, reviews and discusses related studies on the preparation and characterizations of novel magnetic materials and their application in the removal of contaminants. The basic theories and preparation methods of novel magnetic materials are presented in the paper. Physicochemical properties of the magnetic materials were characterized via methods such as transmission electron microscopy, scanning electron microscopy with energy dispersive X-ray spectroscopy, X-ray powder diffraction, superconducting quantum interference device, Fourier transform infrared spectroscopy, solid nuclear magnetic resonance and nitrogen gas adsorption for specific surface areas. With respect to the application of the novel materials as adsorbents, the adsorption behavior of contaminants onto both novel and commercial adsorbents was examined and compared to illustrate the distinct differences in the uptakes of the compounds between magnetic and conventional adsorbents. Furthermore, the relationships between adsorption phenomena and physicochemical properties of the magnetic adsorbents were interpreted. Chapter 4 - Chalcogenide films were able to be used as the recording layer of phase change recording media and in the application of phase change random access memory (PCRAM). The most attractive property of this material is its quick transformation between the amorphous and crystalline states, which phenomenon can accompany huge changes in the optical and electric properties. The reversible transformation between amorphous and crystalline phases was named as Ovonic Memory phenomenon and materials with such kind of properties were also named as the phase change materials. In practical applications, major efforts have been focused on increasing the crystallization speed and improvement on the optical or electrical contrast between amorphous and crystalline state. In the present study, there are two chalcogenide films being deposited on alkali-free glass with RF-sputtering method, pure SbTe films (ST) and Ag/In added SbTe films (AgInSbTe, AIST). In the first part, the microstructure and sheet resistivity of AIST films deposited with different parameters were analyzed. The results show that the as-deposited films possess amorphous structure no matter with what the sputtering parameter being adopted. The sheet resistivity measurement shows the amorphous films possess an extremely high resistivity and the temperature coefficient of resistivity (TCR) is negative. It is worth noting that the

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V. F. Cotler

relationship of amorphous AIST films between the sheet resistivity and film thickness was found to against the classic size effect. In the second part, similar amorphous films were annealed isothermally at different temperatures to obtain their different crystallinity. The sheet resistance of the annealed specimens was measured at room temperature, where the sheet resistance of amorphous films can be 3 × 104 higher than that of the crystalline films. As comparing X-ray diffraction patterns of AIST films to that of ST films, the sheet resistance change of the specimens can be correlated to the crystallization of amorphous phases, which transition temperature of the change in the sheet is at about 433 K for AIST films and 393 K for ST films. Through transmission electron microscopy (TEM) observations and Grazing-incidence X-ray diffraction (GI-XRD), the major phase in the crystalline ST films is the -Sb phase and the mixture of -Sb and AgSbTe2 phases in the crystalline AIST films. Concerning of the thermal activation measurements, the activation energy and crystallization temperature were measured with the differential scanning calorimeter (DSC). The activation energy of AIST films is about 0.92 eV and that of ST films is about 0.82 eV, the crystallization temperature of AIST films is about 475 K and that of ST films is about 445 K. The result reveals Ag/In added SbTe films possesses high room temperature stability. Because the sheet resistance has been proven to change with the crystallinity, an apparatus was developed to estimate the activation energy and the Avrami exponent of crystallization through Johnson-Mehl-Avrami formulism. The activation energy is estimated to be about 0.815 eV and the Avrami exponent (n) is about 1.1 to 1.4. The exponent indicates that the crystal can grow freely and the sheet resistance will decrease dramatically after the impingement effects occurring. A model is proposed to explain why the sheet resistance decreases within a very short period and the homogeneous nucleation and free growth during isothermal annealing in this study. Chapter 5 - Chemical vapor synthesis (CVS) is a process for making fine solid particles by the vapor-phase chemical reactions of precursors. At the University of Utah, this process has been used for the synthesis of the ultrafine powders of titanium and nickel aluminides and more recently aluminum nanopowder, tungsten and tungsten carbide nanopowders, and tungsten carbide - cobalt nanocomposite powder. This CVS process has proved its capability to prepare fine particles of 5-200 nm sizes. An example of the significant features of this technique is its unique capability to produce very uniformly mixed powders of different solid phases. This is possible because the reactants can be perfectly mixed in the gas phase. The chemical vapor synthesis is typically performed in a tubular reactor but more recently it has been carried out in a plasma reactor system. The plasma assisted chemical vapor synthesis adds many other advantages such as a high processing temperature to vaporize all reactants, a high quench rate to form ultrafine powders, and a wide choice of reactants. Thus, it has shown a considerable promise for many applications as a promising method for producing a variety of nanopowders. Chapter 6 - Complex use of products of interaction of MNS SP with water, utilization of heat and effective functioning of power cycle on the basis of nanostructural powder of silicon is the real technology of gaseous hydrogen for hydrogen energy of the nearest future. Application of MNS SP gives undeniable advantages: simple and without expenses of energy from outside technology of hydrogen production in the result of water decomposition: the necessity in storage and transportation of gasiform hydrogen disappears, which promotes fire-

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Preface

xv

and explosionproof of this fragment of hydrogen energy substantially. In the process of the industrial production of nanopowders of silicon, including for other areas of their application their prime price will be reduced some times of their value. The application of MNS SP in the mobile small sources of hydrogen is already expedient. Chapter 7 - Based on our experience of conventional crystal growth of AIIBVI, a techniques has been developed for nanocrystal production of some of these compounds. Further innovations in the technological process appear to significantly reduce the development and production cost of the final product with respect to conventional single crystals, while maintaining a high level of physical properties, which conform to leading standards. New technology of the vapor phase deposition was developed in our laboratory to produce CdTe nanoparticles of average diameter 8 nm. Further development of the technology allows us to produce 10-nm Cd-Zn-Te nanocrystals and to overcome difficulties in obtaining a ternary solid solution Cd1-xZnxTe (x = 0.04—0.1) with a stable chemical composition. Highly dense CdTe and Cd1-xZnxTe ceramics of high mechanical hardness and durability were produced by the room temperature process without any lubricants or binding materials. It was found that CdTe ceramics produced from nanocrystals undergo wurtzite— sphalerite transition under pressure. The polymorphic transition from the hexagonal to the cubic phase in Cd1-xZnxTe ceramics caused by annealing is discussed. The two-component texture composed of axial and {100} components was recorded in the as-compressed materials. The effect of annealing on grain growth and texture in the ceramics is considered. It was found that our compacted ceramics guarantees high transmittance in a wide IR region of 6— 25 m, high specific resistivity on the order of 1010 cm, and high microhardness on the order of 103 MPa. The obtained properties make these materials promising for use in IR optics and for ionizing-radiation detectors.

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In: Nanopowders and Nanocoatings Editor: V. F. Cotler

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Chapter 1

PEROVSKITE NANOPOWDERS: SYNTHESIS, CHARACTERIZATION, PROPERTIES AND APPLICATIONS Xinhua Zhu National Laboratory of Solid State of Microstructures, Department of Physics, Nanjing University, Nanjing 210093, China

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ABSTRACT Perovskite materials display a wide spectrum of attractive properties, such as ferroelectricity, piezoelectricity, dielectricity, ferromagnetism, magnetoresistance, and multiferroics, which make them attractive for applications in ferroelectric random access memories, multilayer ceramic capacitors, transducers, sensors and actuators, magnetic random access memories, and the potential new types of multiple-state memories and spintronic devices controlled by electric and magnetic fields. Following a similar trend to the miniaturization as the conventional CMOS (complementary metal oxide semiconductor) devices, the down-sized electronic devices based on perovskite electronic ceramic materials have also been developed. Advances toward nanoscale electronics have increased interest in the field of perovskite nanopowders. Perovskite nanopowders are versatile matrices for generating transition- and rare-earth metal oxides that exhibit a broad spectrum of properties and functions related to the following characteristics: (a) nearly innumerable combinations of metal cations can be accommodated within perovskite structural systems, (b) by reduction/reoxidation processes, nonstoichiometry (i.e., controlled amounts of ordered oxygen vacancies) can be introduced into the structure. In turn, high oxygen ion mobility or modified electronic and magnetic features can be implemented, and (c) the design of composite structural systems containing perovskite building units (perovskite slabs of different thicknesses) allows fine-tuning electronic and magnetic properties. Conventionally, perovskite nanopowders are prepared by solid-state reactions between the corresponding oxides or oxides and carbonates at temperatures above 1000oC. However, the resulting microstructures of perovskite nanopowders obtained from this method, are not suitable for the miniaturization of electronic devices, due to their significant particle agglomeration, poor chemical homogeneity, and coarse large particle sizes. To resolve the problems and to produce homogeneous and stoichiometric perovskite nanopowders, recently wet chemical

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2

Xinhua Zhu methods have been developed. In this chapter, an overview of the state of art in perovskite nanopowders is presented, which covers their synthesis, characterization, properties and applications. The underlying chemical and physical aspects of the synthesis process of perovskite nanopowders are discussed, focusing on understanding the reaction mechanism and phase transformation from amorphous to the crystalline perovskite phase that occur during the synthesis process. Following a review of the synthesized methods, the microscopic and spectroscopic characterizations of perovskite nanopowders are summarized. And then a wide range of properties and applications of perovskite nanopowders are also addressed. Finally, a perspective on the future outlook of perovskite nanopowders.

Keywords: perovskite nanopowders; synthesis; characterization; properties and applications

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1. INTRODUCTION Perovskite materials are one of the most widely investigated functional materials, which have important properties in ferroelectricity, piezoelectricity, dielectricity, ferromagnetism, magnetoresistance, and multiferroics. The perovskite structure is named for the prototype CaTiO3 mineral called perovskite, which is generally metal oxide with the formula ABO3, where B is a small transition metal cation and A is a larger s-, d-, or f-block cation. In a cubic perovskite, the larger cation A resides on the corners of the unit cell, the smaller cation B is in the center of the unit cell, and the oxygen ions (O2-) are on the centers of the faces (Figure 1a) [1]. The perovskite structure can be also built from three-dimensional corner-sharing BO6 octahedra that are connected through B-O-B linkages. The A-site cation fits in the large cavity at the center of eight corner-sharing BO6 octahedra, and the B-site cation resides in the interstitial site of an octahedron of oxygen anions (Figure 1b) [1]. Interestingly, and of technological importance, a variety of compositions crystallizes in the perovskite structure. Typical perovskite materials of technological importance are ferroelectric BaTiO3, PbTiO3, dielectric (Ba,Sr)TiO3, piezoelectric Pb(Zr,Ti)O3, electrostrictive Pb(Mg,Nb)O3, magnetoresistant (La,Ca)MnO3, and multiferroic BiFeO3. They have attracted interest for several decades, with tremendous applications including ferroelectric random access memories, multilayer ceramic capacitors, transducers, sensors and actuators, magnetic random access memories, and the potential new types of multiple-state memories and spintronic devices controlled by electric and magnetic fields [1-8]. The major challenge in manufacturing these materials is the processing of the materials with reliable and reproducible properties [9,10]. Following a similar trend to the miniaturization as the conventional CMOS (complementary metal oxide semiconductor) devices, the down-sized electronic devices based on perovskite electronic ceramic materials have also been developed. Advances toward nanoscale electronics have additionally increased interest in this field of perovskite nanoparticles [11-13]. For example, to develop high volume efficient multilayered ceramic capacitors (MLCCs), the sizes of BaTiO3 particles with high purity and uniform shape used for fabricating the next generation of MLCCs will be lowered down to tens of nanometers. Therefore, synthesis of high-purity, ultra-fine and agglomerate-free perovskite nanopowders with controlled particle size, morphology and stoichiometry, is the critical step in processing of perovskite ceramics with desirable properties. Perovskite nanopowders are versatile matrices for generating transition- and rare-earth metal oxides that exhibit a broad spectrum

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of properties and functions that are related to the following characteristics: (a) nearly innumerable combinations of metal cations can be accommodated within perovskite structural systems, (b) by reduction / reoxidation processes, nonstoichiometry (i.e., controlled amounts of ordered oxygen vacancies) can be introduced into the structure. In turn, high oxygen ion mobility or modified electronic and magnetic features can be implemented, and (c) the design of composite structural systems containing perovskite building units (perovskite slabs of different thicknesses) allows fine-tuning electronic and magnetic properties.

a

b

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Figure 1. (a) Unit cell of ABO3 perovskite structure, and (b) an ABO3 perovskite structure in which corner-shared oxygen octahedral extending in three dimensions. Reproduced with permission from [1], Schaak, R. E.; Mallouk, T. E. Perovskites by design: a toolbox of solid-state reactions. Chem Mater. 2002, 14, 1455-1471.Copyright © 2002, American Chemical Society.

The evolution of a method to produce perovskite nanopowders with precise stoichiometry and desired properties is much complex. Conventionally, perovskite nanopowders are prepared by solid-state reactions between the corresponding oxides or oxides and carbonates at temperatures above 1000oC [14,15]. However, the resulting microstructures of perovskite nanopowders obtained from this method are not suitable for the miniaturization of electronic devices, due to their significant particle agglomeration, poor chemical homogeneity, and coarse large particle sizes. To resolve the problems arising from the conventional ceramic techniques and to produce homogeneous and stoichiometric perovskite nanopowders, in recent years, wet-chemical routes have been developed [12,16-18]. They can be better controlled from the molecular precursor to the final material to give highly pure and homogeneous materials, allowing for the low reaction temperatures used. The size and morphology of the particles can be controlled, and metastable phases could be prepared [18]. The objective of this chapter is to provide an overview of the state of art in perovskite nanopowders, which covers their synthesis, characterization, properties and applications. First, we review the synthesized methods for perovskite nanopowders, which include the syntheses using solid, liquid or gas phase precursors. The second section deals with the electron microscopic and spectroscopic tools for characterization of perovskite nanopowders. The microstructural features of perovskite nanopowders revealed by electron microscopes and spectroscopic techniques are addressed. In the context of properties we discuss the unique properties of perovskite nanopowders (e.g., ferroelectric and dielectric, electrical, magnetic, optical, and multiferroic properties), and the size effects for these unique properties are also

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discussed. And then a broad range of applications of perovskite nanopowders is addressed. Finally, we provide a perspective on the future outlook of perovskite nanopowders.

2. SYNTHESIS OF PEROVSKITE NANOPOWDERS Due to the powder size, dimensionality, and composition governing the resultant properties of the nanostructured perovskite materials that are assembled from nanopowders as building blocks to achieve certain desired properties, the synthesis of high-purity, ultra-fine and agglomerate-free perovskite nanopowders with controlled particle size, morphology and stoichiometry is the first and perhaps the most crucial step in processing of perovskite ceramics with desirable properties. The major issues for the synthesis of perovskite nanopowders include: (a) the control of particle size and composition, and (b) the control of the interfaces and distributions of the nanobuilding blocks within the fully formed nanostructured perovskite compounds. Over the past several decades, various methods have been developed to prepared perovskite nanopowders and the related nanostructured perovskite compounds. These various methods include synthesis using solid, liquid or gas phase precursors, which come under physical or chemical processing. In the subsequent sections, some important methods for the preparation of perovskite nanopowders are described.

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2.1. Solid-State Reaction Route The solid-state reaction method is the most traditional one for preparing perovskite nanopowders (e.g., BaTiO3, PbTiO3, Pb(Zr,Ti)O3, etc) [14,15]. This process includes weighting starting materials (the corresponding oxides or oxides and carbonates), mixing, milling, and calcining them at elevated temperatures to form the perovskite phase. For example, in synthesis of BaTiO3 nanopowders by solid-state reaction method, the reaction process in air has been proposed to take place in at least three stages and relies on the diffusion of Ba2+ ions into TiO2 [19]. Firstly, BaCO3 reacts with the outer surface region of TiO2 to form a surface layer of BaTiO3 on individual TiO2 grains. Further diffusion of Ba2+ ions into TiO2 necessitates the formation of Ba2TiO4 between the unreacted BaCO3 and the previously formed BaTiO3. After prolonged sintering periods, the intermediate Ba-rich phase Ba2TiO4 reacts with the remaining TiO2 in the core-regions of the TiO2 grains to form BaTiO3. The high temperature calcination produces an agglomerated powder with a coarse particle size which requires additional milling process. However, contamination and other undesirable features during the milling process can create defects in the manufactured products. Furthermore, the more components in the ceramic powders, the more difficult it may be to achieve the desired homogeneity, stoichiometry, and phases. By using nanocrystalline BaCO3 and TiO2 as starting materials, Buscaglia et al. [20] have recently synthesized the perovskite BaTiO3 nanopowders with size of ~ 100 nm and narrow particle size distribution, via a solid-state reaction at calcination temperatures as low as 800oC. The average particle size of powders obtained via this method is essentially determined by the particle size of the used TiO2 because the reaction rate is controlled by the diffusion rate of

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barium ions into the TiO2 lattice [21]. Similar reaction mechanism was also found in the synthesized process of BaZrO3 powders [22]. The morphology of BaZrO3 particles was dependent upon the initial size and shape of the used starting ZrO2 particles. Therefore, fine BaZrO3 powders with particle size of 70-100 nm composing of crystallites of ~ 20-30 nm can be synthesized by using very fine (70-90 nm) starting ZrO2 particles and coarse (~ 1 µm) BaCO3 particles commercially available and calcination at ~ 1000°C. Higher calcination temperatures accelerate the initial stage of reaction but often lead to coarser and moreagglomerated powders.

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2.2. Mechanical Milling Method Recently, many perovskite nanoparticles (e.g., BaTiO3, PbTiO3, PbZrO3) have been successfully synthesized by using several mechanical milling methods [23-25]. As viewed from the energy efficiency, the vibro-mill (or vibratory mill or vibro-energy mill) seems to be more attractive than the ball milling [24]. The vibro-milling enjoys several advantages over the conventional ball-milling produces, such as finer particles, narrower size distribution at a faster rate, simple equipment, low cost starting precursors, and large-scale production of nanopowders [23-25]. This implies that the vibro-milling method can be recognized as a powerful method for producing perovskite nanopowders. By choosing proper milling time and the calcination conditions, high purity perovskite nanopowders such as BaTiO3, PbTiO3, PbZrO3 with the smallest particle size of 100 nm, 17 nm, and 31 nm, can be mass-produced, respectively [26]. During the mechanical milling process, the mechano-chemical activation by the heavy milling is the key step, which alters the physicochemical properties of the starting materials and the mechanism of synthesis. Beauger et al. [27] proposed a multi-step reaction model, to describe the formation of perovskite BaTiO3 nanopowders via the mechanical milling process. According to this model, BaTiO3 is easily formed at the surface of TiO2 particles, which also act as the catalysts for BaCO3 decomposition [28]. When the surface BaTiO3 layer is formed by the decomposition of BaCO3 and its reaction with TiO2, the reaction kinetics is governed by the barium and oxygen ion diffusion through this layer into the virgin TiO2 phase. Moreover, it is expected that starting TiO2 with fine particles is very beneficial to acquiring the final BaTiO3 nanopowder due to the increase in the contact area of reactant particles (high reactivity) and their easy decomposition at low temperature. Because of the excess of barium and oxygen ions in the surface layer, a Ba2TiO4 phase is generally formed at the initial stage [29-31]. Homogeneous BaTiO3 powders can be formed gradually by the reaction between Ba2TiO4 and TiO2 via a multi-step reaction process, as schematically shown in Figure 2 [12]. Welham [33] also demonstrated that nanocrystalline BaTiO3 powders with an average particle diameter slightly larger than 10 nm could directly be obtained by high-energy mechanical milling of BaO and TiO2 (rutile) for several days without additional heat treatment. Although the extremely fine particles can be synthesized by this method, such an approach suffers from the disadvantages of quite small batch sizes and very long processing times. In addition, intensive ball-milling process may result in unfavorable contaminations from the milling media.

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Figure 2. Schematic diagram of the multi-step formation mechanism of BaTiO3 by the solid-state reaction of BaCO3 with TiO2. Reproduced with permission from [12], Yoon, D. H.; Lee, B. I. J. Tetragonality of barium titanate powder for a ceramic capacitor application. Ceram Proc Res. 2002, 3, 41- 47. Copyright © 2002, Journal of ceramic Processing Research.

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2.3. Wet Chemical Routes Perovskite nanopowders prepared by the conventional solid-state reactions usually suffer from the particles with uncontrolled and irregular morphologies, which result in poor electrical properties of the sintered ceramics. In recent years, various wet chemical methods have been developed to replace the conventional solid-state reactions for the synthesis of perovskited nanopowders. The popular wet chemical methods for the preparation of perovskite nanopowders, include sol-gel method [43-49], alkoxide-hydroxide solprecipitation method [50-54], hydrothermal method [55-63], microwave-hydrothermal [6473], solvothermal syntheses [74-78], glycothermal process method [79-81], spray pyrolysis method [82,83], microemulsion synthesis method [84-87], high-gravity reactive precipitation [88-91] and room-temperature biosynthesis [92,93]. The most important advantages of the wet chemical methods include easy controlling the chemical stoichiometry, producing nanopowders with narrow size distribution, and low crystallization temperature due to the constituents mixed at the quasi-atomic level in a solution system. Due to the wet chemical solution process, a dopant such as paramagnetic ions or rare-earth ions could be readily introduced during the preparation of the precursor solution. In the following subsequent sections, various wet chemical methods uased for preparation of perovskite nanopowders are introduced.

2.3.1. Sol-gel (colloidal) processing Sol-gel process is a popular processing route for the synthesis of perovskite nanopowders (e.g., BaTiO3, PbTiO3, BiFeO3) [43-49]. This process involves the formation of a sol by dissolving the metal aloxide, metal-organic, or metal-inorganic salt precursors in a suitable solvent, subsequent drying of the gel followed by calcination and sintering at high temperature to form perovskite nanopowders. Due to the reacting species homogenized at the atomic level in a sol-gel process, the diffusion distances are considerably reduced compared to a conventional solid-state reaction, therefore, the product can be formed at much lower temperatures. In this process, the selection of starting materials, concentration, pH value, and heat treatment schedule play an important role in affecting the properties of perovskite nanopowders. This has been demonstrated in the case of BaTiO3 perovskite nanopowders [94-99]. Barium acetate and titanium isopropoxide are often used as starting materials to

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synthesize BaTiO3 nanopowders. However, the different rates in the hydrolysis and condensation of Ba and Ti precursors often give rise to the chemical component segregation in the obtained gels. To solve this problem, acetic acid or acetylaceton was often used to control the hydrolysis rate of the Ti precursor, since these complexing agents acts as chelating agents to coordinate with Ti species [100,101]. For the obtained gels, a heat treatment at high-temperature over 600C is required to remove the unreacted organics and to crystallize the powders. Several steps involves in the transformation from the precursor to the crystalline BaTiO3 nanopowders, including the transformation from the precursor to the amorphous BaTiO3, and then to the threedimensional nucleation of the crystalline BaTiO3 in the amorphous matrix, and finally to the nanocrystal growth of BaTiO3 via a solid-state reaction [102]. To better control the grain size and its distribution, the heat treatment process parameters of the gels (e.g., post-annealing temperature, time and atmosphere, heating rate) must be optimized [102,103]. Normally, higher annealing temperature or longer annealing time can lead to larger grain size of the powders, while slow heating rate and inert annealing atmosphere can inhabit the aggregated behavior of nanopowders in comparison to air or oxygen atmosphere. That was demonstrated in the synthesis of Pb(Zr,Ti)O3 nanopowders [103]. By using these techniques, monodispersed perovskite nanopowders and related nanostructured materials have been successfully fabricated. The particle size can be adjusted from a few nanometers to micrometers via controlling the sold-state polymerization and the heat treatment process [43,96,102]. In the sol-gel routes based on acetate or double alkoxide precursors, however, noncrystalline BaTiO3 precursors are produced at first, and heating to 800oC or above is required to obtain crystalline BaTiO3 particles, which often removes nano-dimensional morphological characteristics of precursors and produces coarser chemically bonded aggregates.

2.3.2. Alkoxide-hydroxide sol-precipitation synthesis Crystalline BaTiO3 nanopowders can be directly synthesized via an alkoxide–hydroxide sol-precipitation process, which was first proposed by Flaschen [50]. This process has been studied extensively to produce crystalline BaTiO3 nanopowders at a low temperature without further calcination at an elevated temperature. This could reduce the manufacturing costs while maintaining better particle characteristics that could be realized by controlling the precipitation processes. Up to date, numerous publications and patents on the alkoxidehydroxide sol-precipitation process of BaTiO3 powders have appeared in the literature [50,51,104-108]. In this process, the hydrolysis and condensation are the key mechanisms of crystal growth. It was recognized that the amount of water and the method of its addition to the reaction system are decisive factors for controlling the size and the shape of precipitates. Many investigations have shown that BaTiO3 nanopowders can be synthesized at low temperature as 80-100oC via alkoxide-hydroxide method by using aqueous alkaline solution as a starting material [104-108]. This can be ascribed to the hydrolysis-condensation reaction occurring instantly upon mixing of aqueous and alcohol solutions. However, the final products are often highly agglomerated, and their morphological characteristics are quite inhomogeneous and undesirable for subsequent powder processing and sintering. By using the solid barium hydroxide octahydrate as starting material, it is possible to control the hydrolysis-condensation reaction by using water molecules released in situ as Ba(OH)28H2O

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was dissolved in the alcoholic solution. The experimental results have shown that BaTiO3 nanocrystals smaller than 6 nm begin to nucleate at 50oC without forming the intermediate TiO2 anatase, and corner-sharing TiO6 octahedra formed at 60oC. The average size of BaTiO3 precipitates increases up to about 7.5 nm at 80oC, and the BaTiO3 nanopowders show an anomalous lattice expansion with a relatively high tetragonality [52]. Direct formation of the perovskite phase at such low temperatures can be understood from the viewpoint of the coordination chemistry of the transition-metal alkoxide, in which titanium ions are present in unsaturated 4+ oxidation states in tetrahedral coordinations. As the reaction temperature is raised, the hydrolysis and condensation reactions convert them into three-dimensional TiO6 octahedra that share their six corners with other octahedra. The water and hydroxyl ions released from Ba(OH)28H2O convert tetrahedral Ti-isopropoxide into octahedral Ti(OH)62-, which reacts with Ba2+ to form perovskite BaTiO3 directly.

2.3.3 Decomposition of complex double metal salts Perovskite nanopowders such as BaTiO3 can be prepared by decomposition of complex double metal salts at temperatures above 600oC and up to 1300oC via several intermediate phases [34-36]. In this process, complex double salts of Ba and Ti, such as barium titanyl oxalate BaTiO(C2O4)24H2O [34,35] or -citrate BaTi(C6H6O7)36H2O [36], were used as solid precursors. Since both cations are already mixed on an atomic scale in the solid precursor, the thermal treatment leading to the perovskite phase can be performed at much lower temperatures compared with the mixed oxide route. Another advantage, resulting from the intimate mixture, is the possibility to attain very pure and almost exactly stoichiometric compositions under proper synthesis conditions. The reaction mechanism of BaTiO3 nanopowders formed from barium titanyl oxalate has been studied by several techniques. The results suggest that the formation of crystalline BaTiO3 nanopowders includes two steps: (a) the monoclinic crystal structure of the double oxalate initially collapses and converts into an amorphous upon drying, (b) crystallization via the intermediate phases into pseudocubic or tetragonal BaTiO3, depending on the pyrolysis temperature and the dopant content [34-36]. Due to the intimate mixture of barium and titanium ions, the reaction is complete at a much lower temperature in comparison to the conventional solid-state reaction process. Unfortunately, this process does not afford easy control of particle size and interparticle agglomeration: the size and agglomeration of the oxalate are maintained in the decomposed primary barium titanate particles, with a typical average size of approximately 50-250 nm [37,38], depending on the calcination temperature. In addition, the heating rate during calcination is also an effective parameter for controlling the particle size during nonisothermal but rate-controlled thermal decomposition [39]. As an alternative to conventional isothermal calcination, Wada et al. [40,41] proposed a two-step thermal decomposition method using barium titanyl oxalate to synthesize BaTiO3 nanopowders. The first step was carried out at 400oC in flowing O2 for complete dehydration, removal of remaining carbon species, and for preventing the formation of BaCO3 and TiO2. The amorphous phase obtained in the first heating treatment corresponds to the chemical composition of an equimolar mixture of BaCO3 and TiO2. In the second step, upon heat treatment at around 600oC under vacuum, extremely small BaTiO3 crystallites with an average particle size of approximately 17 nm were obtained. Particle size can be easily controlled from 17 nm to 100 nm by changing the temperature during the second annealing

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stage under vacuum [41]. The presence of tetragonality in all of these powders was evidenced by Raman spectroscopy, suggesting that the intrinsic critical size of single crystalline BaTiO3 should be below 17 nm [42].

2.3.4. Hydrothermal routes 2.3.4.1. Hydrothermal process Hydrothermal synthesis involves heating an aqueous suspension of insoluble salts in an autoclave at a moderate temperature and pressure so that the crystallization of a desired phase will take place. The hydrothermal synthesis is a powerful method for the preparation of very fine and homogeneous perovskite powders with a narrow size distribution and spherical morphology. Compared with the routes based on the solid-state reaction or decomposition of the solid precursors, the advantages of hydrothermal crystallization are the reduced energy costs due to the moderate temperatures sufficient for the reaction, less pollution, simplicity in the process equipment, and the enhanced rate of the precipitation reaction. Since there is no necessity for high-temperature calcination in this case, so the additional milling process is eliminated. For an ABO3 perovskite nanopowders, the general hydrothermal reaction can be written as [109]

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A(OH)s+B(OH)s(dissolution)  A(OH)aq +B(OH)aq (precipitation)  ABO3(s)

(1)

The perovskite BaTiO3 nanopowders have been prepared via hydrothermal process from titanium sources (such as oxide, oxide gels, or metalorganic compounds) and an aqueous solution of Ba(OH)2 with NaOH as a mineralizer. Typically, in a hydrothermal process, this reaction involves the reaction of Ba(OH)2 (or some other strong base with a soluble barium salt) and titanium sources (e.g., titanium alkoxide, titanium oxide, titanium oxide gels, or metalorganic compounds). This process leads to powders with a very high purity and little agglomeration. Furthermore, the low cost and easy handling of the reagents, and the fast reaction rate at low temperatures ensure that deagglomerated powders consisting of small particles with narrow size distribution are readily obtained. Although the optimization of hydrothermal conditions for the preparation of nanosized BaTiO3 has often been a matter of empiricism, much improvement has been achieved in the theoretical understanding of the thermodynamics and kinetics of the process, as well as the mechanisms of particle formation. Based on the standard-state thermodynamical properties of the chemical species involved during hydrothermal synthesis of BaTiO3 (e.g., Gibbs energy of formation, partial molal volumes, heat capacity, and etc.), Lencka and Riman [110-112] calculated the phase diagrams which allow ones to predict the required synthesis parameter (e.g., the ranges of reagent concentrations, pH value, and temperature for maximum yield), guiding the experimental hydrothermal process. To control the growth of BaTiO3 powders with the desired size and particle morphology, reaction mechanisms and thermodynamic modeling for BaTiO3 nanopowder formation during hydrothermal processing have been widely investigated [113-116]. Based on the highresolution transmission electron microscopy (HRTEM) observations on the incompletely and fully reacted powders, Pinceloup et al. [113] proposed a dissolution-precipitation model for hydrothermal synthesis of BaTiO3 nanopowders using Ba(OH)2 and TiO2 as precursors. In this model, TiO2 particles are first dissolved to form hydroxytitanium complexes [Ti(OH)n-],

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and then react with barium ions in the solution to precipitate BaTiO3. On the other hand, Hertl [114] and Hu et al. [115] proposed another in situ heterogeneous transformation model, in which TiO2 particles react initially with the dissolved barium to produce a continuous layer of BaTiO3, and the additional barium must diffuse through this layer and reacts with TiO2 until the supply of TiO2 is exhausted. This model was supported experimentally by the hydrothermal conversion from TiO2 microspheres to nanocrystalline BaTiO3 [115]. Eckert et al. [116] also reported on a mechanism evolution from a dissolution-precipitation process at the early stage of the reaction to an in situ mechanism for the longer reaction times. Recently, Walton et al. [117] investigated the hydrothermal crystallization of BaTiO3 by time-resolved powder neutron diffraction methods in situ, using the newly developed Oxford/ISIS hydrothermal cell. They directly observed that the rapid dissolution of the barium source was followed by dissolution of the titanium source before the onset of crystallization of BaTiO3. These qualitative observations strongly suggest that a homogeneous dissolution-precipitation mechanism dominates in the hydrothermal crystallization of BaTiO3 rather than other possible mechanisms proposed in the literatures [114-116]. These contradictive experimental observations reported previously are probably resulted from the different hydrothermal conditions. The crystalline perovskite phase BaTiO3 can be directly synthesized under hydrothermal conditions, however, the resulting products are usually highly defective in their crystallographic structure [118-123]. The crystal symmetry generally does not correspond to the tetragonal modification that is the thermodynamically stable form for BaTiO3 under normal conditions. Rather a cubic modification is often obtained [118-121,124], although the tetragonal phase may be obtained under certain conditions, such as high processing temperatures and prolonged duration time. Vikanandan et al. [119] suggested that the presence of the metastable cubic phase at room temperature is resulted from the compensation of the residual hydroxyl ions in the oxygen sublattice by cation vacancies. Shi et al. [121] also reported the stabilization of the cubic phase of BaTiO3 synthesized by the hydrothermal method, was caused by surface defects including OH defects and barium vacancies. Hennings and Schreinemacher [122] reported on the observation of lattice hydroxyls and the effect of their release on the crystallographic recovery in hydrothermal BaTiO3 particles. Norma et al. [125] also reported that in the as-prepared barium titanate nanopaticles with average size of 66 nm, there was a high concentration of the hydroxyl group and barium vacancy, and its crystal structure was assigned to cubic with an expanded lattice by using the Rietveld. It has been found that the structural defects in the BaTiO3 nanopowders synthesized by hydrothermal method are primarily in the form of lattice OH ions, which are compensated by '' barium vacancies ( VBa ) created on the surfaces of individual particles to maintain the

electro-neutrality [123-129]. The coexistence of high amounts of barium, titanium, and oxygen vacancies provides a rather unstable situation for the BaTiO3 lattice. It is believed that these point defects on the different lattice sites combine and annihilate each other. Therefore, the different charges of the point defects compensate each other to neutrality. As a result, the vanishing vacancies formed upon dehydration are believed to be responsible for the formation of intragranular pores (shown in Figure 3a), which partly disappear upon grain growth above 800oC. In the MLCCs ceramics made from the hydrothermal BaTiO3 powders, a strange expansion called ―bloating phenomena‖ (shown in Figure 3b) was observed at the final stage

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of sintering, which was due to the inherently incorporated hydroxyl ions and protons during the hydrothermal synthetic process [13].

Figure 3. Intragranular porosity of hydrothermal BaTiO3. (a) Powder heat treated for 2 h at 500oC, (b) dielectric X7R ceramics sintered at 1320oC and showing a huge amount of intragranular pores. Reproduced with permission from [13], Pithan, C.; Hennings, D.; Waser, R. Progress in the synthesis of nanocrystalline BaTiO3 powders for MLCC. Int J Appl Ceram Technol. 2005, 2, 1-14. Copyright © 2005, American Ceramic Society.

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In summary, hydrothermally synthesized BaTiO3 can show an adverse effect such as bloating in the final stage of the sintering process due to its inherently incorporated hydroxyl ions and protons during the synthetic process, despite the ideal uniform size and spherical particle shape.

2.3.4.2. Solvothermal process Solvothermal synthesis is defined as a hydrothermal reaction that occurs in a non-aqueous solution (e.g., NH3, methanol, ethanol, and n-propanol). In comparison with the hydrothermal processing, solvothermal synthesis has some advantages [76], such as (a) the reaction occurs under mild conditions and gives cubic-phase perovskite powder; and (b) the powders with particle size on the nanometer scale, exhibiting low agglomeration and a narrow particle-size distribution, due to the differences between the solvents. Up to date, several attempts have been made to synthesize superfine BaTiO3 nanopowder by solvothermal synthesis [17,76-78]. Using benzyl alcohol as solvent, BaTiO3 and BaZrO3 nanoparticles were synthesized by solvothermal process at relatively low temperatures of 200-220oC [17]. An assembly of BaTiO3 nanoparticles with an average particle size of 6 nm is shown in Figure 4a. The lack of any surface protecting layers results in some agglomeration of the particles. According to the randomly oriented lattice fringes, the particles have not coalesced. Based on the selected-area electron diffraction (SAED) pattern (see the inset in Figure 4a), the lattice distances measured from the diffraction rings, are in perfect agreement with the cubic and tetragonal modifications of the BaTiO3 perovskite structure. HRTEM images of two isolated particles oriented along the [110] and [111] directions, are shown in Figure 4b and Figure 4c, respectively. Figure 4d is a fast Fourier transform (FFT) pattern (equivalent to experimental electron diffraction pattern of the local region) obtained from the HRTEM image shown in Figure 4c, which provides an evidence that the particles are well crystallized in the perovskite structure without the presence of defaults. An overview TEM image

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for the BaZrO3 nanoparticles is shown in Figure 5a. It is observed that in most cases the primary particles are not isolated, but form wormlike agglomerates with diameters of 2-3 nm and lengths of up to 50 nm. These worms often assemble into larger, ball-like structures. The SAED pattern of such a spherical assembly (the set in Figure 5a) exhibites some broad rings that match with the BaZrO3 structure. Furthermore, the HRTEM image (Figure 5b) shows that the lattice planes of the individual particles in the ball-like structure oriented randomly with respect to each other. The HRTEM pattern of an isolated elongated particle proved the high degree of crystallinity (Figure 5c). This was further confirmed by the FFT pattern of this particle (Figure 5d), which is characteristic for the BaZrO3 structure without structural defaults. Obviously, the particle was aligned along the [111] direction. By using alcohol-based solvents such as ethanol, methanol and n-propanol, nano-sized (~ 20–60 nm) cubic-phase BaTiO3 powders were obtained [76]. However, the tetragonal BaTiO3 nanopowders with sizes of 50-100 nm were synthesized by using EtOH as a solvent [77]. It was found that the particle size was dependent upon the feedstock concentration (e.g., the precursor concentration). With decreasing the particle size from 89 to 58 nm, the amount of the tetragonal phase in the powder was decreased from 85% to 57%, and the cell parameter ratio (c/a) also decreased from 1.0080 to 1.0071. Recently, BaTiO3 nanopowders with sizes down to 5 nm are synthesized by direct reaction between barium hydroxide octahydrate and titanium (IV) tetraisopropoxide under solvothermal conditions (2-methoxyethanol and absolute ethanol, respectively) [78].

2.3.4.3. Glycothermal process The concepts embodied in hydrothermal processing approaches can be extrapolated to non-aqueous systems. By using glycol media (especially 1,4-butanediol solution) instead of water for the hydrothermal reaction, perovskite BaTiO3 nanopowders can be directly synthesized via glycothermal reaction at temperatures lower than that required by the hydrothermal conversion [130]. Glycothermal reaction of metal alkoxide, acetylacetonate, or acetate is a convenient route for the synthesis of crystalline ceramic powders, avoiding the effect of water. It has some novel features different from the conventional hydrothermal technology. First, glycothermal process does not need the mineralizers for the formation of anhydrous crystalline materials in some cases since 1,4-butanediol could act as an oxidizer [130-134]. Therefore, this process prevents the contamination by alkalis and/or halides, which are commonly used in conventional hydrothermal process. Second, glycothermal process significantly reduces the reaction pressure, which is a critical issue for large-scale production. Finally, it is easy to control the size and shape of the synthesized powders in glycol solution without growth-directing agents [134]. The tetragonal BaTiO3 nanoparticles, have been synthesized at temperature as low as 220oC through glycothermal reaction by using Ba(OH)28H2O and amorphous titanium hydrous gel as precursors and mixture of 1,4butanediol and water as solvent [81]. The glycothermal process, provides a simple low temperature route for producing tetragonal BaTiO3 nanoparticles without alkaline mineralizers, and the molar ratio of Ba/Ti, tetragonality, size and morphology of BaTiO3 nanoparticles can be controlled by adjusting the reaction conditions (e.g., reaction temperature, volume ratio of 1,4-butandiol/water, the Ba/Ti molar ratio of precursor) [81].

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Figure 4. (a) HRTEM image of an assembly of BaTiO3 nanoparticles. The inset is the selected area electron diffraction pattern. (c) and (d) HRTEM images of two isolated particles, and (e) the fast Fourier transform pattern obtained from the HRTEM image shown in Figure c [18]. Reproduced with permission from [18], Niederberger, M.; Pinna, N.; Polleux, J.; Antonietti, M. General soft-chemistry route to perovskites and related materials: synthesis of BaTiO3, BaZrO3, and LiNbO3 nanoparticles. Angew Chem. Int. Edt. 2004, 116, 2320-2323. Copyright Wiley-VCH Verlag GmbH & Co. KGaA.

Figure 5. (a) TEM image of BaZrO3 nanoparticles. The inset is the selected area electron diffraction pattern. (b) HRTEM image of an assembly of particles, and (c) and (d) HRTEM of isolated particle and the corresponding fast Fourier transform pattern, respectively. Reproduced with permission from [18], Niederberger, M.; Pinna, N.; Polleux, J.; Antonietti, M. General soft-chemistry route to perovskites and related materials: synthesis of BaTiO3, BaZrO3, and LiNbO3 nanoparticles. Angew Chem. Int. Edt. 2004, 116, 2320-2323. Copyright Wiley-VCH Verlag GmbH & Co. KGaA.

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2.3.4.4. Microwave-hydrothermal process The microwave-hydrothermal process is often found to be rapid, and has the potential to enhance the crystallization kinetics of hydrothermal process. The term microwavehydrothermal process was coined by Komarneni et al.[135] in 1992, and this process has been used for the rapid synthesis of numerous ceramic oxides, hydroxylated phases, porous materials, and hematite powders [64-73, 136-139]. It offers many distinct advantages over the conventional hydrothermal synthesis, such as cost savings due to rapid kinetics time and energy, rapid internal heating and synthesis of new materials. In the microwave-hydrothermal process, the microwave radiation couples with the material, and the electromagnetic energy is converted into thermal energy, which is absorbed by the material. Therefore, the heat is generated from inside the material, in contrast with conventional autoclave heating methods where the heat is transferred from outside to inside. This internal heat allows very rapid heating to the crystallization temperature, faster kinetics of crystallization by one-to-two orders of magnitude compared to the conventional hydrothermal process, and also saves energy and time. In addition, microwave heating is particularly suitable for perovskite nanopowders because the absorption degree of microwaves by them is much high due to their large dielectric constant and high dielectric loss. Numerous reports have published on synthesis of BaTiO3 nanopowders by microwavehydrothermal process below 200 °C, and these processes were found to be very rapid but they all yielded cubic phase [64-66,140,141]. For example, Khollam et al. [66] obtained submicron-sized BaTiO3 powders (0.1- 0.2µm) at holding time of 30 min. One of the first approaches on the synthesis of the nanosized BaTiO3 powders (about 30 nm) at 30 min, was reported by Jhung et al. [142]. Recently tetragonal BaTiO3 powders are synthesized by microwave-hydrothermal method at typical temperature of 240◦C from hydrous titanium oxide and barium hydroxide, in the absence of chloride ions and alkali metal ions to avoid contaminations. The effects of synthesis conditions, including reaction temperature and time, and reactant composition, on the formation of tetragonal structure and particle size of BaTiO3 powders, have been systematically investigated [143]. The results have shown that the amount of the tetragonal phase and the particle size increased quickly with reaction time, whereas the content of lattice hydroxyl groups decreased. Tetragonal BaTiO3 powder with nearly full tetragonallity (c/a ratio = 1.010) was obtained via the microwave-hydrothermal process performed at 240°C for 20 hours [143]. As the reaction temperature was lowered down to 220◦C, the formation of tetragonal structure and the growth of particles slowed down substantially, showing a critical effect of the reaction temperature on the microwavehydrothermal processing of tetragonal BaTiO3. Higher Ba(OH)2/Ti mole ratio enhanced the formation of tetragonal BaTiO3 and so did higher initial concentration of Ti with fixed Ba(OH)2/Ti ratio. Besides the BaTiO3 nanopowders, Ba1-xSrxTiO3 (x = 0.1-0.4) nanopowders with the average size about 20 nm were also prepared at relatively a short period of time (10 min) via microwave-hydrothermal synthesis [144]. The structure and the average sizes of BST were determined to be in range of 20-50 nm depending on the synthesis time (10-90 min). In conclusion, microwave-hydrothermal process has many advantages over the conventional methods [64-73]. Some of these advantages include, time and energy saving, very rapid heating rates (> 400 K/minute) without damage due to thermal shock, considerably reduced processing time and temperature, and fine microstructures. This process is also environmentally friendly.

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2.3.5. Spray pyrolysis Spray pyrolysis represents a continuous and single-step preparation method for the production of fine homogeneous oxide powders [145-149]. The process of this method covers the following two steps: (a) the precursor solution (sol), which contains the metal ions dissolved in the desired stoichiometry, is sprayed through a nozzle and suspended in gaseous atmosphere (aerosol generator); (b) the suspended droplets are thermally processed to the product phase by allowing the sol droplets to drift through the heated zone of a furnace. Spray pyrolysis has many variations based on the differences in thermal processing step. Some of them are aerosol decomposition, evaporative decomposition, spray roasting, and spray calcinations [150]. Since the conventional spray pyrolysis results in multiple nanosized crystallites that are virtually inseparable, so they form a three-dimensional network [151-153]. Salt-assisted spray pyrolysis has been developed as a novel route to the preparation of nanoparticles below 100 nm [154,155]. This route requires no further thermal treatment of the product, such as calcination or annealing, because metal salts enhance the crystal growth and the homogeneity of the crystals. Compared with the sol-gel method and related precipitation techniques, the powders produced by salt-assisted spray pyrolysis are less agglomerated with improved crystallinity. Nanomter-sized perovskite particles with excellent compositional homogeneity can be prepared by this method. For example, highly crystalline, dense BaTiO3 nanoparticles were synthesized by using a salt-assisted spray pyrolysis method without the need for postannealing [156]. The particles ranged in size from 30 to 360 nm, depending on the synthesis temperature, with a narrow size distribution. The particle size decreased with decreasing operation temperature. The crystal phase was transformed from tetragonal to cubic at a particle size of about 50 nm at room temperature. Salt-assisted spray pyrolysis process can be used to produce high weight fraction of tetragonal BaTiO3 nanoparticles down to 64 nm in a single step. Nano-sized BaTiO3 particles were also prepared by citric acid-assisted spray pyrolysis [157]. It was found that controlling the spray solution with an organic additive made great differences in the structure and morphology of BaTiO3 particles during the calcination. The citric acid additives prevented phase separation of barium and produced phase-pure BaTiO3 particles at the as-prepared state and enhanced the phase transformability of metastable cubic phase to the tetragonal one during calcination. Tetragonal BaTiO3 nanoparticles with size of ~ 150 nm were successfully obtained by simple ball milling the coarse aggregates prepared from the citric acid-assisted spray pyrolysis and calcination at 1050oC. 2.3.6. Microemulsion synthesis Microemulsion synthesis is defined as an isotropic, thermodynamically stable system constituting the micrometer-sized droplets (micelle) dispersed in an immiscible solvent and an amphiphilic surfactant species on the surface of the micelle [84,85]. The crucial aspect of the microemulsion route is the control of the nanoparticle size through suitable selection and addition of a surfactant prior to the hydrolysis of the metal alkoxide sol (reverse micelle of water-in-oil emulsion). The addition of the surfactant molecules creates nanosized domains (nanoreactors) in the range of 0.5 - 10 nm spontaneously, in contrast to the conventional milky macroemulsions, which are only kinetically stabilized and in general prepared by the introduction of mechanical energy. The sizes of nanodomains are only dependent upon the

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composition of the microemulsion, temperature, and the elastic properties of the separating surfactant film. In particular, for the case of water/oil microemulsion with spherical nanosized aqueous micelles dispersed in an oil matrix, the aqueous droplets can be used as nanoreactors and templates for the preparation of solid nanoparticles. Since the reaction is spatially initiated and confined in the originally aqueous micelles, heterogeneous nucleation and crystal growth may be controlled. This method has been recognized as the most appropriate method for the synthesis of various electroceramic compounds (e.g., piezoelectrics [86], varistors [158], superconducting oxides [159-161], and magnetics [162-165]) in the form of nanopowders. Hempelmann et al. [166-168] demonstrated the preparation of nanoparticles for perovskite type materials such as BaTiO3. Figure 6 shows a high-resolution TEM micrograph of a single BaTiO3 nanoparticle obtained by microemulsion-mediated synthesis in combination with the particle size distribution determined by small angle X-ray scattering [169]. Powder particles well below 10 nm in size may be obtained by this technique. Furthermore, it allows the preparation of stable dispersions of nanopowders for the preparation of thin dielectric layers.

2.3.7. High-gravity reactive precipitation High-gravity reactive precipitation (HGRP) can be described as the reactive precipitation taking place under high-gravitational conditions [89]. For the HGRP synthesis, the key part of the rotating packed bed (RPB, Higee machine) is a packed rotator, which is designed to generate acceleration higher than the gravitational acceleration on the Earth. Three typical kinds of reaction systems are often used in the particle syntheses by the reactive precipitation. These are the liquid-liquid, gas-liquid, and gas-liquid-solid reactant phase systems. Recently uniform BaTiO3 nanoparticles are produced at a low temperature (1128 K. Some of the internal pores were released from the particle‘s surface and/or during the grain growth. The presence of the pores affected the density of the BaTiO 3 particle. The behavior of the internal pore was observed in situ with increasing temperature on the thermal stage of a TEM device. The results showed that at >1128 K, some pores move out from the particle‘s surface during TEM observation. This temperature roughly agrees with the temperature at which the density of BaTiO3 powder sharply increases. During observation with increasing temperature, a thin layer appeared on the particle‘s surface at temperature over 573 K and then disappeared at 1193 K.

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Figure 8. Transmission electron microscopy images of (a) as-received BaTiO3 powder (particle size ~ 60 nm) and (b) BaTiO3 powder annealed at 673 K. [180]. Nakano, H.; Urabe, K.; Oikawa, T.; Ikawa, H. Characterization of internal pores in hydrothermally synthesized BaTiO3 particle by transmission electron microscopy. J Am Ceram Soc. 2004, 87, 1594-1597, Copyright © 2004, American Chemical Society.

The hydrothermal BaTiO3 powder with a small particle size are stabilized in a cubic phase at room temperature [42,122,179,180-182], which implies that the distortion of the [TiO6] structure resulting in a cubic-to-tetragonal phase transition as cooled the sample through the Curie temperature is not taken place. A plausible reason is that the small size of the BaTiO3 nanocrystals, which are so small that the structural defects in the particles prevent the completion of the structural transition, leading to high strains within the crystals. The high strains inside the nanoparticles introduced by structural defects (e.g. lattice defects), would make the unit cell distortion (c/a ratio) much smaller than that in the standard BaTiO3. To reveal the high strains in the hydrothermal BaTiO3 nanoparticles by TEM images, Zhu et al. [181] recorded both bright- and dark-field TEM images from the hydrothermal BaTiO3 nanoparticles. Figure 9a is a bright-field TEM image recorded by using a small objective aperture that selects only the (000) central transmitted beam, which shows a narrowdistribution spherical nanoparticles. The dark-field image shown in Figure 9b, was recorded by using a smaller objective aperture that selects the part of the {100} and {110} reflections, as indicated by a circle in Figure 9c. The dark-field image displayed in Figure 9b clearly shows high strains in some BaTiO3 nanoparticles. By using the bright- and dark-field TEM images, Lu et al. [182] also reported several types of TEM contrast variations in an individual BaTiO3 nanocrystal synthesized via hydrothermal method at a temperature of 230C. It is believed that the different types of variations of TEM contrast indicate the existence of different strains in BaTiO3 nanograins. Therefore, in a TEM image, large strain is indicated by a contrast variation across a particle. If a particle is single crystalline and has no strain, it should be uniform in contrast. However, for a single crystalline particle, if the TEM image shows dark-bright variation in contrast, it is likely to have a high strain within the grain. Strain affects the diffraction behavior of the electrons, resulting in dramatic contrast change. The hydrothermal BaTiO3 nanoparticles exhibit a cubic structure (a high temperature phase) at room temperature, such an abnormal crystallographic phenomenon is closely related to the existence of high strains in these BaTiO3 nanoparticles. The strains introduced by a high

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concentration of lattice defects such as OH- ions and barium vacancies, can make the unit cell distortion (c/a ratio) much smaller compared with that of the standard BaTiO3. As a result, no peak splitting was detected in the XRD patterns of the hydrothermal BaTiO3 powders even though they belong to the tetragonal phase.

Figure 9. (a) Bright-field and (b) dark-field TEM images recorded from the hydrothermal BaTiO3 nanoparticles. (c) An selected area electron diffraction pattern from the BaTiO3 particles showing a perovskite structure. The circle indicates the size and position of the objective aperture used to record the dark-field image displayed in (b). Reproduced with permission from [181], Zhu, X. H.; Zhu, J. M.; Zhou, S. H.; Liu, Z. G.; Ming, N. B. Hydrothermal synthesis of nanocrystalline BaTiO3 particles and structural characterization by high-resolution transmission. J Cryst Growth. 2008, 310, 434-441. Copyright © 2008 Elsevier B.V. All rights reserved.

Figure 10. Schematic diagrams for the formation of (a) STEM and (b) HRTEM images. Reproduced with permission from Dmitri Klenov, Melody Agustin and Susanne Stemmer (proviate communication).

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Generally, TEM is the most powerful and appropriate technique for investigating the characteristics of nanoscale particles, however, for detecting the low elemental concentrations that are typically of environmental interest, in which almost the entire concentration of the trace metal is located in just a few nanoparticles, it is almost impossible to find the nanoparticles that contain the metals of interest. When the elemental distribution is widely scattered, the image contrast in the conventional TEM is minimal except under very high magnification, where only a limited number of particles can be examined. However, as the particles of interest consist of relatively heavy elements, as compared with the matrix material, high-angle annular dark field scanning TEM (HAADF-STEM) is a powerful method for finding these nanoparticles of interest, details are described below.

3.1.3. Scanning transmission electron microscopy (STEM) STEM images (also called as Z-contrast incoherent images, or HAADF images) with atomic-resolution, are formed by using incoherent elastically-scattered electrons, as schematically shown in Figure 10a. Normally, a STEM image is formed by collecting highangle (75-150 mrad) elastically-scattered electrons with an annular dark-field detector (see Figure 10a). Such an annular detector captures a large fraction of the high angle intensity, providing an efficient dark-field imaging mode. The simplicity of STEM image is a direct result of the fact that only electrons scattered through large angles are used to form the image, so that interference effects contribute less to the image. The STEM image is consequently far less sensitive (although not immune) to specimen thickness variations, tilt and defocus. In the STEM operational mode, the electron beam is focused to a very fine spot (as small as 0.1 nm or less). By scanning this fine electron beam in a raster across the specimen and collecting the transmitted or scattered electrons, STEM images can be formed. Since the maximum scattering occurs when the electron probe is centered over a column of atoms, the columns appear bright in the image, whereas little scattering occurs when the probe is centered over a channel between atomic columns, therefore these areas appear dark. In a perfect crystal, atoms with higher atomic number, Z, produce more scattering, so that the intensity of the bright spots in the image can be related to the atomic composition of the corresponding column of atoms in the sample. STEM images can be interpreted more directly in terms of atom types and positions. Unlike conventional TEM, HAADF-STEM is based on imaging the incoherent scattering, and the contrast of the image is not reversed by defocusing above and below the point of ―just focus‖[185,186]. As the samples of environmental interest that contain nanoparticles with relatively heavy elements as compared with the matrix material, the contrast of HAADFSTEM image is strongly correlated with atomic number and specimen thickness, which is an appropriate method for finding the nanoparticles of interest [187]. An example for this is given in Figure 11, showing a STEM image of Bi-doped Si bulk crystal viewed from [110] direction [188], which reveals the columns containing individual Bi atoms introduced by ion implantation followed by re-crystallization through solid phase epitaxial growth. Single Bi atoms on lattice sites within the crystal are clearly visible. The density of bright spots correlates with the known dose of the doped Bi atoms. Similarly, in the rare-earth metal ionsdoped perovskite nanopowders, it is also possible to identify their substituted positions in the perovskite structure by using high-resolution HAADF-STEM images. In the near future we will see rapid progress in this direction.

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Figure 11. A STEM image of Bi-doped Si sample viewed from [110] direction. Reproduced with permission from [188], Pennycook, S. J.; Lupini, A. R.; Kadavanich, A.; McBride, J. R.; Rosenthal, S. J.; Puetter, R. C.; Yahil, A.; Krivanek, O. L.; Dellby, N.; Nellist, P. D.; Duscher, G.; Wang, L. G.; Pantelides, S. T. Aberration-corrected scanning transmission electron microscopy: the potential for nano-and interface science. Z Metallkd. 2003, 94, 350-357. Copyright © 2003, Carl Hanser Publisher.

3.1.4. High-resolution transmission electron microscopy (HRTEM) High-resolution TEM (HRTEM) images are formed by using nearly parallel electron beam traveling through the sample, and the direct (transmitted) beam and the diffracted beams are allowed to interfere with one another to form a ―lattice‖ image, as schematically shown in Figure 10b. The image process of a HRTEM involves the following three processes: (a) electron scattering in a specimen; (b) formation of diffracted beams at the back focal plane; (c) formation of a high-resolution image at the image plane. HRTEM images are uniquely capable of providing the information about local atomic structures, which are most useful for identifying individual defects in nanocrystals, studying the atomic arrangements at interface between heterostructures. The formation of an HRTEM image is required to use an aperture large enough to include both the transmitted beam and at least one diffraction beam, in which the transmitted (actually, forward-scattered) beam provides a reference phase of the electron wavefront. As a result, HRTEM images in nature, are interference patterns between the forward-scattered and diffracted electron waves from the specimen. For HRTEM image, the great problem is the identification of the atomic species in the image, due to the inversion of image contrast, which is closely related to the specimen thickness, objective lens defocus, and additional interference effects (Fresnel fringes) at the interface between the crystalline substrate and the amorphous dielectric, in non-intuitive ways. Quantitative interpretations of HRTEM images are simple only as the sample is a weak-phase object (WPO) and the microscope is at Scherzer defocus. In this case, the contrast is related directly to the projected potential of the specimen. That means the atom columns appear as dark spots on a bright background, and the darkness of the spots is proportional to the project potential of the specimen. However, the WPO approximation is especially difficult to satisfy because the

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specimen must be thin enough so that the phase change of Bragg-scattered electrons is small relative to the forward-scattered electrons. Actually, only the thin specimens with low atomic numbers are likely to behave as WPOs. However, in most cases, materials have projected potentials strong enough to cause the phases of the Bragg-diffracted beams to change rapidly with the depth in the specimen. Therefore, most specimens do not behave as a WPO beyond a few nanometers in the thickness direction. Due to the ability of revealing the local atomic structures, HRTEM image is the most useful and appropriate technique for identifying structural defects in perovskite nanocrystals. As an example, the microstructural defects such as anti-phase boundaries (APBs) and edge dislocations in hydrothermal BaTiO3 nanoparticles were revealed by HRTEM images at atomic-level [181]. Figure 12a and Figure 12b show the HRTEM images of two individual nanoparticles, respectively. The (100) and (111) lattice fringes are clearly observed in Figures 12a and Figure 12b, respectively. The corresponding Fourier filtered images and the FFT patterns (see insets) of the selected areas marked by boxes in the two HRTEM images, are shown in Figure 12c and Figure 12d, respectively. The Fourier filtered images clearly demonstrates that how the APBs are formed during the particle growth. There are two crystalline regions marked by I and II in the ellipses with fine white line in Figure 12c and Figure 12d. The two crystalline regions in Figure 12c are deviated from each other by a relative displacement of 1/2d100 (d100: the inter-planar distance between two adjacent (100) planes), whereas in Figure 12d, the relative displacement is 1/2d111 (d111: the inter-planar distance between two adjacent (111) planes). During the solid-phase crystallization, the crystalline growth fronts of part I and part II intersect each other. Then, the intersection of growth fronts accommodates the deviation as APBs. The observed APBs near the edge of a BaTiO3 nanoparticle, were formed by the intersection of two crystalline parts with displacement deviation from each other by 1/2d100 or 1/2d111, as revealed by the HRTEM images. Similar conditions were also observed in a crystalline SrBi2Ta2O9 grain [189].

Figure 12. (a) and (b) HRTEM images of two isolated BaTiO3 nanoparticles. (c) and (d) The corresponding Fourier filtered images and the FFT patterns (see insets). Reproduced with permission from [181], Zhu, X. H.; Zhu, J. M.; Zhou, S. H.; Liu, Z. G.; Ming, N. B. Hydrothermal synthesis of nanocrystalline BaTiO3 particles and structural characterization by high-resolution transmission. J Cryst Growth. 2008, 310, 434-441. Copyright © 2008 Elsevier B.V. All rights reserved.

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Figure 13. High-resolution transmission electron microscopy images of the surface structures at the edges of BaTiO3 nanoparticles viewed from the [001] direction. (a) Both a terrace-ledge-kink (TLK) surface structure and small nucleated and triangular islands with two to three atomic layer thickness are observed. (b) and (c) TLK surface structure with both terraces and ledges lying on the {100} planes; only a small amount of ledges lie on the (110) plane. The inset in (b) is a Fourier-filtered image of the corresponding position, which clearly demonstrates two perpendicular sets of (100) and (010) planes. Reproduced with permission from [183], Zhu, X. H.; Wang, J. Y.; Zhang, Z. H.; J. M.; Zhou, S. H.; Liu, Z. G.; Ming, N. B. Atomic-scale characterization of barium titanate powders formed by the hydrothermal process. J Am Ceram Soc. 2008, 91, 1002-1008. Copyright © 2008, American Ceramic Society.

Figure 14. A surface profile HRTEM image of part of a hydrothermal BaTiO3 particle with size of 80 nm. Reproduced with permission from [184], Zhu, X. H.; Zhu, J. M.; Zhou, S. H.; Liu, Z. G.; Ming, N. B.; Hesse, D. BaTiO3 nanocrystals: hydrothermal synthesis and structural characterization. J Cryst Growth. 2005, 283, 553-562. Copyright © 2005 Elsevier B.V. All rights reserved.

A terrace-ledge-kink (TLK) surface structure was also frequently observed at the edges of the hydrothermal BaTiO3 nanoparticles with rough surface morphology, and in most cases the terrace and ledge lie on the {100} planes [183] . The observed TLK surface structure is shown in Figure 13, which can be well interpreted by the theory of periodic bond chains. Small nucleated and triangular BaTiO3 islands with 3 ~ 4 atomic layer highness, and their outside surfaces faceted as (100) and (010) planes, were also observed in hydrothermal BaTiO3 nanoparticles, as indicated by arrows in Figure 14 [184]. The rarely-seen {110} surface in the

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BaTiO3 nanoparticles were found to be reconstructed so that the surface was composed of corners bound by {100} mini-faces like the triangular small islands. Internal defect textures, such as nanoscale multiple (111) twining and complicated (111) intergrowth defects, were also observed in the BaTiO3 nanopowders synthesized by sol-gel and stearic acid-gel (SAG) methods. They were identified as hexagonal-type BaTiO3 structure [190,191]. Complex arrangements of defects lying on the (111) planes were observed in the SAG-derived BaTiO3 nanocrystal with particle size of 10 nm. The density of the small defects was estimated to be on the order of 1027/m3 in the SAG-derived BaTiO3 nanopowders. These high density of defects could result in the cubic phase structure of SAG-derived BaTiO3 powders even with grain size large up to 3.50µm [190].

3.1.5. Spherical-corrected HRTEM/STEM Traditionally, spherical aberration (Cs) of magnetic lenses limits the resolutions of HRTEM and STEM images. In recent years spherical aberration correctors (e.g., hexapole type Cs-correctors proposed by Rose [192]) have been developed to reduce substantially the effective value of Cs of the objective lens. Its main idea is that multipole lenses such as quadrupoles, sexupoles, and octupoles have lens aberrations with different phase shift errors, W(k), compared with the short solenoids used for the objective lens. Combining these different functional forms makes it possible to make the overall W(k) a more constant function. This is accomplished by placing a set of different lenses along the optical path and tuning their currents. Recently researchers from IBM Thomas J. Watson Research Center and Nion R&D, have taken a step towards reaching the ultimate resolution (sub-angstrom resolution) in an electron microscope [193] by implementing a computer-controlled aberration correction system in a STEM that is less sensitive to the remaining chromatic aberrations. The Cs-corrected STEM mode can provide a sub-angstrom probe with a highbrightness, which offers the prospect of element-selective imaging of single atomic columns using the energy filter. Combined with monochromated HR-EELS, one can further investigate chemistry and electron structure-related properties (e.g., valence state, bonding structure) by single atomic column to column. The Cs-corrected HRTEM mode offers a tunable spherical aberration coefficient from negative to positive values. Properly combining a negative Cs with a positive defocus, at no cost to point resolution, an HRTEM image with bright-contrast of atoms on dark background can be obtained, which can be directly interpreted without image simulation, and light elements such as oxygen atoms and even their vacancies can also be imaged [194-197]. For example, by using the Cs-corrected imaging technique, Jia et al. [197] first performed the atomic-sacle investigations of the electric dipoles near (charged and uncharged) 180 domain walls in thin epitaxial PbZr0.2Ti0.8O3 film sandwiched between two SrTiO3 layers. Figure 15 is an atomic-scale image of the electric dipoles formed by the relative displacements of the Zr/Ti cation columns and the O anion columns in PbZr0.2Ti0.8O3 film, viewed from the [110] direction and recorded under negative spherical-aberration imaging conditions. The local tetragonality c/a and spontaneous polarization inside the domains and across the domain wall were calculated. For the first time, a large difference in atomic details between charged and uncharged domain walls was reported. Such breakthrough would improve our ability to see and thoroughly explore the properties of perovskite nanopowders.We can foresee that the new Cs-corrected HRTEM and STEM will benefit perovskite nanopowder materials research in the new era.

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Figure 15. Atomic-scale imaging of the electric dipoles formed by the relative displacements of the Zr/Ti cation columns and the O anion columns in the approximately 10-nm-thick PbZr0.2Ti0.8O3 layer

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

sandwiched between two SrTiO3 layers. The image was viewed along the [ 110 ] direction and recorded under negative spherical-aberration imaging conditions. The atom columns appear bright on a dark background. The horizontal arrows denote the horizontal interfaces between the PbZr0.2Ti0.8O3 film and the top and the bottom SrTiO3 film layers. The dotted line traces the 180 domain wall between the domain I and domain II. The arrows denoted by ‗P S‘ show the directions of the polarization in the 180 domains. Two insets show higher magnifications of the dipoles formed by the displacements of ions in the unit cells. Yellow symbols denote PbO atom columns seen end-on, red symbols for Zr/Ti columns, and blue symbols for oxygen. Reproduced with permission from [197], Jia, C. L.; Mi, S. B.; Urban, K.; Vrejoiu, I.; Alexe, M.; Hesse, D. Atomic-scale study of electric dipoles near charged and uncharged domain walls in ferroelectric films. Nat Mater. 2008, 7, 57-61. Copyright © 2008, Nature Publishing Group.

3.2. Spectroscopic Characterization 3.2.1. X-ray diffraction (XRD) In X-ray diffraction (XRD), a collimated beam of X-rays (wavelength : 0.5 -2 Å ) is incident on a specimen and is diffracted by the crystalline phases in the specimen according to Bragg‘s law (2dsinθ = , where d is the spacing between atomic planes in the crystalline phase). The intensity of the diffracted X-rays is measured as a function of the diffraction angle 2θ and the specimen‘s orientation. As a primary characterization tool for obtaining the critical features such as crystal structure, crystallite size, and strain, X-ray diffraction patterns have been widely used for perovskite nanopowder research [198-200]. Except for single crystalline nanopowders, the randomly oriented crystals in nanopowders cause broadening of

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the diffraction peaks due to the absence of the total constructive and destructive interferences of X-rays in a finite-sized lattice [201]. This effect becomes more pronounced when the crystallite sizes are in the order of a few nanometers. In addition, inhomogeneous lattice strains and structural faults also lead to the broadening of peaks in X-ray diffraction patterns. Currently, the widely used and simplest method for estimating the crystallite size is based on the Scherrer equation, which can be expressed as [202] 3 d

K  cos 

(2)

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where d is the crystallite size,  is the wavelength of the used X-ray,  is the FWHM (full width at half maximum of a diffraction peak), θ is the diffraction angle, and K is a constant close to unity. The major assumptions are that the sample is free of residual strain and has a narrow grain size distribution. As one example, the crystallite size of PbTiO3 nanopowders obtained from combined polymerisation and pyrolysis route, can be determined based on the obvious broadening of the Bragg reflections or, more precisely, from the line shape of diffractions peaks [203]. However, when the contributions due to strains are taken into consideration, the analysis becomes much more complicated.

3.2.2. Extended X-ray absorption fine structure spectroscopy (EXAFS) Extended X-ray absorption fine-structure (EXAFS) measurements, which are oscillations occurring on the high-energy side of an X-ray absorption edge, can be used to identify interatomic distances in materials [204]. An EXAFS experiment involves the irradiation of a sample with a tunable source of monochromatic X-rays from a synchrotron radiation facility. As the X-ray energy is scanned from just below to well above the binding energy of a coreshell electron (e.g., K or L) of a selected element, the X-ray photoabsorption process is monitored. When the energy of the incident X-rays is equal to the electron binding energy, Xray absorption occurs and a steeply rising absorption edge is observed. For energies greater than the binding energy, oscillations of the absorption with incident X-ray energy (i.e., EXAFS) are observed. EXAFS data are characteristic of the structural distribution of atoms in the immediate vicinity of the X-ray absorbing element. The present consensus is that the accuracy of interatomic distance determined by EXAFS is between 0.01 Å and 0.001 Å , depending on the circumstances [205,206]. However, in certain types of study, such as the differential magnetostriction measurements reported by Pettifer et al. [207], direct comparison should be capable of much higher precision. This is especially true when synchrotron radiation is used from a third-generation source, where the fluxes available can be 1013 photons s-1 eV-1. These should be able to produce relative statistical errors in the absorption spectrum of ~10-6 under optimal conditions in a few hours. With such a statistical accuracy, interatomic strains of the order of femtometres and below should be detectable with EXAFS [207]. The frequency of the EXAFS is related to the interatomic distance between the absorbing and neighboring atoms. The amplitude of the EXAFS is related to the number, type, and order of neighboring atoms. To probe the local structure in perovskite nanopowders, EXAFS has been proven to be an effective technique. Recently, Frenkel et al. [208] carried out EXAFS study on BaTiO3 particles with different average grain sizes of about 20 nm, 35 nm, and 70 nm prepared by

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solution-gelation method. The normalized XAFS spectra were obtained by subtracting the background µ0(k) from the measured absorption coefficient µ(k). The k2 - weighted (k) of the samples with different particle size at room temperature, as well as the data with 10µm particles measured at different temperatures were shown in Figure 16. It was observed that besides a very small difference in amplitude, room temperature EXAFS (k) of all the samples with different particle size did not show significant changes. The changes in the amplitude in Figure 16a were larger than the statistical noise between two measurements of the same sample. However, these changes were smaller than those occurred during heating the sample with a 10µm particle size from 80K to 590K (Figure 16b). In this temperature range, the average structure of BaTiO3 exhibits several phase transitions from rhombohedral to orthorhombic to tetragonal to cubic at elevated temperatures. The fact that the changes between the EXAFS signals measured for different particle sizes were smaller than the changes occurred in the sample with a macroscopic particle size at different temperatures, indicated that the local structure of the samples with all the particle sizes measured was essentially the same within the experimental resolution. The magnitude of the Ti atom offcenter displacement did not depend on the particle size. Petkov et al. [209] have recently demonstrated the use of the pair distribution function (PDF) to understand local structure distortions and polar behavior in BaxSr1-xTiO3 (x = 1, 0.5, 0) nanocrystals. They found that locally, refining over the first 15Å , the tetragonal model was the best fit to the experimental PDF; however, over longer distances (15-28 Å ), the cubic model was the best fit. Their conclusion was that 5 nm BaTiO3 was on average cubic, but that tetragonal-type distortions in the Ti-O distances are present within the cubic structure.

Figure 16. k2 - weighted (k) for the samples with (a) different particle sizes at 300 K and (b) 10 µm particle size at different temperatures. Reproduced with permission from [208], Frenkel, A. I.; Frey, M. H.; Payne, D. A. XAFS analysis of particle size effect on local structure in BaTiO3. J Synchrotron Rad.1999, 6, 515-517. Copyright © 1999, International Union of Crystallography.

3.2.3. Energy dispersive X-ray spectroscopy (EDS) Energy dispersive X-ray spectroscopy (EDS) is an analytical technique used for the elemental analysis or chemical characterization of a sample. As a type of spectroscopy, it relies on the investigation of a sample through interactions between electromagnetic radiation and matter, analyzing X-rays emitted by the matter in response to being hit with charged particles. Its characterization capabilities are due in large part to the fundamental principle that each element has a unique atomic structure allowing X-rays that are characteristic of an

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element's atomic structure to be identified uniquely from each other. An EDS system setup consists four primary components: the beam source; the solid state X-ray detector usually made from lithium-drifted silicon, Si (Li); the pulse processor; and the analyzer. A detector is used to convert X-ray energy into voltage signals; this information is sent to a pulse processor, which measures the signals and passes them onto an analyzer for data display and analysis. Accuracy of EDS spectrum are affected by many variants. Windows in front of the Si(Li) detector can absorb low-energy X-rays (a.k.a. EDS detectors cannot detect presence of oxygen, carbon, boron, etc.). Differing the over-voltage of the EDS will result in different peak sizes - Raising over-voltage on the SEM will shift the spectrum to the larger energies making higher-energy peaks larger while making lower energy peaks smaller. Also many elements will have overlapping peaks (e.g., Ti Kβ and V Kα, Mn Kβ and Fe Kα). The accuracy of the spectrum can also be affected by the nature of the sample. X-rays can be generated by any atom in the sample that is sufficiently excited by the incoming beam. These X-rays are emitted in any direction, and so may not all escape the sample. The likelihood of an X-ray escaping the specimen, and thus being available to be detected and measured, depends on the energy of the X-ray and the amount and density of material it has to pass through. This can result in reduced accuracy in inhomogeneous and rough samples. The main use of EDS is to accurately determine the composition of the sample under investigation. While the TEM images provide real-time pictures of the size and morphology of the nanopowders, the supplementary EDS analysis provides exact composition of the sample. Several examples [210-213] have demonstrated the use of EDS in analysis of oxide perovskite nanopowders, particularly in the determination of the composition of substituted or nanoparticles composite materials.

3.2.4. Electron energy loss spectroscopy (EELS) Electron energy-loss spectroscopy (EELS) based on electron microscopy is a powerful method for investigating electronic structures of nanometer-scale materials. In which a nearly monochromatic beam of electrons is directed through an ultra-thin specimen, usually in a TEM or STEM electron microscope. As the electron beam propagates through the specimen, it experiences both elastic and inelastic scattering with the constituent atoms, which modifies its energy distribution. Each atomic species in the analyzed volume causes a characteristic change in the energy of the incident beam; the changes are analyzed by means of an electron spectrometer and counted by a suitable detector system. The intensity of the measured signal can be used to determine quantitatively the local specimen concentration, the electronic and chemical structure, and the nearest neighbor atomic spacings. The signal in EELS is in the form of ionization edges on a large background. Determinations of chemical concentrations involve a back-ground subtraction to isolate the intensity of the absorption edge (EELS). These isolated intensities are then compared for the different elements in the spectrum, and in many cases are converted into absolute concentrations by use of appropriate constants of proportionality. The accuracy of quantification depends on the reliability of these constants, so significant effort has been devoted to understanding them. More than this, fine details in the EELS spectra can often provide insight into electronic structures. For the state-of-the-art field-emission TEM, an essential feature is its ability to form a nanometer-sized electron probe, which allows for the acquisition of EDS and EELS spectra. This feature for simultaneous structure, composition

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and bonding information at each location is a powerful combination for understanding the structure and chemistry of perovskite nanopowders. Suzuki et al. [214] have used high energy-resolution EELS (energy resolution ~ 0.2eV) to investigate the electronic structures of BaTiO3 nanocrystals synthesized by chemical vapor deposition. The valence excitation spectra of BaTiO3 nanocrystals with average diameters of 34 nm and 6 nm in an energy range from 2 to 40 eV were shown in Figure 17, demonstrating that the onset energies of spectral intensities were 3.2 eV for 34 nm BaTiO3 nanocrystals and 3.5 eV for 6 nm BaTiO3 nanocrystals. This indicated an increase in the bandgap energy of BaTiO3 with a decrease in crystal sizes. Those onset energies obtained from 90 nm specimen areas showed an excellent agreement with those estimated by previously reported optical measurements. Volume plasmon peaks were observed at 26.5eV in 34 nm BaTiO3 nanocrystals and 25eV for 6 nm BaTiO3 nanocrystals. Dielectric functions of the BaTiO3 nanocrystals derived from loss functions by Kramers-Kronig analysis shows not only an increase in the O 2p  Ti 3d (t2g) transition energy, but also a decrease in the peak energy which corresponds to the O 2p Ti 3d (eg) transition. These results show that high energyresolution EELS based on TEM, which provides information of electronic structures from specified small specimen areas, is powerful tool not only for the characterization of new materials but also for the basic research of electronic structures of quantum objects.

3.2.5. X-ray photoelectron spectroscopy (XPS) In X-ray photoelectron spectroscopy (XPS) monoenergetic soft X-rays bombard a sample material, causing electrons to be ejected. Identification of the elements present in the sample can be made directly from the kinetic energies of these ejected photoelectrons. On a finer scale it is also possible to identify the chemical state of the elements present from small variations in the determined kinetic energies. The relative concentrations of elements can be determined from the measured photoelectron intensities. For a solid, XPS probes 2-20 atomic layers deep, depending on the material, the energy of the photoelectron concerned, and the angle (with respect to the surface) of the measurement. XPS is one of the important characterization tools for surface chemical analysis, which has the ability to probe the surface to a few atomic layers deep (0.5 - 5 nm) and obtain a semi-quantitative elemental analysis of surfaces without standards. The valence states of the elements that constitute the surfaces can also be deduced from the XPS characterization. Since the surface chemistry of perovskite nanopowders plays an important role in affecting their densification behavior during the sintering process, a knowledge of their surface chemistry is highly necessary [215,216]. In the case of BaTiO3 nanopowders, the surface barium carbonate (BaCO3) formed by reaction of BaTiO3 with atmospheric and/or solvated carbon dioxide (CO2), or stemmed as a residual from the powder synthesis process, can have a significant effect on the sintering behavior of the nanopowders [216,217]. Bu using XPS numerous investigations have been carried out to analyze the surface chemistry and surface phases of a variety of commercial and laboratory-synthesized BaTiO3 nanopowders [216,218-223]. All these investigations revealed a small contribution to the barium photoemission signal which was invariably attributed to the presence of surface BaCO3; however, only in two instances was this barium contribution accompanied by carbon and oxygen signals which could be positively matched with the carbonate [221,223]. None of these studies considered the presence of water adsorbed at the surface of the powders.

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Bearing these factors in mind, Wegmann et al. [224] characterized the commercial submicrometer BaTiO3 powders by using XPS. Their results showed that the powder particle surfaces were hydrated with physisorbed molecular water and chemisorbed hydroxyl groups. The hydrated surface proved to be stable under the ultra-high vacuum conditions experienced during the XPS experiment and could only be removed by Ar-ion sputtering at 400°C. This behavior suggested that the powder surfaces were thoroughly hydrated under processing conditions experienced during many ceramic forming procedures and that consequently processing aids (e.g., solvents, dispersants and binders) interacted with the adsorbed water layer(s) rather than directly with the ceramic surface. XPS depth profiling by Ar-ion sputtering revealed the powder surfaces to be Ti-rich, confirming the presence of a phase, or phases, to stoichiometrically balance the barium carbonate.

Figure 17. Valence electron excitation spectra of BaTiO3 nanocrystals (BTNCs) with average particle sizes of 6 and 34 nm in an energy range from 2 to 40 eV ( the local spectra in the energy loss from 1 to 7 eV, seen in inset (c)). The insets (a) and (b) are TEM images of BTNCs with average diameters of 34 nm and 6 nm, respectively.Vertical lines in inset (c) indicated the onsets of spectral intensities. Reproduced with permission from [214], Suzuki, K.; Terauchi, M.; Uemichi, Y.; Kijima, K. High energy-resolution electron energy-loss spectroscopy study of electronic structures of barium titanate nanocrystals. Jpn J Appl Phys. 2005, 44, 7593-7597. Copyright © 2005, the Japan Society of Applied Physics.

3.2.6. Infrared (IR) spectroscopy Infrared spectroscopy (IR spectroscopy) is the subset of spectroscopy that deals with the infrared region of the electromagnetic spectrum. It exploits the fact that molecules have specific frequencies at which they rotate or vibrate corresponding to discrete energy levels (vibrational modes). These resonant frequencies are determined by the shape of the molecular potential energy surfaces, the masses of the atoms and, by the associated vibronic coupling. In order for a vibrational mode in a molecule to be IR active, it must be associated with changes in the permanent dipole. The infrared spectrum of a sample is collected by passing a beam of infrared light through the sample. Examination of the transmitted light reveals how much energy was absorbed at each wavelength. This can be done with a monochromatic beam, which changes in wavelength over time, or by using a Fourier transform instrument to measure all wavelengths at once. From this, a transmittance or absorbance spectrum can be

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produced, showing at which IR wavelengths the sample absorbs. Analysis of these absorption characteristics reveals details about the molecular structure of the sample. This technique works almost exclusively on samples with covalent bonds. Simple spectra are obtained from samples with few IR active bonds and high levels of purity. More complex molecular structures lead to more absorption bands and more complex spectra. This technique has been used for the characterization of structural defects in perovskite nanopowders. For example, in the hydrothermal BaTiO3 nanopowders, IR spectroscopy recorded from such powders at different temperatures indicated the presence of internal OH- groups [118,119], which was evidenced by an OH- stretching band at 3200-3600 cm-1. The concentration of internal OH- decreased with increasing temperature at which the hydrothermal synthesis was performed. The defect chemistry explaining the incorporation of chemisorbed water into the crystal structure of hydrothermal BaTiO3, was studied by IR spectroscopy using deuterated powders in order to better distinguish chemically bound water from the adsorbed moisture. It has been found that the OH- groups form hydroxide ions on the regular oxygen sites of the perovskite, which are compensated by the formation of acceptor-type metal vacancies such as barium and titanium vacancies [122,123,129]. Both lattice hydroxyl group and lattice vacancies affect the magnitude of the tetragonal distortion of BaTiO3 powders. Although the high concentration of lattice defects in hydrothermal BaTiO3 powders, like lattice hydroxyl group and barium vacancies, do not completely prevent the cubic-to-tetragonal phase transformation as supposed by the ―lattice defects‖ theory, the strains introduced by the lattice defects make the unit cell distortion (c/a ratio) become much smaller than that in the standard BaTiO3. That is the reason why no splitting of diffraction peaks was observed in the XRD patterns of the hydrothermal BaTiO3 powders even though they are in tetragonal phase.

3.2.7. Raman spectroscopy Raman spectroscopy is a spectroscopic technique used in condensed matter physics and chemistry to study vibrational, rotational, and other low-frequency modes in a system. It relies on inelastic scattering, or Raman scattering, of monochromatic light, usually from a laser in the visible, near infrared, or near ultraviolet range. The laser light interacts with phonons or other excitations in the system, resulting in the energy of the laser photons being shifted up or down. The shift in energy gives information about the phonon modes in the system. In molecules, a molecular polarizability change, or amount of deformation of the electron cloud, with respect to the vibrational coordinate is required for the molecule to exhibit the Raman effect. The amount of the polarizability change will determine the Raman scattering intensity, whereas the Raman shift is equal to the vibrational level that is involved. Typically, a sample is illuminated with a laser beam. Light from the illuminated spot is collected with a lens and sent through a monochromator. Wavelengths close to the laser line, due to elastic Rayleigh scattering, are filtered out while the rest of the collected light is dispersed onto a detector. Raman spectroscopy is the measurement, as a function of wavenumber, of the inelastic light scattering that results from the excitation of vibrations in molecular and crystalline materials. Raman Spectroscopy is sensitive to molecular and crystal structure, which has been extensively used for structure, composition, and phase characterization of materials. They can provide various characteristic vibrational frequencies, such as those associated with lattice defects or surfaces, and derive crucial data on the electronic band structures in solids. As particle sizes in the nanoscale range, new phenomena appear due to the effects of phonon confinement. These include mode wavenumber shifts and

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line broadening, the appearance of zone boundary phonons, surface phonons, and extremely low wavenumber bands due to excitation of bulk resonances of the particle. One must distinguish between purely nanocrystal effects due to phonon confinement and the increased structural disorder that often accompanies extreme reduction in particle size. Application of Raman spectroscopy as a characterization tool requires careful distinction between modifications of Raman line shape due to disorder in the bulk crystal from modifications due to particle size [225]. Recently, a comprehensively review about Raman spectroscopy of nanomaterials (how Raman spectra related to disorder, particle size and mechanical properties) contributed by Gouadec and Colomban is also available [226]. Up to date, Raman spectroscopy has been widely used to probe the local structure and phase transition of perovskite nanopowders (e.g., BaTiO3 [48,78,182,198,227-229], PbTiO3 [226,230-232], SrTiO3 [233-236]) due to its high sensitivity to the lattice vibrations and dynamics, providing important information about the structure, composition, strain, defects, and phase transitions. Based on the crystallography, there are four triply degenerate optical modes of vibration (3F1u+1F2u) in cubic BaTiO3, which is the Pm3m group. When BaTiO3 transforms into the 4mm group tetragonal phase, each of the F1u modes which are infrared active and Raman inactive splits into modes of symmetry A1+E, while F2u modes which are infrared and Raman inactive split into modes of symmetry B1+E. A1 (non-degenerate), B1 (non-degenerate), and E (doubly degenerate) modes are Raman active. There is further splitting of the vibrational modes because of long-range electrostatic forces associated with lattice ionicity. As a consequence, A1 splits into A1(TO1), A1(TO2), A1(TO3), A1(LO1), A1(LO2), and A1(LO3) and E mode splits into E(TO1), E(TO2), E(TO3), E(LO1), E(LO2), and E(LO3), respectively. The Raman scattering process includes first-order scattering, which obeys the above-stated selective rules and involves one phonon Raman scattering. The second or higher order scattering, which does not obey the above-mentioned selective rules and involves multiphonon Raman scattering, forms combination bands or overtone bands. The wave number of combination bands varies continuously, and that of overtone bands is the multiplicity of the first-order bands. When BaTiO3 undergoes from tetragonal structure into cubic structure, the first-order band should decrease gradually. With respect to the secondorder bands, the case is different and the line frequency cannot be accurately estimated [237]. According to the selection rules, all of the optic modes of BaTiO3 with perfect cubic symmetry should be Raman inactive while the same for the polar tetragonal and orthorhombic polymorphous forms should be Raman active [228]. Therefore, in the Raman spectra of the bulk BaTiO3, sharp bands are around 175 cm-1 [A1(TO), E(LO)] and 305 cm-1 [B1, E(TO + LO)] and broad bands around 265 cm-1 [A1(TO)], 520 cm-1 [A1, E(TO)], and 720 cm-1 [A1, E(LO)] are the characteristic peaks of tetragonal phase BaTiO3. Figure 18 shows the crystal structure of as-prepared and heat-treated hydrothermal BaTiO3 powders examined by Raman spectroscopy[238]. The Raman peak at 305 cm−1, the characteristic peak of the tetragonal phase in BaTiO3, is present in all powders, indicating the presence of the tetragonal phase. There is no significant change in the Raman spectra except that there is an intensity jump at 1000oC. This might be related to the migration of lattice defects, because it is believed that 1000oC is the lowest temperature for the multitude migration of lattice defects such as barium and oxygen vacancies [56,129]. Yashima et al. [239] also investigated the size effect on the crystal structure of BaTiO3 nanoparticle with sizes of 40 - 430 nm by Raman spectra along with neutron and high-resolution synchrotron X-ray powder diffraction techniques. They found that the axial ratio c/a of tetragonal BaTiO3 decreased with a decrease in particle size

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from 430 to 140 nm. Barium titanate particles with a size of 40 nm consisted of (a) tetragonal crystals (83 wt %) with a large cell volume and an axial ratio of unity c/a=1.000(5) and of (b) a hexagonal phase (P63mmc, 17 wt %) with a large unit-cell volume. The ferroelectric phase transitions in nanocrystalline PZT with grain sizes of 60 and 40 nm were investigated by Raman spectra [240]. The results show that the E(TO1) phonon mode of PbTiO3 (―soft‖ mode) displays a decrease in frequency and an increase in line width with increasing Zr concentration. A discontinuous behavior in the phonon energy for the soft mode occurs at a morphotropic phase boundary (MPB) of x  0.4 and 0.2 for grain sizes of 60 and 40 nm, respectively, and it can be attributed to a phase transition from ferroelectric tetragonal to ferroelectric rhombohedral phase. The nonzero soft mode frequency near the MPB results from a level repulsion between an additional phonon mode at 10 cm-1 and the soft mode. Raman enhanced behavior was observed for the lowest phonon mode with Zr contents in the range of 0.3 to 0.6. The dependence of Raman phonon modes for PbZr0.3Ti0.7O3 upon grain size indicated a grain-size-induced phase transition at about 13 nm [240].

3.2.8. Secondary ion mass spectroscopy (SIMS) Secondary ion mass spectrometry (SIMS) is a technique used to analyze the composition of solid surfaces and thin films by sputtering the surface of the specimen with a focused primary ion beam and collecting and analyzing ejected secondary ions. In the field of surface analysis, SIMS is usually classified into static SIMS and dynamic SIMS [241]. Static SIMS is the process involved in surface atomic monolayer analysis, usually with a pulsed ion beam and a time of flight mass spectrometer, while dynamic SIMS is the process involved in bulk analysis, closely related to the sputtering process, using a DC primary ion beam and a magnetic sector or quadrupole mass spectrometer. In the static SIMS, the bombarded particles with an energy of typical 1-10 keV, are either ions or neutrals. As a result of the interaction of these primary particles with the sample, species are ejected that have become ionized. These ejected species, known as secondary ions, are the analytical signal in SIMS. The use of a low dose of incident particles (typically less than 5 x 1012 atoms/cm2) in static SIMS, is critical to maintain the chemical integrity of the sample surface during analysis. A mass spectrometer sorts the secondary ions with respect to their specific charge-to-mass ratio, thereby providing a mass spectrum composed of fragment ions of the various functional groups or compounds on the sample surface. The interpretation of these characteristic fragmentation patterns results in a chemical analysis of the outer few monolayers. The ability to obtain surface chemical information is the key feature distinguishing static SIMS from dynamic SIMS, which profiles rapidly into the sample, destroying the chemical integrity of the sample. In the dynamic SIMS, a solid specimen placed in a vacuum, is bombarded with a narrow beam of ions, called primary ions, which are sufficiently energetic to cause ejection (sputtering) of atoms and small clusters of atoms from the bombarded region. Some of the atoms and atomic clusters are ejected as ions (called secondary ions). The secondary ions are subsequently accelerated into a mass spectrometer, where they are separated according to their mass-to-charge ratio and counted. The relative quantities of the measured secondary ions are converted to concentrations, by comparison with standards, to reveal the composition and trace impurity content of the specimen as a function of sputtering time (depth). The SIMS depth profile can provide elemental concentrations in the sample as a function the depth, has a great potential in characterizing the concentration profiles of self-organized or consolidated nanostructures.

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Particularly, SIMS is a powerful analytical technique for determining the elements present in materials, and especially on surfaces, with trace level sensitivity on the order of parts-perbillion (ppb) and sub-nanometer depth resolution and a high spatial resolution (lateral resolution of 5 nm) [242]. A classical SIMS device consists of (a) primary ion gun generating the primary ion beam, (b) a primary ion column, accelerating and focusing the beam onto the sample (and in some devices an opportunity to separate the primary ion species by wien filter or to pulse the beam), (c) high vacuum sample chamber holding the sample and the secondary ion extraction lens, (d) mass analyser separating the ions according to their mass to charge ratio, and (e) ion detection unit [241-245]. SIMS requires a high vacuum with pressures below 10-4 Pa (roughly 10-6 mbar or torr). This is needed to ensure that secondary ions do not collide with background gases on their way to the detector (mean free path), and it also prevents surface contamination by adsorption of background gas particles during measurement.

Figure 18. Raman spectra of hydrothermal BaTiO3 powders: (a) as-prepared, and after heat treatment at (b) 400 oC, (c) 600oC, (d) 800 oC, (e) 1000 oC, and (f) 1200oC for one hour. Reproduced with permission from [238], Wei, X. Z.; Li, Y. L. The influence of lattice defects on the crystal structure of hydrothermal BaTiO3 powders. J Ceram Proc Res. 2005, 6, 250-254. Copyright © 2005, Journal of Ceramic Processing Research.

As an example, SIMS was used to investigate the cation diffusion in perovskite oxides based on lanthanum gallates (LaGaO3) doped with strontium on the A site and magnesium on the B site, which exhibit high oxygen-ion conductivity and represent a promising alternative to YSZ (yttria-doped zirconia) as the electrolyte in solid oxide fuel cells [243]. Although cation diffusion in simple perovskites is known to be very slow, there are several important processes that are determined by the slowest moving species, such as sintering or creep. If the cations exhibit different diffusivities, kinetic demixing of the electrolyte [246] can be an additional origin of long term degradation. It is therefore important to obtain data for cation diffusion in La1-xSrxGa1-yMgyO3-(x+y)/2 (LSGM). By means of SIMS cation impurity diffusion of Y, Fe and Cr and cation tracer-diffusion of La, Sr and Mg in La0.1Sr0.1Ga0.1Mg0.1O2.9 were

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investigated [247]. By combining different modes of SIMS analysis- depth profiling, line scanning and imaging - it was possible to measure diffusion coefficients from about 10-18 cm2 s-1 to 10-6 cm2 s–1and to determine surface exchange coefficients as well. In addition, highresolution SIMS makes it possible to investigate diffusion through space charge layers at surfaces and to distinguish between bulk and grain boundary diffusion in polycrystalline materials, and determine the location of interfaces in solids and their chemical compositional depth profiling [242-244].

3.2.9. Optical absorption-emission spectroscopy Spectacular changes in optical characteristics of nanocrystalline oxides when compared to those of bulk counterparts have triggered tremendous interest among scientists to understand the basic mechanisms responsible for the fascinating optical absorption-emission, which also helps to examine their potential use in variety of optical applications [248-250]. Optical absorption and emission arises as a result of electronic transitions in solids upon exposure to excitation energies in the range of ~ 102 to 103 kJ/mol that cover the near infrared through visible to ultraviolet. There are various types of optical transition in solids. One type of transition is the promotion of an electron from a localized orbital to a higher energy localized orbital of the same atom (d-d, f-f transitions) or from a localized orbital in one atom to a higher energy localized orbital on an adjacent atom (charge-transfer spectra). Another type of transition can be the promotion of electrons from a localized orbital in one atom to the delocalized energy band (conduction band) of the solid as seen in the case of photoconductive materials. The transition energies associated with these processes differ, thereby requiring different excitation frequencies for obtaining their absorption and emission spectra [251]. Understanding the quantum confinement effect on optical absorption and emission characteristics has been the major objective of optical characterization oxides nanopowders. For example, the visible transitions of Nd3+ ions were found in the neodymium ion-doped perovskite hosts powders (Ca, Ba, Sr)TiO3, which were prepared by wet chemical method. The excitation at the band edge of the host at 335 nm generated an intense red emission at 613 nm with a quantum efficiency of 10.8% for Nd0.005(Ca0.97Ba0.01Sr0.015)TiO3 relative to the commercially used red phosphor [252]. With an increase in the dopant concentration of neodymium ions, the excitation and emission intensities both increased up to 0.5 mol% neodymium substitution. There was no shift in the excitation and emission spectra of the samples. The symmetry around the emitting center was more distorted with greater substitution by neodymium ions, resulting in violation of the parity selection rule and thereby producing the red emission. The preliminary photoluminescence properties of polycrystalline powder PbTiO3 were reported by Folkers and Blasse [253]. A broad-band emission in the visible spectral region (also called ―green‖ luminescence) are reported to be universal for ABO3 perovskite-type oxides [254-258]. However, the nature of this wide-band visible luminescence is not well understood, although some mechanisms, such as donor–acceptor recombination [259], transitions in MeO6 complexes [260,261], recombination of electron and hole polarons [262], and charge transfer vibronic exciton [263-266] have been proposed. Recently, Eglitis et al. [267] have performed the quantum chemical calculations and theoretical simulation of the green emission for a PbTiO3 perovskite-type oxides by using the intermediate neglect of differential overlap method combined with the large unit cell periodic defect model. Their results showed that the universal ―green‖ luminescence in the PbTiO3 crystals can be ascribed to the radiative recombination of the self-trapped electrons and holes

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forming the charge transfer vibronic exciton, rather than due to the electron transitions in a MeO6 complex or donor–acceptor recombination, as intuitively suggested earlier. They also demonstrate that well-parameterized semi-empirical quantum chemical methods could be successfully used for the study of optical properties of modern advanced materials such as oxide perovskites.

3.2.10. Mössbauer spectroscopy Mössbauer spectroscopy is a spectroscopic technique based on the Mössbauer effect. In its most common form of Mössbauer absorption spectroscopy, a solid sample is exposed to a beam of gamma radiation, and a detector measures the intensity of the beam that is transmitted through the sample, which will change depending on how many gamma rays are absorbed by the sample. The atoms in the source emitting the gamma rays are the same as the atoms in the sample absorbing them. As can be explained through the Mössbauer effect, a significant fraction of the gamma rays emitted by the atoms in the source do not lose any energy due to recoil and thus have almost the right energy to be absorbed by the target atoms. The gamma-ray energy is varied by accelerating the gamma-ray source through a range of velocities with a linear motor. The relative motion between the source and sample results in an energy shift due to the Doppler effect. In the resulting spectra, gamma-ray intensity is plotted as a function of the source velocity. At velocities corresponding to the resonant energy levels of the sample, some of the gamma-rays are absorbed, resulting in a drop in the measured intensity and a corresponding dip in the spectrum. The number, positions, and intensities of the dips (also called peaks) provide information about the chemical environment of the absorbing nuclei and can be used to characterize the sample. In order to occur Mössbauer absorption of gamma-rays, it is required that the gamma-ray must have the appropriate energy for the nuclear transitions of the atoms being probed, which is almost always achieved by having the same atoms of the same isotope in both the source and the target. Also, the gamma-ray energy should be relatively low, otherwise the system will have a low recoil-free fraction (see Mössbauer effect) resulting in a poor signal-to-noise ratio. Only a handful of elemental isotopes exist for which these criteria are met, so Mössbauer spectroscopy can only be applied to a relatively small group of atoms including: 57 Fe, 129I, 119Sn, and 121Sb. Of these, 57Fe is by far the most common element studied using the technique. The Mössbauer spectrum has been extensively used for the characterization of Fe-containing oxides [268]. In particular, the technique provides crucial information about the local order and associated magnetic properties in nanocrystalline ferrites [269-274]. For example, the Mössbauer spectroscopic characterization of Eu-doped or Mn-doped BiFeO3 powders revealed the addition of Eu or Mn in BiFeO3 induced significant modifications in the Mössbauer hyperfine parameters and the magnetic properties of the powders, whereas no significant microstructural or structural changes were observed [270,271].

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4. PROPERTIES OF PEROVSKITE NANOPOWDERS 4.1. Ferroelectric and Dielectric Properties In recent years studies on ferroelectric and dielectric properties of perovskite nanopwders have been become a major field of research due to their potential applications in memory devices. It has been well documented that a particle size effect on ferroelectricity exists in many perovskite nanopowders of displacive system such as BaTiO3, PbTiO3, and PbZrO3 [275-283]. That means below the critical particle size, the ferroelectricity disappears. An important motivation for study of the size effects on ferroelectric is to determine the ultimate level to which a device based on such systems can be miniaturized. As a typical example, the size effects on the phase transition of perovskite PbTiO3 nanopowders have been investigated [277-282]. The results show that nanocrystalline PbTiO3 particles with size of 20 - 80 nm exhibit a reduction in tetragonal distortion, ultimately transforming to a cubic phase for smaller particles ( critical size ~ 7 nm at room temperature) [279]. The dielectric constant around the Curie temperature ( Tc) can be expressed by a semi-empirical relation [284,285] 1

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



1

 max



1 (T  Tc ) A

(3 )

where  is the dielectric constant, max is the maximum dielectric constant value at the transition temperature Tc, A is a constant, and  is the order of phase transition within the range of 1<  200 V) porous morphology of TiO2 film has been observed in non-fluoride containing solutions [60].However, the pores were not uniform and internal cracks of the oxide film could be observed because of dielectric break down. Anodization at higher potentials, also referred as anodic spark deposition, leads to porous oxide as a result of evolution of oxygen and steam due to plasmachemical and thermal reactions. Very high current density of this process resulted in visible sparks and fusion of oxides. The pores were of 1–5 μm [60]. Whereas, fluoride solutions required much lower voltage (> KV) and can break the bonding of the total magnetic moment to the particles In such circumstances the magnetization vector of the particles can freely move and point in any direction arbitrarily. This phenomenon is called superparamagnism and the corresponding diameter of the particle is called the superparamagnetic size (Dsp). The system and the magnetic particles are referred to as a superparamagnet and superparamagnetic particles, respectively [8]. Moreover, in view of the magnetization curve, the coercivity of single-domain particles is higher than that of multipledomain particles [5]. With decreasing size of single-domain particles, the coercivity of singledomain particle decreases as well. In short, the maximum value of coercivity occurs at particle diameter equal to Ds. When the size of a single-domain particle is further reduced to a certain extent, for which the coercivity becomes zero, this size is referred to as Dsp. In view of Mösbauer spectra, the doublet of the Mösbauer line for superparamagnetic particles is observed instead of the sextet, and their magnetization curve passes through the origin when the magnetic field is zero. For both single-domain and superparamagnetic particles, their main advantages are their large external surface due to the nano-scale diameter of the particles and possible shorter lengths of internal diffusion paths. However, there is an important difference in particle coercivity. For single-domain particles with a remanent magnetization (or coercivity) used in the liquid-solid contacting system, the magnetic particles tend to agglomerate due to the attractive magnetic force if there is no externally intensive mixing, that is, there exists the interparticle interaction. For superparamagnetic particles, the attractive force of anisotropic interactions between particles can be ignored in this case compared with the former case. Under low mixing, superparamagnetic particles can still disperse well in the system and completely contact the aqueous solution and the solutes in the solution. Therefore, although the saturation magnetism of a superparamagnetic particle is lower than that of a singledomain particle for the same species, combined use with a high gradient magnetic separator gives superparamagnetic particle high potential for application as the magnetic core of synthesized micro- or nano-adsorbents which can be easily recovered, regenerated, recycled and reused. The preparation of nanoscale magnetite can be seen in the following sections.

2.3. Synthesis of Magnetic Core, Carrier and Adsorbent/Catalyst 2.3.1. Magnetic core The traditional way to produce micro- and nano-scale magnetite is long-term grinding of macroscale magnetite in a suitable medium until the expected scale of the particles is reached [9]. In addition, many chemical methods were also used in the production of superparamagnetic particles, such as the sol-gel method [10], hydrothermal reaction [11], flow injection method [12], chemical coprecipitation method [13-21], etc. The chemical coprecipitation method is more common than other methods that have previously been proposed due to its usefulness, ease of manufacturing and simple purification process. However, the choice of the method eventually depends on the purpose of the final functionalized materials. Here, we only discuss the coprecipitation method. The coprecipitation method is normally carried out by means of using ferrous (Fe(II)) and ferric

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(Fe(III)) salts with the addition of a basic solution. The characteristics of the resulting precipitate are governed by the molar ratio of Fe(II) to Fe(III), the pH value of the synthesized solution, concentrations of the reagents, sorts of basic solutions, heating time and reaction temperatures [22-33]. Furthermore, additional agitation is necessary to control the homogeneity of particle size. According to the stoichiometric formula, the theoretical ratio of Fe(II) to Fe(III) to produce one mole of magnetite is 1/2, as shown in Eq. [1]. Fe2+ + 2 Fe3+ + 8 OH- → Fe3O4 + 4 H2O

(1)

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However, a partial amount of Fe (II) may be transformed to Fe(III) during precipitation as a result of oxidation by air. This phenomenon is detrimental to the magnetic quality of the products (i.e., less magnetite). For the reason given above, the problem is solved in two different ways. One way is that the commonly used dosage of Fe(II) is in excess of the stoichiometric amount in the system; the other way is using nitrogen stripping to avoid dissolved oxygen and therefore oxidation of the Fe(II) in the solution. With respect to the effect of the basic solution on magnetic quality of the product, Gribanov et al. [23] has provided extensive discussions of applications of various basic solutions as hydrolyzing agents. The optimal precipitator is ammonia (ammonium hydroxide, NH4OH) due to its easy control of pH value and high buffer capacity at a pH range of 8-10, contributing to the production of magnetite with high saturation magnetization. Regarding the reaction time of chemical coprecipitation, the reaction time does not significantly affect the properties of the resulting magnetite [6,19], which mean that the size and uniformity of particles were determined by the very short reaction time when nucleation occurred. The investigation of the effects of ionic strength [34] and reaction temperature [35] indicated that higher ionic strength and reaction temperature lead to smaller particle size. Briefly, particle size below 10 nm with spherical shape can be obtained from this method [19-21].

2.3.2. Magnetic carrier For the sake of stabilizing the naked magnetic particles in solution, the additional forces of either electrostatic or steric repulsion are commonly adopted for this purpose. In view of the application of superparamagnetic particles, much attention has been paid to the further functionalization of such particles in order to match the specific target, such as removal of specific contaminants, magnetic resonance imaging, bioseparation and drug delivery [1921,36]. Therefore, numerous functional groups and chemical compounds were used to coat the naked magnetic particles in order to satisfy different demands, which have been well described in the literature [36]. Fundamentally, the naked magnetic particle was bound or covered by a chemical molecule or film to form the core/shell structure via monomeric and polymeric materials. In respect of the manifold applications of superparamagnetic particles in adsorption, inorganic silica was first considered to be coated on the naked particles for the following reasons. Silicon dioxide (silica, SiO2) is widely used as the carrier of metal catalysts, mobile composition materials, and nano-composites [37]. In addition, silica is rather stable in the environment and does not react with other reductants and oxidants, except fluorides and hydrofluoric acid, characteristics which are of paramount importance for the selection of the protective film for the magnetic core. This property makes it possible to regenerate the adsorbents under either acidic or basic conditions. Silica, therefore, possesses high potential to be applied as a passive film. Furthermore, the silica film also provides

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benefits, such as inhibition of sintering of the magnetically finished products, dispersion of the magnetic particles, enhancement of wear-resistance and stability of colloids in solution [38]. Such superparamagnetic particles with core/shell (SiO2/Fe3O4) structure can also easily be further functionalized by means of silylation based on the silanol group of surfaces and surface coatings. In sum, coating with a passive film of silica on the naked magnetic particle possesses the advantages of protecting the magnetic core, stabilizating the magnetic particles, improving the physicochemical stability of magnetic particles and supplying a basis for further functionalization. The superparamagnetic particle (SiO2/Fe3O4) core/shell structure is called a magnetic carrier in this study, as shown in Figure 1.

O O

Si

O

O

Si

Si O

O O

O

Si O

OH

O

Si O

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Fe3O4

O

Silica/Fe3O4

O

Si

O

Al

O

O Si

Si

Q3 structure

O

Si(3Si,1Al)

O

Aluminosilicate/Silica/Fe3O4

Figure 1. Diagrams showing the concept of nanocomposites comprised of magnetite, silica and aluminosilicate synthesized in this study. The nature and structures of the synthesized composites in this study were based on analyses of solid-NMR spectra.

The silica film can be formed by acidifying (acidulation) [39-42], sol-gel [39-45], emulsion [46], and thermal methods [47]. The emulsion used for the synthesis of core/shell magnetic carriers is composed of water, solvent (e.g., saturated benzene [42] and oil [46]), surfactants and TEOS/TMOS. In the thermal method, the mixture of TEOS, iron salt and methanol was directly put into the furnace for the pyrolysis process under various temperatures, which play a significant role in determining the size of the final particles. Here, we come to the point at which it is necessary to deal in more detail with the acidifying and sol-gel methods due to the wide use of these two methods. For the acidifying method, the aqueous sodium silicate (Na2SiO3) solution, with pH value greater than 12, is employed as a precursor. The addition of an acidic solution (acidification agent) gives rise to the aggregation of silica particles, which results in the formation of silica gel when the pH value is below 7. The reactions of the silica sol-gel process have been widely studied and discussed since Stöber et al. [15] described the systematic study of syntheses of monodisperse silica in 1968. The sol-gel process normally consists of four main stages, which

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are sol formation, gelation or aging, drying and calcination or sintering. The basic chemistry of a sol-gel process normally is based on reactions of hydrolysis and condensation of the alkoxide precursors (e.g., tetramethylorthosilicates (Si(OCH3)4, TMOS) and tetraethylorthosilicates ((Si(OC2H5)4, TEOS)), whereas condensation is the only reaction to be considered for the acidulation method. The starting solution of the sol-gel method is the mixture of the alkoxide precursor, solvent (e.g., water), co-solvents (e.g., methanol and ethanol) and catalysts (e.g., acid and base). The composition of the mixture determines the relative rates of the overlapping hydrolysis and condensation processes, which are regarded as the sole factor in determining the structure of the final products [39]. The influence of the acidic or basic catalysts on the silica sol-gel process determines the relative rates of the hydrolysis and condensation processes. For the acid-catalyzed gel, condensation is the ratelimiting reaction which results in the generation of many small particles. For the basecatalyzed gel, the rate-limiting reactions are hydrolysis and particle nucleation, which leads to the formation of larger particles with low specific surface area and high porosity. The higher concentration of the ammonia makes hydrolysis and condensation reactions faster, resulting in larger particles [48]. In addition, the higher the initial concentration of alkoxide precursor is, the larger are the particles formed [49]. Higher reaction temperature promotes the rate of nucleation and inhibits the growth of nuclei so that it, by contrast, produces smaller particles [50]. Although it is not possible to accurately predict the microscopic structure of the materials due to the complexity of sol-gel chemistry, the mixture of the sol-gel process still can be specially designated in advance in order to obtain the particular magnetic carrier desired. It is possible to combine the acidifying method and sol-gel method in sequence to control the particle size and distribution, according to the report of Philipse et al. [18]. Furthermore, ultrasonic treatment has been shown to improve the dispersion of magnetic particles in aqueous solution, so as to adjust the amount of magnetic core in the magnetic carrier [17].

2.3.3. Magnetic adsorbent/catalyst The magnetic adsorbent/catalyst can be further synthesized from a magnetic carrier for various purposes. The methods to carry out the synthesis of a magnetic adsorbent/catalyst are similar to the ways in which solid catalysts and metal oxides are prepared, such as chemical vapor deposition, chemical reduction and coprecipitation, sonochemical reactions, gel-sol, microwave and high frequency plasma, low energy cluster beam deposition, ball milling, hydrothermal, microemulsion and self-assembly [51-55]. In considering the convenience of operation, requirement of equipment, development of synthesizing technique, and scale-up of the experimental procedure and performance of the products, homogeneous precipitation (or deposition-precipitation) and sol-gel methods are chosen and shortly described in this study. The following examples will explain how to prepare the superparamagnetic adsorbent/catalyst through these two methods. 2.3.3.1. Homogeneous precipitation method Precipitation is the most frequently used method for catalyst preparation due to the advantages of high purity and flexibility [51]. The resulting products are affected by the raw materials, pH value of the solution, precipitating agent, temperature, aging, etc. With respect to quality of the resulting products (precipitates), one important point should be made about the homogeneity. The first and final formed precipitates commonly differ from each other,

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mainly caused by the different concentrations of the ions and aging time in the solution during reaction. Therefore, homogeneous precipitation is designed to overcome the drawbacks of the above-mentioned phenomenon by means of the addition of the precipitating agent (e.g., urea, CO(NH2)2), which gradually and slowly raises the pH level of the solution. The chemical reactions of urea decomposition can be divided into two stages in the solution, as shown in Eqs. 2-3. In addition, the overall reaction of urea decomposition is presented in Eq. 4, which produces ammonium, hydroxyl ion and carbon dioxide. The following serves as an example: Al-MCA. The synthesizing procedures of superparamagnetic Al-MCA can be broken down into three steps. The first step begins with loading the magnetic carrier in the homogeneous solution of aluminum nitrate and urea, the amount of which used in the reaction is recommended to be 1.5-2.5 times the amount stoichiometrically needed [53]. Next, the solution with suspended magnetic carrier is heated to enhance the further decomposition of cyanate for a suitable period until the precipitate of aluminum hydroxycarbonate forms and deposits (or precipitates) on the magnetic carrier. After easy separation by means of a permanent magnet for a while, the magnetic solids are finally calcined to obtain the resulting magnetic adsorbent. The overall reaction is expressed in Eq. 5, in which the crystal type of the final product depends on the calcination conditions. Chang et al. [19] and Luther et al. [56] have successfully coated bayerite on a magnetic carrier and alumina on silicon nitride (Si3N4) through the following methods, respectively. CO (NH2)2 → NH4+ + CNO-

(1)

CNO- +3 H2O →NH4+ + OH- + CO2

(2)

CO(NH2)2 + 3 H2O → 2 NH4+ + 2 OH- + CO2

(3)

calcination

2 Al(NO3)39H2O + 5 CO(NH2)2  

(4)

Al2O3 + 5 CO2 +28 H2O +8 N2

(5)

With respect to the interaction between the precursors and magnetic carrier, it can be evaluated by means of measuring either the pH values versus time in the system or the amount of alkali used in the system both with and without magnetic carriers. Extensive information can be seen in the study of Geus and Van Dillen [57].

2.3.3.2. Sol-gel method Compared with the precipitation method, the sol-gel method possesses the advantage of compositional homogeneity at a microscopic or molecular level. There are two common routes to carry out the synthesis of metal oxides through the sol-gel process [52]. The choice of route mainly depends on the sort of precursors (one is inorganic salts and the other is metal alkoxides). The synthesizing procedures of superparamagnetic adsorbents can be divided into three subsequent steps. Firstly, the magnetic carrier loaded into the solution is well suspended in the admixture under mixing. After suitable aging time, the drying of the solvent is carried

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out under mild condition, such as irradiation with an infrared lamp. Finally, the gel is calcined to obtain the resulting magnetic adsorbent. The sol-gel method for silica has been described earlier, and it was mentioned that relative rates of the overlapping hydrolysis and condensation processes can be viewed as the single factor affecting the properties of the products. For the metal alkoxides, the selection of the types of precursors can effectively control the reaction rates [58]. In addition, the amount of solvent also plays an important role in the characteristics of the gel. Concerning the effect of the solvent, it can be judged from the hydrolysis ratio (h, defined as the ratio of moles of water (H2O) to the alkoxide precursor (M(OR)Z) used in the system). According to the study of Livage et al. [53], there are three categories of h: the first is h < 1; the second is 1  h  Z; the third is h > Z. For h < 1, an infinite network seldom develops, and gelation or precipitation cannot occur due to the low functionality toward condensation and no local excess of water, respectively. When 1  h  Z, polymers can form in this category. Concerning the third category (h > Z), cross-linked polymers, particulate gels, or precipitates can form as long as an excess of water is added to the alkoxide. Therefore, the particle size and form can be controlled by the value of h through the adjustment of the ratio of water and alkoxide content. Regarding the aging and drying processes, the purposes of these two stages are to allow a gel to undergo changes in properties and to remove the solvent from a gel, respectively. The important parameters affecting the aging process include time, temperature and composition of solvent. Several methods can be used to execute the drying process, including evaporation, supercritical drying, freeze drying etc. The final stage of sol-gel (i.e., calcination or sintering) can make the materials crystallize into different structural types by means of altering parameters such as final temperature, heating rate and gaseous environment. We have successfully synthesized the superparamagnetic adsorbent/catalyst through sol-gel methods, such as aluminum-type [19], zirconium-type [20] and aluminosilicate-type [21].

3. APPLICATIONS Here, there are examples of superparamagnetic adsorbents to illustrate the synthesis, characterization and applications of such superparamagnetic particles, which we have successfully synthesized to apply in the removal of environmental contaminants. Activated alumina is the widely used adsorbent in drinking and wastewater treatments for the elimination of fluoride compounds [59-60], due to the formation of inner-sphere complexation [61]. Furthermore, zirconia also has high affinity to fluoride so that it is also often used in chromatographic analysis and also the removal of fluoride [62-64]. The removal of fluoride anions takes place through adsorption at a positively charged surface group at low pH value or ligand exchange with hydroxyl ions at high pH value [65-66]. With respect to the cationic contaminants such as Cu(II), Zn(II), and NH4+, the surface of negative charge is necessary to produce electrostatic attraction for removal of positively charged compounds. Previous studies [19-21] have successfully synthesized three effective types of superparamagnetic adsorbents of bayerite/SiO2/Fe3O4 (Al-MCA), zirconia/SiO2/Fe3O4 (ZrMCSG), and aluminosilicate/SiO2/Fe3O4 (Alsi-MCSG) via three sequential steps: chemical precipitation of Fe3O4, coating of SiO2 on Fe3O4 using acidifying (denoted as MCA) and sol-

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gel (denoted as MCSG) methods, and further coating of bayerite, zirconia and aluminosilicate on MCA or MCSG via the sol–gel method. In the following content, we discuss the procedures in detail from the previous studies. The commercial activated alumina (anhydrous -Al2O3, CA) and nano-sized zirconia (NCZ) discussed below were from Merck and Sigma, respectively. Prior to experiments, CA was first washed with 450 cm3 of 0.1 N NaOH solution and then washed with 3.610-4 N HNO3 until the pH value of the effluent was kept at 4. After overnight drying at 373±5K, the resulting adsorbent was activated alumina (CA).

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3.1. Synthesis of Magnetic Core, Carrier and Adsorbents The magnetite cores were prepared using ferric and ferrous chloride salts (FeCl3 and FeCl2) at a stoichiometric molar ratio of 2:1 (as seen in Eq. 1) by adding ammonia as precipitator in an inert atmosphere. The experimental temperature was kept at 353 K. After strong mixing for 3 minutes, the resulting magnetic cores (magnetite) were immediately separated from the solution by means of a permanent magnet and then washed several times with distilled water. The particle size from this simple synthesizing procedure is around 7 nm and the reaction time has no significant influence on the particle physicochemical properties judged by the TEM and XRD patterns. Furthermore, the magnetic carrier (SiO2/Fe3O4) was obtained via the acidifying method (denoted as MCA) and sol-gel method (denoted as MCSG). In the acidifying method, the obtained magnetic core was put in a sodium silicate solution, which was prepared with a volume ratio of sodium silicate to ultra-pure water equal to 1:2. Then hydrochloric acid (6 wt.%) was added dropwise into the solution and lowered the pH value of the solution down to 6 in 3 h. After a simple magnetic separation and washing several times, the obtained magnetic carrier MCA was dried at 373±5 K over night. On the other hand, the mixture for sol-gel synthesis, was composed of ammonium hydroxide, isopropanol, ultra-pure water and tetraethylorthosilicate ((Si(OC2H5)4, TEOS)). The ratio of isopropanol : water : NH4OH (25 vol. %) : TEOS was 30 : 2 : 7.8 : 1 with a total volume of 2 dm3. After strong mixing for 5 hours at a temperature of 313 K, the magnetic silica-coated particle was dried with an infrared-ray lamp overnight. It was then calcinated at 673 K for eight hours to obtain a magnetic silica-coated carrier of MCSG.. Finally, magnetic adsorbents were prepared from a magnetic carrier by means of the solgel method. For the synthesis of an aluminum-type adsorbent of AlMCA, first, 8 g MCA, 40 g aluminum isopropoxide, and 800 cm3 ultra-pure water were mixed in a 2-L reactor under ultrasonic shaking for 20 min. Subsequently, 213 cm3 of isopropanol–water mixture (33 vol.% of isopropanol in ultra-pure water) was introduced dropwise in the reactor at the rate of 3 cm3 min-1. Finally, the obtained slurry was dried overnight under an infrared-ray lamp and then calcined from room temperature to 823 K at a heating rate of 10 K min-1 under a constant N2 stream of 100 cm3 min-1. After the temperature reached 823 K, the temperature was maintained at the final temperature for 2 h. For magnetic zirconia/SiO2/Fe3O4 (ZrMCSG), the calcined MCSG was then further contacted with the other mixture, which comprised nitric acid, isopropanol, ultra-pure water and zirconium (IV) tert-butoxide (Zr(OC(CH3)3)4). The volumetric ratio of isopropanol:water:nitric acid:Zr(OC(CH3)3)4 was 88:3:1:21 with a total volume of 130 cm3.

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After vigorous mixing for 5 h at a temperature of 313 K, the obtained particles were dried with an infrared-ray lamp overnight (denoted as preZrMCSG). Then, the magnetic particle of preZrMCSG was calcinated at 673K for 4 h to obtain a superparamagnetic zirconiumtype adsorbent of ZrMCSG.. For aluminosilicate/SiO2/Fe3O4 (AlSiMCSG), the mixture was composed of aluminum isopropoxide, TEOS, isopropanol, ultra-pure water, and MCSG. The ratio of aluminum isopropoxide:TEOS:water:isopropanol was 0.6g:21 cm3:140cm3:55cm3. After a reaction time of 20 h, the particles were dried at 373K in the oven overnight. All the magnetic adsorbents (i.e., AlMCA, ZrMCSG and AlSiMCSG) were washed with ultra-pure water until the specific electric resistance of the effluent was equal to the original value (i.e., 18.3M-cm), adopting a permanent magnet for liquid-solid separation. After overnight drying at 373±5K, the resulting particles were the magnetic adsorbent and ready for use.

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3.2. Characterization of Adsorbents 3.2.1. Average particle size and the composition of the adsorbent The particle size of Fe3O4 formed by the coprecipitation method is around 7 nm, as illustrated in Figure 2. The particle size and shape are rather homogeneous and uniform, indicating that the coprecipitation method is an effective method to produce nano-sized Fe3O4. The characteristics of the magnetic particles determined by nitrogen gas adsorption at 77 K are illustrated in Table 1. The BET specific surface areas of Al-MCA, Zr-MCSG and AlSi-MCSG are 42, 79.6 and 158.5 m2 g-1, respectively. From the analysis of the pore volume, the meso- and macro-pores consitute the majority of the total pore volume. Combined with the TEM analysis, the increment of specific area in magnetic particles seems to represent the void between the aggregates rather than the intraparticle pore volume . Take the previous study [19] for example. The particle sizes of Fe3O4, MCA and Al-MCA are around 7, 31.3 and 33 nm, respectively, so that it is reasonable to assume that Al-MCA is a non-porous particle. Further, adopting more assumptions that the particle is spherical and the true density equals the apparent density, the average particle size can be calculated using the formula dp=6/(ABP). However, the values obtained from TEM micrographs supported by a statistical method (log-normal distribution) and specific surface area may not be coincident due to the aggregation of particles during the specific surface area measurement procedure. Regarding the effect of the concentration of sodium silicate on the magnetic carrier, three volume ratios of sodium silicate to ultra-pure water (i.e., 0.03, 0.2 and 0.5) were investigated. The results show that the lower the ratio is, the more serious the phenomenon of agglomeration is. In addition, when the acidifying rate is raised to 6 cm3 min-1, the enclosure of the magnetic core with silica is still effective. Also, the average particle size of the magnetic carrier at an acidifying rate of 6 cm3 min-1 is larger than that at an acidifying rate of 3 cm3 min-1 judged from TEM results. The size distribution of nano-sized particles has been determined mostly by TEM up to now. However, if the nanoparticles are packed rather tightly, it is difficult to distinguish the entire surface morphology and particle size of the individual nanoparticles with TEM. Therefore, for the individual cluster of Zr-MCSG and Alsi-MCSG, an SEM picture is used to simply show the approximate particle size of the strong aggregate. The spectra from

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SEM/EDX analyses can show the approximate particle size of a single grain or aggregate and the composition of superparamagnetic adsorbents, as shown in Figure 3. The approximate particle size of a cluster is under 100 nm judged by the SEM picture. In addition, the elemental ratio by weights or moles can tell the abundance of Al on the surface of AlsiMCSG when compared with that of other superparamagnetic aluminosilicates synthesized by other methods and conditions.

Figure 2. Transmission electron microscopy (TEM) of magnetite.

Figure 3. SEM/EDX pictures of superparamagnetic Alsi-MCSG. (From Chang et al., Colloid Surf. A: Physicochem. Eng. Asp. 2009, 336 (1-3), 159-166. (doi:10.1016/j.colsufa.2008.11.042).

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Table 1. Characteristics of N2 adsorption at 77K of various superparamagnetic adsorbents Property BET specific surface area (m2 g-1) Pore volume (cm3 g-1) Total, Vt Micro-, Vi Meso-, Ve Macro-, Va

Al-MCA 42.34 0.1013 0.0017 0.0741 0.0272

Zr-MCSG 79.60 0.1651 0.0015 0.1231 0.0405

Alsi-MCSG 158.5 0.2147 0.0307 0.0508 0.1332

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(Data integrated from: Chang et al., Colloid Surf. A: Physicochem. Eng. Asp. 2006, 280 (1-3), 194-202 (doi:10.1016/j.colsufa.2006.02.011); Chang et al., Colloid Surf. A: Physicochem. Eng. Asp. 2008, 327 (1-3), 64-70 (doi:10.1016/j.colsufa.2008.06.006); Chang et al., Colloid Surf. A: Physicochem. Eng. Asp. 2009, 336 (1-3), 159-166. (doi:10.1016/j.colsufa.2008.11.042)).

3.2.2. The crystal structure and elemental environment of the adsorbents XRD (X-ray Diffraction) is commonly used to reveal the crystal structure and also the chemical composition of materials. The crystal structures of magnetic particles was examined by a powder X-ray diffraction (XRD) analyzer (MXP diffraction collector, MAC Science, Japan) with Cu K(wavelength λ = 1.5418 Å) and the Joint Committee on Powder Diffraction Standards (JCPDS) database. The positions of peaks and the relative intensities of XRD patterns can be used to verify the crystal structures of the powder, as illustrated in Figure 4. According to the characteristic lines in the JCPDS database, the synthesized particle was verified to be magnetite. Another example was: the crystalline zirconia on the Zr-MCSG was a mixture of tetragonal and monoclinic types. Sometimes, it is difficult to judge the crystal structure of magnetic adsorbents due to the overlap of many peaks produced by the core/shell particles. Since the functional group is the on the top layer of the core/shell structure, it is reasonable to individually synthesize the top layer compound for XRD characterization purposes. The crystal of aluminum hydroxide of Al-MCA is a typical example for this purpose [19]. Regarding the Alsi-MCSG, the angular ranges between 15o and 30o represent the characteristic lines of not only silica but also the aluminosilicate so that it failed in discriminating them from each other due to the strong peak of amorphous silica. Therefore, additional analyses were necessary to determine the crystal structure of the aluminosilicate on Alsi-MCSG. The common techniques used are Fourier transform infrared spectroscopy (FTIR), electron spin resonance (EPR) and nuclear magnetic resonance (NMR). Take AlsiMCSG for example, as illustrated in Figure 5. The FeO, SiFeO, asymmetric and symmetric SiOSi, AlO4 and AlO6 can be identified through the FTIR spectra, while, in contrast, the bond of SiOAl was not judged due to the overlap with the peak of SiOSi. Therefore, the exact connection between Si and Al needs further investigation by means of other techniques. However, the analysis, which is performed under an additional magnetic field, normally encounters difficulty due to the strong peaks arising from the magnetic property of magnetic particles, such as EPR and NMR. Due to the good enclosure by the silica film, the elemental environment of the final functional group on the magnetic carrier was hypothetically not affected by the magnetic core. Therefore, investigating the elemental environment of the final group can be done through the synthesis of the functional layer on

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the silica surface without using the previous synthesis procedure of the magnetic core. The example of Alsi-silica referred to as Alsi-MCSG was taken for this purpose. The solid NMR was used to analyze the crystal structure of aluminosilicate, as illustrated in Figure 6. The combination of the results obtained from Si29 and Al27 MAS NMR spectra was able to clarify the incorporation of aluminum into the silica framework rather than the aggregate of alumina on the silica surface. Therefore, a combination of several analytical techniques can well define the crystal and elemental environment of the magnetic materials.

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Figure 4. Powder x-ray diffraction (XRD) patterns of various Fe3O4 nano-particles synthesized under different reaction times. (I), (II), and (III) correspond to reaction times of 3, 15, and 30 min, respectively. (From Chang et al., Colloid Surf. A: Physicochem. Eng. Asp. 2006, 280 (1-3), 194-202 (doi:10.1016/j.colsufa.2006.02.011))

3.2.3. Elemental analysis and magnetization curves For the sake of understanding the elemental compositions of the synthesized particles, it is necessary to enforce the digestion process prior to the analysis of inorganic compositions with inductively coupled plasma atomic emission spectrometry (ICP-AES). The microwave digestion equipment (ETHOS Touch Control, Milestone, Italy) was applied to digest the magnetic particles. According to the application note 039 provided by Milestone company, around 0.1 g of sample was mixed with the reagents (5 cm3 H2SO4 of 98 wt.%, 2 cm3 HNO3 of 65 wt.% and 2 cm3 HF of 40 wt.%). After the digestion process, the samples were examined by ICP-AES (JY-24, Jobin Yvon, Lonjumeau, Paris, France). The results are shown in Table 2. Table 2 indicates that the elemental composition of the magnetic carrier obtained from an acidifying rate of 6 cm3 min-1 is similar to that gained from an acidifying rate of 3 cm3 min-1. However, the weight percentage of bayerite in the magnetic adsorbent synthesized from the magnetic carrier with an acidifying rate of 6 cm3 min-1 is higher than that synthesized from the magnetic carrier with an acidifying rate of 3 cm3 min-1. This may be due to the fact that a faster acidifying rate possibly makes larger silica particle size and more porous volume. These phenomena may cause and promote more metal hydroxides deposition or combination with silanol groups on the surface of the silica on the magnetic carrier.

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Figure 5. FTIR spectra of Alsi-MCSG. (From Chang et al., Colloid Surf. A: Physicochem. Eng. Asp. 2009, 336 (1-3), 159-166. (doi:10.1016/j.colsufa.2008.11.042)

The magnetization curve is commonly used to understand the magnetic properties of materials, such as saturation magnetization, existence of hysteresis and magnetization remanence. The magnetization curves of the magnetic core, carrier and adsorbent obtained from the SQUID magnetometer are shown in Table 3 and Figure 7. The saturation magnetization (MS) of the magnetic core (Fe3O4) is 59 emu g-1. With respect to the magnetic carrier MCA, three volume ratios of sodium silicate to ultra-pure water (i.e., 0.03, 0.2 and 0.5) were investigated and the results show that materials of lower volume ratio possess higher saturation magnetization but worse enclosure of the magnetic core. In addition, the acidifying rate does not significantly affect the saturation magnetization of the resulting materials. Regarding the magnetization under various magnetic fields, zero coercive force and being free of hysteresis are observed for the magnetic core, magnetic carrier and magnetic adsorbent, indicating that all the materials synthesized are superparamagnetic. Based on the saturation magnetization of Fe3O4, the mass percentages of Fe3O4 content for magnetic carrier (C) and AlMCA are 14%, and 14%, respectively. Comparing the mass percentages obtained from the elemental analyses, the values of both analyses are coincident. Furthermore, although the silica is a diamagnetic substance, the susceptibility of silica is rather small (i.e., -10-6 emu mol-1 Oe-1 [5]), indicating that the existence of a silica film does not significantly counteract the magnetism gained from the magnetic core. For magnetic particles, the value of MS is a very important factor in the application because it is a key

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factor in the control of magnetic particles under the magnetic field [67-68]. Therefore, it is necessary to determine the magnetic properties under the designated conditions.

Figure 6. MAS NMR spectra of Alsi-MCSG. (From Chang et al., Colloid Surf. A: Physicochem. Eng. Asp. 2009, 336 (1-3), 159-166. (doi:10.1016/j.colsufa.2008.11.042)

Table 2. The chemical composition of various magnetic particlesa

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Sample SiO2/Fe3O4c SiO2/Fe3O4d MASGe MASGf

Mass ratio b, % Fe3O4 SiO2 1 6.2 1 6.15 1 5 1 5.31

Al(OH)3   1.2 0.55

Content, wt. % Fe3O4 SiO2 14 86 14 86 14 70 15 77

Al(OH)3   16 8

a

Only main components are considered in this table. Mass relative to that of Fe3O4. c Magnetic carrier is obtained with the acidifying rate of 6 cm3 min-1. d Magnetic carrier is obtained with the acidifying rate of 3 cm3 min-1. e Magnetic adsorbent is obtained from the magnetic carrier with acidifying rate of 6 cm3 min-1. f Magnetic adsorbent is obtained from the magnetic carrier with acidifying rate of 3 cm3 min-1. b

Table 3. Saturation magnetization of magnetic particles Magnetic particles Magnetic core (Fe3O4) Magnetic carrier (A)b Magnetic carrier (B)c Magnetic carrier (C)d Al-MCA-1e

Saturation magnetization, emu g-1 59.24 48.49 15.02 8.50 8.27

%a 100 82 25 14 14

a

Mass percentages of Fe3O4 content based on the saturation magnetization of Fe3O4. Volume ratio of sodium silicate to ultra-pure water and acidifying rate are 0.03 and 3 mL min-1. c Volume ratio of sodium silicate to ultra-pure water and acidifying rate are 0.2 and 3 mL min-1. d Volume ratio of sodium silicate to ultra-pure water and acidifying rate are 0.5 and 6 mL min-1. e Synthesized from the magnetic carrier (C). b

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60 50 40

Magetization, emu g-1

30 20 10 0 -10 -20 -30 -40 -50 -60 -10

-8

-6

-4

-2

0

2

4

6

8

10

Magnetic field, kOe Figure 7. Magnetization curves of various magnetic particles. , , ,  and : Fe3O4, magnetic carrier A, magnetic carrier B, magnetic carrier C and Al-MCA-1. The specifications of the magnetic particles are given in Table 3.

30 20

Zeta potential, mV

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40

10 0 -10 -20 -30 -40 2

3

4

5

6

7

8

9

10

11

12

pH, Figure 8. The zeta potential of Al-MCA-1 as a function of pH value under an ionic strength of 0.01N NaClO4.

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Figure 9. Titration of solution with suspended Al-MCA-1. Ο and : ion strength of 0.1 N and 0.01 N NaClO4. a. acidimetric-alkalimetric titration; b. charge calculated from the titration curve (charge balance); c. extrapolation to zero charge yielding intrinsic pKa1,int and pKa2, int.

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Figure 10. Titration of solution with suspended CA. Ο and : ion strength of 0.1 N and 0.01 N NaClO4. a. acidimetric-alkalimetric titration; b. charge calculated from the titration curve (charge balance); c. extrapolation to zero charge yielding intrinsic pKa1,int and pKa2, int. Nanopowders and Nanocoatings: Production, Properties and Applications : Production, Properties and Applications, Nova Science Publishers,

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3.2.4. Zeta potentials of magnetic particles The zeta potential is the potential measured at the surface of hydrodynamic shear or slipping plane of the particles. The magnitude of this value reveals the stability of the colloidal system. One of the most important factors affecting the zeta potential of particles is the pH value of the solution for the solid-liquid system. For the adsorption system, the surface charge of the adsorbent plays an important role in the adsorption capacity of ionic compounds. The point of zero charge of solid particles is the pH at which the surface charge of the particle is zero (i.e., pHPZC). The pHPZC of magnetic adsorbent in aqueous solutions can be determined using electrophoretic mobility combined with Henry‘s equation (Zeta Sizer 300, Malvern) or titration methods. Figure 8 shows that the zeta potential of Al-MCA-1 is around 5.9 obtained from electrophoresis under various pH values of solutions. Regarding the titration experiments for AlMCA-1 and CA, the preparation method followed Hao and Huang‘s research [65] and the calculation adopted Stumm‘s method [61], as illustrated in Figures 9-10. The results show that the pHPZC values of CA and Al-MCA-1 are 6.8 and 6.3, respectively. Comparing the pHPZC values obtained from electrophoresis and titration methods for Al-MCA-1, they are very close to each other. Although it may be inadequate to assume that the double layer of nanoparticles is flat [69], the calculation still can provide useful information. According to results of the electrophoresis and titration methods, the surface charges of CA and Al-MCA-1 are positive when the pH value of the solution is under 6. Furthermore, the surface charges of such nano-sized magnetic adsorbents are very sensitive to the pH value of solutions, especially under high ionic strength, indicating that the adsorption capacity of the specific compound may be significantly dependent on the pH value of the solutions. For adsorption by electrostatic attraction, the surface charge dominates the adsorption capacity for the removal of contaminants.

3.3. Adsorption Capacity The adsorption of a specific compound on the surface of nano-sized particles involves three sequential mechanisms: external mass transfer, internal mass transfer (i.e., surface and pore diffusions) and adsorption [2]. Due to the dual properties of nano-sized particles and magnetism, the application of such magnetic particles can be exactly estimated on the basis of adsorption capacity and equilibrium time. Compared with the traditional fixed bed column, the application of magnetic particles possesses advantages of high flexibility, easy control and unsophisticated calculation. Therefore, the application of nano-sized particles normally adopted the batch type reactor, which can easily control the reaction time to reach equilibrium for maximum adsorption capacity. The Langmuir and Freundlich isotherms are commonly used to describe the adsorption behavior of compounds in solutions due to the good agreement with experimental data so that these two equations are adopted here. The mathematical relationships of Langmuir and Freundlich isotherms are shown in Eqs 6 and 7, respectively. qe =

q L K L Ce , 1+K L Ce

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(6)

Preparation, Characterization and Potential Application of Magnetic... 1

q e  k F  Ce nF

117 (7)

In Eqs. 6 and 7, qe and Ce are the adsorbate concentrations in solid and liquid phases at equilibrium, respectively. The qL and KL are the Langmuir isotherm constants, in which the first one is on behalf of the monolayer adsorption capacity. The kF and nF are the Freundlich equilibrium constants, which represent the adsorption capacity and strength of adsorption, respectively. 20 18 16

qe, g kg-1

14 12 10 8 6 4 2 0 0

10

20

30

40

50

60

70

80

90

100

-3

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Ce, g m

Figure 11. Description of equilibrium data by the Freundlich (---) and Langmuir (—) isotherms under constant pH value of 6 and ionic strength of 0.1 N NaNO3. ○ and : Experimental data of Al-MVSG-1 and CA at 298 K.

The adsorption of fluoride from aqueous solution by Al-MCSG-1 and CA is shown in Table 4 and Figure 11. Not only the Freundlich but also the Langmuir isotherm equations can well describe the adsorption behaviors of fluoride on CA and Al-MCA-1 judged by the high correlation coefficients (all values of r2 > 0.9). Comparing isotherm constants predicted from Freundlich and Langmuir equations with CA and Al-MCA-1, qL of Al-MCA-1 (i.e., 24.33 g kg-1) is higher than that of CA (i.e., 10 g kg-1), indicating that Al-MCA-1 possesses higher monolayer adsorption capacity. In addition, the surface charge is very sensitive to the pH value so as to significantly affect the adsorption capacity, especially at the pH range near the point of zero charge (pHPZC). Take Alsi-MCSG for example. At a pH value of 6, a series of different magnetic aluminosilicates were used to adsorb Cu(II) ions from solutions. The Alsi-MCSG which possesses the highest zeta potential of -24 mV has the highest adsorption capacity for Cu(II). In the case of Zr-MCSG, the value of qL is comparable to that of NCZ. In addition, the adsorption capacity of sulfate at a pH value of 4 is almost twice as big as that at a pH value of 55.8. In solutions with a pH higher than pHPZC, the uptake of sulfate ended. Therefore, the adsorption capacity of sulfate is highly dependent on the pH value of the solution, which is the dominant influence in the adsorption behavior of sulfate. In brief, the results show that the novel magnetic adsorbents have high potential to be applied in the adsorption removal of

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contaminants from aqueous solution due to the advantages of comparable or higher adsorption capacity than commercial adsorbents, easy control and high flexibility. Table 4. Values of adsorption isotherm parameters and correlation coefficients (r2)a for adsorption of fluoride on various adsorbents System CA (-Al2O3) Al-MCSG-1

qL (g kg-1) 10.00 24.33

Langmuir isotherm KL rL2b (m3 g-1) 0.05 0.9360 0.02 0.9668

Freundlich isotherm kF nF ((g kg -1 ) (g m-3 ) -1/n F) 1.34 2.76 1.65 1.78

rF2b 0.9712 0.9914

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The units for Ce and qe are g m-3 and g kg-1, respectively. rL2and rF2 are correlation coefficients for linear regression fittings of experimental data via Langmuir and Freundlich isotherms, respectively.

Figure 12. Correlation of experimental data by means of the Freundlich (---) and Langmuir isotherms ( ─) for the adsorption of Cu (II) on various superparamagnetic materials at a fixed pH value of 6. △, ○ , □, ▽ and ◇: Experimental data of C2, C21, Alsi-MCSG, C1 and C11, respectively. The definitions of the other symbols were referred to those published in Chang et al., Colloid Surf. A: Physicochem. Eng. Asp. 2009, 336 (1-3), 159-166. (doi:10.1016/j.colsufa.2008.11.042)

4. CONCLUSION The preparation and characterization of nano-sized superparamagnetic materials have been introduced in this study. Such superparamagnetic material can be easily designed and synthesized toward a specific treatment purpose. Application of aluminum-type, zirconiumtype and aluminosilicate-type superparamagnetic materials to the removal of contaminants from aqueous solution in heterogeneous system demonstrated their high flexibility, easy control and high adsorption capacity in the water and wastewater treatment. Furthermore, the

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combination of several analytical techniques can well define the crystal and elemental environment of the magnetic materials. The magnetic material can be easily separated from aqueous solution with a simple permanent magnet or commercial high gradient magnetic separator, also representing the great advantage of the heterogeneous system.

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[22] Khalafalla, SE; & Reimers, GW. IEEE Trans. Magn., 1980, 16 (2), 178-183. [23] Gribanov, NM; Bibik, EE; Buzunov, OV; Naumov, VN. J. Magn. Magn. Mater, 1990, 85(1-3), 7-10. [24] Tominaga, M; Matsumoto, M; Soejima, K; Taniguchi, I. J. Colloid Interface Sci., 2006, 299(2), 761-765. [25] Weissleder, R. US. Patent 5,492,814, 1996; Chem. Abstr., 1997, 124, 283285. [26] Sjorgren, CE; Briley-Saebo, K; Hanson, M; Johansson, C. Magn. Reson. Med., 1994, 31(3), 268-272. [27] Itoh, H; Sugimoto, T. J. Colloid Interface Sci., 2003, 265(2), 283-295. [28] Thapa, D; Palkar, VR; Kurup, MB; Malik, SK. Mater. Lett., 2004, 58(21), 2692-2694. [29] Pardoe, H; Chua-anusorn, W; St Pierre, TG; Dobson, J. J. Magn. Magn. Mater, 2001, 225(1-2), 41-46. [30] Babes, L; Denizot, B; Tanguy,G; Le Jeune, JJ; Jallet, P. J. Colloid interface Sci., 1999, 212(2), 474-482. [31] Vayssières, L; Chanéac, C; Tronc, E; Jolivet, JP. J. Colloid Interface Sci., 1998, 205 (1-2), 205-212. [32] Jiang, WQ; Yang, HC; Yang, SY; Horng, HE; Hung, JC; Chen, YC; Hong, CY., J. Magn. Magn. Mater, 2004, 283(2-3), 210-214. [33] Jolivet, JP. Metal Oxide Chemistry and Synthesis. From Solution to Solid State; Wiley: Chichester, U.K., 2000. [34] Qui, X. Chin. J. Chem., 2000, 18, 834. [35] Sun, S; Zeng, H. J. Am. Chem. Soc., 2002, 124, 8204. [36] Laurent, S; Forge, D; Port, M; Roch, A; Robic, C; Elst, L.V; Muller, R.N. Chem. Rev., 2008, 108(6), 2064-2110. [37] Jehng, JM; Hu, H; Gao, X; Wachs, IE. Catal. Today, 1996, 28(4), 335-350. [38] Katz, E; Shipway, AN; Willner, I. In Nanoscale Materials; Liz-Marzán, L.M; Kamat, P.V; Eds; Kluwer Academic Publishers: Boston, MA, 2003; pp 5-78. [39] Watton, SP; Taylor, CM; Kloster, GM; Bowman, SC. In Progress in Inorganic Chemistry; Karlin, K.D; Ed; Volume 51; John Wiley & Sons, Inc.: Hoboken, NJ, 2003; pp 333-420. [40] Legrand, AP. The Surface Properties of Silicas; John Wiley: New York, NJ, 1998. [41] Iler, RK. The Chemistry of Silica-solubility, Polymerization, Colloid and Surface Prosperities, and, Biochemistry; Wiley: New York, NY, 1979. [42] Alcala, MD; Real, C. Solid State Ion., 2006, 177(9-10), 955-960. [43] Brinker, CJ; Scherer, GW. Sol-gel Science: the Physics and Chemistry of Sol-gel Processing; Academic Press: Boston, MA, 1990. [44] Hench, LL; West, JK. Chem. Rev., 1990, 90(1), 33-72. [45] Cannas, C; Casula, MF; Concas, G; Corrias, A; Gatteschi, D; Falqui, A; Musinu, A; Sangregorio, C; Spano, G. J. Mater. Chem., 2001, 11(12), 3180-3187. [46] Tartaj, P; Serna, CJ. J. Am. Chem. Soc., 2003, 125(51), 15754-15755. [47] Tartaj, P; González-Carreño, T; Serna, C.J. Langmuir, 2002, 18(12), 4556-4558. [48] Matosukas, T; Gulari, E. J. Colloid Interface Sci., 1998, 124(1), 252-261. [49] Satoh, T; Akitaya, M; Konno, M; Saito, S. J. Chem. Eng. Jpn. 1997, 30(4), 759-762. [50] Kim, KD; Kim, HT. J. Sol-Gel Sci. Technol, 2002, 25(3), 183-189. [51] Schüth, F; Unger, K. In Preparation of Solid Catalysts; Ertl, G; Knözinger H; Weitkamp, J; Eds.; Wiley-VCH: New York, NJ, 1999; pp 60-84.

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In: Nanopowders and Nanocoatings Editor: V. F. Cotler

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Chapter 4

EFFECTS OF AG/IN ADDITIVES AND CRYSTALLIZATION KINETICS ON THE RESISTIVE CHARACTERISTICS OF AMORPHOUS SBTE CHALCOGENIDE FILMS Chung-Wei Yanga*, Chien-Chih Choua and Truan-Sheng Luia a

Department of Materials Science and Engineering, National Cheng Kung University No. 1 University Road, Tainan 701, Taiwan, R.O.C.

ABSTRACT Copyright © 2009. Nova Science Publishers, Incorporated. All rights reserved.

Chalcogenide films were able to be used as the recording layer of phase change recording media and in the application of phase change random access memory (PCRAM). The most attractive property of this material is its quick transformation between the amorphous and crystalline states, which phenomenon can accompany huge changes in the optical and electric properties. The reversible transformation between amorphous and crystalline phases was named as Ovonic Memory phenomenon and materials with such kind of properties were also named as the phase change materials. In practical applications, major efforts have been focused on increasing the crystallization speed and improvement on the optical or electrical contrast between amorphous and crystalline state. In the present study, there are two chalcogenide films being deposited on alkali-free glass with RF-sputtering method, pure SbTe films (ST) and Ag/In added SbTe films (AgInSbTe, AIST). In the first part, the microstructure and sheet resistivity of AIST films deposited with different parameters were analyzed. The results show that the as-deposited films possess amorphous structure no matter with what the sputtering parameter being adopted. The sheet resistivity measurement shows the amorphous films possess an extremely high resistivity and the temperature coefficient of resistivity (TCR) is negative. It is worth noting that the relationship of amorphous AIST films between the sheet resistivity and film thickness was found to against the classic size effect. In the second part, similar amorphous films were annealed isothermally at different temperatures to obtain their different crystallinity. The sheet resistance of the annealed *#

Corresponding author: email id: [email protected], Telephone: 82-53-660-6890, Fax: 82-53-422-9631

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Chung-Wei Yang, Chien-Chih Chou and Truan-Sheng Lui specimens was measured at room temperature, where the sheet resistance of amorphous films can be 3 × 104 higher than that of the crystalline films. As comparing X-ray diffraction patterns of AIST films to that of ST films, the sheet resistance change of the specimens can be correlated to the crystallization of amorphous phases, which transition temperature of the change in the sheet is at about 433 K for AIST films and 393 K for ST films. Through transmission electron microscopy (TEM) observations and Grazingincidence X-ray diffraction (GI-XRD), the major phase in the crystalline ST films is the -Sb phase and the mixture of -Sb and AgSbTe2 phases in the crystalline AIST films. Concerning of the thermal activation measurements, the activation energy and crystallization temperature were measured with the differential scanning calorimeter (DSC). The activation energy of AIST films is about 0.92 eV and that of ST films is about 0.82 eV, the crystallization temperature of AIST films is about 475 K and that of ST films is about 445 K. The result reveals Ag/In added SbTe films possesses high room temperature stability. Because the sheet resistance has been proven to change with the crystallinity, an apparatus was developed to estimate the activation energy and the Avrami exponent of crystallization through Johnson-Mehl-Avrami formulism. The activation energy is estimated to be about 0.815 eV and the Avrami exponent (n) is about 1.1 to 1.4. The exponent indicates that the crystal can grow freely and the sheet resistance will decrease dramatically after the impingement effects occurring. A model is proposed to explain why the sheet resistance decreases within a very short period and the homogeneous nucleation and free growth during isothermal annealing in this study.

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1. INTRODUCTION Chalcogenide films refer to alloys containing at least one Group VI element such as the alloys of GeSbTe and SbTe compounds [1-3]. A specific characteristic of the chalcogenide films is that they can be quickly and reversibly transformed between an amorphous state and a crystalline state by using a focused laser beam irradiating a tiny zone or an electrical current at different pulses. It is recognized that Sb-Te binary compound alloying of Ag/In (the AgInSbTe compound, AIST) possesses better cycling stability of the amorphous phase and acquires sensitive ability of crystallization [4, 5]. The optical contrast and the electric properties corresponding to the change of crystal structures are expected to be significantly different [3-6]. This phase change technology has been successfully applied in optical data storage such as the recordable discs [4, 7-9]. Figure 1 shows the equilibrium phase diagrams of binary Sb-Te alloys and AgInTe2-Sb-Te alloys [7], and Figure 2 represents the schematic illustrations of the multilayer structure and the record/erase process of chalcogenide phase transformation recording media. In the last decade, many researchers have put their efforts into optical properties and alloys with a focus on disc performance. It is postulated that the reversible transformation of amorphous-to-polycrystalline phases might be correlated with the optical applications [8, 10, 11]. Previous study investigated the crystallization of glass through in situ capacitance measurement and found that the size of nano-crystallites after an abrupt change in capacitance change was in the range of 10 nm [12]. The nano-sized crystallites and related structures of the crystallized AIST film may make it difficult to recognize the effects of the crystalline phases on the electrical properties. Other than its use in optical recordable discs, the phenomenon results in the crystallization accompanies an electrical conductivity/resistivity change for most chalcogenide materials. Recently, researchers have been trying to develop a

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new function of the chalcogenide phase change material, non-volatile memory [3, 13, 14]. Two stable structural states of chalcogenide alloys, the amorphous state and polycrystalline state show a significant difference in sheet resistance. Information therefore can be recorded and erased in the memory cell through resistive heating. Although it is a common phenomenon that most amorphous materials possess higher resistance, the stability of amorphous chalcogenide films and their faster switching speed still attract much attention. Sb2Te3 or GeSbTe materials were used to fabricate chalcogenide memory recently and the literatures reported that the crystallization degree can be programmed to have different resistance states [14, 15]. Several articles reported that GeSbTe films can be successfully used to integrate Ovonic Unified Memory (OUM), which is a kind of non-volatile memory that utilizes the reversible amorphous-to-crystalline characteristics in chalcogenide alloys material [3, 6, 13]. The SbTe films were also reported to have the potential for the fabrication of chalcogenide non-volatile memory by using the focus ion beam method [14], and AgInSbTe (AIST) films are also expected to be used in OUM chips. This characteristic of the amorphous chalcogenide alloy films can be applied for the phase change random access memory (PCRAM) [5, 15, 16]. Such a technology can satisfy the requirements for non-volatile memory because both the amorphous and polycrystalline structures are stable at room temperature [3, 6, 13, 14]. In practical applications, OUM exhibits a high/low resistance ratio of about 10 to 100, depending on its programmed states, film structures and material properties [13, 15]. A key physical property for OUM applications is the sheet resistance with dependence on different crystalline structures. The thermoelectric power and electrical resistivity of crystalline SbTe films have been studied. The sheet resistivity and Seebeck coefficient of SbTe films can change with the crystalline state and the appearance of the second phase [17, 18]. However, there are few evidences reported on the relationship between microstructure characteristics and its corresponding electrical properties. The structural change of as-deposited SbTe films caused by post-heat treatment is investigated in terms of crystallization temperature and thus, activation energy in the previous study [2, 17-23].

Figure 1. Equilibrium phase diagrams of (a) the binary Sb-Te alloy and (b) the AgInTe2-Sb-Te alloy systems.

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Figure 2. The schematic illustrations of the multilayer structure and the record/erase process of chalcogenide phase transformation recording media.

Previous study indicated that SbTe chalcogenide films with Ag and In additives have a higher crystallization temperature and better stability than the SbTe binary system [24]. However, few reports focus on the kinetics of sheet resistance change of chalcogenide films and the relationship with the crystallization behavior. Therefore, it is important to understand the difference between kinetics of sheet resistance change and those of crystallization during the development of chalcogenide memory. The aim of this chapter is to clarify the electrical properties for both SbTe and AgInSbTe films in association with the microstructural characteristics with relation to the critical crystallization temperature, and the differences between SbTe and AgInSbTe films are also discussed. Since crystallization-dependent changes in sheet resistance are usual for chalcogenide films, an alternative method by measuring the sheet resistant variation with time is probably applicable to study the crystallization kinetics of amorphous AIST films. Apparatus has been developed to perform isothermal sheet resistance measurements to obtain the crystallization activation energy and the Avrami exponent, which is subsequently compared with the results of DSC measurement. The crystallization incorporation with sheet resistance changes in AIST films and the corresponding kinetics are discussed in the present study.

2. VARIATION OF MICROSTRUCTURE AND ELECTRICAL CONDUCTIVITY OF AMORPHOUS SBTE (ST) AND AGINSBTE (AIST) FILMS DURING CRYSTALLIZATION Alkali-free glass sheets of dimensions 30 (l) × 20 (w) × 1 (t) mm3 were selected as substrates for the radio-frequency (RF) magnetron sputtering process. Prior to sputtering, the substrates were degreased and ultrasonically cleaned in acetone, and then they were dried with clean compressed air. Figure 3 shows the schematic apparatus of RF-sputtering system, which was used to deposit films on the glass substrates in this study. The substrate-to-target distance was fixed at 52 mm. A 13.56 MHz RF power supply was connected to the target in a coater, which was pre-pumped to a pressure of less than 6.67 × 103 Pa. The sputtering pressure was maintained at 0.533 Pa with an Ar gas flow. A sputtering parameter of 200 W sputtering power and 60 s sputtering duration were performed to deposit SbTe (ST) and silver/indium (Ag/In) additives AgInSbTe (AIST) chalcogenide films using two different targets. The chemical compositions of SbTe and AgInSbTe targets are listed in Table 1.

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Figure 3. The schematic apparatus of RF-sputtering system used in this study.

Table 1. Chemical compositions of SbTe and AgInSbTe sputtering targets and as-deposited films (at. %).

Target

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Film

SbTe AgInSbTe SbTe AgInSbTe

Ag  5.0 3.4

In 5.5 5.2

Sb 59.8 59.0 68.9 60.8

Te 30.2 30.5 31.1 30.6

Film thickness was measured within five analyzed specimens by the -step equipment (VEECO Dektop3 ST surface profiler). The film thickness was recognized as approximate 150 ± 10 nm and the thickness difference between ST and AIST specimens was less than 10% due to the similar deposition rate of 3.8 nm/s. The chemical compositions were analyzed by induction-coupled-plasma spectroscopy (ICP) and recognized similar Sb/Te atomic ratio for both ST and AIST films as listed in Table 1. The crystalline structures and phase composition were identified by grazing-incidence X-ray diffractometer (GI-XRD, Rigaku D/MAX2500), using Cu K radiation at 30 kV, 20 mA with an X-ray incident angle of 1° and a scan speed of 1° (2)/min. Figure 4(a) shows the GI-XRD spectra of ST and AIST films. The broadened patterns appear at low diffraction angles prove that as-deposited ST and AIST films possess amorphous states. Figure 4(b) is the TEM (JEOL 3010) microstructural feature of asdeposited ST film. It can be recognized that the crystalline state belongs to amorphous -Sb structure through the analysis result of selected area electron diffraction (SAD) pattern. Considering to the GI-XRD pattern of the AIST film displayed in Figure 4(a), the broadened peak shift to about 30º (2) after the Ag/In additives. In addition, the AIST film exhibits similar amorphous microstructural feature and SAD pattern compared with the ST film as shown in Figure 4(c). According to the research results of the EXAFS analysis method by H. Tashiro et. al. [11], the atomic bonding state of Sb element is slightly changed with the Ag/In additives. The atomic bonding measuring results show the bonding distance is almost the same for the amorphous state of Sb and Te elements. Therefore, it is recognized that the as-deposited AIST films should be close to the amorphous -Sb structure.

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Figure 4. (a) GI-XRD spectra of RF-sputtered ST and AIST films. FE-TEM images and selected area electron diffraction (SAD) of as received (b) ST and (c) AIST films.

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2.1. Crystallization Characteristics of Sb and AIST Films The differential scanning calorimeter (DSC) was used for the thermal analysis of AIST and ST films with a heating rate of 0.17 K•s1. The analysis results are shown in Figure 5. Exothermal peaks appeared at 445 K for the ST curve and at 476 K for the AIST curve, which representing the crystallization temperatures of ST and AIST films. This result is compatible with other studies [5, 8], but the crystallization temperature of ST film is found to be lower than that for AIST film in the present experiment. Figure 6 shows the GI-XRD spectra of ST and AIST films before and after crystallization at various heating temperatures, where the broad halo peak at low diffraction angles reveals that the structures of as-deposited specimens are amorphous structures. The ST film remains amorphous structure at the annealing temperatures below 403 K. Once the annealing temperature rises to above 413 K, the -Sb phase appears and the crystallinity significantly increases with an increase in annealing temperature. The same phenomenon is occurred in AIST films but the initial crystallization temperature is 20 K higher than ST films. In Figure 6(a), there is only the -Sb phase existing in the crystallized ST films after heat treatment. The same phase is recognized in the crystallized AIST film as shown in Figure 6(b). However, some previous studies reported that AgSbTe2 phase can precipitate in the SbTe film with Ag additives after annealing at a temperature above 423 K [5, 8, 25-27]. Based on the experimental data, the crystalline structure of AgSbTe2 is cubic and that of -Sb phase is rhombohedral structure, and the lattice parameters of both phases are similar. Consequently, the diffracted peaks of -Sb phase are not easy to distinguish from those of AgSbTe2 phase. Some literatures also reported that the In atoms can substitute within Sb sites to become as a solid solution if the In content is less than 5 at.% [5, 8, 18]. According to the diffraction patterns shown in Figure 6(b), it represents a good agreement with the mentions that no In-

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related compound is found. Furthermore, AgSbTe2 and -Sb phases are expected to be formed in AIST film after annealing at 523 K in this study. Since the as-deposited ST and AIST films can be crystallized through heat treatment at above 523 K, this heating temperature is adopted as the full-annealing temperature in the present study. In order to evaluate the crystallization degree of the ST and AIST films, we adopted a commonly used index of crystallinity (IOC) defined from the ratio of the strongest main peak intensity of the films (Ic) and the strongest main peak intensity of the standard specimen which was full-annealed at 523 K for 1 hour (Io) according to the relationship IOC = (Ic/Io) × 100% [28-30]. The IOC of the standard specimen is defined as 1.0. The specimens with higher IOC index yield better degree of crystallinity.

(a)

(012)

(012)

(104) (110)

(104) (110)

(202) (202)

523 K

433 K

413 K

403 K

(200) (200) (012) (012) (220) (220) (104) (104)

(222) (222) (202) (202)

(110) (110)

  453 K

 

433 K

  

 

  

 

423 K

393 K

413 K

as-deposited

20

(b)

Intensity, / a.u

523 K

:AgSbTe2

-Sb :Sb AgSbTe2

(b)

:Sb -Sb

(a)

Intensity, / a.u

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Figure 5. Differential scanning calorimeter measurements of ST and AIST films at a heating rate of 0.17 K•s‒ 1.

as-deposited

30

40

2Theta 2

50

60

20

30

40

50

60

2Theta 2

Figure 6. Grazing incidence X-ray diffraction spectra of (a) ST and (b) AIST films after post-heattreatment at different temperatures (◇ denotes AgSbTe2 phase and ▽ denotes -Sb phase).

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Figure 7. (a) The IOC index and (b) the average crystal size versus annealing temperatures for ST and AIST films after annealing at a given temperature.

Figure 7(a) shows the variation of IOC indices for as-deposited films and crystallized films annealing from 393 K to 523 K. Considering the results from Figure 6 and Figure 7(a), the crystallinity of the ST and AIST films after annealing at 403 K and 423 K did not significantly increase significantly (only from 0.1 to 0.15) in comparison with the asdeposited films. After annealing at temperatures higher than 413 K and 433 K for the ST and AIST films, respectively, the IOC indices of the films dramatically increased to about 0.8. Moreover, the main contribution to the increase in the IOC indices can be recognized from the crystallization of -Sb phase from the amorphous matrix according to GI-XRD patterns (Figure 6). The crystal size of annealed ST and AIST films was also estimated with the highest XRD diffraction peak according to Scherrer‘s formula of Eq. (1) [31]. The symbol Dhkl represents the crystal size, the factor  represents the full width at half maximum (FWHM),  is the Bragg‘s angle, and  (1.54 Å ) is the wavelength of Cu K radiation X-ray used in this study. In addition, Stokes and Wilson further derived a relationship at K0 = 1 between the average crystal size  (ACS, Å ) and the FWHM value according to Eq. (2). Figure 7(b) shows the variation of ACS value for the crystallized ST and AIST films annealing from 393 K to 523 K. It is recognized that the ACS is increased with increasing annealing temperatures.

Dhkl 



K0  cos

  cos

(1)

(2)

Figure 4 shows a random atomic modular patterns and SAD patterns refer to the amorphous state as-deposited films. After crystallization treatments, Figure 8 shows the TEM images of ST and AIST films after annealing at 433 K. Some crystallites (denoted as the C region with a black contrast) are observed for both ST and AIST films in the amorphous

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matrix (denoted as the A region). The crystallite size is about 10 nm from the high resolution images. Fine crystallites are also observed in the images of the films after annealing at 523 K as shown in Figure 9, and the AIST film seems to have a larger amount of crystallites than the ST film. Figure 9(b) shows that the AIST film after heat treatment possesses a similar morphology with the crystallized ST film (Figure 9(a)) but the size of the crystallites is smaller.

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Figure 8. TEM images of (a) ST film and (b) AIST film after annealing at 433 K for 1 h.

Figure 9. TEM images of (a) ST film and (b) AIST film after annealing at 523 K for 1 h.

To identify this crystalline phase, the Fast Fourier Transformation (FFT) diffraction patterns using high resolution lattice images were investigated. The FFT pattern of the crystallite denoted D(012) in Figure 10(a) is shown in Figure 10(b), which reveals that the crystallite D(012) should be Sb phase with a zone axis of [121] (d(012) = 3.1 Å ) and a rhombohedral structure. The size of a single -Sb crystallite in the crystallized ST film is estimated to be larger than 20 nm. A partial amorphous state could still remain in the crystallized ST film, as denoted A in Figure 10(a). In

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addition to the amorphous and -Sb phases, a secondary crystallite, denoted T(200), is also observed in the crystallized AIST film as represented in Figure 11(a). The FFT diffraction pattern as shown in Figure 11(b) shows the structure of the AgSbTe2 phase with a cubic structure of a zone axis [100] (d(200) = 3.03 Å ) [27]. It is worth noting that a discontinuous interface boundary exists between the adjacent -Sb crystallites and this may indicate that each crystalline -Sb can nucleate and grow independently. The nucleation and growth of -Sb phase can be easy because the literature has reported that there is little volume fraction change associated with the formation of a nucleus for crystallization, and less strain and surface energy need to be overcome when a nucleus grows from the amorphous matrix [32-34]. The crystallization rates of amorphous AIST and ST films can be expected to be small and be close to each other because the same major phase (-Sb) is formed.

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A

Figure 10. High resolution TEM image of ST film after annealing at 523 K for 1 h: (a) the lattice image and (b) FFT diffraction pattern (A denotes the amorphous phase and D(012) denotes the -Sb phase).

Figure 11. High resolution TEM image of AIST film after annealing at 523 K for 1 h: (a) the lattice image and (b) FFT diffraction pattern (A denotes the amorphous phase, D denotes the -Sb phase and T(200) denotes the AgSbTe2 phase).

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Figure 12. Sheet resistances versus annealing temperatures for ST and AIST films after annealing at a given temperature for 1 h.

The crystallization of amorphous chalcogenide films has been proven by other researchers to be a process involving nucleation and growth [8, 32, 35]. The secondary phase precipitated in the early stage can be expected to enhance the heterogeneous nucleation but the higher crystallization temperature of the AIST film reveals that crystallization may be more difficult for the movement of Ag and In atoms in the amorphous matrix. It is believed that the formation of a small amount of AgSbTe2 phase slightly affected the activation energy and thus, crystallization rate of the amorphous AIST film. TEM observation and GI-XRD analysis showed that no indium compounds precipitated in the crystallized AIST films.

2.2. Electrical Properties of Sb and AIST Films Figure 12 shows that the sheet resistance, which measured using the standard four-pointprobe method [36], decreases with an increase of annealing temperature for AIST and ST films, where the sheet resistance of as-deposited film is 3  104 times larger than the film after annealing at 523 K. Furthermore, the sheet resistance of the amorphous AIST film (asdeposited) is approximately twice as high as that of the amorphous ST film and the ratio is maintained for specimens annealed at different temperatures. The transition temperature from a high resistive state (amorphous state) to a low one (crystalline state) is about 393 K for ST film and 433 K for AIST film. Hence, it seems that the higher sheet resistance of as-deposited AIST film is attributed to Ag and In additives. Furthermore, it is obviously that the drastic change of sheet resistance from a high value to a low value is due to crystallization of amorphous films, as shown in Figures 7 and 12. The temperature coefficient of resistivity (TCR) and the Hall coefficient of amorphous and crystalline AIST films listed in Table 2 indicate that the amorphous AIST film can be recognized as a kind of p-type semiconductor similar with the crystalline film [17, 19, 27].

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Another difference between the above-mentioned amorphous and crystalline films demonstrated that the negative dependence of resistivity on the temperature associated with the characteristics of carrier transportations within the films. The higher carrier density of the amorphous film than that of crystalline film suggests more carriers as a result of the point defects within the amorphous microstructure. It should be noted that the carrier mobility of amorphous films is much smaller than that of crystalline films. This phenomenon implies that more scattering of the carriers in the amorphous film due to the randomly oriented atoms. Table 2. Various electrical properties of amorphous and crystalline AIST films

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Amorphous Crystalline

TCR (ppm/K) 9636 2246

Hall coefficient (cm3/C) 1960 5  104

Carrier density (1/cm3) 1.52  1022 1.41  1015

Carrier mobility (cm2/Vs) 0.384 113.6

The influence of additives on the higher resistance of amorphous and crystalline AIST films than ST films can be mainly resulted from the In element. This phenomenon can be explained by the reduction of the mobility of carriers after increasing the In concentration for In2Te3-Sb2Te3 solid solution, which is because In atoms occupy the lattice position of Sb and concentration of antisite defects is reduced [18, 37]. Concerning of the effect of Ag, the addition of Ag was reported to change the local bonding structure from amorphous to crystalline state, which results in a higher energy needed for crystallization. It means that the amorphous state will be thermally stabilized by the existence of Ag [11]. However, the appearance of the AgSbTe2 phase in the crystalline AIST films did not further change the sheet resistance as compared with the crystalline ST films, which indicates the minor effect of the AgSbTe2 phase on the electrical resistivity. The electronic transportations in chalcogenide glass are different from the amorphous semi-conductors, which is generally attributed to changes occurring in the concentration of so-called valence alternation pairs [23, 38]. The chalcogenide alloy in the polycrystalline state shows a drastic increase in free electron density, which gives rise to a difference in resistivity and reflectivity [1, 13]. That is the reason why the sheet resistance of AIST films can show a drastic decrease from the amorphous to crystalline state. So far, there were two chalcogenide films utilizing such electrical property in Ovonic Unified Memory (OUM) applications: (1) Ge-Sb-Te ternary alloy system, in which the sheet resistivity of amorphous film is around 103 to 104 times compared to that of crystalline one [13, 39]. (2) Sb-Te binary alloy system, in which the sheet resistivity of amorphous film is around 30 times higher than that of crystalline one [14]. In the present study, the sheet resistance of amorphous AIST film is around 104 times compared to that of crystalline film, the value of which is close to the ternary alloy system. Larger difference of sheet resistance between amorphous and crystalline states can be positive to the OUM application because more programmed states (various resistive states) can be achieved for higher capacity of memory device [13].

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3. ACTIVATION ENERGY OF AIST CHALCOGENIDE FILM THROUGH ISOTHERMAL SHEET RESISTANCE MEASUREMENTS According to Kissinger‘s method, the activation energy can be derived from Eq. (3) by assuming that the crystallization rate reaches the maximum value in the transformation process [35, 40].

ln



T

2

C

E 1  R T

(3)

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where  is the heating rate, E is the activation energy, T is the absolute peak temperature in Kelvin (K), parameters R and C are constants. A straight line was obtained by plotting ln( / T2) versus 1/T. DSC measuring results at a constant heating rate (Figure 5) are applied to obtain the crystallization activation energy. The activation energy of crystallization was calculated from the slope of the straight line [40] through the least-squares fitting method, and Figure 13(a) shows the plots calculated from Eq. (3). The activation energies of crystallization for AIST and ST films were obtained as 0.94 eV and 0.82 eV, respectively. The higher crystallization temperature of AIST films indicates that the addition of the elements of Ag and In improved the stability of the amorphous structure of SbTe films [25]. However, the crystallization activation energies of ST and AIST films are fairly similar, which reveals that the crystallization rate is slightly affected by Ag and In additives.

Figure 13. (a) The activation energy obtained using least-squares fitting plots of ln(/ T2) versus 1/T for ST and AIST films. The slope represents the activation energy E over gas constant R. (b) DSC curves measured at different heating rate of 5, 10, 15, 20 and 30 K•min1 from room temperature to 673 K for the AIST films.

In addition, Figure 13(b) shows the results of DSC measurements at five different heating rates of 5, 10, 15, 20 and 30 K•min1 from room temperature to 673 K for the AIST films. There is only one exothermic peak corresponding to the crystallization temperature of amorphous AIST film at each heating rate. For example, the crystallization temperature was 465 K when the heating rate was 10 K•min1, as shown in Figure 13(b). According to GIXRD analysis, the exothermic peak should be subject to the crystallization of -Sb phase. Besides, the crystallization temperature increased with the heating rate as shown in Figure 13(b) and the activation energy of the -Sb crystallization was determined to be 0.92 eV. This value is quite close to the previous data which obtained from Figure 13(a). The activation

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energy of crystallization for -Sb phase can attribute to the movement of Sb and Te atoms in a short range and the rearrangement of the crystalline structure because the lattice site is occupied by all the constituent elements at random [11, 33, 41]. Consequently, the activation energy of -Sb crystallization can be affected by the following factors: rearrangement of the Sb atoms, diffusion in a short range and the additive atoms as an obstacle to the movement of the Sb or Te atoms [32, 42, 43]. Ar gas thermal couple

recorder

sample

voltage meter

power supply

vacuum PID controller

Figure 14. The schematic apparatus of the isothermal sheet resistance measuring system.

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3.1. Isothermal Sheet Resistance Measurements The crystallization activation energy of AIST films can be measured by the DSC method, but the Avrami exponent cannot be obtained from Kissinger‘s method because the equation of Kissinger‘s method is derived regardless of reaction order [40]. A measurement of optical properties is interpreted in terms of Johnson-Mehl-Avrami kinetics for crystallization of amorphous GeTe films by isothermal heating [44]. Time- and temperature-dependent changes in the sheet resistance indeed provide a method to study the crystallization of the amorphous films. Therefore, a new apparatus was designed to perform the isothermal sheet resistance measurements in situ at different temperatures. Figure 14 represents the schematic apparatus of isothermal sheet resistance measuring system. Four tiny probes with a tip diameter smaller than 0.1 mm were placed on the specimen and pressured by a loaded-spring to ensure the ohmic contact between probes and specimens. A fixed current (2 mA) was injected through two outside probes from a DC power supply. The apparatus was installed in a sealed chamber and heated in an Ar gas atmosphere. The temperature was controlled by a PID controller to maintain the temperature deviation within 3 K, and the sheet resistance was obtained in situ by recording the voltages between two inside probes with a multi-meter during the isothermal annealing. Three specimens were selected to be quenched by interrupting annealing and blowing Ar gas to preserve the crystalline state at designated heating durations. Since the sheet resistance of AIST films has been proven to change with the crystallization degree or IOC indices, the crystallization rate can be evaluated through monitoring the variation in sheet resistance during the isothermal heat treatment. This method, isothermal sheet resistance measurement, was the first time to be adopted to

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represent the crystallization kinetics of amorphous AIST film in this study. Figure 15(a) shows a curve of AIST film obtained by an isothermal sheet resistance measurement performed at 433 K. The total heating duration was about 1200 s and the corresponding temperature profile in the heating duration is shown as a broken line in Figure 15(a). The

(a)

3K

Fig. 17

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(b)

Figure 15. (a) The plot of sheet resistance versus heating duration obtained by isothermal sheet resistance measurement performed at 433 K. (b) Select-area electron-diffraction patterns of specimen A, B and C corresponding to Stage I, II and III respectively. The specimens were performed the isothermal sheet resistance measurement at 433 K and quenched after the heating duration = 100, 650 and 900 s.

three stages were as follows: Stage I, II and III are defined in Figure 15. In the first stage (Stage I), the temperature was still rising and had not reached 433 K yet. The dramatic decrease of sheet resistance with increasing temperature corresponds to the negative temperature coefficient of resistance (TCR) of amorphous AIST film in Stage I. The sheet resistance could be reversed if annealing was interrupted and the specimen temperature returned to the room temperature. This phenomenon indicates that a thermal activation mechanism of mobile carriers from valence band is responsible for the change of sheet

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resistance in Stage I. When the annealing duration exceeds 100 s, the temperature profile (broken line) changes its slope to a lower value and the heating rate slows down. Meanwhile, the change of the sheet resistance curve becomes flat. When the annealing duration exceeds 400 s, -Sb phase starts crystallizing significantly in Stage II. Most of the crystallization would be occurred in this stage. A secondary sheet resistance transient was subsequently observed when the annealing duration was 750 s. Afterward the sheet resistance curve becomes stable again in Stage III. The crystallization might carry on, but the rate would be slow. Three specimens A, B and C were quenched after different durations denoted in Stage I, II and III, respectively. Figure 15(b) shows the SAD patterns of three quenched specimens. An amorphous feature of specimen A reveals that crystallization had not occurred in the early period of Stage I. The SAD pattern of specimen B shows crystalline structures formed in the specimen but the retaining diffused nature of the diffraction rings indicates the crystal size is small. The disappearance of the diffuse ring pattern indicates an improvement in crystallinity for specimen C. Figure 16 shows the TEM images of specimen C quenched after sheet resistance transient. A lot of crystallites, denoted O, had formed in the film but some amorphous areas, denoted A, still retained. The d spacing calculated with the high resolution images referred the crystallite to -Sb phase as shown in Figure 16(b). TEM observations ensured the crystallization of -Sb was associated with sheet resistance transient in Stage II.

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(a)

(b)

Figure 16. TEM observation of specimen C, which was quenched after heating duration = 900 s. (a) TEM image, and (b) high resolution image of the area circled in Figure 17(a), O denotes the area with ordered structures and A denotes the area with amorphous structures. The SAD pattern of O referred to the crystallite -Sb phase.

3.2. Activation Energy and Avrami Exponent for Isothermal Crystallization For further analysis, the local curve of Stage II in Figure 15(a) was re-plotted in Figure 17. The sigmoid feature shown in Figure 17 can be incorporated into the Johnson-MehlAvrami formalism [43-45]. The transformed volume fraction of AIST film is correlated to the change in isothermal sheet resistance. Isothermal sheet resistance measurements were repeated at annealing temperatures of 423, 433 and 443 K to obtain three JMA curves as

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shown in Figure 18. The duration of 50% sheet resistance change (t1/2) for each annealing temperature was determined by the method described in Figure 17 and listed in Table 3.

Figure 17. Demonstration of how to determine t1/2 with the curves obtained by isothermal sheet resistance measurement, where R1/2 R2 = R1 R1/2.

Table 3. The durations of t1/2 obtained at various annealing temperatures in isothermal sheet resistance measurement. t1/2 (s) 895 750 410

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Annealing temperature (K) 423 433 443

Figure 18. Three curves of sheet resistance versus heating durations obtained using isothermal sheet resistance measurement. The specimens were performed at 423, 433 and 443 K for curve I, II and III respectively. Nanopowders and Nanocoatings: Production, Properties and Applications : Production, Properties and Applications, Nova Science Publishers,

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A quantitative description of the isothermal crystallization kinetics of amorphous solids is able to be made using the classical JMA formalism. The classical JMA formalism can be presented as Eq. (4), where F is the volume fraction transformed (equivalent to the fraction of sheet resistance change) at time t, k is the rate constant, and n is the Avrami exponent [43-45].

F  1  exp( kt n )

(4)

Equation (4) can be represented as Eq. (5) by taking nature logarithm,

ln[ln(1  F ) 1 ]  ln k  n ln t

(5) In a condition of 50% volume transformed, F = 1/2 is introduced to Eq. (5) and derives Eq. (6), where k is the rate constant which is mainly affected by the annealing temperature and the activation energy.

ln[ln 2]  ln k  n ln(t1 / 2 )

(6)

The rate constant k can often be represented in a form of the Arrhenius equation with an activation energy Q [45, 46]:

k  k0 exp(Q/RT)

(7)

Equations (6) and (7) can be combined as Eq. (8),

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

C  n ln t1 / 2  

Q1 RT

(8)

where C is a constant related to k0, and R is the gas constant. The activation energy Q of crystallization can be obtained through regressive plots of ln(t1/2) versus 1/T. However, the Q still depends on the Avrami exponent n, Eq. (5) can be used to obtain a plot of ln[ln(1F)1] versus ln(t) with constant k (isothermal transformation). The plot gives a straight line with a slope of n. The regressive plots for Avrami exponents at annealing temperatures of 423, 433 and 443 K are shown in Figure 19(a). The slope of each straight line was subjected to the Avrami exponents of each isothermal annealing temperature. The range of n value was from 1.1 to 1.4 within the isothermal annealing temperatures from 423 to 443 K. Another regressive line for calculating activation energy using Eq. (8) was subsequently plotted with different annealing temperatures as shown in Figure 19(b). The slope of Q/R, gives an activation energy of 0.815 eV when adopting average n = 1.25. The activation energy of the AIST film obtained by the isothermal sheet resistance measurement (0.815 eV) was similar to the result using DSC measurements (0.92 eV). The main difference between the Kissinger‘s method by DSC measurement and isothermal sheet resistance measurement is the heating manner, which of DSC method was performed at a constant heating rate while isothermal measurement was performed at a fixed temperature. For DSC measurement, the maximum exothermic energy correlated to the maximum crystallization rate is used to obtain the activation energy using Kissinger‘s method. However, the duration was as a result of most

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crystallites forming in amorphous films while the maximum variation in the sheet resistance occurred in isothermal sheet resistance measurement. This phenomenon indicates that the sheet resistance change was able to be used as a physical parameter, like the exothermic peak in DSC measurement, to calculate the relative kinetics of crystallization. The Avrami exponent n which characterizes crystallization behavior can only be determined by isothermal measurements and cannot be obtained through DSC measurement because the equation from Kissinger‘s method is derived with regardless of reaction order [40]. The Avrami exponent typically depends on the dimensions associated with crystal growth and the time dependence of nucleation [43-45]. Theoretical analysis suggests that n = 2.0 to 3.0 if transformation occurs in a thin sheet of solid material, in which the average dimensions of a transformed region will be much larger in planar surface than in thickness on account of the ―thin‖ nature [46]. However the Avrami exponent was only n = 1.1 to 1.4 in this study. In fact, many anomalous Avrami exponents have been observed in the crystallization processes of various metallic glasses and anomalous exponents can be caused by various reasons such as: (i) different mechanisms controlling the crystallization process, (ii) a time dependent nucleation rate, (iii) possibility of simultaneous grain growth of the crystallized region during crystallization, and (iv) impingement effect, which is important especially at the final stage of crystallization if the anisotropy of the crystals is sufficient high [43, 46-49]. Some experiments and mathematic simulations have been done for local and average Avrami exponents in the crystallization process of various amorphous alloys but not for AIST films [43, 48, 50]. A theoretical Avrami exponent n is suggested to be 1.5 for transformations with parabolic growth and the volume transformation rate for spherical growth in the absence of impingement (free growth) can be obtained with the assumption of a random distribution of pre-existing nuclei or constant nucleation rate [46].

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(a)

(b)

Figure 19. The regressive plots of Avrami exponent and activation energy for AIST film. (a) ln[ln(1F)1] versus ln(t) for different annealing temperatures. The slope of each straight line represents the Avrami exponent n. (b) n[ln(t1/2)] versus the inverse of annealing temperature. The slope of the straight line was subjected to Q/R, where Q is activation energy and R is gas constant.

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4. EFFECT OF CRYSTALLIZED PHASES IN SHEET RESISTANCE OF AMORPHOUS AIST CHALCOGENIDE FILM The above-mentioned results of TEM lattice images and FFT diffraction patterns for the AIST films before and after isothermal annealing at 523 K for 1 h can help to provide the correlation with electrical-crystallization and crystalline phases for AIST films. From the lattice images, Figure 11(a) provides further evidence to verify the microstructural features, which also confirms that there are still a little amount of amorphous areas, as denoted by A. Furthermore, different lattice planes are observed in Figure 11(a), the d spacing between the (012) lattice planes of the crystallite denoted in D(012) is 3.1 Å , which is very close to the d spacing as also reported in the literature (3.1042 Å ) for the rhombohedral structure of -Sb phase [51]. Based on experimental evidence as shown in Figure 11(a), a coherent interface between T(200) and D(012) crystallites can be confirmed. The interface is indicated by the triangle in Figure 11(a). The diffraction pattern of the FFT image as shown in Figure 10(b) confirms the -Sb phase has a rhombohedral structure with a zone axis [121] (d(012) = 3.1 Å ).

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Figure 11(b) shows the AgSbTe2 phase has a cubic structure with a zone axis [100] (d(200) = 3.03 Å ). It is worth noting that there are several discontinuous -Sb crystallites, as marked by D in Figure 11(a), in which the -Sb phase possesses different lattice planes. It is appropriate to suggest that AgSbTe2 crystallites play an important role as nucleation sites for the crystallization of -Sb phases due to the AgSbTe2 is suggested to precipitate at a low temperature [8, 26]. Figure 20 shows the schematic description of the crystallization when the isothermal sheet resistance of the amorphous AIST films was measured. The sheet resistance drastically changes with the progress of crystallization, which describes the parabolic growth of the individual crystallites from the amorphous matrix. The sheet resistance can be expressed as a combination of the amorphous state and the crystalline phase within an AIST film (i.e. Rtotal = Ramorphous + Rcrystalline as shown in Figure 20(a)) [52]. Generally, the sheet resistance of the amorphous AIST film is much higher (about 3  104 times) than that of the crystallite state, and the phenomenon of resistance change resulted from a current channel that developed after a critical amount of the crystallites are adjacent. Consequently, the sheet resistance of the film can be high when the crystallites are small and few. Once the crystallites grow and make contact with each other, the sheet resistance decreases dramatically due to the high current paths being connected. The critical volume fraction of the crystallites for developing above mentioned current channels is about 78% for the 2-D growth of thin films. Based on the experimental results as shown in Figure 7(a) and Figure 12, the IOC data shows a consistent result with the above calculation. It is obvious that the contact of each crystallite is responsible for the increase of high current paths and the dramatic decrease of sheet resistance (Figure 20). The Avrami exponent is derived from the assumed highest changing rate of sheet resistance in the present study (Figure 17) and hence, before the changing rate becomes significantly slow, the free-growth-like crystallization could propose an Avrami exponent of 1.1 to 1.4 in AIST film. When the high current paths are well established, which corresponds to the maximum volume percentage of the crystallites, the sheet resistance reaches a stable low-resistance state.

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Figure 20. A schematic description of the crystallite growth and transient sheet resistance change after crystallizations. (a) The parabolic crystallite growth, the sheet resistance of the film is a combination of Ramorphous and Rcrystalline, and (b) a current channel developed by the impingement of the crystallites.

5. CONCLUSION

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The effects of Ag/In additives and crystallization kinetics on the resistive characteristics of amorphous SbTe chalcogenide films has been characterized. The results can be summarized as follows: (1) The structures of the as-deposited ST and AIST films are a single amorphous state and their sheet resistances show an extremely high value. No indium (In) related compound is found in the crystallized films when the In content is about 5 at.%. (2) The major phase -Sb forms in the crystallized AIST films when the annealing temperature is above 433 K. After annealing at 523 K, the -Sb phase is formed in the crystallized ST films, while AgbSTe2 and -Sb phases are formed in the crystalline AIST film through the XRD analysis and TEM observations. Besides, the precipitation of secondary AgSbTe2 phases in AIST film result in the production of fine -Sb phases. (3) The sheet resistance decreased with increasing the annealing temperature for ST and AIST films, where the sheet resistance of as-deposited film is about 3  104 times larger than the film after annealing at 523 K for 1 h. The variation of sheet resistance can be correlated to the crystallization of amorphous phases, which transition temperature of the changes for ST films and AIST films is at about 393 K and 433 K, respectively. Moreover, the sheet resistance of SbTe films with Ag and In additives is twice as high as that of ST films before and after crystallization. (4) The activation energy obtained using the isothermal sheet resistance measurement is 0.815 eV, which is similar to that obtained by DSC method. The Avrami exponent (n) was found to be about 1.1 to 1.4 by the isothermal sheet resistance measurement and this value is slightly smaller than 2.0 to 3.0 for typical transformation in thin

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Chung-Wei Yang, Chien-Chih Chou and Truan-Sheng Lui films. It is suggested that impingement effects are the cause of the smaller Avrami exponent in isothermal sheet resistance measurement. (5) A current channel can be developed because annealing the crystallization phase up to 78% is a dominant factor that drastically reduces the sheet resistance of AIST films.

ACKNOWLEDGMENTS The authors appreciate Prof. Li-Hui Chen for the discussion and the assistance in the conception of crystallization kinetics modeling. This study was financially supported by the National Science Council of Taiwan (Contract No. NSC 95-2221-E-006-118) for which we are grateful.

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[15] Hosaka, S., Miyauchi, K., Tamura, T., Yin, Y. & Sone, H. (2003). In Proc. 15th Symposium on Phase Change Optical Information Storage; Ide, T., Ed., The Society of Phase Change Recording, pp 52-55. [16] Maimon, J., Hunt, K., Rodgers, J., Burcin, L. & Knowles, K. (2002). Proceedings of 2002 Non-Volatile Memory Technology Symposium; Nov. [17] Das, V. D. & Soundararajan, N. (1989). J. Appl. Phys., 65, 2332-2341. [18] Dhar, S. N. & Desai, C. F. (2002). Philos. Mag. Lett., 82, 581-587. [19] Kemper, M. J. H.& Oosting, P. H. (1982). J. Appl. Phys., 53, 6214-6217. [20] Raaijmakers, I. J. M. M., Van Ommen, A. H., Reader, A. H. (1989). J. Appl. Phys., 65, 3896-3906. [21] Sanchez, E. M., Prokhorov, E. F., Hernandez, J. G. & Galvan, A. M. (2005). Thin Solid Films, 471, 243-247. [22] Palmer, W. & Marinero, E. E. (1987). J. Appl. Phys., 61, 2294-2300. [23] Pattanaik, A. K. & Srinivasan, A. (2004). Semicond. Sci. Technol, 19, 157-161. [24] Tashiro, H., Harigaya, M., Kageyama, Y., Ito, K., Shinotsuka, M., Tani, K., Watada, A., Yiwata, N., Nakata, Y. & Emura, S. (2002). Jpn. J. Appl. Phys., 41, 3758-3759. [25] Chen, Y. M. & Kuo, P. C. (1998). IEEE Trans. Magn., 34, 432-434. [26] Zayed, H. A., Ibrahim, A. M. & Soliman, L. I. (1996). Vacuum, 47, 49-51. [27] Lakshminarayana, D. (1991). Thin Solid Films, 201, 91-96. [28] Yang, C. Y., Wang, B. C. & Wu, J. D. (1995). J. Mater. Sci.: Mater. Med., 6, 249-257. [29] Kweh, S. W. K., Khor, K. A. & Cheang, P. (2000). Biomaterials, 21, 1223-1234. [30] Yang, C. W. & Lui, T. S. (2008). J. Euro. Ceram. Soc., 28, 2151-2159. [31] Cullity, B. D. & Stock, S. R. (2001). Elements of X-ray Diffraction; Prentice-Hall, Inc., New Jersey, pp 170. [32] Fujimori, S., Yagi, S., Yamazaki, H. & Funakoshi, N. (1988). J. Appl. Phys., 64, 1000-1004. [33] Arun, P., Tyagi, P., Vedeshwar, A. G. & Paliwal, V. K. (2001). Physica B, 307, 105-110. [34] Barrett, C. R., Nix, W. D. & Tetelman, A. S. (1973). The Principles of Engineering Materials, Prentice-Hall, Inc., New Jersey, pp163-170. [35] Chiang, D., Jeng, T. R., Huang, D. R., Chang, Y. Y. & Liu, C. P. (1999). Jpn. J. Appl. Phys., 38, 1649-1651. [36] Smith, D. L. (1995). Thin Film Deposition, Principles and Practices, McGraw Hill, pp 575-577. [37] Horak, J., Stary, Z., Lostak, P. & Pancir, J. (1988). J. Phys. Chem. Solids, 49, 191. [38] Mott, N. F. & Davis, E. A. (1979). Electronic Process in Non-crystalline Materials, Clavendon Press, Oxford. [39] Park, T. J., Choi, S. Y. & Kang, M. J. (2007). Thin Solid Films, 515, 5049-5053. [40] Kissinger, H. E. (1957). Anal. Chem., 29, 1702-1706. [41] Matsunaga, T.& Yamada, N. (2003). In Proc. 15th Symposium on Phase Change Optical Information Storage; Ide, T., Ed., The Society of Phase Change Recording, pp 7-12. [42] Men, L., Jiang, F. & Gan, F. (1997). Mater. Sci. Eng. B, 47, 18-22. [43] Sun, N. X., Liu, X. D. & Lu, K. (1996). Scripta Mater, 34, 1201-1207. [44] Libera, M. & Chen, M. (1993). J. Appl. Phys., 73, 2272-2282.

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[45] Christian, J. W. (1981). The Theory of Transformation in Metals and Alloys, Pergamon, Oxford, pp 540-542. [46] Henderson, D. W. (1979). J. Non-Cryst. Solids, 30, 301-315. [47] Greer, A. L. (1982). Acta Metall, 30, 171-192. [48] Ghosh, G., Chandrasekaran, M. & Delaey, L. (1991). Acta Metall, 39, 925-936. [49] Shepilov, M. P. & Baik, D. S. (1994). J. Non-Cryst. Solids, 171, 141-156. [50] Holzer, J. C. & Kelton, K. F. (1991). Acta Metall, 39, 1833-1843. [51] Ghosh, G. (1994). J. Phas. Equilib. Diff., 15, 349-355. [52] Verhoeven, J. D., Downing, H. L., Chumbley, L. S. & Gibson, E. D. (1989). J. Appl. Phys., 65, 1293.

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Chapter 5

CHEMICAL VAPOR SYNTHESIS (CVS) OF INORGANIC NANOPOWDERS H. Y. Sohn and Taegong Ryu Department of Metallurgical Engineering; University of Utah Salt Lake City, Utah 84112, US.A

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ABSTRACT Chemical vapor synthesis (CVS) is a process for making fine solid particles by the vapor-phase chemical reactions of precursors. At the University of Utah, this process has been used for the synthesis of the ultrafine powders of titanium and nickel aluminides and more recently aluminum nanopowder, tungsten and tungsten carbide nanopowders, and tungsten carbide - cobalt nanocomposite powder. This CVS process has proved its capability to prepare fine particles of 5-200 nm sizes. An example of the significant features of this technique is its unique capability to produce very uniformly mixed powders of different solid phases. This is possible because the reactants can be perfectly mixed in the gas phase. The chemical vapor synthesis is typically performed in a tubular reactor but more recently it has been carried out in a plasma reactor system. The plasma assisted chemical vapor synthesis adds many other advantages such as a high processing temperature to vaporize all reactants, a high quench rate to form ultrafine powders, and a wide choice of reactants. Thus, it has shown a considerable promise for many applications as a promising method for producing a variety of nanopowders.

INTRODUCTION Many metals in the form of powders, especially ultrafine powders (UFP) or nanopowders, display useful physical properties. With a large specific surface area, they are desirable raw materials for powder metallurgical processing. They also possess other exceptional properties, including light absorption [Nikklason, 1987], magnetism [Okamoto et *

Corresponding author: Tel: +886-6-2757575 ext. 62931, Fax: +886-6-2380698, E-mail: [email protected] (Chung-Wei Yang)

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al., 1987], and superconductivity [Parr and Feder, 1973]. For much the same reasons, the powders of intermetallic compounds are also offer promising possibilities. Several methods have been practiced in the production of metallic and intermetallic powders. Here, we will summarize developments in the synthesis of such powders from metal chlorides, with an emphasis on the synthesis of intermetallic powders. A reaction between a metal halide and hydrogen can in general be written as follows: MXn(g) + 0.5nH2(g) = M(s) + nHX(g)

(1)

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where M and X represent the metal and the halogen, respectively. The hydrogen reduction of single-metal chlorides for the preparation of metallic UFP has been studied for tungsten and molybdenum [Lamprey and Ripley, 1962], cobalt [Saeki et al., 1978], and nickel, cobalt, and iron [Otsuka et al., 1984]. By reducing vaporized FeCl2, CoCl2, and NiCl2 by hydrogen at 1200 to 1300 K, Otsuka et al. [1984] were able to prepare corresponding metal particles in the size range 52 to 140 nm with up to 99.7 % metal chloride conversion. The synthesis of metal carbide UFP has been practiced by the vapor-phase hydrogen reduction [Zhao et el., 1990; Hojo et al., 1978]. Hojo et al. [1978] produced the UFP of tungsten carbide (WC, W2C) of 40 to 110 nm size by vapor-phase reaction of the WCl6–CH4–H2 system at 1000 to 1400°C. Magnesium is a much stronger reducing agent for chlorides than hydrogen. Thus, the reduction of titanium and aluminum chlorides by magnesium is feasible at temperatures around 1000°C, whereas the reduction of these chlorides by hydrogen is not feasible up to nearly 3000°C. Titanium sponge is prepared by the chlorination of rutile, followed by the reduction of the resulting titanium chloride by liquid or gaseous magnesium [Barksdale, 1966]. A reaction between a metal halide and magnesium vapor can be written as follows: MXn(g) + 0.5nMg(g) = M(s) + 0.5nMgX2(l,g)

(2)

where M and X are the metal and the halogen, respectively. In recent years, Sohn and coworkers [Sohn and PalDey, 1998a - 1998d; Sohn et al., 2004] applied the basic concepts of the above chloride reduction methods to the chemical vapor synthesis of intermetallic and metal alloy powders. These reactions can in general be written as follows, when hydrogen is used: mMClx(g) + nNCly(g) + 0.5(mx + ny)H2 = MmNn(s) + (mx + ny)HCl(g)

(3)

where M and N represent two different metals, x and y being the valences, and MmNn the intermetallic compound formed. Sohn and PalDey [1998a] synthesized fine powder (100-200 nm) of Ni4Mo at 900 to 1100°C using hydrogen as the reducing gas. Sohn and PalDey [1998b] also synthesized nickel aluminide (Ni3Al) particles (50-100 nm) at 900 to 1150°C using hydrogen as the reducing agent. The fact that aluminum chloride is reduced by this reaction scheme is very significant, because the reduction of AlCl3 alone by hydrogen is thermodynamically unfavorable at these temperatures. The negative free energy of formation of the intermetallic compound makes the overall reaction feasible. Using the same chemical vapor synthesis process, Sohn et al. [2004] prepared ultrafine particles of Fe-Co alloys by the hydrogen reduction of FeCl2-CoCl2 mixtures. Sohn and PalDey [1998c; 1998d] also

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synthesized ultrafine powders of the aluminides of titanium and nickel using magnesium as the reducing agent. More recently, this process has been applied to the synthesis of nanosized powders of aluminum [Choi et al., 2007; Sohn et al., 2007], tungsten [Ryu et al., 2009a; Ryu et al., 2009b] and tungsten carbide [Sohn et al., 2007; Ryu et al., 2008a; Ryu et al., 2008b; Ryu et al., 2009c], and nanocomposite tungsten carbide - cobalt [Sohn et al., 2007; Ryu et al., 2008c; Ryu et al., 2008d]. This work has demonstrated that it is possible to prepare fine particles of 20 - 200 nm by CVS. Further, it has also been shown that this technique has a unique capability to produce very uniformly mixed powders of different solids. An example of this feature will be presented later conjunction with the CVS of WC-Co composite powder. Additionally, the process easily allows the incorporation of minor amounts of additives or doping agents, again very uniformly distributed throughout the product powder. The purpose of this work is to summarize the CVS of nanosized powders, which have been done in this research group under an electrically heated tubular reactor and a non-transferred plasma reactor system.

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1. THE SYNTHESIS OF ALUMINUM NANOPOWDER AS A PRECURSOR OF HYDROGEN STORAGE MATERIALS IN A TUBULAR REACTOR A number of hydrogen storage materials with high storage capacities that contain aluminum have been identified recently [Seayad and Atonelli, 2004]. Their hydrogen release and uptake kinetics must be improved and thus the search for new hydrogen storage materials based on nano-grained materials has received much attention. These nano-grained materials can improve hydrogen storage properties by the extremely large surface area to volume ratio and modified surface structures that can significantly change reaction kinetics and increase reversibility [Zaluski et al., 1999]. Nano-grained aluminum powder is an important starting material for numerous aluminum-containing compounds such as NaAlH4 [Í ñiguez and Yildirim, 2005], LiAlH4 [Züttel et al., 2003], Mg(AlH4)2 [Fichtner and Fuhr, 2002], and AlH3 [Graetz et al., 2006] that have been identified to have high potentials. The chemical vapor synthesis (CVS) process has been applied to the preparation of aluminum nanopowders that can enhance the efficiency of the hydrogen storage materials by reducing the path lengths for hydrogen diffusion [Choi et al., 2007; Sohn et al., 2007]. The apparatus used for this purpose is shown in figure 1, which consists of a vertical tubular reactor, an electrically heated furnace, two separate powder feeders, volatilizers, a powder collector, and an off-gas scrubber. The powder feeding system is an entrained-flow feeder. The reactor is an alumina tube with an inner diameter of 7.9 cm and a length of 1.52 m. The external powder feeders were specially designed to deliver fine AlCl3 powders and Mg powders with sizes between 200 and 400 μm into the reactor. Entrained-flow powder feeders independently control the feeding rates of the powders. The input amount is determined by measuring the weight change of the feed before and after the experiment. The fed AlCl3 and Mg powders are evaporated in two separate alumina volatilizers placed inside the reactor with the position of the AlCl3 volatilizer being 15 cm above that for Mg. Ar gas is supplied to control the concentration and residence time of the reactants and to prevent back flow. The produced powder is collected in a collector containing ethanol, and the off-gas is neutralized

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by a NaOH solution. The collected powder is rinsed with pure ethanol to remove unreacted AlCl3 and the MgCl2 produced from the reaction and filtered using a vacuum filter with a pore size of 40 nm. The filtered powder is dried in a glove box under an Ar atmosphere and characterized by means of SEM (Topcon SM-300), EDX (EDAX), and XRD (Siemens D 5000).

Figure 1. Schematic diagram of the experimental apparatus: (1) Ar gas, (2) flow meters, (3) AlCl3 powder feeder, (4) Mg powder feeder, (5) furnace, (6) powder collector, and (7) scrubbing solution [Adapted from Sohn et al. (2007)].

The overall reaction of AlCl3 and Mg in the gas phase is represented by AlCl3(g) + 1.5Mg(g) = Al(g) + 1.5MgCl2(g)

(4)

To investigate the effect of temperature on the equilibrium composition in the AlCl3-Mg-Ar system, thermodynamic calculations have been performed by the use of HSC Chemistry software developed by Outokumpu Research Oy, which is based on the principle of Gibbs free energy minimization. Figure 2 shows that AlCl3 reduction by Mg is favorable in the temperature range evaluated. The experimental conditions were temperature of 1000oC, both volatilizer temperatures of 1000oC, AlCl3 feeding rate of 0.1 g/min, Mg feeding rate of 0.03 g/min, and total Ar flow rate of 8.5 L/min (25ºC, 86.1 kPa total pressure at Salt Lake City).

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Figure 2. Equilibrium composition vs. temperature for a mixture of 1 mol AlCl3-1.5 mol Mg-10 mol Ar. [Adapted from Sohn et al. (2007)].

Figure 3 shows the SEM micrograph and chemical composition of the synthesized powder. The shape of the particles is seen to be spherical with sizes between 100 and 500 nm. From the EDX and XRD analysis shown in figures 3 and 4, the product is seen to be essentially pure Al. The grain size calculated from the XRD pattern using the Scherrer equation [Cullity, 1978] was around 40 nm, indicating that the particles are agglomerates of nanograins. Based on the results presented here, the CVS process has proved to be a useful method for producing nano-grained aluminum powder that is an excellent precursor for many promising hydrogen storage materials.

Figure 3. SEM micrograph and EDX result of the product powder. [Adapted from Choi et al. (2007)].

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Figure 4. XRD pattern of the product powder. [Adapted from Sohn et al. (2007)].

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2. THE SYNTHESIS OF WC-CO COMPOSITE POWDER IN A TUBULAR REACTOR Among the hard alloys and refractory carbides, cemented tungsten carbide has broad industrial applications in such fields as metal working, drilling, and mining industries under severe conditions, thanks to its superior properties of high hardness and excellent wear resistance. Because of its brittleness, tungsten carbide is typically used in the form of a metal matrix composite with cobalt as the matrix. Cobalt contents in WC-Co composites are typically limited to the range of 3 - 30 wt% for commercial applications. The fine size of cobalt powder is also shown effective in the reduction of cobalt pooling during consolidation process. The mechanical properties such as hardness, compressive strength and rupture strength depend on the composition and microstructural parameters [Upadhaya, 2002; Wahlberg et al., 1997; Zhu and Manthiram, 1996]. Nano-grained powders have been produced by various methods, such as the thermochemical spray drying process [McCandlish et al., 1992], mechanical alloying [Puszynski, 2001], and chemical vapor condensation [Chang et al., 1994]. In this research group, the CVS method has been applied to produce WC-Co nanocomposite powder to utilize the many advantages the method provides, including the uniformity of composition when applied to the preparation of a mixture of different solid particles and the small grain size of the synthesized particles. In the work by Sohn and coworkers [Sohn et al., 2007; Ryu et al., 2008c; Ryu et al., 2008d], WCl6 (99.9%) and CoCl2 (99.7%) were used as the precursors, and hydrogen (99.9%) and methane (99.9%) were used as the reducing and carburizing agents, respectively. The apparatus consists of a powder feeder, a vertical reactor, powder collectors, off-gas scrubber,

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and burner, as shown in figure 5. Argon gas is flowed through the powder feeder as the carrier gas as well as to flush the reactor. The reactor is an alumina tube of 5.4 cm inner diameter and 1.50 m length and placed in an electrically heated vertical furnace. The powder produced is collected using a Teflon-coated polyester filter with a pore size of 1 μm. The product is analyzed by XRD (Siemens D 5000), and the morphology is examined by TEM (JEOL, JEM-2000FXII) with energy dispersive X-ray spectrometry (EDS) that also give distribution of the elements. The total carbon contents of the powders are measured with a Carbon Determinator (LECO, CS-444). The grain size of WC is calculated from the XRD pattern using the Scherrer equation calibrated by using standard samples. The average particle size, as opposed to the grain size, is measured by using a Brookhaven Instruments ZetaPALS unit, which is a zeta potential analyzer based on phase analysis light scattering. The molar ratio of the different phases in the product is estimated from XRD patterns using the internal standard method, combined with the following calibration equation [Cullity, 1978];

I X   I X

(5)

where Iα = integrated intensity of α phase peak; Iβ = integrated intensity of β phase peak; Xα = mole fraction of α phase; Xβ = mole fraction of β phase; and κ = constant. The overall reactions for synthesizing the WC-Co mixtures using metal chloride powders are as follows:

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WCl6(g) + xCoCl2(g) + CH4 (g) + (x+1)H2 (g) = WC (s) + xCo (s) + (2x+6)HCl (g) (6)

Figure 5. Schematic diagram of the experimental apparatus: (1) WCl6 powder feeder, (2) CoCl2 powder feeder, (3) thermocouple, (4) vertical furnace, (5) alumina reactor, (6) powder collector, (7) scrubber, and (8) Bunsen burner. [Adapted from Ryu et al. (2008d)].

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The equilibrium composition diagram for the WCl6-CoCl2-H2-CH4 mixture was obtained by the use of HSC Chemistry software developed by Outokumpu Research Oy, as shown in figure 6. The diagram shows that it is thermodynamically feasible to synthesize WC and Co from the WCl6-CoCl2-H2-CH4 mixture.

Figure 6. Calculated equilibrium composition in WCl6-CoCl2-H2-CH4 system with input amounts of WCl6, CoCl2, H2 and CH4 of 1, 0.5, 30, and 3 kmol, respectively.

Production of WC from WCl6: The apparatus was used first to determine the conditions for the synthesis of WC powder alone and the detailed experimental conditions are described in the references by Sohn and coworkers [Sohn et al., 2007; Ryu et al., 2008c]. Figure 7 shows the XRD patterns of the products obtained with three different reaction temperatures. The main product was WC but some W2C formed with the decrease in reaction temperature. As the reaction temperature increased from 1200°C to 1400°C, the grain size of WC increased from 12 ± 1 to 18 ± 1 nm.

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Figure 7. X-ray diffraction patterns of the product obtained at (a) 1400°C, (b) 1300°C, and (C) 1200°C under the following conditions: CH4 feeding rate of 0.1 L/min (25°C, 86.1 kPa), H2 feeding rate of 0.25 L/min (25°C, 86.1 kPa), total flow rate of 1.1 L/min (25°C, 86.1 kPa), and WCl6 feeding rate of 0.04 g/min. [Adapted from Sohn et al. (2007)].

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Production of WC-Co Composite Powder from WCl6 and CoCl2 A mixture of WCl6 and CoCl2 powders was fed to produce WC-Co composite powder [Sohn et al., 2007; Ryu et al., 2008c; Ryu et al., 2008d]. The main product obtained by this system was W2C and even Co3W3C was produced as shown in figure 8(a). Thus, the volatilizers for WCl6 and CoCl2 were placed where the temperatures were 440°C and 1400°C, respectively. With this arrangement, the main product was a mixture of WC and Co with some free carbon as shown in figure 8(b). The grain size of WC was 24 ± 1 nm. The particle size obtained with an input C/W molar ratio of 2.3 was examined using ZetaPALS and TEM. ZetaPALS gave an average particle size of 380 nm but TEM examination showed that the size of individual particles was 25 ± 5 nm, although this measurement is based on a rather limited number of particles. Figure 9 shows the TEM photograph of the WC and Co powders. The difference in the particle sizes obtained by the two methods can be explained by the agglomeration of the particles as well as their motion in liquid media that contributed to errors in measurement by ZetaPALS. In addition, excess carbon was coated on the surface of produced particles, as shown in the TEM micrograph. However, the excess carbon can be removed by post-treatment such as reaction with hydrogen as will be discussed later [Ryu et al., 2008a; Ryu et al., 2008d].

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Figure 8. X-ray diffraction patterns of the product obtained with (a) mixture feeding (WCl6/CoCl2 molar ratio: 1.1 and feeding rate of mixture: 0.053 g/min), and (b) modified arrangement (feeding rate of WCl6: 0.06 g/min and feeding rate of CoCl2: 0.02 g/min) under the following conditions: reaction temperature of 1400°C, CH4 feeding rate of 0.1 L/min (25°C, 86.1 kPa), H2 feeding rate of 0 L/min, and total flow rate of 0.6 L/min (25°C, 86.1 kPa). [Adapted from Sohn et al. (2007)].

Figure 9. TEM photograph of WC-Co nanopowder synthesized under the following conditions: reaction temperature of 1400°C, total flow rate of 1.1 L/min (25°C, 86.1 kPa), CH4 feeding rate of 0.01 L/min (25°C, 86.1 kPa), no H2, WCl6 feeding rate of 0.06 g/min, and CoCl2 feeding rate of 0.02 g/min. [Adapted from Sohn et al. (2007)].

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Since excess methane was used in this work, free carbon was always present in the synthesized composite powder. Thus, in order to evaluate experimental variables on the carbon content in the product, carbon content analysis should be conducted and it was necessary to scale up the precursors feeding rate because the amount of product obtained at the previous feeding rate of precursors, WCl6 of 0.06 g/min and CoCl2 of 0.02 g/min, was too small to perform the carbon content analysis. Therefore, in the subsequent experiments, other factors that also affect the degree of carburization as well as conditions to minimize the amount of the free carbon in the product at an increased precursors feeding rate were explored. The effect of reactants concentration was investigated by varying the feeding rate of WCl6 from 0.06 g/min to 0.6 g/min and that of CoCl2 from 0.02 g/min to 0.2 g/min. Figure 10 shows the XRD patterns of the products obtained with different feeding rates of precursors. The product under these conditions was WC and Co together with small amounts of W2C and/or W depending on the feeding rate of each reactant. No intermediate phases such as Co3W3C or Co6W6C were present in the product. The feeding arrangement in the modified system prevented the unreacted precursor gases from coming into contact with each other and forming intermediate phases, even when the concentration of precursors was increased.

Figure 10. X-ray diffraction patterns of the products obtained with different feeding rates of precursors in the feed stream at (a) 0.06 g/min for WCl6 and 0.02 g/min for CoCl2, (b) 0.3 g/min and 0.1 g/min, and (c) 0.6 g/min and 0.2 g/min, respectively. under the following conditions: reaction temperature of 1400oC, WCl6 volatilizer temperature of 440oC, C/W ratio of 2.3 in the feed stream, no H2, and total flow rate of 1.1 L/min (25°C, 86.1 kPa). [Adapted from Ryu et al. (2008d)].

The mole fractions of tungsten carbides phases are shown in figure 11. The molar ratio was calculated from the XRD peak intensity of WC at 2θ of 35.6, that of W2C at 2θ of 39.5, and that of W at 2θ of 40.2, according to Equation (5). The W2C phase was represented as WC0.5 to better show how much of W is in each carbide phase. As the feeding rate of WCl6 was increased from 0.06 to 0.3 g/min, W2C disappeared and a small amount of W was produced. The amount of produced W increased somewhat as the feeding rate was increased

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to 0.6 g/min. The degree of carburization was slightly decreased with the increase of reactants concentration considering that the carburization occurs from W to Co3W3C and W2C and then WC in the presence of Co. The grain size of WC, 25 ± 1 nm at the lowest feeding rate, increased to 30 ± 1 nm at the highest feeding rate. It is likely that the increase in WCl6 concentration leads to increased grain size of the reduced W, which resulted in a slower carburization rate. The morphology of the WC-Co composite powder is presented in figure 12. EDS mapping of the composite powder confirmed that the WC and Co particles were uniformly mixed in the product, which presents an example of the significant feature of the CVS process to produce very uniformly mixed powders, as mentioned earlier.

Figure 11. Effect of precursor feeding rates on the mole fractions of tungsten carbides with CoCl2/WCl6 molar ratio = 1. [Adapted from Ryu et al. (2008d)].

The effect of CH4 concentration in the feed stream was explored at an increased precursor feeding rate by controlling C/W molar ratio from 2.7 to 2.1. Figure 13 shows the mole fraction of tungsten carbides depending on different C/W molar ratios in the feed stream, indicating that the degree of carburization decreased with the decrease of CH4 concentration in the feed stream. The weight percentage of carbon in the product was converted to percent excess carbon, which is defined as the percentage of the excess carbon over the corresponding stoichiometric amount if all W was present as WC. The % excess carbon in the product decreased from 67.8 % to 3.13 % when the C/W molar ratio decreased from 2.7 to 2.1, as shown in figure 14. The grain size of produced WC, 30 ± 1 nm was not affected by carbon concentration in the feed stream within the range tested.

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Figure 12. SEM image of WC-Co composite powder with associated EDS maps for tungsten and cobalt. (a) WC-Co composite powder, (b) EDS map of W (white), and (c) EDS map of Co (gray) under the synthesis conditions: reaction temperature of 1400oC, WCl6 feeding rate of 0.6 g/min, CoCl2 feeding rate of 0.2 g/min, C/W ratio of 2.3 in the feed stream, no H2, and total flow rate of 1.1 L/min (25°C, 86.1 kPa). [Adapted from Sohn et al. (2007)].

Figure 13. Effect of C/W molar ratio in the feed stream on the mole fractions of tungsten carbides under the following conditions: reaction temperature of 1400oC, WCl6 volatilizer temperature of 440oC, WCl6 feeding rate of 0.6 g/min, CoCl2/WCl6 ratio of 1, no H2, and total flow rate of 1.1 L/min (25°C, 86.1 kPa).

From the results, free carbon was still present in the product even at the lowest C/W molar ratio, which can be controlled in the feed stream. The ratio of H2 produced from CH4 to H2 required to reduce both precursors was unity at C/W molar ratio of 2.1. Thus, other methods were further searched to reduce excess carbon in the product as described in the following sections.

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Figure 14. Effect of C/W molar ratio in the feed stream on % excess carbon.

Figure 15. Effect of H2/CH4 molar ratio in the feed stream on the mole fractions of tungsten carbides under the following conditions: reaction temperature of 1400oC, WCl6 feeding rate of 0.6 g/min, CoCl2/WCl6 ratio of 1, C/W molar ratio of 2.7, and total flow rate of 1.1 L/min (25°C, 86.1 kPa). [Adapted from Ryu et al. (2008d)].

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Hydrogen was added as the reducing gas as well as to reduce the amount of free carbon in the reaction. The effect of hydrogen concentration in the feed stream was investigated on the product composition and carbon content in the product by controlling the H2/CH4 molar ratio from 0 to 4.3. Figure 15 shows the mole fractions of tungsten carbides obtained with different H2/CH4 molar ratios, indicating that the degree of carburization decreased as the H2/CH4 molar ratio was increased. It is likely that hydrogen addition in the feed stream inhibited the decomposition of CH4 into C and H2. Thus, as the feeding rate of H2 increased, the C potential decreased, which resulted in a lower degree of carburization and lower carbon content in the product as will be shown in figure 16. The % excess carbon in the product decreased from 67.8 % to -17.5 % as the H2/CH4 ratio in the feed stream was increased from 0 to 4.3, as shown in figure 16. The grain size of produced WC decreased from 30 ± 1 nm with no H2 to 26 ± 1 nm with H2/CH4 ratio of 4.3.

Figure 16. Effect of H2/CH4 molar ratio in the feed stream on % excess carbon. [Adapted from Ryu et al. (2008d)].

Tungsten carbide powder is used mainly to make bulk components by a consolidation process. Therefore, the presence of incompletely carburized W2C, W, or Co3W3C phases can be tolerated in this work because they can be fully carburized during the subsequent sintering process. These phases can also be fully carburized during the separate post-treatment of the produced composite powder as described in the following section [Ryu et al., 2008d]. Free carbon was always present in the produced powders since excess methane was used in the reaction. Thus, the produced powders were subjected to the post-treatment to remove free carbon as well as to fully carburize unreacted tungsten carbide phases to the WC phase. The produced composite powder was placed in a ceramic boat, which in turn was placed in a tube furnace under hydrogen. The effect of hydrogen heat treatment on the product composition, grain growth, and carbon content was investigated. The experimental conditions were as follows: The amount of powder treated was 2 g, the treatment temperature was

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900oC, which was low enough to prevent rapid grain growth of particles, and hydrogen was flowed at a rate of 0.5 L/min (25ºC, 86.1 kPa) after the temperature reached 900ºC. The result showed that the unreacted W2C and W phases were carburized to the fully carburized WC phase after the powder was treated by hydrogen, as shown in figure 17. It is likely that the carburization of W2C and W phases to the fully carburized WC phase is caused by highly mobile hydrocarbons formed during the hydrogen heat treatment, mainly CH4, which react with W2C and W.

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Figure 17. X-ray diffraction patterns of the products obtained with different hydrogen treatment times at 900ºC using the powder produced from the tubular reactor: (a) before treatment, (b) 1 hour, (c) 2 hours, and (d) 4 hours of holding time. [Adapted from Ryu et al. (2008d)].

The % excess carbon in the product, 112 % before the treatment, was down to 7 % after 1 hour of post-treatment, 3 % after 2 hours, and 0 % after 4 hours, as shown in figure 18. Up to 4 % excess free carbon is not only tolerated but often required in the subsequent consolidation process to remove the small amount of oxygen present on the particle surface. Therefore, a post-treatment of 2 hours under the conditions used in this work is sufficient. The grain size of WC, 30 ± 1 nm before the treatment, increased to 36 ± 1 nm after 4 hours of treatment. The particle size of composite powder before and after the hydrogen heat treatment was examined by TEM photographs, as shown in figures 19(a) and (b). EDS maps of the composite powder obtained after the hydrogen heat treatment are shown in figures 19(c) and (d). The actual particle size, as opposed to the grain size, of the composite powder was less than about 40 nm before the treatment and less than about 70 nm after the treatment for 4 hours in hydrogen based on measurements on a rather limited number of particles. EDS maps also indicate that WC and Co particles were uniformly mixed in the product, which verifies the significant feature of the CVS process investigated by Sohn and coworker [Sohn and PalDey, 1988a – 1988d, Sohn et al., 2004; 2007] to produce very uniformly mixed powders, as mentioned earlier.

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Figure 18. Effect of hydrogen heat treatment time at 900ºC on % exces s carbon. [Adapted from Ryu et al. (2008d)].

Figure 19. TEM photographs with EDS maps of WC-Co composite powders: (a) WC-Co composite powder obtained from the tubular reactor, (b) WC-Co composite powder after hydrogen heat treatment for 4 hours at 900ºC, (c) EDS map of posttreated sample (Tungsten, black area), and (d) EDS map of posttreated sample (Cobalt, black area). [Adapted from Ryu et al. (2008d)].

3. WC-CO IN A PLASMA REACTOR Thermal plasma process provides many advantages on the synthesis of nanosized powder such as a high processing temperature to vaporize all the reactants completely and rapidly, a high quench rate to form fine powders, and the versatility of a wider range of reactants to choose from. Numerous reports [Tong et al., 2005; Moriysohi et al., 1997; Fukumasa et al., 2003; Chazelas et al., 2007; Lee et al., 2007] on the thermal plasma synthesis of metals, ceramics, and composites in recent years have shown that thermal plasma synthesis has given a new direction and impetus to many industrial applications as one of the most promising

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methods for producing nanosized powders. Thus, a plasma reactor has been used to prepare WC and WC-Co composite powders in this laboratory [Sohn et al., 2007; Ryu et al., 2008a; Ryu et al., 2008b; Ryu et al., 2009c; Ryu et al., 2009d]. The plasma system used in their laboratory is equipped with a plasma generator with a downward plasma torch, a power supply unit, a cylindrical reactor, a quenching chamber, a cooling system, a precursor feeding system, a powder collector, a gas delivery system, an off-gas scrubber, and an offgas exhaust system, as shown figure 20. The plasma torch consists of a water-cooled tungsten cathode and a copper anode nozzle operating at atmospheric pressure. The reactor consists of a vertical water-cooled stainless steel tube of 15 cm inner diameter and 60 cm length and an inner graphite cylinder of 7.6 cm inner diameter and 60 cm length. Graphite felt is placed between the graphite tube and the inner wall of the water-cooled stainless-steel tube for the insulation of the reactor. The quenching chamber connected to the bottom of the reactor is a watercooled two-layer stainless-steel box to cool the outgoing gas to a temperature lower than 150oC. A data acquisition system records the temperatures at the reactor exit, the input and output cooling water, and outgoing gas from the quenching chamber. The precursor feeding system for this plasma reactor is the same as that for the tubular reactor except for a watercooled delivery line through which the precursor is fed toward the outside boundary of the visible plasma flame (7 mm diameter) from a distance of 15 mm near the exit of the torch. Ar and H2 gases are used as the primary and secondary plasma gases, respectively. Before delivering precursor into plasma flame, the reactor is heated by the plasma flame generated until its temperature reaches a steady level. The pressure in the reaction chamber was always 86.1 kPa. The metal chlorides and metal oxides are used as the reactants in the plasma reactor system. The mixture of Ar, H2, and CH4 flows through precursor feeding system to carry precursor powder into the plasma flame. The powder produced is collected using a Tefloncoated polyester filter with a pore size of 1 μm. Argon gas is passed through the powder feeder as the carrier gas as well as an inert gas to keep the atmosphere in the sample container inert.

Production of WC from WCl6 The plasma system was first tested to find optimum conditions at which tungsten carbide could be produced from WCl6-H2-CH4 mixture by a thermal plasma process and the detailed experimental conditions are described in the references by Sohn and coworkers [Sohn et al., 2007; Ryu et al., 2008a; Ryu et al., 2008b]. The main product under these conditions was WC1-x with small amounts of W2C and WC phases, as shown in figure 21. This result is consistent with the phase diagram [Sara, 1965], which indicates that WC1-x is likely the first stable solid tungsten carbide phase formed as the W-C liquid solution cools towards solidification at 2710oC. After the initial formation the WC1-x phase may be carburized to WC. Further, this phase decomposes into WC and W2C phases below 2530oC. However, the rapid quenching of the synthesized particles in this system, which provides little time for either decomposition or further carburization, resulted in the WC1-x phase remaining as the major phase in the collected product. Figure 21 also shows the XRD patterns of the products obtained with different CH4/H2 molar ratios in the feed stream. The amount of W2C and WC decreased as the CH4/H2 ratio was increased and when CH4 alone was used as the reaction gas, the product was WC1-x. The grain size of WC1-x, 13 ± 1 nm, decreased to 9 ± 1 nm when no H2 was used.

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Figure 20. Schematic diagram of apparatus: (1) entrained-flow powder feeder for precursors, (2) plasma gun, (3) cylindrical reactor, (4) cooling chamber, (5) powder collector, (6) scrubber, and (7) offgas exhaust system. [Adapted from Sohn et al. (2007)].

Figure 21. X-ray diffraction patterns of the products obtained from different CH4/H2 molar ratios at (a) [CH4/H2] = 0.5, (b) [CH4/H2] = 1, and (c) only CH4 under the following conditions: WCl6 feeding rate of 3.5 g/min, the flow rate of Ar-H2-CH4 mixture to carry the WCl6 powder of 4 L/min (25ºC, 86.1 kPa), applied plasma torch power of 13 kW, and the flow rate of plasma gas (Ar) of 38 L/min (25ºC, 86.1 kPa) with no secondary plasma gas (H2). [Adapted from Ryu et al. (2008a)].

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The results of these experiments indicated that methane alone was enough to reduce and carburize tungsten hexachloride to tungsten carbide, in which the subsequent experiments were performed without the hydrogen addition in the feed stream. The results obtained from the preliminary experiments conducted in the thermal plasma process showed a considerable promise compared with other conventional processes, owing to the high temperature generated by the plasma flame to rapidly volatilize the precursor (WCl6) and the rapid quenching of the product to yield nanosized tungsten carbide powder. Other factors that significantly affect the WC1-x formation were further tested, as described below. The effect of plasma torch power on the product composition and grain size was investigated by varying it from 11 kW to 32 kW. The main product was WC1-x and the amounts of WC and W2C were small in all cases. The grain size of WC1-x obtained at a power level of 11 kW was 8 ± 1 nm and increased to 16 ± 1 nm as the applied power was increased to 19 kW or higher, up to 32 kW. The particle size of WC1-x was also examined using TEM photographs. Figure 22 shows that the particle size of WC1-x obtained at a power level of 11 kW was less than 10 nm and it increased to about 20 nm at an increased plasma torch power of 32 kW. The morphology of produced particles was mostly round. Since excess methane was used in the reaction, free carbon was also present in the product.

a

b

Figure 22. TEM photographs of WC1-x nanopowder synthesized at different power levels of plasma torch of (a) 11 kW and (b) 32 kW under the following conditions: flow rate of plasma gas (Ar) of 57 L/min (25ºC, 86.1 kPa) with no secondary plasma gas (H2), WCl6 feeding rate of 3.5 g/min, C/W ratio of 6.3, and the flow rate of Ar-CH4 mixture to carry the WCl6 powder of 2.5 L/min (25ºC, 86.1 kPa) with no H2. [Adapted from Ryu et al. (2008a)].

The effect of plasma gas flow rate (Ar) on the product composition and grain size was also investigated by varying it from 29 L/min (25ºC, 86.1 kPa) to 75 L/min (25ºC, 86.1 kPa) with no secondary plasma gas (H2). The main product under these conditions was WC1-x with a small amount of WC. Although small in all cases, the amount of WC decreased as the plasma gas flow rate was increased. This reduced the time for WC1-x to be further converted

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to WC. The grain size of WC1-x obtained at the plasma gas flow rate of 29 L/min (25ºC, 86.1 kPa) was 15 ± 1 nm and decreased to 9 ± 1 nm when the plasma gas flow rate was increased to 75 L/min (25ºC, 86.1 kPa). The addition of secondary plasma gas (H2) into the plasma flame plays an important role to increase the plasma flame temperature as well as to enlarge the high temperature region [Choi, 2006]. Thus, the effect of adding hydrogen in the plasma flame on the WC1-x formation and grain size was investigated. The secondary plasma gas (H2) was added at a composition of Ar-2 mol % H2 or Ar-7 mol % H2 in the plasma flame, which resulted in a H2 flow rate of 0.66 L/min (25ºC, 86.1 kPa) or 2.1 L/min (25ºC, 86.1 kPa), respectively. The applied power of the plasma torch was 13 kW when no secondary plasma gas (H2) was added and increased to about 18 kW when the secondary plasma gas (H2) was added at 0.66 L/min (25ºC, 86.1 kPa) or 2.1 L/min (25ºC, 86.1 kPa). The predominant phase obtained under these conditions was WC1-x with a small amount of WC, as shown in figure 23. The grain size of WC1-x was 15 ± 1 nm when no secondary plasma gas (H2) was added and increased to 18 ± 1 nm when the secondary plasma gas (H2) was delivered at 0.66 L/min (25ºC, 86.1 kPa). However, the grain size decreased to 16 ± 1 nm when the secondary plasma gas (H2) was further increased to 2.1 L/min (25ºC, 86.1 kPa). The maximum amount of WC in the product was observed when the secondary plasma gas (H2) was delivered at 0.66 L/min (25ºC, 86.1 kPa). The addition of the secondary plasma gas (H2) into the plasma flame promoted further conversion of WC1-x to WC, which resulted from the increased reaction time provided by the expanded flame length. However, when the excess hydrogen was added, it also could have inhibited the decomposition of CH4 into C and H2, which resulted in a decreased C concentration in the reaction. Free carbon was always present in the product since excess methane was used in the reaction. Thus, the effect of methane concentration in the feed stream was investigated by varying the C/W ratio from 6.3 to 1.5. In the case of the lowest C/W ratio, the ratio of H2 produced from CH4 to H2 required to reduce precursor was unity. The main product obtained under these conditions was WC1-x with a small amount of W2C or WC. Although small in all cases, the amount of WC decreased and the formation of W2C was observed as the methane concentration in the feed stream was decreased. The grain size of WC1-x was 13 ± 1 nm at C/W ratio of 6.3 and it was not affected by decreasing the methane concentration in the feed stream within the range tested. The % excess carbon in the product obtained at a C/W ratio of 6.3 in the feed stream was 219 %. It decreased to 26 % as the C/W ratio in the feed stream was decreased to 1.5. The presence of free carbon was still observed even at the lowest C/W ratio in the feed stream. However, the free carbon also can be removed by post-treatment using hydrogen gas as described in the following section [Ryu et al., 2008a]. The produced powders were subjected to the post-treatment to remove free carbon as well as to fully carburize unreacted tungsten carbide phases, WC1-x, W2C, or W phases to the WC phase. The powder obtained from the plasma reactor was placed in a ceramic boat, which in turn was placed in a tube furnace under H2 atmosphere. The effect of hydrogen heat treatment on the product composition, grain growth, and carbon content was investigated. The amount of powder treated was 10 g, the treatment temperature was 900ºC, which was low enough to prevent rapid grain growth of particles, and hydrogen was flowed at a rate of 0.5 L/min (25ºC, 86.1 kPa) after the temperature reached 900ºC. The result showed that the unreacted WC1-x and W2C phases were fully carburized to WC phase during this posttreatment, as shown in figure 24. The carburization of WC1-x and W2C phases to the fully

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carburized WC phase occurred at this temperature by highly mobile hydrocarbon, mainly CH4 formed during the hydrogen post-treatment, which reacts with WC1-x and W2C particles, as mentioned earlier.

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Figure 23. X-ray diffraction patterns of the products obtained with various flow rates of the secondary plasma gas (H2) of (a) 0, (b) 0.66 L/min (25ºC, 86.1 kPa), and (c) 2.1 L/min (25ºC, 86.1 kPa) under the following conditions: flow rate of the primary plasma gas (Ar) of 29 L/min (25ºC, 86.1 kPa), C/W ratio of 6.3, WCl6 feeding rate of 3.5 g/min, and the flow rate of Ar-CH4 mixture to carry the WCl6 powder of 2.5 L/min (25ºC, 86.1 kPa) with no H2. [Adapted from Ryu et al. (2008a)].

Figure 24. X-ray diffraction patterns of the products obtained with different hydrogen treatment times at 900ºC using the powder produced from the plasma reactor: (a) before treatment, (b) 1 hour, (c) 2 hours, (d) 3 hours, and (e) 5 hours. [Adapted from Ryu et al. (2008a)].

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The excess carbon in the product, 213.5 % before the treatment was completely removed after the powder was treated for 5 hours. The grain size of WC1-x was 13 ± 1 nm before the treatment and WC powder with grain size of 36 ± 1 nm was obtained after the treatment. Figure 25 shows the TEM photographs of WC1-x powder obtained at 13 kW torch power and the WC powder obtained after the hydrogen heat treatment for 5 hours. The particle size of WC1-x powder was less than 20 nm before the treatment and the particle size of WC obtained after the treatment was less than 100 nm from measurements on a limited number of particles. As shown in the TEM photographs, grain growth and agglomeration of particles were observed after the treatment. Because the grain growth occurred during the treatment at this temperature, the powder obtained from the plasma reactor was also treated at 800ºC. Figure 26 shows the XRD patterns of the products obtained from 2 and 5 hours of hydrogen treatment at 800ºC. The main product after the treatment was W2C with a small amount of WC, which already existed before the treatment. It can be seen that WC1-x was decarburized to the W2C phase rather than being carburized to the WC phase, even though excess carbon was present during the hydrogen heat treatment at this temperature. The % excess carbon as defined in the previous section, in the product before the treatment was 213 % and decreased to -39 %, after the powder was treated for 5 hours at 800ºC. From the results, it can be seen that the post-treatment temperature of powders consisting of mainly WC1-x phase obtained from the plasma reactor should be higher than 900ºC to produce the fully carburized WC phase in this work.

a

b

Figure 25. TEM photographs of powders: (a) WC1-x powder obtained from the plasma reactor and (b) WC powder obtained after the hydrogen treatment for 5 hours at 900ºC. [Adapted from Ryu et al. (2008a)].

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Figure 26. X-ray diffraction patterns of the products obtained with different hydrogen heat treatment times at 800ºC using the powder obtained from the plasma reactor: (a) before treatment, (b) 2 hours, and (c) 5 hours. [Adapted from Ryu et al. (2008a)].

Ammonium paratungstate (APT, (NH4)10W12O41·xH2O) and cobalt oxide (Co3O4) were also used as the precursors to synthesize WC1-x-Co composite powder in this plasma system. These materials can be volatilized in the plasma flame owing to the high temperature, and followed by subsequent vapor phase reactions. These materials are also suitable for the industry application rather than metal chlorides. Also, a post-treatment of synthesized powders using these materials to produce WC-Co composite powder was tested, as will be discussed later.

Production of W and WC from APT Ammonium paratungstate as a W precursor was first tested to evaluate whether it could be reduced to tungsten before attempting to produce tungsten carbide [Ryu et al., 2009a]. When APT powder was delivered with only Ar as a carrier gas into the plasma flame, WO3 was produced, as shown in figure 27(a). Tungsten powder was produced by the addition of H2 gas as the reducing agent in the feed stream together with APT powder as received without pre-treatment, as shown in figure 27(b). XRD results showed that the product was tungsten without any detectable oxides. When CH4 was fed as the carburizing agent with H2 in the feed stream together with APT powder, tungsten carbide (WC1-x) was produced and a small amount of W2C phase was mixed in the product, as shown in figure 27(c). Equations given below represent the process of thermal decomposition of APT, and subsequent hydrogen reduction and methane carburization in the thermal plasma process. - Thermal decomposition

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(NH4)10W12O41·xH2O → 12WO3 + 10NH3(g) + (5+x)H2O(g)

171

(7)

- Hydrogen reduction WO3 + 3H2(g) → W + 3H2O(g)

(8)

- Methane carburization

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WO3 + H2(g) + CH4(g) → W2C, WC1-x, WC + H2O(g)

(9)

Figure 27. X-ray diffraction patterns of products obtained by thermal plasma process of APT ((NH4)10W12O41·xH2O) with different reactant gases of (a) Ar, (b) Ar and H2, and (c) H2 and CH4 under the following conditions: APT feeding rate of 3.5 g/min, the flow rate of carrier gas to feed the APT powder into the plasma flame of 6 L/min (25ºC, 86.1 kPa), the flow rate of plasma gas (Ar) of 48 L/min (25ºC, 86.1 kPa) with no secondary plasma gas (H2), and applied power of 13 kW. [Adapted from Ryu et al. (2009c)].

Both XRD and EDS analyses, as shown in figures 27 (b) and 28 showed only the W peaks. The produced W powder had a grain size of less than 25 nm determined by applying the Scherrer equation to the XRD data. The particle size of W obtained at this condition was also examined using ZetaPALS and TEM. ZetaPALS gave an average particle size of 88.5 nm, as shown in figure 29 but TEM examination showed that the actual particle size of W was less than 45 nm, as shown in figure 30.

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a

b

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Figure 28. SEM micrograph and EDS result of the W nanopowder obtained at the plasma torch power of 13 kW and H2 feeding rate of 2.9 times the stoichiometric value: (a) SEM micrograph and (b) EDS micrograph. [Adapted from Ryu et al. (2009a)].

Figure 29. Particle size distribution in log-normal scale from ZetaPALS measurement of W nanopowder obtained at the plasma torch power of 13 kW and H2 feeding rate of 2.9 times the stoichiometric value. [Adapted from Ryu et al. (2009a)].

The difference in the particle sizes obtained by the two methods can be explained by the agglomeration of the particles as well as their motion in liquid media that contributed to errors in particle size measurement by ZetaPALS. The surface area of this W powder was measured using BET to be 2.8 ± 0.1 m2/g. The average size of particles can be obtained from the surface area by following equation under the assumption of spherical particles:

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D

6 A

173

(10)

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where, D is the average diameter of W particles produced, A is the surface area of the powder per gram (m2/g), and ρ is the density of W. The average particle size of W thus determined was about 111 nm. It is larger than what was observed from the TEM photograph, which yielded an average W particle size of about 28 nm with the size ranging from 13 to 45 nm based on the measurement on a rather limited number of particles. This difference also could be explained by the agglomeration of particles. The effect of plasma torch power on the tungsten carbide synthesis was investigated by varying it from 7 kW to 29 kW. Figure 31 shows the XRD patterns of the products obtained with different power levels of the plasma torch. The product obtained under these conditions was WC1-x with a small amount of W2C phase. The product composition was not affected by plasma torch power and the grain size of WC1-x increased as the plasma torch power was increased within the range tested. The grain size of WC1-x obtained at 7 kW of plasma torch power was 5 ± 1 nm and it increased to 13 ± 1 nm as the plasma torch power was increased to 29 kW. The particle sizes of WC1-x were also examined using TEM photographs. Figure 32 shows that the particle size of WC1-x obtained at the power level of 7 kW was less than 10 nm and it increased but mostly less than 20 nm at the increased plasma torch power of 29 kW.

Figure 30. TEM photograph of W nanopowder synthesized at the plasma torch power of 13 kW and H2 feeding rate of 2.9 times the stoichiometric value. [Adapted from Ryu et al. (2009a)].

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Figure 31. X-ray diffraction patterns of the products obtained with different power levels of plasma torch at (a) 7 kW, (b) 10 kW, (c) 16 kW, and (d) 29 kW under the following conditions: flow rate of plasma gas of 38 L/min (25ºC, 86.1 kPa) without secondary plasma gas (H2) addition, APT feeding rate of 3.1 g/min, C/W ratio of 5.7 in the feed stream, and the flow rate of CH4-H2 mixture to carry APT powder of 4 L/min (25ºC, 86.1 kPa) with no Ar was used as a carrier gas. [Adapted from Ryu et al. (2009c)].

a

b

Figure 32. TEM photographs of WC1-x nanopowders synthesized at different power levels of plasma torch: (a) 7 kW and (b) 29 kW. [Adapted from Ryu et al. (2009c)].

Production of WC-Co from APT and Cobalt Oxide The plasma system used for the production of W and WC was applied to synthesize tungsten carbide and cobalt composite powder using premixed ammonium paratungstate (APT) and cobalt oxide [Ryu et al., 2009d]. Tungsten and cobalt composite powder was first synthesized from APT-Co3O4-H2 mixture as shown in figure 33(a). In subsequent experiment, CH4 was added with pre-mixed APT and Co3O4 powder in the feed stream to synthesize

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tungsten carbide-cobalt composite powder and the product was a mixture of WC1-x and Co, which contained small amounts of W2C and W phases, as shown in figure 33(b). The effect of gas composition was also tested by varying the CH4/H2 ratio in the feed stream.

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Figure 33. X-ray diffraction patterns of the products obtained by thermal plasma process from (a) APTCo3O4-H2 mixture and (b) APT-Co3O4-CH4 mixture under the following conditions: Plasma torch power of 9 kW, the flow rate of plasma gas (Ar) of 28 L/min, feeding rate of pre-mixed APT-Co3O4 with a Co/W ratio at 0.7 of 3.5 g/min, and the flow rate of H2 or CH4 to carry APT-Co3O4 mixture of 4 L/min (25oC, 86.1kPa). [Adapted from Ryu et al. (2009d)].

However, adding H2 in the feed stream resulted in the formation of substantial amounts of W2C and W phases in the product. When CH4 alone was used as a reaction gas, the highest degree of carburization was obtained. The grain size of WC1-x calculated from XRD patterns by applying the Scherrer equation was 9  1 nm. EDS mapping of the composite powder was also conducted and it confirmed that the WC1-x and Co particles were uniformly mixed in the product, as shown in figure 34. This presents an example of the significant feature of the vapor phase reaction to produce very uniformly mixed powders, as mentioned earlier.

a

b

c

Figure 34. SEM image of WC1-x-Co composite powder with associated EDS maps for tungsten and cobalt: (a) WC1-x-Co composite powder, (b) EDS map of W (light area), and (c) EDS map of Co (gray area). [Adapted from Ryu et al. (2009d)].

By controlling the CH4/H2 ratio in the feed stream the amount of free carbon could be reduced. However, it also results in the formation of substantial amounts of W2C and W phases. The composite powder obtained from the plasma reactor was subjected to the postNanopowders and Nanocoatings: Production, Properties and Applications : Production, Properties and Applications, Nova Science Publishers,

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treatment under hydrogen as discussed earlier. The amount of powder treated was 5 g, the treatment temperature was 900ºC to prevent rapid grain growth of particles, and hydrogen was flowed at a rate of 0.5 L/min (25ºC, 86.1 kPa) after the temperature reached 900ºC. The unreacted WC1-x and W2C phases were carburized to the fully carburized WC phase after the powder was treated by hydrogen for 2 hours, as shown in figure 35. The % excess carbon in the product, 138 %, before the treatment, was down to 37 % after 1 hour of post-treatment, and 0 % after 2 hours. The grain size of WC in the composite powder obtained after the treatment for 2 hours was 40  1 nm.

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Figure 35. X-ray diffraction patterns of the products obtained with different hydrogen treatment times at 900oC using composite powder synthesized from a plasma reactor (a) before treatment, (b) 1 hour and (c) 2 hours.

CONCLUDING REMARKS The CVS process has shown to be capable of producing nanosized powders of 5-200 nm sizes. Recently, this method was applied to the preparation of nanosized Al, W, WC, and WC-Co composite powders in a tubular reactor system and a plasma reactor system by Sohn and coworkers [Sohn and PalDey, 1988a – 1988d, Sohn et al., 2004; 2007; Ryu et al., 2008a – 2008d; Ryu et al., 2009a-2009d]. From their research work, it was also confirmed that very uniformly mixed nanopowder of different solid phases less than 100 nm in size can be obtained by this method. The plasma reactor system is suitable for the industrial application because it is easy to scale up. Further work is continuing in this laboratory to optimize the experimental conditions for the synthesis of various nanosized powders and to apply the CVS process to other nanopowders.

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REFERENCES Barksdale, J. Titanium Its Occurrence, Chemistry and Technology; Ronald Press Co.: New York, NY, 1966, pp 213-223, pp 400-79. Chang, W.; Skandan, G.; Hahn, H.; Danforth, S. C.; Kear, B. K. Nanostruct. Mater. 1994, 4, 345-351. Chazelas, C.; Coudert, J. F.; Jarrige, J.; Fauchais, P. J. Eur. Ceram. Soc. 2007, 27, 947-950. Choi, J. W.; Sohn, H. Y.; Choi, Y. J.; Lu, J.; Han, G.; Fang, Z. Z. 8th Global Innovations Symposium: Metal Powders for Energy Production and Storage Applications, ed. by Fang, Z. Z. and Sears, J. pp. 1-8, Collected Proceedings: Emerging Materials, TMS (The Minerals, Metals & Materials Society), Warrendale, PA, 2007. Choi, S. I.; Nam, J. S.; Lee, C. M.; Choi, S. S.; Kim, J. I.; Park, J. M.; Hong, S. Curr. Appl. Phys. 2006, 6, 224-229. Cullity, B. D. Elements of X-ray diffraction; Addison-Wesley Pub. Co., Reading, Mass, 1978, 99-106, 415-417. Fichtner, M.; Fuhr, O. J. Alloys Compd. 2002, 345, 286-296. Fukumasa, O.; Fujiwara, T. Thin Solid Films. 2003, 435, 33-38. Graetz, J.; Reilly, J.; Sandrock, G.; Johnson, J.; Zhou, W. M.; Wegrzyn, J. Advanced Materials for Energy Conversion III, TMS. (2006), 57-63. Hojo, J.; Oku, T.; Kato, A. J. Less Common Metals. 1978, 59, 85-95. Í ñiguez, J.; Yildirim, T. Appl. Phys. Lett. 2005, 86, 103-109. Lamprey, H.; Ripley, R. L. J. Electrochem. Soc. 1962, 109, 713-715. Lee, S. H.; Oh, S. -M.; Park, D. -W. Mater Sci. Eng. C. 2007, 27, 1286-1290. McCandlish, L. E.; Kear, B. H.; Kim, B. K. Nanostruct. Mater. 1992, 1, 119-124. Moriysohi, Y.; Futaki, M.; Komatsu, S.; Ishigaki, T. J. Mater Sci. Lett. 1997, 16, 347-349. Nikklason, G. A. J. Appl. Phys. 1987, 62, 258-265. Okamoto, Y.; Koyano, T.; Takasaki, A. Japan J. Appl. Phys. 1987, 26, 1943-1945. Otsuka, K.-I.; Yamamoto, H.; Yoshizawa, A. J. Chem. Soc. Japan. 1984, 6, 869-878. Parr, H.; Feder, J. Phys. Rev. 1973, 7, 166-181. Puszynski, J. A. Powder Materials: Current Research and Industrial Practices. 2001, 89-105. Ryu, T.; Sohn, H. Y.; Hwang, K. S.; Fang, Z. Z. J. Mater. Sci., 2008a, 43, 5185–5192. Ryu, T.; Sohn, H. Y.; Hwang, K. S.; Fang, Z. Z. High Temp. Materials and Processes, 2008b, 27, 91-96. Ryu, T.; Sohn, H. Y.; Han, G.; Kim, Y.-U.; Hwang, K. S.; Mena, M.; Fang, Z. Z. Metall. Mater. Trans. B. 2008c, 39B, 1-6. Ryu, T.; Sohn, H. Y.; Hwang, K. S.; Fang, Z. Z. Ind. Eng. Chem. Research, 2008d, 47, 93849388. Ryu, T.; Sohn, H. Y.; Hwang, K. S.; Fang, Z. Z. Int. J. Refractory Metals and Hard Materials. 2009a, 27, 149-154. Ryu, T.; Hwang, K. S.; Choi, Y. J.; Sohn, H. Y. Int. J. Refractory Metals and Hard Materials, 2009b, 27, 701-704. Ryu, T.; Sohn, H. Y.; Hwang, K. S.; Fang, Z. Z. J. Amer. Ceram. Soc., 2009c, 92, 655-660. Ryu, T.; Sohn, H. Y.; Hwang, K. S.; Fang, Z. Z. J. Alloys Compd., 2009d, 481, 274-277. Saeki, Y.; Zaki, R. M.; Nishikara, H.; Ayoama, N. Denki Kagaku. 1978, 46, 613-617. Sara, R. V. J. Am. Ceram. Soc. 1965, 48, 253.

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Seayad, A. M.; Atonelli, D. M. Adv. Mater. 2004, 16, 765-777. Sohn, H. Y.; PalDey, S. Mater Sci. Eng. A. 1998a, 247, 165-172. Sohn, H. Y.; PalDey, S. J. Mater Res. 1998b, 13, 3060-3069. Sohn, H. Y.; PalDey, S. Metall Mater Trans. B. 1998c, 29B, 457-464. Sohn, H. Y.; PalDey, S. Metall Mater Trans. B. 1998d, 29B, 465-469. Sohn, H. Y.; Zhang, Z.; Deevi, S.; PalDey, S. High Temp. Mater Processes. 2004, 23, 329333. Sohn, H. Y.; Ryu, T.; Choi, J. W.; Hwang, K. S.; Han, G.; Choi, Y. J.; Fang, Z. Z. JOM. 2007, 59, 44-49. Tong, L.; Reddy, R. G. Scripta Mater. 2005, 52, 1253-1258. Upadhaya, G. S. Cemented Tungsten Carbide; Noyes Publications: New Jersey, NY: 2002, 1. Wahlberg, S.; Grenthe, I.; Muhammed, M. Nanostruct. Mater. 1997, 9, 105-108. Zaluski, L.; Zaluska, A.; Ström-Olsen, J. O. J. Alloys Compd. 1999, 290, 71-78. Zhao, G. Y.; Revenkar. V. V. S.; Hlavacek, V. J. Less Common. Metals. 1990, 163, 269-280. Zhu, Y. T.; Manthiram, A. Composites Part B. 1996, 27, 407-413. Züttel, A.; Wenger, P.; Rentsch, S.; Sudan, P.; Mauron, P.; Emmenegger, C. J. Power Sources. 2003, 118, 1-7.

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In: Nanopowders and Nanocoatings Editor: V. F. Cotler

ISBN: 978-1-60741-940-2 © 2010 Nova Science Publishers, Inc.

Chapter 6

NANO- AND MICROSTRUCTURAL SILICON POWDERS IN THE SYNTHESIS AND STORAGE OF HYDROGEN А. А. Kovalevskii, А. S. Strogova, V. А. Labunov and А. А. Shevchenok Institute of Applied Mathematics and Mechanics, Donetsk, Ukraine

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ABSTRACT Complex use of products of interaction of MNS SP with water, utilization of heat and effective functioning of power cycle on the basis of nanostructural powder of silicon is the real technology of gaseous hydrogen for hydrogen energy of the nearest future. Application of MNS SP gives undeniable advantages: simple and without expenses of energy from outside technology of hydrogen production in the result of water decomposition: the necessity in storage and transportation of gasiform hydrogen disappears, which promotes fire- and explosionproof of this fragment of hydrogen energy substantially. In the process of the industrial production of nanopowders of silicon, including for other areas of their application their prime price will be reduced some times of their value. The application of MNS SP in the mobile small sources of hydrogen is already expedient.

INTODUCTION One of the most essential problems of modern society is a search of alternative energy sources. The most perspective secondary type of fuel is hydrogen – ecologically clean power carrier, besides practically inexhaustible. In connection with the catastrophic worsening of the ecological state of the planet and exhaustion of resources of hydrocarbon raw materials harmless fuel is most preferable. In the modern production of hydrogen the electrolysis of water and its varieties is mostly used [1-7]. Such process requires heavy expenditures of energy at the production of hydrogen. Many laboratories in the world are trying to solve the problem of diminishing of expenses of energy on the production of hydrogen from water, but

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no substantial results are obtained yet. Some directions in the solution of the problem of its production in their own way were set in the modern production of hydrogen [2-4]. The reaction of water with micro- and nanostructured silicon powders (MNS SP) is of a particular interest [8-11]. A great interest to materials of nano- and microstructures is steadily noticeable for some last years already due to remarkable properties typical for the micro- and nanostructural state [12-16]. The size effects, determining the high level of properties, are mostly developed in the interval approximately up to 100 nm and are determined, at least, by the followings four circumstances:   

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

With the decreasing of particles sizes the role of surfaces division increases substantially; Characteristics of surfaces and their composition in a nanometric interval may be different from the same of ordinary coarse-grained and large-dimension objects; The size of particles in the process of their decreasing may become comparable with characteristic sizes for some physical phenomena (for example, with the length of free run of transmitters in the processes of transfer); The size of crystallites (crystal grain) appears to be comparable with the de-Broyle wavelength – quantum effects occur.

All above-mentioned influences the possible unmonotonous motion of properties change in a nanometric interval and occurrence on dependences of properties such as the size of particles of the special points, predicting of the presence of which is not always possible. Application of nanomaterials broadens, although the problem of realization of optimum nanostructures from the point of providing of optimum properties and their stability is in the stage of solution. In this connection it appears to be interesting to consider the peculiarities of synthesis and storage of hydrogen with the use of MNS SP. Last years micro- and nanostructural powders of certain types of materials attract special attention of different groups of researchers [1-3], which is related with the perspective of their use in the synthesis and storage of hydrogen [8-11, 21-22]. Also as an adsorbent, on the surface of which matters, lowering its surface-tension in relation to an environment are adsorbed. Taking into account the fact that silicon resources in the earth's crust take the second place, and in the production of semiconductor devices and integrated circuits has accumulated and is constantly restocking a large amount of silicon wastes, utilization of which is a serious problem we‘ll try to use it for the synthesis and accumulation of hydrogen. Production and storage of hydrogen is quite a standard task for modern industry, in this field it is carried out safely and is provided with necessary service. In addition it is possible to accumulate, large volumes of hydrogen in cylinder vessels and underground storehouses. However as for the vehicles a substantial breach in the technology of side board storage of hydrogen in order to attain a run between re-fillings, comparable with the cars of petrols or diesels is needed. In principle new developments and ideas in this direction are already being conducted [17-22], and new transport systems are passing the demonstration stage. The latest methods, based on sorption of hydrogen in the hydrides of metals and in chemical hydrides, and also on carbon, require further research and comparative analysis. The

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hydrides of inter metallic connections are used, first of all, as systems, accumulating hydrogen. Mainly, formation of hydrides of such metals as magnesium, zirconium, titan takes place at the chemical reactions of these metals with hydrogen at the elevated pressure. They are expensive, and their resources are exhaustible. More effective method is hydrogenation of the so-called mechanical alloys, i.e. materials, got by the mechanical activating of mixture of metals and other materials. In this work the features of production and use of MNS SP are in the synthesis and storage of hydrogen are examinated.

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METHODS OF EXPERIMENT Nanostructural powders of silicon with the size of particles (crystalline grains) from 20 to 70 nm were produced in the result of decomposition of monosilane (SiH4) in HF–plasma in the volume of chamber at pressure 20-40 Pa in the modernized plasma unit «Plasma-600». Microstructural powders of silicon with the size of particles from 70 to 6000 nm were produced in the result of grinding of wastes of single-crystalline silicon in the vortex stream acoustic mill of VIM-8. Particles sizes and their quantitative distributing by size were defined by pictures from the scanning electronic microscope of S-4800 (Hitachi, Japan). Also middle sizes of nanoparticles were determined by means of diffraction of x-rays on the crystalline kernel of particle. Saturation of silicon powder, both by water and by hydrogen was carried out in the carbidized cuvettes of graphites (MPG-8) with the volume of separate cell of 1 cm3. Saturation by hydrogen was carried out in the unit of epitaxial increase UNES-2П-КА at a temperature from 297 to 723К. Hydrogen used with the dew point of 198-208К was used. Quantitative content of hydrogen was estimated by the difference of weight of cuvette before, and after the saturation. Weighing was conducted with the help of the analytical scales SETRAL-200S with 10-6 accuracy of measurement.

RESULTS AND DISCUSSION With the purpose of investigation of the process of water decomposition in micro- and nanostructural powders at the first stage of investigation the study of absorption of water by silicon powders, I.e. their saturation by water was conducted. In this case we investigated the influence of pressing (in a vacuum, in the air, with the use of polyvinyl alcohol and without it, with different dispersion of powder) at a temperature of 250 - 1200оС for the purpose of saturation of powders by water before and after their treatment in HNO3:HF=1:4 and KOH, NaOH (Figure 1, 2). It is determined that pressing of microstructural powders in the air and saturating them by polyvinyl alcohol cause suppression of their absorption ability. After such pressing at the temperature of 250 °С microstructural silicon powders with dispersion from 70 to 6000 nm before their treatment in acids and hydroxides absorb not more than 18% of water, but after their treatment in the mixture of nitric and hydrofluoric acids, chosen in correlation 1:4 to 24%, and in KOH to 40%.

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, %

vacuum with PVA vacuum without PVA on air with PVA on air without PVA

32 30 28 26 24 22 20 18 16 14 12 10 8 6 4 200

400

600

800

1000

1200

vacuum with PVA vacuum without PVA on air with PVA on air without PVA

36 34 32 30 28 26 24 22 20 18 16 14 12 10 8

0

ТТО, С

200

400

600

800

1000

1200

0

ТТО, С

Figure 1. Conformities to the law of change of absorption ability of MNS SP with the temperature of heat treatment at the different terms of pressing (a, b) and those passed the additional treatment in HNO3-HF (b).

, %

, %

HNO3:HF (1:4)

50

HNO3:HF:CH3COOH (2:7:2)

45

NaOH:H2O (1:3)

40 35 30 25 20

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15 10

20

30

40

50

60

70 R, nm

40 38 36 34 32 30 28 26 24 22 20 18 16 14 12 10

HNO3:HF (1:4) HNO3:HF:CH3COOH (2:7:2) NaOH:H2O=1:3

0

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Figure 2. Conformities to the law of change of the absorbing ability of MNS SP with the size of the nano- (а) and microstructural (b) particles.

Pressed MNS SP at the same temperature in a vacuum without polyvinyl alcohol absorb already up to 30% of water. However the increase of temperature of heat treatment in the process of pressing brings to the sharp falling of their sorption ability and at the temperature of pressing of 1200°С it does not exceed 5% and 10% accordingly. It is determined by affecting of oxygen and carbon of polyvinyl alcohol on the surface properties of MNS SP. Air, as the source of oxygen, brings to acidifying of MNS SP. Dioxide of silicon which influences the etching process of etching of MNS SP in acid etchings appears on surface, as forming dioxide of silicon comes forward as masking coverage from the influence of acids and hydrooxides. In case of polyvinyl alcohol decay carbon is being formed which is together with MNS SP able to form silicon carbide also affecting the process of etching. Acting as a mask silicon dioxide as well as silicon carbide restrain the etching of separate MNS SP crystallites thus in a small degree influencing the change of their micro- and nanostructures in the result of chemical treatment. Additional injection into the solution composition of strong acids of glacial acetic acid in proportions of ingredients

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HNO3:HF:CH3COOH as 2:7:2 promotes increasing of MNS SP sorption capacity at the expense of raising of the efficiency of separate crystallites (grains) etching up to pure silicon. It should be noted that the most sorption ability of powder of silicon is observed on samples, pressed in a vacuum with the most minimal content of polyvinyl alcohol. Opposite, MNS SP, pressed in the air, differ by considerably lesser sorption ability. As after baking in the air in MNS SP basic phases are silicon and dioxide of silicon (SixOy), they strongly influence the sorption ability of powders on the whole. It is known from literature [25, 26, 28], that dioxide of silicon exists at atmospheric pressure in seven modifications: α- quartz, β-quartz, γ-tridimite, β-tridimite, α-tridimite, βcrystobalite, α- crystobalite and two amorphous modifications: quartz glass or loshatelerite and gel of SiO2. All forms of SiO2 are related to the latticed structures, although formally the structure of SiO2 is identical to the structure of ortosilicate silicon of Si[SiO4]. The grate of SiO2 is more or less closed. It is formed by the infinitely proceeding rings of tetrahedrons of SiO4. Rings in the different modifiers of SiO2 differ by the number of the linked tetrahedrons and form. Two forms of rings are distinguished. One of them is characterized by a presence of only sixnominal rings, and the other simultaneously with sixnominal and by eightnominal rings. Rings form spirals with a step equal to the vector of Byurgers. Polymorphic forms also differ by the value of the angle of connection of Si-O-Si and by the length of connection of Si-O. In extreme terms SiO2 has some tens of modifications. Including: fibred silica W, silica O, stishovite, coesite, kinite and others. Such structural imperfection influences the hydrophobicity (hydrophilic behavior) of nano- and microstructural powder. One of possible methods of pressed micropowders upgrading, along with the high-power loading, is the use of powder-like mixtures with the polimodal distribution of particles by size. In this case large pores between larger matrix particles in the volume of the pressed material should be filled with the nanostructural particles of SP that multiplies a contact quality and consequently, flowing speed and intensity of diffusive processes at heat treatment, promoting the growth of density of baked MNS SP and improvement of their physical and chemical descriptions [27]. Real powder-like mixtures and pressed powders from them have substantial heterogeneity form and in size of particles, and their structure frequently cardinally differs from a model one. As it turned out, in the process of numerous researches maximum saturation both by water and by water with ethyl spirit is observed directly for powders, synthesized in the process of decomposition of monosilane and passing treatment in mixture of nitric, hydrofluoric and vinegar acids, chosen in correlation 2:7:2 or in KOH and NaOH. It achieves 45-50% from the weight of initial mass. Microstructural powders produced by the grinding of single-crystalline silicon, take up to 34 masses % of water i.e. a little less than micropowders, got as the result of monosilane decomposition (Figure 2). Investigations on the estimation of influence of temperature of heat treatment on the saturation of MNS SP of different dispersion showed how in the conditions of pressing temperature influences the saturation of the pressed powder of silicon water (Figure1). Increase of temperature of heat treatment from 250°С to 1200°С at pressing of nanostructures powders of silicon, got in the process of decomposition of monosilane, results in the decrease of absorbing ability of the pressed samples from 18% to 5%. In its turn increase of temperature of heat treatment from 250°С to 1200°С at pressing of MNS SP, created as a result of grinding of single-crystalline silicon wastes, results in the decrease of absorbing ability of the pressed samples from 24% to 10%. Treatment of both

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those and other samples of pressed powder of silicon in acids and hydroxide enables in the area of low (250 - 500°С) temperatures of pressing providing sorption ability at the level of 24 - 34%. The increase of temperature of heat treatment in the process of pressing from 500 to 1200°С results accordingly in the decrease of sorption ability of the pressed material to 9 10%. It is possible to explain such conduct of the pressed powders with changing of temperature of pressing by the conformities to the law of change of size of crystalline particles with changing temperature. Thus, pressing environment, temperature of heat treatment in its process, presence of penetrating agent and dispersion of MNS SP and their subsequent treatment in acids and hydrooxide are fundamental factors, influencing the conformity to the law of their saturation by water. Taking into account the results of the research obtained we undertook an attempt to use them in the process of the synthesis of hydrogen in the result of decomposition of water to MNS SP. It is found that the features of the thermal mode of process of co-operation of nanostructural silicon powder with water result in appearance of new effects which were not known for reaction with participation of large particles of silicon. Above all things it is the self-heating of nanoparticles to temperatures, exceeding the temperature of water by hundreds of degrees. Speed of hydrogen emission at co-operation of water with MNS SP at a room temperature makes 0,1-0,3 dm3·с-1. The next advantage of the use of nanostructural powders of silicon in this reaction is that the degree of transformation of silicon makes up 95.100% depending on the degree (volume) of moistening. Moreover, introduction to the distilled water of small quantities of potassium or sodium hydrooxides results in considerable growth of speed of reaction. Findings testify in behalf of MNS SP from compact material at their interaction with water. At their use they interact with water at high speed and the degree of transformation is ~100% and exactly their application will allow getting hydrogen with sufficient speed at ordinary terms. The reaction of silicon with water is exothermic. In this case during interaction warmth is evolved, and there is no need of energy supply from outside. The process of chemical decomposition of water on the MNS SP in the presence of hydrooxides of potassium and sodium is represented by the following reactions: Si + 2 KOH + H2O = K2 SiO3 + 2H2 (газ), Si + 2 NaOH + H2O = Na2 SiO3 + 2H2 (газ). Financial balance of reaction shows, for example, that from 28 gr of silicon, 112 gr of potassium hydrooxides and 18 gr of water there appears 44.8 dm3 (liters) of hydrogen. So, if we to take 1 dm3 of water and silicon and potassium or sodium hydrooxide in the proper proportions, we‘ll get 2450 dm3 of hydrogen. Such amount of appearing hydrogen, undoubtedly, is of interest in case of use of small-size and middle- size stationary installations for its synthesis. Important feature of the process of chemical synthesis is the speed of formation of hydrogen. In the basis of reaction as it was already specified interaction of three components lies. Thus the reaction is exothermic and it flows with excluding of thermal energy in an amount of 196 kilojoules/mole and does not require the preliminary warming up of components of reaction. The reaction flows at a room temperature. Its speed depends on dispersion of powder of silicon (Figure 3), it is evidently, that at the use of both

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nanostructural and microstructural powders of silicon with multiplying the size of particles there is an increase of speed of formation of hydrogen. It is very important to mark that for nanostructural powders there is quite a different character of speed change of formation of hydrogen as compared to the microstructural powders. Multiplying of the size of nanostructural particles of silicon from 30 to 70 nm results in the growth of speed of formation of hydrogen from 0.1 to 0.3 dm3·с-1, while passing to the microlevel of 100-5000 nm results in insignificant growth of speed of formation of hydrogen from 0.30 to 0.40 dm3·с-1, moreover, with an output on a plateau at the size of particles of 3000 nm. In the area of nanosize of particles of powder of silicon r=30–70 nm the speed of formation of hydrogen is related to the size of particles of a nonlinear function, described by a concave curve (Figure 3a), and in case of microstructural particles of powder of silicon with a r=100–5000nm the speed change of forming is described by the mirror inverted nonlinear function with the insignificant change of the output parameter. 3 -1 3 -1 , dm  с , dm  с 0,30

а)

b)

0,40 0,38

0,25

0,36 0,20

0,34 0,15

0,32

0,10 30

0,30 40

50

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r, nm

60

70

0

1000

2000

3000

4000

5000

r, nm

Figure 3. Dependence of speed of hydrogen formation ν from dispersion of powders of silicon of r: a – nanostructural powders; b - microstructural powders. ν, dm3·с-1; r, nm.

It is clear the size of particles of nano- and microstructural particles regulates the supply of water and hydrooxide of potassium (sodium) in the volume of powder and accordingly abstraction of hydrogen from the volume on the whole. It is possible to assume that there is a considerable increase of density of micropowder at the use of nanostructural particles, more dense mass of fragments-conglomerates, which retentive the intensive penetration through them of basic components of reaction in the volume of layer of powder is created as a result. Speed of forming and abstraction of hydrogen go down in this case. Diminishing of size of nano- and microparticles results in the considerable diminishing of the volume of the unit of bulked weight of powder. The eventual state of powder is determined by the middle size of cavities between nanostructural particles by which the components of reaction in the volume of powder-like material penetrate. Multiplying of the size of particles forming the layer of nano- and microstructural powder, results in creation of less number of nano- microstructural particles and diminishing of the area of interphase boundaries with surplus energy, which are

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instrumental in penetration of liquid into the volume layer and in the improvement of penetrating ability for the components of reaction. Sorption ability of powders increases as a result, consequently, the process of forming and abstraction of hydrogen is intensified. Micrometer size of particles of powder of silicon ≥ 3000 nm result in stabilizing of the process of formation of hydrogen. Influence of nano-, microstructural particles with size more than 3000 nm is unimportant. In this case it is possible to come to the conclusion that for the small nanostructural particles of 100 nm it is a deposit of volume. The most densely organized nanostructural silicon powders are produced from the smallest and nanoispersion particles. Thus, for multiplying the duration of the process of formation of hydrogen in the system silicon powder-water-hydroxide it is more expedient to use nanostructural ≤70 nm silicon powders, which allow us to manage the reaction of hydrogen synthesis by a chemical method, more efficiently and for its decline nanostructural particles should be held in the weighed state. In Figure 4 pictures are resulted, where hydrogen torches are shown at different speed (stream) of appearing hydrogen: length of laminar torch is multiplied with growth of speed of formation of hydrogen. Apparently, intensity of processes of burning of hydrogen, appearing in the process of chemical reaction of interaction of MNS SP with water and hydroxide of potassium, chosen in correlation 3.0 gr of powder of silicon, 11.5 gr of hydroxide of potassium and 2.0 gr of water depending on the speed of formation of hydrogen (size of particles) differ in principle.

Figure 4. Dependence of length of the laminar torch of burning of hydrogen on the speed of its forming: a – ν = 0.1 dm3·с-1; b – ν = 0.2 dm3·с-1; c– ν =0.26 dm3·с-1; d – ν = 0.30 dm3·с-1; i – ν = 0.35 dm3·с-1; f – ν = 0.40 dm3·с-1.

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A small peak-like torch (Figure 4, a) appears at the speed of formation of hydrogen of 0.15 dm3·с-1 (size of particles to 50 nm), duration of its burning by 240 sec with, the volume of the excluding hydrogen of 5 dm3. Multiplying of speed of formation of hydrogen up to 0.3 dm3·с-1 (nanostructure particles ≤70 nm) results in multiplying of a peak-like torch and in reduction of duration of its burning up to 180 sec. At the speed of formation of hydrogen a 0.3 dm3·с-1 (transitional size of particles from nano to micro 70–100 nm) the torch takes the shape of a long sharp pick. Burning of hydrogen goes on quickly enough: from the beginning of inflammation to the completion of burning the process flows in within 120 sec and is accompanied by an intensive sound. In the peak of its burning the flame is torn away from a tube and achieves its maximum size in intensity, and then begins to diminish up to its disappearance. It confirms that the initial products of reaction are used up and the reaction stopped. At the speed of formation of hydrogen of 0.37-0.40 dm3·с-1 (size of particles of 100-5000 nm) the torch of burning of hydrogen changes its form: it takes the thickened shape of either cylinder or a bent pick (Figure of 4i,f). The process of burning flows within 60-90 seconds. Luminescence is bright, as well as in all previous cases. It allows to assume that ''clean'' hydrogen appears in all cases. Premature firing of the appearing hydrogen immediately results in an explosion with a strong shock wave and a loud sound. In this case the flame is sucked in to the internal volume of a conclusion tube, and the torch becomes ''running''. If the tube is thrown out from a head vessel the force of the shock wave, as well as the sound, attenuates. Otherwise a vessel transforms into a ''bomb''. In this case the force of the shock wave and the sound achieves its maximal size and the explosion of an enormous force at the volume of appearing hydrogen just about 5 dm3 is heard. The intensive burning of hydrogen begins after the complete moistening of micro - and nanopowders of silicon by water and by hydrooxide of potassium (sodium) and stops at the complete using up of initial products of reaction. In sinking there is a silicate of potassium (sodium) which may be used as liquid glass. Thus, the processes of burning of hydrogen at different speed of its formation (different size of nano– and microstructural powders) testifies the sufficient high efficiency of the process of chemical synthesis of hydrogen as an alternative thermal energy source [8]. Storage of hydrogen through absorption of different materials is an important problem. Numerous experimental results showed that absorption of hydrogen MNS SP depended on a thickness of the layer and on the bulked weight of powder. In Figure 5 (curve 1) information on absorption of hydrogen by MNS SP is presented. The volume of bulked powder made up 0,785–1,570 cm3. From this information it is clean, that MNS SP with the increasing of the bulked weight promotes the hydrogensorption capacity. Increasing of the bulked weight, as well as of the volume of the powder results in the increase of its absorbing capacity to hydrogen. However multiplying the density of micro- and nanostructural powders in the set volume, on the contrary, reduces their absorbing ability a little (picture 5 (curves 2, 3)). From our point of view it is explained by the decrease of permeability of hydrogen into the volume of the made more compact powder at the set terms of hydrogenation. Activating of MNS SP at the temperature of hydrogenation of 623К during 30 minutes at pressure of hydrogen in a running reactor 0,25 MPa results in absorption of hydrogen in an amount necessary for formation of hydride of silicon of composition of SiH2. As turned out, in the result of the researches conducted in any case it is necessary to tend to the density of powders being within the limits of 0,8–1,0 g·cm-3. In this case the size of

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particles of the grinded up powder is within the limits of 500–6000 nm. It is expedient to conduct the process of hydrogenation of such powders at the temperature of 423–473К and pressure of hydrogen about 0,1 MPa. Bulked weight, gr

Hydrogensorption capacity, mass %

0,7

0,8

0,9

1,0

1,1

1,2

1,3

4,0

3,5

3,0

1 2 3

2,5

2,0

1,5 0,7

0,8

0,9

1,0

1,1

Specific density, gr/сm

1,2

3

4,5 4,0

Mass growth, mass %

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Figure 5. Conformity to the law of change of the hydrogensorption capacity of MNS SP depending on bulked weight (1), and on the specific density of powder (2, 3). 1,2-microstructural powder of silicon, got as a result of grinding down of wastes of semiconductor single-crystalline silicon, 3- nanostructural powder of silicon, got as a result of plazmochemical decompositions of monosilane.

3,5 3,0 2,5 2,0

1 2

1,5 1,0 0,5 0,0 0

1

2

3

4

5

6

7

8

9

10

Time of hydrogeneration, min

Figure 6. Absorption of hydrogen by microstructural powders made of single-crystalline silicon: without chemical treatment (1) and with chemical treatment (2) of the grinded up powder of silicon in 50% solution of potassium (KOH) hydrooxide during 30 minutes. Nanopowders and Nanocoatings: Production, Properties and Applications : Production, Properties and Applications, Nova Science Publishers,

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The maximal hydrogensorption capacity of such powders may achieve 4,0 masses % (Figure 6). At preliminary treatment of micro- and nanostructural powders in the mixture of acids or hydrooxides speed of the hydrogenation appeared to be higher than the speed of hydrogenation of powders without preliminary chemical treatment. It is evident than the treatment by acids and hydrooxides (by alkalis) MNS SP mixture creates the greater number of imperfect centers on every particle (grain), that reduces energy of adsorption of hydrogen on its surface. High enough sorption of hydrogen, as well as water, is conditioned by physical and chemical interaction, i.e. the processes are intermediate between processes on the surface of particles (physical process) and in the volume of particles (chemical process). Nanostructural particles as hydrogen absorbing medium are more suitable than microstructural particles, as they, due to a greater surface, are able to retain hydrogen at a room temperature. Moreover, «copulas» appear at the receipt of homogeneous nanostructural particles, being close-packed quadrangular grates from the parallel put cubes, distance between which makes 3 micrometer. Such piling of particles of silicon multiplies heat-sinking ability of the system due to arising up cavities in it. The geometrical packing of molecules of hydrogen inside multi-layered nanostructural particles provides the accumulation of it up to 4 masses %. This fact testifies that silicon powders grinded up to the size of particles of 500–6000 nm silicon and subjected to the chemical «grinding» or «polish» in mixture of acids or hydrooxides are good absorbers of hydrogen. We will especially mark two interesting phenomena found in the process of hydrogenation of MNS SP in the conditions of low pressure of hydrogen. At absorption of hydrogen at the temperature of 373-773К its pressure does not go down below 0,1 МPа, and at the initial size of pressure below this value formation of hydrides does not take place. Thus, the area with pressure of hydrogen more than 0,1МPа must correspond to steady existence of hydride of SiH2. It is necessary to adopt the temperature of 623К as the optimum temperature of hydrogenation, at which maximal mobility of atoms of silicon in the process of forming of hydride is evidently achieved. The second feature of reaction in these terms is the formation of high temperature form of Si2H4 (SiH2), stable at a room temperature. At a temperature higher than 723К for the hydride of such composition there is a reverse phase transition. However this phenomenon finds its explanation. Appearing at a low temperature ≤ 573К hydride is in the equilibrium state on conditions of hydrogenation in a running reactor. With terminating of thermal influence it passes to the metastable state, transformation from which into a new stable state requires overcoming of a certain potential barrier, for example, as a result of the further heating.Possibly, here will be some temperature hysteresis, i.e. a temperature higher than573К will be required. In this case the content of hydrogen in samples does not exceed 4 masses % in any case. Thus, a principally new method of hydrogenation by hydrogen of MNS SP through the hydrides of silicon in the conditions of relatively low pressures of hydrogen and temperature of hydrogenation is presented. Moreover, this method supposes limitation of temperature of hydrogenation to 723К, at temperature higher than that the reverse process of dehydrogenation begins (Figure 7).

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Consequently, manipulating a temperature, we can, both introduce hydrogen in MNS SP and extract it from their volume. Hydrogen entering the volume of MNS SP in the range of temperatures 373 – 723 K due to a diffusive process, can react not only with the atoms of silicon with formation of hydrides of SiH2 type but also with polymerization of hydrides into more complicated compositions of Si2H4, Si2H6 type. It is possible to come to the conclusion, that the initial MNS SP fully hydrogenates at the grains boundaries to its equilibrium states. Naturally, in this case in hydrogenation of MNS SP active centers, arising up at introduction of hydrogen also take part. These are particles of H, Si2+, SiH2, SiH3, etc.

3

Spee of hydrogen passing flow, dm /min

4,0 3,5 3,0 2,5 2,0 1,5 1,0 0,5 0,0 0

5

10 15 20 25 Duration of effusion process, min

30

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Figure 7. Conformity to the law of hydrogen extraction from powder of silicon by mass 28 gr and by the density of 0,8 g·cm-3.

The temperatures of starting and completion of hydrogenation of MNS SP are in certain dependence on duration of the process and on dispersion of powder. We found that the process has two obviously expressed stages, carried on temperatures: temperature of hydrogenation up to 723 to and temperature of de hydrogenation higher than 723К. The transition temperature between the stages is determined by the features of dispersion of MNS SP and structures of hydrides on the surface of particles. The researches of process of transporting of hydrogen into silicon powder in dependence of its dispersion and mass showed new conformities to the law of the process. Depending on dispersion of powder and its specific weight (density) the constituents of the product are divided into the rows of hydrides of SiH, SiH2, SiH3, SiH4 and Si2H4 and Si2H6 type. It is set that on the whole the results of hydrogenation are determined by dispersion of powders of silicon, temperature of process of hydrogenation, density of powders and speed (by pressure) of stream of hydrogen. Our researches showed that the speed of saturation of MNS SP by hydrogen practically does not depend on powder‘s nature, but depends on size of powder and its specific density. The advantage of microstructural powders with a specific less than 1,2 g·cm3 is found out (Figure 8). The maximal hydrogensorption capacity of such microstructural powders approaches 4 masses %. The low density of MNS SP allows to reduce the temperature of hydrogenation up to 473 – 573К and thus to remove the negative influence of the process of baking of powders.

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Speed of the process of hydrogenation is related to the composition structure of MNS SP It is set that in the process of transition from nano- to microsize an accumulation of hydrogen, both in volume and on the surface of powder takes place [11].

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Figure 8. Microrelief of powders with the specific density of 1,2 g·cm3.

CONCLUSION Multiplying of volume and of the bulked weight of MNS SP results in multiplying of water- and hydrogensorption capacities of powders. Chemical treatment of MNS SP in mixture of acids or concentrated hydrooxide (alkalis) is instrumental in an increase of their water- and hydrogensorption capacities. Maximal water- and hydrogensorption capacity of powders of silicon is achieved at their density of 0,81,0 g·cm-3, at the temperature of hydrogenation of 373-673К and pressure of hydrogen in the running system of 0,1- 1,0 MPa. In this case it is approaching to 4 masses %. Powders of grinded silicon with the size of particles of 500 – 6000 nm and subjected to the chemical «grinding» («polish») in the mixture of acids or hydrooxides are good sorbents for hydrogen. In the light of information of calculations on enthalpy of hydrogenation of silicon powders with different content of hydrogen it is set that with the increase of hydrogen content in the hydride compound sharp diminishing of exothermic process of hydrogenation takes place. It is desirable to execute the process of hydrogenation of silicon powders at the temperature 423 – 473К and pressure of hydrogen in the running system of 0,1 – 0,2 MPa. Complex use of products of interaction of MNS SP with water, utilization of heat and effective functioning of power cycle on the basis of nanostructural powder of silicon is the

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real technology of gaseous hydrogen for hydrogen energy of the nearest future. Application of MNS SP gives undeniable advantages: simple and without expenses of energy from outside technology of hydrogen production in the result of water decomposition: the necessity in storage and transportation of gasiform hydrogen disappears, which promotes fire- and explosionproof of this fragment of hydrogen energy substantially. In the process of the industrial production of nanopowders of silicon, including for other areas of their application their prime price will be reduced some times of their value. The application of MNS SP in the mobile small sources of hydrogen is already expedient.

REFERENCES

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[1]

Ponomarev-Stepnoi N.N., Pahomov V. В. Hydrogeneous economy and the future of manking // V mire nauki. 2006. №7 // Published in Neorganicheskie Materialy, 2008, Vol. 44, No. 5, pp. 519–523.3-7. [2] Apresyan L. A., Vlasov D.V., Vlasova T.V. and etc. Reactor with activated hydrogen for the synthesis of carbon nano-tubes // Published in Jurnal technicheskoi phiziki. 2006. v. 76. № 9. Pp.140-142. [3] Vinogradov D.V. Modern state of hydrogeneous power engineering // Voprosi atomnoi nayki i tehniki. Seriya vakyym, chistie materiali, sverhprovodniki. 2006. № 1. Pp.153155. [4] Miller E.L., Rocheleau R.E., Deng X.M. Design consideration for a hybrid amorphous silicon/photoelectrochemical multijunction cell for hydrogen production // International journal of Hydrogen Energy. 2003. № 28. Pp. 615-623. [5] Rocheleau R.E., Miller E.L. Photoelectrochemical production of hydrogen: engineering loss analysis. // Int. J. Hydrogen Energy. 1997. v. 22. №8. Pp. 771-782. [6] Tarasov B.P., Fokin V.P., Borisov D.N. and etc. Hydrogen accumulation by alloys of magnesium and rare metals with nickel // International scientific journal for Alternative Energy and Ecology. 2004. v. 15. №1. p.47-52. [7] Goidin V.V., Molchanov V.V., Buyanov R.A. Synthesis of intermetallic compounds hydrides at mechanochemical activation and increased hydrogen pressure // Published in Neorganicheskie Materialy. 2004. v. 40. № 11. Pp.1328-1332. [8] Kovalevskii A.A., Labunov V.A., Dolbik A.V., Saurov A.N., Basaev A.S., Strogova A.S. Investigation of characteristic properties of hydrogen synthesis at water decomposition by micro- and nanostructuried silicon powders // Published in Injenernophizicheskii jurnal. 2008. v. 81. № 3. Pp.587-591. [9] Kovalevskii A.A., Dolbik A.V., Strogova A.S. Investigation of influence of thermal treatment conditions in the process of compacting on the degree of micro- and nanostructured silicon powders saturation by water // Published in Microsistemnaya tehnika, 2007, v.10 (86), pp. 19-22. [10] Kovalevskii A.A., Strogova A.S. Generating and storage of hydrogen by micro- and nanopowders of silicon // XVIII Mezhdunarodnaya krimskaya konferentsiya «SVCHtechnika I telekommenikamynikascionnie technologii» 8-12 sentyabrya 2008г., Sevastopol, Ukraina, (September 8-12, 2008, Sevastopol, Ukraine (18th

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Nano- and Microstructural Silicon Powders in the Synthesis…

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[18] [19] [20]

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InternationalCremean Conferense ―Microwave and TelecomunicationTechnology‖) pp.608-609. Kovalevskii A.A., Dolbik A.V., Strogova A.S. Investigation of sorption ability of micro- and nanosized silicon powders // Published in Materiali. Technologii. Instrumenti. 2008. v.13. № 1, Pp.35-37. Klanchar M., Hughes T.G. System for Generating Hydrogen: - U.S. Patent № 5.634.341 (3.07.1997) and №5.867.978 (9.02.1999). Borislav Bogdanović, Richard A. Brand, Ankica Marjanović, Manfred Schwickardi and Joachim Tölle. Metal –doped sodium aluminum hydrides as potential new hydrogen storage materials // Journal of Alloys and Compounds. 2000. v. 302. № 1-2. pp. 36-58. Zaluska A., Zaluska L., Strom-Olsen O.L. Nanocrystalline magnesium for hydrogen storage // Journal of Alloys and Compounds. 1999. № 288. P.217-225. Malishenko S.P. Hydrogen as a power accumulator in electric power engineering // Published in Rosiiskii himicheskii jurnal. 1997. v. XLI. № 6. Pp.112. Andrievsky R.A. Hydrogen in nanostructures // Publishes in Uspehi fizicheskix nauk. 2007. v.177. №7. Pp.721-735. Kopcitanchuk I.G., Ivanov E.U., Boldirev V.V. Interaction of alloys and intermetallides produced by mechano-chemical methods with hydrogen // Uspehi himii. 1998. v. 67. № 1. Pp.75. Antonova M.M. Magnesium compounds as hydrogen accumulators / Kiev, Ukraina. IPM. 1993. Schwarz R.B. Hydrogen storage in magnesium-based alloys. MRS Buletin. 1999. v. 24. №11. p.40. Verbetsky V.N., Klymkin S.N. Interaction of magnesium alloys with hydrogen. Pergamon: Hydrogen Energy Progress. VII. 1998. vol.2. p.1319. Geigir‘ev T.F., Baranov A.P. Lyahov N.Z. Mechano-chemical synthesis of intermetallic compounds // Publishes in Uspehi himii. 2001. v. 70. № 1. Pp.52. Tarasov B.P., Shilkin S.P. On the possibility of emission and accumulating of highpurity hydrogen with the help of hydroforming intermetallic compounds // Publishes in Jurnal prikladnoi himii. 1995. v. 68. № 1. Pp.537. A.A. Kovalevskii, A.A. Shevchonok, V.A. Labunov, A.S. Strogova. Investigation of mechanism of hydrogen transportation into powder silicon // Publishes in Nano- I microsistemnaya tehnika. 2008. т.4(93). pp.13-16. A.A. Kovalevskii, A.A. Shevchonok, V.A. Labunov, A.S. Strogova. Characterictic properties of hydrogen diffusion into micro- and nanostructured silicon powders // Publishes in Materialy. Technologii. Instrumenti. 2008. v. 13, № 2, Pp.58-61. Korovskii Sh.J. Novie materiali electroradiotehniki i avtomatiki. Riga. 1980. 48 p. Pasinkov V.V., Sorokin V.S. Publishes in Materialy elektronnoi tehniki. / S.-Peterburg. 2001. 364 p. Strelov K.K. Theoretic basis of refractory materials technology // М.: Nayka, 1985. 480 p. A.A. Kovalevskii, A.A. Shevchonok, A.S. Strogova. Oxidation Behavior of Micro- and Nanostructured Silicon Powders // Published in Neorganicheskie Materialy, 2008. V. 44. № 5. pp. 519–523.

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In: Nanopowders and Nanocoatings Editor: V. F. Cotler

ISBN: 978-1-60741-940-2 © 2010 Nova Science Publishers, Inc.

Chapter 7

SEMICONDUCTOR CERAMIC MATERIALS PRODUCED FROM AIIBVI NANOPOWDERS N.N. Kolesnikov, E.B. Borisenko, V. V. Kveder, D.N. Borisenko, A.V. Timonina and B.A. Gnesin1 Institute of Solid State Physics, Russian Academy of Sciences, Institutskaya 2, Chernogolovka, Moscow distr., 142432 Russia

ABSTRACT

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Based on our experience of conventional crystal growth of AIIBVI, we have developed the techniques for nanocrystal production of some of these compounds. Further innovations in the technological process appear to significantly reduce the development and production cost of the final product with respect to conventional single crystals, while maintaining a high level of physical properties, which conform to leading standards. New technology of the vapor phase deposition was developed in our laboratory to produce CdTe nanoparticles of average diameter 8 nm. Further development of the technology allows us to produce 10-nm Cd-Zn-Te nanocrystals and to overcome difficulties in obtaining a ternary solid solution Cd1-xZnxTe (x = 0.04—0.1) with a stable chemical composition. Highly dense CdTe and Cd1-xZnxTe ceramics of high mechanical hardness and durability were produced by the room temperature process without any lubricants or binding materials. It was found that CdTe ceramics produced from nanocrystals undergo wurtzite— sphalerite transition under pressure. The polymorphic transition from the hexagonal to the cubic phase in Cd1-xZnxTe ceramics caused by annealing is discussed. The two-component texture composed of axial and {100} components was recorded in the as-compressed materials. The effect of annealing on grain growth and texture in the ceramics is considered. It was found that our compacted ceramics guarantees high transmittance in a wide IR 

Corresponding author. Tel.: +1-801-581-5491; fax: +1-801-581-4937 Email address: [email protected]

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N.N. Kolesnikov, E.B. Borisenko, V. V. Kveder, D.N. Borisenko, et al region of 6—25 m, high specific resistivity on the order of 1010 cm, and high microhardness on the order of 103 MPa. The obtained properties make these materials promising for use in IR optics and for ionizing-radiation detectors.

INTRODUCTION

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Cd1-xZnxTe (CZT) single crystals traditionally grown from melt are used in infrared technique and as materials for ionizing radiation detectors. The development of nanotechnologies provides possibilities for fabrication of materials with new properties and broadens the conventional application fields. Nanotechnological methods are promising for fabrication of ceramic materials from AIIBVI binary and ternary compounds. Among these materials very interesting are CdTe and Cd1-xZnxTe compounds which have numerous applications in infrared optics and ionizing radiation detectors. Particularly interesting is Cd0.9Zn0.1Te composition as one, used in ionizing radiation detectors operating at room temperature without cooling. As compared to the traditional melt growth of Cd1-xZnxTe crystals [1], the developed vapor deposition technique is simpler, because no high pressures are required to depress decomposition below the melting point and to decrease loss of volatile components. Furthermore, this technique provides production of a material with a definite uniform stoichiometric composition. This is an advantage, because a composition of a melt-grown bulk, especially Cd1-xZnxTe ternary solution, varies along the growth axis [1]. The objective of this study is to investigate the effect of external parameters, such as temperature, pressure of compaction, annealing temperature and time on phase composition, microstructure, and crystallographic texture which determine optical, electrical and mechanical properties of a material.

EXPERIMENTAL The CdTe and CZT nanopowders were deposited using the technique we have develpoped for vapor deposition of AIIBVI nanopowders [2]. To produce a nanopowder, pure CdTe single crystal or synthesized CZT were evaporated from a fused silica boat in helium gas flow. The powders were deposited on a gold substrate (Fig. 1). The temperature in the source zone of evaporation was 750—990ºC in the case of CdTe and 800—850ºC in the case of CZT, the helium flow velocity was 1000—1500 ml/min, the temperature in the deposition zone was 540—610º C. These temperature regimes and helium flow velocity are chosen to provide deposition of 8-nm CdTe and 10-nm CZT particles (Fig. 2). There is a dependence between a composition of Cd1-xZnxTe nanopowder and a solid source. It was found experimentally which composition of the evaporation source is required to obtain a definite composition of the nanopowder. For example, in order to obtain Cd0.9Zn0.1Te nanopowder composition, the source composition should be Cd0.5Zn0.5 Te. A difference between the source and the nanopowder compositions is due to different partial vapor pressures of the components and diffusion coefficients in helium. The problem of

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Semiconductor Ceramic Materials Produced from AIIBVI Nanopowders

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variations in chemical composition under thermal evaporation of AIIBVI compounds is discussed elsewhere [3, 4].

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Figure 1. Setup for vapor deposition of AIIBVI nanopowders.

Figure 2. Cd1-xZnxTe nanopoweder, TEM image.

A one-die scheme was used to compress CdTe or CZT in an Instron machine. The pressure ranged from 400 to 700 MPa. The temperature of compression varied in the range of Td = 20-200C (0.2-0.35 of the melting point Tm). After the required pressure had been achieved, the specimen was kept under pressure for 0.5-120 min. The lower temperatures in the compaction process are preferred for producing large amounts of material, since it ensures the simplest and most economical scheme to obtain material with an acceptable density. The

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upper temperature limit is restricted due to the tendency of AIIBVI compounds to oxidize and partially decompose, particularly at temperatures above 300C [2]. The density of the compressed specimens measured by weighing them in water was 92-97% of the calculated density, depending on the compression regime. The relative error of the measurements was ~0.2 %. X-ray phase analysis was carried out using a Siemens D500 diffractometer. A sample was fixed in a ring-shaped holder with a 2-mm accuracy of alignment. Either FeK or MoKwas used as the excitation source Crystal texture of the compacted samples was studied using the reflection (Schultz) technique. The tilt-angle range was 0-50 , the azimuth and the polar angle steps were 5 , and the exposure time at each point was 3 s. To obtain a pole figure (PF), the X-ray detector was set in a certain position based on the previously recorded diffractogram of the compacted CdTe. The background and the defocusing were taken into account when processing the obtained data. The microhardness was measured using a standard Vickers hardness testing technique. The load was 100 g. The light-transmission spectra were recorded within the wavelength range of 2.5-25 m using a spectrophotometer Specord 75 IR. The specific electrical resistivity was measured using the four-point-probe technique and direct current. The current-voltage characteristics were obtained using a bias voltage source produced by Ortec.

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RESULTS AND DISCUSSION Microstructures and textures of CdTe crystals compacted at temperatures in the range of 20-200 C have been studied. The highest density of the ceramics (97% of the calculated density) was achieved under pressure p = 650 MPa at room temperature.. The compacted ceramics consisted of grains with the average diameter 8 m (Fig. 3). In the compacted CdTe, grains are elongated along a certain crystallographic direction (Fig. 3). The average grain sizes of the specimens show little or no change under a constant load at room temperature, regardless of time under pressure. The volume fraction of pores was reduced by 2-3 % in a specimen kept for an hour under pressure. An increase in temperature of compaction caused grain growth. The volume fraction of recrystallized grains reached approximately 90 % in a specimen compressed for 60 min at 200 C under p = 650 MPa. The mean size of new grains was 60 m (Fig. 4). At the same time, micron and sub-micron pores continued to heal under pressure. A layer-by-layer etching of the specimens shows that grains grew throughout the whole volume. Pole figures {220} obtained for the basic cubic phase show a sharp axial texture. In addition, it was shown that 4-fold symmetry is observed in the {220} PF of the CdTe samples compacted at room temperature (Fig. 5). It is associated with another textural component {001}. It is quite unusual to observe such a sharp texture in a ceramic, since it is more common for compacted materials to produce rather scattered textures. It was found that though the texture of CdTe specimens compressed at 200 C remains somewhat axial (Fig. 6), it becomes increasingly more scattered than the texture of the specimens compressed at room temperature. Thus, grain growth and textural evolution in the CdTe ceramics compacted at temperatures in the range of 20-200 C, together with the fact that no

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strain-hardening was observed in the deformation curves, confirm that dynamic recrystallization takes place during the compression of the nanopowder.

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Figure 3. CdTe ceramics produced at T = 20C, p = 650 MPa, time under pressure t = 10 min.

Figure 4. CdTe ceramics produced at T = 200C, p = 650 MPa, time under pressure t = 10 min.

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Figure 5. Pole figure {220} of CdTe ceramics produced at T = 20C, p = 650 MPa.

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Figure 6. Pole figure {220} of CdTe ceramics produced at T = 200C, p = 650 MPa.

Figure 7. Diffractogram of CdTe nanopowder.

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Figure 8. Diffractogram of CdTe ceramics produced at T = 20C, p = 400 MPa.

Figure 9. Diffractogram of CdTe ceramics produced at T = 20C, p = 650 MPa.

The CdTe nanopowder consisted of particles of 8 nm in diameter, which was produced through the vapor phase. CdTe usually crystallizes as a cubic phase with sphalerite structure (a = 6.481 Å ) [5], which is stable up to near the melting point (Tm = 1092 C). However, hexagonal structure (wutrtzite) also exists under standard atmospheric pressure. A metastable hexagonal phase with wurtzite structure (a = 4.58 Å , c = 7.50 Å ) has been observed in films grown on substrates [6]. The diffractogram in Fig. 7 shows that the CdTe nanopowder is a mixture of two phases, wurtzite and sphalerite, and that the wurtzite phase dominates. X-ray diffractograms show that intensities of sphalerite lines increase, while the lines of the hexagonal phase weaken as pressure increases (Fig. 8). This means that the hexagonal phase transforms into the cubic phase under pressure. Wurtzite has almost completely converted to sphalerite at a pressure of 650 MPa at Td = 20ºC (Fig. 9). Grain boundaries were revealed by etching the compacted CdTe in a chemical solution containing 7 vol % Br and 93 vol % CH3OH. In the samples compacted under pressure p =

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400 MPa and then subjected to mechanical polishing and etching we observed plates of 20-50 m in diameter containing twins on the surfaces (Fig. 10 ). The volume fraction of such plates was about 2 % in the specimens that mainly consisted of grains of approximately 8-m diameter. Twin growth is typical of wurtzite-to-sphalerite phase transition in AIIBVI and AIIIBV crystals [7], so the observed structure gives additional evidence that the phase transition develops under pressure in the compacted CdTe. Further increase in pressure resulted in decrease in volume fraction of the twin plates, which disappear at p = 650C. They are replaced by grains with high-angle boundaries. The observed changes in the microstructure agree with the data of x-ray phase analysis.

Figure 10. CdTe ceramics produced at T = 20C,: (a) p = 400 MPa grains contain twin plates; (b) p = 600 MPa twin plates disappear.

Figure 11. Diffractogram of Cd0.9Zn0.1Te ceramics produced at T = 20C, p = 400 MPa. Nanopowders and Nanocoatings: Production, Properties and Applications : Production, Properties and Applications, Nova Science Publishers,

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Figure 12. Diffractogram of Cd0.9Zn0.1Te ceramics produced at T = 20C, p = 400 MPa and annealed at Tan = 300C, tan = 40 min.

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Oppositely, in Cd0.9Zn0.1Te ceramics made of the nanopowder the wurtzite-sphalerite transformation is not completed under pressure. According to the x-ray phase analysis, 5-10 % of the compacted ceramics is still in the hexagonal phase after compression (Fig. 11). The systematic shift of diffraction lines, with respect to pure CdTe, is caused by change in the unit cell parameter from а = 6.481 Ǻ to а = 6.442 Ǻ, which corresponds to the mentioned Cd0.9Zn0.1Te composition. The ceramics produced through compaction of CZT nanopowder at room temperature was then annealed in a sealed silica tube at 300ºC for 10—40 min. After 40 min of the annealing, hexagonal lines disappear from diffractograms (Fig. 12).

Figure 13. Grains containing twin plates in Cd 0.9Zn0.1Te ceramics produced at T = 20C, p = 400 MPa,. Nanopowders and Nanocoatings: Production, Properties and Applications : Production, Properties and Applications, Nova Science Publishers,

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Figure 14. The same place on the sample surface as in Fig. 13 before annealing of Cd0.9Zn0.1Te ceramics produced at T = 20C, p = 400 MPa, after annealing at Tan = 300C, tan = 40 min.

In parallel, some change were also observed in the microstructure. Grains containing twins were observed immediately after compaction (Fig. 13). Their volume fraction was about 5%. Twin grains firstly grew during annealing and, as the annealing time increased, the twin grains were replaced by grains with high-angle boundaries (Fig. 14). In parallel, some change were also observed in the microstructure. Grains containing twins were observed immediately after compaction (Fig. 13). Their volume fraction was about 5%. Twin grains firstly grew during annealing and, as the annealing time increased, the twin grains were replaced by grains with high-angle boundaries (Fig. 14). The microstructure of the ceramics compressed under 400 MPa is shown in Fig. 15. Grain size distribution in the ceramics is heterogeneous. According to the quantitative estimation, grain size distribution is bimodal with two extremes at 2 and 20 m, respective volume fractions are 52 and 40%. Fine grains are settled on grain boundaries of more coarse grains. The prevailing grain size after annealing was 40 m (Fig. 16).

Figure 15. Microstructure of Cd0.9Zn0.1Te ceramics produced at T = 20C, p = 400 MPa. Nanopowders and Nanocoatings: Production, Properties and Applications : Production, Properties and Applications, Nova Science Publishers,

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Figure 16. Microstructure of Cd0.9Zn0.1Te ceramics produced at T = 20C, p = 400 MPa, annealed at Tan = 300C, tan = 40 min.

Figure 17. Pole figure {220} of Cd0.9Zn0.1Te ceramics produced at T = 20C, p = 400 MPa.

Figure 18. Pole figure {220} of Cd0.9Zn0.1Te ceramics produced at T = 20C, p = 400 MPa and annealed Tan = 300C, tan = 40 min. Nanopowders and Nanocoatings: Production, Properties and Applications : Production, Properties and Applications, Nova Science Publishers,

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It has been found that in the compacted CZT ceramics the axial texture and weak cubic texture present (Fig. 17). However, both are less pronounced than in the CdTe ceramics fabricated from nanopowder. As the annealing time increases, the textures become weaker and more scattered. After annealing at 300ºC for 40 min no texture was observed (Fig. 18). Thus, though the structures of CdTe and Cd0.9Zn0.1Te ceramics made of nanopowders have much in common, the observed differences are substantial. In CdTe ceramics the phase transition to the stable cubic phase is completed during compaction, whereas in Cd0.9Zn0.1Te ceramics it is not. This material should be annealed to attain the stable crystallographic state. As was mentioned, the texture in the ceramics made of the three-component solid solution is weaker and more scattered than that of binary composition. Both results can be caused by the compaction conditions: the CdTe ceramics of 97% density was produced under pressure p = 650 MPa, while to produce Cd0.9Zn0.1Te of the same density, p = 400 MPa was required. Consequently, mechanical properties, in particular, microhardness of CdTe and Cd0.9Zn0.1Te ceramics also differ. The microhardness of CdTe ceramics as a function of pressure is shown in Fig. 19. The measuring technique has an accuracy of about 2 %. The microhardness increases both with time and pressure. It should be noted that the microhardness of melt-grown CdTe single crystals is about H = 430 MPa, which is more than three times lower than the corresponding value for the compacted CdTe H= 1500 MPa. Despite a decrease in defect concentration caused by the recrystallization, the microhardness of the specimens compressed at 200 C remains rather high H = 1400 МПа. This can probably be explained by the increase in density due to pore healing.

Figure 19. Dependence of microhardness on compaction pressure for CdTe ceramics.

The microhardness of CZT ceramics produced through nanopowder is lower, but as compared to the Cd0.9Zn0.1Te single crystal (H = 450 MPa) it is still rather high, H = 1000 MPa. The described annealing after compaction does not cause deterioration of mechanical hardness. The specific electrical resistivity () of the compacted CdTe specimens was tested using the four-point-probe technique, and a value of 5·109-5·1010 ·cm was obtained, which is much higher than the typical value of 2·108 ·cm for CdTe single crystal. This value is in

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good agreement with the values (approximately 1010 ·cm) obtained from current-voltage characteristics. Typical current-voltage characteristics are shown in Fig. 20. The infrared transmittance spectra of the compacted CdTe (Fig. 21) show significant light scattering in the short wavelength region (wavelengths < 6m). However, light transmission is rather high in the range of 6-25 m, but worse than that of CdTe single crystal. This is probably caused by rougher surface of the ceramics, because a proper polishing procedure for these materials has not been developed yet. Thus, the developed technique provides production of materials through the method which is far less complicated and less expensive than traditional melt growth. There is another advantage in production of CZT ceramics as compared to the single crystals: the nanopowder composition is constant and stable and is determined by the source composition, contrary to the ternary composition of Cd-Zn-Te single crystal, which varies along the crystal growth direction.

Dark current (nA)

100

10

sample2-direct sample2-reverse

1

sample3-direct sample3-reverse sample4-direct

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sample4-reverse

10

100

Applied bias voltage (V) 1000

Figure 20. Current-voltage characteristics of CdTe ceramics produced at T = 20C, p = 650 MPa.

Figure 21. Transmittance spectrum of CdTe in IR range: 1 – single crystal; 2 – ceramics. Nanopowders and Nanocoatings: Production, Properties and Applications : Production, Properties and Applications, Nova Science Publishers,

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CONCLUSION 1. A new technique was developed to produce CdTe and Cd1-xZnxTe (x = 0.04—0.1) nanopowders consisting of 8-10-nm particles through vapor deposition. 2. Ceramic materials of high 95-97% density were produced from these powders at room temperature 3. It has been shown that recrystallization goes both during compaction of the powders and during annealing after deformation. Grain size can be modified by external parameters, such as temperature and time of compression and annealing. 4. The produced ceramics materials have high mechanical hardness, the specific resistivity is higher than that of the single crystals, its transmission in infrared range is acceptable for optical applications.

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REFERENCES [1] Kolesnikov, N.N.; James, R.B.; Berzigiarova, N.S.; Kulakov, M.P. HRB and HPVZM shaped growth of CdZnTe, CdSe and ZnSe crystals. In: James RB, Franks LA, Burger A, Westbrook EM, Durst RD. Proceedings of SPIE: X-Ray and Gamma-Ray Detectors and Applications IV; 2002; v. 4784, pp. 93-104. [2] Kolesnikov, N. N.; Kveder, V. V.; James, R. B.; Borisenko, D. N.; Kulakov, M. P. Growth of CdTe nanocrystals by vapor deposition method, Nucl. Instr. and Meth. in Phys. Research, 2004, v. A527, pp. 73-75. [3] Kolesnikov, N.N.; Kulakov, M.P.; Fadeev, A.V. Changing of the ZnSe composition during zone melting. Izvestiya Akademii Nauk SSSR-Neorgan Mater, 1986, v. 22, no 3, pp. 395-398. [4] Kulakov, M.P.; Fadeev, A.V.; Kolesnikov, N.N. Determination of some properties of ZnSe melt and calculation of it composition. Izvestiya Akademii Nauk SSSR-Neorgan Mater, 1986; v. 22, no. 3, pp. 399-402. [5] Physics and Chemistry of AIIBVI Compounds, Ed. by M. Aven, J. S. Prener, NorthHolland Publishing Company, Amsterdam, 1967. [6] Pandeya, S.K; Tiwaria, U.; Ramana, R.; Prakasha, C.; Krishnab, V.; Duttab, V. and K. Zimik, Growth of cubic and hexagonal CdTe thin films by pulsed laser deposition, Thin Solid Films, 2005, v.473, no. 1, pp.54-57. [7] M. P. Kulakov, I. V. Balyakina, Solid State wurtzite-sphalerite transformation and phase boundaries in ZnSe-CdSe, J. of Crystal Growth 113 (1991) 653-657.

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In: Nanopowders and Nanocoatings Editor: V. F. Cotler

ISBN: 978-1-60741-940-2 © 2010 Nova Science Publishers, Inc.

Expert Commentary

THE ELECTROCHEMICAL SYNTHESIS OF THE TUNGSTEN CARBIDE NANOPOWDERS AND CARBON NANOTUBES Kh.B. Kushkhov and Kh. M. Berbekov Kabardino

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Balkarian State University, Nalchik, Russian Federation The principally new methods of electrochemical synthesis of the tungsten carbide nanopowders and carbon nanotubes have been developed in the centre of the collective use ―The X-ray diagnosis of the materials‖ of the Kabardino – Balkarian State University. The research is a part of the priority project ―The industry of the nanosystems and nanomaterials‖. The electrolysis of the oxide and oxide – halogenide ionic solutions is employed for obtaining the tungsten carbide nanopowders. The basis of the electrochemical synthesis of the tungsten carbide is the simultaneous electrodeposition of atoms of both the tungsten and carbon on the cathode which may later join to form the nanoparticles. The developed scheme is closed, ecologically safe and makes it possible to produce the nanodispersed powders of high purity and gives the opportunity to regulate the structure of the material. Comparing to the other known methods of manufacture, the high – temperature electrochemical synthesis allows to reduce the energy expenses and to simplify the process of production, because it consists of two stages only. The first is the very electrolysis while the separation of the acquired product from the electrolyte happens on the second one. The tungsten carbide nanopowders obtained with the electrochemical method could be pressed and sintered without introducing the linking additives. The exclusion of cementing agents results in the harder ceramics of the same viscosity. Besides, the tungsten carbide nanopowders possess the high electrocatalytic activity in the reactions of the electroemission and electrooxidizing of the hydrogen and can be used in various electrochemical devices instead of platinum (batteries and fuel cells, hydrogen energetics). Our tungsten carbide nanopowder can be characterized by: the specific area up to 20 square meters per gram, the particle size less than 200 nanometres, the impurity content varying in the limits of 0.01 to 0.001%. The electrochemical synthesis of the carbon nanotubes is processed in two regimes. In first one the graphite cathode is being electrochemically eroded during the electrolysis of the

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individual chloride melts of alkali and alkali – earth metals and their mixes. Nanotubes are being produced after intercalating of the alkali and alkali – earth metals into the graphite. In the second one the chloride melts of alkali metals containing the lithium carbonate are being electrolysed at the excess pressure of the carbon dioxide (up to 15 atmospheres). The ―building‖ unit of the nanotubes in the last case is the carbon after the electroreduction of the carbon dioxide dissolved in the melt.

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INDEX

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A absorption, 14, 30, 31, 32, 35, 39, 40, 44, 147, 181, 182, 187, 189 absorption coefficient, 31 absorption spectroscopy, 40 acceleration, 16 acceptor, 35, 39 accuracy, 30, 32, 181, 198, 206 acetate, 6, 7, 12, 17 acetic acid, 7, 52, 182 acetone, 126 achievement, 16 acid, xii, 7, 15, 28, 52, 70, 82, 84, 90, 101, 103, 106, 182 acidic, 74, 78, 82, 95, 101, 102 acidification, 102 acoustic, 47, 181 activation, xiv, 5, 88, 124, 125, 126, 133, 135, 136, 140, 141, 143 activation energy, xiv, 124, 125, 126, 133, 135, 136, 140, 141, 143 active centers, 190 actuators, xi, 1, 2 additives, 15, 126, 127, 128, 133, 134, 135, 143, 149, 209 adhesion, 72, 73, 74, 75, 80, 87, 89, 91, 94 adhesion strength, 73 adjustment, 105 adsorption, xiii, 38, 72, 80, 88, 90, 98, 99, 101, 105, 107, 109, 116, 117, 118, 189 aerosol, 15 Ag, ix, xiii, xiv, 73, 123, 124, 126, 127, 128, 133, 134, 135, 143 age, 74, 75, 76, 77, 79, 81, 87, 90, 91 agent, 102, 103, 148, 170, 184 agents, 7, 12, 101, 149, 152, 209

aggregates, 7, 15, 19, 107 aggregation, 102, 107 aggression, 87 aging, 74, 103, 104, 105 aging process, 105 AIIIBV, 202 air, 4, 7, 43, 51, 71, 73, 79, 101, 126, 181, 183 alcohol, 7, 11, 182 alkali, xiii, 14, 88, 104, 123, 210 alkaline, 7, 12, 43, 48, 72, 88 alloys, xii, 69, 71, 92, 93, 94, 95, 124, 125, 141, 148, 152, 181, 193 alpha, 71 alternative, 8, 38, 48, 49, 126, 179, 187 alternative energy, 179 alters, 5 aluminosilicate, 102, 105, 107, 109, 118 aluminosilicates, 108, 117 aluminum, xiv, 71, 95, 104, 105, 106, 107, 109, 110, 118, 147, 148, 149, 151, 193 ambient pressure, 19 ammonia, 44, 101, 103, 106 ammonium, 74, 101, 104, 106, 174 ammonium hydroxide, 101, 106 amorphous, xii, xiii, xiv, 2, 7, 8, 12, 19, 25, 71, 86, 87, 88, 109, 123, 124, 125, 126, 127, 128, 130, 131, 132, 133, 134, 135, 136, 137, 138, 140, 141, 142, 143, 183, 192 amorphous phases, xiv, 124, 143 amplitude, 30, 31 analytical techniques, 110, 119 anatase, 8, 86, 87, 88, 89, 91 anisotropy, 45, 141 annealing, xiv, xv, 7, 8, 15, 45, 52, 74, 87, 88, 124, 128, 130, 131, 132, 133, 136, 137, 138, 139, 140, 141, 142, 143, 144, 195, 196, 203, 204, 206, 208 anode, 164 anomalous, 8, 141 antiferromagnetic, 44, 47, 52

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Index

apatite, xii, 69, 72, 73, 75, 82, 87, 91, 93, 95, 96 aqueous solution, 9, 11, 94, 98, 100, 103, 116, 117, 118, 119 aqueous solutions, 94, 116, 119 aqueous suspension, 9 Arrhenius equation, 140 ASI, 66 aspect ratio, xii, 70, 71, 83, 88 assignment, 78 assumptions, 30, 107 atmosphere, 7, 15, 75, 76, 77, 95, 106, 136, 150, 164, 167 atmospheric pressure, 164, 183, 201 atomic emission spectrometry, 110 atomic force, 95 atomic force microscopy (AFM), 41, 43, 95 atomic orbitals, 47 atoms, 19, 20, 24, 28, 30, 32, 34, 37, 40, 44, 128, 133, 134, 136, 189, 209 attachment, 80, 81, 89, 90, 94

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B back, 25, 32, 91, 149 baking, 183, 190 band gap, 45 bandgap, 33, 51 barium, 5, 6, 7, 8, 9, 10, 12, 14, 15, 16, 17, 18, 23, 27, 33, 34, 35, 36, 43, 50, 53 barrier, 52, 83, 84, 85, 189 basic research, 33, 53 baths, 74 batteries, 82, 209 beams, 25 behavior, xiii, 7, 18, 21, 22, 31, 33, 37, 41, 44, 45, 46, 52, 53, 70, 93, 96, 97, 98, 116, 117, 126, 141, 183 benefits, 102 benzene, 102 bias, 67, 198 bi-doped, 24, 25 binding, xv, 30, 76, 78, 195 binding energy, 30, 76, 78 bioactive materials, 96 biocompatibility, xii, 69, 70, 71, 72, 82, 86, 93 biological activity, 72, 81 biological systems, 17 biomaterial, 72, 80, 90, 95 biomaterials, 96 biomedical applications, xii, 44, 69, 82, 92 biomineralization, 17 biomolecules, 18 bioseparation, 101

biosynthesis, 6, 17, 18 bismuth, 46 blocks, 4 blood, 88 blood plasma, 88 body fluid, 88, 94, 96 body temperature, 73 bonding, 28, 33, 73, 100, 127, 134 bonds, 35, 44 bone growth, 90, 91, 96 breakdown, 83 broad spectrum, xi, 1, 2 buffer, 49, 101 building blocks, 4 bulk crystal, 24, 36 bulk materials, xi burning, 186, 187

C CAD, 62 calcination temperature, 4, 8 calcium, xii, 69, 70, 72, 73, 74, 75, 78, 79, 80, 88, 92, 93, 94, 95 calibration, 153 CAP, 66 capacitance, 48, 53, 124 capsule, 70 carbide, xiv, 147, 148, 149, 152, 157, 161, 164, 166, 167, 170, 173, 174, 182, 209 carbides, 152 carbon, 8, 32, 33, 79, 104, 119, 153, 153, 155, 157, 158, 159, 160, 161, 162, 163, 166, 167, 169, 175, 180, 182, 209 carbon dioxide, 33, 79, 104, 210 carbon nanotubes, 209 carbonates, xii, 1, 3, 4, 78 carburization, 157, 158, 161, 162, 164, 167, 170, 171, 175 carrier, 50, 101, 103, 104, 106, 107, 109, 110, 111, 112, 113, 134, 153, 164, 170, 171, 174, 179 catalysis, 92 catalyst, 103, 105 cathode, xii, 20, 42, 69, 73, 74, 164, 209 cathode ray tube, 20 cathodic process, xii, 69 cation, 2, 10, 28, 29, 38, 48, 85, 86 cavities, 19, 185, 189 CdZnTe, 208 cell, xii, xiii, 2, 3, 10, 12, 22, 35, 37, 39, 47, 48, 49, 69, 70, 72, 80, 81, 82, 89, 90, 91, 92, 94, 125, 181, 192, 203 cell culture, 80, 81

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Index cell growth, 81, 90, 91 cellulose, 96 ceramic, xi, 1, 2, 3, 4, 6, 12, 14, 16, 19, 34, 47, 48, 51, 161, 167, 196, 198 ceramics, xv, 2, 4, 6, 10, 11, 17, 45, 72, 163, 195, 198, 199, 200, 201, 202, 203, 204, 205, 206, 207, 208, 209 CH3COOH, 183 CH4, 148, 153, 154, 155, 156, 158, 159, 160, 161, 162, 164, 165, 166, 167, 168, 170, 171, 174, 175 channels, 142 charged particle, 31 chelating agents, 7 chemical etching, 83 chemical interaction, 189 chemical properties, 20, 72 chemical reactions, xiv, 104, 147, 181 chemical stability, 43, 82 chemical vapor deposition, 33, 103 chemical vapor synthesis, xiv, 147, 148, 149 chloride, 14, 106, 148, 153, 210 chlorination, 148 chromatography, 48 clusters, 19, 37 CMOS, xi, 1, 2 Co, 13, 43, 51, 99, 148, 149, 152, 153, 154, 155, 156, 157, 158, 159, 162, 163, 164, 170, 174, 175, 176, 177 CO2, 33, 75, 76, 104 coatings, xi, xii, xiii, 69, 70, 72, 73, 74, 78, 79, 80, 81, 92, 93, 94, 95, 96, 102 cobalt, xiv, 43, 99, 147, 148, 149, 152, 159, 170, 174, 175 collagen, 72 colloids, 102 commercialization, 16 communication, 23 compaction, 196, 197, 198, 203, 204, 206, 208 compatibility, 43 compensation, 10 complexity, 103 components, xiii, xv, 4, 32, 48, 97, 112, 161, 184, 185, 195, 196 composites, 45, 101, 102, 152, 163 compounds, xiii, xv, 4, 9, 16, 37, 46, 53, 73, 98, 99, 101, 105, 116, 124, 133, 149, 195, 196, 197, 198 compressive strength, 152 concentration, 6, 10, 12, 14, 23, 24, 32, 35, 37, 39, 45, 50, 78, 79, 103, 107, 134, 149, 157, 158, 161, 167, 206 conception, 144 condensation, 7, 8, 19, 72, 85, 103, 105, 152 condensed matter, 35, 119

213

conduction, 39 conductive, 41, 43, 49 conductivity, 38, 42, 43, 48, 71, 124 confinement, 35, 39, 45 conformity, xii, 69, 184 connective tissue, 70 consensus, 30 consolidation, 152, 161, 162 constant load, 198 constituent materials, 51 consumption, 47, 49 contaminants, xiii, 97, 98, 99, 101, 105, 116, 118 contamination, 4, 12, 38 content analysis, 157 control, 4, 7, 8, 9, 12, 15, 42, 47, 49, 51, 52, 53, 72, 73, 75, 81, 101, 103, 105, 112, 116, 118, 149 conversion, 10, 12, 148, 167 cooling, 164, 165, 196 copper, 164 copulas, 189 core-shell, 30 correlation, 44, 117, 118, 142, 181, 183, 186 correlation coefficient, 117, 118 corrosion, xii, 69, 71, 82, 87 cost saving, 14 costs, 7, 9 couples, 14 coupling, 34, 44, 46, 47, 52 covalent, 35 covalent bond, 35 critical behavior, 52 critical current, 84 critical current density, 84 cross-linked polymers, 105 crust, 180 crystal growth, xv, 7, 15, 16, 76, 77, 80, 95, 141, 195, 207 crystal structure, 8, 10, 29, 35, 36, 38, 71, 88, 109, 124 crystal structures, 109, 124 crystalline, xii, xiii, xiv, 2, 7, 8, 10, 12, 15, 17, 19, 22, 25, 26, 29, 35, 44, 45, 53, 55, 57, 86, 88, 89, 95, 109, 123, 124, 125, 127, 128, 131, 133, 134, 136, 138, 142, 143, 145, 181, 183, 188 crystallinity, xiv, 15, 19, 73, 85, 90, 94, 123, 124, 128, 129, 130, 138 crystallites, 5, 8, 15, 78, 124, 130, 132, 138, 141, 142, 143, 180, 182 crystallization, xiii, xiv, 6, 8, 9, 10, 14, 24, 26, 123, 124, 125, 126, 128, 129, 130, 132, 133, 134, 135, 136, 138, 140, 141, 142, 143, 144 crystallization kinetics, 14, 126, 137, 143, 144

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214

Index

crystals, 15, 22, 29, 37, 39, 51, 75, 76, 78, 79, 80, 81, 88, 89, 141, 196, 198, 202, 208 culture, 80, 81 cycling, 124 cylindrical reactor, 164, 165 cytotoxicity, 81

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D database, 109 decay, 45, 77, 85, 182 decomposition, xiv, 5, 8, 9, 15, 104, 161, 164, 167, 170, 179, 181, 183, 184, 191, 196 defaults, 11 defects, 4, 10, 16, 20, 21, 22, 25, 26, 28, 35, 36, 38, 45, 134 deformation, 35, 72, 199, 208 degenerate, 36 degradation, xi, 38 degree of crystallinity, 12, 129 degrees of freedom, 46 dehydration, 8, 10 dehydrogenation, 189 delivery, 101, 164 density, 18, 19, 21, 24, 28, 32, 70, 75, 78, 79, 80, 82, 84, 85, 88, 107, 134, 173, 183, 185, 187, 188, 190, 191, 197, 198, 206, 208 dental implantations, 70 dental implants, 70 deposition, xii, xiii, xv, 19, 33, 50, 69, 70, 71, 72, 73, 74, 75, 76, 77, 78, 80, 82, 88, 89, 92, 93, 94, 95, 96, 103, 110, 127, 195, 196, 197, 208 deposition rate, 127 deposits, 75, 78, 79, 80, 104 detection, 38, 44 deviation, 26, 136 diamagnetism, 99 dielectric constant, 14, 41, 48, 50, 51, 85 dielectric materials, 48, 49, 50 dielectrics, 49, 53 differential scanning, xiv, 124, 128 differential scanning calorimeter, xiv, 124, 128 diffraction, xiii, xiv, 10, 11, 21, 22, 25, 29, 30, 35, 36, 77, 98, 99, 109, 110, 124, 127, 128, 130, 131, 132, 137, 138, 142, 181, 203 diffusion, xiii, 4, 5, 6, 38, 77, 85, 97, 98, 100, 136, 149, 196 diffusivities, 38 dimensionality, 4 dipole, 34, 46, 48 dipole moment, 48 discs, 78, 124 dislocations, 26, 95

disorder, 36, 61 dispersion, xi, 98, 102, 103, 181, 183, 184, 185, 190 displacement, 26, 31, 49 dissolved oxygen, 101 distilled water, 106, 184 distortions, 31 distribution, 4, 5, 6, 7, 9, 11, 15, 16, 17, 19, 20, 22, 24, 30, 31, 32, 44, 45, 48, 103, 107, 141, 153, 172, 183, 204 distribution function, 31 division, 180 domain structure, 47 domain walls, 28, 29 donor, 39 dopant, 6, 8, 39 dopants, 43 doped, 24, 38, 39, 40, 42, 45, 46, 50, 52, 70, 94, 193 doping, 42, 149 drinking, 98, 105 drinking water, 98 drug delivery, 101 drying, 6, 8, 79, 103, 104, 105, 106, 107, 152 DSC, xiv, 124, 126, 128, 135, 136, 140, 141, 143 DSC method, 136, 140, 143 durability, xv, 195 duration, 10, 126, 137, 138, 139, 140, 186, 187, 190

E earth, xi, 1, 2, 6, 24, 43, 44, 45, 48, 50, 210 ECM, 80 ecological, 179 effluent, 106, 107 electric charge, 49 electric field, 42, 46, 47, 48, 52, 73, 85 electrical conductivity, 42, 43, 48, 124 electrical properties, 6, 44, 124, 125, 126, 134 electrochemical deposition, xii, 69, 76, 77, 94, 95 electrocrystallization, 94 electrodeposition, 78, 79, 88, 89, 94, 209 electrodes, 48, 49, 52, 95 electrolysis, 78, 179, 209 electrolyte, xiii, 38, 42, 48, 70, 78, 79, 83, 85, 88, 90, 95, 209 electrolytes, xii, 69, 70, 78, 79, 82, 83, 94, 96 electromagnetic, 14, 21, 31, 34 electron, xiii, xiv, 3, 9, 11, 17, 18, 20, 22, 24, 25, 27, 28, 30, 32, 34, 35, 39, 98, 99, 108, 109, 124, 134, 137 electron beam, 20, 24, 25, 32 electron density, 134 electron diffraction, 11, 21

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Index electron microscopy, xiii, xiv, 9, 17, 18, 20, 22, 24, 25, 27, 32, 98, 99, 108, 124 electron spin resonance, 109 electronic circuits, 48 electronic structure, 32, 33, 34, 45 electronic systems, 50 electrons, 20, 21, 22, 24, 26, 32, 33, 39, 42, 43, 46 electrophoresis, 93, 116 electroreduction, 210 electrostatic force, 36, 41 emission, 20, 32, 39, 45, 110, 184 endurance, 49 energy, xiii, xiv, 5, 9, 14, 15, 19, 21, 28, 30, 32, 33, 34, 35, 37, 39, 40, 45, 72, 76, 98, 99, 124, 125, 126, 133, 134, 135, 136, 140, 141, 143, 153, 179, 184, 185, 189, 191, 209 energy efficiency, 5 energy supply, 184 environment, 40, 44, 48, 72, 80, 82, 101, 105, 109, 119, 180, 184 environmental contaminants, 105 environmental protection, 51 epitaxial growth, 24 EPR, 109 equilibrium, 76, 82, 85, 116, 117, 124, 150, 154, 189, 190 equilibrium state, 189, 190 estimating, 30 etching, 182, 198, 201 ethanol, 11, 43, 103, 149 ethylene glycol, 83 evaporation, 19, 105, 196 evolution, 3, 10, 50, 82, 198 EXAFS, 30, 127 excitation, 19, 33, 34, 35, 39, 52, 198 exciton, 39, 45 exclusion, 209 expenditures, 179 experimental condition, 150, 154, 161, 164, 176 exposure, 39, 43, 48, 198 external magnetic fields, 99 extracellular matrix, 80 extraction, 38, 190 extrapolation, 114, 115

F feeding, 149, 150, 155, 156, 157, 158, 159, 160, 161, 164, 165, 166, 168, 171, 172, 173, 174, 175 feedstock, 12 ferric oxide, 99 ferrimagnets, 52, 99 ferrite, 44

215

ferroelectrics, 41, 46, 47, 51, 52 ferromagnetic, 18, 44, 46, 47, 119 ferromagnetism, xi, 1, 2, 46, 47, 52, 99 ferromagnets, 99 FFT, 11, 26, 131, 132, 142 fibronectin, 80 fibrous tissue, 70 field-emission, 32 film, xiv, 16, 28, 29, 49, 52, 70, 73, 82, 83, 84, 85, 101, 102, 109, 111, 123, 124, 125, 127, 128, 131, 132, 133, 134, 135, 137, 138, 140, 141, 142, 143 film formation, 84 film thickness, xiv, 73, 84, 123, 127 films, xiii, xiv, 19, 29, 37, 50, 67, 78, 93, 95, 123, 124, 125, 126, 127, 128, 129, 130, 132, 133, 134, 135, 136, 141, 142, 143, 144, 201, 208 filopodia, 81, 89 finite size effects, 45 fire, xiv, 179, 191 fixation, 70 flame, 164, 166, 167, 170, 171, 187 flexibility, xii, 69, 103, 116, 118 flight, 37 flow, 16, 100, 126, 149, 150, 155, 156, 157, 159, 160, 165, 166, 167, 168, 171, 174, 175, 196 flow rate, 16, 150, 155, 156, 157, 159, 160, 165, 166, 167, 168, 171, 174, 175 fluid, 90, 98 fluoride, xii, 70, 82, 83, 84, 85, 86, 87, 105, 117, 118, 119 fluoride ions, 82, 84 fluorides, 101 focusing, xii, 2, 38 Fourier, xiii, 11, 13, 26, 27, 34, 98, 99, 109, 131 Fourier transform infrared spectroscopy, xiii, 98, 99, 109 fragmentation, 37 free energy, 148, 150 Freundlich isotherm, 116, 118 FTIR, 78, 99, 109, 111 FT-IR, 79 fuel cell, 38, 42, 43, 48, 209 functionalization, 101 fungus, 17 fusion, 82 FWHM, 30, 130

G gadolinium, 99 gallium, 43 gamma, 208 gamma radiation, 40

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216

Index

gamma rays, 40 gas, xiii, xiv, 3, 4, 16, 19, 38, 43, 48, 64, 72, 74, 82, 98, 99, 107, 126, 135, 136, 140, 141, 147, 148, 149, 150, 152, 161, 164, 165, 166, 167, 168, 170, 171, 174, 175, 196 gas chromatograph, 48 gas phase, xiv, 3, 4, 19, 147, 150 gas sensors, 43, 48, 72, 82 gases, 19, 38, 43, 48, 157, 164, 171 gel, xii, 6, 7, 12, 15, 28, 44, 45, 70, 72, 100, 102, 103, 104, 105, 106, 120, 183 gelation, 31, 103, 105 gels, 7, 9, 96, 105 generation, 30, 46, 48, 72, 74, 82, 85, 86, 103 generators, 52 Gibbs energy, 9 Gibbs free energy, 150 glass, xiii, 45, 123, 124, 126, 134, 183, 187 glasses, 141 glycerol, xiii, 70, 71, 83 glycol, 12, 52, 83 gold, 43, 196 grain, xii, xv, 7, 10, 16, 19, 20, 21, 22, 26, 28, 30, 37, 39, 45, 46, 69, 70, 72, 108, 141, 151, 152, 153, 154, 155, 158, 161, 162, 164, 166, 167, 169, 171, 173, 175, 176, 180, 189, 195, 198, 204 grain boundaries, 19, 72, 204 grains, xii, 4, 69, 72, 78, 81, 181, 183, 190, 198, 202, 204 graphite, 164, 209 gravity, 6, 16 grazing, 127 groups, 14, 34, 35, 37, 78, 90, 98, 101, 110, 180

H H2, 73, 74, 96, 110, 148, 153, 154, 155, 156, 157, 159, 160, 161, 164, 165, 166, 167, 168, 170, 171, 172, 173, 174, 175 halogen, 148 handling, 9 hardening, 199 hardness, xv, 152, 195, 198, 206, 208 hazards, 48 healing, 206 health, 48 heat, xiv, 5, 6, 7, 8, 9, 11, 14, 36, 38, 44, 71, 72, 73, 86, 88, 125, 128, 129, 131, 136, 161, 162, 163, 167, 169, 170, 179, 182, 183, 189, 191 heat capacity, 9 heating, 7, 8, 9, 14, 31, 44, 101, 105, 106, 125, 128, 129, 135, 136, 137, 138, 139, 140, 184, 189

heating rate, 7, 8, 14, 105, 106, 128, 129, 135, 138, 140 height, 42 hematite, 14 heterogeneity, 183 heterogeneous, 10, 16, 74, 118, 133, 204 heterostructures, 25, 50 high pressure, 196 high resolution, 131, 138 high temperature, 4, 6, 22, 47, 73, 166, 167, 170, 189 histological, 93 homogeneity, xii, 1, 3, 4, 15, 101, 103, 104 homogenized, 6 homogenous, 84 HRTEM, 9, 11, 13, 17, 18, 23, 25, 26, 27, 28 HSC, 150, 154 hybrid, 48, 52, 192 hydration, 92 hydride, 187, 189, 191 hydrides, 180, 189, 190, 193 hydro, 162, 183 hydrocarbon, 168, 179 hydrocarbons, 162 hydrochloric acid, 106 hydrodynamic, 116 hydrofluoric acid, 82, 101, 181 hydrogels, 96 hydrogen, xiv, 72, 73, 82, 85, 148, 149, 151, 152, 155, 161, 162, 163, 166, 167, 168, 169, 170, 176, 179, 180, 181, 184, 185, 186, 187, 188, 189, 190, 191, 192, 193, 209 hydrogen gas, 167 hydrogenation, 181, 187, 188, 189, 190, 191 hydrolysis, 7, 8, 15, 17, 103, 105 hydrophilic, 183 hydrophilicity, 72 hydrophobicity, 183 hydrothermal, 6, 9, 10, 11, 12, 14, 21, 22, 23, 26, 27, 35, 36, 38, 94, 100, 103 hydrothermal process, 9, 11, 12, 14, 27 hydrothermal synthesis, 9, 14, 27, 35 hydroxide, 6, 7, 12, 14, 35, 73, 79, 84, 101, 106, 109, 184, 186 hydroxides, 73, 79, 110, 181 hydroxyapatite, xii, xiii, 69, 70, 74, 75, 88, 89, 93, 94, 95, 96 hydroxyl, 8, 10, 11, 14, 34, 35, 88, 104, 105 hydroxyl groups, 14, 34 hyperthermia, 44 hysteresis, 111, 189

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Index

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I IBM, 28 id, 123 identification, 25 images, 11, 13, 17, 18, 20, 21, 22, 23, 24, 25, 26, 27, 28, 32, 34, 42, 91, 128, 130, 131, 138, 142 imaging, 21, 24, 28, 29, 39, 101 immersion, 72, 87 implants, xii, 69, 70, 72, 82, 93 IMS, 37 in situ, 7, 10, 21, 124, 136 in vitro, 71, 81 in vivo, 93 inactive, 36 incidence, xiv, 124, 127, 129 inclusion, 21 incubation, 91 indices, 130, 136 indium, 126, 133, 143 induction, 72, 86, 127 industrial, xv, 16, 17, 152, 163, 176, 179, 192 industrial application, 17, 152, 163, 176 industrial production, xv, 179, 192 industry, 48, 53, 170, 180, 209 inelastic, 32, 35 inert, 7, 106, 164 infinite, 105 infrared, 34, 35, 36, 39, 105, 106, 107, 196, 207, 208 infrared light, 34 infrared spectroscopy, xiii, 98, 99, 109 inhibition, 102 initiation, 88 injection, 49, 100, 182 inorganic, 6, 17, 98, 101, 104, 110 inorganic salts, 104 insertion, 49 insight, 32 instability, 85 instron, 197 insulation, 164 insulators, 53 integrated circuits, 180 integrity, 37, 42 interaction, xiv, 19, 37, 42, 44, 45, 70, 71, 98, 100, 104, 179, 184, 186, 191 interaction process, 19 interactions, xiii, 21, 31, 45, 72, 80, 97, 99, 100 interface, 25, 49, 71, 85, 120, 132, 142 interference, xiii, 24, 25, 42, 98, 99 intermetallic compounds, 148 interphase, 185 interstitial, 2, 42

217

interval, 180 intrinsic, 9, 90, 114, 115 inversion, 25 Investigations, 183 IOC, 129, 130, 136, 142 ion beam, 37, 38, 125 ion implantation, 24 ion mass spectroscopy, 37 ion transport, 42 ionic, 42, 88, 101, 113, 116, 117, 209 ionic solutions, 209 ionicity, 36 ionization, 32 ionizing radiation, 196 ions, 2, 4, 5, 6, 8, 10, 11, 14, 15, 23, 24, 29, 35, 37, 38, 39, 45, 46, 49, 70, 73, 74, 76, 79, 81, 82, 84, 85, 88, 89, 90, 104, 105, 117 IR spectroscopy, 34, 35 iron, 99, 102, 119, 148 irradiation, 30, 105 isothermal, xiv, 8, 124, 126, 136, 137, 138, 139, 140, 141, 142, 143 isothermal crystallization, 140 isothermal heating, 136 isotherms, 116, 117, 118 isotopes, 40 isotropic, 15

J JEM, 153 judge, 109 Jung, 56, 59

K kernel, 181 kinetics, 5, 9, 14, 87, 88, 89, 98, 126, 136, 137, 140, 141, 143, 144, 149 KOH, 181, 183, 184, 188

L lamellae, 81 lamina, 186 laminar, 186 Langmuir equations, 117 lanthanum, 38, 51 laser, 19, 35, 124 laser ablation, 19

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218

Index

lattice, 5, 8, 10, 11, 14, 16, 17, 20, 21, 22, 24, 25, 26, 30, 35, 36, 38, 42, 46, 85, 128, 131, 132, 134, 136, 142 lattice parameters, 128 law, 29, 182, 184, 188, 190 leakage, 49, 50 lens, 25, 28, 35, 38 lenses, 21, 28 ligand, 52, 105 light scattering, 35, 153, 207 light transmission, 207 likelihood, 32 limitation, xi, 189 linear, 40, 78, 84, 118 linear regression, 118 liquid phase, xiii, 97, 98, 117 lithium, 32, 210 loading, 104, 183 low energy cluster beam deposition, 103 low power, 47 low temperatures, 8, 9, 11, 45 lubricants, xv, 195 luminescence, 39, 45

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M maghemite, 99 magnesium, 38, 148, 149, 181, 193 magnesium alloys, 193 magnet, 104, 106, 107, 119 magnetic, viii, xi, xiii, 1, 2, 3, 19, 28, 37, 40, 44, 46, 47, 52, 53, 79, 97, 98, 99, 100, 101, 102, 103, 104, 106, 107, 109, 110, 111, 112, 113, 116, 117, 119 magnetic field, xi, 1, 2, 44, 46, 47, 52, 99, 100, 109, 111 magnetic materials, xiii, 97, 99, 110, 119 magnetic moment, 44, 100 magnetic particles, xiii, 97, 99, 100, 101, 103, 107, 109, 110, 111, 112, 113, 116 magnetic properties, xii, 1, 3, 40, 44, 99, 111 magnetic resonance, 101 magnetic resonance imaging, 101 magnetic structure, 52 magnetism, 46, 98, 99, 100, 111, 116, 147 magnetite, 99, 100, 101, 102, 106, 108, 109 magnetization, 44, 45, 46, 47, 99, 100, 101, 110, 111, 112 magnetizations, 44, 99 magnetoresistance, xi, 1, 2, 44, 67 magnetostriction, 30 magnetron sputtering, 126 manganites, 44, 46, 47, 52

manifold, 101 manipulation, 47 manufacturing, 2, 7, 100 mapping, 158, 175 MAS, 110, 112 masking, 182 mass spectrometry, 37 mass transfer, xiii, 88, 97, 98, 116 material surface, 80 matrix, xi, 7, 16, 24, 45, 51, 72, 80, 92, 130, 131, 132, 133, 142, 152, 183 MCA, 104, 105, 106, 107, 109, 111, 112, 113, 114, 116, 117 measurement, xiii, 33, 35, 38, 107, 123, 124, 126, 136, 137, 139, 140, 141, 143, 155, 172, 173, 181 measures, 32, 40 mechanical energy, 15 mechanical properties, xii, 36, 61, 69, 78, 96, 152, 196, 206 media, xiii, 5, 12, 99, 123, 124, 126, 155, 172 melt, 196, 206, 207, 208, 210 melting, 196, 197, 201, 208 melts, 210 membranes, 46, 82 memory, xiii, 41, 47, 49, 50, 123, 125, 126, 134 Merck, 106 metal hydroxides, 110 metal ions, 14, 15, 24, 46 metal oxide, xi, 1, 2, 46, 49, 50, 55, 57, 103, 104, 164 metal oxides, xi, 1, 2, 46, 55, 57, 103, 104, 164 metal salts, 8, 15 metals, 24, 42, 85, 147, 148, 163, 180, 210 methane, 152, 157, 161, 166, 167, 170 methanol, 11, 102, 103 micelles, 16 microemulsion, 6, 15, 17, 103 micrometer, 15, 34, 189 microorganisms, 18 microparticles, 185 microscope, 20, 25, 28, 32 microscopy, xiii, 17, 18, 20, 22, 32, 41, 95, 98, 99, 108 microspheres, 10 microstructure, xiii, 123, 125, 134, 196, 202, 204 microstructures, xii, 1, 3, 14, 19, 53, 180 microwave, 6, 14, 44, 103, 110 microwave heating, 14 microwave radiation, 14 microwaves, 14 migration, 36 miniaturization, xi, 1, 2, 3, 41, 53 mining, 152

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Index

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mirror, 185 mixing, 4, 7, 48, 100, 104, 106, 107 mobility, xii, 1, 3, 16, 19, 116, 134, 189 modeling, 9, 144 modern society, 179 modulation, 50, 51 moisture, 35 molar ratio, 12, 101, 106, 153, 155, 156, 157, 158, 159, 160, 161, 164, 165 molar ratios, 158, 161, 164, 165 mole, 14, 101, 153, 157, 158, 159, 160, 161, 184 molecular structure, xi, 35 molecules, 7, 15, 19, 34, 35, 189 molybdenum, 71, 148 monochromatic light, 35 monochromator, 35 monoenergetic, 21, 33 monolayer, 37, 117 monolayers, 37 monomeric, 101 morphological, 7 morphology, 2, 3, 4, 5, 9, 12, 15, 16, 17, 20, 27, 32, 72, 76, 79, 82, 83, 88, 90, 91, 107, 131, 153, 158, 166 Mössbauer effect, 40 motion, 40, 155, 172, 180 motivation, 41 movement, 133, 136 MRS, 66, 193 multiplicity, 36

N nanoclusters, xi nanocomposites, 45, 102 nanocrystal, xv, 7, 22, 28, 36, 46, 195 nanocrystalline, 4, 5, 10, 11, 17, 19, 23, 26, 37, 39, 40, 41, 43, 45, 46, 79, 81, 87, 95 nanocrystals, xv, 8, 22, 25, 26, 27, 31, 33, 34, 45, 46, 195, 208 nanodevices, 19 nanolayers, 46 nanomaterials, 18, 20, 36, 61, 180, 209 nanomedicine, 94 nanometer, 11, 19, 32, 38, 44, 53, 72, 82 nanometer scale, 11, 53, 82 nanometers, xi, 2, 7, 21, 26, 30, 70 nanoparticles, xi, xv, 2, 5, 11, 12, 13, 15, 16, 17, 18, 19, 22, 23, 24, 26, 27, 32, 41, 43, 44, 45, 46, 52, 53, 64, 94, 107, 116, 181, 184, 195, 209 nanoreactors, 15 nanostructured materials, 7, 72 nanostructures, 37, 180, 182, 183

219

nanosystems, 209 nanotechnologies, 196 nanotechnology, xi nanotube, 71, 72, 82, 88, 89, 90, 91, 92 nanotubes, xii, 70, 82, 83, 86, 87, 88, 89, 90, 91, 92, 94, 95, 96, 209 NASA, 119 NATO, 66 Nb, 2 Nd, 42, 45 needles, 75, 77, 78 neglect, 39 neodymium, 39 network, 15, 105 next generation, 2, 48, 50, 53 nickel (Ni), xiv, 99, 147, 148 niobium, 71 nitrate, 43, 74, 104 nitrates, 52 nitric acid, 106 nitride, 50, 104 nitrogen, xiii, 94, 98, 99, 101, 107 nitrogen gas, xiii, 98, 99, 107 NMR, 99, 102, 109, 112 noise, 31, 40 non-crystalline, 7 non-destructive, 49 normal, 10, 80, 107, 172 normal conditions, 10 normal distribution, 107 novel materials, xiii, 98 NSC, 144 n-type, 43 nuclear, xiii, 40, 98, 99, 109 nuclear magnetic resonance, xiii, 98, 99, 109 nucleation, xiv, 7, 16, 74, 77, 78, 80, 83, 84, 88, 89, 95, 101, 103, 124, 132, 133, 141, 142 nuclei, 19, 40, 77, 103, 141 nucleus, 19, 78, 132

O observations, xiv, 9, 124, 138, 143 oil, 15, 102 optical, xiii, 3, 28, 33, 36, 39, 45, 46, 47, 52, 53, 123, 124, 136, 196, 208 optical properties, 40, 45, 46, 124, 136 optics, xv, 196 optimization, 9 oral, 70 organic, 6, 15, 98 organic compounds, 98 organism, 44

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220

Index

orientation, 29, 47, 52 orthopaedic, 90 orthorhombic, 31, 36 oscillations, 30, 52 osteoblastic cells, 81 osteoblasts, xii, 69, 80, 87, 92 oxalate, 8 oxidants, 101 oxidation, xii, 8, 70, 72, 82, 101 oxide, xii, 2, 8, 9, 14, 15, 19, 20, 32, 38, 40, 42, 43, 48, 50, 51, 69, 70, 71, 72, 82, 83, 84, 85, 86, 87, 88, 90, 92, 93, 95, 96, 170, 174, 209 oxide nanoparticles, 19 oxide thickness, 72 oxides, xi, 1, 2, 3, 4, 14, 16, 39, 40, 42, 43, 44, 45, 46, 49, 50, 51, 55, 57, 71, 82, 86, 89, 90, 92, 94, 95, 99, 103, 104, 119, 164, 170 oxygen, xii, 1, 2, 3, 5, 7, 10, 19, 28, 29, 32, 33, 35, 36, 38, 42, 43, 45, 46, 73, 82, 85, 101, 162, 182

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P parabolic, 141, 142, 143 paramagnetic, 6 parameter, xiii, 8, 9, 12, 44, 72, 83, 123, 126, 141, 185, 203 particle morphology, 9 particle nucleation, 103 particle shape, 11, 48 passivation, 84 passive, 83, 101 patents, 7 pathways, 90 Pb, 2, 4, 7, 45, 50, 52 PCs, 51 peptide, 19, 42, 43 peptides, 19 periodic, 27, 39, 81, 94 periodicity, 51 permeability, 187 perovskite, xi, 1, 2, 3, 4, 5, 6, 7, 8, 9, 10, 11, 12, 14, 15, 16, 18, 19, 20, 21, 23, 24, 26, 28, 29, 30, 32, 33, 35, 36, 38, 39, 41, 42, 43, 44, 45, 46, 47, 48, 49, 50, 51, 52, 53, 64 perovskite oxide, 38, 42, 43, 46, 50, 51 perovskites, 13, 38, 40, 50, 52 perturbations, 84 pH, 6, 9, 17, 70, 71, 73, 74, 78, 79, 81, 88, 89, 93, 94, 95, 101, 102, 103, 104, 105, 106, 113, 116, 117, 118 pH values, 71, 104, 116 phase boundaries, 26, 208 phase diagram, 9, 124, 125, 164

phase transformation, xii, 2, 35, 124, 126 phase transitions, 31, 36, 47 phonon, 35, 36 phonons, 35, 36 phosphate, 70, 72, 73, 74, 75, 76, 77, 78, 79, 80, 88, 89, 90, 92, 93, 94, 95 phosphates, 78, 79, 96 phosphor, 39 phosphorus, xii, 69, 73, 74 photoabsorption, 30 photocatalysis, 94 photoemission, 33 photographs, 162, 163, 166, 169, 173, 174 photoluminescence, 39, 45 photoluminescence spectra, 45 photon, 20 photonic, 51, 52 photonic crystals, 51 photonics, 54 photons, 30, 35, 51 photovoltaic, 72, 82 physical properties, xi, xii, xv, 69, 147, 195 physicochemical, xiii, 5, 98, 99, 102, 106 physicochemical properties, xiii, 5, 98, 99, 106 physicochemistry, 95 physics, 35, 119 physiological, 71, 73, 79, 94 piezoelectric, 2, 47 piezoelectricity, xi, 1, 2 planar, 26, 141 plasma, xiv, 19, 88, 93, 103, 110, 127, 147, 149, 163, 164, 165, 166, 167, 168, 169, 170, 171, 172, 173, 174, 175, 176, 181 platinum, 209 play, 6, 45, 70, 82, 87, 102, 142 point defects, 10, 134 polarity, 86 polarizability, 35 polarization, 28, 29, 41, 43, 46, 47, 48, 73, 74, 78, 84, 87 pollutants, 98 pollution, 9, 48 polycrystalline, 39, 44, 124, 125, 134 polyester, 153, 164 polymer, xi polymer matrix, xi polymeric materials, 101 polymerization, 7, 189 polymers, 105 polystyrene, 51 polyvinyl alcohol, 181, 182, 183 poor, xii, 1, 3, 6, 40

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Index pore, 19, 21, 82, 83, 84, 85, 86, 98, 107, 116, 150, 153, 164, 206 pores, 10, 11, 20, 21, 22, 82, 83, 85, 86, 89, 91, 98, 107, 183, 198 porosity, 11, 70, 103 porous, 14, 72, 73, 82, 83, 84, 94, 96, 107, 110 porous materials, 14 potassium, 17, 184, 185, 186, 187, 188 power, xiv, 47, 49, 125, 126, 136, 164, 165, 166, 167, 169, 171, 172, 173, 174, 175, 179, 183, 191 precipitation, 6, 7, 9, 15, 16, 71, 101, 103, 104, 105, 143 pre-existing, 141 pressure, xv, 9, 12, 19, 73, 126, 150, 164, 181, 183, 187, 189, 190, 191, 195, 196, 197, 198, 199, 201, 203, 206, 210 probe, 18, 20, 24, 28, 30, 32, 33, 36, 42, 45, 133, 198, 206 production, xiv, xv, 5, 12, 15, 16, 44, 99, 100, 101, 143, 148, 174, 179, 180, 181, 191, 192, 195, 196, 207, 209 proliferation, 81, 89, 90 proportionality, 32 prosthesis, 94 protection, 51 protein, 82, 93 proteins, 72, 80, 90 protons, 11 prototype, 2 p-type, 133 pulse, 32, 38, 49, 51, 81, 94 pulsed laser, 19, 93, 208 pulsed laser deposition, 208 pulses, 124 pure water, 106, 107, 111, 112 purification, 100 pyrolysis, 6, 8, 15, 30, 102

Q quadrupole, 37 quantitative estimation, 204 quantum, xiii, 33, 39, 45, 98, 99, 180 quantum chemical calculations, 39 quantum confinement, 39 quantum objects, 33 quartz, 183

R R&D, 28 radiation, xv, 14, 30, 31, 40, 45, 60, 127, 130, 196

221

radiation detectors, xv, 196 radio, 126 rain, 208 Raman scattering, 35, 36, 46 Raman spectra, 36, 38 Raman spectroscopy, 9, 35, 36, 61 random, xi, xiii, 1, 2, 45, 47, 49, 50, 123, 125, 130, 136, 141 random access, xi, xiii, 1, 2, 49, 50, 123, 125 range, xi, xii, 2, 4, 14, 15, 20, 31, 33, 34, 35, 36, 39, 40, 41, 42, 44, 45, 52, 53, 71, 73, 78, 79, 81, 83, 84, 101, 117, 124, 136, 140, 148, 150, 152, 158, 163, 167, 173, 189, 197, 198, 207, 208 rare earth, 45 rat, 93 raw materials, 16, 103, 147, 179 Rayleigh, 35 reactant, 5, 14, 16, 157, 171 reactants, xiv, 16, 147, 149, 157, 158, 163 reaction mechanism, xii, 2, 5, 8, 9 reaction order, 136, 141 reaction rate, 4, 9, 105 reaction temperature, 3, 8, 12, 14, 16, 101, 103, 154, 156, 157, 159, 160 reaction time, 10, 14, 73, 101, 106, 107, 110, 116, 167 reactivity, 5, 42 reading, 49 reagent, 9 reagents, 9, 101, 110 recombination, 39, 45 recovery, 10 recrystallization, 199, 206, 208 recrystallized, 198 redox, 74 refining, 31 reflection, 198 reflectivity, 134 refractive index, 51 refractory, 19, 152 regenerate, 101 regenerated cellulose, 96 regression, 118 regular, 17, 35, 51 relationship, xiv, 123, 125, 126, 129, 130 relationships, xiii, 98, 99, 116 reliability, 32 reparation, 95, 96, 103 resistance, xi, xii, xiv, 50, 69, 71, 73, 82, 87, 102, 107, 123, 124, 125, 126, 133, 134, 136, 137, 138, 139, 140, 142, 143, 144, 152 resistive, 125, 133, 134, 143

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222

Index

resistivity, xiii, xv, 123, 124, 125, 133, 134, 196, 198, 206, 208 resolution, 9, 16, 17, 20, 23, 24, 25, 26, 27, 28, 31, 33, 34, 36, 38, 39, 131, 132, 138 resources, 179, 180, 181 response time, 49 retention, 49 rhombohedral, 31, 37, 128, 131, 142 rings, 11, 138, 183 rods, 88 room temperature, xiv, xv, 10, 15, 16, 17, 18, 19, 22, 31, 41, 42, 44, 45, 47, 53, 67, 71, 73, 74, 75, 93, 106, 124, 125, 135, 137, 184, 189, 195, 196, 198, 203, 208 room-temperature, 6, 17, 18, 19, 45, 47, 53 roughness, 70, 72, 82, 88 Russian Academy of Sciences, 195 rutile, 5, 86, 87, 148

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S SAD, 127, 128, 130, 138 salt, 6, 9, 15, 102 salts, 8, 9, 16, 101, 104, 106 sample, 20, 21, 22, 24, 25, 30, 31, 32, 33, 34, 35, 37, 38, 40, 43, 83, 88, 110, 163, 164, 198, 204 saturation, 100, 101, 111, 112, 181, 183, 184, 190 scanning electron microscopy, xiii, 20, 98, 99 scanning electronic microscope, 181 scattering, 16, 17, 21, 24, 25, 32, 35, 36, 46, 134, 153, 207 Schmid, 65 search, 53, 149, 179 segregation, 7 selected area electron diffraction, 13, 23, 127, 128 selecting, 53 selectivity, 43, 48 self-assembly, 52, 103 SEM, 20, 32, 75, 80, 89, 90, 91, 99, 107, 108, 150, 151, 159, 172, 175 SEM micrographs, 90 semiconductor, xi, 1, 2, 43, 49, 50, 133, 180, 188 semiconductors, 43 sensing, 43, 64, 95 sensitivity, 36, 38, 43, 48 sensors, xi, 1, 2, 44, 48, 72, 82 series, 21, 117, 119 serum, 80 shape, 2, 5, 7, 11, 12, 17, 30, 34, 36, 48, 73, 79, 101, 107, 151, 187 shock, 14, 187 short period, xiv, 14, 124 Si3N4, 49, 50, 104

Siemens, 150, 153, 198 sigmoid, 138 sign, 85, 99 signals, 20, 21, 31, 32, 33 signal-to-noise ratio, 40 signs, 83 silanol groups, 110 silica, 45, 101, 102, 105, 106, 107, 109, 110, 111, 183, 196, 203 silicate, 102, 106, 107, 111, 112, 187 silicon, xiv, 32, 46, 49, 104, 179, 180, 181, 182, 183, 184, 185, 186, 187, 188, 189, 190, 191, 192 silicon dioxide, 182 silver, 92, 126 simulated body fluid (SBF), 88, 94, 96 simulation, 28, 39 simulations, 141 single crystals, xv, 53, 195, 196, 206, 207, 208 single-crystalline, 19, 44, 181, 183, 188 sintering, 4, 6, 7, 11, 33, 38, 73, 102, 103, 105, 161 SiO2, 49, 50, 51, 101, 105, 106, 107, 112, 183 sites, 10, 24, 35, 42, 78, 81, 83, 88, 128, 142 sodium, 88, 89, 102, 106, 107, 111, 112, 184, 185, 187, 193 software, 150, 154 solar, 72, 82 sol-gel, 6, 7, 15, 28, 45, 72, 100, 102, 103, 104, 105, 106 solid oxide fuel cells, 38, 42, 43, 48 solid phase, xiv, 24, 147, 176 solid state, 32, 48 solid surfaces, 37 solidification, 164 solid-state, xii, 1, 3, 4, 6, 7, 8, 9, 53 solubility, 120 solvent, 6, 11, 12, 15, 102, 103, 104, 105 solvents, 11, 34, 52, 103 solvothermal synthesis, 11 sorbents, 191 sorption, 180, 182, 183, 184, 189 191 spatial, 38 species, 6, 8, 9, 15, 25, 32, 37, 38, 78, 100 specific surface, xiii, 72, 98, 99, 103, 107, 109, 147 spectroscopy, xiii, 30, 31, 32, 34, 35, 36, 37, 39, 40, 98, 99, 127 spectrum, xi, 1, 2, 30, 32, 34, 37, 40, 79, 207 speed, xiii, 123, 125, 127, 183, 184, 185, 186, 187, 188, 190 spin, 44, 46, 52, 109 spintronic devices, xi, 1, 2 sputtering, xiii, 19, 34, 37, 50, 123, 126, 127 SQUID, 99, 111

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Index stability, xiv, 43, 52, 71, 82, 102, 116, 124, 125, 126, 135, 180 stabilization, 10, 18, 43 stages, 4, 77, 102, 104, 105, 137, 190, 209 stainless steel, 164 standards, xv, 33, 37, 109, 195 statistics, 90 steady state, 84 steel, 164 STEM, 23, 24, 25, 28, 32 steric, 101 stoichiometry, 2, 3, 4, 6, 15, 16, 85 storage, xiv, 49, 50, 54, 79, 124, 149, 151, 179, 180, 181, 191, 193 strain, 22, 29, 30, 36, 44, 132, 199 strains, 22, 30, 35 streams, 98 strength, 43, 70, 71, 73, 85, 88, 101, 113, 114, 115, 116, 117, 152 stress, xii, 69, 85, 88 stretching, 35, 79 strong interaction, 99 strontium, 38, 50, 92 structural changes, 40 structural defects, 10, 22, 26, 35 substitution, 39, 43 substrates, 48, 70, 71, 73, 75, 78, 94, 126, 201 subtraction, 32 success rate, 70 sulfate, 117 superconducting, xiii, 16, 98, 99 superconductivity, 148 supercritical, 105 supply, 10, 49, 126, 136, 164, 185 suppression, 44, 49, 181 surface area, 44, 72, 78, 82, 84, 90, 107, 149, 172, 173 surface chemistry, 33 surface energy, 18, 132 surface layer, 4, 5 surface modification, 70, 72 surface properties, xi, 182 surface region, 4, 20 surface roughness, 88 surface structure, 27, 149 surface tension, 19 surface treatment, 93 surfactant, 15 surfactants, 102 surgical, 93 surplus, 185 susceptibility, 99, 111 switching, 47, 49, 52, 125

223

symbols, 29, 118 symmetry, 10, 36, 39, 46, 47, 198 synchrotron, 30, 36 synchrotron radiation, 30

T tantalum, 50 targets, 126, 127 TCR, xiii, 123, 133, 134, 137 teflon, 153, 164 TEM, xiv, 11, 13, 16, 17, 20, 21, 22, 23, 24, 25, 32, 33, 34, 75, 77, 99, 106, 107, 108, 124, 127, 128, 130, 131, 132, 133, 138, 142, 143, 153, 155, 156, 162, 163, 166, 169, 171, 173, 174, 197 TEOS, 102, 103, 106, 107 terbium, 45 terraces, 27 thermal activation, xiv, 124, 137 thermal analysis, 128 thermal decomposition, 8, 170 thermal energy, 14, 99, 184, 187 thermal evaporation, 19, 197 thermal plasma, 163, 164, 166, 170, 171, 175 thermal treatment, 8, 15, 19 thermodynamic, 9, 71, 150 thermodynamic calculations, 150 thermodynamic stability, 71 thermodynamics, 9 thin film, 19, 37, 50, 67, 84, 93, 142, 144, 208 thin films, 19, 37, 50, 67, 93, 142, 144, 208 three-dimensional, 2, 7, 8, 15, 21, 51, 76, 77, 78, 80 tin, 71 tissue, xii, 69, 70, 71, 94 titania, 82, 87, 88, 89, 93, 95, 96 titanium, xii, xiv, 6, 8, 9, 10, 12, 14, 16, 35, 69, 70, 71, 72, 73, 74, 75, 78, 81, 82, 83, 84, 86, 87, 89, 90, 91, 92, 93, 94, 95, 96, 147, 148, 149, 177 titanium alkoxide, 9 titanium isopropoxide, 6 titanium oxide, 90, 91 titration, 114, 115, 116 topological, 90 toxicity, 51 transfer, 39, 86, 98, 116, 180 transformation, xii, xiii, 2, 7, 10, 35, 94, 123, 124, 126, 135, 140, 141, 143, 184, 189, 203, 208 transformations, 141 transistor, 49, 50 transition, xi, xiv, xv, 1, 2, 8, 16, 18, 22, 33, 36, 39, 41, 42, 44, 46, 71, 124, 133, 143, 189, 190, 195, 202, 206 transition elements, 71

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224

Index

transition metal, 2, 42, 46 transition metal ions, 46 transition temperature, xiv, 41, 44, 124, 133, 143, 190 transitions, 37, 39, 40 transmission, xiii, xiv, 9, 20, 22, 23, 24, 25, 26, 27, 98, 99, 124, 198, 207, 208 transmission electron microscopy, xiii, xiv, 9, 20, 22, 24, 25, 27, 98, 99, 124 transmittance spectra, 207 transport, xi, 42, 78, 98, 180 transportation, xiv, 85, 86, 179, 191 tubular, xii, xiv, 69, 82, 83, 84, 85, 86, 90, 96, 147, 149, 162, 163, 164, 176 tungsten, xiv, 147, 148, 149, 152, 157, 158, 159, 160, 161, 164, 166, 167, 170, 173, 174, 175, 209 tungsten carbide, xiv, 147, 148, 149, 152, 157, 158, 159, 160, 161, 164, 166, 167, 170, 173, 174, 209 twins, 202, 204 two-dimensional (2D), 51, 76, 78, 80, 85

vibration, 36, 78 vibrational modes, 34, 36 Vickers hardness, 198 viscosity, 83, 209 visible, 24, 35, 39, 52, 82, 84, 164 voiding, 12 voids, 52, 82, 85 vortex, 181

W wastes, 181, 188 wastewater treatment, 98, 105, 118 water, xiv, 7, 8, 12, 15, 33, 35, 83, 87, 96, 98, 102, 103, 105, 106, 107, 111, 112, 118, 164, 179, 180, 181, 182, 183, 184, 185, 186, 187, 189, 191, 198 wave number, 36 wavelengths, 34, 207 wear, 102, 152 worms, 12 writing, 47

U ultra-fine, 2, 4, 98 ultra-thin, 32 ultraviolet (UV), 35, 39, 45, 54 uniform, 2, 11, 16, 22, 44, 48, 82, 85, 107, 196 urea, 104

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V vacancies, xii, 1, 3, 10, 23, 28, 35, 36, 45, 85, 86 vacuum, 8, 21, 34, 37, 38, 73, 150, 181, 182, 183 valence, 28, 33, 134, 137 values, 28, 71, 84, 99, 104, 107, 111, 116, 117, 207 van der Waals, 19 vanadium, 71 vapor, xiv, xv, 19, 33, 103, 147, 148, 149, 152, 170, 175, 195, 196, 197, 201, 208 vapor phase deposition, xv, 195 variables, 52, 157 variation, 22, 85, 126, 130, 136, 141, 143 vector, 47, 100, 183 vehicles, 180 velocity, 40, 196 versatility, 163 vessels, 180

X X-ray absorption, 30 x-ray diffraction, 110 X-ray diffraction (XRD), 23, 29, 35, 46, 75, 81, 86, 87, 99, 106, 109, 110, 124, 127, 128, 129, 150, 151, 152, 153, 154, 155, 156, 157, 162, 164, 165, 168, 169, 170, 171, 173, 174, 175, 176, 177 X-ray photoelectron spectroscopy (XPS), 33 x-rays, 29, 30, 31, 32, 33, 181

Y yield, 9, 16, 19, 86, 98, 129, 166 YSZ, 38

Z zeta potential, 113, 116, 117, 153 zirconia, 38, 105, 106, 109 zirconium, 105, 106, 118, 181 Zn, xv, 105, 195, 207

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